If?.. aaaaSg iiriB S w wl M i cornel, university Library "•^ ''°""\n introduction tomesjuav Metaliurgvian "troau „.„„i,,iiiii;iili™i CORNELL UNIVERSITY LIBRARY Bequest of Prof, w'jldar D. Bancroft ENGINEERING llBt^^^"^ Cornell University Library The original of tiiis book is in tine Cornell University Library. There are no known copyright restrictions in the United States on the use of the text. http://www.archive.org/details/cu31924004614859 AN INTKODUCTION TO THE STUDY OF PHYSICAL METALLUKGY METALLUEGY A TEEATISB XTNDER THE GENEEAL EDITOESHIP OF WALTEE EOSENHAIN, B.A., D.SC, F.E.S. CoMros;TE Photo-Microcraph op Comtui ;-Si-:cTio\ OF A CAsi:-HvRDENF.n Steel Rod. niACixiFiCATioN 20 diameters. METALLURGY AN INTRODUCTION TO THE STUDY OF PHYSICAL METALLURGY BY WALTER ROSENHAIiN, B.A., D.Sc, F.R.S. SUPERINTENDENT METALLURGY DEPARTMENT OF THE NATIONAL PHYSICAL LABORATORY, NEW YORK D. VAN NOSTRAND CO TWPiNTY-mVE PARK PLACE 1914 PEEFACE The present volume is intended to serve as an introduction to the subject of Physical Metallurgy for those who are in- terested in the physics and physical chemistry of metals, whether as scientific investigators, manufacturers or users of metals, or students of metallurgy, engineering and allied subjects in which a knowledge of metals plays an important part. The volume also serves as an introduction to the Metallurgical Series which is in course of publication by Messrs. Constable under the editorship of the author. In the various volumes of this series will be found that more detailed treatment of many of the principal subjects touched upon in this book which is required by those particularly interested in any one metal or alloy or in some special aspect of the subject. With these more detailed volumes in view, the treatment of the whole subject in the present work has been intentionally kept some- what general, the object of the author being to awaken interest and to stimulate thought and ideas rather than to communicate a great mass of detailed data. In writing such a book the author has necessarily relied very largely upon the work of others, and in the tables of references at the end of each chapter he has endeavoured to make the only real acknowledgement which is possible in such circumstances. These references lay no claim to being an exhaustive bibliography of the subjects touched upon, but they constitute a fair guide to introductory reading. In regard to illustrations also, the author is indebted to many others for the examples he has chosen. As far as possible, the micrographs particularly have been taken from the author's own work or from that of members of the staff of the Metallurgy Department at the National Physical Laboratory. Many of these have vi . PREFACE already appeared as illustrations of the author's various papers (many of them joint papers), and he is indebted to the Councils of the various Societies and Institutions concerned for per- mission to reproduce these illustrations here. Specific acknowledgment in regard to these illustrations is made below. Photo-micrographs and diagrams by other authors have been very kindly placed at the present author's disposal in a number of cases, and he desires to express his indebtedness to the following : — Sir J. Alfred Ewing, F.R.S., for Figs. 110 to 113 inclusive, and for permission to reproduce a number of illustrations first pubUshed in joint papers by Sir Alfred Ewing and the present author. Dr. J. E. Stead, F.R.S., for the micrograph. Fig. 17. Professor W. E. Dalby, F.R.S., for the diagrams. Figs. 87 and 88. The author also desires to record his indebtedness to Dr. R. T. Glazebrook, C.B., F.R.S., for permission to reproduce here a number of photographs of apparatus and apphances installed at the National Physical Laboratory and a large number of photo-micrographs taken there in the course of the regular work of the Laboratory. Particular thanks are also due to various members of the staff of the Metallurgy Department, who have rendered most valuable assistance in the preparation of the illustrations, particularly to Messrs. W. H. Withey, D. Ewen, S. L. Archbutt and J. L. Haughton. To the two last-named the author is also indebted for much valuable help in the reading of proofs. By kind permission of the Councils of the various Societies and Institutions concerned, Figs. 18, 19, 23, 24, 25, 26, 46, 51, 52, 53, 100, 106, 110, 111, 112, 113, 117, are reproduced from various papers in the Philosophical Transactions and Pro- ceedings of the Royal Society of London ; Figs. 99, 101, 109, 114, 116, from the Journal of the Iron and Steel Institute ; Figs. 13, 15, 72, 73, 75, 96, 119, from the Proceedings of the Institution of Mechanical Engineers ; Figs. 97, 98, 107 from the Journal of the Institute of Metals ; and Fig. 133 from the Journal of the Staffordshire Iron and Steel Institute. TABLE OF CONTENTS CHAPTER I Intboductokt Definition of " Physical Metallurgy " Scope of Physical Metallurgy . Relation to Pure Science Practical Importance Achievements due to recent Advances in the Knowledge of Metals Importance of Physical Metallurgy to Makers of Metals Control of Manufacturing Processes Importance to Users ... Control of Materials Systematic Testing and Investigation and Bad Eesults . Bapid Growth of Physical Metallurgy Early History of Metallography Early History of Mechanical Testing References ..... of Causes of both Good fAOB 1 1 2 3 4 4 4 5 5 6 10 12 13 15 PART I The Stetjctttee akd Constitution of Metals and Alloys CHAPTER ri The Microscopic Examination of Metals ImpracticabiUty of Transparent (thin) Sections Preparation of Specimens for the Microscope Choice of Specimens Cutting, Grinding and Fine Grinding . Polishing ..... The Physical Nature of Polishing and Grinding Operations Formation of Amorphous Surface Layers . 17 18 18 20 23 24 25 vm CONTENTS Development of the Structure Etching with Weak Solvents Polish Attack ...... Heat-Tinting and Selective Electro-Deposition Preservation of Specimens .... Beferences ....... CHAPTER III The Metallubgical Microscope Adaptation of the Microscope for Metallurgy Illumination of Metal Surfaces Oblique and Normal Illumination Simple Metallurgical Microscope Special Metallurgical Microscopes Eosenhain Microscope Optical System of the Microscope Monochromatic Light .... Stopping down ..... Aperture and Depth of Focus . Curvature of Field .... Diffraction and Eesolving Power Limit of Useful Magnification . Rules for the Use of the Microscope Sources of Illumination .... Critical Illumination Intense Illumination for Photography Projection Apparatus and Photography . Binocular Stereoscopic Microscope (Greenough) Work Accessories ...... Mounting and Levelling Devices References ...... for Metallurgical CHAPTER IV The Miceo-Steuctuee of Ptjee Metals and of Allots Typical Structure of a Pure Metal . Formation of a CrvstalUne Aggregate Etch Figures and Oriented Lustre of Crystals Freezing of a Liquid, and Growth of Dendrites Crystal Boundaries or Junctions Interlocking at Crystal Boundaries . Crystallographic Characters of Metals Structure of Alloys : — Mutual Solubility of most Liquid Metals Alloys which form Solid Solutions Dendritic Cores in Solid Solutions 30 30 33 34 36 37 38 38 38 41 43 43 44 46 47 48 50 51 52 53 54 54 56 58 59 59 59 60 61 62 66 67 69 71 71 72 73 74 CONTENTS ix Structure of Alloys — continued. page Alloys which do not form Solid Solutions .... 74 Eutectiferous Alloys . . . . . . .76 Eeferences .......... 77 CHAPTEE V The Thermal Study of Metals and Allots Thermal Data as Basis for Constitutional Diagrams ... 78 Methods of Heating and Cooling Specimens . . . .78 Bosenhain Vertical Tube Furnace with Temperature Gradient 79 Methods of Temperature Measurement Thermo-Couples Cold- Junction Arrangements Calibration Galvanometers and Potentiometers . Heating and Cooling Curves . Inverse Kate Curves Differential and Derived Differential Curves Typical Thermal Curves of — Pure Metal ....... Solid Solution ...... Eutectiferous Alloy ...... Correlation of Thermal Curves with Micro-Structures Thermal Curves and " Freezing-Point " Curves Constitutional Diagram and Thermal Curves Equilibrium — Stable and Meta-Stable Conditions Definition of " Liquidus" and "Solidus" Constitutional Diagram of the Ideal Simple Eutectiferous System 91 Constitutional Diagram of a System forming Solid Solutions . 97 The Copper-Nickel System 98 Constitutional Diagram of System forming Solid Solutions, with a EuteetiterouB Eange ....... 99 The Copper-SUver System 100 Constitutional Diagram of System in which a Definite Compound is formed 100 The Magnesium-Tin System ...... 101 Thermal Phenomena in Solid Metals and Alloys . . . 102 Correlation with Micro-Structure by Quenching. . . 102 Kosenhain Quenching Apparatus .... 102 Thermal Analysis ......... 103 Application to Determination of Eutectio Point . . . 104 Micro -Structure and the Constitutional Diagram . . , 105 Determination of Limits of Solid Solubility . . . 105 Determination of Composition of a Compoimd . . . 104 Method of Prolonged Annealing, with and without Quenching . 105 Determination of the " /SoWdits " ..... 106 Eeferences 106 80 80 81 81 82 83 83 84 85 96 88 89 91 91 93 94 CONTENl:S CHAPTER VI The Constitutional Diagkam and the Phtsical Pkopertibs of Allots FAOB Correlation between Constitutional Diagram and Physical Proper- ties 107 As a Guide to the Utmty of AUoys 107 As a Means of Studying the Constitution of Alloys . . 107 Constitution and Specific Volume or Density of Alloys . . 108 Magnetic Properties of Alloys ...... 109 Electrical Conductivity . . . . . • .110 Curves of Specific Conductivity for Alloy Systems. . Ill Ideal Eutectiterous System . . . . .112 Solid Solution System 112 Intermediate, partly Eutectiferous System . .112 Systems containing Definite Compounds . . 114 The Copper- Antimony System . . .114 Temperature Co -efficient of Conductivity . . . .115 Experimental Difficulties of Electrical Methods of Studying Alloy Constitution ......... 116 Constitution and Other Physical Properties . . . .116 Electro -Chemical Properties . . . . . .117 Concentration-Potential Curves . . . . .117 Study of Alloy Constitution by the Method of Residue Analysis . 118 Difficulties of the Method 118 Abel's Discovery of " Cementite " . . . . .119 Constitutional Diagrams of Alloys in the Light of the Phase Rule 119 Formulae of the Phase Rule ...... 120 Meta-Stable Conditions and the Limitations in the Applica- tion of the Phase Rule . . . . . . .121 True Equilibrium and its Relation to Practical Conditions . . 122 Ternary Alloy Systems ........ 123 Graphical Representation . . . . . .124 The Ternary Base-Triangle and the Constitutional Model . 126 References .......... 128 CHAPTER VII Typical Allot Systems Simplest Type of Binary System . . . . . .131 Lead-Antimony ........ 131 Slightly more Complex . . . . . . . .132 Lead -Tin System . . . . . . . .132 Micro-Structures and Constitution according to Eosenhain and Tucker ........ 132 Mazzotto's Views ....... 132 CONTENTS xi PAGE 135 135 136 Constitutional Diagram aocordingto Eosenhain and Arohbutt 136 139 140 Ternary and Complex Alloys derived from the Lead-Tin System Type-Metals and Bearing-Metals .... Alloys of Aluminium and Zinc Micro - Structures Alloys of Zinc and Copper Constitutional Diagram according to Shepherd, with Modifi cation by Carpenter ...... Constitution and Micro -Structure of Alloys rich in Copper Hot and Cold Working of Brass as related to the Constitu tional Diagram ....... Alloys of Tin with Copper ...... Constitutional Diagram as modified by Hoyt Constitution and Micro -Structure of AUoys rich in Copper Effects of Heat-Treatment ..... Ternary Alloys of Zinc, Tin and Copper .... Ternary Constitutional Diagram according, to Hpyt . . Ternary Alloys containing Phosphorus .... Constitution and Micro-Structure .... Ternary Alloys containing Manganese .... Alloys of Aluminium with Copper ..... Constitution and Structure of the Alloys rich in Copper Constitution and Structure of the Alloys rich in Aluminium Ternary Alloys of Aluminium and Copper with Tin and with Man ganese ......... Alloys rich in Aluminium, containing Magnesium References ......... CHAPTER VIII The Iron-Carbon System Constitutional Diagram — The Meta-Stable System Constitutional Diagram — The " Stable " System The Compound FesC and Limits of the Diagram . The Liquidus of the Diagram ..... The Solidus of the Iron-Carbon Diagram .... The Lower Lines of the Iron-Carbon Diagram . ■y-Iron and its Transformations ..... The Euotectoid " Pearlite " Allotropy of Iron and its Connection with the Decomposition of the y-Iron Solid Solution ..... Thermal Curves of Pure Iron after Burgess. AUotropic Changes in Iron related to Changes in Hardness Thermal Changes in Steels of Various Carbon Content Quenching Experiments on Steel ..... Hot Etching of Iron and Steel . . 155 156 157 158 158 160 162 162 162 163 164 164 165 166 168 167 171 173 174 xu CONTENTS Critical Points of Steel ....... Definition of Aci, Aca, Acj, and of An, Ar^, Ars . The Points Acia, Acjs and Aru, Aras and Arua Micro -Structures of Carbon Steels when slowly Cooled Micro-Structures obtained by Quenching and Tempering . Austenite, Martensite, Troostite .... The Hardening and Tempering of Steel .... Various Theories of Hardening ..... Bearing of the Constitutional Diagram on Hardening Practice Constitution and Micro -Structure of Iron-Carbon AUoys contain ing more than 2 per cent, of Carbon Constitution and Structure of White and Grey Cast-iron Graphite and Temper Carbon .... Cast-iron as " Steel plus Graphite " plus Impurities Malleable Castings ....... Case-Hardening ....... References ......... 172 172 174 172 173 178 178 180 184 186 187 189 188 190 190 192 PART II The Pbopebties of Metals as Related Stbucttjre and Constitution. TO their CHAPTER IX The Mechanical Testing of Metals General Principles of Mechanical Testing . Difficulty of " Imitating " Practical Conditions . Fundamental Principle of Isolating and Testing a Physical Property ..... Tensile Testing. General Considerations . Principle of Machines Used .... Direct Counterpoise Machines Hydraulic Measurement Machines Dalby's Form of " Weigh Bar " Apparatus . Data from Tensile Tests : — Elastic Modulus and Elastic Limit Extensometers — Ewing Martens Stress-Strain Diagrams Compression Testing References .... Singli 193 194 195 197 198 198 200 202 203 204 204 207 211 212 CONTENTS xui CHAPTER X The Mechanical Testing of Metals — continued Torsion Testing General Considerations Machines Used Other " Pure Shearing " Tests. Punching Tests Bending and Folding Tests Ball Hardness Test Definition of " Hardness " Benedicks' Formula for Brinell Test Machines for Making the Ball-Test Martens Hardness Tester Cone Indentation Test . Other Hardness Tests Turner's " Sclerometer " . Shore's " Soleroscope " Fatigue and Alternating Stresses Bauschinger and Wohler's Work Wohler's Test. Comparison of Materials Direct Alternating Stress Tests Machines of Osborne- Eeynolds and Smith and of Stanton and Bairstow Alternating Shearing Tests Alternate Bending Tests Arnold's Machine Sankey's Machine Impact Testing. General Considerations Izod Impact Tester . Charpy Impact Testing Machine Tensile Single-Blow Impact Testing Machine (Stanton) Alternating Impact Testing Stanton's Machine . Notched Bars in Impact Testing Comparison of Results for a Series of AUoys obtained by Various Forms of Static and " Dynamic " Tests References ...... 213 214 214 215 216 217 218 219 219 220 220 222 222 222 222 224 224 225 228 228 230 231 231 231 232 233 233 234 235 236 236 238 239 CHAPTER XI The Effect of Strain on the Stkucture of Metals Elongation of Crystals when Metal is Strained .... 241 Persistence of Crystalline Structure 242 Mechanism of Plasticity in a CrystaUine Aggregate . . . 243 Shp-bands ; their Character in Various Metals . . . 243 The Nature of " Shp-Bands " 244 XIV CONTENTS Slip-Bands by "Vertical and Oblique Light and in Transverse Section 245 Disturbance of Crystalline Arrangement under Severe Strain . 246 Formation of Amorphous Layers ..... 246 Hardening of Metals by Plastic Strain .... 247 Kaising of Elastic Limit in Tension accompanied by Lowering of Elastic Limit in Compression ..... 248 Temporary Destruction of Elasticity in Iron by Plastic Strain, and its Recovery ........ 248 Twinning in Metals under Strain ...... 251 Doubts whether it occurs ....... 251 Fracture and its Mechanism under Different Conditions . . 252 Tensile, Shock and Alternating Stress Fractures . . . 252 Relation of Fracture to Micro -Structui-e .... 252 Explanation of " Fatigue " Fracture by the Formation of Local Shp-Bands ........ 253 Non-Occurrence of Crystallisation under Fatigue or Vibration 254 Crystal Boundaries as Sources of Strength in a Metal. . . 256 Mechanism of Cohesion at a Crystal Boundary . . . 257 The Amorphous Cement Theory ..... 257 Brittle Fracture of Ductile Metals at High Temperatures . 258 Behaviour of Duplex Alloys under Strain ..... 260 Carbon Steel as a Typical Duplex AUoy .... 263 Modes of Fracture of Steel ...... 264 References .......... 264 CHAPTER XII The Thermal Treatment of Metals Annealing ....... Meaning of the Term " Anneahng Temperature Softening and Re-Crystallisation Two Stages of Annealing . Annealing of Certain Metals . Lead, Tin, Cadmium and Zinc Copper, Gold . Wrought Iron and Steel . Excessive Annealing Oxidation and Decarburiflation of Steel " Gassing " of Copper Coarsening of Structure . Influence of Coarse Structure on Mechanical Properties " BaUing up " of Duplex Alloys under Prolonged Annealing Effect on Mechanical Properties .... Heat-Refining of Steel ....... Influence of the Critical Points . 265 266 267 268 269 269 270 271 272 272 272 273 274 276 277 282 279 CONTENTS XV PAQK Practical Bearing of Crystal Growth and Ke-CryataUiBation at the Critical Points of Steel 282 Burnt Steel 283 Heat Treatment of Special Steels 284 References .......... 285 CHAPTER XIII The Mechanical Treatment of Metals, Including Casting Casting in Ingots and in Moulds ...... 287 ChiU and Sand Moulds 288 Influence of Rate of Cooling ....... 288 Casting Temperature ....... 288 Over-Heating of Molten Metals 289 Contraction Stresses in Castings ..... 289 Influence of the Shape of a Casting on the Crystal Structure . 290 Radial Structure of Circular Ingots 292 Growth of Crystals at Right- Angles to Isothermal Surfaces. . 292 Isothermals of a Circular Ingot ..... 290 Isothermals and Crystals at a Re-Entrant Angle . . . 292 Hoat-Reflning of Steel Castings 294 Soundness of Ingots ........ 294 Steel Ingots 295 Methods of Producing Sound Steel Ingots . . . .296 Segregation in Ingots . . . . . . . .297 Effects on the Distribution of Impurities in the Final Products 298 Prevention of Segregation. ...... 298 Hot Working defined 299 Cold Working and Strain-Hardening ..... 300 Forging and Rolling Temperatures, "Finishing " Temperature 300 Structure of Forged or RoUed Metal 302 Longitudinal Distribution of Constituents .... 303 Effect of Annealing 304 Metals Ductile when Cold, but Brittle when Hot. Brass and German Silver ......... 305 Strain Hardness and Excessive Cold Work .... 306 Cohen's " Strain Disease " in Metals 309 " Season Cracking " and Re-crystaUisation .... 308 Cutting Operations. ........ 309 Shearing, Punching, DriUing . . . . . .310 Behaviour of Metals and Alloys under Cutting Tools . . 311 References .......... 312 CHAPTER XIV Defects and Failures in Metals and Allots Impurities in Metals and Alloys . . . . . .313 Manner of Occurren?© ....... 314 XVI CONTENTS Impurities in Metals and Alloys — continued. Typical Impurities .... Phosphorus in Steel . Arsenic in Copper Antimony and Bismuth in Copper " Slag " Enclosures in Steel Errors of Composition of Alloys Liquation and Deposition of Heavy Crystals General Tabulation of Defects in Metals and Alloys and Causes .... Failures of Metals from " Abuse " Failures of Metals from " Fatigue " Failures of Metals from " Wear " Corrosion. In Iron and Steel . Prevention In Brass Condenser Tubes In Aluminium and Light Alloys References ..... their PAGE 315 315 316 316 317 320 321 322 327 327 327 328 329 330 332 335 LIST OF ILLUSTEATIONS IN TEXT. TIQ. 3. Diagram of " Vertical " or Normal Illuminators . . 39 6. Diagram of Image Formation in the Microscope . . 44 7. Eelation of Aperture to Depth of Focus .... 48 8. Diagram of " Critical lUumination " .... 54 9. Diagram of Gas Lamp for Microscope .... 56 10. Electric Microscope Lamp ...... 56 16. Diagram of Normal and Obhque Eays faUing on an Etched Surface ......... 66 28. Cooling Curve of Pure Metal (Zinc) showing sharp Freezing- Point 87 29. Cooling Curve of Eutectiferous Alloy showing Initial Freezing and Eutectic Arrest ..... 89 30. Ideal Equilibrium (Constitutional) Diagram of a simple Binary Eutectiferous System . . . . .91 31. Diagram illustrating the Eelation between Arrest-points on Thermal Curves and the Lines of the Constitutional Diagram ......... 92 32. Coohng Curve showing the Freezing-Eange of an AUoy forming a Solid Solution ...... 97 33. Cons,titutional Diagram of the AUoys of Copper and Nickel, typical of an uninterrupted Series of Sohd Solutions . 98 34. Constitutional Diagram typical of AUoy Systems partly eutectiferous but forming Solid Solutions at each End of the Series ......... 99 35. Constitutional Diagram of the AUoys of Magnesium and Tin, typical of a Series forming a definite Inter-metaUio Compound ........ 101 37. Diagram of Thermal Analysis Curves .... 104 38. Curve of Specific Electric Conductivity typical of an unbroken Series of SoUd Solutions. (AUoys of Gold and Copper) . . . . . . . .112 39. Curve of Specific Electric Conductivity typical of a Eutecti- ferous Series. (Alloys of Cobalt and Copper) . .112 40. Diagram Ulustrating Need for numerous Observations in lajdng down Curves ....... 113 41. Curve of Specific Conductivity for the Alloys of Copper and Antimony, showing Discontinuities corresponding to definite Compounds . . . . . . .114 42. Triangular Diagram for Plotting the Composition of Ternary AUoys 124 P.M. b xviii LIST OF ILLUSTRATIONS pia. 43. lAqwi&us Surfaces for the Constitutional Model of a Ternary Alloy System. (Alloys of Lead, Tin and Antimony) . 126 45. Constitutional Diagram of the Lead-Tin Alloys . . 133 48. Constitutional Diagram of the Zinc-Aluminium Alloys . 136 49. Thermal Analysis Curves of Zinc-Aluminium Alloys . 139 50. Typical Cooling Curves of Zinc -Aluminium AUoys . . 140 54. Constitutional Diagram of the Zinc-Copper Alloys . . 141 57. Constitutional Diagram of the Tin-Copper AUoys . . 147 61. Diagram of the Constitution of the Ternary Alloys of Tin, Zinc and Copper, rich in Copper 150 62. Constitutional Diagram (tentative) of the Aluminium - Copper Alloys ........ 154 64. Constitutional Diagram of the Iron-Carbon System . . 161 65. Constitutional Diagram (tentative) of the Iron-Graphite System 162 66. Typical Inverse-Eate Heating and Cooling Curves of Pure Iron (Burgess and Crowe) ...... 168 74. Temperature-Tenacity Curve for very Soft Steel at High Temperatures . . . . . . . .179 80. Diagram of the Single Lever Tensile Testing Machine . 199 83. Diagram of the Ewing Extensometer .... 204 84. Diagram of Martens' Extensometer .... 205 85. Stress-Strain Curve showing Elastic Limit . . . 206 86. Diagram of the Dalby Apparatus ..... 208 87. Stress-Strain Curve of Mild Steel, showing Yield-Point . 210 88. Stress-Strain Curve of EoUed Brass, showing Gradual Yielding 211 89. Diagram of Shearing Test ...... 215 91. Diagram of the Wohler Test 225 92. Alternating Stress Test Curves for Two Materials . . 227 93. Diagram of Testing Machine for Eeversals of Direct Stress . 228 94. Diagram of Izod's Single-Blow Impact Testing Machine . 233 95. Diagram showing Comparative Eesults of various Tests applied to a Series of Alloys. (Zinc-Aluminium) . . 238 102. Diagram of the Formation of Slip Bands . . . 243 103. Diagram illustrating the Optical Behaviour of Slip Bands ........_ 244 121. Diagram of Isothermals in a Cylindrical Eod while Cooling 290 124. Diagram of Isothermals at a Ee-entrant Angle in a Cooling Casting 292 130. Eod passing through a Drawing-Die . . . _ 307 131. Diagramjof the Longitudinal Section of HoUow-drawn Eod or Wire 308 LIST OF PLATES. Ceoss-sbction of Case-hardened Steel Rod . Frontispieoe PAOE Plate 1 26 Fig. 1. Ordinary Polished Surface, x 150 N. „ 2. Fine-ground Surface, x 150 0. „ 13. Ferrite, Deeply Etched. X 150 N. „ 15. Ferrite, Deeply Etched. X 150 O. Plate II 42 Fig. 4. Simple Metallurgical Microscope. „ 5. Eoseuhain Metallurgical Microscope. Plate III 58 Fig. 11. Micrographic Projection Apparatus. „ 12. Leitz Metallograph. Plate IV 62 Fig. 14. Diagrams of Model of Growth of Crystalline Aggre- gate. Plate V 66 Fig. 17. Etching Figures on Silicon Steel (Stead). „ 18. Etching Figures on Tin. x 1,000 N. „ 19. Geometrical Cavities in Cadmium, x 1,000 N. Plate VI 74 Fig. 20. Dendrites in Ammonium Chloride. „ 21. Dendritic Cores in Cast Brass, x 100 N. „ 22. Steel with 0-1 per cent. Carbon, x 150 N. „ 23. Tin-Lead, containing 95 per cent. Sn. x 150 N. Plate VII 76 „ 24. Tin-Lead, containing 85 per cent. Sn. X 150 N. „ 25. Tin-Lead, containing 74 per cent. Sn. x 150 N. „ 26. Tin-Lead, containing 45 per cent. Sn. x 160 N. Plate VIII 80 Fig. 27. Vertical Tube Furnace for Thermal Curves. „ 36. Quenching Apparatus. XX LIST OF PLATES PAGE Plate IX 126 Fig. 44. Model of Liquidus Surface of Portion of the Alloys of Manganese- Alumiaium-Copper. „ 46. Lead-Tin, containing 63 per cent. Sn. x 1,000 N. „ 47. Tin-Antimony, showing Crystals of SbSn. x 150N. Plate X 139 Pig. 51. Aluminium-Zinc, showing Dendrite of AlaZng. X 150 N. „ 52. Aluminium-Zinc, Homogeneous Structure of AlaZn, (78 per cent. Zn). x 150 N. „ 63. Aluminium-Zinc, showing Eutectoid Structure from Decomposition of AlaZuj. X 600 N. Plate XI 144 Fig. 55. Brass (70 per cent. Cu), showing Twinned Structure. X 150 N. „ 56. Muntz Metal (60 per cent. Cu). x 150 N. Plate XII 149 Fig. 58. Copper-Tin (18 per cent. Sn), slowly cooled, x 150N. „ 59. Copper-Tin (18 per cent. Sn), annealed and quenched. X 150 N. „ 60. Copper-Tin (18 per cent. Sn), annealed at 450° C. X 150 N. Plate XIII Fig. 63. Manganese - Aluminium - Copper, showing Acicular Structure, x 150 N. „ 67. Steel, containing 0-2 per cent. Carbon, x 150 N. „ 68. PearHte in Carbon Steel, x 1,000 N. Plate XIV Fig. 69. Steel, containing 0-6 per cent. Carbon, x 160 N. „ 70. Steel, Quenched, showing Martensite, Troostite and Ferrite. x 1,000 N. „ 71. Steel, showing Fine Martensite. x 1,000 N. „ 72. Steel, containing 1-2 per cent. Carbon, showing Cementite and PearUte. x 150 N. Plate XV Fig. 73. Austenite-Martensite in severely Quenched High- Carbon Steel. X 1,000 N. „ 75. Troostite in Steel Quenched during Critical Rana-e X 1,000 N. * 156 175 182 LIST OF PLATES xxi PAOE Plate XVI 187 Fig. 76. White Cast Iron (ChiUed). x 150 N. „ 77. Grey Cast Iron, x 150 N. „ 78. Cast Iron showing Temper Carbon, x 150 N. „ 79. Case and Core of Case-hardened Steel, x 150 N. Plate XVII 200 Fig. 81. Single-lever Vertical Testing Machine. Plate XVIII 202 Fig. 82. Amsler Vertical Tensile Testing Machine. Plate XIX. 220 Fig. 90. Martens Direct-reading Ball-hardness Tester. „ 96. Strained (Elongated) Ferrite. x 150 N. Plate XX 242 Fig. 97. Ferrite, before Straining, x 160 N. „ 98. Same Field of Ferrite after Straining, x 150 N. Plate XXI. 243 Fig. 100. Slip-bands in Lead, x 1,000 N. „ 99. Slip-bands in Lead. X 150 N. „ 101. Slip-bands in Nickel-Steel, x 150 N. Plate XXII 245 Fig. 104. Slip-bands in Iron, x 150 N. „ 105. Slip-bands in Iron, x 150 0. „ 106. Slip-bands in Cross-section, x 1,200 N. Plate XXIII 251 Fig. 107. Twinning in Silver, x 400 N. „ 108. Slip-bands and Twinning in Copper, x 150 N. Plate XXIV 254 Figs. 110 — 113. Ferrite under Successive Stages of Fatigue. X 1,000 N. Plate XXV 256 Fig. 109. Tensile Fracture of Iron, x 400 N. „ 114. Shock Fracture of MUd Steel. X 400 N. „ 115. Shock Fracture of Defective MUd Steel, showing Inter-Crystal Weakness, x 150 N. „ 117. Tensile Fracture of Mild Steel, x 150 N. xxii LIST OF PLATES Plate XXVI 259 Fig. 116. Deformation in y-Iron at 1,000°C. X 400 N. „ 118. Shock Fracture of Brittle Wrought Iron. X ISDN. „ 119. "Balled up" Cementite in Damaged Mild Steel (Sodium Picrate Etching). X 400 N. „ 120. Over-heated Mild Steel, x 150 N. Plate XXVII 292 Fig. 122. Structure of Circular Ingot of Lead. „ 125. Eadial Crystals on Tin-plate. „ 123. Fringe Crystals at a Re-entrant Angle in Zinc. Plate XXVIII 304 Fig. 126. Longitudinally-arranged Pearlite with Equi-axed Ferrite. X 150 N. „ 127. Phosphoric Banding in MUd Steel. X 20 N. „ 128. Phosphoric Banding, showing Impression of Brinell BaU-test. x 12 N. „ 129. Same Specimen, after Annealing, x 12 N. Plate XXIX 310 Fig. 132. " Boot " of a Tool-out in Mild Steel. X 150 N. „ 133. Phosphide Eutectic in Cast Iron, x 150 N. „ 134. Iron and Sihcon in Aluminium, x 150 N. Plate XXX 318 Fig. 135. Slag Enclosure in Steel Bail, x 150 N. „ 136. Slag Enclosure in Gun Tube, x 150 N. Plate XXXI 323 Fig. 137. Forged Steel, Fine Structure, x 150 N. „ 138. Forged Steel, Coarse Insufficiently Worked Struc- ture. X 150 N. „ 139. Severe Local Cold Working in Steel, x 150 N. „ 140. Oxide Enclosure in " Burnt " Steel, x 150 N. AN INTRODUCTION TO THE STUDY OF PHYSICAL METALLURGY CHAPTER I INTRODXrCTOBY Since the term " Physical Metallurgy," which appears in the title of this book, is not one which is as yet widely used, some definition is perhaps required. Yet the term speaks largely for itself, since it obviously connotes that branch of science which deals with metals in their physical aspect. The term is justified by the need for a distinguishing name for that great branch of the knowledge of metals which has to a large extent grown up during the last fifty years — a branch which concerns itself with the nature, properties and behaviour of metals and of alloys as such, as distinct from the far older branch of metallurgy which deals with the reduction of metals from their ores. Hitherto the term " Metallurgy " has indeed been almost entirely confined to this latter meaning, and those who have grown old in this idea are a little apt to resent an innovation which gives to the old term " Metallurgy " a wider and more general meaning than it formerly bore so as to include the newer knowledge of metals. This inevitable widening of the old term, however, demands a subdivision, so that the department of metallurgy which relates to the reduction of metals and their refining may well be termed " Process " or " Chemical " Metallurgy, leaving to the younger branch of the science the newer term " Physical Metallurgy." Thus defined and understood, the scope of Physical Metal- lurgy is an exceedingly wide one, and one which brings it well over the border-land of several sister-sciences — such as chemistry on the one side, physics on another, and that branch of know- ledge generally known as "strength of materials" in yet P.M. ^ 2 STUBY OF PHYSICAL METALLUEGY another direction. Besides these, crystallography bears largely on our subiect. But while our young science necessarily draws largely upon the resources of these her elder sisters, she is not without gifts in return. Perhaps this purely scientific aspect of our subject may with advantage be dealt with first. While the greatest practical importance obviously attaches to a deeper knowledge of metals and a better understanding of the phenomena which are met with in connection with them, yet from the point of view of pure science also, the better knowledge of metals is of con- siderable importance. In the first place — ^to take the lower ground first — ^metals and alloys are in constant use for the construction of scientific instruments and apphances. In the uses to which metals are put in this connection some very extraordinary demands are frequently made upon their properties — and a deeper knowledge of those properties and of the factors which govern them must lead to a sounder method of application to the difficult problems of instrument con- struction. Only too frequently, even at the present time, those who design, make and use scientific instruments are obviously only dimly aware that anything reaUy useful is known about metals ; consequently the use of relatively unsuitable materials continues broadcast. New and valuable materials are neglected, others of doubtful utility are at times greedily employed. The contents of the present volume, although they can but touch the fringes of the subject in many direc- tions, should serve to supply a little of this much -needed knowledge to those who often are obHged to employ metals under conditions of special difficulty {^). But it is not only in connection with instruments that Physical Metallurgy has an important bearing on what may be termed scientific practice. A great many experiments, both of a physical and a chemical nature, are made upon metals. Yet in a great many cases, and until recently in all cases, the exact nature and condition of the metal employed has not been fully stated or — probably — ^understood by the experimenter. We have endless examples of determinations of physical con- stants made on metals simply described as " copper " or " soft INTRODUCTORY 3 iron," or at most we find a distinction made between metals " cast " and " wrought." A perusal of the following pages will show very clearly how inadequate and even futile such a description really is. The newer knowledge of metals which forms the basis of this volume, therefore, supplies the means for rendering precise and definite all determiaations of physical constants made on metallic materials. In view of the import- ance which is rightly attached to ever -increasing accuracy in scientific measurements, this consideration is really one of fundamental importance. There is yet another point of view from which the whole subject-matter of Physical Metallurgy is of considerable scientific importance. The study of the soUd state of matter and of molecular physics in general is attracting an increasing amount of attention, both from pure physicists and from those primarily interested in physical chemistry and crystallography. For all those interested in this subject, however, metals possess a very special importance because they are — in aU probability — the simplest form of soUd matter. This follows from the fact that they are in the pure state simple elementary substances. Accordingly it is in the study of metals that advances in our knowledge of the constitution of sohd matter are likely to be made. For that reason the results and theories of Physical Metallurgy may well be commended to the careful attention of pure physicists — an attention which they have hitherto received in only a few instances. Turning to the more immediately practical aspects of our subject, the importance of Physical Metallurgy scarcely requires either explanation or emphasis. To all the industries which are concerned with metals, and at the present time there are very few indeed which are not to a large extent thus concerned, every real advance in our knowledge of metals must be a real and important advantage. This makes itself felt most directly in the great industries which deal directly with the production and treatment of metals ready for technical or industrial use — such as the iron and steel industries and the non-ferrous metal industries. Every process which the metals imdergo, from the moment when they first leave the refining furnaces as metals, b2 4 STUDY OF PHYSICAL METALLURGY until they reach their ultimate user, forms the direct subject- matter of our science — ^for in each and all of such processes, are involved those properties of metals and those phenomena associated with them which it is the object of this science to study and to elucidate. And already there are many great achievements which stand to the credit of this comparatively young science, for although the source of many of these improve- ments in our metal products cannot be directly traced to any one scientific pubMcation or to the results of any one research laboratory in which the methods of Physical Metallurgy have been pursued, yet these things stand as the outward and visible tokens of the new knowledge of metals which has grown up in recent years. We may call to mind a few of the more striking examples, such as the alloy steels and the high-speed cutting steels, modern guns and armour-plate and armour-piercing projectiles, non-corrodible alloys for the construction of vessels for chemical manufactures, new light aUoys of remarkable strength and durabihty utilised in our air -craft, and many others which, if enumerated, might weU fill several pages of this book with a catalogue of modern achievement in the immediate province of our science. Nor is it only in the evolution of new and strikingly important materials or products that the new knowledge has made itself felt ; in the improvement of manufacturing processes its influence has already been great, and is destined to become still greater, since it opens up entire new ranges for methods of control and new understanding of what is really vital and what is accidental or unimportant in the processes now in use or in others whose use is suggested. This is a benefit which the leaders of our technical industries have been somewhat slow to avail themselves of, particularly in the industries connected with metals. These are necessarily old industries, in which modern ideas have had a somewhat severe struggle for supre- macy. Even now that supremacy has not yet been won as fully as it should so far as Great Britain is concerned. Manu- facturers are still too slow to make use of the methods which scientific research along the lines of Physical Metallurgy has rendered available. On the other hand, too, it must be remem- INTRODUCTORY 5 bered that our science is still in its infancy and that, much as it has already achieved, great regions of darkness still lie over the majority of our industrial and technical processes. The new science can lay no claim even to proximate completeness, and for that reason we have said that the manufacturer should utilise the methods of research developed in this subject — ^for in many cases the results he requires for his practical guidance are not yet available and must be found in his own laboratory or, at all events, at his own expense. But the science has progressed far enough to furnish the capable investigator with methods powerful enough to give a prospect of success in an attack upon the most difficult practical problem, provided that the requisite means and the requisite time are allowed for obtaining the solution. But it must be borne in mind that many of these " practical " problems involve the very deepest questions of our science, and they cannot be solved with anything but the best aids to research, nor by any but the best-trained of investi- gators. For that reason, perhaps, the rather halting efforts at the practical application of our science which have been made in a good many Works have led to somewhat disappointing results. Once the conditions of the case are fully understood and properly met, however, far-reaching success is certain to follow, and, as a matter of actual fact, has already been obtained in many cases. There is another special aspect of the practical importance of Physical Metallurgy which perhaps deserves emphasis at this point. The importance of the subject to those who are con- cerned with the production and manipulation of metal has just been discussed, but the subject is of at least as high a degree of importance to those who have to deal with the uses of our metals. In the widest sense this embraces the whole of civilised mankind ; in the narrower sense the " users of metals " are engineers and constructors generally. To these a better under- standing of their most important materials of construction is obviously of paramount importance. At the present time, engineers in all branches of design are forced to adopt what are known as large " factors of safety " — i.e., to employ quantities of material from three to eight or even ten times in excess of 6 STUDY OF PHYSICAL METALLUEGY what -would really be required to afford the necessary strength if all the factors of the case were completely known and the behaviour of the metal fully understood and its consistent performance entirely to be relied upon. In spite of this expensive precaution, however, failures and breakages still occur at intervals in engineering practice, thus proving that even these " factors of safety " do not really entirely ensure the degree of safety which their numerical value would suggest. Of course to some extent uncertainties exist in all engineering work, and these are by no means confined to questions of the quaUty and exact mode of behaviour of metals. Yet this latter factor is a very important one — in part because the precise behaviour of metals is reaUy subject to considerable apparent vagaries and uncertainties, and partly because in the absence of complete knowledge or of searching investigation it is easy to throw the responsibility for any given failure upon the material of construction, and thus to overlook or to evade the real lessons of the failure from the point of view of design. From this point of view the development of the methods of Physical Metallurgy has already placed at the service of engineers and metallurgists powerful weapons for the exhaus- tive investigation of cases of failure or breakage. These methods are not at present utiHsed as fuUy as in the interests of the advancement of technical knowledge they should be. The reason lies partly in want of acquaintance, on the part of those vinder whose jurisdiction practical failures occur, with the nature and extent of the investigation which is now possible, and partly also from a desire to avoid or shirk searching inves- tigation for business reasons. The maker of a metal article which has failed frequently shows an easily understood anxiety to melt up the broken parts, and dislikes having the whole case thoroughly studied, perhaps by an independent expert. Yet such a course — of covering up traces of failures as rapidly as possible instead of studying them carefully and drawing valuable conclusions from them — is not in the long run wise even in the immediate interests of the manufacturer himself, and is certainly inimical to technical progress generally. For unless the manufacturer has knowingly supplied a defective INTRODUCTORY 1 article, lie himself is generally ignorant of the cause of these failures, and it is only by the careful study of such cases that his knowledge of his own processes, and of the risks of failure attaching to them, can be materially increased. It may, however, be hoped that, as the value of the methods of Physical Metallurgy becomes more widely known and appreciated, those concerned with cases of failure or breakage will avail themselves of these means of investigation to a steadily increasing extent, either by themselves estabhshing properly majuied and equipped research laboratories or availing them- selves of the services of public laboratories or of private con- sultants competent to deal with such matters. It may perhaps be pointed out here that the systematic utilisation of the methods of Physical Metallurgy need not, and, indeed, should not by any means be confined, even from the practical point of view, either to the direct control of Works practice in the manipulation of metals, or to the investigation of failures. Apart from definite research of an exploratory nature, another very important function remains to be men- tioned. While the study of failures occurring either under test or in practice indicates very clearly the sources of trouble, and thus shows what must be avoided by rational practice, its indications are still of a somewhat negative character, and something more is required to point out in a definite and systematic manner the path towards the best possible results. The real test of " best results " lies, of course, in the continued test of actual use, and experience quickly shows that few, if any, industrial metallic products are so uniform in quaUty that this searching test of practical service does not distinguish widely between different batches or even between different individual pieces which are all nominally exactly alike. We have in this circumstance a further most valuable point of attack to which the methods of Physical Metallurgy could be most profitably applied, although at the present time a vast amount of valuable material of this nature is allowed to go to waste. The problem which is suggested by these variations of behaviour so often met with in practice is simply that of ascer- taining what are the real, although possibly minute, differences 8 STUDY OF PHYSICAL METALLURGY between the articles which have given the best results and those which have behaved either in a merely average or even in a barely satisfactory manner. There are two ways in which the methods of Physical Metallurgy could be readily applied to the solution of this problem. The first consists in making a very thorough examination, by the most searching methods avail- able, of a considerable number of articles as delivered by the maker or as sent out for use. This is not a question of tests to be made merely for the purpose of determining whether the goods in question come up to the standard required by the specification under which they have been purchased, but some- thing much more far-reaching. It is a question of determining as many of the properties and constants of the material as can reasonably be examined. These data should then be recorded and the behaviour of the individual articles followed through their career of actual service ; these service results will then be correlated, by systematic tabulation and comparison, with the investigatory test data obtained before the articles were put into use. It will not be long, if such a system is followed in a thorough manner, before relations between certain test results and service value become evident, thus furnishing guidance of the highest value in framing future specifications. To the user of metals the value of such a course is obvious, even though the labour and cost involved is not inconsiderable. We may take an example from railway work, such as rails and tyres. The service conditions and useful life of such objects could be followed up, in a selected number of cases, with moderate ease, even though a considerable number of cases would require to be studied in order to ehminate, as far as possible, the disturbing effects of local variations in service conditions. But what railway engineer would not be glad to have before him to-day, when called upon to draw up a specifica- tion for a fresh supply of rails or tyres, the data of chemical analyses, full mechanical tests, the micro -structure, macro- structure and thermal data covering several hundreds of rails or tyres whose subsequent service behaviour was also recorded ? At present only a few comparatively isolated data of this kind are available, and it may easily happen that the testa which INTRODUCTORY 9 have been made under various specifications did not really test just that property or combination of properties which is of primary importance for these very articles. Nor need the manufacturers of metal objects fear the results of such a systematic correlation between preMminary exhaustive tests and service behaviour. What a manufacturer whose aim it is to supply sound and reliable goods and thus to build up a reputation in the engineering world, has to dread are chiefly two things — ^the first and most serious is a faulty specification which may, perhaps, impose very severe tests of the wrong kind, tests which the maker finds it difiicult to comply with, but which depend upon properties not really vital to the particular object for which the goods are intended. The second is the risk that, either through lack of adequate systematic data or through prejudice, goods which are in reahty inferior may attain a spurious reputation. Both these dangers would be steadily eliminated by the systematic appUcation of the methods of Physical Metallurgy which is here suggested, so that ultimately the manufacturer, as well as the user, must benefit materially by the advancement of technical knowledge which would result. The method of preliminary investigatory testing which has just been discussed has one material drawback — that its results can only be arrived at gradually in the course of time, as experience of practical behaviour is obtained in correlation with preliminary data. For that reason the second way of applying the methods of Physical Metallurgy in correlation with the results of practice is worthy of some consideration, although its results are obtained on a far less satisfactory and conclusive basis. This second method consists in the syste- matic study and examination of what may be described as " favourable examples." In the case of railway work already cited, rails and tyres which have behaved particularly well in service might be utihsed in this way, and investigations under- taken in order to ascertain not only the causes of failure or unsatisfactory behaviour, but to determine the principal characteristics of metal which has been found particularly successful in the searching test of actual practice. To some extent this method also requires time to mature its results, 10 STUDY OF PHYSICAL METALLURGY since many " favourable examples " would need to be studied before definite conclusions could safely be drawn. But at least the long period of waiting for the tested objects to wear out or fail would be eliminated. This latter method also has another obvious advantage, which lies in the fact that in a great many cases the exhaustive physico -metallurgical study of an object implies its destruction so far as future practical use is concerned. This does not apply to such a thing as a rail, from which relatively small pieces can be cut at either end, but it applies very markedly to the case of a tyre. In many cases, if not in all, this difficulty can be overcome by the use of special methods of testing and examination, but in the case of " favour- able examples " withdrawn from service for purposes of study no such difficulty arises, and in many cases much more thorough examination can then be undertaken. The growing extent to which the importance of physico- metallurgical methods in connection with engineering is being recognised may be seen from the rapid growth and vigorous work of various societies and associations which are largely concerned with Physical Metallurgy. Thus the International Association for the Testing of Materials, which was founded by Bauschinger in 1884, has shown a steadily increasing member- ship and rapidly increasing vigour. While perhaps at the outset this Association was primarily interested in the purely mechanical testing of metals, work of a more or less distinctly metallurgical type soon made its appearance in its Proceedings, and this side of its activity has grown very much in more recent years. At the present time the scope of the work of the Association (Metals, Section A) is practically co -extensive with the subject of Physical Metallurgy as here defined, and the questions of mechanical testing, both commercial and scientific, are rightly considered as closely related to the study of metals by microscopic, thermal and other physical methods. The foundation and rapid growth of the Institute of Metals, whose scope is exactly co-extensive with that of Physical Metallurgy so far as the non-ferrous metals are concerned, is another indication of the increasing recognition of the value and importance of our subject. INTRODUCTORY 11 In dealing with so wide a subject in an introductory volume it is obvious that it will not be possible to attempt an exhaustive treatment, either in regard to the subjects to be touched upon or even in the treatment of those subjects which are dealt with. All that can be aimed at is to present a general survey, not so much for the purpose of giving detailed information or instruction as with the object of arousing an interest in the whole field of inquiry and of conveying some idea of what has already been achieved and in what directions further advance is to be sought. For those who are only indirectly interested in metals, such a general survey will perhaps prove sufficient ; for others who wish to take up the subject in detail, much fuUer study is required, and the present volume can only serve literally as an introduction to the literature of the subject. Indeed, although Physical Metallurgy as a whole may be looked upon as the branch of a specialist, yet for the worker in that field still further specialisation is required, so that each subdivision of the subject requires separate treatment m an individual volume. The general survey of Physical Metallurgy which is con- templated in this book may perhaps be best undertaken if the whole subject is approached from some distinct single point of view. The one which it is proposed to adopt here is based upon the view that the fuller understanding of metals is dependent upon a loiowledge of their internal structure and constitution. The methods of studying that structure and constitution, and the resulting knowledge, will therefore be dealt with first, the useful and interesting properties of metals and alloys being subsequently dealt with — so far as space permits — ^from the point of view of their structure. In the first instance we shaU confine our attention to substantially pure or " simple " metals, since their nature and behaviour is better understood than that of alloys in many respects. Sub- sequently the nature of alloys will be considered, and our knowledge of pure metals will be applied to them so far as that is possible. As far as possible it will be well to deal with general ideas and principles rather than with detailed facts, but in order to offer to the reader a survey of the subject 12 STUDY OF PHYSICAL METALLURGY possessing any degree of completeness, it wUl be necessary to include a section dealing directly with the more important groups of metals and alloys. In dealing with the subject in the manner just indicated it will not be possible to retain any direct connection with the chronological order in which the various branches of the subject have been developed. Historical references, however, are of little direct importance to those making their first acquaintance with the subject, and little space can be devoted to them in the present volume. The history of the subject is, however, largely contained in the references to original papers which are given at the end of each chapter. A very brief general outline of the history of the subject is all that can be given here. Observations on metals in some form or other go back as far as the uses of metals themselves, and metal objects of very great age have been found in the East. In that very vague sense, therefore. Physical Metallurgy may well lay claim to great antiquity. On the other hand, as a more or less self- contained science, it is of quite recent date. The great modern growth of interest in the detailed study of metals has, in fact, arisen from the remarkable results which have flowed in the first instance from the application of the microscope and the pjnrometer to the examination of metals. The first systematic application of the microscope to the study of metals was made by Henry Clifton Sorby, of Sheffield, in the year 1864. Sheffield is rightly proud of her distinguished son, who thus laid one of the foundation stones of our science. There is, however, less ground for local pride in the fact that, although Sorby lived among workers in steel who should have been the first to value the work which he did in their midst, yet that work was allowed to fall into forgetfulness. The whole matter was, indeed, completely neglected, both in Sheffield and elsewhere in England, until it had been indepen- dently taken up by Martens in Germany and Osmond in France, neither of whom were aware of Sorby's earher work. Among other early workers in this subject may be mentioned Werth, Grenet, Charpy and Le Chatelier in France ; Heyn, Wiist INTRODUCTORY 18 andTammanninGtermany ; Andrews, Arnold, Roberts-Austen and Stead in England ; and Howe and Sauveur in America. The fact that the present author was privileged to count Roberts-Austen and Osmond among his personal friends, and that Arnold and Stead are still actively at work in this field, serves to show how very recent this whole development has been. Another sign of the same kind is that ten years ago there was not only no text-book of Physical Metallurgy, but not even a book or manual on metallography in existence, so that students of the subject were restricted to the reading of original papers, many of which — ^as is often the case with a new and rapidly developing subject — ^were of a decidedly contro- versial character. The mechanical properties of metals have, of course, been studied both by engineers and physicists for a very long time, and in that direction also our science is of no recent origin. Galileo was probably one of the first to make a strength test on a metal, since he determined the length of a copper bar which, if suspended vertically, would just break under its own weight C). Curiously enough, this method of stating the strength properties of a metal has recently been revived under the term " specific tenacity " {*), for the purpose of allowing comparisons on a sound basis between light alloys of consider- able strength and other materials, such as steel, whose greater strength is more or less outweighed by their higher density. As a matter of fact, however, the men of the earher middle ages had little interest in the strength of metals, except possibly in regard to arms or armour. The great architectural achieve- ments of the " Gothic " period, although dependent upon some fair knowledge of the laws of equihbrium and stabihty, called for little knowledge of the strength of materials, since in these great structures of stone the material is entirely in compression, and stone possesses great strength against this type of stress. Consequently, even in the absence of accurate knowledge, the Gothic builders rarely overloaded their material so far as compressive strength goes, although they did frequently overload the foundations of their buildings. It was really the requirements of bridge-building that began 14 STUDY OF PHYSICAL METALLUEGY to lead engineers to undertake mechanical tests of their materials, and early efforts in that direction are associated in France with the names of Perrouet, Rondelet and Labardie, and in this country with Brunton and Brown, Bramah,Fairbaim and Hodgkinson. Probably the first tensile testing machine provided with an hydraulic ram for the application of the load and a lever for its measurement was installed at Woolwich Arsenal before 1837, the early applications of the hydrauUc press being, of course, associated with the name of Bramah. Somewhat later came the work of Eaton Hodgkinson, who described his experimental researches on the strength of pillars of cast iron and other materials in 1840, while the work of Fairbairn followed in 1850. Since that time there has been a steady development in mechanical testing appliances, some of those most prominently associated with this movement being Kennedy, Unwin, Ewing, Baker and Kirkaldy in this country, and BauschLnger and Martens in Germany. With regard to the magnetic, electrical and other properties of metals, the diversity of subjects is so great that it is difficult to indicate even outstanding events in the history of the subject. The whole trend of events, however, can be briefly summarised to this effect, that all the properties of metals — ^including their mechanical properties — ^have been studied and investigated in the first place essentially for their own sake or for the light which they might throw upon that particular branch of know- ledge with which they were concerned. It is only quite recently that the grouping together and the correlation of all these properties has been undertaken and that these things have begun to be studied not so much for their own sake as for the light which they could throw upon the nature, structure and constitution of metals. It is this manner of regarding the properties of metals and the methods of studying them that has brought into being the new science of Physical Metallurgy. But on the other hand this very development owes its inception very largely to the fact that the exploration of the internal structure and constitution of metals and alloys which consti- tutes modern metallography has itself thrown so much light upon the nature and behaviour of metals that it has awakened INTRODUCTORY 15 interest in all the properties of metals and has stimulated efforts in their study by holding out promises — ^which in many cases have already reached fulfilment — of still further valuable light, where light is so much needed, in that region of compara- tive mystery and imperfect knowledge which stiU lies much too near the whole field of our most important materials of construction and technical use. Befzbences. (1) Eosenhain. "Some Alloys Suitable for Instrument Work." Proceedings of the Optical Convention, Vol. II., 1912. (2) Un'w'in. Presidential Address to the Institution of Civil Engineers, 1912. (3) Eosenhain and Archbutt. " Tenth Report to the Alloys Eesearch Committee — on the Alloys of Aluminium and Zinc." Proc. Inst. Mech. Engineers, April and May, 1912, pp. 359 —362. PAKT I. THE STRUCTUEE AND CONSTITUTION OF METALS AND ALLOYS CHAPTER II THE MICKOSCOPIC EXAMINATION OF METALS Since the whole subject of Physical Metallurgy is to be treated in the present book largely from the point of view of the internal structure of metals, it becomes essential to under- stand at the outset the manner in which that structure can be studied. This study is undertaken principally by the aid of the microscope, and the present chapter will therefore deal with the preparation of specimens of metal for microscopic examination. In this branch of the subject, as indeed through- out the book, more attention wUl be devoted to the principles underlying the various operations and the laws which govern them than to the detailed description of apparatus or of experimental methods. Looked at in the ordinary way, metallic objects possess an appearance of complete homogeneity, but the existence of a structure at once becomes obvious when a piece of metal is broken. It is not surprising, therefore, that Martens began his work on the apphcation of the microscope to the examina- tion of metals by the microscopic study of fractures (^). This, however, proved a very limited field at the outset, as the irregularities of fractured surfaces make it impossible to use more than very moderate magnifications, while the path of the fracture itself is not necessarily a fair cross -section of the material — i.e., the fracture may, and in some cases does, pick out a path through a weak constituent. The necessity of THE MICROSCOPIC EXAMINATION OP METALS 17 using cut sections having a plane surface had been recognised by Sorby, who worked on the strong analogy which undoubtedly exists between metals and igneous rocks so far as micro - structure is concerned. Sorby, having developed the methods of microscopical petrology by the study of thin sections or slices of rocks, attacked the corresponding problem for metals. The use of thin sections through which Ught could be trans- mitted in the ordinary manner employed for the microscopic study of organic tissues or of most rocks was, however, soon found to be impracticable. The earhest objection arose no doubt from the practical difficulty of grinding and polishing such sections, since, owing to the opacity of metals to hght, they would have to be extremely thin — a slice of iron, for example, in order to transmit even 10 per cent, of the Mght falling upon it would require to be less than 1 X 10~* milhmetres thick. Even with more modern resources it would be difficult and expensive to prepare such sections ; the attempt to do so systematically for microscopic study, however, has never been made, since it is now recognised that the very operations of cutting, grinding and polishing alter the structure of the metal to a certain depth below the surface, which varies with the amount of force employed in the processes ; even with very gentle rubbing, however, alterations of structure to a depth of one thousandth of a millimetre are produced in all but the hardest metals (^), and for that reason we should anticipate that the structure of the very thin shces of metal which would be required for use with transmitted light would be entirely altered throughout their thickness by the process of preparing them. The use of Rontgen rays for the purpose of penetrating greater thicknesses of metals has been tried by Heycock and Neville (^) in the case of aUoys of gold and aluminium, which contain constituents of widely different densities, and thus afford reasonable contrasts in their power of transmitting X-rays. When those experiments were made, however, the technique of radiology was yet in its infancy, and it seems probable that with modern appliances it might be well worth while to study this method more thoroughly, particularly with a view to the use of moderate magnifications. P.M. 18 STUDY OF PHYSICAL METALLURGY For practical purposes, at the present time at all events, the microscopic study of metals is only possible by the process of looking at suitably prepared surfaces by reflected light ; but this method has proved sufficiently powerful to yield an entirely new insight into the structure of metals and alloys. If any ordinary metallic object is looked at under a microscope, the surface will generally be found to be covered with markings — ^grooves, scratches, smears, etc. — ^which have obviously no connection whatever with its internal structure, but which simply owe their origin to the treatment which that particular surface happens to have received. Even if an ordinarily " pohshed " article be examined in this way it will be found to be covered with such accidental markings. The photograph (Fig. 1, Plate I.) shows the appearance of the " polished " brass case in which a microscope objective is kept — ^its surface shows some resemblance to a ploughed field. It is obvious that if the true structure of the metal is to be observed, the surface must be treated in such a way as to leave it as free as possible from any accidental surface markings. This is accomphshed by the process of polishing by methods which have been evolved and perfected for this special purpose ; these methods, however, differ widely from the relatively crude and rapid processes employed in commercial polishing — a difference which arises not only from the more perfect character of the poUsh required when the surface is intended for examina- tion under the microscope, but also from the fact that, while for microscopic study the metal should be treated as gently as possible in order to avoid disturbance of its structure to any considerable depth below the actual sTirface, in the commercial processes rapidity is the principal object, and actions of a " burnishing " nature are freely utilised without regard to the internal disturbances which they produce. The preparation of a specimen of metal for microscopic examination necessarily begins by the choice of the surface to be examined. This is a matter in which the knowledge and powers of observation of the worker come into play to a very large extent, for the choice of the section must depend in every case upon the particular object to be attained and upon the THE MICROSCOPIC EXAMINATION OP METALS 19 conditions of the case. Thus where an investigator is studying the structure of a series of alloys in order to assist him in understanding their constitution, he will, in general, be dealing with small experimental ingots, and will, as a rule, find it desirable to examine an entire vertical cross -section of such an ingot in order to detect variations arising from any possible lack- of uniformity of composition from top to bottom, or effects of variations in the rate of cooling from the centre to the outside of the ingot. The whole of such a cross -section may, of course, be inconveniently large for the purpose, since, as a rule, it is not wise to undertake the preparation of areas much larger than four or five square centimetres (about one square inch) in a single piece. Where larger areas are required, the specimen is best cut up into two or more pieces and each polished separately. More complex conditions affecting the choice of specimens arise where the material to be examined forms part of a larger mass of metal, whether cast or wrought. In castings there are the effects of variations in cross-section and rate of cooling, as well as possible variations in chemical composition, to be con- sidered, while in rolled or forged metals the influences of mechanical treatment always produce differences of structure according to the direction in which the section is cut relatively to that in which work has been applied. As a rule, therefore, it win be necessary to cut several sections, in different planes, from each larger piece of metal to be examined. Where breakages are under investigation it will further be necessary to compare the structure close to the fracture, and in the neigh- bourhood of any fine cracks, with that of the material at a distance from the fracture. No general rule can be laid down for the selection of the surfaces to be examined, but it is prob- ably the mistake which is most frequently made in this con- nection to examine too smaU a number of specimens, or too small a total sectional area of the material. Quite apart from the possibility of finding, by the aid of numerous specimens and extensive examination, some local feature or defect in the metal, this extended examination has the further advantage that it serves to impress on the mind of the observer certain features 02 20 STUDY OF PHYSICAL METALLURGY of the structure which, while really most significant, are not very prominent in any single area. Often enough some such feature is disregarded, even by careful observers, as being accidental or unimportant, but when the same pecuMarity is repeatedly observed on a number of sections it begins to make its proper impression on the mind, and thus saves the investi- gator from overlooking an important line of evidence. It has several times been the experience of the present author that, on examining the first few sections of a piece of metal which had shown entirely abnormal mechanical behaviour, no ab- normal structural feature could be detected, but that by the prolonged examination of numerous sections the somewhat less obvious, but none the less real, abnormality was noticed and ultimately traced to its origin. The rule thus indicated, of examining the largest possible number of specimens, need not, of course, be followed in a con- siderable number of cases — where either the examination of the very first section shows an obvious and unmistakable feature of a kind which must repeat itself throughout the mass of metal, or in cases where the microscopic examination is imdertaken for the specific purpose of determining whether the material exhibits some special feature, such as segregation, or the typical coarse angular structure due to overheating. The whole question of the choice and number of sections to be examined is thus a matter essentially for the judgment of each worker, and what has just been said can only serve as a rough guide. The manner in which a specimen of convenient size for purposes of polishing should be cut off from the rest of the material will, of course, depend upon the size and shape and nature of the object in question. This is, however, a simple question of a kind which constantly arises in all engineering workshops, and requires no further consideration here beyond the remark that in every case care must be taken to avoid any- thing which may affect the structure or constitution of the metal to be examined. Thus the modern methods of cutting metal by the aid of the oxy-acetylene blow-pipe must be avoided, or at least carefully controlled, if employed for the THE MICROSCOPIC EXAMINATION OF METALS 21 present purpose, since the temperature attained by the metal, even at a considerable distance from the line on which the blow -pipe has been used, is quite high enough to alter the struc- ture appreciably, and thus — ^possibly — ^to efface the very features which the microscopic examination is intended to detect. The use of such a method of cutting must, therefore, always be looked upon with suspicion, particularly as it is not impossible that the process may be intentionally resorted to by some interested party with the direct object of interfering with the microscopic diagnosis of the case. Violent mechanical methods, such as shearing or punching, or breaking off by blows after preliminary nicking, are also to be avoided, since these processes are also liable to introduce disturbing features into the microscopic appearances presented by the metal. , When a piece of convenient size has been obtained in some satisfactory way — and in a great many cases it is simply a question of the skilful use of a simple hack-saw for a few minutes — that surface of the piece which corresponds to the section selected for study must first be cut down to the nearest possible approach to flatness. This may be done either by means of the file or the grinding wheel. There is no doubt whatever that the latter is far preferable. The necessity of gripping the specimen in the vice — with its attendant risk of mechanical injury and distortion — ^is entirely avoided ; but a much greater advantage lies in the fact that the process of grinding, particularly in the case of the softer metals, causes much less disturbance of the metal below the surface. It is found that the marks, particularly of a rough file, are difficult to remove entirely by the subsequent rubbing on comparatively fine emery papers, and, even when apparently completely effaced, there appear to be residual regions of disturbed structure for a slight depth below the deepest file-cuts. The individual cuts produced by the grains of emery or carborun- dum in a wheel are much smaller and more numerous than the cuts of a file, and the disturbance which they cause does not penetrate so far into the substance of the metal. A disadvantage which attaches to the use of the grinding wheel in the ordinary way is that the specimen rapidly becomes 22 STUDY OF PHYSICAL METALLURGY heated by the energy dissipated in friction at the grinding surface. This heating is not intense enough to affect a great many materials, but in others — such as hardened steel— serious changes may easily be caused by this rise of temperature. In any case, however, it is far safer to guard against this heating by constantly dipping the specimen into cold water and never holding it in contact with the grinding wheel long enough at a time to allow it to become seriously heated. This process of repeatedly removing the surface from contact with the wheel, however, makes it difficult to grind anything like a flat surface, since the angle of contact is apt to change shghtly every time. This results in the formation of a fresh facet on the surface every time it is brought against the wheel, and, finally, a some- what convex surface is produced. This is extremely imdesir- able, since it not only renders the subsequent polishing opera- tions more difficult, but also interferes with the definition of the images furnished by the microscope from the finished specimen. In order to avoid some of these difficulties and at the same time to render the whole grinding operation less violent, the author has adopted the use of large, slow-running grinding wheels. Those actually employed have a diameter of ten inches and run at speeds of approximately seventy revolutions per minute. Grinding with these is, of course, much slower than with the ordinary wheel running at several thousand revolutions per minute, but the sKght loss of time is more than compensated by the fact that, first, the specimens do not get seriously heated even after several minutes of continuous grinding, and, second, it is easy to replace the specimen in contact with the grinding wheel — ^if it has been taken away for purposes of examination — by the simple method of feeling for the position of best contact. With a fast-running wheel a fresh facet is formed at the very first instant of contact, and any endeavour to adjust the position of the specimen merely results in grinding a curved surface. With the slow-running wheel one has time to get the specimen adjusted in its former position before material grinding has taken place. A further advantage of the slow-running wheels is that the " cut " of the THE MICROSCOPIC EXAMINATION OE METALS 23 abrasive is not so deep and does not cause so much internal disturbance as that of a wheel of the same grade running at a high speed. In actual practice the use of these slow-ruiming wheels has been found to effect a material saAring of time over the whole process of grinding and pohshing. Although the grinding process by the use of rotating discs armed with abrasives can be carried right down to the finest available grade of emery, etc., this is not found desirable. The objection to mechanical violence becomes greater as the grade of abrasive becomes finer, for the simple reason that less of the material remains to be removed in the subsequent operations, so that any disturbances produced are hkely to remain behind, only to be detected when the poUshed surface is chemically attacked in that stage of the process known as " etching." For that reason many workers, including the author, employ the finer grades of abrasive entirely by hand, using the well- known device of turning the specimen through a right -angle on passing from one grade of emery to the next finer one. In this way it is possible to see that the visible scratches left by each grade are completely effaced by those due to the finer grade. Ultimately the surface of the specimen is left covered with a system of more or less parallel scratches or marks from the finest available grade of emery, and these are then removed by the process of polishing proper, which is usually carried out by means of a rapidly revolving disc, covered with soft cloth or leather, and fed with water and some form of " pohshing powder." If the emery grinding has been correctly carried out at every stage, the final removal of the emery marks by the aid of the wet disc should not occupy more than four or five minutes with specimens of steel of the size indicated above. If a longer pohshing process is needed some of the earher stages have not been properly carried out, and a really satisfactory metallo- graphic poMsh will not be obtained. With softer metals the conditions may be somewhat different, and success in pohshing these is a matter requiring the most scrupulous care ia regard to cleanliness throughout the entire operation, and also a certain degree of skill, which is only acquired by prolonged experience and requires something of a natural gift of touch. 24 STUDY OF PHYSICAL METALLURGY Extreme difficulties of this kind are only met with in dealing with such metals as lead or thalUum, but to a lesser extent they are also encountered with aluminium and, in a decreasing degree, with copper, brass and bronze. Although in the first place the operations of grinding and poHshing are of interest from the point of view of experimental practice, careful study of the phenomena connected with them has led to most interesting results — in the first instance at the hands of Beilby (<) — and although these can only be considered in detail at a later stage, when the whole theory of the amor- phous condition in metals will be discussed, some reference to them is required at this point in order to afford some insight into the real meaning and nature of the operation of polishing. The nature of the processes involved in ordinary cutting operations are now fairly well understood. Essentially the action of an edged cutting tool consists in bringing to bear upon a very small area of metal a stress sufficiently intense to produce rupture. Obviously, however, although actual rupture is confined to a single line or surface, severe strain will be pro- duced in the immediate vicinity, so that in an ordinary machined or filed surface the visible grooves are accompanied by corre- sponding sub -surface regions of strained metal. This region will be deeper, the deeper the " cut " which has been taken and the greater the force which has been employed, and also the blimter the tool — i.e., the larger the area to which the intense pressure of the tool has been appHed. When a grinding -wheel or emery paper is substituted for a file, then the same process continues, at all events so far as the coarser grades are con- cerned, only that the actual cutting edges of the individual grains are exceedingly minute and sharp, while the depth of cut taken by each of them is very small. Consequently the depth of disturbed metal beneath the surface is correspondingly slight, particularly if the emery grinding is done gently. There are, however, other actions which accompany the cutting and grinding process, and these — although unimportant when a coarse cut is being taken — become predominant when the finest grades of abrasive are employed. A fuller under- standing of these actions will be obtained when the chapter on THE MICROSCOPIC EXAMINATION OF METALS 25 the efEects of strain on the structure of metals is studied, but for the present it will be sufficient to state that severe mechanical strain may in certain conditions break up the essentially crystalline structure and convert the metal locally to an intensely hard and brittle amorphous condition. In the process of such transformation from the crystalline into the amorphous, however, it appears that the metal passes through a temporary or transition stage in which it is capable of under- going a certain small amount of flow, much hke a viscous liquid such as pitch or ordinary oil paint. Now when oil paint is somewhat thickly appUed with a brush, the bristles of the brush as they pass along the surface leave behind them a system of ridges and furrows similar to those after filing or rubbing on emery paper. In the case of the paint, how- ever, the viscous hquid, acting under the influence of surface tension, gradually rounds out and smoothes over these ridges and grooves, finally assuming the smooth surface typical of a liquid at rest before it sets hard enough to prevent further flow. Now, as Beilby has shown, this analogy is so far complete that in metals, too, the sharp angular grooves left by the abrasive grains show a considerable tendency to become smoothed out. When the grooves are still fairly coarse the distances involved are too large and the metal does not retain its mobihty long enough, nor is a sufficient quantity of dis- turbed metal available to smooth out the furrows to a notable extent. When, however, the grooves become very fine — as they do with the use of the specially fine grades of emery paper employed for metaUographic work — ^then the effects of this flow- ing action become important and quite easily visible under high magnifications. The grooves left by these very fine abrasives " run together " and show the pecuhar broken-up, drop-like forms which are typical of the efEects of surface tension. Fig. 2, Plate I., shows a photograph, imder a magnification of 150 diameters, of such a system of grooves or scratches. This photograph was taken from a surface as obtained direct from a No. 000 emery paper, without subsequent treatment of any kind. Now when such a surface is subjected to the action of a THE MICROSCOPIC EXAMINATION OF METALS 27 with rouge it occasionally happens that particles of rouge are subsequently exuded from the pohshed surface. It is a further significant fact in this connection that not every finely-divided substance can be successfully employed as a polishing agent. This might perhaps be ascribed to variations in hardness or in the shape of the ultimate particles, but such suppositions do not account for the fact that all the good poBshing materials — such as alumina, magnesia, rouge, putty powder, oxide of chromium, etc., belong to one particular chemical class, viz., that of oxides. Now many metals are known to possess the power of dissolving their own oxides, and this property may have some considerable bearing on the polishing action of these bodies. At all events, if the pohshing agent is forced to interpenetrate the surface layers of the metal, the " forced solution " thus formed may well exhibit a greater degree of mobihty than the pure metal — just as the mobihty of a metal is increased by the addition of a second alloying metal which lowers its melting point. By the combined mechanical and chemical actions of the pohshing disc and powder the surface layers of molecules are given a certain degree of temporary freedom, and are thus able to form a smooth " polished " surface, such as that of a hquid at rest. Indeed, if we look at the really smooth surfaces found in nature, we see at once that, apart from the faces of single crystals, they are all surfaces of liquids — even ordinary sheet glass has a natural smooth surface, known as " fire poUshed," because it is simply the congealed smooth surface formed on the glass when it was fluid. It is thus interesting to see that poHshed surfaces are in reality also formed by the smoothing out of an extremely thin fluid layer under the action of surface tension. From the point of view of practical metallography this view is of particular importance, because it serves to explain a number of features which are met with in polished surfaces. One of the most striking and, at flrst sight, perplexing of these is the manner in which an apparently perfectly poHshed surface shows, even after very shght etching, a number of scratches and holes which were certainly not visible before. The beginner 28 STUDY OF PHYSICAL METALLURGY is apt to assume that he has somehow managed to scratch the dehcate polished surface during the handling of the specimen while he was etching it, but, although this may occur to some extent, yet such scratches and grooves still make their appear- ance even if the utmost precautions are taken to avoid all risk of injury to the surface. The real cause of this phenomenon is that these grooves and scratches were really present all the time, but that they were filled up and covered over by the amorphous layer which was formed during polishing. This layer is at once removed, more or less completely, when the surface is attacked by some etching reagent, and the original markings are thus laid bare. This fact can be readily verified by photographing a marked place on a specimen which shows some well-defined features of this sort after a first etching. Then let the surface be re-poMshed on the disc (without the intervention of any emery grinding) until all signs of the mark- ings have disappeared ; then, if the surface be again etched, it will be seen that the deepest and most prominent of the original markings have reappeared. The extent to which this recrudescence of scratches, etc., will take place will depend entirely upon the depth of the markings left on the surface before polishing was begun, and it is in order to minimise annoy- ance and delay, or even error, from this cause that the use of gentle methods of grinding and pohshing is advocated here. The appliances to be employed for pohshing, being of a simple mechanical kind, need not be described here, but a few words are required in regard to the material of which the pohshing disc, or rather its working surface, should be made, as well as to the pohshing powders to be used. For the great majority of ordinary work the pohshing disc is best covered with some soft material, such as fine wash- leather, " Selvyt," or best quahty wooUen cloth of the kind used for liveries. The latter material gives much the longest wear, and, if the right quahty is obtained, gives very satis- factory results. " Selvyt " pohshes very well, but wears out rather rapidly, while wash-leather is not particularly well suited to the modern polishing powders, such as alumina or magnesia, although it worked excellently when rouge was THE MICROSCOPIC EXAMINATION OF METALS 29 universally employed. For extremely fine work, where the erosion of soft materials and the rounding off at the edges of minute holes which results from the use of thick, soft fabrics for poUshing is objectionable, the best material is old, well- worn and washed calico, used in several thicknesses. It may be possible to " age " this material artificially, but, if purchased new, it is never successful — ^the fibre appears to be hard enough to scratch the metal, even in the case of steel. A still less jdelding bed for polishing purposes is formed by pitch, which is widely employed for the pohshing of glass. Glass, however, is much easier to polish than metals, and in the author's hands pitch has never proved very satisfactory. In the early days of the microscopic study of metals rouge was practically the only polishing material employed. It is the merit of Le ChateHer to have drawn attention to the fact that certain other substances yield a polish at least as good as that given by rouge and attain that object with much greater speed. Le Chatelier (^) at the same time described an elaborate method for the preparation of these new and superior pohshing media, such as alumina, oxide of chromium, etc. While the materials prepared by his method are very excellent, the methods them- selves are lengthy and expensive, as the yield of useful pohshing powder is decidedly small compared with the cost of such materials as ammonia alum, ammonium chromate, etc. Many of these materials may at the present time be purchased com- mercially, while the best quality of " heavy oxide of magnesia," if carefully sifted and kept from the air and moisture, serves as a very good substitute. Alumina may, however, be prepared by the direct oxidation of metalhc aluminium. This can be brought about by inoculating the metal with a little mercury, when oxidation proceeds rapidly for a time. The product, kept free from dust or other contamination, is an excellent pohshing powder. When a properly pohshed surface has been duly obtained by the processes which have just been discussed, it will be found that under the microscope such a surface appears a nearly featureless blank — the more nearly perfect the pohsh, the fewer are the features to be seen on the surface. If the surface is 30 STUDY OF PHYSICAL METALLURGY examined under a low magnification — say fifty or one hundred diameters — and with obUque illumination, a very severe test of the perfection of the poHsh is obtained, since under these conditions, especially if the specimen is rotated, even the minutest scratch or hole becomes apparent. For most pur- poses, however, no such searching test is required, and a surface can generally be regarded as fairly satisfactory if under a magnification of one himdred diameters and " normal " illumination the field of view does not show more than one or at most two fine scratches and few or no holes. The few remaining scratches and defects do not as a rule interfere to any material extent with the subsequent examination of the structure, and it is a useless expenditure of time and energy to strive after greater perfection of pohsh except for very special purposes. When the metal under examination is an alloy con- taining constituents which differ very considerably in hardness, the poHshed surface is not generally a featureless blank, since the process of polishing usually erodes the softer constituent and leaves the harder one sUghtly in rehef — this " rehef polishing " is seen as a faint pattern under the microscope, but its indica- tions are not always to be rehed upon, and it is to be regarded as rather an undesirable feature. Li order to afford any insight into the structure of the metal, the pohshed surface of a specimen must be treated in some way calculated to produce a surface pattern corresponding to the section of the internal structure presented by the pohshed surface. As we have seen, over the poHshed surface itself is spread a very thin film of amorphous metal, fiUing up and smoothing over not only the surface irregularities originating from the mechanical treatment, but also the irregularities of structure of the metal itself. This surface film must, there- fore, first be removed more or less completely by the action of some suitable reagent. Accordingly we find that the universal practice is to " etch " the surface by means of some reagent which slowly attacks and dissolves the metal. The amorphous surface film is more readily attacked than the rest of the metal, and is thus rapidly and easily removed. The subsequent action of the etching agent then depends upon the nature of THE MICROSCOPIC EXAMINATION OP METALS 31 the agent and upon that of the metal to whioh it is apphed. The reagents most frequently employed are weak organic acids or dilute solutions, in alcohol or water, of mineral acids. In the early days of the science dilute nitric acid, varying from a strength of 1 per cent, in alcohol to 0-1 per cent, in water, was very widely employed. The tendency of nitric acid to cause uneven oxidation and consequent tarnishing of the surface, particularly of iron or steel, has, however, led to its almost complete abandonment in favour of more rehable and satis- factory reagents. For iron and all its alloys a solution of picric acid in alcohol (either 4 per cent, or saturated) is now almost universally employed, although nitric acid much diluted with amyl alcohol, and hydrochloric acid in ethyl alcohol, are also useful. For other metals and alloys special reagents are employed in nearly every case. Copper and its alloys are frequently etched with a solution of ferric chloride acidified with hydrochloric acid, while for aluminium and its alloys either caustic soda, hydrochloric acid containing a small addition of nitric acid, or hydrofluoric acid are frequently employed. The exact choice of an etching reagent, although a matter of considerable importance, is a question of the chemical relationships of the metal or aUoy under investigation, and it is not possible to consider these matters in detail here. The reagents just mentioned all act in the same way, viz., by the gradual solution of the metal — a process which is of the nature of a gradual unbuilding of the structure of the metal from the surface downward. It is obvious at the outset that such a process must not be carried very far, otherwise con- siderable surface irregularities will be produced and the plane character of the surface, so necessary for microscopic examina- tion, wiU be destroyed. The extent to which etching may be carried depends, in fact, upon the magnification with which it is desired to examine the metal. In the case of mild steel to be examined under a magnification of 1,000 diameters, an etching with picric acid for a period of five seconds is satis- factory ; if the same steel is to be examined under a magnifica- tion of one hundred diameters, it is preferable to carry the etching on for thirty seconds in order to obtain a clear, brighj. 32 STUDY OF PHYSICAL METALLURGY image, but the surface thus treated would be too deeply attacked to be readily focussed under the higher power. When the surface is to be examined under very low powers indeed, or with the unaided eye, then the etching process must be carried very much further, and it is usually preferable to employ an agent which acts more rapidly, such as — ^in the case of steel — dilute sulphuric acid or a 10 per cent, solution of copper- potassium chloride. The chemical attack, or structural unbuilding of a metal by means of an etching reagent, however, does not take place uniformly over the entire surface ; the exact manner in which the attack takes place depends upon the character of the metal, i.e., whether we are dealing with a simple or a duplex material. The " simple " material here understood is one in which only a single constituent is present and in which all parts of the metal are of exactly the same chemical com- position ; in the case of a duplex structure, on the other hand, the surface is made up of two different constituents which differ in chemical composition. Li the latter case we usually find that one of the two is electro -positive to the other to such an extent that chemical action is almost entirely confined to the one, while the other remains unaffected until the etching process has been carried to a very considerable depth. In these circumstances we find that the more easily attacked constituent is roughened by the etching reagent, with the result that, as will be more fully explained below, it appears dark under the microscope, while the more resistant constituent remains smooth, and consequently appears bright under the microscope. The result is the formation of a sort of mosaic pattern, which is really a sectional plan or map of the structure of the metal. In the case of a pure metal, or of an alloy having only a single constituent, there can be no such preferential attack on one constituent with corresponding protection for the other, and one would at first sight expect such a pure metal to be uniformly attacked all over without the formation of any pattern at all. Actually this is the case to this extent, that an etching reagent which produces a well-marked pattern on a THE MICROSCOPIC EXAMINATION OP METALS 33 duplex alloy in a few seconds produces no visible pattern at all on a pure metal in the same period of time. If the attack is continued, or if a more rapidly acting reagent is used, how- ever, a very definite pattern is produced — the uniform surface of the specimen is found to become divided up into more or less polygonal, but somewhat irregular, areas bounded, appar- ently, by black Hnes when seen under normal illumination. We shall see in a later chapter that these polygonal areas are in reaHty sections of the crystals of which the metal is built up, and that each of them has been dissolved to a different depth according to the direction in which the plane of the section has cut through the crystal system in each individual case. For the moment it is sufficient to note that the slower rate of etching observed with pure or " simple " metals as compared with duplex alloys arises from the fact that, while in the latter case there are chemical differences between the two constituents, the various crystals of the pure metal are chemically all alike, but differ merely in crystalhne position or " orientation." In etching a pure metal, therefore, we are dependent upon very minute differences of chemical activity, and there is a corresponding difficulty in preventing the inter- ference of accidental circumstances which tend to produce stains or even spurious patterns on the surface. Much greater care must, therefore, be taken in both the polishing and etching operations when pure metals or " simple " alloys are dealt with than in the case of duplex alloys. The method of gradual chemical attack or solution is not the only one open for the purpose of developing a structural pattern, as it will be readily understood from what has been said above that any process which acts differently upon the various constituents or upon the individual crystals of a single constituent will develop the structural pattern. The process of " rehef pohshing " has already been mentioned, but by itself it does not jdeld satisfactory results. It has however, been employed successfully, in the first place by Osmond (*), in combination with a very mild form of chemical attack, thus constituting what is known as " polish attack." Here the pohshing process is carried on in the presence of a sHghtly P.M. D 34 STUDY OF PHYSICAL METALLURGY corrosive liquid which would not, by itself, produce satisfactory etching in any reasonable time. Combined with the polishing action, however, an exceedingly uniform and very gradual etching is obtained. The resulting patterns generally consist entirely of fine hnes indicating differences of level or depth, and there is Httle roughening or darkening of the attacked surface. Another method of developing the structural pattern of a poUshed surface is by the formation of surface films, whose thickness varies according to the constituent of which each particular element of the surface consists. These surface films may be formed in various ways. One of the best known is the process of " heat tinting," first developed by J. E. Stead ('), where the specimen is heated in contact with air, and surface films are formed by the gradual oxidation of the metal. The depth of film formed at any point varies according to the chemical composition of the metal at that point, and the sur- face becomes more or less deeply coloured accordingly. Stead has employed this method for a variety of purposes, but primarily for detecting the distribution of phosphorus in iron and steel. The presence of very small proportions of this metalloid produces a marked effect on the rate of oxidation of iron, with the result that very clear and beautiful patterns are obtained. A word of caution is, however, needed. The use of the heat-tinting method requires, in the first place, a per- fectly clean surface — soHd or liquid matter of any kind remaia- ing on the surface affects the rate of oxidation and spoils the resulting pattern. Beyond this, however, the heat-tinting method is peculiar in that there is no material removed from the " etched " surface. As a consequence of this pecuHarity it follows that the amorphous surface film which, as we have seen, is formed during the pohshing operation, is not removed by the process, and if there has been any material amount of " smearing " of the amorphous forms of the various consti- tuents over one another, the boundaries of these constituents, as revealed by the oxide films, may not strictly correspond to the internal structure. It is, therefore, always safer to etch a polished surface very Ughtly with a weak or dilute acid before undertaking heat-tinting, since such a proceeding certainly THE MICROSCOPIC EXAMINATION OF METALS 35 removes the amorphous surface film. The thickness of such a film being very small in the case of hard materials hke iron or steel, this precaution is not so essential with these metals as it is with softer ones like copper or its alloys. Surface films somewhat similar to those produced by the oxidation process which is employed in heat-tinting may be obtained in other ways. With metals containing silver or copper, exposure to an atmosphere containing hydrogen sulphide gas is sufficient to bring about a development of sur- face colours, but this is not very easily graduated to the right extent. Another and much more promising method was first employed in deahng with the alloys of copper and tin by F. GioUitti {^), and has recently been worked out more particu- larly for the case of steel ; it consists in bringing about upon the surface a very gradual deposition of metalhc copper by a process of electro-chemical replacement. When a piece of iron or steel is exposed to contact with a solution containing a copper salt, the iron is attacked, a portion of it passes into solution, and an equivalent quantity of copper is deposited as a film on the surface of the iron. As a mere method of etching, employing copper salt solutions as solvents for attacking iron or steel, this process has long been used, Hejoi (^) having employed solutions of copper-ammonium chloride for revealing more especially the details of crystalline structure, and also for that deep form of etching used when the surface is to be examined for its broad features without the use of magnifica- tion. The present author, together with J. L. Haughton (^''), has quite recently found that, when a suitable solution con- taining ferric chloride, hydrochloric acid, cupric chloride and stannous chloride is employed, the copper is deposited in a very interesting manner. In a very pure, mild steel, for instance, it is found that the appearance of the constituents as seen after " etching " with this new reagent is exactly the reverse of that seen after etching with picric acid ; when, however, the steel contains impurities — as all commercial steel does — ^which are not present as separate micro-constituents, but are diffused in solution throughout the metal, the effect of the new reagent varies according to the amount of such impurity present at d2 36 STUDY OP PHYSICAL METALLURGY each point of the metal. The purest part of the iron appears to receive the deepest deposit of copper. Accordingly, in a steel containing phosphorus, the deposition of copper affords a picture of the phosphorus distribution exactly similar to that obtained by heat-tinting. The copper deposition method is, however, very much quicker than heat-tinting. The process of etching is not quite the last which the speci- men must undergo before it is ready for microscopic examina- tion. When removed from the etching bath, the specimen must be very thoroughly washed in order to remove all traces of the etching reagent. This is essential, not only in order to arrest the etching action at the desired point, but also in order to prevent the surface of the specimen from being subsequently disfigured by the formation of crystals of the etching reagent or its salts. This is very apt to occur in the case of specimens containing holes or cracks into which the etching solution finds its way ; there it evades the washing, unless that is carried out very vigorously, and then it creeps out and spreads a fringe of crystals around the edge of the crack or cavity. This can be prevented only by washing the speci- men under a strong jet of water for several minutes. For many purposes it is then quite sufficient to dry the surface of the washed specimen by gently wiping it with a clean, old handkerchief, although this is liable to produce a few scratches on the surface. Where greater care is required, the specimen should be washed in absolute alcohol to remove the water, and it may then be dried by simple evaporation, aided perhaps by a fan or, better stiU, by a current of heated air. The alcohol sometimes leaves a slight film where it has last dried, and this can be avoided by washing the specimen, after it has been rinsed in alcohol, in clean, pure ether. This is, however, a somewhat difficult matter, as the ether is such a good solvent for aU manner of substances that it is very liable to take up some impurities in the course of a few weeks' use, and then the ether leaves a worse deposit than the alcohol used alone. As a rule specimens for microscopic examination are of somewhat transient interest and need not be kept very long ; the properly washed and dried surfaces remain useful for THE MICROSCOPIC EXAMINATION OF METALS 37 examination for several days or even weeks if kept in a well- made desiccator provided with a suitable drying agent — calcimn chloride is one of the best for this purpose. The specimens can, however, be given a considerable degree of protection from corrosion or tarnish by coating them with not too thin a layer of collodion or celluloid, dissolved in the form of a varnish in amyl acetate or acetone respectively. The majority of varnishes of this kind are too thin for this purpose, the tendency being to jrield a coating of such tenuity that it shows the interference colours of Newton's rings. By using a rather stronger solution this can be avoided, and the coat of varnish is still thin enough and transparent enough to aUow the specimens to be examined and photographed — even under the highest powers of the microscope — ^without loss of definition and without removal of the varnish. Specimens thus protected will keep for several months without tarnishing materially. In the case of iron and steel specimens permanent preservation is possible in another way, viz., by storing the specimens immersed in a solution of chromic acid or of one of the other reagents which render iron " passive." Not only do the specimens remain free from rust while immersed in such a liquid, but, if carefuUy handled, they retain their passivity for some time after they have been removed from the liquid. Eeferences. (1) Martens. " The Microscopical Examination of Iron." Zeitschr. Vereins Deutscher Ingenieure, Vol. XXI., January, May and November, 1878. (2) Beilby. " The Hard and Soft State in Metals." May Lecture, 1911, Jour. Inst, of Metals, Vol. VI. (3) Heyoock and Neville. Trans. Chem. Soc, 1898, 73, 714. Phil. Trans., 1900, 194a, 201. (4) Beilby. See note (2) above and note (2) in Chapter XI. (5) Le Chatelier. Bev. de Metallurgie, 1905, 2, 528. The Metallographist, Vol. IV., January, 1901, No. 1, p. 1. (6) Osmond, i^tude des Alliages, 1901, 277. (7) Stead. Journ. Iron and Steel Inst., 1900, II., 137. (8) Giollitti. Gazzetta, 1906, 36, II., 142 ; 1908, 38, II., 352. (9) Heyn. Verhandlungen Vereins zur Beforderung des Gewer- befleisses, 1904, 235. (10) Kosenhain and Haughton. Journ. Iron and Steel Inst., 1914. I. CHAPTER III THE METALLXTEGICAL MICEOSCOPE Whilb it would lie beyond the scope of this volume to offer any attempt at a complete account of the theory of the microscope or a detailed guide to practical microscopy, it is perhaps desirable to deal with this aspect of our subject to a certain extent. A rudimentary knowledge of the principles which underhe the working of the modern microscope is essential if that instrument is to be used to full advantage, and while there are many admirable treatises dealing with the microscope as such, they generally approach the subject from the standpoint of the more usual uses of the microscope where the apphcation of the higher powers is confined to the study of thin transparent sections. The apphcation of the instru- ment to the study of metals, where we have to examine approxi- mately flat surfaces by reflected light, introduces a series of different conditions. A further justification — if such were needed — ^f or the introduction of a chapter on the microscope into the present volume is provided by the numerous evidences of the abuse or inadequate use of the microscope which is furnished by the photo -micrographs published with some otherwise valuable papers on metals. We may begia by considering the manner of illuminating the metal surfaces which have been polished and etched for micro- scopic examination. We are confined to reflected light, but the incident fight may be made to fall upon the surface either obfiquely or normally, and these two alternative methods are both employed. Obfique illumination, produced when a suitably arranged beam of light falls upon the metal surface from some direction outside the lenses of the microscope, is in many ways the simplest, since no special appfiances or fittings are required. Its use is of considerable importance and value THE METALLURGICAL MICROSCOPE 39 particularly where it is a question of recognising the differences of level and other features of the surface configuration. On the other hand, the use of oblique light is bounded by certain limitations. In the first place, the character of the optical effects produced by this mode of lighting is pecuhar and in some respects exaggerated or distorted ; minute surface defects are apt to become unpleasantly prominent, and slight differences of surface texture appear as vivid contrasts of brightness. This exaggerated character becomes increasingly apparent as the light becomes more oblique. Accordingly, the use of obUque lighting is confined to the lower powers of the microscope (the 16 mm. is the highest power of the Zeiss apochromatic objectives which can be usefully employed in this way). With higher powers the free distance between the mount of the lens and the surface under examination is too small to allow an oblique beam to reach the surface at a reasonable angle. A further difficulty in the use of oblique light lies in the fact that the proportion of the incident light which is reflected into the microscope is very small, so that the images are not as a rule bright enough to lend themselves to photography. The extreme contrasts of light and shade also render photography difficult and the result unsatisfactory, except in special cases. The second, and most widely used, mode of lighting metal specimens under the microscope is generally known as " verti- cal " illumination, although it is more correct to caU it normal illumination. By this method the light is caused to fall upon the surface of the specimen in a direction at right angles to that P^ B Fig. 3. — ^Diagram of " Vertical " or Normal Illuminators. 40 STUDY OF PHYSICAL METALLURGY surface, i.e., if the microscope is used in the vertical position (which is — ^by the way — a very undesirable practice), then the light falls upon the specimen vertically from above — hence the term " vertical illumination." The manner in which this mode of illumination is obtained is shown diagrammatically in Fig. 3, A and B. Both figures represent sectional views of a microscope and metal specimen, with indications of the paths of light-rays. In A we see at xy the section of a small disc of very thin, plain glass placed at an angle of 45° across the axis of the microscope, directly opposite an aperture in the side of the tube. The light enters the instrument at this aperture and falls upon the disc of glass, which allows the greater part of the light to pass through ; a certain small proportion, however, is reflected by the surfaces of the shp of glass, and is thus directed downward upon the back of the objective lens of the micro- scope, shown at in the figure. This lens condenses the light and produces a relatively bright spot of light upon the surface of the specimen. The light reflected by the surface of the specimen is collected by the objective and passes upward on its way to form the image in the instrument ; in its upward path, however, it again encoimters the thin shp of glass and, although this time the greater part passes through unchanged to the image above, a part is reflected and returned to the source of Ught outside the microscope. It will be seen that this method of illumination is very wasteful of light, but for ordinary eye observation no great intensity of illumination is really wanted and for photographic purposes a very powerful source of light — such as an arc lamp — can be used. Apart from the question of intensity of light, the use of such a plain glass reflector has the advantage that it is possible to obtain perfectly axial (central) illumination, with no tendency to throw shadows towards one side more than another. On the other hand, all the rays that go to form the image have to pass through the slip of thin glass, and this undoubtedly affects the perfect sharpness of the image, so that for the highest magnifica- tions and the greatest clearness the use of the glass reflector is inadmissible. Another and perhaps more widely used form of reflector is shown in Fig. 3, 6. In this case the line pq indicates THE METALLURGICAL MICROSCOPE 41 the section of an opaque reflector, such as a sUp of silvered glass, or what is known to opticians as a " totally reflecting prism," placed not centrally in the tube of the microscope, but in such a position as to cover one half the area of the tube. Of the Hght that enters the aperture at the side of the instrument, nearly all that portion which falls upon the reflector is sent down through the lens and thus on to the specimen. If the beam of Ught is properly directed, the resulting spot of hght on the specimen can still be circular — although originating from a semi -circular reflector — and of very nearly equal bright- ness all over. In practice, however, it is very difficult to secure perfectly even illumination by this means unless the reflector itself can be moved about in all directions ; usually there is a decided falhng off in brightness towards one side of the field of view. From the specimen, the rays of hght in this system pass upwards through the objective ; all those rays, however, which come from one half of the objective are stopped and sent out- wards by the reflector and are thus lost to the image, but the remaining half of the rays are amply sufficient to form a very perfect image which is free from the disturbing effect of the thin glass shp which is present in the other system. While this system has thus the advantage of giving more perfectly -defined images, it has the disadvantage of less uniform illumination and of somewhat one-sided lighting, as the beam can never be sent down upon the specimen in a strictly axial direction. It will thus be seen that both systems have their special advantages, and the true solution of the question as to which should be used is to employ an arrangement whereby both kinds of reflectors can be used and readily interchanged with one another. The methods of lighting metal specimens which have just been described carry in their train certain mechanical require- ments in the construction of the microscope which have been met more or less completely in those instruments which have been specially designed for this work ; these are generally known as " metallurgical microscopes." One of the simpler of these instruments is shown in Fig. 4 (Plate II.) . This instrument is, in many respects, very similar to the simpler forms of ordinary microscope employed for biological and other studies ; THE METALLURGICAL MICROSCOPE 43 The type of microscope just described suffers from a number of serious disadvantages which become evident in use — particu- larly when the instrument is to be used for the study of metals from the mechanical rather than the chemical point of view. For the purpose of overcoming these disadvantages, the present author has designed a special form of metallurgical microscope, which is shown in Fig. 5 (Plate II.), Both in appearance and in mechanism this instrument differs from the older forms, but the optical system is the same and the same lenses are employed. To begin with, the stand is of a rigid form, which avoids the tendency to vibration and displacement found in other micro- scopes — ^the instrument can be carried about without risk of disturbing the focus of the lenses ; for work with the highest magnifications this is very important, and where measurements are to be made it becomes vital. This rigidity is further assisted by the fact that the body tube, 66, is rigidly attached to the stand or " limb," ss, all focussiag movements — coarse and fine adjustment — ^being attached to the stage. This arrangement has the additional advantage that the delicate fine adjustment is placed actually in the axis of the microscope, so that there is no overhang to magnify minute defects in the slides or screw ; the stage itself can be completely rotated and presents a clear surface with no projections to obstruct beams of oblique light in any position. The milled heads which actuate the mechanical movement of the stage and the fine and coarse focussing screw are, by this arrangement, placed within comfortable reach of one hand, and the operator soon learns to work these screws together with different fingers in such a way as to keep the image in sharp focus while the specimen is being moved about beneath the lenses. This arrangement greatly facilitates the systematic examination of large areas, which is often required. Another important feature of this microscope is the arrangement of the illuminator ; this is not a separate fitting to be screwed to the end of the tube — that arrangement involves an extra risk of unsteadiness and also makes it im- possible to interchange different forms of illuminator without removing the lens entirely. In the author's microscope the objective is screwed direct to the end of the body-tube, and the 44 STUDY OF PHYSICAL METALLURGY illuminator — or rather the reflector — is carried on a small fitting which slides into an aperture in the lower end of the body-tube ; this is so arranged that the reflector can be moved at will both across the tube and along it, thus enabling the operator to find the position which best suits the particular lens he is using — and in practice this position is found to differ considerably with different types of objective. While this adjustabihty is of very considerable use to the expert worker, 1/ / virtual I W tmaae Objective \ fO&Ject Fig. 6. — Diagram of Image Formation in the Microscope. it is apt to be a little awkward for the beginner, as the latter will probably place the reflector in the worst possible positions to begin with. For that purpose, however, the instrument can be supplied with an illuminator, fitting into the same aperture but in a fixed position, which is reasonably good for some of the most usual lenses. Turning now to the optical system, we find that in the modern compound microscope — which is the instrument always understood when we speak of " the microscope " — the images are formed by two lenses, or systems of lenses, known as THE METALLURGICAL MICROSCOPE 45 objectives (or object glasses) and eye-pieces respectively. We need not enter into the details of their design or construction, but it vidll be well to consider some questions which affect their performance and their limits of usefulness. The diagram of Fig. 6 shows in a general way how images are produced in such a microscope. The objective, which is always a lens of very short focus, is placed close to the object and produces an image, I, I, at a certain plane in the tube of the instrument ; this image is already magnified in a definite proportion which depends upon the relative distances between lens and object and lens and image, the power or focal length of the lens itself determining what each of these two distances must be in order to yield a sharp image. Theoretically the actual choice of these distances should be immaterial, so long as the object to be examined is always kept outside the front focal distance of the objective. With the highly-corrected modem microscope objectives, however, the best results are only obtained if the lens is used so as to form its real primary image at one particular distance from the lens. This distance is determined by the actual length of the microscope tube and is generally spoken of as " the tube-length " — this determines the distance between the objective and the eye-piece. It is accordingly important that each objective used should be employed at the proper tube-length for which it has been designed. This length is usually marked on the mount of the objective. The primary real image produced by the objective is then examined by means of the eye-piece, which acts as a magnifying glass or — ^in the case of photo-micrography — as a projection lens. The point which is of importance here is that the per- formance of the object glass is of much greater importance than that of the eye-piece. The latter merely enlarges to a con- venient size or projects to a convenient distance the image produced by the former ; the detail and its sharp definition must aU be present in the initial image produced by the objective — the best of eye-pieces cannot improve the image, while the enlargement produced by the eye-piece serves as a severe test for the perfection of the image to which it is applied. The quaUty of the image produced by the objective depends 46 STUDY OF PHYSICAL METALLURGY upon the general excellence of the construction of the system of lenses — the completeness with which the optical defects, such as chromatic aberration, spherical aberration, and others are removed or " corrected," and also upon what is known as the " aperture " of the lens. As regards the optical corrections of his lenses, the micro- scopist is largely in the hands of the opticians, but there can be no doubt that the best of modern lenses, such as the " apo- chromatic " objectives of Zeiss, leave little or nothing to be desired in the completeness with which the defects inherent in simple lenses have been overcome. There are, however, circumstances in the use of the microscope which sometimes place the objectives under very severe conditions, and then the smaU residual defects become apparent. This can largely be avoided if the operator possesses the necessary knowledge of optics ; here we can only indicate two devices which are frequently useful, particularly where photo-micrographs of difficult subjects have to be prepared. One of these is the use of monochromatic light. Such light can readily be obtained either by isolating the light of one particular colour from a beam of white light by means of a prism or other apparatus which is used so as to split the white light into a reasonably wide spectrum, or a similar, although less perfect, isolation may be obtained by means of coloured solutions placed in glass cells in the path of the beam of light. A very efEective solution of this kind, giving a nearly monochromatic blue light very suitable for photographic work is obtained by preparing a saturated solution of copper-ammonium acetate. A still better approximation to monochromatic light may be obtained by the use of thin sheets of gelatine, mounted between sheets of clear glass. Such gelatine light filters, stained with suitable dyes, are now obtainable from Messrs. Wratten and Wain- wright, specially for use in photo-micrographic work. By the use of such monochromatic light it is possible to eliminate from the image the effect of a small amount of residual chromatic aberration which is sometimes found in microscope objectives, particularly in regard to the parts of the image lying away from the centre of the field of view. This chromatic aberration THE METALLURGICAL MICROSCOPE 47 arises from the fact that the lenses affect the light of different colour (i.e., of different wave-length) in a different degree, so that the images produced by the red rays, for instance, do not quite coincide with the images produced by the blue rays. In a simple lens this would be the case to such an extent that the resulting image would be badly blurred by blue and red edges, but the sldll of the optician has so balanced the action of different kinds of glass as almost entirely to overcome the formation of these coloiired edges ; by the use of light of a single colour, however, the formation of coloured edges can be entirely avoided. The use of such light has the further ad- vantage, especially for photography, that there is no longer any possibihty of a difference between the visual and the photo- graphic focus of the image. The rays of hght which most affect the eye are of much longer wave-length than those which principally affect the photographic plate, and consequently it sometimes happens, if white light is used, that an image which appears perfectly sharp to the eye is slightly out of focus when photographed ; if monochromatic light (preferably blue) is used, and the final focussing is done by that light, there is no risk of this occurring — since only the rays which have been used for focussing are then allowed to act upon the photographic plate. The other device to be mentioned here as an aid to obtaining the best results from a given lens is that known as " stopping down." Without entering into the somewhat intricate theory of the question, it may simply be stated that some of the defects which exist in optical images produced by lenses arise from the fact that the rays which have passed through the outer edges or zones of a lens are not readily brought into coincidence — at the image — with the rays coming from the same point of the object which have passed through the lens nearer to its centre. As a result, when the outer zones of a lens are covered, so as to cut off the hght passing through them, the sharpness of the image tends to increase, although the amount of hght which is transmitted by the lens is reduced and the image thus loses in brightness. This process of reducing the effective diameter of the lens is known as " stopping down," and produces two other 48 STUDY OF PHYSICAL METALLURGY important efEects besides that of eliminating the aberrations due to the outer regions of the lens. The first of these is a decided advantage, while the other entails such serious dis- advantages as to put a practical hmit to the application of the process. The advantageous effect of stopping down is to increase the " depth of focus " and also the " flatness of field " —in other words, to cause the image to appear sharp over a larger area. The manner in which a reduction of the effective diameter of the lens affects depth of focus can best be seen from the dia- gram of Fig. 7, in which the objective is represented by 00 ; p is a point in the object from which a bundle of rays, rrr, diverges, passing through the lens, 00, and being again con- q®-i = Fig. 7. — Relation of aperture to depth of focus. verged to the point P in the image. If we think of a second point, s, in the object situated at the same distance from the lens as the first point, p, then S, the image of s, will be situated at the same distance behind the lens as P, the image of p. If the lens is so placed that the points P and S fall upon the focussing screen or into the focal plane of the eye -piece, then the images represented by these two points will appear per- fectly sharp. But now consider a point, q, situated a little further from the lens than either p or s — such a point might represent one of the lower or deeper places in a somewhat deeply-etched specimen. The image of such a point will lie at Q, nearer the lens than P and S. The result wiU be that if the screen or focal plane is set to suit P and S, that part of the image representing points hke Q will not be sharp — ^the bundle THE METALLURGICAL MICROSCOPE 49 of rays coming from q does not meet at a point where it crosses the focal plane but constitutes a cone of small angle, 1(^1, and the section of this cone on the focal plane is a small circle or disc and not a point. So soon as these discs are of appreciable size, the image becomes blurred. Now the size of such a disc depends on two factors — ^the distance of the true image point, Q, in front of or behind the focal plane, and the angle of the cone of rays, aQa. In practice all those parts of the image appear to be sharp in which these little discs do not exceed a certain minute size (about "005 inch), and consequently an objective possesses a small " depth of focus " — ^that is, a range of distances for points on the object for which sharp images are still possible. This range of distances or depth of focus decreases rapidly as the width, aa, increases, but it can be very considerably increased by " stopping down." In the diagram, if the lens be covered by a stop in such a way that aa is reduced to a' a', then the angle of the cone of rays meeting at an image point is reduced to a'Qa', and the distance at which Q may be situated from the focal plane without causing a blurred image is very considerably increased. The same effect makes itself felt in yet another way. In the diagram of Pig. 7 we have supposed the points p, q and s to lie near that part of the object opposite the middle of the lens, aa, but in the case of microscope lenses the objects to be examined are often large compared with their perpendicular distance from the lens itself and we then find a systematic variation in the distajice from a point on the object to the centre of the lens, arising from the fact that the rays of Ught travel in an increas- ingly oblique direction when they come from points at increasing distances from the axis pf the instrument. The result is that the image of a perfectly flat object assumes a concave form, the image -points representing the outer edges of the object lying nearer the lens as compared with those which represent the central points of the object.^ This is known as the 1 This account of the causes of " curvature of field " is very elementary, and is true only for an ideal simple lens ; in actual lenses the combined effects of the components sometimes reverse the direction of curvature. P.M. B 50 STUDY OF PHYSICAL METALLURGY "curvature of the field " and results in a difficulty which makes itself felt whenever microscopic images are to be photographed or even merely projected upon a flat plate ; when the image of the centre of the field is sharply focussed the outer edges are blurred. The image-points of the outer portions of the field of view Me in front of or behind the plane of the plate or screen and the cones of rays where they intersect the plate produce discs instead of points. These discs can again be reduced in size, and the sharpness of the image extended over a larger area, by stopping down. With these considerations in view it would be natural to suppose that the best results could be obtained by reducing the effective diameters of our lenses to mere pin-holes. The student will be readily cured of this idea if he will try the experiment of focussing a good microscope on an object showing some rather small features, commencing with the iris of his illuminator wide open and gradually closing it. At first the image will certainly improve in definition and contrast, although rapidly diminishing in brightness, but when the aperture is reduced below a certain amount the appearance of the image begins to change ; narrow lines spread out into bands, minute points into circular or irregular patches, and every dark edge is bordered by alternate dark and fight bands. This change in the image is due to what is known as diffraction, which renders the images obtained with unduly small apertures misleading and useless. The reason for these effects is to be found in the undulatory nature of light and the mutual interfer- ence of Mght waves when these are caused to pass through small apertures. Interference always occurs, whether the small aperture is there or not, but the presence of a small aperture cuts off the other light -waves which previously obhterated the effect by " averaging up " the illumination ; when these other waves are stopped by the iris or the edges of the lens, the interference or diffraction effects become very prominent. That they are always present more or less can be readily verified by using eye -pieces of very high magnifying power ; when this is done minute rings and bands can be seen sur- rounding the dark parts of the image even when large THE METALLURGICAL MICROSCOPE 51 apertures are used. Here, then, we have a very important practical limit to the extent to which the image given by a lens can be improved by stopping down — ^this process must never be carried so far as to bring about any very marked change in the character of the image. It has already been pointed out that the effects of diffraction are always present in microscope images and, from what has just been said, it will be seen that the actual diameter of the objective lenses — or their " aperture," as it is technically caUed — ^is of vital importance in governing the degree to which the resulting images are affected by these disturbing influences — ^for the best results, lenses of large aperture must be employed and, indeed, it is the aperture of the objective which governs the amount of magnification which can be usefully employed with it. We have already seen that an objective of given focal length used in the ordinary way gives a real image magnified in a definite ratio ; in the case of very short-focus lenses, such as the so-called oil-immersion objectives, this initial magnifica- tion may be as high as eighty or one hundred times. But this initial image may further be viewed with a magnifying eye- piece, or even with a second complete microscope, and thus magnified to almost any desired extent, although with a single eye -piece a further magnification of eighteen or twenty times is the most that can be conveniently applied. The question is, how far is it worth while to carry this magnification ? The existence of the disturbing effects of diffraction makes it evident that there must be a hmit of useful magnification for any lens ; the mere enlarging of an image is of no advantage unless additional detail is thereby rendered visible, and such detail can no longer be seen or " resolved " when the distances between the images of adjacent features of these details are so close together that the diffraction bands or rings formed around the image of one feature overlap and blur the image of the other. With the best modern apochromatic objectives this limit of resolving power is attained when the total magnification of the image reaches 1,500. This figure — sometimes expressed as 1,500 diameters — indicates that the length of any line in the object is multiplied 1,500 tiraes as seen in the image. Increas- £ 2 52 STUDY OF PHYSICAL METALLURGY ing magnification beyond this point, whether by direct projec- tion or by means of photography, yields no additional informa- tion and only shows the diflEraction figures more clearly. This whole question of " resolving power " is too large to be considered fully here, but the knowledge that such a definite Hmit exists, and of the nature of the causes affecting it, is very necessary to every microscopist if only on account of the caution which it should inspire in the interpretation of the images seen under the highest magnifications. The correct interpretation of these images, especially when the peculiar character of the illumination employed in metaUographic work is considered, is a matter requiring some care and caution. A striking example of this kind of difficulty is found in cases where a small area on a specimen is obviously at a different level from the rest ; in such cases some care is required to determine whether we are dealing with a mound or a pit in the surface. This question is further comphcated by the fact that we see such images right and left reversed, and it is not easy to interpret shadows. An adjustable reflector in the illuminator is of great assistance in such cases, but the final resource is the micrometer screw of the fine adjustment of the microscope — by observing whether the object and lens have to be moved closer together or further apart to obtain focus at different levels, questions of difference of level can always be settled. A more difficult case arises when features are under observation which lie close to the hmit of what the microscope can distin- guish. Thus in mild steel, when treated in certain ways, a series of structures can be developed ranging from beautifully laminated " pearhte " on the one side to a dark, almost feature- less constituent on the other. The extreme cases are, of course, quite distinctive, but the intermediate gradations are not easily distinguished. Where fine details, and particularly laminations, have to be resolved the influence of the opaque reflector (Fig. 3 B, p. 39) must be borne in mind, since its presence reduces the working aperture, and therefore the resolving power to one half in one direction. Thus, if laminations are placed parallel to the edge of the reflector, they may be completely blurred, while THE METALLURGICAL MICROSCOPE 53 they appear sharp and clear when turned through a right-angle. This case is particularly instructive, as showing the care re- quired in interpreting such images and the need for a rotating stage on metallurgical microscopes in order that the specimens may be examined in every azimuth — ^no specimen showing fine detail should be examined without passing through a complete rotation of the stage and subjecting it to varying iUumiaation by moving the reflector of the illuminator — shghtly oblique illumination often reveals details which are invisible in the full glare of central illumination. The general considerations of the conditions which govern the resolving power of the microscope lead to certain definite rules in the use of the instrument which may be summed up thus : For a given magnification it is desirable to obtain as much of the magnification as possible by means of the objec- tive — for most purposes an eye -piece magnification of eight or ten times is the best. The objectives used should possess a large aperture, and for extremely high magnifications a lens of larger aperture is preferable, even if its focal length is somewhat greater and its initial magnification consequently somewhat less. Finally, to obtain the best resolution in examining minute objects, the aperture of the lens should not be unduly diminished by the use of the stop or iris, and when the stop is closed down very much the observer must expect serious interference from diffraction effects. For most purposes of visual microscopy on opaque objects it is found desirable so to regulate the beam of light entering the microscope that about two-thirds of the full diameter of the objective is utiMsed. This can be readily judged by removing the eye- piece and seeing how much of the back lens of the objective appears to be " filled " by the beam of light. In order to obtain the best results from good microscope lenses it is essential, beyond the observance of the guiding rules just given, to secure a proper illumination of the object. The manner in which opaque objects are illuminated by reflected light has already been indicated in general terms, but the actual illumiaants and their best arrangement deserve some attention. The best methods of lighting differ somewliat according to the S4 STUDY Oi? PHYSICAL METALLURGY Eye Piece circumstances, i.e., whether the object is to be examined visually or whether a photograph is to be taken. For purposes of visual examination the source of light should not be very minute, but should preferably possess an equally illuminated area about one haK-inch square at least. Perhaps the most satisfactory source of Hght for this purpose is produced by means of an incandescent gas-mantle placed inside a white opal chimney. A somewhat similar effect can be obtained with a metaUic -filament electric lamp having a large number of loops of filament and provided with an opal bulb or an opal shade placed close to it. The illuminated surface of the opal glass is then to be regarded as the real source of hght, and the best manner of utilising this light is to place the front surface of the opal glass in the position indicated on the diagram, Fig. 8, by the letter 0. In this diagram the point F repre- sents the focal plane of the eye- piece and I and m are the distances of F and respectively from the reflector, R, of the vertical illu- minator. If now O is so placed that m = I, then a sharp image of surface SS of the specimen when IS brougnt to focus. The reason for this Fig. 8. — Diagram of lUuinination. ' Critical O will be formed on the the objective is brought result Ues in the fact that when the objective is in focus the image of SS is formed at F, and, as all optical processes are reversible, it foUows that the lens would form an image at SS of any object situated at F ; but is situated not at F, but at its image in the mirror of the reflector, and consequently the objective forms a sharp image of O on the surface of the specimen. This condition gives what is known as " critical illumination," and is far the best mode of Ulumination possible, particularly as it does not require the intervention of any condensing lenses or other elaborate appliances. Two points only have to be remembered: since a sharp image of the THE METALLURGICAL MICROSCOPE 55 source of light is formed on the specimen, the source of light itself must be uniform and free from detail, which would other- wise be superposed on the structure of the specimen in a confusing and irritating manner. For this reason a bare incandescent mantle or incandescent electric lamp, Nernst lamp or similar illumiaant cannot be employed, nor can ordinary ground glass be used in place of the opal glass, since the details of the surface of the ground glass are seen very clearly under the microscope in these circumstances. A second consideration arises from the fact that the lamp with its opal shade must be placed quite close to the microscope in order to fulfil the require- ment m = i, and if an unprotected or uncovered lamp is used the general lighting of the microscope and its surroundings becomes unpleasantly intense, while direct light from the lamp may reach the eye of the observer at the eye-piece. In order to avoid such serious inconvenience it becomes desirable to enclose the whole lamp in a metal or other opaque chimney, provided with an aperture on the side opposite the C microscope and at a suitable level to allow n MM 100 00( light to reach the opening of the illumi- Fig. 9. — Diagram of nator of the instrument. If arranged in ^crosTOpe °^ this way this mode of lighting not only gives the best possible illumination of the specimen, but also affords great comfort to the eye of the observer. The author has used several simple forms of lamp for this purpose. In one of these an ordinary upright gas burner with incandescent mantle is placed inside a vertical brass tube. At a suitable level an aperture is provided, and this is fitted with a side tube about one inch in diameter, running out at right angles. The light from the opal shade of the gas burner passes along this side tube and, if the lamp is so placed that the open end of this tube is close to the aperture of the illuminator of the microscope, the desired result is at once attained, as the length of the side tube is so adjusted that it acts as a distance-piece for focussing 66 STUDY OF PHYSICAL METALLURGY the light. The arrangement is shown in the diagram, Fig. 9. For electric light the author uses the arrangement shown in the diagram of Fig. 10, where N is the glower of a Nernst lamp, while R is a rod of optical glass having one end finely groimd. The hght entering this rod at one end is repeatedly reflected at the periphery of the rod and, finally, leaves the rod at the other end. This end of the rod then acts as a uniformly illuminated source of light, and may be so placed as to yield " critical illumination," while the arrangement of tubes shown in the diagram is sufficient to screen the light from the eye of the observer. The light from this lamp is rather bright and some- what yellow in colour ; it is easily softened and rendered whiter by the interposition of a screen of glass having a slight blue tint. For photographic purposes this mode of illumination is not, unfortunately, avail- able at present ; even when a piece of opal glass is illuminated by the concentrated beam from a powerful elec- tric arc lamp, the Hght which it emits is not sufficiently powerful for photographic purposes ; the images formed upon the ground glass of the camera are too feeble to be readily or accurately focussed, and thus much of the advantage of this form of hghting would disappear. It is, in fact, in order to facihtate the arrangement and focussing of the image upon the ground glass or other screen of the camera that extremely power- ful illuminants are to be recommended for purposes of photo -micrography ; the photographic plate itself is amply sensitive to take good photographs with very feeble illumina- tion. The illuminants actually used are, as a rule, either the electric arc or the limelight. Both these sources of light consist of small areas of highly heated incandescent matter Fig. 10. — Electric Microscope Lamp. THE METALLURGICAL MICROSCOPE 57 giving out a powerful beam of light. These minute bright areas cannot well be employed by placing them in the position of O in Fig. 10, since the images which they would yield upon the surface of the specimen would not be large enough to cover the field of view, while even the area covered by the image would not be uniformly lighted. The plan is, therefore, adopted of concentrating as much as possible of the light emanating from the limelight or the arc by means of a condensing lens into a parallel, or nearly parallel, beam of light. This beam is then passed through a water jacket in order to arrest the heat-rays which accompany the light coming from such a source, and then the beam is further concentrated by means of a lens until a small but intensely briUiant image of the craters of the arc, or of the limelight, is formed. The most favourable arrangement is such that this image of the source of Mght falls upon the iris of the vertical illuminator. Theoretically the best point would be the " back principal focus " of the objec- tive, since in that case the light would leave the lower face of the object-glass in an approximately parallel beam, and would then be reflected as such from the face of the specimen back into the lens. In practice, however, it is not worth while to locate this point with any great care, since the manner in which microscope objectives are constructed does not render them capable of emitting a satisfactorily parallel beam of light even in the best circumstances, and the point for focussing the light which has been suggested above is sufficiently near the theoreti- cally correct point to give practically equally good results. At all events a patch of very brilhant illumination will be seen upon the specimen immediately beneath the objective, and by this means images of such brightness can be produced that they can not only be photographed and focussed with the greatest ease, but they can even be projected upon a screen and ex- hibited to audiences of moderate size. This method of showing microscope images by projection is of great use whenever it is desired to explain the detailed features of a given micro - structure to two or more persons, but it must be remembered that in clearness of definition and width of view the projected image cannot approach the image as seen visually under the THE METALLURGICAL MICROSCOPE 59 that enters the microscope and the rays that leave the objective to form the magnified image are reflected through a right -angle. The appearance of such an instrument is shown in Fig. 12, Plate III. There can be no doubt that these microscopes offer certain advantages which make them particularly convenient for quick working at moderate magnifications. The specimen requires no mountitig or levelling, but on the other hand, as the poHshed surface lies on the stage, injury by scratching is very apt to result. The multiplication of reflecting surfaces is also a disadvantage, which makes itself felt at the highest magnifications, and for this class of work the author has not found this type of instrument very successful. Another type of instrument which is often useful, especially for the examination of fractures, and of small objects generally, by oblique illumination only, is the " Greenhough " stereoscopic binocular microscope, in which the specimen is looked at through two convergent microscopes. The stereoscopic effect obtained is very beautiful, and for magnifications up to about seventy diameters the instrument 'is very useful. For the ordinary type of metallurgical microscope (as distinct from the Le Chatelier inverted type) it is usually necessary to mount the specimen on a slip of glass, wood or metal in such a way that when this slip is laid on the stage of the microscope, the etched surface of the specimen shall lie exactly at right angles to the optic axis. Most of the older appUances used for thus " levelling " the specimens involve laying the specimen face downward on some level surface, and then adjusting the carrying slip parallel to that surface and attaching the specimen to it either by wax, plasticine, seahng wax, plaster, etc. Among these is the device formerly used by the author, in which the mounting slip is brought down upon the back of the specimen by a parallel motion guided by four accurately-made links. A much more perfect device, not requiring the polished surface to be touched in any way, has, however, been devised and adopted more recently. In this device a telescope is used with an internal reflector, somewhat like the clear -glass illuminator. Light from a suitably placed lamp is sent down the telescope tube and leaves the objective as a parallel bundle of rays. 60 STUDY OF PHYSICAL METALLURGY These strike the polished surface of the specimen lying upon the stage of the instrument, and are reflected back into the telescope. Matters are so adjusted that if the surface of the specimen is accurately at right angles to the axis of the tele- scope the observer at the eye-piece would see the image of the source of light reflected by the specimen falling accurately upon the centre of a pair of cross-wires placed in the eye-piece of the instrument. As a rule this will not be the case when the speci- men is first looked at, but if it is mounted on soft wax or plasti- cine, its position is easily adjusted with the fingers untU the image of the source of light faUs exactly on the right place. Li this way it is possible very quickly to level the specimen to a very high degree of accuracy. The arrangement has the advan- tage that it can be adjusted to work with a particular micro- scope in such a way as to aUow for any want of complete accuracy in the adjustment of the stage of that microscope. A further advantage lies in the circumstance that if a specimen has a rounded surface, so that the whole of it cannot be set at the proper angle, the setting can be made for any particular small portion from which a photograph is to be taken. For that purpose it is only necessary to cover the surface of the specimen, while it is being adjusted, with a piece of paper or other mask in which a small hole has been cut in such a position as to expose the small area in question. References. (1) Rosenliain. Journ. R. Microscopical Soc, 1906, pp. 146 — 155. (2) Le Chatelier. 6tude des Alliages, 1901, 421. The MetallograpMst, Vol. IV., January, 1901, No. 1, p. 1. For other papers on Metallurgical Microscopes see also the following : Martens and Heyn. " Mitteilungen der Konigl. Technlschen Versuchsanstalt," 1899, XVII., 73. Stead. Journ. Iron and Steel Inst., 1908, III., 22. Jouin. Iron and Steel Inst., 1897, I., 42. Benedicks. Metallurgie, 1909, VI., 320. CHAPTER IV THE MICRO -STRUCTURE OE" PURE METALS AND OF ALLOYS When a specimen of any pure metal, in either the cast or the " annealed " state, is prepared and examined under the micro- scope in the manner indicated in the previous chapters, a typical appearance is always seen — an appearance typical of the structure of pure metals as a class rather than of any individual metal. The appearance thus presented by the purest iron (good Swedish charcoal iron) is shown in Fig. 13, Plate I., and also in Fig. 97, Plate XX., where the structure is photographed under a magnification of 150 diameters. The appearance of other pure metals is very similar ; and, indeed, the similarity of all pure metals as seen under the microscope serves to show that microscopic examination cannot generally be employed in place of chemical analysis for the purpose of distinguishing different metallic elements from one another, for although the practised observer might be able to distinguish some of the more typical metals by the aid of minor peculiarities, the similarities between many of them are so great that much uncertainty would remain. On the other hand, in many cases, as will be evident later, microscopic examination will at once decide whether a given specimen consists entirely of one metal or is an alloy of two or more. For the present, however, the important fact is that all pure metals, when free from the disturbing effects of mechanical treatment, show a strikingly similar appearance under the microscope. As the figures show, this " structure " consists of a number of mutually adjacent roughly polygonal areas, separated from one another by narrow dark lines, but otherwise practically free from any marked features. These polygonal areas are readily recognised as representing a sectional view of roughly polyhedral grains, which evidently constitute the entire mass of such a metal. 62 STUDY OP PHYSICAL METALLURGY Microscopic research has, however, elucidated the nature of these grains, and they are now recognised to be true crystals, the metal being thus an agglomerate of crystals, much as a mass of rock-salt or of granite is an agglomerate of minute crystals. It is true that these crystals lack one of the most striking features of many crystals, and that is the regular geometrical outline which we see in specimens of Iceland Spar or Rock Crystal ; but this absence of geometrical form is due to the fact that none of these crystals have had the opportunity for free growth which is requisite for the development of geo- metrical forms ; when a crystal of a salt is formed in a solution its surfaces remain free to grow in the liquid surrounding them, but in the crystalHsation of a metal the growth of each crystal has been stopped by the interference of a neighbouring crystal, so that each of these grains is bounded, not by any regular geometrical outline, but by the more or less irregular surface upon which two adjacent crystals have met in the process of their initial formation. The manner in which this occurs may be made clearer by an analogy with the way in which children's building blocks might be piled up to cover a given area. If the blocks were all exactly alike in size and shape it would, of course, be possible to lay them all exactly parallel to one another in a single continuous pattern all over the surface to be covered. But for this purpose it would be necessary to begin the process at a single point and to work outward from that alone. If the work were begun by several operators at the same time, each starting from their own chosen centres and working each to his own scheme, without reference to that adopted by any of the others, a very different result would follow : each man would build up a regular pattern, and all the patterns would be essentially alike, but they would not be parallel or " similarly oriented," so that where they met they would form irregular joints. The diagram of Pig. 14, Plate IV., shows five stages of such a process, and the final result, with the blocks themselves left out and only the outlines of the joints traced where the various differently-oriented areas have met. This outline is exactly like the boundary between adjacent crystals in the micro -section of a pure metal — it is the " acci- PLATE IV. (a) {b) H ^^uiL '^'' ' "^^^l (<;) {d) (e) Fig. 14. {/) [To face p. 62. PURE METALS AND ALLOYS 63 dental " meeting-line of adjacent growing crystals of different orientation. The essence of crystalline character does not reside in geometrical outlines or shapes, but in the regular arrangement of the molecules or groups of molecules within the mass. The geometrical outline is merely one result of that internal struc- ture — a result which is only apparent in favourable circum- stances. This regular arrangement or orientation is fully present in the crystals which constitute the mass of any specimen of pure metal. Some of the evidence upon which this statement is based is so striking and interesting that a brief account of it is desirable. The photo -micrographs typical of pure metals given in Figs. 13 and 97 are, as already indicated, taken under normal illumina- tion, but these same specimens yield a different appearance when viewed under oblique light. The characteristic beauty and lustre of this appearance must be seen direct through the microscope to be appreciated ; Fig. 15, Plate I., gives a photographic representation of the same field of view as that shown imder normal light in Fig. 13. The polygonal outMnes of the crystals are still visible, but now they appear as bright lines on a dark background ; the areas of the polygons, which in the former view appeared almost uniformly white, are now greatly differentiated ; some appear perfectly dark, while others shine out brilliantly. If a specimen thus lighted be slowly rotated on the stage of the microscope, a very striking phenomenon is witnessed — the crystals which at first appeared bright rapidly wane and become dark, while others flash out brightly. This phenomenon may be briefly called " the rotation effect," although it is more correct to refer to the " oriented lustre " of metals. To understand its nature and meaning we must consider for a moment the manner in which the surface of the specimen has been prepared. The important point in the preparation for the present purpose lies in the etching process, since before that was applied the surface of the metal was practically a uniform plane. Now the etching reagent — such, for example, as dilute nitric acid — ^is, as we have seen, merely a weak or slow solvent of the metal, and it effects 64 STUDY OF PHYSICAL METALLURGY this solution by gradually unbuilding or taking down the structure of the metal exposed to its action ; if now the different crystals which have been cut through by the plane of the polished surface are each built up on separate systems or arrangements of their own, each will also be unbuilt according to this system. If this unbuilding process be stopped at any given moment, the part of the surface belonging to each crystal wiU be covered with some more or less distiact traces of the plan upon which the unbuilding was proceeding. These traces actually consist of a number of very minute facets, differing in size and position, but all of the same shape and similarly oriented, i.e., all facing the one way over the entire part of the surface belonging to the same crystal, but — as a rule — differing in shape and orientation from one crystal to another. Further, it is found that the unbuilding of the crystals by the process of gradual solution does not go on at the same rate upon aU the crystals exposed in a section — in fact the rate of attack differs shghtly, but distinctly, from each crystal to its neighbour, according to the different orientation of the molecules or groups of molecules of which the crystal is buUt up. The etched surface, therefore, consists of a series of polygonal areas, corresponding to the various crystals intersected by the surface, each at a slightly different level, and thus necessarily connected by short, steeply-sloping surfaces, and each covered with a system of very minute facets which are similar and similarly oriented over each of these polygonal areas, but differ in shape and orientation when we pass from one area to another. What, then, are the microscopic appearances to be derived from a surface of such configuration ? Under normal illumination, if the etching has not been very deep, the facets just referred to are too minute in depth to be readily seen except under the highest magnification, but the differences of level between the adjacent polygonal areas become very markedly evident on account of the short, steep connecting surfaces. These are so steeply inclined to the horizontal that light falhng downward upon them, instead of being reflected back into the lens of the microscope, is thrown out to one side, and the short surface thus appears, as seen from above, as a PURE METALS AND ALLOYS 65 fine black line ; these are the black lines which bound the poly- gonal areas of the crystals. When examined under oblique illumination, however, these steep surfaces, or at aU events those of them which face the direction from which the light happens to be coming, catch the obliquely incident light and reflect it into the objective, so that these bounding surfaces now appear as narrow bright lines. The diagram of Fig. 16 will make this clear, if it is remembered that only those surfaces will appear bright to an observer looking into the eye -piece of the microscope which reflect rays of light into the objective of the instrument. From these diagrams it will readily be seen that the comparatively hori- zontal surfaces of the polygonal areas them- selves will, as a rule, appear dark when viewed under obUque illumination. There is, however, an important exception to this rule, and this occurs when the minute facets with which each of these areas is covered are so ^ig. turned that they catch the incident light and reflect it into the objective. Among the multitude of crystals of differing orientations found in a single field of view, this will happen here and there in almost any position of Hght and specimen, with the result that several crystals wiU generally shine out brightly in these circumstances. This brightness wiU naturally differ somewhat from one of these crystals to another, according as the facets direct the whole of the hght full into the objective or so direct it that only a portion is caught by the lens. As the specimen is rotated under the obhque illumination, the facets which at first reflected light into the microscope are gradually turned in such a way as P.M. F 16. — Diagram of normal and oblique rays falling on an etched surface. PURE METALS AND ALLOYS 67 next. A striking example of this kind is given in Fig. 19, Plate V. With this brief outline of the evidence, the reader must be asked to accept the view — ^now generally recognised — that the polygonal grains seen in pure metals are true crystals so far as their internal structure is concerned. We should perhaps hint at once that the somewhat natural tendency to regard crystals as necessarily brittle in their nature is a misconception based upon ordinary experience of such crystalline bodies as sugar or spar ; recent discoveries have, in fact, revealed the existence of crystals so plastic as to be practically fluid, so that they are usually called " liquid crystals," while among metals the crystalline character is perhaps most readily observed in one of the softest and most plastic, viz., lead. We now turn to the question of the manner in which the crystalline structure of a pure metal is originally formed. At all events in the case of artificially prepared metals, excepting only those prepared by electro-deposition, the genesis of the micro -structure is to be sought in the process of sohdification from the molten or fluid state. Judging from the anaJogy of the process of freezing as we see it in the case of water or other pure hquids, which can be more readily observed than molten metals, the course of the process is somewhat as follows : as the temperature falls no very definite change occurs until the freezing-point is reached ; then at a number of points in the liquid, solidification com- mences — ^minute crystals are formed, and each of these rapidly grows in size. As a rule, this growth does not take place simply by the accretion of concentric layers ; the crystal is seen to throw out a number of arms into the surrounding hquid, from these arms numerous secondary branches are thrown out, and these in turn throw out other spines. The familiar figures made by the frozen moisture on window panes in winter are examples of this process of " dendritic " crystallisation, only that on the window pane the process has been arrested at an early stage for want of further material, ti the interior of a mass of water — or of molten and freezing metal — there is no lack of material, but there are a number of crystals growing f2 68 STUDY OF PHYSICAL METALLURGY out from adjacent centres which vie with one another in appro- priating both the material and the space in which it exists. In this way the crystalline arms which are thrown out by adjacent centres soon meet and thus put an end to further growth, the further development of the process then taking the shape of a gradual fillin g in of the space occupied by the dendritic arms which formed a species of advance guard for the final progress of the fully -formed crystal. The actual rate of progress of the arms of each of these growing crystals is in reality determined by the rate at which heat is abstracted from the different parts of the mass. Thus a crystal arm which grows outwards in a direction away from the coolest and towards the hottest part of the liquid will not readily meet with any other crystal arm, since in the hotter portions of the mass crystaUisation has not yet begun — an arm growing in that direction will, therefore, grow on without interruption for a much longer time, and there- fore for a much longer distance, than an arm which endeavours to grow in a direction at right angles to the direction of the flow of heat. It has already been indicated that the dendritic arms which grow rapidly outward from the centres or nuclei of crystallisa- tion, which are formed when a molten pure substance cools below its freezing-point, continue their growth until they are checked by encountering similar dendritic arms emanating from adjacent centres. Similarly, when the interstices of the dendrites come to be filled up by the growing crystals during the completion of the solidification or freezing process, this growth also continues until stopped by meeting with the advancing edge of an adjacent crystal. This process has b^en illustrated diagramatically in Fig. 14, Plate IV., but that figure also illustrates a further point. Where the two systems of cubical blocks meet, a series of interstices are necessarily left, each vacant space being too small to allow of the introduction of an additional block. One may well ask if there is anything of the same sort where two crystals meet, and if so, how the adjacent crystals adhere to one another ? The view that there really are such interstices has been held by a number of in- vestigators, but the author cannot accept it, for several reasons. PURE METALS AND ALLOYS 69 Among these may be mentioned the fact, to be considered more closely in a later chapter, that the cohesion across the bounding surfaces of two crystals is actually stronger, in normal pure metals, than the cohesion across any of the surfaces within the mass of a crystal, so that when such a crystaUine aggregate is broken, the crystals are not pulled apart from one another, but are actually broken across. Normally the inter - crystalline boundaries, in fact, behave hke strengthening ribs, and not as surfaces of weakness. It would be hard to account for this fact if the existence of actual interstices were assumed, but one may none the less assume that there are interstices between the adjacent crystals, but that these are filled with matter which has not become crystalline. This hypothesis leads to a series of most interesting con- clusions, and some of these the author and his colla- borators have been enabled to verify experimentally. At the present time, however, it cannot yet be claimed that the theory of the existence of an amorphous or non-crystaUine cement between the adjacent crystals in a metal has received general acceptance. Its bearing on the behaviour of metal when undergoing deformation or fracture will, however, be discussed when that portion of the subject comes to be dealt with. Here it is only necessary to draw attention to the special conditions which exist at the crystal junctions. The process of sohdification as an aggregate of crystals all having the same chemical composition and all formed at one particular temperature {i.e., the freezing or solidifying point) is tjrpical of pure substances, i.e., of substances which are chemically homogeneous and are neither mechanical mixtures nor solutions. It is particularly tjrpical of pure metals, and it follows that the micro -structure of aU pure metals will show a strong similarity. This is true of the general type of structure — as we have already seen — ^but there are some minor difEerences due to the peculiar properties of each metal. For instance, the size or scale of the entire structure habitually varies enormously between different metals, although it may be accepted as a rule that the size of crystals depends upon the rate of cooling during the process of solidification. For this reason alone metals 70 STUDY OF PHYSICAL METALLURGY having a very high freezing-point, such as copper, silver, platinum and iron, will usually present a very minute structure, while metals which solidify at more moderate temperatures — such as tin and lead — usually possess a much coarser structure. By suitably modifying the conditions of cooling, the size of the crystals in any metal can, however, be enormously altered, a reduction of the rate of sohdification always resulting in the formation of larger crystals. In many cases, as we shall see later, other processes, such as prolonged heating at tempera- tures considerably below the melting-point, wiQ cause the crystals to grow, and may thus ultimately produce a structure on a larger scale than that which existed in the metal as first formed on freezing. It is important to note that a large and coarse structure invariably accompanies undesirable mechanical properties. For this reason the temperature at which metal is cast very materially affects its strength, for the rate of solidification is necessarily reduced if the metal is cast at an unduly high temperature ; in those circumstances the surplus heat of the metal has raised the temperature of the mould to an undesirable extent before the metal itself has cooled to the freezing temperature, with the result that the rate at which the mould is able to abstract heat is unduly diminished and the rate of cooling of the metal during the freezing process is thereby reduced. Apart from the mere size of the crystals which result from gradual solidification, the rate of crystallisation and the precise manner in which it occurs has a very iinportant effect on the mechanical properties of the metal. This is due to an effect which occurs at the boundaries of the crystals forming the mass of the metal, and arises from the greater or less degree of mutual interpenetration of the dendritic arms or branches which shoot out from the centres of crystallisation in the early stages of the process. Fig. 20, Plate VI., is a photograph of such inter- penetrating dendritic arms sent out from centres of crystallisa- tion — ^in this case not of a metal, but, for convenience of observa- tion, of a salt (ammonium chloride). Metals, however, as Figs. 21, Plate VI., and 51, Plate X., show, crystallise in a precisely similar manner. The extent to which this interpene- PURE METALS AND ALLOYS 71 tration occurs depends upon both the rate and the mode of crystallisation, and the strength of the metal and its power to resist external forces which tend to tear the crystals apart from one another will in turn depend upon the extent to which adjacent crystals are locked together by the mechanism of such interpenetration. Where there is much " interlocking " the bounding surfaces of the adjacent crystals will not be straight and smooth, but serrated and indented, thus providing the metal with a much greater length of junction surface and with the additional strength which appears to reside in these junctions. It is interesting to consider in this connection what factors other than rate of cooling, which has already been mentioned, are likely to affect the formation and interpenetration of dendrites. In the case of crystals forming in salt solutions it has been observed that the presence of quite small traces of an impurity is often enough to alter the crystalline " habit," and the shape and arrangement of the dendrites very con- siderably, even though the impurity does not enter into the composition of the crystals themselves. It seems probable that some such effect as this may also occur in metals, and that in this way we can account for the very large effects produced by certain minute additions to some metals. Thus the presence of less than O'l per cent, of vanadium in steel produces a dis- proportionately large effect on the mechanical properties, and yet there is no special micro -constituent which is formed as the result of the addition of this minute quantity of vanadium. In other metals similar, although rather less marked, effects can be observed. In some of these cases the effect of the added material on the micro -structure can be seen in a slight altera- tion in the appearance of the crystal boundaries ; the presence of vanadium in steel, for instance, renders the whole structure decidedly less distinct — an observation which lends some support to the explanation here suggested. Before leaving the subject of the crystallisation of pure metals it should be remarked that the crystaUine system according to which the molecules or groups of molecules within each crystal are arranged has been definitely determined in some cases. In 72 STUDY OF PHYSICAL METALLURGY the case of iron, for example, an examination of the etching figures leads to the view that the metal has crystallised in the cubical system, since aU the etching figures present shapes which are readily recognisable as sections of either cubes or octahedra, and these are the tjrpical forms of the cubic system. This view has been fully estabhshed by the work of Osmond {^) and his collaborators, who obtained actual crystals of iron by chemical processes which allowed these crystals to attain their natural external geometrical shapes. Other metals have received less attention, but in the majority of cases the evidence available is at all events consistent with the view that they also belong to the cubic system. Marked exceptions are, however, presented by cadmium and zinc, which have been shown to belong to the hexagonal system (*), and by bismuth and antimony. Leaving for the moment the subject of the micro-structure of pure metals, we turn to the much larger group of metallic bodies known as alloys. For the sake of simplicity, and in order to remain on ground which has now been fairly well explored, we wiU confine our attention for the present to alloys containing only two elementary metals — such alloys being known as " binary alloys " — the whole range of alloys between two metals being known as a " binary system." In order to understand the nature and constitution of alloys, it is preferable to commence our consideration with reference to the simplest state in which such bodies exist, viz., that of homogeneous fusion. Although certain pairs of metals cannot be made to mix in all proportions while in the molten con- dition, this is an exception, for in the great majority of cases two molten metals can be mixed with one another in the fluid state in any relative proportion. In this respect these fluids resemble such Hquids as water and alcohol, or water and sul- phuric acid. We may, in fact, safely carry this analogy much further, and regard mixtures of molten metals, i.e., all molten binary alloys, as simple solutions of the two ingredient metals in one another. Sometimes, it is true, the two ingredients of such a solution enter into chemical combination with one another, and in that case it might be more correct to regard PURE METALS AND ALLOYS 73 the resulting liquid as a solution of the excess metal and of the compound metal in one another, but this is a distinction of minor importance in the present connection. Since our interest naturally centres in the solid metal which results from the solidification of molten alloys, the question which Mes before us is : " What happens to the mutual solution of two molten metals when the temperature is so far lowered that the metal undergoes sohdification ? " The answer is that, apart from the formation of definite inter-metallic compounds, there are two opposite modes of sohdification adopted by binary alloys and a range of intermediate modes connecting these extremes. These extremes are simply (a) the case in which the state of mutual solution remains undisturbed by the process of sohdification and the alloy crystallises while still remaining a solution, i.e., each of the crystals which are formed ultimately attains the same average composition as the molten Hquid from which it was deposited — such crystalhsed solutions are usually termed " solid solutions." (The German term for these sub- stances, " Mischkrystalle," is sometimes translated — ^most un- desirably — as " mixed crystals " ; since this term suggests a mixture of crystals of different kinds, it is liable to convey an entirely erroneous impression, and should be avoided in the interests of clearness.) The other extreme (6) is that the state of solution existing in the liquid condition is entirely destroyed by the passage into the soHd state, and the two constituents separate during the process of crystallisation ; in this case the solid alloy can finally attain the condition of a mixture of crystals of the two pure metals. Intermediate between these two extremes are those alloys — and these are by far the most numerous — ^in which the state of mutual solution is partially maintained in the sohd state, i.e., alloys containing up to a certain Hmiting proportion of the second constituent crystalMse according to group (a) as sohd solutions without any separation of the two constituent metals, but alloys containing a higher proportion of the second constituent than the limiting propor- tion just referred to undergo a partial separation during freezing, the excess of the second metal above the Hmit of " sohd solu- bihty " being separated as crystals of a different constituent. 74 STUDY OF PHYSICAL METALLURGY These alloys resemble those of group (6), except that the two kinds of crystals present in them are not those of the pure con- stituent metals, but consist of the saturated solid solutions of each of the metals in the other. The micro-structure of any given alloy depends both upon the type to which the system of the particular metals in ques- tion m"ay belong, and also on the place of the alloy in its binary system. Alloys of group (a), of which those of gold-silver, iron- manganese, and copper-nickel are examples, when allowed to crystallise sufficiently slowly to attain their condition of final equilibrium, present a micro-structure exactly similar to that of a pure metal — the homogeneous character of the molten solution is in that case so completely maintained that the microscope cannot detect the presence of any second consti- tuent. As a rule, however, this state of complete homogeneity is not attained by solid solutions ; for reasons which will presently be explained, a certain degree of temporary separa- tion occurs during the freezing process, and usually persists to an extent readily recognised under the microscope. An example of this kind is shown by Pig. 21, Plate VL, which represents the micro-structure of an aUoy of copper and zinc containing about 25 per cent, of zinc (a kind of brass, therefore), which when very slowly cooled appears entirely homogeneous, but when rapidly solidified exhibits the dendritic structure shown in the figure. The characteristic feature of such structures, however, is that the arms or branches of the dendrites are usually ill-defined, and that they lie whoUy within polygonal boundaries which readily recall the crystal boun- daries of pure metals. The micro-structure of alloys of class (6) is rather more com- plex. Theoretically the addition to the pure metal, forming one end of a binary system of this type, of a minute trace of a second element should result in the appearance in the micro- structure of traces of crystals of a second constituent. In some cases this is true to a surprising degree of accuracy. Thus the addition of five parts of carbon to 10,000 parts of pure iron at once becomes visible under the microscope, while the presence of one part in 1,000 is readily seen, as shown in the photograph PLATE VI. Fio. 20. Pio. 21. V :',;,. ■:^/> -:v>v f^v. ^-.- ^-\'^-.* \.^-. /^ v^>;-:: ., * v' k „ >- i--' i'^ '.•"^W5! ^^^i^Jg^-Hs y 't^i » 1 > < / ' 1 \ ; ^ W _4' u 1 / ■r" , ) i •A V'-^ ^ _, t jy ■C\ -■'L ^ m / y ' "" / ^ v^ i^' •,-' /i^ / ^r I: / *■ 1 f i I 1 Pig. 22. Pia. 23. \Tofacep. 74. PURE METALS AND ALLOYS 75 Fig. 22, Plate VI. The polygonal grains of the pure iron crystals are readily recognised, but here and there between them are seen small dark patches which are entirely absent in the pure metal. These are patches of the carboniferous con- stituent, and in steels of increasing carbon-content the relative area occupied by these dark patches rapidly increases. The reason why the second constituent appears dark in this case, as in the majority of others, is simply due to the fact that, when the polished surface is attacked with an etching reagent, the more soluble, i.e., more readily attacked, constituent is affected almost exclusively, while the other remains practically un- affected until the attack is carried to a considerable length. In this way the areas occupied by the more readily soluble con- stituent are eroded and roughened by the action of the etching reagent, and consequently scatter the hght which falls upon them ; under normal illumination they therefore appear darker than the other constituent, while under obHque hght they shine out brightly on a dark background. The effect produced upon the micro-structure of an alloy of this type by successive additions of the second metal are illustrated by a series of micrographs of the alloys of tin and lead, commencing from the tin end of the series. Figs. 23 to 26, Plates VI. and VII., represent the micro-structures of alloys containing respectively 95, 85, 74 and 45 per cent, of tin by weight. In these alloys the white areas represent the crystals of tin, which are much less readily attacked by reagents than the accompanying crystals of lead. In Figs. 23 to 26 it wiU be seen that increasing lead content leads to an increase of the dark constituent, which at first appears merely in the form of thin veins in the inter-crystalline boundaries of the tin, thicken- ing up into wider patches and bands until the crystals of tin appear as islands in a ground-mass composed of a mixture of small crystals of tin and of the dark constituent. Ultimately, when the lead content has risen to 37 per cent., the crystals of free tin disappear entirely and the whole alloy consists of a mixture of minute patches of dark and Hght constituents (see Fig. 46, Plate IX.). A further increase of lead content causes 76 STUDY OF PHYSICAL METALLURGY the appearance of isolated larger patches of the dark constituent. But while at that end of the series from which we started we were dealing with alloys rightly described as belonging to class (6), at the lead end of the series we find that the white constituent disappears — ^in alloys which have been allowed to approach the condition of final equihbrium — while some 16 per cent, of tin are still present ; at this end the lead-tin system is therefore of the intermediate type possessing a considerable, but still limited, power of forming sohd solutions. The series of photo-micrographs of the lead-tin system (Figs. 23 to 26) show an additional feature deserving of special attention, because it is a very striking characteristic of the micro-structure of all alloys of the type (6) and of most of the alloys of the type intermediate between (a) and (6). In these figures we see a second constituent appearing, but when we examine this constituent in the alloys in which it is present to a larger extent, we find that it is not a simple body, but itself possesses a duplex structure, although in the present case, and, indeed, in the majority of cases, this structure is very minute. The cause of this duplex structure of the second constituent lies in the fact that during the crystallisation of the alloys of this type the separation of solvent and dissolved metal from one another only occurs finally at the moment of complete solidification ; that portion of the alloy which crystallises last contains both metals, and these, on freezing, separate from one another in the form of minute crystal plates or granules, thus forming together the fine-grained duplex constituent shown in the photographs. It will be seen that the alloy containing 37 per cent, of lead and 63 per cent, of tin consists entirely of this duplex constituent. For reasons which will appear when the thermal phenomena connected with the freezing of metals and alloys are discussed, this duplex constituent, and particu- larly the alloy which is composed entirely of this constituent, is known as the " eutectic " alloy of the two metals in question. In many respects the micro-structures of all binary eutectic alloys are very similar to one another. Having now passed in rapid review the most important types of micro-structure met with in metals and alloys in their normal PLATE \U. Fig. lM. Fig. L'."(. Fig 26. [Tufuce p. H], PURE METALS AND ALLOYS 77 state (i.e., when not disturbed by mechanical deformation), we shall pass in the next chapter to a consideration of the modes of solidification of metals and alloys as revealed by a study of the thermal phenomena which accompany the process of crystallisation. RSFEBEMCES. (1) stead. Joum. Iron and Steel Inst., 1898, 1., 145. (2) Ewing and Rosenhain. Phil. Trans. Roy. Soc, 1899, Vol. 193a, pp. 353—375. (3) Osmond and Cartaud. Annales des Mines, August, 1900. Rev. de Metallurgie, 1906, III., 653. Journ. Iron and Steel Inst., 1906, III. (4) Rosenhain and Tucker. Phil. Trans. Roy. Soc, 1908, Vol. 209a, pp. 89—122. CHAPTER V THE THERMAL STtTDY OF METALS AND ALLOYS The study of the phenomena which occur when metallic bodies are heated and cooled has thrown a flood of light on the nature and constitution of metals and alloys, so that the phenomena in question, which relate to fusion and solidification and to the many internal changes which occur in metals or alloys while they are wholly or partly sohd, constitute an important branch of Physical Metallurgy. The importance of this subject lies primarily in the fact that our most compre- hensive and satisfactory method of describing the nature and constitution of alloys of a given system consist of a diagram — known as the " Equilibrium Diagram," or, perhaps, better as the " Constitutional Diagram " — which is based primarily on thermal data. We shall see in a later chapter how it is possible to deduce from the constitutional diagram or to correlate with it a number of important properties of alloys, so that an under- standing of the manner in which these diagrams are established, and the validity of the evidence upon which they are based, is essential. We will begin by considering the manner in which suitable specimens of metal can be heated and cooled for the purposes of these observations ; we shall then briefly consider the instruments used for measuring the temperatures involved, and, finally, the methods of representing the data in the form of " heating and cooling curves." The means of heating a comparatively small specimen of metal steadily to any desired temperature and allowing it to cool down again appear at first sight to be very simple — any laboratory furnace capable of reaching the desired temperature would seem to be adequate. This is only correct where observations of a very rough character are to be taken. The difficulty lies in securing a sufiicient degree of steadiness in both the heating and cooling process, while uniformity of rate THERMAL STUDY OF METALS AND ALLOYS 79 of heating and cooling is still more difficult to attain. With most laboratory furnaces it is necessary to begin by heating with considerably reduced power — ^the gas or the electric current being increased step by step as the temperature rises. But every step of this kind produces a sudden change in the rate of heating and brings with it a disturbance in the heating curve which is being determined. During cooling, on the other hand, there need be no such steps if the furnace is simply allowed to cool down " naturally " ; here, however, the difficulty arises that at a high temperature the furnace cools with excessive rapidity, while at the lower end of the tempera- ture range there is a wearisome delay, due to the extremely slow rate of cooling which obtains there. Both these difficulties can be overcome to some extent by the use of suitable devices, such as controUing the temperature of the furnace by an electric current passing through a suitable heating coU or winding, regulated by some form of automatic rheostat whose resistance is steadily changed in order to keep the heating or cooling of the furnace as steady and uniform as possible. These devices are, however, always elaborate and expensive and never completely satisfactory. The author has, therefore, recently adopted a different plan, which gives excellent results with very Uttle elaboration. In all the older methods of heating and cooling specimens of metal for purposes of thermal study, the specimen has been placed in a furnace and allowed to heat up and to cool down with the furnace. In the author's new arrangement the tem- perature of the furnace is kept steady throughout the entire experiment. The furnace consists of a long vertical tube, which is kept at the highest temperature required at one end, while the other end is cold or nearly so. The specimen is then lowered into this tube or raised out of it in such a way as to pass in a perfectly steady and uniform manner from the cold to the hot part of the furnace, or vice versd, at any desired rate within the range of the raising or lowering mechanism em- ployed. With a Mttle care the tube furnace can be arranged to have a perfectly uniform temperature distribution from one end to the other. By means of this arrangement it becomes 80 STUDY OF PHYSICAL METALLURGY possible to obtain heating and cooling curves at any desired rate, over any range of temperature within the range of an electric resistance furnace, with great ease and regularity, as the temperature of the hot end of the furnace can be controlled very accurately and kept quite constant by means of a thermo- couple permanently fixed in the hot part of the tube. A photograph of the apparatus employed for this purpose is shown in Fig. 27, Plate VIII. An additional advantage of this apparatus is the facility which it provides for submitting a specimen of metal to heating and coohng in any desired atmosphere. By placing the specimen in the closed end of a tube of porcelain or sihca of suitable length and diameter, the upper end being closed by a stopper and tap, the atmosphere in which the specimen is placed can be varied at will, from a high vacuum to air, hydro- gen, nitrogen, or other gas. The influence of gases on the thermal phenomena, particularly of steel, has lately received some consideration and cannot be left out of account in present day research work. The instruments employed for the measurement of the temperature of metal specimens next require consideration. For the purpose of taking heating and coohng curves the thermo-electric couple is in almost universal use, as it possesses indisputable advantages over any other device. Most fre- quently the couple employed consists of a fine wire of pure platinum and one of platinum alloyed with 10 per cent, of either rhodium or iridium. The latter has the advantage of being rather more sensitive, and in the author's hands it has proved extremely constant in its indications, even under difficult conditions. The opinion is, however, widely held that the rhodium couple is more reMable in its indications. For temperatures above 500° or 600° C. the platinum type of thermo-couple is the only one sufficiently permanent to be suitable for laboratory use. For lower temperatures, however, " base metal " couples can be employed with considerable advantage — such a couple as that of copper with the alloy known as " Constantan " being very satisfactory as regards permanence and constancy of indications, while it yields more PLATE VIII. Fk;. 27. THERMAL STUDY OF METALS AND ALLOYS 81 than three times the electro-motive force per degree of tempera- ture which is obtainable from a platinum-iridium couple. The action of the thermo-couple consists in setting up an electro-motive force which is proportional to the difference between the temperatures of the hot and cold junctions. In the case of cooling or heating-curve work in the laboratory, the hot junction is placed in the specimen of metal under examina- tion, care being taken to insure that the junction is properly- placed as nearly as possible at the centre of the mass of metal. The " cold junction " is always kept at a temperature of 0° C. by being placed in a glass tube, which is itself plunged into a mass of melting ice. The ice-box usually employed in the author's laboratory consists of an ordinary (cheap) glass vacuum vessel in which the ice only melts very slowly. Such an " ice-box " in a horizontal position is seen in the photo- graph (Fig. 27, Plate VIII.), where the cold junction is allowed to travel up and down with the moving specimen. The whole of the thermo-couple and the wires leading to it must be carefully insulated electrically so as to remain entirely unaffected by the electric currents employed for heating the furnaces ; the hot junction must further be protected from actual contact with the metal specimens in order to avoid contamination of the couple wires. This protecting sheath — ^generally of fire-clay or of "alundum," or sihca, undoubtedly introduces a sUght " lag " of temperature between the couple and the metal, but, provided that the thermo-couple has been calibrated with the intervention of precisely the same kind of sheath, no error is introduced by this lag. The actual calibration of the thermo-couples employed for this work is usually carried out by the aid of determinations of the freezing-points of certain metals, whose true temperatures are well known. These generally include the following : — Copper (under charcoal) . 1,083° C. Silver (under charcoal) . 960-5° C. Aluminium (99-7 per cent., pure) 658-7° C. Zinc (99-98 per cent., pure) 419-4° C. Lead ..... 327-4° C. Tin 231-9° C. P.M. o 82 STUDY OF PHYSICAL METALLURGY The thermo-couple furnishes an electro-motive force which is proportional to the temperature difference between the hot and the cold junctions. To utilise this indication the thermo- couple must be attached to some form of electrical measuring instrument by means of which the electro-motive force can be read or registered. The simplest form of such an instrument is a galvanometer of high resistance whose deflections are read as measuring the indications of the couple. The resistance of the galvanometer must be high in order to avoid serious inter- ference from changes of temperature in the couple wires themselves and in the connecting wires, for the galvanometer used in this way merely measures the small electric current which the thermo-couple is able to send through the circuit ; this current depends both upon the electro-motive force of the couple and on the resistance of the circuit. In the latter factor the resistances of the couple mres and of the connecting leads play a part which becomes important if the resistance of the galvanometer is low. Beyond this, the simple deflection method has two very serious disadvantages. The first is that, with most sensitive galvanometers, the position of the zero is not constant ; it follows that the actual value — in temperature or in " micro-volts " — of a given scale reading is not constant, and the observer has no means of telling, during the course of a series of observations for a heating or cooling curve, whether the zero has shifted from its original position. The second disadvantage is that the sensitiveness of the simple deflection method is somewhat limited. A scale one metre long is prob- ably as much as can be conveniently employed, and if the curves are to be traced over a range of 1,000° C. then the scale value is 10° C. per millimetre. By one of the devices mentioned below, the effective length of scale for a range of 1,000° C. can easily be extended to as much as eight metres. This increase of sensitiveness — and consequently, under proper conditions, of accuracy — is obtained first by the use of a more sensitive galvanometer, and in the second place, by balancing the greater part of the electro-motive force of the couple by a known electro-motive force taken from an instru- ment known as a " potentiometer," In this arrangement if. THERMAL STUDY OF METALS AND ALLOYS 83 for example, the thermo-couple indicates an E.M.F. of 12,456 micro-volts (= -012456 volts) then the stops of the potentio- meter are set to give a counter E.M.F. of 12,400 micro-volts and only fifty-six micro-volts are read off by deflection. The process can be carried further so that the entire E.M.F. of the couple is balanced at each temperature and the galvanometer always kept at zero, but where the temperatures are changing somewhat rapidly, as in the taking of thermal curves, this is not convenient. The details of the electrical measuring instruments (potentiometers, etc.) employed for this purpose cannot be given here, but the reader will find full details in the literature of the subject (^) and in the catalogues of the instru- ment makers. Only the outhnes of the methods have been given here in order to enable the reader to appreciate the principles adopted. With one of the better class of thermo-electric apparatus just referred to, the observer can readily determine accurately the temperature of the specimen of metal which is under thermal observation at any time during the heating and cooling process. The observations thus made can be recorded in various ways. The simplest of these consists in taking tempera- ture readings at fixed intervals of time and then plotting the results with temperatures as ordinates and times as abscissae. A " time-temperature " curve is obtained which indicates the behaviour of the metal in the most direct way. So long as the metal is simply being raised or lowered in temperature at a steady rate, this curve follows a smooth course ; a departure from this smoothness indicates that there has been either an evolution or an absorption of heat within the specimen. The only objection to this type of curve is that the irregularities produced in it by comparatively small evolutions or absorp- tions of heat are themselves extremely small unless the curve is plotted to an impracticably large scale. Consequently a method of plotting the observations is adopted which is more economical of space and, therefore, allows of the use ol a very much more open scale, with the result that even minute thermal phenomena appear quite clearly. This result is obtained by the use of what is known as the " inverse rate " o 2 84 STUDY OF PHYSICAL METALLURGY curve first adopted by Osmond («) in 1887. For this purpose the observer notes the intervals of time which are occupied by the metal in rising or falling through successive equal differences of temperature. Thus one might take the times occupied by successive rises of 3° C. and plot these as abscissae against the actual temperature of the metal at each observation. In this way is obtained a curve whose ordinates are T (temperature) and whose abscissae are ™i where t is time. With a uniform ai. rate of heating or cooHng, this curve becomes a vertical straight line ; an evolution of heat during cooling or an absorption of heat during heating causes the curve to deflect outwards and to form a hump or " peak." It can be shown that, if the rate of heating is uniform, the area of this peak is proportional to the quantity of heat evolved or absorbed. Some examples of curves of this kind are shown in connection with special metals later in this book (Figs. 50, 66). Mention must also be made of another type of heating and cooling curve, known as the " differential " and the " derived differential." These are obtained by a method devised by Roberts-Austen (^), which consists in comparing the rate of heating or of cooling of the metal under experiment with a standard piece of metal placed in the same furnace and heated and cooled along with the experimental specimen. This method has the advantage that no clock or chronograph is required and that the results are capable of a high degree of accuracy, even when no extremely sensitive galvanometer is available. This is attained by the use of a " differential " thermo-couple which possesses two thermo-junctions arranged to oppose one another and placed one in the experimental piece and the other in the standard piece of metal — usually a platinum cyhnder. The readings of this couple simply indicate the difference of temperature between the platinum cyhnder and the specimen, and if these are plotted against the actual temperature of the specimen as obtained from the readings of an independent thermo-couple placed in the speci- men, the " differential " curve of Roberts- Austen is obtained. Since the uniform cooUng or heating of the platinum cylinder, THERMAL STUDY OF METALS AND ALLOYS 85 which serves as a basis for this type of curve, is in reality proportional to the time which elapses between successive readings, this " differential " curve is practically identical with a time-temperature curve, the clock having been replaced by the platinum cyHnder. Consequently, in order to make the differential curves comparable with the inverse-rate curves, we must plot, not the actual readings of the differential thermo- couple at each temperature, but the change in its reading since the previous observation. If these are plotted against the temperature of the specimen, we obtain the curve first employed by the author (*) and termed the "derived differ- ential " curve. This is very similar to the inverse-rate curve and is practically identical with it in physical meaning. We may now consider the forms which the thermal curves of metals and alloys will assume in various circumstances, in order to arrive at an understanding of the use which may be made of data furnished by these curves.^ Beginning with the simplest case, the coohng of a pure metal from the molten, Uquid state may be considered. If thermal observations — which we may assume to be made and plotted on the " inverse rate " principle — are taken, the resulting curve will run down regularly and smoothly until a certain temperature — known as the freezing-point — ^is reached. Down to this point the rate of coohng has remained constant or nearly constant and the time occupied by each successive fall of, say, 1° C. has remained the same or has merely undergone a gradual increase as the metal has cooled. The resulting inverse-rate curve therefore is a smooth, nearly straight and nearly vertical line (as in Fig. 28). As soon, however, as the metal reaches the " freezing- point " the fall of temperature is arrested. The process of sohdification which then commences is accompanied by a very considerable evolution of heat, i.e., the energy which had been stored in the liquid metal in the form of movement and mutual separation of the molecules, is liberated in the form of heat when the molecules assume the relatively close juxtaposition and comparative immobility of the solid state. Since the metal ' Examplea of actual thermal curves are given in Fig. 60, p. 140, and Pig. 66, p. 168. 86 STUDY OF PHYSICAL METALLURGY is losing heat byradiation and conduction to surrounding bodies, this evolution of heat does not lead to a rise of temperature, in fact, if a rise of temperature were to occur the process would be arrested, since the metal would then be raised above the temperature at which freezing can begin. Consequently the freezing process goes on at just such a rate that the heat set free by the process exactly balances the heat lost by the metal to external objects, and the temperature of the metal remains constant during the whole process. This constancy of tempera- ture during the entire process of solidification is typical of only three classes of bodies, viz., pure metals, pure compounds of two or more metals, and the special class of alloys known as "eutectic alloys " which we have already mentioned ; these three classes of bodies aU share this characteristic — that they consist entirely of a substance or substances which freeze at the same temperature. The freezing, then, of a pure metal takes place entirely at the temperature known as the " freezing-point," and in spite of the continuous external loss of heat the temperature of the metal remains constant for the whole duration of the process. During the freezing in this case, therefore, the rate of cooling, i.e., the rate of fall of temperature, becomes either zero or at all events very slow, and the inverse rate therefore very large ; in other words, when there is a prolonged arrest in the cooling process the time occupied by a fall of temperature of 1° C. may attain a value of several minutes, and the inverse- rate curve takes a sudden jump outwards from the axis (c. Fig. 28). As soon, however, as the solidification of the whole mass of metal is complete, the evolution of heat ceases and the temperature again begins to faU. During the time occupied by the freezing process, however, the heated vessel in which the cooling metal is contained has continued to cool down, while the temperature of the metal itself has been kept constant, with the result that at the end of the freezing process the metal is left considerably hotter than its immediate surroundings. The rate of coohng is, therefore, considerably accelerated and the inverse-rate curve shows a corresponding deflection inwards {de, Fig. 28). These small differences of temperature, however, soon become equalised and the cooling process THERMAL STUDY OF METALS AND ALLOYS 87 continues smoothly and regularly until the ordinary tempera- ture is reached. The inverse-rate curve thus produced is shown in Fig. 28 which is reproduced from an actual coohng curve of pure zinc. It will be seen that the peak of the curve is very sharp and narrow, indicating that the retardation of coohng, i.e., the evolution of heat within the metal, w£is con- fined to a very short range of temperature. We now turn to consider the more comphcated phenomena a «0 b . T s " 400 ' \ ^350 3 2 \ \ 250 ■ \ \ Fig. 28. — Cooling Curve of Pure Metal (Zinc) showing sharp freezing point. met with when the cooling of a binary aUoy is observed. Up to the point where sohdification commences the process is the same as that already described in the case of a pure metal, except for the important fact that the temperature at which solidification commences in an alloy (apart from some special exceptions) is always lower than that at which the principal constituent metal of the alloy would begin to sohdify. This is simply a case of the universal phenomenon of the lowering of the freezing-point of a hquid when a foreign element is dissolved 88 STUDY OF PHYSICAL METALLURGY in it, for — as we have already seen — ^the liquid alloy is simply a solution of two liquids in one another. As soon as the process of soHdification commences, however, important divergences between the behaviour of a pure metal and an alloy make themselves felt in the time observations and in the resulting inverse-rate curve. What actually occurs depends on the type of alloy in question. We wiU first take the case of the type referred to as (6) in the previous chapter, i.e., that in which the metals are mutually insoluble in the solid state and, therefore, separate out completely from one another in the process of freezing. Let the constituent metals of the aUoy be called A and B and let us suppose that the alloy under observa- tion contains a large proportion of A. In that case the freezing process will begin by the crystaUisation, at one particular temperature, of a small quantity of A. The soHdification of this small quantity of metal sets free a certain quantity of heat and the curve shows a sudden outward bend, resembhng the outward sweep at the beginning of freezing in a pure metal ; but in the present case the quantity of metal sohdified forms only a portion of the entire mass, and the correspondingly smaller quantity of heat liberated causes a smaller deflection of the curve. The freezing of a relatively small quantity of pure A will, however, leave behind a hquid which contains relatively more of the second metal, B, than the original molten alloy, and this richer hquid wiU only begin to deposit further quantities of A at a lower temperature. As the temperature gradually falls, therefore, successive quantities of A will crystaUise out, leaving the hquid successively richer in B. As solidification is gradually taking place during this entire process, heat is being continuously Mberated and the natural rate of coohng of the mass is thereby continuously retarded — the inverse-rate curve therefore remains well out to the right (as at bb, Fig. 29) only gradually returning towards its normal position. This process of selective freezing, however, does not go on indefinitely, as might perhaps be supposed, i.e., until the whole of the metal A present in the aUoy has crystallised, leaving the liquid as pure molten B which might then crys- tallise separately. On the contrary, the concentration of B THERMAL STUDY OF METALS AND ALLOYS 89 in the liquid cannot pass a certain definite limit, and this limit is arrived at when the temperature of the cooling alloy has also reached a Mmiting value ; when these conditions are attained, the whole of the remaining liquid sohdifies at one temperature, giving rise to a considerable evolution of heat and producing upon the inverse-rate curve a sharp peak resembHng that of the initial freezing of a pure metal. If we compare this description of the sohdification process Fig. 29. — Cooling Curve of Eutectiferous Alloy showing Initial Freezing and Euteotic Arrest. with the account of the micro-structure of alloys of this type as given in the previous chapter, the correlation of the two will at once be apparent. The separation of the constituent A during the earlier process of selective freezing represents the formation of the crystals of pure metal seen in the micro- structure, while the solidification of the remaining Hquid at the single lower temperature represents the sohdification of the second, duplex constituent which we have already named the " eutectic " alloy. This word " eutectic," derived from the 90 STUDY OF PHYSICAL METALLURGY Greek, simply means " most fusible," and indicates that this alloy, or this constituent of an alloy, is the most fusible mixture, i.e., the mixture of lowest freezing-point, which can be prepared out of the binary system in question. We have already indicated that this eutectic constituent has for each system of alloys a definite composition and also a definite freezing-point. It is obvious that an alloy consisting wholly of this eutectic constituent can be prepared, and this, hke a pure metal, wiU solidify entirely at one temperature, viz., at that temperature at which solidification of all other alloys of the same series is completed. From the present point of view we may trace the effect of successive additions of a metal B to a metal A much in the same manner as we did in the case of the micro-structure, and this will serve to bring out the correlation more clearly. The addition of B to A brings about a lowering of the tem- perature at which soKdification begins, and also produces an extended range of temperature over which the freezing process is spread out ; when the proportion of B is still small, the freezing process begins by the ehmination of pure solid A, and this continues until nearly the whole alloy is sohd. The portion which is left liquid, however, attains the eutectic composition and finally freezes as a whole when the temperature of the freezing of the eutectic is reached. With successive further additions of B the temperature of the commencement of solidification is successively lowered, but the temperature of final soUdification remains unaltered, while the hquid which freezes last always attains the composition of the eutectic. Finally, by further additions of B the eutectic composition is reached ; here the temperature of the initial freezing has been lowered to the temperature of eutectic freezing and the whole mass soUdifies at one temperature. Further additions of B gradually reverse the series of changes just described, with this difference that the substance which now separates in the pure state, when freezing commences, is pure B where it was formerly pure A. The composition of the residual liquid and the temperature of final solidification remain unaltered until the whole mass of metal consists of pure B, which, of course, again sohdifies in the manner already described for pure metals. THERMAL STUDY OF METALS AND ALLOYS 91 This entire process or group of processes, although somewhat complicated when thus verbally described, is very readily represented by the aid of a diagram. If we plot as abscissae the percentage compositions of a series of alloys, and then, using temperature as ordinate, insert the points corresponding to the beginning of freezing in each alloy and also the points corresponding to the freezing of the residual metal or euteotic, we obtain a diagram which represents the entire set of pro- cesses in a simple manner. Such a diagram is shown for the imaginary case of two metals perfectly insoluble in one another ^ "" ^-"""^^^ ^ y Fig. 30. — Ideal Equilibrium (Constitutional) Diagram of a simple Binary Eutectiferous System. in the soUd state, in Fig. 30. Such a diagram is in reality a summary of the indications of the cooling-curves of all the alloys of the system, and from the diagram the nature of the cooling-curve of any of these alloys can be readily foretold. This relationship is indicated in the diagram (Fig. 31) where the cooling-curves of some of the alloys are shown. In looking at this diagram it must be remembered that the " peaks " of the cooHng-curves really lie in a plane at right-angles to that of the diagram, since the co-ordinates of the cooling-curve are temperature and time, while those of the diagram proper are temperature and composition of the alloy. The diagram 92 STUDY OF PHYSICAL METALLURGY shown in Fig. 30, while thus representing a summary of a whole series of cooling-curves, has a further meaning, since it also serves as a species of map or chart of the various conditions in which alloys of this particular system can exist at difEerent temperatures. Thus any point in the diagram corresponds to an alloy of definite composition existing at a definite tempera- ture, the composition being that corresponding to the abscissa of the point and the temperature that represented by the ordinate. It wiU also be seen that the Unes joining the various freezing-points, viz., the fines AEB, CEF and EL, divide the Fig. 31. — Diagram illustrating the relation between arrest-points on Thermal Curves and the lines of the Constitutional Diagram. whole area of the diagram into a number of separate fields or areas. If we look at the matter more closely we shall see that each of these areas represents a definite state of the alloys. Thus, if we take any point such as P, lying above the line AEB, the alloy corresponding to that point will be completely liquid ; any point lying above the Hne AEB represents an aUoy at a temperature above that at which solidification can commence. Next, if we take a point Q lying within the fines AE, EC, CA, such a point must represent an aUoy partly sofid, partly Mquid, and from what has been said above in describing the process of freezing of such alloys we see that the part which is already THERMAL STUDY OF METALS AND ALLOYS 93 solid Mdll consist of the pure metal A. The whole area or field AECA thus represents alloys consisting of mixtures of solid A with a residue of liquid, but it should be noted that the pro- portion of solid to liquid will vary very widely for different portions of this area. If now we take a point below the line EC, we find that the alloy must be completely soUd, since the line EC passes through the points of final solidification, or freezing of the eutectic. The alloys in the region CEZX con- sist of a mixture of solid A with sohd eutectic. Near the hne CX the A constituent predominates, and the eutectic occupies only a small proportion of the whole volume, but, as the line EZ is approached, the eutectic is present in increasing proportions until, when the line EZ itself is reached, the entire alloy consists of the eutectic. This, of course, corresponds with the fact indicated by the meeting of the various lines in the point E, which indicates that for an alloy of this com- position, i.e., for the eutectic alloy, the temperature of the commencement of freezing is also the temperature of final complete solidification, the eutectic alloy having thus a simple freezing-point hke that of a pure metal. The condition of the alloys corresponding to the remaining areas of the diagram will readily be understood on the basis of those already described ; the region BEDB corresponds to mixtures of solid B with liquid, and the region ZEDY to mixtures of solid B with sohd eutectic. A diagram of this kind is generally known as an " equiU- brium diagram " of the system of alloys in question ; this name is apphed to it because the diagram indicates the conditions which the alloys would assume for any given composition or temperature, provided that they were allowed to attain equili- brium. This latter proviso is very important, because in a great many cases alloys can be brought to temperatures and compositions indicated on such diagrams without assuming the conditions implied by the diagram. These other conditions are always of a more or less temporary and artificial nature — exposure to suitable temperatures which allow of more ready adjustments of molecular conditions always results in allowing the alloy gradually to revert to the condition of stable equili- 94 STUDY OF PHYSICAL METALLURGY brium. On the other hand, these other " meta-stable " condi- tions are of very considerable importance, since most of the metals in ordinary use are to some extent held in such a condi- tion by the difficulty which their molecules experience in rearranging themselves at ordinary temperatures. A striking example of this class of phenomenon is found in hardened steel, and, indeed, all industrial steel comes under this description to some extent. The eqmlibrium diagram, however, shows the "natural" state of the alloys, i.e., that state to which they revert whenever the opportunity is given them to do so. In many cases, however, where true equilibrium conditions are not thoroughly known or cannot be even approximately realised in the laboratory, the diagram represents partially meta-stable conditions and is better termed a " constitutional " rather than an "equilibrium" diagram. Certain of the lines seen in the diagram of Fig. 31 have received names which it will be convenient to define here. Thus the lines AE, EB form the lower boundary of that region in which the alloys are completely liquid ; they are sometimes called the " melting-point curve," but since most of the alloys do not possess a single " melting-point," it is preferable to adopt the Latin term " liquidus " for this curve. Correspond- ing to the line which bounds those regions of the diagram which represent entirely fluid alloys, we have a line which bounds those regions within which the alloys are completely soUd, and this line is known as the " solidus " curve. In the diagram of Fig. 30 the solidus curve is represented by the simple straight line CED, but it must be remembered that this diagram corresponds to an ideally simple case, the diagrams of actual alloys being always much more complex. The above definitions, however, hold in all cases. We have now to consider some of the compUcations which are introduced into both coohng-curves and constitutional diagrams when the simple, ideal condition laid down above — that the two metals shall be mutually completely insoluble in the solid state — ^is departed from in varjdng degrees. When the state of mutual solution of two metals remains unchanged by the process of solidification, we have seen that THERMAL STUDY OF METALS AND ALLOYS 95 the result, so far as the microscope is concerned, is the pro- duction of a homogeneous mass of crystals quite similar to the aggregate of crystals which constitutes a pure metal. Corre- sponding to this micro-structure we should expect the cooMng- curve of such a " soHd solution " to resemble that of a pure metal. To a certain extent this is true, but there is an im- portant difference, and one which is very apt, unless extremely gradual cooling is resorted to, also to introduce a disturbing factor into the resulting micro-structure. Experiment soon shows that such soHd solutions do not possess a simple freezing- point like that of pure metals, but that they show a " freezing range," i.e., a range of temperature of definite extent over which the freezing process extends. If the coohng process were artificially arrested at any temperature lying within this range, the alloy would be found to consist of a definite propor- tion of soKd and liquid — a proportion which would be quite fixed for any given temperature. The freezing process of a sohd solution is therefore obviously different from that of a pure substance, and we must now examine this difference. In the first place, the temperature at which freezing com- mences will, as a rule, be lower than that at which the pure metal would freeze, and we find, further, that while in the case of the pure metal during the freezing process the portion which has already solidified is exactly the same in composition as that which is still liquid, this is not and, indeed, cannot be the case where the formation of a solid solution is concerned. For if the part already frozen and the part still Hquid were exactly alike in composition, their freezing temperature would also be the same, and the rest of the hquid would soUdify without any fall of temperature ; this is exactly what does occur in a pure metal, but does not occur in a solid solution. Let us consider an alloy containing a small proportion of B dissolved in A. When freezing commences, the sohd which separates first contains less B than the liquid alloy as a whole. The remaining hquid is thus left richer in B than it originally was, and it therefore possesses a lower freezing-point. As soon as the temperature falls, a further quantity of soUd is formed, con- taining more B than the first portion of solid, but less B than 96 STUDY OF PHYSICAL METALLURGY the liquid from which it is deposited. At this stage, therefore, the alloy consists of cores of sohd containing least B, shells of soHd surrounding them, containing rather more B, and a residual liquid still richer in B. This, however, is not the stable con- dition, for, if time be allowed, the composition of the whole solid wiU become the same, viz., like that of the fresh solid formed at the temperature in question; but this can only occur by the slow process of diffusion, some of the excess of B travelling from the Uquid through the already solid shell richer in B to the core which contains the lowest proportion of B. The process of sohdification, extended over a range of tempera- tures as thus indicated, continues until the whole of the alloy has become solid, provided that the composition of the alloy is such that the limit of solubility of sohd B in solid A is not passed at any time. If time be allowed, as already indicated, the whole of the solid thus formed will attain one and the same composition, and an aggregate of homogeneous crystals will result. When, however, the rate of coohng has not been slow enough to allow the process of diffusion to equahse the com- position, each of the crystals of the sohd solution will consist at the end of the freezing process of a core of metal relatively poor in B, surrounded by successive shells or layers increasing in their content of B from the centre outward. When alloys of this kind, which have been cooled somewhat rapidly from fusion, are examined under the microscope, very distinct traces of the cores and shells just described are often found ; in some cases the changes of composition between adjacent layers are very strongly marked, and the aUoy almost looks as if it possessed a true duplex structure, while in other cases the changes of composition are very gradual, and the cores then present a shadowy, ghost-hke appearance which is very readily recognised whenever met with. An example of such a " core " structure has been shown above in Fig. 21, Plate VI. The shape of the cooling-curve of an alloy of this kind is shown in Fig. 32. The commencement of freezing is well marked, but there is no definite arrest — the peak in the inverse rate curve is spread over a range, and, although the maximum departure of the curve is always found near the THERMAL STUDY OF METALS AND ALLOYS 97 beginning of the procjess, the curve only returns to its normal position at a temperature at which the whole of the alloy has soUdified. The precise position of the end of the freezing process, however, is not usually sharply defined on such curves. As in the case of the class of alloys formed by metals which are mutually insoluble in the solid state, the cooling-curves of a system of alloys of metals which are mutually soluble in the solid state can also be grouped together in the form of a constitu- tional diagram, although the shape of this diagram is not so well defined. A typical example is given in Fig. 33, which represents the alloys of copper and nickel. The Hne ACB is the " liquidus" forming the lower boundary of the region in which the aUoys are com- pletely Hquid ; the " solidus " is indicated by the hne AsB, and the relations of the coohng-curves to the diagram are indicated by the dotted lines representing the inverse- rate cooling curves of a few of the alloys of the series. The shape of the liquidiis curve in such a system of alloys, however, is not always of the kind indicated in Fig. 33. In certain cases the curve either droops or rises in the centre, although it is not quite certain whether in such cases special disturbing causes may be at work, such as the formation of a chemical compound between the two metals, the compound, however, being completely soluble in either of the constituent metals. With such questions we need not, however, concern ourselves here — ^the typical form of the liquidus curve of a series of alloys forming solid solutions is that of a smooth, continuous curve showing no sharp break or angle. With this continuous form of curve the properties P.M. H Fig. 32. — Cooling Curve showing the Freezing- Range of an Alloy form- ing a Solid Solution. 98 STUDY OP PHYSICAL METALLURGY of such alloys correspond very closely, also forming a continuous series without sharp breaks or changes of properties. While the class of alloys consisting of metals entirely insoluble in one another in the solid state may safely be regarded as an ideal case never met with in practice, the class of alloys whose constituent metals are mutually soluble in all proportions in the soHd state is met with in quite a number of cases of actual binary systems. In the greater number of cases, however, Fig. 33. — Constitutional Diagram of the Alloys of Copper and Nickel, typical of an uninterrupted series of Solid Solutions. metals behave towards one another in an intermediate manner, the metals being soluble in one another in the soHd state to a definitely limited extent. The behaviour of their alloys during freezing and the shape of their constitutional diagrams is, therefore, also of an intermediate nature. Alloys whose composition is such that their solidification implies no dis- turbance of the state of mutual solution of the two metals behave as if they belonged to class (a) described above, solidify- ing in the manner typical of sohd solutions and exhibiting the corresponding cooling curve and micro-structure. Thus, if THERMAL STUDY OF METALS AND ALLOYS 99 metal A is capable of retaining say 8 per cent, of metal B in solid solution, then alloys of A and B containing less than 8 per cent, of B will, at aU events if slowly cooled, solidify as simple solid solutions. On the other hand, those alloys of the series whose composition places them beyond the limit of soHd solubility of the metals at either end of the series behave in most respects in a manner very similar to alloys of class (6) described above. The only difference is that the solid which Fig. 34. — Constitutional Diagram typical of Alloy Systems partly eutectifeious but forming solid solutions at each end of the series. first crystallises is not pure metal, but the saturated soHd solution of one metal in the other, according to the end of the series to which the particular alloy belongs. The cooling- curves of alloys belonging to this portion of the series are in every way like those of alloys of class (a), and the constitu- tional diagram is correspondingly similar. The diagram of the whole of such a system is, therefore, made up of three distinct portions — a central portion BO resembling the diagrams of class (a) and portions AB and CD at each end, typical of alloys of class (6). A diagram of this type is shown in Fig. 34. The h2 100 STUDY OF PHYSICAL METALLURGY alloys of silver and copper belong to this class. The liquidus is, in a general way, very similar to that of a system entirely belonging to class (a), but the solidua differs widely — ^the line of eutectic solidification does not extend to the two sides of the diagram, the solidus being completed by two curved branches leading from the end of the eutectic line up to the freezing- points of the two metals. The diagram is again divided up into fields which represent groups of alloys in different conditions, but we need not enumerate them in detail, since the reader will be able, by comparison with the two simpler cases, to ^discover the interpretation of each field, if it is not at once obvious to him from what has already been said. A further class of alloys, possessing a somewhat different constitutional diagram from those already described, are those systems in which definite inter-metallic compounds are found. These, however, need but little separate discussion here, since the whole matter can be regarded in a very simple manner. From the point of view of cooling-curves and constitutional diagrams, as well as from that of micro-structure, any definite inter-metaUic compound may be regarded as being simply another pure metal — ^it is in fact a pure substance of metallic character, so that this method of regarding it is by no means strained. If, however, this idea be kept steadily in mind, the facts concerning inter-metallic compounds are much simpUfied. A system of alloys of two metals in which such compounds occur is simply broken up into a series of systems. Thus, if the metals A and B form the compound A„B„, the series of alloys between A and B may be looked upon as two successive series, one consisting of the alloys of A with the compound metal A,„B„, and the other of the compound and the metal B. Each of these systems has its own characteristic constitutional diagram usually with series of soUd solutions, a eutectic alloy, etc., and to obtain the complete diagram of the A-B system it is merely necessary to juxtapose these two diagrams. In doing this, we see at once that, since the compound A„B„ forms the right- hand end of the first diagram and the left-hand end of the second (see Fig. 35, relating to the alloys of magnesium and tin) THERMAL STUDY OF METALS AND ALLOYS 101 the liquidus and solidus^ of both diagrams will meet at the point which represents the freezing-point of the pure com- pound A„B„, and — ^what is more — ^both solidus and liquidus fall away from the point corresponding to the compound. It is thus clear that in any complete constitutional diagram the presence of a pure compound will be marked by (1) a maximum point in the liquidus, and (2) a meeting-point of solidus and liquidus. Correspondingly, in the micro-structure we find the compound represented by the typical structure of a pure metal, soo 4 BOO t^ 700 y^ \ '^~~~'^~~~-~^ f \ \soo ■ \ ^400 ■ \ 300 ■ \ 200 1 1 1 \^ to 20 30 40 SO 60 70 80 90 lOO/iSn. Mg Concentration Sn. Fig. 35. — Constitutional Diagram of the Alloys of Magnesium and Tin, typical of a series forming a definite Inter-metallic Compound. viz., an aggregate of crystals aU of one kind, although these, as we have already seen, are not always to be distinguished from the crystals of a slowly-cooled soHd solution. It may here be remarked at once that, from the point of view of the mechanical properties of alloys, inter-metallic compounds and alloys consisting to any notable extent of such compound bodies are entirely useless. It is a very striking fact that 1 The vertical part of the solidus in Fig. 35 should leallybeshown as two lines, very close together and sloping very steeply towards the maximum of the liquidus ; they are shown merged into a single vertical line. 102 STUDY OF PHYSICAL METALLURGY these compound bodies are always brittle and, as a rule, very weak mechanically. Their presence in small proportions m many cases serves to strengthen and stiffen metals or alloys which would otherwise be unduly soft, but as soon as a notable proportion of such a body is present the ductihty of the alloys disappears. The same action, of course, also occurs where pure metals, or solid solutions of pure metals, of a brittle nature enter into the structure of an alloy, but the interesting fact at the present point is the almost universal occurrence of brittleness in the case of inter-metaUic compounds. Another very important feature which is frequently met with on constitutional diagrams has yet to be discussed ; this feature is a horizontal line (or lines) representing a series of changes occurring in some (or all) of the alloys at a temperature below that of complete solidification. Such changes occur in many solid alloys, in some cases at moderately high tempera- tures, while in others the changes take place at quite low temperatures. When the cooling-curves of such alloys are followed to temperatures below the end of the freezing process, fresh evolutions of heat, represented by peaks on the coohng- curves, are met with. When these are observed in successive members of a series of alloys they usually lie on a horizontal Hne in the diagram, although in a certain number of cases curved hues are met with. In the great majority of cases, also, it is possible to trace some change of micro-structure which is associated with the evolution of heat. Such changes of structure are usually observed by means of a process known as " quenching." For this purpose small pieces of the alloy are raised to a temperature just above that at which the evolu- tion of heat takes place and are kept at that temperature for a sufficient length of time to attain the condition of equilibrium which corresponds to that temperature. Then the specimen of alloy is cooled as rapidly as possible, usually by immersion in cold water. A special apparatus for effecting this operation without either removing the specimen from the furnace in which it has been heated, or exposing it to oxidising gases, has been devised by the author (*), and by its aid this operation can be effected with great precision. This apparatus is illus- THERMAL STUDY OF METALS AND ALLOYS 103 trated in Fig. 36, Plate VIII. The very rapid cooling undergone by the specimen of naetal in these circumstances is intended to prevent the occurrence of the change which gives rise to the evolu- tion of heat under study, and it is found in practice that, although as a rule such changes cannot be entirely suppressed even by the most rapid cooUng, they are reduced to such an extent that a good idea of the true structure of the metal as it existed at the quenching temperature can be arrived at. A comparison with the micro-structure of a piece of the same alloy when slowly cooled, then serves to reveal the nature of the transformation which the metal has undergone in passing through the change accompanied by the heat-evolution. Photo-micrographs of an alloy of copper and tin containing 13-4 per cent, of tin, one slowly cooled, another quenched from a temperature of 700° C, and a third annealed at 480° C, are shown in Figs. 58, 59, and 60, Plate XII. Perhaps the best known of such changes are those occurring in carbon steels, but these will be dealt with in a later chapter. The data for the establishment of a constitutional diagram which are furnished by heating and cooHng-curves are not, however, exhausted when the temperatures of all the observed arrest-points are plotted on the diagram. The quantities of heat evolved or absorbed at each of these arrest-points are also of importance. Were it possible to measure these quantities with any great degree of accuracy they would furnish data of the highest importance, but in practice the quantitative indications of thermal curves must be regarded as decidedly approximate. Whether the simple duration of an arrest be observed, as has been done by Tammann (®) and his collaborators, or whether the more elaborate and accurate method of measuring the areas of the peaks of carefully determined inverse-rate or derived differential curves be adopted, as has been done by the author ('), no very high degree of accuracy can be attained, principally because the end-point of an arrest is never sharply defined, and also because the conditions of heating and coohng cannot be kept absolutely constant from one alloy to another. Yet even the approximate data are useful, provided that care is used in their interpretation ; this is particularly necessary 104 STUDY OF PHYSICAL METALLURGY where there is any possible doubt whether the heat-evolutions which are being measured take place under conditions of equilibrium or whether the alloys are in a partially meta-stable condition. Where the latter is the case, the quantity of heat evolved or absorbed in a given instance may vary owing to causes outside those under immediate investigation, and false conclusions may be arrived at. With these difficulties in mind, the indications of a quantitative nature obtained from thermal curves may be utilised to determine the concentration at which a eutectic point occurs, since the heat evolved by the freezing of the eutectic will be at a maximum in a specimen containing nothing but eutectic, and will fall off regularly in alloys whose composition lies on either side of the eutectic composition. If the quantities of heat evolved during cooUng are plotted as ordinates against the composition of the alloys as abscissae the resulting curve shows a peak or maximum which indicates the eutectic composition, in the manner shown in Fig. 37. In the same way, the end of a eutectic hne where it either abuts against a vertical line indicating a compound or a pure metal, or where it merges into the curved " solidus " of a series of sohd solu- tions, as at F and H in Fig. 48, may be inferred by con- tinuing such a curve as that of Fig. 37 to the point where it crosses the axis of composition, but this extrapolation is always rather uncertain. Similarly, where a compound is formed by a reaction between solid and liquid portions of an alloy, the composition of the compound may be determined by finding much as in the case illustrated in Fig. 49, the composition corresponding to the maximum heat-evolution ; but here it is particularly important to be certain that in every case the Composition of Alloy Fig. 37. — Diagram of Thermal Analysis Curves. THERMAL STUDY OF METALS AND ALLOYS 105 reaction has had time to complete itself, otherwise the quantity of heat evolved may depend more on the rate of cooling than on the true composition of the compound. The constitutional diagram as discussed in the preceding pages would almost appear to be httle more than a summary of the purely thermal data obtained from heating and coohng curves. Yet while these are extremely useful and important in assisting an investigator to establish such a diagram, they are not in themselves quite sufficient for that purpose, and their indica- tions must be checked and amplified by microscopic research. In many cases this is merely a question of observing whether a new constituent appears where the supposed lines of the diagram would lead one to anticipate its coming, and whether the typical eutectic structure occurs, as the diagram would suggest. In other cases more elaborate investigation is required. The use of annealing at a temperature just above a line of thermal arrest - points, followed by quenching, has already been mentioned. In some cases, however, lines of the diagram cannot be located at all by the pyrometer and there microscopic research is the principal resource. For instance, since thermal curves only represent vertical sections of the diagram, these curves cannot be used to fix the position of vertical hnes in the diagram. Where these fines indicate the limits of solid solubility their exact position can best be settled by means of a series of small specimens of the alloys ranging a short way on either side of the supposed line. These must then be " annealed " for a long period — in some cases amount- ing to many weeks — at exactly known temperatures, and subsequently examined with the microscope. Up to a certain hmit in composition they are found to have become homo- geneous — i.e., after the prolonged annealing they consist solely of a single constituent. Beyond that composition no amount of prolonged heating will render them entirely homo- geneous — i.e., traces of the second constituent persist to the end. The Hmit of solubifity is thus fixed at the boundary between these two different results. In the same way, quench- ing experiments often serve as the only true guide to the position of the " solidus " curve in a range of sofid solutions, 106 STUDY OF PHYSICAL METALLURGY We have already seen that thermal ciirves yield no satisfactory indication of the exact end of the process of solidification. This point can, however, be ascertained with considerable accuracy by quenching small specimens of suitably chosen alloys from a series of temperatures Ijdng slightly above and shghtly below the temperature at which the " solidus " is likely to be met. The specimens quenched from a temperature shghtly above the solidus must contain, at the moment of quenching, a very small amount of Hquid (molten) metal, and, during the rapid coohng which occurs in quenching, this hquid metal soHdifies very quickly and thus produces minute regions of very fine structure. These " fusion spots " are readily noticed when the specimen is afterwards pohshed and etched, so that it becomes fairly easy to say which of a series of speci- mens have been quenched from temperatures above and which from below the " solidus." Ebfeeenoes. (1) Stansfleld. Phil. Mag., 1898, V., 46, 59. KurnakofE. Zeitsohr. Anorg. Chemie, 1904, 42, 184. White. Physical Review, 1907, 25, 334. Carpenter and Keeling. Journ. Iron and Steel Inst., 1904, I., 224. (2) Osmond. Memoirs de I'ArtiUerie Marine, 1887, XV., 673. Journ. Iron and Steel Inst., 1890, I., 38. (3) Roberts- Austen. Fifth Report, Alloys Research Committee, Proc. Inst. Mech. Engineers, 1899. (4) Rosenhain. Proc. Physical Soc, 1908, XXI., 180. Burgess. Bull. Bureau of Standards, 1908, V., 199. (5) Rosenhain. Journ. Iron and Steel Inst., 1908, I., 87. (6) Tammann. Zeitschr. Anorg. Chemie, 1903, XXXVII., 303 ; 1905, XCV., 24 : 1905, XCVII., 289. (7) Rosenhain and Archbutt. Phil. Trans. Roy. Soc., 1911, Vol. 21lA, pp. 315—343. CHAPTER VI THE CONSTrrtTTIONAL DIAGRAM AND THE PHYSICAL PBOPEETIES OF ALLOYS The " Constitutional Diagram " has so far been considered almost entirely as a representation of the behaviour of the aUoys at various temperatures so far as solidification and fusion are concerned, but the diagram in reality relates to all the physical properties of the materials — ^indeed, in the case of the lines of a diagram which relate to a transformation occurring in the soUd alloys there is nothing in the nature of a change of state connected with the passage from one of the fields of the diagram to another. It is, however, a universal rule that passage across the boundary of any of the fields of a properly constructed constitutional diagram involves a change in most of the physical properties of the alloys — ^whether the transition from one field to another be due to change of temperature or change of chemical composition. This coimection between the constitutional diagram and the properties of alloys is, of course, not only useful for the purpose of forecasting the probable properties of an alloy of given composition from its position in the diagram, but it is also available in the reverse way, for by the study of the changes in physical properties of a group of alloys with changing com- position and temperature, it is sometimes possible to fix with accuracy the position of certain lines of the diagram which are not readily ascertained by either microscopic or thermal methods. Almost every physical property, and even some essentially chemical or electro-chemical properties, have been utilised in this way, and although none of them have as yet proved as powerful for this purpose as the use of the pyrometer and the microscope, yet valuable results and confirmations have been obtained in this way. 108 STUDY OF PHYSICAL METALLURGY It is not possible to enter into a detailed account of these varied investigations here, and Httle more than an enumeration will be attempted. The references to the Hterature wiU, how- ever, enable the reader to follow up any of these paths of inquiry as far as it has hitherto been explored. One of the most obvious physical properties which might be correlated with the constitutional diagram for the present purpose is the specific volume (^) or density of metals and alloys, and more particularly the changes which the specific volume undergoes as the alloy is heated and ultimately melted, or during the reverse process. This method has been employed by Charpy and Grenet (*) for the study of the critical points in steel, but it does not lend itself well to the exact study of changes involving fusion or solidification. The author has attempted an investigation of this kind for the more fusible alloys, such as those of tin and lead, and has developed a method of differen- tial weighing against a " neutral " body for that purpose (*) ; unfortunately, however, the difficulty of finding a hquid in which the specimens of alloys could be weighed at temperatures up to and above their melting points has not yet been satis- factorily solved. A certain amount of work has, however, been accomplished on very fusible alloys — such as amalgams and alloys of the alkah metals — ^in Germany (*). The whole subject has been very fully reviewed recently by Guertler, who has worked out the theory of the whole matter very clearly. Mention must also be made of the efforts of Turner (^) and Ms collaborators (Murray, Ewen, Haughton, Chamberlain, and others) who endeavoured to trace the volume changes in metals and alloys by pouring the molten material into a sand mould of a T shape and following the movements of the "free " end of the metal by means of a dehcate indicator or " extenso- meter." A very large amount of work has been done in this way and some interesting results have been obtained, but, unfortunately, the method is open to serious and fundamental objections, and it is now admitted that the results, although in certain cases they appear to be closely related to the form of the constitutional diagram, cannot attain a degree of precision sufficient to be of use in the determination of these diagrams. PHYSICAL PROPERTIES OF ALLOYS 109 The relationship between magnetic properties and the constitution of alloys is not as yet sufficiently well understood to allow of magnetic measurements being satisfactorily em- ployed for the study of alloy constitutions, although, of course, the magnetic properties of iron and its aUoys are often con- sidered in discussions of their constitutional diagrams, particu- larly in relation to the much-discussed question of the " beta " phase in pure iron and mild steels. Although these magnetic properties have been very widely studied {*), it does not appear to the author that they can be very safely used as critical data in such questions ; the magnetic behaviour of any given piece of iron or steel appears to be influenced in a very complex way by its entire past history — thermal, mechanical and even magnetic — on the one hand, while, on the other hand the magnetic pro- perties would appear to be directly related to the nature of the atoms of the element rather than to the arrangement — crystalline or molecular — in which those atoms may be present at any given time. Apart from the alloys of iron there is only a limited range of metals and alloys in which magnetic pro- perties are sufficiently strongly marked to invite careful study. Among these the alloys of copper with aluminium and man- ganese, and some allied materials, generally known as " Heussler alloys " (^), are of special interest since they constitute materials showing comparatively strong magnetic properties while they appear to contain no strongly magnetic element. The fuller study of these questions, however, has shown that several compounds of manganese behave as if that element were magnetic when present in a certain state of combination, and similar properties probably exist in other alHed elements. Another interesting feature which connects the magnetic properties of alloys with their constitution is the effect produced on the magnetic hysteresis of pure iron by the addition of various alloying elements in small quantities. It was thought at one time that all impurities or intentional additions to pure iron tended to increase its magnetic hysteresis, and for such purposes as transformer sheets the aim was to produce the purest possible commercial iron. In this respect the effect of impurities or additions would have been strictly analogous to 110 STUDY OF PHYSICAL METALLURGY their effect on the electrical conductivity of metals. This analogy, however, does not hold, since it has been found, that the addition of certain elements, notably aluminium and silicon, so far from increasing the hysteresis, markedly reduces it {^). Study of the matter has shown that the effect of an addition or impurity in this respect is a function of the relative atomic volume ; elements whose atomic volume is less than that of iron tend to increase the hysteresis, while those whose atomic volume is greater have the opposite effect. This observation has proved of considerable practical value in the production of sheet iron of remarkably low hysteresis for magnets and transformers, while it is also of considerable theoretical interest in connection with the theories of magnetism in metals, such as that of Ewing (^), and the " magneton " theory of Weiss ("). The electrical properties are of far more immediate im- portance in connection with the constitutional study of metals and alloys. This method of investigation has now attained such a degree of importance that no investigation of the constitution of a system of alloys can be regarded as really complete until a study of electrical conductivities and their temperature coefficients has been carried out. A very full statement of the relation between these electrical properties and the constitution of alloys has been given by Guertler (") in a paper read to the Institute of Metals in 1910 and only a brief indication of the wider aspects of the subject can be given here. In regard to electrical conductivity, observation has shown most definitely that the highest conductivity is always found in a pure metal, and that the presence of any alloying element reduces the conductivity very materially, particularly if the second element forms a soUd solution with the pure metal. This occurs even in cases where the added metal itself is a far better conductor than the metal to which it has been added. A satisfactory physical explanation of these phenomena has not yet been put forward, although an attempt is made to account for the phenomena of electrical conductivity in metals by the electron theory which supposes that the electric current PHYSICAL PROPERTIES OF ALLOYS 111 in a metal is constituted by a stream of moving electrons ; the electrical resistance of the metal then becomes a simple fric- tional resistance to the movement of these electrons, but it is not easy to see why the simultaneous presence of two or more kinds of atoms or molecules — all of whose attendant electrons are supposed to be aUke — should produce such striking effects. The correlation between the phenomena of electrical con- ductivity and the constitutional diagram of a system of aUoys however, a very close one. If the specific electric conductivity is plotted for a binary system it is found that some very typical curves are obtained. Taking the various types of binary systems in the same order as before, we have first those in which there is a complete and unbroken series of solid solu- tions between the two metals. For such a system the curve of conductivities takes the form of a deep U, falling rapidly from the pure metals at the ends of the series and becoming more or less horizontal for the middle of the series. The depth of the U is, however, very great ; thus in the copper-nickel series, the aUoy known as " constantan " has a conductivity of only one-thirtieth of that of copper, while it consists of 60 per cent, copper and 40 per cent, nickel. In alloys of our second group, where no solid solutions are formed, and in which, with the exception of the pure metals forming the ends of the series and the eutectic alloy itself, aU the members of the series consist, in the solid state, of crystals of one of the metals embedded in the eutectic, the curve of electrical conductivities assumes a very different shape. Here the two metals are present in a state of simple mechanical mixture or juxtaposition — there is no intermingling of the molecules, and thus each retains its original electrical conductivity unaltered, and the alloy exhibits a behaviour which is the arithmetic mean of its two constituents. The curve of conductivities is thus — ^if plotted in relation to composition by volume — a simple straight line joining the values for the two simple metals. Such an actual case is probably never met with in practice, and, indeed, it can be shown that there must be a slight degree of mutual sohd solubility between any two metals — complete absence of molecular interpenetration cannot really exist. 112 STUDY OF PHYSICAL METALLURGY Consequently, although several binary systems of alloys — such as those of cadmium and zinc, and of lead and antimony — closely approach the Umiting case, in the majority of binary 10 80 90 /00%Ca Cu 20 30 40 SO 60 Au Composition. Fig. 38. — Curve of Specific Electric Conductivity typical of an unbroken series of Solid Solutions. (Alloys of Gold and Copper.) Co 30 100% Cu Cu 40 SO 60 Composition. Fig. 39. — Curve of Specific Electric Conductivity typical of a Euteoti- ferous Series. (Alloys of Cobalt and Copper.) alloys of the " eutectiferous " type, the curve of conductivity drops sharply at either end for a short distance and then becomes horizontal. This form is, of course, typical of the intermediate type of alloys in which there is a limited range in which the two metals form soHd solutions, the remainder of the systems PHYSICAL PROPERTIES OF ALLOYS 113 being eutectiferous. Here we have a drop of the conductivity curve at either end very marked, with a straight line forming the intermediate part. Typical curves of this kind are given in Figs. 38 and 39, representing data obtained from the alloys of gold and copper, which form an unbroken series of soUd solutions, and those of copper and cobalt, which are eutecti- ferous over the greater part of their range. Reference to these figures serves to show that, where the measurements of electrical conductivity have been made on a sufiicient number of alloys, it becomes possible to determine the type of the binary Fig. 40. — Diagram illustrating Need for numerous Observations in laying down Curves. system in question from the shape of the conductivity curve. The reservation as to a sufficient number of measurements is essential, however, as with a few points only it would not be possible to distinguish between the " sohd solution " and the " intermediate eutectiferous " types. This is illustrated diagrammatically in Fig. 40, where it is clearly shown how two different types of curve may agree in fitting points representing only a few determinations. The same remark also apphea, by the way, to thermal and other data and, indeed, actual mistakes in laying down a constitutional diagram have been made in several instances owing to the fact that the observed P.M. I 114 STUDY OF PHYSICAL METALLURGY points were too far apart to indicate all the important features. In the case of more complex alloys, containing either definite compounds or series of solid solutions which are based upon definite compounds, the conductivity curve assumes more complex shapes. In these alloys new sohd constituents make their appearance at various points along the axis of concentra- tion, and corresponding to these sudden changes in structure l^ CuJBi)^ 30 40 50 60 70 80 30 /00XS6 ^u Composition. Sb Fig. 41. — Curve of Specific Conductivity for the Alloys of Copper and Antimony, showing Discontinuities corresponding to definite Compounds. and constitution, the curve of conductivities generally, but not always, changes its direction. As a rule, therefore, a change of direction in the conductivity curve may be taken as indicating that one has passed into a new field or region of the equihbrium diagram, but it is not safe to conclude that, because the conductivity curve shows no deflection, no line of the dia- gram can have been crossed. Definite inter-metallic com- pounds resemble pure metals in their conducting properties to some extent — i.e., they conduct better than either mixtures PHYSICAL PROPERTIES OF ALLOYS 115 or solid solutions containing them — ^but they do not conduct nearly so well as their component pure metals, nor does their conductivity bear any numerical relation to the pure metals and the proportions in which they are present in the compound. As a rule, but again by no means universally, the existence of a definite compound in a binary system is indicated by a break in the conductivity curve, as shown at P in Fig. 41, which relates to the alloys of copper and antimony, where two compounds are indicated. Where such a break is met with the existence of a corresponding compound can be inferred with certainty, but the absence of a break is no proof of the absence of a compound. Beyond the study of the electrical conductivities of metals and alloys it has also been found interesting to determine the rate at which the conductivities change with variations of temperature. The conductivities of pure metals always decrease, in many cases with some rapidity, with rise of tem- perature, while in the case of alloys it is generally found that the " temperature coefficient " varies in accord with the electrical conductivity itself, so that a pure metal having a high conductivity exhibits a large temperature coefficient, whUe an alloy of the soUd solution type which has a very low conductivity also shows a low temperature coefficient. In- cidentally this property is of importance in connection with several well-known high-resistance alloys which possess the additional advantage of exhibiting only shght changes in resistance with variations of temperature. The alloy " Con- stantan," containing copper 60 per cent., nickel 40 per cent., which has already been mentioned, is a typical example. The use of curves representing the variation of the temperature coefficient of electrical conductivity or resistance in connection with the determination of the hnes of constitutional diagrams has been suggested, and certainly possesses some important advantages. Before leaving the subject of electrical measurements on alloys in connection with the estabUshment or interpretation of constitutional diagrams, the experimental aspect of the subject requires brief mention. Measurements of electrical i2 116 STUDY OF PHYSICAL METALLURGY properties, such as conductivity, resistance or temperature coefficient, require a knowledge of the actual dimensions of the specimen of metal employed. In the ordinary way such measurements are made on drawn wires, whose length and diameter can be readily measured to a high degree of accuracy. Even relatively short, thick rods can be used for this pur- pose, provided that the electrical instruments employed are sufficiently sensitive. In every case, however, the accuracy of the resulting measurement must depend upon the accuracy of the measurements of the dimensions of the test-piece and these in turn depend upon the regularity of the shape of the piece of metal. Where metals and alloys can be obtained in the form of drawn wires, or even of rods which can be accurately turned in the lathe, no difficulty need arise. In the middle regions of our binary series, however, the great majority of alloys are too weak and brittle, not only for the purpose of drawing into wires, but even to allow of being turned in a lathe. The attempt is sometimes made to make electrical measurements on such alloys on rods cast as nearly as may be to a circular shape. Sometimes special devices are adopted for this purpose, such as drawing the molten metal into a tube either by the appHcation of suction or pressure. In most of these cases the degree of regularity attained is insufficient for exact work, principally because in a cast metal, and particularly if it is cast in the form of a thin rod, there is no guarantee that there may not be cavities or blow-holes, either of a relatively large size or of microscopic dimensions, and their presence would interfere vitaUy with the measurements. It would require an excessive amount of labour to determine by subsequent sectioning and microscopic examination that any given specimen had really been free from cavities. An accurate density determination is somewhat simpler, but is not quite conclusive, as no really reliable standard of comparison is available. A number of other physical properties of metals have also been utilised in connection with the study of alloy systems. These can only be briefly mentioned. The mechanical pro- perties, such as hardness, tensile strength, elastic properties, etc., are aU closely related to the position of the aUoy in the PHYSICAL PROPERTIES OP ALLOYS 111 constitutional diagram, but, with the possible exception of hard- ness, these properties are not measured with sufficient ease and rapidity to invite their use in exploratory work. There is the further difficulty that the mechanical properties of any metal depend to a very large extent upon the mechanical and thermal treatment which it has undergone, so that numerous factors apart from chemical composition and constitution enter into these matters. The study of the mechanical properties has, therefore, been undertaken, in the great majority of cases, for the sake of the value attached to a knowledge of those important properties for their own sake rather than as a guide to the constitution of a system of alloys. This whole question is therefore treated in a separate chapter. Other physical properties, such as specific heat, thermal conductivity, etc., have been employed for the study of metals and alloys, but not to an extent which would justify their treatment here. The measurement of thermo-electric power is a more promising method, but labours under certain diffi- culties which are hkely to prevent its wide adoption. Rather more consideration must, however, be given to some electro- chemical and purely chemical methods of studying alloys. The electro-chemical method consists in making determina- tions on a series of alloys of the electric potential which is set up when the specimen is used as an electrode in a specially- arranged voltaic cell. By this means a concentration-potential curve (") can be set up, and it is claimed that on this ciu:ve the exact concentrations at which certain constituents appear or disappear are clearly marked, and that the existence or otherwise of definite compounds can be readily inferred. The method has not as yet been widely applied, but it appears to offer some advantages, although it suffers from one serious difficulty in that the potential observed in a gifen case must depend solely upon the constituents which are exposed, at the time of experiment, to contact with the electrolyte of the cell. If the alloy is of uniform composition and the structure is fine, then the surface in contact with the electrolyte is almost certainly adequately representative of the whole, but if the constitution is coarse in structure, and particularly if one of the 118 STUDY OF PHYSICAL METALLURGY constituents exhibits a tendency to be surrounded by or embedded in another, errors may arise. The objection just raised applies with still greater force to the purely chemical method of studying alloys which consists in the endeavour to isolate the crystals of any definite com- pound which they may contain. It frequently happens that an inter-metaUic compound is considerably less soluble in some reagent than the rest of the alloy, and prolonged exposure to such a reagent should produce a residue consisting entirely of the isolated crystals of the insoluble compound. This method was extensively used before the modern methods of thermal and microscopic analysis had been developed, and a very large crop of so-called " definite alloys " or compounds was obtained by the older workers. Lists of these may still be found in many of the older text-books of chemistry. The modern methods have, however, shown that the great majority of these com- pounds were entirely fictitious, and that no place for them can be found in any well-established constitutional diagram. This state of affairs has rightly led to the method being largely discredited, and the reasons for its failure are not far to seek. A study of the micro-structure of alloys containing compounds serves to show that we could not expect to separate them from the surrounding metal by such a process of differential solution. The crystals in many cases completely enclose particles of the matrix, which is thereby entirely protected from solution. In other cases, particularly where a chemical reaction occurs between the crystals first deposited from solution and the residual mother-liquor, the resulting compound or product forms a relatively thin sheath around a core of totally different composition. In stiU other cases, where an alloy consists of a soUd solution which has not been given time to attain equili- brium during the coohng process, the outer margins of crystals may be more or less soluble than the cores, and this difference will vary with the size of each individual crystal. If an attempt is made to separate a " compound " from such a material a certain residue will be obtained, but its composition wiU not be that of any definite compound found in the alloys in question. The whole method is thus beset with serious PHYSICAL PROPERTIES OF ALLOYS lid difficulties and used by itseK must be regarded as decidedly weak, although in careful combination with the study of the micro-structure valuable data can be obtained by its aid. The discovery of the composition of the carbide of iron known as " Cementite " by Abel, in 1885, was made by this process of residue analysis and must stand as a notable achievement to its credit. Any general review of the subject of the constitutional study of alloys and of the construction of equilibrium diagrams would be incomplete without some reference to the thermo- dynamic principle of physical chemistry known as " the phase rule " (^^). This principle is sometimes employed as the basis for the study of the whole subject treated in the present chapter, but the writer has thought it preferable to approach the subject from the experimental and observational side rather than from a purely theoretical basis. The phase rule is a simple arithmetical formula, deduced from the first principles of thermo-dynamics, which enables us to state what will be the behaviour of a mixture of any kind under given changes of conditions. For this purpose the constituents of an alloy are classified as "phases," any constituent which is spacially distinct from the surrounding region being termed a " phase." Each phase within itself must be regarded as completely homogeneous and of one and the same chemical composition throughout. Thus in a liquid or molten alloy (except in the rare cases where two metals are present which do not mix when molten) the entire system consists of a single phase, viz., the Hquid. A perfectly pure metal after solidification again consists of only one single pheise ; during the process of solidi- fication, however, since both hquid and solid are present simultaneously, the system consists of two phases. But two phases present at the same time need not differ in their physical state ; thus every eutectic alloy, containing juxtaposed crystals of two metals or of two solid solutions, consists of two phases, since the two kinds of crystals are spacially distinct from one another and are of different chemical composition. The " phase rule " tells us how many such phases can be present at any one time, when the alloys are in complete equiH- 120 STUDY OF PHYSICAL METALLURGY brium, under given conditions. These " conditions " are determined by the three factors which affect the constitution of an alloy, viz., chemical composition or " concentration," temperature and pressure. If any one of these conditions can be changed without bringing about the complete disappearance of an existing phase or the appearance of a new phase, then the alloy is said to " possess a degree of freedom " in respect to that condition. An alloy can, at most, possess three degrees of freedom, i.e., one with regard to each of the conditions " con- centration," "temperature," and "pressure." If we denote this " degree of freedom " by the symbol /, the number of phases present by the letter p, and the number of component elements {i.e., metals) as n, then the " phase rule " may be written f=n-p + 2, from which it will be seen at once that the greater the number of phases present the fewer will be the remaining degrees of freedom. It should also be noted at once that while the value of / can never exceed 3, it can also never become negative, i.e., less than 0. This sets a limit to the possible number of phases which can ever be present in a mixture containing a definite number of components. In a binary alloy, where n = 2, p can consequently never be greater than 4. In the special case of alloys, with which we are solely con- cerned here, one of the " conditions " may be disregarded, as we do not usually vary the pressure far from that of the atmo- sphere. When we consider the effects of stress, of course, this may no longer be true, but for purposes of equihbrium diagrams the pressure factor may as a rule be disregarded. If this is done the system can never have more than two degrees of freedom, and the formula of the phase rule becomes f = n — p + 1. For the simple binary system, where n = 2,p can consequently never exceed 3 when the alloys are in equilibrium. This number of phases, three, can, however, co-exist only when/= o and that implies that a mixture of three phases in a binary system possesses no degree of freedom, and can, therefore, exist at one definite temperature and concentration only. In the equih- PHYSICAL PROPERTIES OF ALLOYS 121 brium diagram this fact expresses itself in the circumstance that four regions or " phase fields " can only meet at a single point which represents a fixed temperature as well as a fixed composition or concentration. In a simple binary eutectiferous series the only point where three phases co-exist is at the melting-point of the pure eutectic, since there we have hquid in equihbrium with solid crystals of both component metals at the same time. In the case of a pure metal, where n = 1, f cannot exceed 2, and this number only occurs at the melting- point. This subject need not be pursued further here, especially as its intricacies would require much space for their adequate elucidation. A tendency exists in some quarters to attach exaggerated importance to the " phase rule " as a guiding principle in the study of alloys, but, although of undoubted value, there are certain definite hmitations. The greatest of these is that the rule is only fully apphcable to alloys in com- plete equilibrium, and such complete equilibrium is rarely, if ever, strictly attained. If, therefore, observation indicates that the number of phases which appear in given circumstances is at variance with the deductions from the phase rule, the explanation is at once available that the alloys as examined are not in final equihbrium. As a rule it must be admitted that efforts to bring the alloys to a true equihbrium condition always show at least a tendency towards the ehmination of one of these redundant phases. In practice, however, many of the most interesting, and also the most puzzhng, features in alloy systems are associated with conditions of imperfect equihbrium. Indeed, it may fairly be said that aU materials which find practical appHcation in the arts depart more or less from a state of physico-chemical equihbrium. So weU recognised has such a condition become that its persistence is admitted in the term used to describe it, viz.. " meta- stable." The alloys employed in a meta-stable condition include every kind of steel and almost all the more complex alloys. The persistence of these meta-stable conditions is, of course, due to the fact that most metals at the ordinary temperature are sohd bodies possessing a liigh degree of rigidity or viscosity. 122 STUDY OF PHYSICAL METALLURGY i.e., offering very great internal resistance to molecular re-arrangement of any kind. If, therefore, an alloy has once been cooled down to the ordinary temperature at such a rate that full equilibrium has not had time to estabUsh itself during the cooUng process — and excessively slow rates of cooling would in most cases be required in order to allow equihbrium to be attained — the molecules and atoms of the material are no longer sufficiently free to undergo re-arrangement and the alloy retains its condition of meta-stabihty. Examination of ancient metal objects, particularly of some from Egypt, indicates that this persistence of meta-stable conditions may last over many thousands of years and may for practical purposes be regarded as permanent. It is only by raising the temperature of the metal to a point high enough to afford the molecules sufficient freedom to imdergo changes of arrange- ment, and maintaining the metal at such a temperature for a sufficiently long time — possibly many months — that equihbrium can be attained. These considerations naturally raise the question whether our constitutional diagrams should rightly relate to these ultimate conditions, only attainable in ideal circumstances, of complete physico-chemical equilibrium, or to the practically far more important meta-stable conditions. There can be no doubt as to the answer : equihbrium diagrams must relate only to really stable conditions of complete equihbrium. This is necessary, since in those circumstances only is the constitution of an alloy a fixed and definite thing ; in the case of meta-stable condi- tions, however permanent they may be, there is no definitely fixed condition — ^the alloy will approximate more and more closely to the ideal equihbrium condition the slower has been the rate of coohng or the more prolonged the period of heating or anneahng. Consequently, the true equihbrium diagram, indicating the final equilibrium condition, stands as the limit towards which the alloys tend and in terms of which it is alone possible to express or to describe their intermediate meta- stable conditions. Indeed, the skilful interpretation of any equilibrium diagram will serve as a guide to the structures and constitutions which the aUoys are likely to assume in their PHYSICAL PROPERTIES OP ALLOYS 123 more usual meta-stable state resulting from ordinary practical rates of cooling or even from intentionally accelerated cooling, such as quenching. An endeavour to put this interpretation of the equilibrium diagram upon a correct quantitative basis has recently been made by Gulliver ("), and if his methods can be sufficiently simplified for practical purposes, they should lead to a still wider and more direct usefulness of the equili- brium diagrams. Our consideration of the constitution of alloys has so far been confined, for the sake of simphcity, to binary systems — i.e., to alloys of two metals only, In practice, of course, alloys only rarely belong to a simple binary system, since the majority of industrial materials contain three or four component metals or metalloids, in addition to impurities present in greater or lesser amount. From the scientific point of view also, the study of alloys must be pushed forward to include systems of three component metals (ternary aUoys), and even of four or more. The difficulty, however, Hes in the rapidly increasing complexity of all constitutional and structural conditions when the number of components is increased. For the ternary alloys we still have available a method of clear and compara- tively simple graphical representation, but for alloys of four or more metals this aid fails us so far as any comprehensive view is concerned, and the representation can only be made piecemeal. In addition to this difficulty the fact must also be faced that for the study of a ternary system an enormously larger number of experimental determinations and observa- tions must be made. If for the adequate investigation of a binary system the study of fifty separate alloys is found necessary, then a proportionately close study of a ternary 50 X 50 system would require the examination of ^ — 1)250 alloys — ^in itself a stupendous task. In spite of this difficulty the study of ternary systems is now being undertaken vigor- ously, but it must be admitted that the closeness of observa- tions is still far behind that which has been found desirable in the best investigations of binary systens, so that the results on ternary aUoys obtained by the majority of workers up to the 124 STUDY OF PHYSICAL METALLURGY present must be regarded rather as preliminary surveys than as exhaustive studies. The graphical representation of the constitution of a ternary system of alloys is accomphshed by means of a " constitutional model " in three dimensions in place of the plane constitutional diagram which suffices for systems of two components. Where the binary diagram is erected upon a single line as a basis, points along the Hne indicating the composition of the mixture, the ternary model is erected upon an area the points in which Fig. 42.- -Triangular Diagram for Plotting the Composition of Ternary Alloys. indicate the composition of the triple mixture. This can be most conveniently done by employing for the basis of the equilibrium model an equilateral triangle drawn to such a scale that its height is taken as 100. It is a well-known property of such a triangle that the sum of the perpendicular distances of any point within it from the thiee sides is equal to the height of the triangle and is, therefore, constant over the whole area of the triangle. These perpendicular distances may thus be used to indicate the percentage composition of every alloy of a ternary system. In the triangle ABC of Fig. 42 PHYSICAL PROPERTIES OF ALLOYS 125 perpendiculars have been dropped upon the three sides from each of the three corners of the triangle, and each of these perpendiculars has been divided into ten equal parts. If now Unes are drawn through each of these divisions they will mark all the points lying at equal distances of 10, 20, 30, etc., from one of the three sides. This distance in every case represents the percentage of that metal present in the alloy which is placed in the diagram at the opposite corner of the triangle. Thus the perpendicular Ax dropped from the corner A upon the opposite side has its second division, representing 20 per cent., at the point P. The hne RPT drawn through this point parallel with the side BC of the triangle thus connects all the points in the triangle whose composition is such that they contain 20 per cent, of the metal A. Similarly the hnes drawn through other points of that perpendicular, or through corre- sponding points of the two other perpendiculars, connect aU the points having a constant proportion of one metal. The outside Unes of the triangle itself thus obviously represent the three simple binary systems formed by the three component metals taken two at a time, the corners representing the pure metals themselves, exactly as the end vertical hnes of a binary diagram represent the pure metals. Since the composition of every possible member of a ternary system can thus be represented by a point in an equilateral triangle, the constitution of the ternary system can be repre- sented in exactly the same way as that which we have seen employed in the case of the binary systems, by the erection of verticals upon the points representing alloys which have been studied. In this way melting or freezing-points, eutectic points, ^transformations and, indeed, every feature in the equihbrium conditions of an alloy can be represented by vertical plotting above the point representing its composition. But while in the simple binary case these verticals were erected on a single Une and could thus be shown on a single sheet of paper, in the ternaries the result will be a model in which surfaces take the place of the simple hnes of the binary diagram. Thus we may erect verticals representing the tem- perature at which solidification begins, and the ends of these PHYSICAL PROPERTIES OF ALLOYS 127 tion of the ternary model in the first instance. Thus one might begin with a binary series of metals A and B and then determine a quasi-binary system — ^in reality a section of the ternary model — corresponding to the hne of the triangle representing the system of A and B plus 5 per cent, of C, and following this throughout the whole system by successive fines lying parallel to AB. The data thus obtained will, of course, also make it possible to plot the sections of the model along fines paraUel to either of the two other sides. The series of sections thus obtained may then be combined into a sofid model ; a very convenient method is to erect sheets of thin, clear glass or celluloid verticaUy upon a board graduated to represent the basal triangle and then to draw the successive sections upon the corresponding sheets of glass. The transparency of the whole is then sufficient to enable the observer to obtain a good idea of the true ternary equifibrium surfaces. The apparent complexity both of these methods and of the resulting surfaces might perhaps be regarded with some mis- giving from the point of view of the utifity of such results, but fortunately in those cases hitherto investigated, and particularly in those relating to the more important afioys, considerable simpfification is found applicable. We have already indicated that in binary aUoys the more interesting and important members always fie near the two ends of the series. Similarly it is found that the majority of ternary and of complex afioys consist predominatingly of one metal, or at most diverge only sUghtly from a binary system. Within these fimits it has been found that in the majority of cases the form of the ternary surfaces is fairly simple and of such a nature that the sections of the ternary model lying near the outer fines of the triangle are not very widely different from the constitutional diagrams of the fimiting binary series. The actual temperatures of the various curves are, of course, altered more or less materiaUy, and the various phase-fields are somewhat changed in shape and extent ; but it is rare to find new phases entering into the constitution of such aUoy groups. For that reason it becomes possible to regard these ternary aUoys as modifications of the nearest similar binary afioys by merely applying certain 128 STUDY OF PHYSICAL METALLURGY modifications or " corrections " to allow for the effect of the third metal. The conclusion to be drawn from what has just been said is that the detailed study of ternary alloy systems from the equihbrium point of view is not as yet of vital importance except from the purely theoretical point of view. It is, of course, far too early to suggest that the fuUer study of the more complex systems may not reveal the existence of materials of unexpected interest. Indeed, some such materials have already been found, as, for instance, the Heussler magnetic alloys and — ^far more important — the whole series of ternary alloy steels. Where such interesting materials are to be met with, the full study of the ternary and even of more complex equUibria is urgently wanted. Thus our whole understanding of -the alloy steels is at the present time vague and our know- ledge scrappy for want of a complete and systematic series of investigations of these difiicult ternary systems. The only reason why these investigations have not been undertaken hitherto Hes in the fact of their great experimental difficulty and expense. The complete and accurate survey of such systems as iron-carbon-manganese, iron-carbon-nickel, and other similar ones would undoubtedly clear up many of the existing doubts and difficulties ; it is, however, for those commercially interested in these materials to see that the necessary time and money are made available for their study. Kefekences. (1) Guertler. Journ. Inst. Metals, X., 1913, 2, p. 175. (2) Charpy and Grrenet. Bull. Soc. d' Encouragement, 1903, CII., 464, 883. (3) Kosenhain. Journ. Inst. Metals, X., 1913, No. 2, p. 190. (4) Maey. Zeitsohr. Phys. Chemie, 1899, 29, 119; 1901, 38, 289, 292 ; 1904, 50, 200. Vincentini and Omodei. Atti. E. Accademia deUe Science di Torino, 1887, XXIII., p. 8. Retgers. Zeitschr. Phys. Chem., 1889, III., 497. Topler. Annalen der Physik, 1894 (III.), LIII., p. 343. Liideking. Annalen der Physik, 1888 (III.), XXXIV., p. 21. (5) Turner. Journ. Iron and Steel Inst., 1906, 1., p. 48. Turner and Murray. Journ, Inst, Metals, No. 2, 1909, II , p. 98. PHYSICAL PROPERTIES OF ALLOYS 129 Ewen and Turner. Journ. Inst. Metals, No. 2, 1910, IV., p. 128. Haughton and Turner. Journ. Inst. Metals, No. 2, 1911, VI., p. 192. Chamberlain. Journ. Inst. Metals, No. 2, 1913, X., p. 193. (6) Curie (P.). Annales de Chemie et de Physique, 1895 (7), V., 289. Morris. Phil. Mag., 1897 (V.), 44, 213. Guertler and Tammann. Zeitschr. Anorg. Chemie, 1904, 42, 353. Hadfield. Proo. Inst. Civil Engineers, 1888, 93, III., 1. Curie (S.). Etude des AlUages, 177. Hopkinson (J.). Proc. Roy. Soc, 1890, 47, 23. Guillaume. Comptes Rendus, 1897, 124, 176, 1515 ; 125, 238 ; 1898, 126, 738. Etude des AUiages, 459. Dumas. Comptes Rendus, 1900, 130, 357. Dumont. Comptes Rendus, 1898, 126, 741. Tomlinson. Proo. Roy. Soc, 1884, 56, 103. (7) Heussler. Verh. Deutsoh. Physikalisch. Gesellschaft, 1903, V., 219. Stark and Haupt. Deutsch. Physikalisch. GeseUsohaft, 1903, V., 222. Heusler. Zeitsohr. Angew. Chemie, 1904, 260. Take. Verh. Deutsoh. Physik. GeseU., 1905, VII., 133. (8) Hadfield, Barrett & Brown. Roy. Dublin Soc, 1900 and 1902. Inst. Electrical Engineers, 1902, XXXI. (9) Ewing. Proc. Roy. Soc, 1890, 48, 342. (10) Weiss. Journ. de Physique, November, 1911. Archives des Sciences Physiques et NatureUes, 4, XXXI. Trans. Faraday Soc. VIII., 1912. Soci6t6 Fran^aise de Physique, 1913. (11) Guertler. Journ. Inst. Metals, No. 2, 1911, VI., p. 135. (12) Laurie. Trans. Chem. Soc, 1888, 53, 104 ; 1889, 55, 667 ; 1894, 65, 1031. Phil. Mag., 1892 (V.), 33, 94. Zeitschr. Phys. Chemie, 1909, 67, 627. Herrsohkowitz. Zeitschr. Physik. Chemie, 1898, 27, 123. Bijl, Zeitschr. Physik. Chemie, 1902, 42, 641. Reinders. Zeitschr. Physik. Chemie, 1903, 42, 225. Pnshin. Journ. Russian Phys. Chem. Soc, 1907, 39, I., 13, 353, 528, 869. Zeitschr. Anorg. Chem., 1907, 66, 1. Pushin and Pashky. Journ. Russ. Phys. Chem. Soc, 1908, 40, 826. Pushin and Laschtschenko. Journ. Russ. Phys. Chem. Soc, 1909, 41, 23. Zeitschr. Anorg. Chem., 1909, 62, 34. (13) Gibbs. Amer. Journ. Science, XVI., 1878. Jiiptner von JornstorS. Stahl u. Eisen, 1899, XIX., 23. P.M. K 130 STUDY OP PHYSICAL METALLURGY Le Chatelier. Comptes Rendua, 1900, 130, 85 ; 1899, 30, 385, 413. Roozeboom. Zeitschr. Phys. Chem., 1900, 34, 437. Journ. Iron and Steel Inst., 1900, II., 311. (14) GuUiver. Journ. Inst. Metals, 1913, No. 1, IX., p. 120 ; 1914, No. 1, XI. (15) Charpy. The Metallographist, II., 1899, p. 26. BuU. Soc. d'Encouragement, June, 1898. (16) Rosenhain and Lantsberry. Ninth Report to the Alloys Research Committee. Proc. Inst. Mechanical Engineers, 1910. CHAPTER VII TYPICAL ALLOY SYSTEMS In the preceding chapters the constitution and structure of alloys has been treated in a very general way, with only passing reference to individual series of alloys for purposes of illustra- tion. The whole subject will, however, be rendered much clearer by treating in somewhat greater detail the constitution and structure of a few selected systems. Not only will this serve to show the appHcation of the principles and methods described in the previous chapters, but it will introduce the reader to some of the more important facts concerning those alloy systems with which the physical metallurgist is most frequently called upon to deal. If one were to select the alloy systems solely from the latter point of view, there can be no doubt that the alloys of iron and carbon would call for attention before any of the others. That particular system is, however, of a very complex kind, and for that reason it is preferable to postpone our detailed consideration of the iron-carbon alloys until a shghtly later stage, when our acquaintance with other alloys will make it easier to follow the intricacies of that most important system. The simplest type of binary alloy system, as we have seen in the last chapter, would be one of the purely eutectiferous kind in which the two metals are completely mutually insoluble in the solid state. The equilibrium diagram for an ideal case of this kind has already been given (Fig. 30, p. 91), and it is beUeved that the alloys of lead and antimony approximate very closely to this ideal limit. But even in those cases it is probably only a question of sufficiently close and accurate study to -discover that the real case departs more or less widely from the ideal hmit. In the older books and papers dealing with this subject it was customary to quote the alloys of lead and tin as typical examples of «. binary series of this simple k2 132 STUDY OF PHYSICAL METALLURGY type. More careful study has, however, shown that at the lead end of this series at all events there is a considerable degree of sohd solubility, lead forming with tin a series of solid solutions containing up to about 16 per cent, of tin (^). The earher mistaken ideas about these alloys arose from the fact that when lead-tin alloys lying near the lead end of the series are allowed to cool from fusion even at a slow rate, both the thermal arrest due to the soHdification of eutectic and the corresponding eutectic structure in the sohd alloy can be detected. This, however, is solely due to the slowness with which these alloys approach their true equihbrium condition. In the author's experiments on this series it was found that heating at a tem- perature of 175° C. for a period of six weeks was required in order to bring the aUoys approximately into their final state. The fuller study of the lead-tin series has also shown that the simple ideal diagram requires modification in yet another way. The aUoys are found to exhibit an evolution of heat on cooling and a corresponding absorption of heat on heating at a tem- perature which lies, for alloys containing from 16 to 63 per cent, of tin, at 149° C. An explanation of this fine of arrest-points has been put forward by Mazotto (^) to the effect that they are due to a species of " under coohng " in consequence of which the eutectic does not completely sohdify at the true eutectic temperature, but remains in a meta-stable Hquid condition untU the hmiting temperature of 150° C. is reached. It is difficult to see how such a process of under cooUng can possibly account for the absorption of heat at the same temperature when the alloys are again heated, and there are several other vital objections to this view. In the same way the view of Degens (*) that a chemical compound is formed in the alloys at the temperature in question will not bear critical examination in the fight of the data obtained by the author. If a compound were involved then the amount of heat evolved on cooUng through this critical temperature would attain a maximum value for alloys whose composition approached that of the compound in question ; a maximum does occur, but the com- position of the alloys near this maximum is not consistent with any rational formula for a compound. The only remaining TYPICAL ALLOY SYSTEMS 133 explanation is that put forward by the author — that the critical temperature in question arises from a change in the metal lead which at that temperature alters in its power of dissolving tin in the solid state. This explanation is embodied in the consti- tutional diagram given in Fig. 45, which is slightly modified from that first published by the author and Tucker, in order to bring it nearer to the requirements of the phase rule. Further study of these alloys may well throw fresh fight on these phenomena and result in a more satisfactory explanation, so 3S0 "^ISO too fi \ (i ■^ Eutectic O Pb. W * Eubectic Sn*Eulectic Sa +Eutectic 30 «<7 SO 60 Composition 70 80 90 IOO%Sn. Pig. 45. — Constitutional Diagram of the Lead- Tin Alloys. that the diagram below the eutectic fine must be regarded as somewhat tentative. The micro-structures associated with the various fields of the lead-tin equifibrium diagram are very typical of the struc- tures met with in all eutectiferous alloys. They have already been illustrated in Figs. 24 to 27. A further example of these alloys is given in Fig. 46 (Plate IX.), which shows one of the structures typical of the lead-tin eutectic, having the laminated character often associated with eutectics. This laminated character, however, although often beautifully marked, is not to be regarded as an essential feature of these substances. If a specimen of a well-laminated eutectic is heated for a long time 134 STUDY OF PHYSICAL METALLURGY to a temperature a few degrees below its melting-point, the structure gradually changes, various laminae coalescing to form lumps or globules, and the result is a much coarser granular structure. Such a granular structure may also be obtained if the molten eutectic is very slowly cooled. The eutectic alloy is the result, so far as its structure is concerned, of the simultaneous crystallisation or freezing of the two component metals. As a rule one of these acts as the " predominant partner," and its own crystalline form or habit determines the way in which the whole structure is arranged. It has been shown that eutectic aUoys, like pure metals, consist of an aggregate of juxtaposed crystals. These crystals are, however, merely skeletons formed of one of the metals — the " predominant partner " — with the interstices filled in by the other metal. In the case of the lead-tin alloys the tin is the predominant metal, and each of the crystals of the eutectic is in reahty a radiating structure, known as a " spherulite," of tin carrying the lead in its interstices. The micro-structures illustrating the lead-tia series are typical of other similar alloys, such as those of lead and anti- mony. An interesting practical apphcation of certain alloys of this latter series has recently been suggested by Hannover (*) in the production of porous metal, principally for use in the construction of electric storage batteries. For this purpose he utihses an alloy somewhat Uke that illustrated in Fig. 25 while it is at a temperature lying between the lines AE and CE of the diagram (Fig. 30) — i.e., within the temperature range where the crystals of one metal (in this instance lead) are already soHd but the eutectic is still fluid. By centrifuging the alloy at this temperature Hannover has found it possible to drive out the fluid eutectic, thus leaving behind a porous mass of lead. The constitutional diagram of the lead-tin series is also of interest from the practical point of view, owing to the use of some of these alloys for soft-soldering. The most fusible and also the most homogeneous solder is formed by the eutectic alloy of the series, containing 37 per cent, of lead ; owing to the high price of tin, however, solders nearer the lead end of the TYPICAL ALLOY SYSTEMS 135 series are more frequently employed, and for certain purposes offer a decided advantage. The diagram shows that the eutectic alloy soUdifies at a single definite temperature, but an alloy containing equal parts of lead and tin has a solidification range of 60° C. between the temperature at which crystals of lead first begin to separate and that at which sohdification is completed by the freezing of the residual eutectic. In such work as the " wiping " of joints as practised by plumbers, this range of temperature, in which the alloy (or solder) is in a pasty stage, is utilised by the workman to bring the mass gradually into the desired shape around his joint. From the lead-tin series of alloys a number of ternary and complex alloys are derived, some of which are of considerable practical importance. Among these are the aUoys of lead and tin with antimony, which are largely employed for type-founding and for the production of bearing-metals. We cannot go into the structure and constitution of these ternary alloys in any detail, but it may be mentioned that these alloys containing antimony largely owe their value to the existence of a definite compound between tin and antimony, having the formula SnSb {^). This is a hard, brittle substance which appears in the alloys in the form of angular crystals. The presence of these hard crystals tends to make the alloys much stronger, stiffer and harder, but also more brittle. For type-metal it is necessary to secure adequate hardness and at the same time to have a metal which will fill the minute interstices of a mould with great accuracy. The actual composition of type-metal employed varies widely with the special purpose for which it is required and the consequent limitation which has to be imposed upon its price. This consideration governs the amount of tin which can be intro- duced. The structure of the aUoy is always rather complex, but consists of primary crystals of either antimony or of the antimony-tin compound, embedded in two layers of eutectic. In the case of white bearing-metals the requirements are diffe- rent — what is wanted is a metal sufficiently soft and plastic to accomodate itself to the inequalities of the moving parts and thus to afford them an even bearing, while it must also present 136 STUDY OF PHYSICAL METALLURGY a surface of sufficient hardness to prevent rapid abrasion. The crystals of the tin-antimony compound serve to provide the resisting hard surface, while the soft and plastic eutectics in which they are embedded yield to the pressure and adapt themselves to the needs of the bearing in question. The typical structure of an alloy consisting of the hard crystals of SbSn embedded in eutectic is shown in Fig. 47, Plate IX. We now pass on to consider the constitutional diagram and the correlated micro-structures of a decidedly more complex system of alloys, viz., those of aluminium and zinc. This 10 20 30 40 SO 60 70 80 30 WO%Cl Zn Composition. Al Pig. 48. — Constitutional Diagram of the Zinc-Aluminium Alloys. system still shows a weU-defined eutectic, but at the one end of the series there is a long range of sohd solutions, while the series also shows one definite compound possessing somewhat remarkable features. This series is of interest £is forming the basis of what are probably the best available light alloys of aluminium, some of which combine a comparatively low density (about 3* 25) with the strength and toughness of mild steel. The constitutional diagram is shown in Fig. 48. In this case also a much simpler diagram was for a long time generally accepted ; the author and Archbutt (*), however, established the diagram here shown which is itself not quite complete in TYPICAL ALLOY SYSTEMS 137 regard to the region occupied by the dotted continuations of the Unes CGH and IJKL. The diagram consists first of all of the freezing-point curve or liquidus ABCD. Starting from the zinc end, this runs down to a eutectic point B at a concentration of about 5 per cent, of aluminium. From that point the liquidus runs up smoothly to the point C, where a small break or kink is found, and then again smoothly to the point D, corresponding to pure aluminium. The " solidus " or curve bounding the region in which the alloys are completely solid is given by the lines AEBFGHD. The curious feature in this curve is the sudden step-up at the point F corresponding to a concentration of about 78 per cent, of zinc and 32 per cent, of aluminium, which represents the compound AlgZuj. The reason for this step-up becomes quite plain when the manner in which the compound AljZug is formed in these alloys, is understood. Along the branch of the liquidus between C and D the alloys begin to solidify, on coohng, by depositing crystals of a solid solution of zinc in aluminium, this substance being called 7 in the diagram. As the alloys cool further they continue to deposit an increasing quantity of this y body until they reach the temperature 443° C, which is indicated in the diagram by the horizontal line CGH. Alloys lying to the right of the point H, i.e., containing more than 60 per cent, of aluminium, become completely solid before they reach this temperature, and in that case it is possible that no further change takes places in them, although the dotted continuation of the line GH is intended to suggest that something does occur even there. To the left of the point H, however, the alloys are stiU partially liquid when the temperature 443° C is reached and then a chemical reaction sets in, resulting in the absorption of the residual liquid and the formation of the compound Al^Znj. This reaction may be written 7 -f Liquid = AljZug. The crystals of 7, as a result of this reaction, become coated with a layer or sheath of the newly-formed compound, and as soon as such a sheath has become completed the reaction in the ordinary way comes to an end, even though the residual liquid is not used up. Li this case the alloy remains in a meta- 138 STUDY OF PHYSICAL METALLURGY stable condition and cools down to the temperature of the line BPP, where the residual liquid solidifies as eutectic. This is indicated by the dotted line FP which was formerly regarded as the true eutectic line. If, however, these alloys are kept at a temperature just below 443° C. for a considerable length of time, the formation of the compound M^Zn^ is completed, by the aid of slow diffusion, in spite of the protecting sheaths above described, and when an aUoy so treated is further cooled there is no residue of liquid and consequently no further freezing takes place on the line FP. This statement holds good for all alloys which contain less zinc than the quantity required to convert the whole of the alloy into the compound AlgZug — i.e., for alloys lying to the right of the line FG. To the left of that line, however, there is an excess of zinc, and no amount of prolonged heating will cause the alloys to become completely solid at temperatures just below 443° C. In other words, even after the whole of the 7 crystals have been converted into compound, a residue of liquid remains and this solidifies as a eutectic of zinc and the compound along the line BF. It is for this reason that the " solidus " steps up at FG from the line BF to the line GH. The line GH reaUy indicates the temperature which limits the stable existence of the compound AlgZug, so that when alloys cool down through that temperature the compound is formed, and when they are heated up through it the compound is again decomposed. At the point C this line cuts the liquidity curve and we see that to the right of C the compound has decomposed, on heating, before the alloys are completely hquid, while to the left of C the compound melts without having previously been decomposed. It follows that, con- versely, on freezing, the alloys to the left of the point G begin to solidify by depositing crystals, not of the 7 solid solution, but of the compound AlgZnj. This difference accounts for the small break in the liquidus curve at C. The diagram contains another horizontal line IJKL, lying at a temperature of 256° C. This line marks the lower limit of stability of the compound AlgZng. On cooUng down past this temperature the compound breaks up into a duplex structure, PLATE X. Fk;. ol. -^vj' "','';• ;•■ ■■-'■iff*:' ■^i ,: ■?i; <■< ^7/f/ 'oyco, }laim gSS, lP.Zn ^ ^ \ •( ■ ' i ■ - CO- 1 \wo [\ 1 fo/5 ? ^ ^ \ / \ /y7 ' « JJ; ?/.;„ ^ c: \ \ \ \ bS^ FN I OIOO ^500 \ \ ^ s, \ *^ ^ ^^ \ ^ '*- , Ki3. sCj: ^oZt2, 200 - ,_.v^ * C •^ loi. // 20 W 60 20 'to 60 W 60 80 W 60 60 100 Time Interval (seconds) Fig. 50. — Typical Cooling Curves of Zino-Aluminium Alloys. exhibit the typical polyhedral structure of a simple metal or of a solid solution. Having described thus briefly some of the principal features of the constitutional diagrams of two very typical systems, we now pass on to consider the diagrams and constitution of some of the more important alloys of copper which in industrial importance, at the present time, rank next to steel itself. The alloys of copper with zinc, tin and aluminium all afford examples of somewhat complex constitutional diagrams which it will not be possible to discuss in any detail in the present volume, and TYPICAL ALLOY SYSTEMS 141 attention will be very largely confined to those regions of the diagrams which relate to the more important alloys of the series in each case. For fuller details reference must be made to special works dealing with each group of alloys. The constitutional diagram of the aUoys of copper with zinc is shown in Fig. 54. Essentially this is the diagram estabhshed by Shepherd (') and practically universally accepted ; a modification due to the work of Carpenter {^) on the yS phase 1 1 Id 20 30 40 50 , GO 7d 80 Cu Composition. Fig. 64. — Constitutional Diagram of the Zinc-Copper Alloys. 90 IOO%ln Zn of these alloys has, however, been included, affecting the Unes in the region occupied wholly or in part by the j8 phase. The diagram shows in all six phases, but we shall only deal with the first three of these, viz., those denoted in the diagram by the letters a, 13 and y. The first of these bodies is the a phase ; this is a sohd solution of zinc in copper, and may contain as much as 37 per cent, of zinc. The position of the point 6 in the diagram indicates the limiting composition of the alloys which never contain anything but the a body — i.e., 142 STUDY OF PHYSICAL METALLURGY from the moment of incipient solidification down to the ordinary temperature, whether cooled slowly or quickly, the alloys containing less than 30 per cent, of zinc always consist entirely of the pure a body, but it does not follow from this that the alloys are necessarily entirely homogeneous in all circumstances. If we recall the process of freezing, as already described for a solid solution, we see that in the case of an alloy containing, say, 20 per cent, of zinc, the solid which first separates when the alloy begins to freeze wiU have approximately the com- position 8 per cent, zinc, 92 per cent, copper. If the cooling at this stage is not slow enough to allow of the attainment of complete equilibrium, the central portion of each crystal, i.e., the portion first formed, will remain to the end considerably richer in copper than the outer or later-formed portions of the crystals. Each crystal of metal of such composition will, in these circumstances, consist of a core which is richer in copper and an outer portion which is poorer in copper, than the average composition of the alloy. A quickly-cooled alloy in the cast condition thus shows an apparent duplex structure, although consisting of only one phase, for although the layers of soUd solution of varying concentration which form these crystals are all forms of the same a phase, their different con- centration renders them susceptible to the attack of etching reagents to different degrees ; by most reagents the regions richest in copper are less attacked than those which contain a larger proportion of zinc. Since those portions of the crystals which are first formed consist of the dendritic arms and branches already described as occurring in the early stages of the freezing of a metal, the etched pattern of such an aUoy reveals this dendritic structure, and in some circumstances this may be almost as clearly defined as if the outer regions of the crystals really consisted of a different phase. An example of this kind has already been shown in Fig. 26, Plate VII. When, however, such an aUoy is slowly cooled, or, after rapid initial cooling, is subsequently heated so as to allow of the attainment of equifibrium by the slow process of diffusion, the metal becomes entirely homogeneous, the dendritic cores disappear, and we have again the familiar aggregate of homo- TYPICAL ALLOY SYSTEMS 143 geneous crystals such £is are seen in all pure metals. It should be noted that the mechanical work and subsequent annealing . which occur in the working of brass as usually carried out are also sufficient to render the metal homogeneous in this sense, with the result that wrought brasses containing less than 30 per cent, of zinc do not as a rule show any trace of dendritic structure ; their structure in fact is very similar to that of pure copper when treated in the same way, i.e., it consists of an aggregate of crystals whose size depends upon the manner of heating and working, but these crystals show the strongly marked characteristics of a freely " twinned " structure. An example is shown in Fig. 55, Plate XI. When we pass beyond the range of alloys terminated by the point b of the diagram and examine alloys containing from 30 to 37 per cent, of zinc, we meet with a different state of affairs, as we have now entered a region in which the /3 phase plays a part. In these alloys crystallisation commences, on cooUng past the hne AB, in much the same manner as in the first class described above, but this only continues for a short time, until in fact the temperature of the line bB (about 880° C.) is reached. The solid which is formed above this temperature is still the a body, but at this temperature the remaining liquid freezes in the form of the second phase, j9. While the a body, consisting as it does chiefly of copper in the free state in which a relatively small proportion of zinc is held in solution, retains most of the properties of copper, and is, therefore, a compara- tively soft and very ductile body, which can be readily worked in the cold, the phase is a much harder, stronger, but also much less ductile body, and its presence in the alloy at once makes itself felt by a decided increase in the strength and hard- ness, but also in an equivalent decrease in ductility, which soon makes cold working impossible. For this reason the subsequent cooling-process of this particular group of aUoys is of special interest and importance. Owing to the shape of the hne be, these alloys on cooHng pass from the region in which both o and y8 can exist side by side into a region where a only is stable, and where, therefore, jS either does not exist at all or only occurs as a meta-stable form, i.e., a form which has been 144 STUDY OF PHYSICAL METALLURGY preserved in existence by cooling the alloy at a rate too rapid to allow of the completion of a change which is necessary for the attainment of equilibrium. In the present case the /8 body tends to change into the a body when the alloy passes through the temperature represented by that point on the hne be, which concerns that particular alloy. If the alloy is slowly cooled this change takes place and the relatively hard /3 body disappears, and the alloy is reduced to an homogeneous aggregate of a crystals. If, however, the coohng is too rapid to allow this change to occur or to be completed, the alloy retains a certain amount of the /3 body in its constitution and is thereby rendered harder and more brittle. The anneahng and subsequent cooling of such alloys is, therefore, capable of modifying their properties in a very decided manner. When the proportion of zinc indicated by the point 6 is exceeded, i.e., in alloys containing more than 37 per cent, of zinc, the /3 phase remains stable at all temperatures down to 470° C, and the slowly-cooled alloys possess a duplex structure which is well seen in ordinary Muntz metal (approximately 40 per cent, of zinc), as in Fig. 56, Plate XI. Rapid coohng from a suitably high temperature also affects the structure and constitution of these alloys, for rapid cooling from temperatures lying above the hne jBc^ wlQ result in the suppression of the /8 to a change, and the alloys may even be obtained as a homo- geneous mass of the /3 phase. Here, again, heat-treatment is capable of profoundly modifying the structure and constitution, and consequently the properties, of the alloys ('). The region of true stabihty of the fi body is conjfined, accord- ing to Carpenter {^), to temperatures above 470° C. At this temperature a hne of small heat-evolutions occurs in the thermal curves of these alloys, and Carpenter considers that this indicates the decomposition of the phase into a and y. In accordance with this view the hnes of the diagram starting from the points B and C respectively are drawn to meet at the point c^ on the hne djd2. The decomposition of the ^ phase which occurs along this hne is not readily seen under the microscope, nor can the decomposition products of the pure jS phase be caused to coalesce into lumps or laminae visible under moderate magnifi- PLATE XI. Fi( Fig. 56. [To face p. \H. TYPICAL ALLOY SYSTEMS 145 cation, although such coalescence appears to occur in the presence of certain third elements. How far this supposed decomposition of the yS body affects the mechanical or other physical properties of the alloys is also not yet ascertained, although it has been suggested that it plays a part in the spon- taneous cracking of certain brass articles. Further evidence must, however, be awaited before much weight can be attached to this whole matter. In the region of the diagram enclosed by the lines BC, Cc^, CjB, the alloys consist entirely of the /8 phase, and it is interest- ing to note that it is this phase which, while comparatively hard and brittle in the cold, lends itself to hot roUing and forging, while the a phase, which is soft and ductile at the ordinary temperature, is generally regarded as being too weak and friable when hot to withstand hot working.' At the ordinary temperatures the alloys lying between the points d^ and Ci appear to consist of a mixture of the a with the /3 body, while those between the point Ci and d^ appear to consist of a mixture of the )8 and 7 phases. This latter phase is exceedingly hard and brittle and its presence in the alloys renders them useless for any purpose where strength and toughness are required. This is a tj^ical example of a law very widely apphcable to aUoys, viz., that those phases of a binary system which contain the two elements in anything like equal proportions are hard and brittle, only the alloys near the ends of a series being as a rule sufficiently strong and ductile to be of practical utihty. We have already seen that the /3 phase is harder and more brittle than the a, so much so that the best brasses, in which strength and ductility are of importance, are generally made with a zinc-content of approximately 30 per cent., this being the cheapest alloy which does not contain the /8 phase. For many purposes the 60/40 brass (Muntz metal) is still sufficiently ductile, but with the appearance of the 7 phase the strength of the brass diminishes rapidly and its ductility is still further reduced, so that brasses containing such large proportions of ' The author has recently seen a 70/30 brass hot-rolled quite success- fuUy, and it therefore seems that the hot-shortness of the brasses may not be an inherent property of the alloys. P.M. L 146 STUDY OF PHYSICAL METALLURGY zinc are of little importance for engineering purposes. It is perhaps worthy of notice at this point that the zinc-copper alloys are particularly noticeable on account of the wide range of composition over which they show useful mechanical pro- perties. In the tin-copper and aluminium-copper series a content of 15 per cent, of the added metal is sufficient to destroy the ductility of the alloys. The fact that zinc is itself a cheap element, and the further unique fact that 30 per cent, of it can be added to copper before a second phase makes its appearance, act together to bring about the wide utihty and numerous applications of the brasses as compared with all other copper alloys. The shape of the constitutional diagram has thus a most direct practical bearing on the scope and utihty of the system of alloys which it represents. The micro-structure of the zinc-copper alloys has already been referred to in describing their constitution, and the structure of the o alloys has been illustrated in the cast " cored " condition and in the worked and annealed condition in Figs. 21 and 55, Plates VI. and XI. The duplex structure of the alloys corresponding to the fields of the diagram in which the a and yS phases appear together is shown in Fig. 56, Plate XI., which is typical of Muntz metal, zinc 40 per cent., copper 60 per cent. The /8 shown in this micrograph is, of course, " apparent /3," and, according to Carpenter's view, would consist of an almost ultra-microscopic mixture of a and 7. We now pass on to the consideration of the constitution and structure of another important group of the alloys of copper, viz., those with tin. The constitutional diagram of this system has been very carefully worked out by Heycock and Neville (^o) in a classical research, but their diagram was admittedly incomplete, and it has to a certain extent been modified by the subsequent work of Shepherd and Blough (^^), and of GioUitti (^^). Even now it is doubtful whether the intricacies of this complex system have been completely worked out. The diagram as given in Fig. 57 also embodies modifica- tions suggested by Hoyt (^^), but our discussion of the alloys must again be confined to those near the copper end of the series, which are the most important for all practical purposes. TYPICAL ALLOY SYSTEMS 147 These may be taken as having a tin-content of less than 20 per cent. In these alloys, according to the diagram, we shall only meet with three constituents, viz., those denoted by o, |3, and B. The very complex changes undergone by the alloys ranging from 20 to 40 per cent, of tin do not extend to the range of Cu 80 30 4-0 50 60 Composition . Fig. 57. — Constitutional Diagram of the Tin-Copper Alloys. ioo%s„ on . bronzes ordinarily employed for engineering purposes. The transformation which occurs along the line PQ is, however, of great imporance in regard to aU bronzes containing more than 12 per cent, of tin. The micro-structure of tin-copper alloys containing less than 10 per cent, of tin is very simple and exactly similar to that of the corresponding zinc-copper alloys, being that of a simple l2 148 STUDY OF PHYSICAL METALLURGY solid solution of a tin-copper compound in copper. When rapidly solidified these alloys exhibit the dendritic structure characteristic of solid solutions which have not attained equihbrium, but slower freezing, or subsequent annealing at a high temperature, obliterates the difEerences of composition which exist between the cores and the peripheral regions of the crystals, and results in the formation of the uniform crystal- line aggregate with which we are already familiar. This change, however, very materially affects the strength and other properties of the metal — the rapidly-cooled material being very much superior from the mechanical point of view to that obtained by slow cooMng. Hot rolling, of course, further improves the strength of the material. Cold working — ^in these as in other metals — hardens the material, increasing its ultimate strength, but seriously reducing its ductiUty. In the case of bronzes, with the exception of those containing very little tin, the hardening effect of rolhng in the cold is much more marked and rapid than is the case with copper-zinc alloys. Bronzes somewhat richer in tin cannot be rolled or worked to any considerable extent in the cold. The analogy between the a body of the tin-copper series and the a body of the zinc-copper aUoys is in many respects very close ; under the influence of mechanical work and annealing hoih these bodies assume the rectilinear forms characteristic of twinned crystals, such as that illustrated in Fig. 55, Plate XL The a body of the aluminium-copper series behaves in a similar manner. Alloys lying in the short range between the points E. and S in the diagram correspond to the second group of the zinc- copper alloys to this extent, that while their soHdification begins by the deposition of the a body, it is completed, along the line RS, by deposition of /3. Only those near the point R, however, lying vertically above the curve RT, are completely trans- formed into a, even by very slow coohng. Those lying to the right of the composition corresponding to the point T undergo a transformation whereby the ^ originally present is trans- formed finally into a mixture of a and S phases. According to the diagrams of Heycock and Neville, and of Shepherd and PLATE XII. I-'IG. 59. J' W!^f Jr.: Mm • ' « ^%,A ♦ it- ? Fig. 60. [To face p. 149 TYPICAL ALLOY SYSTEMS 149 Blough, this is a direct transformation at the temperature of the Hne PR according to the equation ^ = a + b. Hoyt, however, has found evidence for the existence of an intermediate transformation along the line TU, where /8 is transformed into an intermediate phase, 7. In either case the finally resulting phase, S, is much harder and more brittle than /8, so that the mechanical properties of the alloys are improved by quenching in such a way as to retain the ^ phase. The various conditions of micro-structure found in an alloy of this group are illustrated in Figs. 58, 59 and 60, Plate XII., which represent the structure of an alloy containing 18 per cent, of tin in various conditions, as slowly cooled in the ordinary way. Fig. 58, Fig. 59, as annealed at 750° C. for thirty minutes and quenched (thus being in the a -{- ^ state), and Fig. 60 after annealing at 450° C. (just below the hne PQ) in order to complete the transformation, thus bringing the aUoy into the a + B condition. The alloys richer in tin, particularly those covering the intricate part of the diagram relating to alloys containing between 20 and 40 per cent, of tin, although of great theoretical interest, cannot be dealt with here, as a discussion of their constitution and micro-structure would occupy too much space. Such a discussion is further rendered difl&cult by the fact that the exact details of the constitution are not yet defi- nitely settled. Only one point of particular interest must be mentioned ; this relates to the fact that owing to the peculiar form of the constitutional diagram the alloys containing from 33 to 37 per cent, of copper exhibit a very unusual behaviour, which consists in first undergoing sohdification in the usual manner, but afterwards — on cooling to a lower temperature — again becoming hquid, and only then undergoing final sohdifica- tion at a still lower temperature. Such behaviour can, of course, only be explained by the occurrence of a chemical reaction as the result of which the crystals of a soHd solution are transformed into a compound — or a solution of a compound in an excess of one of its constituents — ^which is more fusible than the original soHd solution, with the result that the forma- 150 STUDY OF PHYSICAL METALLURGY tion of the compound on cooMng through a critical temperature is accompanied by the fusion of the alloy which had just previously soUdified. In the preceding sections the constitution and structure of typical members of the two binary series — zinc-copper and tin-copper — ^have been very briefly discussed. A very large number of important alloys, however, although closely related to these two series, are essentially ternary alloys in which both Fig. 61. — Diagram of the Constitution of the ternary alloys of Tin, Zinc and Copper, rich in Copper. zinc and tin are present. In view of the complexity of both the two binary systems, it might be anticipated that the ternary equilibrium would be excessively complicated. Actually the systematic investigation of these ternary alloys has only been undertaken quite recently by Hoyt (^^), but although that author has obtained interesting and valuable data, it cannot be considered that these alloys have been fully explored as yet. As far as it goes, Hoyt's ternary diagram is reproduced in Fig. 61. The diagram relates only to the corner of the ternary triangle representing the alloys rich in copper ; the figure shows TYPICAL ALLOY SYSTEMS 151 the binary diagrams of the zinc-copper and the tin-copper systems folded down to either side of the triangle, while the maimer in which the various regions of the binary diagrams are connected in the ternary diagram are shown by the lines drawn on the ternary triangle. From these it will be seen that there is a large area, marked a in the diagram, in which the alloys solidify as a simple a soMd solution, whose structure and properties resemble those of the a constituent of either of the two binary systems, the tin and zinc simply replacing one another, although this replacement is not without influence on the mechanical properties. This field is marked Cu 8, 8 in Hoyt's diagram. Next to this lies the region 8, 8, 7, 7, in which the alloys contain two phases, one the soHd solution, and the second the y phase, which is similar to the corresponding phase of the zinc-copper system. Within this region, however, the structures of the alloys vary according to their mode of solidi- fication and subsequent transformations ; in those near the fine 7, 7, and particularly near the zinc-copper side of the diagram, a duplex structure, consisting of the a and y phases, is found, and this is regarded as the product of decomposition of the ^ phase of the zinc-copper system, this decomposition being rendered more evident by the presence of tin. The ternary system just described is by no means the only one which is of interest and importance in connection with the alloys of copper, but the matter cannot be pursued here. It should be pointed out, however, that the influence of a third metal or metalloid on a system of alloys cannot be fully under- stood until the constitution of the ternary system formed by that third element with the primary system has been studied and elucidated. As an approximation, it is sometimes sufiicient to determine " the effects of small additions " of the third element to the binary system, but this is in some cases an un- satisfactory proceeding, chiefly because traces of new consti- tuents which may appear are Ukely to be overlooked or their true significance misunderstood if only the alloys lying very close to the true binary system are studied. It is to be hoped, therefore, that the difl&cult and tedious exploration of ternary systems will be steadily pursued. 162 STUDY OF PHYSICAL METALLURGY Mention may be made here of certain alloys which belong really to the class of ternaries just referred to, but whose true constitution is not yet fuUy worked out. Their practical importance, however, demands some reference to them. The addition of phosphorus to copper (") and its alloys is well known to improve the mechanical qualities of the metal, particularly if the quantity is carefully adjusted. This is not in reality due to the formation of any new alloy, since the phosphorus serves simply to remove the oxides present in the molten metal, the phosphoric oxide passing into the slag or scum on the surface. When melted in the ordinary way, copper and its alloys are always more or less seriously contaminated with oxides. In the case of copper the presence of the oxide is easily recognised under the microscope (^), and its influence on the properties of the metal is well known. For electrical purposes the presence of oxide is particularly un- desirable, but the use of phosphorus for its removal is not satisfactory, owing to the fact that any slight excess of the metalloid remaining in the copper would seriously affect its conductivity and its mechanical properties. For that reason a more satisfactory deoxidisor for copper has been sought, and it is believed to have been found in the substance known as boron sub-oxide (^•). In the zinc-copper and the tin-copper alloys oxide is also habitually present. In the tin-copper series it can be readily detected both analytically and under the microscope, as Heyn and Bauer (") have shown very clearly. In the zinc-copper series its detection by either method is very difficult. The use of phosphorus for the removal of oxide is principally practised in the case of the tin-copper alloys. When the quantity of added phosphorus — which is usually introduced in the form of either a phosphorus-copper or a phosphorus-tin alloy — is so adjusted that the phosphorus is entirely or almost entirely oxidised, the resulting material is not reaUy a ternary alloy, but merely a purified tin-copper alloy, which is, therefore, somewhat misdescribed by the term " phosphor bronze " (^^). When larger quantities of phos- phorus are introduced, the metalloid is found in the resulting alloy in the form of a definite compound, a phosphide of copper. TYPICAL ALLOY SYSTEMS 153 which is hard and brittle. The presence of this substance in the alloys, where it is readily distinguished under the micro- scope, renders them much harder and stiffer than the pure tin- copper alloys, but at the same time renders them comparatively brittle. As a consequence, these alloys — containing up to 1-5 per cent, of phosphorus — are employed principally for bear- ing-metals, where the copper-phosphide plays much the same part which the tin-antimony compound plays in certain white- metal bearings. Metallic manganese may also be employed as a deoxidising agent in copper aUoys, and although it is a less active reducing agent than phosphorus, it has the considerable advantage that if used in excess it does not render the alloys hard and brittle. Manganese does not form any definite compounds with copper, the copper-manganese constitutional diagram being that of an unbroken series of solid solutions, although this particular system exhibits a minimum in the liquidus which somewhat recalls a eutectiferous system. The presence of small quantities of manganese in alloys of copper with zinc or tin, therefore, does not as a rule give rise to the formation of any new constituents ; the manganese enters into solid solution, and merely adds its quota to the hardening effect of the dissolved metals. A more powerful deoxidising agent than manganese is furnished by aluminium, but this differs from manganese in two vitally important respects. In the first place, the oxida- tion product of aluminium is a particularly refractory substance — alumina — which has a strong tendency to remain in the molten metal in suspension as fine particles. These, of course, tend to lessen the strength and toughness of the alloy. Beyond this, aluminium exerts a very powerful hardening effect on copper, owing to the fact that it enters into chemical combination with that metal. Both these reasons militate against the use of aluminium as a deoxidising agent for alloys of copper, although it is employed for that purpose in the metallurgy of steel. Returning to the binary aUoys of copper, those with aluminium remain to be considered. These are of interest in several directions ; their constitutional diagram presents a number of peculiar features, some of which have not yet been 154 STUDY OF PHYSICAL METALLURGY fully elucidated, although the subject has received much study at the hands of Carpenter and Edwards ("), Curry C^), and others (2*). A constitutional diagram based on their results is reproduced in Fig. 62, but this must still be regarded as somewhat tentative. From the point of view of practical apphcation the alloys at the copper end of the series, con- taining up to 10 per cent, of aluminium, are of considerable interest, both on account of their valuable mechanical properties and also on account of their marked powers of resisting Certain difficulties which have been experienced corrosion. cc*cc in P 6 + Eutect. 07, Ca A/. 10 60 50 40 Composition. Fig. 62. — Constitutional Diagram (tentative) of the Aluminium-Copper Alloys. in their use and working are likely to be overcome when their advantages are fully realised. At the other end of the series is a group of light aUoys consisting principally of aluminium, but containing up to 4 or 5 per cent, of copper, which are of interest, not so much for their own sake, as on account of an important alloy which is derived from them. This alloy, known by the trade name of " Duralumin," will be referred to. again in connection with the ternary system manganese- aluminium-copper. At the other end of the series we will again confine our attention to the alloys containing less than 10 per cent, of aluminium. Up to an aluminium content of slightly more than 7 per cent. TYPICAL ALLOY SYSTEMS 165 the alloys are simple solid solutions consisting entirely of the a phase, which is very similar in character and appearance to the a phase of the zinc-copper or the tin-copper system. Like the two latter the a body of the aluminium-copper series readily undergoes twinning when subjected to plastic deforma- tion followed by annealing. In its mechanical properties, however, the aluminium-copper a phase differs widely from that of the other copper alloys. This series of solid solutions, however, exhibits one peculiarity ; it will be seen from the diagram that, unlike the two other copper systems discussed above, there is no long freezing -range for the a phase ; indeed, the solidus and liquidus lie so close together that it is difficult to distinguish them from one another. It is a direct conse- quence of this circumstance that the portions of the a body which crystallise first do not differ very much in composition from the portions which solidify last, and as a result the cast alloys do not exhibit that well-marked structure of dendritic cores which is so frequently met with in solid solutions. This range of pure a alloys corresponds to the first branch of the liquidus curve, AB of the diagram. At the point B a definite minimum in the liquidus is reached and in the majority of cases such a minimum is associated with the formation of a definite eutectic alloy. In the aluminium-copper system, however, this is not the case, and the minimum seems to be analogous to that which occurs in the manganese-copper series, being simply formed in an unbroken series of solid solutions. It would, however, probably be more correct to draw the curve continuously through this minimum rather than to show a definite cusp. From this minimum the liquidus curve rises to a maximum at D, which corresponds to the definite compound CugAl. Corresponding to this new branch of the liquidus there is a striking change in the micro-structure of the aUoys, which now becomes duplex. Since the field of the a phase is bounded by the sloping line BO, it will be seen that an alloy whose composition lies just to the left of the point P (i.e., containing about 12-5 per cent, of aluminium) should still be homogeneous when in a state of equilibrium. The change from /3 to a in TYPICAL ALLOY SYSTEMS 157 material which can be readily broken to powder, and is consequently very useful in the preparation of alloys in the foundry, as it can readily be weighed out exactly to the desired quantity. If a small ingot of this alloy be allowed to set superficially, and if the skin is then broken and the residue of liquid metal is poured off, a mass of beautiful interlacing crystalline needles is left behind. These, although they consist of equal proportions of copper and aluminium, are of a brilliant white colour and extremely brittle. The ternary alloys of aluminium and copper with other metals have received a good deal of attention. Edwards and Andrews {^) have studied the equilibria of the aluminium-tin- copper system, while the author and Lantsberry C) have investigated the ternary system containing manganese. The relationships found are, however, too complex to be described here, and reference for particulars must be made to the original memoirs. One or two points of interest may, however, be mentioned. The addition of manganese to the aluminium- copper system, so far as the alloys near the copper end of the series are concerned, appears to result in a sHght increase of strength without any reduction of ductihty, provided that the manganese addition is small or that the alloy in question lies well below an aluminium content of 10 per cent. Apart from this effect, however, these alloys containing manganese appear to possess a still greater resistance to corrosion than the binary aluminium copper alloys, and they are also pecuharly resistant to abrasion, so much so that they excel tool-steel in that respect. This combination of properties offers some advantages in certain branches of instrument construction, particularly where the use of non-magnetic material is essential. When severelj' cold- worked these alloys attain a very high degree of strength and hardness, so that it has been possible to make cutting tools of these materials which are quite capable of cutting wood and even stone. In the Hght alloys containing less than 4 per cent, of copper the addition of manganese was found to be an advantage, but the quantity that can be added is limited by the fact that aluminium and manganese form the definite compound AlgMn, 168 STUDY OF PHYSICAL METALLURGY which is a hard and brittle body, and one, moreover, which — if present in any large proportion — tends to cause the alloy to fall to pieces spontaneously. Additions of manganese are therefore confined to 2 per cent. The aUoy of aluminium with copper 3 per cent, and manganese 1 per cent, was at the time when it was investigated the best light alloy known, and by the addition of 0-5 per cent, of magnesium it is converted into the alloy already mentioned, known as " Duralumin." ^ The presence of this small percentage of magnesium appears to confer on this material, and indeed upon any alloy consisting largely of aluminium, a peculiar property of hardening slowly in the course of several days after it has received an appropriate heat treatment (**). The exact nature of this hardening process is not understood, but it results in very nearly doubling the tensile strength of the aUoy without very seriously reducing its ductility. This remarkable behaviour constitutes one of the most striking unsolved problems in present-day Physical Metallurgy. References. (1) Rosenhain and Tucker. Phil. Trans. Roy. Soc, 1908, 209a, p. 89. (2) Mazotto. Internat. Zeitsohr. f. Metallographie, I., 5, p. 289. (3) Degons. Zeitschr. Anorg. Chem., 1909, 63, p. 207. (4) Hannover. Metall. and Ciem. Engineering, No. 9, September, 1912, X., p. 609. (5) Charpy. Metallographist, 1899, II., p. 9. Bull. Soc. d'Encouragement, June, 1898. (6) Rosenhain and Archbutt. PM. Trans. Roy. Soc, 211a, 1911, p. 315. Journ. Inst. Metals, No. 2, 1911, VI., p. 236. (7) Shepherd. Journ. Phys. Chem., 1904, No. 6, VIII., p. 421. (8) Carpenter and Edwards. Journ. Inst. Metals, No. 1, 1911, V., p. 127. Carpenter. Journ. Inst. Metals, No. 1, 1912, VII., p. 90; No. 2, 1912, VIII., p. 51. (9) Bengough and Hudson. Journ. Soc. Chem. Ind., 1908, XVII., pp. 43 and 660. Journ. Inst. Metals, No. 2, 1910, IV., p. 92. Stead and Steadman. Journ. Inst. Metals, No. 1, 1914, XI. ' A group or range of alloys is known under the trade name " Dura- lumin " ; the composition stated above is typical of one of these. TYPICAL ALLOY SYSTEMS 159 (10) Heycook and Neville. Phil. Trans. Roy. Soc, 1903, Vol. 202a, p. 1. (11) Shepherd and Blough. Journ. Phys. Chem., X., p. 630. (12) Giolitti and Tavanti. Gazzetta Chemica Itahana, 1908, 38, 2, p. 209. (13) Hoyt. Journ. Inst. Metals, No. 2, 1913, X., p. 235. (14) Heyn and Bauer. Zeitaohr. Anorg. Chem., 1907, 52, p. 129. Huntington and Desch. Trans. Faraday Soc, 1908, IV., p. 51. (15) Heyn and Bauer. Zeitschr. Anorg. Chem., 1904, Vol. XXXIX., p. 11. (16) Weintraub. Amer. Electrochem. Soc. ; Engineering, August 11th, 1911, XCII.,p. 203. (17) Heyn and Bauer. Zeitschr. Anorg. Chem., 1905, XLV., p. 52. (18) Hudson and Law. Journ. Inst. Metals, No. 1, 1910, III., p. 161. (19) Carpenter and Edwards. Proo. Inst. Mech. Engineers, 1907. (20) Curry. Journ. Phys. Chem., 1907, 11, 425. (21) Gwyer. Zeitschr. Anorg. Chem., 1908, LVII., 113. (22) Edwards aad Andrew. Journ. Inst. Metals, No. 2, 1909, II., p. 29. (23) Eosenhain and Lantsberry. Proc. Inst. Mech. Engineers, 1910. (24) Wilm. MetaUurgie, VIII., p. 225. CHAPTER VIII THE IRON-OARBON SYSTEM On account both of their intrinsic importance and of their complexity of constitution and structure, the iron-carbon alloys require treatment in a special chapter. Such separate treat- ment must not, however, be taken to indicate that iron and steel differ in any fundamental way from other aUoy systems. Such an idea appears to have been prevalent in the minds of some of the earher workers (^) in the subject, and even now there is an unfortunate tendency to discriminate too strictly between " ferrous " and " non-ferrous " metallurgy or metallo- graphy. It is, therefore, well to emphasise that iron and steel are alloys subject to precisely the same laws and to be studied in exactly the same way as other aUoy systems. A good deal of the difficulty and confusion which has at times clouded the view of the constitution of the iron-carbon system has, indeed, arisen from the circumstancethatthis subject has often attracted workers who came to it without previous acquaintance with the more general methods of studying alloys, and often even without an adequate knowledge of the fundamental principles of physical chemistry. On the other hand, the iron-carbon system has probably been subjected to more minute study than any other series of aUoys ; now, as each binary system is more closely studied, it is found that constitutional diagrams which on first investigation appeared to be simple have turned out to be more and more complex. The complexity of the iron-carbon system as we know it to-day may thus simply be the result of this minute study, so that even in this respect the system does not differ very widely from other alloy series. An approximate constitutional diagram of the iron-carbon system is given in Fig. 64. This is not, strictly speaking, an equiHbrium diagram, since it represents what is probably merely a very persistent meta-stable system. The diagram of THE IRON-CARBON SYSTEM 161 the completely stable system is not, however, thoroughly known and is of minor importance, since it deals with conditions never met with either in practice or in the laboratory. So far as it is known it is given in Pig. 65. The constitutional diagram of Fig. 64, while, strictly speaking, representing only the author's views on the constitution of iron ISOO'C 1400 /300 1200 J-> JIOO •a I. ■S woo u S 300 600 ■ 700 eoo ■ soo \ \ ^-\ C ■ ^■^--^ D ^^■\^£ \/ 1 < / \ F H \ / J Fe Composition Fig. 6i.— Coastitutional Diagram of the Iron-Carbon System. sr-c c and steel, is very similar to the diagram which is accepted by the great majority of physical metallurgists, and is practically identical with that pubUshed by the International Testing Association in connection with the Report of the International Committee on the Nomenclature of the Microscopic Constituents of Iron and Steel in 1912 (2). Some minor points on which differences of opinion exist wiU be mentioned as the details of P.M. M 162 STUDY OF PHYSICAL METALLURGY mo J200. Liquitf Austenlte ^' ^ < Graphite ■^Litfuid kmo AustenitS •*-£uteclic Graphite *£utectic \I. the diagram are discussed. The diagram as a whole — broadly speaking — may be regarded as accepted by all but one small and isolated group of metallurgists, who have not, however, put forward an alternative diagram for discussion. Beyond the present warning that such a group exists, therefore, their pecuhar views will not receive further consideration here. The constitutional diagram of the iron-carbon alloys, as shown in Fig. 64, is pecuhar, in the first place, on account of the fact that it does not extend beyond a carbon concentration of 6 per cent. A few alloys of rather higher carbon content have been prepared experimentally by the aid of the electric furnace, and attempts have been made to carry the diagram somewhat further to the right (*), but in the intro- ductory treatment of the subject which is contem- plated here that region of the diagram will not be con- sidered. The reason why the diagram is not continued to the full extent on the carbon side lies, of course, in the fact that the system in question is one of alloys between a metal and a metalloid, the latter — carbon — being, moreover, a substance which cannot be melted in any ordinary way. We find similar limitations in the alloy systems of other metalloids, such as those of copper and arsenic, where the volatility of the arsenic limits the diagram to the vicinity of the copper end. In reaUty the " iron-carbon " system is the " iron-cementite " system, the name " cementite " being given to the definite iron-carbon compound having the formula FeoC, and containing 6-6 per cent, of carbon. The liquidus of the iron-carbon diagram is given by the lines ABC in the diagram (Fig. 64). The branch AB represents the deposition from the molten alloys of a solid solution of carbon — or more probably of iron carbide — ^in iron. The ■ eoo I 2 3 4 S 6 no Composition Fia. 65. — Constitutional Diagram (tentative) of the Iron-Graphite System. THE IRON-CARBON SYSTEM 163 branch BC corresponds to the solidification of crystals of iron carbide (cementite), but the formation of such crystals from fusion is frequently accompanied or very shortly followed by their decomposition into iron and graphite. This change is in reality a transition from the meta-stable system iron-iron carbide to the stable system iron-graphite. It is of great importance in connection with the structure and constitution of cast iron, and will be referred to in that connection. The point B of the diagram is obviously a eutectic point similar to those found in the lead-tin and aluminium-zinc systems ; in the present case the eutectic is that of the sohd solution mentioned above — which we may refer to as the 7-iron solid solution — with cementite. Here also there is reason to believe that there is an alternative mode of solidifica- tion according to the stable iron-graphite system, a eutectic of that kind being indicated by the lines of the diagram of the stable system in Pig. 65. The solidus of the iron-cementite system is given by the lines AD, DB, the point D indicating the limit of solid solubility of cementite in iron at a temperature of about 1,135° C. The system thus belongs to the general type, in which there is a considerable range of solid solubility between the two com- ponents, with a central eutectiferous range. The eutectic temperature in this case lies at or near 1,135° C. The position of the solidus between A and B cannot be ascertained from the data derived from thermal curves, and its position as given in the diagram is based on the work of Gutowsky (*), who deter- mined the position of the solidus by the method of quenching small specimens of steel of various composition and thus ascertaining by the microscope at what temperatures the first traces of the presence of liquid could be discovered. This work is the best as yet available on this line, and it is pro- bably sufficiently accurate to warrant the general shape and position of the line AB as drawn in the figure, but confirmatory work of the most exact nature is required in order finally to fix this important line. Like all binary diagrams of this type, that of the iron-carbon system must possess a line limiting the region of the stable M 2 164 STUDY OF PHYSICAL METALLURGY existence of the 7-iron solid solution which is formed by the final solidification of the alloys along the line AD. This line must start from D, the point at which, under conditions approximating to equilibrium, the eutectic first makes its appearance, and run downward to the base of the diagram. If the limit of solid solubility of cementite in 7-iron were the same at all temperatures, this bounding line would be a simple vertical. Actually, however, the amount of cementite which 7-iron can retain in solid solution decreases with falling tem- perature, and consequently the line DI slopes backwards towards the iron end of the diagram. This is in itself an unusual feature in alloys, as the bounding line of the region of the a phase in both the zinc-copper and the tin-copper systems slopes away from the copper end of the diagram. The line DI also exhibits another pecuharity in that it does not extend to the base of the diagram. This arises from the circumstance that the 7-iron solid solution does not exist at the ordinary tem- perature in a state approaching equilibrium, but undergoes decomposition at temperatures lying on or above the line HIJ (near 700° C). The line DI, therefore, ends where it cuts the line HIJ, that being the limit of existence of the 7-iron phase. In one sense the Hne DI may be regarded as strictly analogous to one branch of a " liquidus " curve in an ordinary eutectiferous alloy system. Such a branch of a liquidus curve indicates the temperatures at which the liquid solution of two components begins to deposit crystals of one of them, i.e., the limit of solubility of the one component in the liquid solution of the two. Similarly the Hne DI indicates the temperatures at which the solid solution, 7, begins to deposit crystals of cementite, which is one of the components of the solid solution. Now if we pursue this analogy a little further we shall expect to find that, just as a downward sloping branch of a liquidus generally meets another downward sloping branch at a eutectic point, so the line DI, representing the temperatures at which the 7 solid solution deposits cementite, should meet, at a point corre- sponding to a eutectic point, a branch or curve corresponding to the deposition from the sohd solution of the other component. THE IRON-CARBON SYSTEM 165 viz., iron. This branch is found in the diagram in the shape of the two lines EG and GI. The fact that there are two lines instead of a single one arises from precisely the same kind of cause as that which produces a break in the liquidus curve of the aluminium zinc alloys at the point C of that diagram (Fig. 48), or at the point B in the zinc-copper diagram (Fig. 54) — it is simply that the phase deposited from the solution along the branch before the small break is different from that which is deposited along that part of the curve lying beyond the break. At all events, we certainly have at I a point corre- sponding, for the decomposition of the y solid solution, to the eutectic point of a liquid solution, while the hnes EGI and ID, which meet at this point, I, correspond to the two branches of a liquidus curve in the case 6i an alloy solidifying from fusion. The point I is consequently called the " eutectoid " point, and the body formed along the line HIJ is known as a " eutectoid." The region in which the y solid solution is stable is thus bounded by the lines AD, DI and IGE. We may now follow somewhat more closely the manner in which this solid solution undergoes decomposition. As we have already seen, along the branch EGI we may expect to find that the solid solution deposits crystals of iron, while along ID it deposits crystals of cementite or iron carbide. The question naturally arises, why should the 7-iron solid solution deposit its burden of dissolved cementite, or its excess of iron at certain temperatures, instead of carrying them permanently in solution as the a phase in most of the copper alloys carries the zinc or tin or aluminium with which it is associated ? Here also the analogy vnth the Mquid solution of two metals, i.e., with a molten alloy, will aid us in understanding what takes place. The deposition of a solid metal from a liquid solution is inti- mately connected with the natural freezing-point of the metal in question. In the pure metal itself, solid crystals are deposited from the liquid (i.e., from the molten metal) at one definite temperature, which we call the freezing-point of the pure metal. Now that branch of the liquidus which relates to the deposition of crystals of that metal from the hquid alloy starts from the freezing-point of the pure metal and 166 STUDY OF PHYSICAL METALLURGY slopes downward to the eutectic point. The whole of such a branch of the liquidus may, therefore, be regarded as represent- ing the depression of the natural freezing-point of the pure metal by successively increasing additions of another metal. AVhat is really happening along such a liquidus curve is thus simply the freezing of the metal, modified, principally in regard to temperature, by the presence of the second element. If the analogy between the decomposition of the y solid solution and the freezing of a series of alloys is correct we should expect to find that the decomposition of the 7 solid solution is a phenomenon which we could trace as something analogous to a change of state in the pure metal, i.e., in iron itseK. And this is actually the case. The curve EGI starts from a point in pure iron itself — a point which does not, it is true, represent the passage of the material from the Hquid to the sohd state, but which indicates a profound change taking place in the iron, and accompanied by evolutions of heat on cooling and absorptions of heat on heating which are comparable with those which occur on freezing or melting. In the case of a change of state — i.e., a change from liquid to solid or from liquid to gas — ^we see a phenomenon which involves a profound change in the arrangement of the molecules of the substances. In the gas the molecules are free from one another and move about unrestrictedly in all directions ; in the liquid state the molecules are much less free to move, although stiU com- paratively free in contrast to the relatively rigid manner in which they are arranged in the crystalline system of a solid substance. We now come to the conception that, within the range of the sohd state, there may be, and indeed there are, various possible arrangements of molecules, or even of the constituent atoms of a molecule, correspondin to differences almost as great as those between two different " states." Such different conditions are known as " allotropic " conditions or modifications, and their existence is well known in such elements as carbon, where graphite, diamond and " amorphous" carbon are recognised as allotropic modifications of one another. So also in sulphur, selenium, oxygen, tin, antimony and many other elements, well-known allotropic varieties exist. The THE IRON-CARBON SYSTEM 167 occurrence of well-defined heat evolutions on cooling and heat absorptions on heating even in the purest obtainable iron are in themselves strong evidence that iron is capable of existing in at least three diverse or " allotropic " conditions. These are generally known as the 7, ^ and a forms of iron, and they have formed the subject of much controversy. This has, however, principally turned upon the question whether one of these modifications was in itself extremely hard, and whether its existence could adequately account for the hardness of hardened steel. That question does not concern us here, and it may fairly be said that the existence of two allotropic varieties of iron (the 7 and the a) is universally accepted. The j8 form is still the subject of discussion, but the question of its inclusion or otherwise only leads to a very slight modifica- tion of the constitutional diagram. The main point which is required for the comprehension of the constitutional diagram is that 7-iron in the pure state undergoes an allotropic trans- formation at a temperature close to 900° C, marked by the point E in the diagram of Fig. 64. In consequence of this transformation the iron, now in the a or the ^ state, can no longer hold any appreciable amount of carbide in soUd solution. Accordingly, along the line EGI, the solid solution deposits crystals of practically pure iron free, or almost free, from carbon. The deposit of such iron crystals, however, produces upon this solid solution precisely the same result as the deposition of crystals of lead, for example, produces upon a liquid alloy of lead and tin — the residual solution becomes impoverished in lead (or iron), and the temperature at which further deposition takes place thus corresponds to a point a little further down the Hne of deposition. Right down to the temperature of the line HIJ, therefore, there wiU always be a residue of 7-iron sohd solution, and the amount of this residue will be greater the nearer the aUoy under consideration lies to the point I. When the alloy finally reaches the temperature of the hne HIJ, the residual 7-iron solid solution undergoes transformation en masse, and the " eutectoid " body known as " pearKte " is formed, with a large evolution of heat. The precise meaning of the portion of the constitutional 168 STUDY OF PHYSICAL METALLURGY diagram of the iron-carbon system relating to the group of alloys generally known as "steel," which may be taken as lying to the left of the point D in the diagram, will be best understood by following a few typical aUoys through the cooling and heating 1000 Time Intervals. Pig. 66. — Tjrpical laverse-Rate Heating and Cooling Curves of Pure Iron (Burgess and Crowe). processes. At the same time the micro-structure of these alloys can be considered. We shall begin with pure iron. A typical set of heating and cooling curves of this material, reproduced from the work of Burgess {^), is shown in Fig. 66. It will be seen that there are two peaks on both the heating and the cooHng curves. On cooling we commence with the iron in the -y state, and nothing THE IRON-CARBON SYSTEM 169 occurs until the point E of the diagram is reached, at or near 900° C. At this point there is a marked evolution of heat on cooling, and the heating curve shows a corresponding absorp- tion of heat on heating, but at a slightly higher temperature. It has been shown that the exact interval between the tempera- tures at which this transformation occurs on heating and on cooling depends partly upon the rate of heating and cooling and on the maximum temperature which has been attained ; but even when these factors are allowed for there probably remains a certain " lag " between the two, representing the tendency of an existing state to persist for a short distance on either side of its true equilibrium temperature. This is a species of under-cooling or of super-heating which is of constant occurrence. Phenomena of under-cooling are frequently met with in the crystallisation of salt solutions, and in that case it has been shown that below the true freezing-point there is a limiting range within which the solution can be made to crystallise instantly if it is brought into contact with a solid crystal of the same kind to act as a nucleus. In the absence of a nucleus, however, spontaneous crystallisation only occurs after the limit referred to has been passed. Now in a solid solution no nuclei can be introduced, and consequently it is not perhaps surprising to find that transformation does not occur until some corresponding limiting temperature has been passed, this limit lying very sUghtly above the true equihbrium temperature on heating and considerably below it on cooling. There is, however, no rise of temperature during the trans- formation, such as occurs when an under-cooled liquid freezes. This may, however, be due to the fact that the total quantity of heat evolved is much smaller, and also that the circum- stances of the change, occurring as it does in a fairly rigid solid, do not permit of rates of transformation sufficient to bring about an actual rise of temperature. On further cooling below the point E, the pure iron exhibits another " critical point " or evolution of heat on cooling, with a corresponding absorption on heating at the point marked F. This is a much smaller point than the one at E, and attempts have been made to discredit its existence (*) in pure iron. 170 STUDY OF PHYSICAL METALLURGY According to the views of Benedicks ('), this thermal point does not represent an allotropic transformation of iron at all, but merely indicates the final disappearance of 7-iron molecules from the metal. This idea is based upon the assumption that the whole of the y-iron is not at once transformed into another modification at E, but that a certain number of 7-iron mole- cules persist in a state of solution in the other form of iron, and that the last of these undergo sudden transformation a,t the point F. This view is stated here in the desire to represent current opinion in as fair a manner as possible, but the author believes that aU the weight of evidence — into which it is not possible to enter here — goes to show that a /3 phase of iron, intermediate between the 7 phase, which exists above the point E, and the a phase, which exists below the point F, really exists. Not only is there a definite thermal point at F, but it persists through the range of steels between F and G, and there is a shght break in the Une EGI at the point G. There are also some small but very definite discontinuities in the physical properties of iron corresponding to the temperature of the point F, the most important of these being the sudden disappearance of ferro-magnetism at that temperature. Another is a sudden change in hardness or tenacity at that temperature, as shown by the author and Humfrey. Provisionally, at aU events, we may, therefore, think of iron passing from the 7 to the /3 state at or near 900° C. on cooKng, and from the /3 to the a at or near 750° C. Pure iron undergoes no further transformations on cooUng down to the ordinary temperature. The question now arises whether any changes of micro- structure can be correlated with the two allotropic trans- formations of iron. As a rule, changes which occur at high temperatures can be more or less completely inhibited by suffici- ently rapid cooUng, i.e., by quenching. That this is not the case for iron is at once obvious from the fact that while iron above 750° C. is practically non-magnetic, no known method of quenching pure iron renders it non-magnetic after cooling. As regards micro-structures, it must be remembered that we are dealing with a pure metal and that pure metals of most widely divergent properties stiU show practically identical THE IRON-CARBON SYSTEM 171 micro-structure. It is not surprising, therefore, to find that iron quenched from temperatures above 900° C. does not differ markedly in micro-structure from the same material slowly cooled. That the mechanical properties also remain practically unaffected is more remarkable, since it is not Ukely — even if we had not definite proof to that effect — that 7-iron should resemble a-iron in that respect. The inference which we are forced to accept is that in pure iron the aUotropic transformations cannot be prevented by rapid cooling or quenching. There is, however, good evidence to show that the properties of iron do undergo marked changes on passing through the critical temperatures, but this will be referred to, in connection with the effects of plastic strain on metals, in Chapter XI. The cooling and heating curves of really pure iron, free from carbon, as reproduced in Fig. 66, show no sign of any thermal change corresponding to the point H of the diagram of Fig. 64. This is, of course, in accordance with the indications of the diagram itself, since the reactions represented by the line JIH are due to the presence of the dissolved carbide in the 7-iron of steels ; as the pure iron end of the series is approached, these reactions diminish in intensity and disappear entirely in pure iron itself. We may next consider the case of an iron-carbon alloy or steel containing about 0*2 per cent, of carbon. In the y region this is a homogeneous soUd solution, and decomposition only begins, by the deposition of crystals of ^-iron, when the alloy cools down to the temperature of the sloping line EG. We have thus a first arrest or critical point, on cooling, corresponding to the commencement of the deposition of iron, in the ^ condition, from the solution, at a tem- perature which for a 0-2 per cent, carbon steel lies at 840° C. As the alloy cools further, the quantity of ;8-iron separated from the solid solution steadily increases until the line FG is crossed at a temperature near 750° C. Here all the free j8-iron present in the steel is transformed into the a form and, whereas above the fine EG we had a mixture of residual 7-iron sohd solution with crystals of ]3-iron, below the Une EG 172 STUDY OF PHYSICAL METALLURGY we have a mixture of residual y-iron solid solution with crystals of a-iron. With further cooling, the quantity of the a-iron crystals increases steadily until the temperature of the line HIJ is reached ; there the residual 7-iron solid solution is transformed into the eutectoid mixture of a-iron and cementite. In this steel we thus have three thermal critical points indicated by absorptions of heat on heating and evolutions of heat on cooling. These " arrest-points " have received special names derived from their relative position, counting from the ordinary temperature upwards. The point connected with the line HIJ is called the first arrest, or " A^," the one which occurs on the crossing of the line FG is the second arrest, or " Ag," and the arrest related to the sloping line EG is the " third arrest," or " A3." Now in these steels, as also in pure iron, as has already been indicated, the temperature at which these transformations take place on coohng are not identical with those at which they occur on heating, and the arrest-points on heating and cooling are, therefore, distinguished by writing " Ac " for the points on heating and " Ar " for the points on cooling. Thus " Acj " is the arrest observed on crossing the line EG on heating, while " Arg " is the arrest obtained when the line EG is crossed on cooling. These special terms are convenient in use, but intro- duce an additional complexity into a nomenclature already somewhat needlessly involved. In the case of the metallo- graphy of steel, the analogy with petrography has been followed all too closely, so that names — and often personal names- have been given to everything. The equally complex metallo- graphy of the tin-copper system is quite as clear without the use of such names, while in steel the use of personal names has undoubtedly introduced an element of personal acrimony into what should be purely scientific discussions. The micro-structure of a steel of 0-2 per cent, carbon content is found to be quite in accordance with the indications of the diagram and of the cooling curves. When slowly cooled in the ordinary way, such steel is found to possess a duplex structure, such as that shown in Fig. 67, Plate XIII., where the light areas represent the crystals of iron, while the dark area^s are the eutectoid body formed at the line HIJ. THE IRON-CARBON SYSTEM 173 The simple white constituent is readily seen to be identical with the crystals of pure iron, and is consequently known as " ferrite." In commercial steels, however, the " ferrite '' always holds some silicon and phosphorus in solid solution. The eutectoid is known as " pearlite " owing to the fact that it frequently exhibits a finely-laminated structure like that of mother-of-pearl and, under suitable illumination, displays an iridescent lustre of the same kind. That this body is typically laminated is illustrated in Fig. 68, Plate XIII., which represents such pearlite as seen under high magnification. In this respect the eutectoid " pearlite " bears out its analogy with the normal eutectics, which are also typically laminated. Both kinds of bodies, however, can easily be obtained in a non-laminated state, so that the laminated structure must not be regarded as a fundamental characteristic of eutectics or eutectoids. The fact that the laminated eutectoid body — pearhte — ^is really formed on crossing the line HIJ is readily proved by quenching experiments. Such steel quenched from a tem- perature lying between the lines HI and EG still exhibits a typical duplex structure, one of whose constituents is ferrite, but the area of ferrite is markedly less than in the slowly-cooled sample, particularly if the quenching temperature lies near the line EG. The second constituent, however, is seen to be fundamentally different in character from the pearlite of the slowly-cooled material. The second constituent still appears darker than the ferrite when etched with the usual reagents, such as picric acid, but under moderate magnifications it appears to be perfectly homogeneous, and only under the highest magnification can any structure be found in it — but this is not a lamination or granulation, but a tracery of inter- lacing needles. If the quenching temperature is taken stiU higher, i.e., above the fine EG, then the whole — or very nearly the whole — area of the specimen is found to be covered by the darker-etching constituent having the faint acicular structure. We thus recognise in this constituent the representative of the 7-iron solid solution which exists in these steels at the moment of quenching. That this constituent is not really the solid solution itself preserved unchanged by the act of quenching is 174 STUDY OF PHYSICAL METALLURGY shown by several considerations, which, however, are best discussed in relation to steels of higher carbon content. The results of quenching experiments on this grade of steel, however, are obviously sufficient to justify the existence of the various lines of the diagram, with the possible exception of the line FG. whose influence on the structure has not so far been estabhshed, Additional evidence in support of the correctness of the diagram is also afforded by the work of BaykofE (*), who succeeded in etching specimens of steel at high temperatures by the action of hydrochloric acid gas. This method, even more definitely than the earher efforts of Saniter (^») at hot-etching by means- of fused calcium chloride baths, reveals the fact that, at tem- peratures above the line EG, the steel has a structure of simple polyhedra, such as those which we are accustomed to meet in homogeneous solid solutions. Turning now to a steel of slightly higher carbon content, such as 0-6 per cent., we find a somewhat simpler tj^e of thermal curve. On cooling, the lines GI and HI are successively crossed, and the steel thus shows only two Ar points. The lowest of these is obviously identical with Ar^ as we saw it in the steel of 0-2 per cent, carbon ; the higher one was originally regarded as being the result of the merging of the Ar^ and Aj-j points of the lower carbon steels, and the point is consequently known as Ar23 or Ac23 respectively. This nomenclature is not, however, consistent with the constitutional diagram. If we accept the existence of the /3 phase in the region EFG, then the line GI represents a transformation — ^from the y to the a state direct — which is not present in either Ar^ or Arj, while, if we deny the existence of the /3 phase, we must also wipe out the break in the line EGI at the point G, and in that case the point corre- sponding to the line GI should be simply Ar^. In any case, therefore, the term Ar^^ is illogical, but, since it is frequently used, some reference to it was essential at this point. As we have already remarked that the line FG is without any apparent influence on the micro-structure of steel quenched above it or below it, the absence of this line from the portion of the diagram relating to the steel now under discussion (0-6 per cent, carbon) will not affect the range of micro- PLATE XIV. Fi(i. (>9. Via. 70. Fig. 71. Fig. 72. [Tofncej). 175. THE IRON-CARBON SYSTEM 175 structures met with in this steel as compared with one of lower carbon content ; only, in accordance with the indications of the diagram, the higher-carbon steel will, when slowly cooled, contain a much larger proportion of the eutectoid pearUte than is the case with a lower carbon content. This is well brought out by comparing Fig. 69, Plate XIV., relating to this steel, with Fig. 67, Plate XIII., relating to the other. On quenching at higher temperatures, the features met with in the milder steel are largely reproduced, again with the modifications due to the larger proportion of the 7-iron sohd solution and the smaller excess of free a-iron or ferrite. The higher carbon content also seems to favour the development of the acicular structure of the quenched soUd solution, which is much more marked in the present steel than in the previous one. There features are shown in the photo-micrograph (Fig. 70, Plate XIV.), which shows the structure of this steel after quenching from 750° C. We now have to consider the steel of eutectoid composition — a point determined by Arnold {") as lying at a carbon content a very Uttle below 0-90 per cent, of carbon. This figure, how- ever, will vary very perceptibly if any other elements, such as sihcon or manganese, are present to an appreciable extent. Such a eutectoid steel will, in accordance with the diagram, exhibit only a single arrest-point on either heating or cooling. Following the somewhat illogical method already referred to above, this point is called Ac^gj and Ar^gj on heating and cooling respectively. It is a very strong evolution or absorption of heat — so strong, indeed, that if a piece of such steel is allowed to cool in air, when the evolution of heat which occurs at the transformation of the sohd 7-iron solution to the eutectoid pearHte takes place the steel is seen to glow visibly. This phenomenon received the name " recalescence " long before its true nature was understood, and the thermal curves of steel are for that reason sometimes called " recalescence curves " — a misleading term whose use is to be deprecated. Microscopically, steel of approximately 0-9 per cent, of carbon consists, when slowly cooled, entirely of the pearhte con- stituent, thus corresponding to the pure euteotic of ordinary binary alloys. Immediately above the temperature of the line 176 STUDY OF PHYSICAL METALLURGY HIJ, however, such steel consists entirely of the homogeneous solid solution of cementite in 7-iron. Ordinary quenching is obviously unable to retain this homogeneous solid solution completely unchanged, for the acicular constituent, and fre- quently some other forms, make their appearance. A typical structure is shown in Fig. 71, Plate XIV. This is, of course, a steel hardened by quenching, for while even mild steels, con- taining as Httle as 0-2 per cent, of carbon, are, in the strict sense of the word, " hardened " as the result of quenching from tem- peratures above the line EGI, this hardening is very slight until a carbon content of 0-6 per cent, is reached, and it only attains its full value at or near the eutectoid composition. For carbon contents beyond the eutectoid composition, an upper thermal point, corresponding to the crossing of the line ID, again makes its appearance, and the structure of the slowly- cooled steel again shows two weU-defined constituents. One of these consists of angular crystals of an obviously hard material which lie embedded in the pearUte. We have here crystals of cementite separated from the 7-iron solid solution on crossing the hne DI, embedded in the pearUte resulting from the decom- position of the residual sohd solution on crossing the line IJ. A typical example of this structure is shown in Fig. 72, Plate XIV. If such a steel is quenched from a temperature above the Une ID, provided that time enough has been allowed for the somewhat slow solution of the free cementite in the soMd solution, we should expect to obtain a homogeneous mass of the unchanged 7-iron solid solution. In the quenching of lower carbon steels this has never been attained, the incipient decomposition or degradation typified by the acicular structure being always found. In these higher carbon steels it appears that a portion of the 7-iron solid solution is actually retained in its original homogeneous state, and we find streaks of this entirely un- changed sohd solution lying among the acicular degradation product. This gives to severely-quenched high-carbon steels a very striking appearance, which is illustrated in Fig. 73, Plate XV. The account of the micro-structures and constitution of the iron-carbon alloys up to a concentration of about 2 per cent, of THE IRON-CARBON SYSTEM 177 carbon might be regarded as adequately concluded at this point, at all events for the purposes of an introductory survey, were it not for the fact that the importance of hardened and tempered steel lends very great interest to the structures met with in quenched steel. It has already been indicated that quenching never quite results in the retention of the structure as it exists at the instant when rapid cooling is begun, but that the very changes which it is intended to inhibit always occur to a small extent — the actual extent depending upon the circumstances of quenching and on the nature of the metal in question. In the case of steel the products of the incipient decomposition which takes place in these circumstances are particularly complex and interesting, although it is probable that similar complexity could be discovered in other alloys if they were studied in the same way. At any temperature and composition corresponding to the area ADIGE of the constitutional diagram the alloys consist of the simple, homogeneous 7 solid solution, having, as we have seen, the typical polyhedral structure of a " simple " alloy. The character of this y solid solution, however, must vary from one side of this region to the other, owing to the difference of carbon content — the 7 phase being pure iron in the 7 condi- tion at one side of its range of existence, and a soUd solution containing as much as 2 per cent, of carbon at the opposite extreme. Its behaviour on quenching is, therefore, correspond- ingly different, but these differences may be readily understood if it is realised that the presence of the carbon or iron carbide in solution must materially affect the manner, and still more the rate, at which the solid solution can undergo transformation. Where there is Mttle or no carbon present, the 7-iron molecules, in becoming transformed into the other allotropic form, have merely to re-arrange themselves in situ — there is no need for any transportation of matter through relatively considerable distances. Where the solid solution contains carbon, its normal decomposition involves the special separation of a carbon-rich constituent (pearhte) from a constituent free from carbon (ferrite), and this means that the carbide molecules must undergo a definite change of place, so that this decomposition implies P.M. N 178 STUDY OF PHYSICAL METALLURGY the movement of matter (iron carbide) through distances which are very large compared with molecular dimensions. Such movement necessarily requires time, and we consequently are led to expect that while it is difficult to inhibit the transforma- tions of pure iron by quenching, yet the rate of cooling attained by such means is sufficient to prevent the completion of the transformation in the presence of carbon, and more completely the higher the carbon content. This conclusion is in striking accord with the observed facts. While even the most severe quenching produces little effect in pure iron, it produces very marked effects in high-carbon steels and effects of intermediate intensity in steels of intermediate composition. We thus meet with a series of intermediate stages, between the theoretical hmit on the one side — ^never attained in practice — of the com- plete preservation of the original 7-iron solid solution, and the complete and unhindered transformation of pure iron. These intermediate or transition products have a special interest on account of their typical occurrence in hardened and tempered steels, and have, therefore, received much special attention and a series of personal names. The complete preservation of the 7-iron solid solution by quenching is, as already indicated, never attained in pure carbon steels. In steels of the highest carbon content, quenched violently from very high temperatures, traces of the undecom- posed 7 solid solution remain as white, structureless streaks, running across the acicular constituent already mentioned. This structure has received the name of " Austenite," and by that term, as adopted by the International Association, we now understand the 7-iron solid solution when preserved as such down to the ordinary temperature — either as the result of quenching or by the presence of a third alloying element, as described below. The first stage in the decomposition of the 7-iron soHd solu- tion consists in the formation of the acicular constituent already described and illustrated (see Figs. 70 and 71, Plate XIV.). This constituent has received the name of " Martensite," and much speculation has been offered as to its true nature. It is un- doubtedly the hardest constituent of hardened steels, but it THE IRON-CARBON SYSTEM 179 varies very widely in the details of its structure. When steel has been quenched from very high temperatures, the Martensite is found to occur in large grains showing a rather coarse acioular structure. When, on the other hand, a steel of eutectoid com- position is quenched at a temperature just above the critical point Arj, then the structure of the resulting Martensite is exceedingly fine — so fine, indeed, that some workers, not perhaps 20,000 EOO 1000- 1100 FlQ. 100 BOO 300 Temperature in Degrees Centigrade 74. — Temperature-Tenacity Curve for very Soft Steel at high Temperatures. provided with the best of microscopic apphances or employing unsatisfactory etching methods, have failed to detect the acicular structure, and have claimed that the best hardened steel contains a structureless " Martensite," which they have some- times distinguished by the term " hardenite " — a purely local term whose general use is not to be recommended. The careful examination of samples of the best and most carefully hardened steel, however, has convinced both the author and the majority of impartial observers that Martensite can always be shown to n2 180 STUDY OF PHYSICAL METALLURGY have an acicular structure, although in " properly hardened " steel that structure is exceedingly minute. The question, " What is Martensite and to what is its great hardness really due ? " cannot as yet be answered quite con- clusively, particularly as several rival views are in the field. These may, however, be briefly indicated. The first and simplest is the purely " allotropic " view according to which the hardness of Martensite is due to the presence in it of a notable proportion of " hard " /3-iron. Although attempts have been made to discredit this view recently, there is still much to be said in its favour. Thus it has been shown that the transformation from a to /3-iron in nearly carbonless iron is accompanied by a sudden step-up in strength. This is indicated in Fig. 74, which reproduces one of the temperature strength curves of an iron containing 01 per cent, of carbon from the work of the author and Humfrey {^). This step-up is not so large as the author had at one time supposed, but its existence is none the less significant, and is quite consistent with the view that, if it could be retained in that condition down to the ordinary temperature, ^-iron would be very hard indeed. On this view, then, the acicular structure of Martensite would be due to the formation, on the cleavage planes of the original homogeneous y-iron soMd solution, of needles of /3-iron. Normally such /8-iron would be obliged to expel the carbon whicli it had held in solution before the transformation, but, during quenching, time for such separation would not be available, and the)8-iron would be compelled to retain in "forced solution," or in very fine suspension, the carbon thus unavoidably retained in situ. It may well be that the presence of this " retained " carbon is the real cause of a development of great hardness in the Martensite, and in that case the transformation of the y-iron might be regarded as taking place direct to the a form without reference to the occurrence of the /3 phase, whether hard or otherwise. An alternative theory of hardening, which may be briefly termed the " amorphous " theory, has recently gained much ground. Like the ^S-iron theory, it explains the hardness of quenched steel by postulating the existence of an intrinsically THE IRON-CARBON SYSTEM 181 hard but unstable transition product, which is formed when the transformation of the homogeneous 7-iron solid solution into ferrite and carbide is hindered by quenching. Instead of identifying this hard substance with /3 iron, which has only a short range of stabiUty in pure iron and low-carbon steels, the new theory ascribes hardening to the presence in the steel of amorphous layers similar to those which are believed to be the cause of the strain-hardening of ductile metals. This hard amorphous phase has, of course, no stable existence below the solidus curve of the alloy system, but if its existence is admitted it serves to explain the phenomena of hardening in a very simple manner and to correlate the hardening of steel by quenching with the hardening of ductile metals by plastic deformation. On the question of the manner in which amorphous layers are formed in steel during quenching, several rival views have been put forward. The first ascribes the genesis of the amor- phous metal to the same cause as that which is operative when metals are strained ; it has been suggested that the interlacing needles of Martensite are merely extremely fine twin lamellae and that Austenite and Martensite are merely the twiimed portions of the same constitutent. This view meets with the insuperable difficulty that although quenching does set up very severe internal stresses in steel, it does not cause any serious internal flow or movement. The strain-hardening of metals, however, only becomes marked when severe plastic flow has occurred, and the conclusion is unavoidable that the deformations which occur during the quenching of steel are altogether too minute to produce the severe straining required to render the metal partially amorphous. A more rational view is that the y-iron crystals in passing through the transition temperature break up, leaving the molecules temporarily in a chaotic (amorphous) condition pending their re-organisation into crystals of a-iron and cementite, and that it is only this re-arrangement which is stopped by the rapidity of coohng. The present author is inclined to accept this view, which has been put forward by Humfrey, with the modification that the breakdown of the 182 STUDY OF PHYSICAL METALLURGY y-orystals only occurs in thin films around the boundaries of growing crystals of a-iron. This modification avoids the necessity for assuming that there are at any time present in the steel considerable quantities of the essentially unstable amorphous phase. According to this view, which can as yet only be regarded as a working hypothesis, the y-iron solid solution on reaching the temperature of Ars begins to undergo decomposition, by the formation of a large number of minute crystals of a-iron. These minute crystals will be formed on the planes of slip and of cleavage of the y-crystals, as some little movement must take place at these points, thus causing disturbance which is favourable to transformation. Each of these extremely numerous but very minute a-crystals will be surrounded by a film of amorphous metal which is at any moment in the act of passing to the growing a-crystal and is continuously renewed from the mass of the surrounding y-crystals. By rapid coohng, however, the growth of these a-crystals is stopped at a very early stage and the quenched steel is arrested in a condition in which it consists of numberless minute a-crystals surrounded by layers of hard and strong amorphous metal and possibly embedded in some unchanged y-iron. The hardness of the quenched steel is then ascribable to the presence of an extremely minute network of amorphous layers. The amorphous material of these layers will not only possess the hardness of amorphous iron, but wUl be rendered still harder by the presence of carbide in a high state of concen- tration. The minute a-crystals, as they are formed, must reject the carbide which was present in solution in their y-iron mother-crystals and this carbide will be rejected into the surrounding amorphous film. This film, being in character identical with the liquid phase, will take up all the carbide that is thrown out, and will retain it on subsequent cooling and congeahng, so that the crystals of a-iron will be surrounded by films — ^not of amorphous iron, but of an amorphous iron- carbide solution. The amorphous theory as thus outlined, affords explanations of many of the more important phenomena connected with the behaviour of quenched steels and of alloy steels, but the matter cannot be pursued here, partly because PLATE XY. Fig. 7;i. Fig. 75. [To face jj. 182. THE IRON-CARBON SYSTEM 183 the whole conception is still so new that opinion upon it must necessarily rematn tentative for some time to come. The Martensite stage found in quenched carbon steels, although theoretically of the greatest interest, is by no means the last of the transition stages met with. If the rate of cooling is rather slower, we find that the edges of the Martensite grains, more particularly where they border upon ferrite, in steels containing less than the eutectoid proportion of carbon (some- times called hj^o-eutectoid steels) appear to have undergone a further change. Instead of the faintly-coloured interlacing needles, this part of the structure appears, under most etching reagents, very deeply coloured and shows characteristic rounded or " woolly " outlines. By a suitable rate of cooHng, or by subsequent re-heating to a suitable temperature — i.e., by " tem- pering " — the whole of the Martensite can be transformed into this very dark-etching constituent to which the name of " Troostite " has been given. Controversy has turned a good deal upon the constitution of this " constituent " also, but we cannot enter into that question here. It is sufficient for our present purpose to regard Troostite as a further step in the degradation of Austenite into ferrite plus pearUte. The dark colour on etching appears to be due to the condition of the carbon in this stage ; the separation of the dissolved carbide from the iron must have begun as soon as the first particles of 7-iron became transformed into the /3 or a condition. In the Martensite stage this separation only makes itself felt very slightly by the small differences of colour between the various systems of interlacing needles. When the Troostite stage is reached we are probably deahng with a material in which the separated iron carbide is present in slightly larger masses. According to Benedicks {^^), the carbide in Troostite is present as a colloidal suspension, and it is quite probable that Troostite only differs from Martensite by the sMghtly greater segregation of the carbide particles. Troostite, at all events, is most frequently found in association with Martensite ; its typical appearance is shown in Fig. 75, Plate XV. When the rate of cooling is further lessened or the temperature of re-heating or tempering is further increased, the steel is permitted to progress somewhat further 184 STUDY OF PHYSICAL METALLURGY towards the final state of ferrite plus cementite. Instead of the Troostite described above, we find a constituent still devoid of any visible detailed structure under the highest available magnifications, but not etching to such a deep colour aa Troostite and without the peculiar rounded outUnes. This is often called " Sorbite," but it may be regarded as simply a variety of pearlite in which the two constituents are still so finely divided that they cannot be microscopically separated. The dividing Une between pearlite and Sorbite, indeed, is simply a question of microscopic resolving power. As between Martensite and Troostite on the one side and pearMte and Sorbite on the other, however, there is a very important distinction ; the two former are found essentially in steels quenched from temperatures above the line HIJ, while the latter are only met with if the rate of coohng through that temperature has not been unduly accelerated. Troostite is most readily obtained by oil-quenching from temperatures just above HIJ, or by water- quenching while the steel is actually in the critical range Ar^. Rapid air-cooling of steel specimens from temperatures above Aj-^ is sufficient to produce Sorbite, which merges into pearhte as the rate of coohng is decreased. Sorbite is as typically associated with pearhte as is Troostite with Martensite, although under special conditions all four may be met with in a single piece of steel. Having dealt with the theory of hardening steel and with the micro-structures obtained by quenching, a few words may be said at this point on the practical bearing of the constitu- tional diagram and our interpretation of it on the hardening operations which are so widely employed in practice. In the first place, it is evident that in order to harden a carbon steel it must first be heated to a temperature high enough to ensure the entire passage of the steel beyond the Une HIJ. In the case of steels lying near the eutectoid composition it is safe to go a little further and to state that, for satisfactory hardening, the passage of the entire steel into the region of the 7-iron soHd solution is desirable. For this purpose the steel must be heated high enough to pass through the critical points Ac^ and Acjg, ■which in steels used for hardening generally he close together. THE IRON-CARBON SYSTEM 185 It must not be forgotten, however, that the temperature of Aci is from 30° to 40° C. higher than that of Ar^, so that the indications of heating curves rather than of cooling curves must be followed in deciding the temperature to which the steel must be raised. Next, for reasons which wUl be explained in the chapter on heat-treatment, the steel must not be kept at this maximum temperature longer than is necessary to ensure its complete transformation. The steel must then be allowed to cool down, the rate being moderate, until it reaches a temperature just above that at which the critical change Ar^ would take place, and then it must be quenched. Exactly at what temperature above the critical point quenching must take place will, of course, depend upon the size and shape of the object and the nature of the quenching bath. For small tools, good results are generally obtained by first heating to 760° C. and then cooling down and quenching in water from a temperature of 700° C. The resulting steel should consist of fine-grained Martensite and should be fully hardened, i.e., in the condition sometimes described as " gleiss hard." Such steel then requires " tempering " by re-heating to various temperatures, generally below 400° C, resulting in the transformation of more or less of the Martensite into Troostite with a corresponding softening and toughening effect. The treatment depends, of course, upon the purpose for which the steel is required. While this method gives very good results, it still exposes the steel, in most cases, to unnecessary strain by allowing it to be first hardened fully, i.e., to a degree not required for the final purpose. Now hardening involves a risk, owing to the fact that rapid coohng sets up severe internal stresses in the steel, and these may, and frequently do, lead to cracking and warping, particularly in objects of considerable size or com- plicated shape. For that reason it is found desirable to quench in the first place from the lowest temperature which will ensure the attainment of full hardness, i.e., just above Ar^. Where a lower " temper " is sufficient, however, the same object can be attained by coohng from that temperature (just above Ar^), but not by the vigorous method of quenching in a large bath of cold water. A slower rate of cooling, which allows the steel to 186 STUDY OF PHYSICAL METALLURGY acquire the desired temper at once without re-heating or " letting down," can be secured by the use of oil, or of hot water, or even by quenching in measured small quantities of water which reach the boiUng point a certain short time after the steel has been immersed in them. Such methods, if care- fully appHed, not only reduce the risk of loss from cracking or warping, but also save the time and cost of subsequent tempering. The constitutional diagram of the iron-carbon system has so far been considered only in regard to aUoys whose composition places them to the left of the point D. We have now to consider the remainder of the diagram, relating to the alloys richer in carbon. Very roughly speaking, the alloys so far considered may be classed together as " steel," while the remaining members of the system include principally those materials known as " cast-iron " — ^but the division is by no means accurate or satisfactory, since the essential difference between " steel " and " cast-iron " is not one of carbon content. From the discussion of the diagram already given, the constitution of the various phases found and, consequently, the micro-structures of the alloys lying to the right of the point D, can be to a large extent inferred. The slowly-cooled alloys of concentration lying between the point D and the line BJ should consist, according to the diagram, of crystals of cementite embedded in pearhte. This cementite, while not primarily formed by crystallisation from the molten metal, is derived from the eutectic which forms along the line DB, and also from the solid y-iron solution which has formed the primary crystaUisation, but wLich, during the range of cooling from DB to IJ, has rejected some of the cementite which it originally held in solution. The cementite thus rejected from the solid solution would, with moderate rates of cooling, aggregate itself about the cementite crystals of the eutectic in such a way that these would grow at the expense of the crystals and lamellise of the solid solution. What finally remained of this soKd solution would decompose at the temperature of the line IJ, and the cementite of the alloy would thus be surrounded by pearlite. This structure is not often met PLATE X\l. Fi(i. 7(i. Yui. Fig. 7s. Fig. 79 ['/'ofin'{> p. 187. THE IRON-CARBON SYSTEM 1B7 with in actual samples, since " a moderate rate of cooling " generally results in a partial reversion to the mode of crystaUisa- tion of the stable system, with the resulting formation of graphite. On the other hand, perfectly " white " irons, both of this group and of the final group lying to the right of the line BJ, can be obtained by fairly rapid cooHng, and these consist of cementite and pearlite, as indicated by the description given above. An example of this structure is given in Fig. 76, Plate XVI. If, however, the coohng has been somewhat more rapid, the decomposition of the solid solution along the line IJ may be more or less completely prevented, and the resulting structure is then a mixture of Martensite, or of the transition forms, such as Troostite, with cementite. With regard to the highest carbon alloys lying to the right of the line BJ, it need only be said that the tendency to graphite formation decidedly increases with the carbon content, and these irons can only be maintained in the " white " condition, i.e., free from graphite, by very vigorous chilling or rapid cooUng from the molten state. When obtained in the graphite-free condition their structure consists of crystals of cementite embedded in a matrix derived from the soUd solution which formed one of the constituents of the eutectic formed along the Une DB. This may take any form ranging from the Martensite- Austenite structure of quenched high-carbon steel to that of pearhte. As has already been indicated, there is a strong tendency in the alloys richer in carbon to foUow, in part at least, the mode of sohdification of the stable iron-graphite system indicated in the diagram of Fig. 65. This appears, however, to affect only the regions of higher temperature in the diagram ; it would seem that the formation of graphite in the earlier stages of solidification practically results in the formation of an alloy of much lower carbon content — corresponding to the carbon removed in the form of graphite — ^which then follows the meta-stable mode of further cooUng and transformation of these lower-carbon alloys. The result is that the structure of grey iron — as graphitic iron is generally called — ^is practically that of steel with which a certain amount of graphite appears 188 STUDY OF PHYSICAL METALLURGY to be mixed as a species of impurity (in addition to other im- purities frequently present). It is not strictly correct, however, to regard these irons simply as impure steels, for the reason that the presence of the graphite plays an important part in the behaviour of these materials when heated, affecting the results of such operations as annealing and quenching, while, of course, it also reduces the strength and destroys the ductihty of the material to a very considerable extent. For the purposes of describing the micro-structure of the slowly-cooled iron, how- ever, the similarity to steel mixed with a certain proportion of graphite is very useful. The grade of steel to which any given specimen of cast-iron or pig-iron corresponds when examined under the microscope depends on its initial carbon content, and also upon the amount of graphite which has been deposited in the earlier stages of cooling. Generally spealdng, the irons containing the lower proportions of carbon tend to resemble the lower carbon steels, while the irons richer in total carbon tend to resemble the high-carbon steels, but this rule is subject to very wide variations according to the rate of cooKng and to the presence or absence of impurities — such as silicon and sulphur — ^which affect the tendency of the metal to deposit graphite. The final structure will accordingly depend upon the amount of " com- bined carbon " present in the iron. Slowly-cooled iron containing less than 0-9 per cent, of combined carbon will, on the basis of the analogy with steel which has been stated above, consist of a mixture of ferrite and pearlite, graphite being interposed in plates or nodules accord- ing to circumstances — an example of this kind of structure is given in Fig. 77, Plate XVI. An iron containing more than 0-9 per cent, of combined carbon, on the other hand, will consist of cementite and pearlite mixed with graphite, but it sometimes happens that the pearUte — as a result of slow cooling — separates out into very widely-differentiated layers or bands, so that the ultimate structure is really better defined as ferrite and cementite mixed with graphite. The structures just described result from slow cooHng carried on from the molten state down to the ordinary temperature ; THE IRON-CARBON SYSTEM 189 if the cooling is rapid throughout, the meta-stable system will be followed and no graphite will be formed, but a further condition of the iron must be considered which results from subsequent re-heating of iron which was originally cooled rapidly. The re-heating of white iron consisting of cementite-pearlite, if carried to a sufficient temperature, always results in the liberation of carbon in the free state — ^really a form of graphite although it is generally known as " temper-carbon." This free carbon is, of course, derived from a partial transition of the alloy from the meta-stable to the stable system, some of the cementite being broken up into ferrite and carbon. As a " phase " this carbon is identical with graphite, but, being formed in the solid metal, it occurs in a very finely-divided condition which reveals itself in a very characteristic manner when such iron is dissolved in certain reagents. The micro- scopic appearance of " temper-carbon " is shown in Fig. 78, Plate XVI., which shows the structure produced in a perfectly white cast-iron by heating to 900° C. for several hours in vacuo-. The fact that the heating in the above example took place in vacuo is of importance, since the phenomenon in question is liable to be masked by others if the iron is heated in an oxidising atmosphere or in other oxidising surroundings. The result in the latter circumstances is a tendency for the iron to become more or less decarburised by the oxidation of the carbon which it contains. The first effect of such an oxidising annealing process is probably the conversion of combined carbon into temper-carbon or finely-divided graphite, but, either by the penetration of gases or by the diffusion of solid oxides of iron and their re-action with the finely-divided temper-carbon, the latter is oxidised and removed from the metal. If this process is carried far enough, the structure of the resulting iron is simply that of comparatively pure steel (ferrite and pearlite) with a few patches of temper- carbon. This process is used to a considerable extent in the production of " malleable castings " — the castings in question are produced from cast-iron in the ordinary way, but they are Subsequently packed in oxide of iron and exposed to a high 190 STUDY OF PHYSICAL METALLURGY temperature for a long time — the resulting decarburised iron is thus brought to a condition more or less resembling that of very- mild steel, and it is thus rendered much softer and more ductile than the original cast material, but its strength cannot and does not become equal to that of actual steel, more particularly of worked or forged steel. There are two reasons for this differ- ence ; in the first place, the material is in a condition of crystaUisation formed by casting alone and without that refining which results from forging or rolling, and this relatively weak and coarse structure is further inevitably affected in an unfavourable manner by the prolonged heating to which the castings are subjected. In the second place, it must be borne in mind that the removal of the temper-carbon or graphite is never complete and the residual graphitic areas diminish the strength of the material. Further, the chemical composition of a malleable casting, particularly as regards the impiu-ities which are generally kept within very low Mmits in the case of steel — such as sulphur, phosphorus and oxygen — ^is rarely such as to conduce to favourable mechanical properties in these decarburised castings. The process of decarburising cast-iron by prolonged heating in oxidising conditions is that known as Reaumur's ; in modern practice, however, there is a tendency to prefer the American or " black heart " process. In this there is no real decarburisa- tion; all the carbon originally present in the castings is precipitated as temper-carbon, whose presence gives these "malleable castings" their characteristic black fracture. This process has the advantage that oxidation is avoided and some surprisingly good results have been obtained, where pure materials are treated in oarefuUy-controIled stoves. The process whose theoretical foundation we have just men- tioned, i.e., the softening or rendering malleable of hard iron castings by heating under oxidising conditions, finds an exact converse in the widely-used process of " case hardening," or cementation, by which a soft low-carbon steel is rendered hard by the addition of carbon resulting from prolonged heating in carburising surroundings. Such carburising sur- ivoundings may coijsist merely ,of an atmosphere of carbon THE IRON-CARBON SYSTEM 191 monoxide gas, or the steel may be packed in a carbonaceous powder, although it has been found that the process is always facilitated by the presence of nitrogenous matter, the gas cyanogen beii^g particularly powerful as a carburising agent. The process of case hardening consists in the absorption of carbon by the outer layers of the steel, thus forming a coating or " case " of high-carbon steel. When the piece thus coated is subsequently quenched at a suitable temperature, the low- carbon core is but little affected thereby, while the high- carbon case is left in the Martensitic condition and is therefore exceedingly hard. Provided that the transition from the one region to the other is not unduly abrupt, the two portions of the piece support one another very well, and a very valuable class of article is produced. The absorption of carbon by heated low-carbon steel can, of course, only occur when the steel is at a temperature above Ar^, and it does not occur readily until the whole of the steel is in the condition of 7 solid solution. In that condition the steel readily takes up carbon, which slowly diffuses inwards. The complete cross-section of a steel rod case-hardened in this way is shown in the frontispiece of this volume, which is taken from a composite photograph built up of some 140 separate photographs. The dark ring sur- rounding the mild-steel core consists of the high-carbon " case." The transition from " case " to " core " in a hardened steel is shown in Fig. 79, Plate XVI. In concluding this chapter it must be pointed out that the constitution and structure of the iron-carbon alloys have only been very briefly surveyed, and that many matters of funda- mental importance have either been barely mentioned or have not been dealt with at aU. Some of these, such as the effects of both good and bad heat treatment, are closely related to that aspect of Physical Metallurgy which is treated in the second part of this book, viz., the relation between micro- structure and mechanical and other physical properties, so that their consideration is left to that part of the book. Another matter of fundamental practical interest, although of com- paratively httle direct interest from the point of view of general principles, ig the nature, maimer of occurrence and influence 192 STUDY OF PHYSICAL METALLURGY of impurities — not only in the iron-carbon alloys, but in the pure metals and in all systems of alloys. To a slight extent this question is dealt with in the chapter on the defects of metals, but for fuUer treatment reference must be made to separate treatises deaUng with the various groups of alloys. It must be further pointed out that our treatment of " steel " has been confined to a very brief consideration of the pure or nearly pure binary aUoys of carbon. The ternary and quaternary alloys of this system, including all the " special " or " alloy " steels whose introduction into metallurgical and engineering practice has constituted the greatest of recent metallurgical advances, cannot well be treated, even super- ficially, in an introductory volume such as this. Their adequate treatment requires, not a separate chapter, but a separate volume. References. (1) Arnold. Proc. Inst. Mech. Engineers, 1891, p. 683. (2) Howe and Sauveur. Proc. International Testing Assoc, VI., Congress, New York, 1912. (3) Euff. Metallurgie, 1911, VIII., 456 and 497. (4) Gutowsky. MetaUurgie, 1909, VI., 731. (5) Burgess and Crowe. Trans. Amer. Inst. Mining Engineers, XL VI., 1913. (6) Carpenter. Journ. Iron and Steel Inst., No. 1, 1913. (7) Benedicks. Journ. Iron and Steel Inst., No. II., 1912. (8) Eosenhain and Humfrey. Proc. Roy. Soc, 1909, Vol. 83a, p. 200. Journ. Iron and Steel Inst., No. I., 1913. (9) Baykoff. Revue de Metallurgie, 1909, VI., 829. (10) Saniter. Journ. Iron and Steel Inst., No. II., 1897 ; No. I., 1898. (11) Arnold. Proc. Inst. Civil Engineers, 1895, CXXIII.. p. 127. (12) Benedicks. Zeitsohx. f. Chem. u. Industrie der EoUoide, 1910, p. 290. PART II. THE PROPERTIES OP METALS AS RELATED TO THEIR STRUCTURE AND CONSTITUTION CHAPTER IX THE MECHANICAL TESTING OF METALS In our brief survey of the constitution and structure of some of the important and typical alloy systems, references to the physical properties of the materials, and particularly to their strength and ductility, have frequently been made. Since the large-scale production of metals depends essentially on their engineering uses, those properties which are of primary import- ance for the use of metals or alloys in structures or machines are also of the highest importance in the production of the metals. Weak and brittle alloys have, in special cases, found uses of more or less importance for special reasons and for special purposes, but these uses are insignificant compared with the engineering uses of steel or even of brass or bronze. The methods of ascertaining the mechanical properties of metals are, therefore, of the highest interest alike to the metallurgist and the engineer. These methods have, how- ever, been very largely developed by the engineer for the purpose of controlling the products of the metallurgist and of ascertaining how the materials sold to him are likely to behave in actual use. It will probably be readily admitted that the only completely reliable test for that purpose is actual use itself, and it is upon the results of that very exhaustive test that engineers endeavour to base their future specifications. The test of actual use has, however, the serious disadvantage that it is generally slow and always extremely expensive — indeed, in the majority of circumstances where tests are called P.M. o 194 STUDY OF PHYSICAL METALLURGY for in order to ascertain the quality of a given material, that test is totally inappUcable. As a result, engineers have elaborated various systems of tests intended to furnish the information as to how the material wiU behave in practice by means of simple experiments which can be completed in — at most — a few hours. A test of this kind is frequently intended to imitate the conditions of actual service. Where such an imitation can be attained on a small scale very valuable results are found by its aid. The testing of ship-models in modern ship-tanks, or of aeroplane or airship models in wind-channels constitute favourable examples of this kind, but their nature also illus- trates the Hmitations of such a method. This limitation may be readily expressed by sajring that the small scale model wiU give a reasonable result if the law of similarity is observed. Whether this condition is fulfilled can be ascertained in most cases by varying the scale of the experiments ; if the results have a real meaning it should be possible to arrive at concordant results for the behaviour of the full-size object from the data of experiments, whatever the actual scale of the experiments may have been, within reasonable hmits. Applying this criterion leads to the conclusion that in a mechanical test the results should either be independent of the actual dimensions of the test-piece, or at least related to those dimensions by some simple geometrical law. If this criterion is appHed to the majority of the more empirical tests sometimes advocated in the testing of metals, the results of the tests frequently fail to come up to the requisite standard. The reason is not far to seek. In the conditions of practical use engineering materials are exposed to complex conditions and to forces whose combined eefEcts it is not generally possible to calculate. When experiments or tests are devised to imitate these complex practical conditions the imitation is never perfect and only rarely " to scale." Usually it is necessary to reduce the time factor in the experiments, and in order to do this — i.e., to secure the fairly rapid failure of the test-object in the experiment — some other factor must be enormously exaggerated. Where the material might, if THE MECHANICAL TESTING OF METALS 195 unsatisfactory, fail in actual service in the course of months or years, in a test it must be made to fail in a few hours or at most a few days. The difficulty lies in bringing about such rapid failure without radically altering the conditions under which the material works. Beyond this difficulty lies the fact that in these " imitation " tests the conditions of the material under the test itself are ill-defined or incompletely loiown or understood. The consequence is that even duplicate apparatus fails to give strictly comparable results, and comparability is entirely lost if the size of apparatus or test-piece is changed. A recognition of these difficulties leads, fortunately, to a very clear and simple principle which is coming to be recognised and appreciated in the science of mechanical testing — as it has long been recognised and followed in all other scientific measurement. This is that, in order to obtain results which have a definite meaning, we must measure one property of our materials at a time. This is simply the principle of isolating the various factors of a problem and deahng with them one at a time, which has long been followed in all other branches of science. A single, definite physical quantity is chosen for measurement and the conditions of the experiment are carefully designed, so as to eliminate, as far as possible, the influence of any other factors. In this way measurements are obtained which can not only be accurately repeated at any time and place by any suitable instrument, but the values found are quite independent of the size of test-piece or other incidental circum- stances. Two difficulties arise in the apphcation of this principle to mechanical testing. The first is that the various physical properties which constitute " strength " are not yet fully under- stood and analysed, so that it is not as yet easy to say which and how many physical properties are of direct importance from the point of view of the engineer. Once the correctness of the principle is appreciated, however, this analysis will rapidly follow — the accumulation of really accurate and definite data will make it possible to unravel the physical relations involved in " strength " under various circumstances. Mean- while, the criterion that the ^esijlts must, subject to a simple 02 196 STUDY OF PHYSICAL METALLURGY formula ,be independent of the dimensions of the test-piece and of the testing apparatus, affords a guide in the whole matter. The second difficulty arises from the fact that in many cases the results of tests made on the principle here advocated are likely to furnish data which the engineer is not easily able to translate into practice. The answer is that such an objection really apphes with far greater force to the vague and often meaning- less results of some of the current " imitation " tests. Even among existing and well-established tests the engineer is only able to interpret those whose indications he has, through long experience, learnt to correlate — ^however imperfectly — with the results of actual use. If future development along the line here indicated takes place, the same process must be gone through — the results of tests must be correlated as accurately as possible with the data of experience ; if the tests are in themselves based on a sound principle and determine the reaUy fundamental properties of our materials, however, this correla- tion is bound to be far more satisfactory than can be the case with tests which are in themselves vague in their indications, owing to the complex mixture of properties which determine their results. Although the adoption of the principle of determining the individual fundamental properties of metals has been referred to above as forming the basis for future development in the science and art of testing, this must not be regarded as imply- ing that there are not, among the tests at present habitually employed, a number which conform to this principle. The process of natural selection has, in fact, been at work and, in spite of frequent voluntary departures from the guiding principle recognised above, those methods which have stood the test of time and have shown that they can and do yield really valuable and reliable data, have permanently established themselves in the favour both of practical engineers and of those concerned with the investigation of metals. In this and the following chapter we shall briefly describe and review the principal forms of mechanical test now in use or contemplated, and we shall see how far these tests follow or depart from our fundamental principle. No attempt will be made to give THE MECHANICAL TESTING OF METALS 197 working details of the numerous testing appliances which are now in use all over the world — experimental details wiU only be considered as far as they are essential to a discussion of the results yielded by any form of test. The best known and most widely used of all mechanical tests is undoubtedly the tension test, in which a piece of metal of known dimensions is firmly held at the two ends and exposed to a tensile force which is increased until rupture occurs. So firmly is this test established that for a time there was a strong tendency on the part of engineers to rely upon it exclusively ; more recently, however, a strong movement has made itself felt in the direction of requiring, in addition to the tension test, some form of test in which the metal shall be subjected to shock, vibration or " fatigue." That movement has so far justified itseK that few will be found at the present time who do not admit the need for some form of " dynamic " test, but the exact form of test to be adopted is still a matter for wide divergence of opinion. In accordance with what has been said above, the permanence of the tension test, and the favour which it has found in engineering practice, leads us to anticipate that its results have a definite physical meaning and that the test is sufficiently simple at least to approximate to our fundamental criterion of measuring one single property of the material. It is quite true that, within the elastic range, mathematical analysis shows that a tensional stress is equivalent to a pure shear plus a negative hydrostatic pressure, but it is doubtful how far that analysis really applies to the important stages of a tensile test which occur after the metal has passed the elastic Umit and is extending plastically. Not only that, but the component stresses of tension — shear and hydrostatic pressure in the negative direction — are not readily realised in any form of testing apparatus, although approximately pure shearing tests can be made. Therefore, although a tensile test may not be quite strictly a test of a single isolated physical property — i.e., the property of resisting a simple statical stress — yet it is a test under a very simple combination of stresses and for that reason may with some precautions be made to 198 STUDY OF PHYSICAL METALLURGY yield results having a very definite meaning and com- parable between different machines and different sizes of test-pieces. The machines used for the purpose of tensile testing are very varied in form and size, the latter ranging from a " toy " machine, such as that employed in the research of the author and Humfrey {'■) for measuring the tensile properties of thin little strips of red-hot steel, up to huge machines capable of exerting tensile forces of several thousand tons. All these machines, however, work to a certain extent on one common principle in that they possess three main organs ; the first of these serves to apply to one end of a test-piece of suitable shape the force required to produce the desired tensile stress. The second organ serves to " take up " the extension in the length of the test-piece which occurs when large stresses are applied to it ; in the case of ductile metals this extension may be very large indeed, and the machine must be capable of maintaining the load on the test piece in spite of these large extensions. Finally, the machine must possess an organ or apparatus for measuring the force applied to the test-piece in an accurate and reliable manner. As a rule, the first and second organ are combined by the use of an hydraulic ram, which applies the desired load and at the same time takes up any extension of the test-piece. The various types of testing machine differ principally in regard to the third organ, i.e., the apparatus by which the tension or load on the test-piece is measured. Broadly speaking, there are only two ways in which this is done, viz., by direct counterpoise through the medium of one or more levers and weights, and by some form of hydraulic pressure measurement. In this country the lever and weight form of testing machine is almost exclusively used, a result of the refusal of the Board of Trade to recognise tensile tests made upon any other type of testing machine. This is, however, a restriction no longer justified by facts, and one which therefore cannot long be maintained in view of the circumstance that the lever type of machine has been entirely superseded in Germany by one of the other types, which owes its great development largely to Martens (^). In THE MECHANICAL TESTING OP METALS 199 America, also, the splendid Emery (^) machines do not depend upon direct lever counterpoising. What we may term the " English " type of tensile machine is shown diagrammatically in Fig. 80. In that figure the test- piece, T, is attached to the two shackles or holders, S S'. The upper shackle hangs vertically downward from a stirrup, R, which rests, by a knife-edge, K, on the lever, L. This lever is free to turn in a vertical plane on the knife-edge, I. On the Pig. 80. — ^Diagram of the Single Lever Tensile Testing Machine. lever, L, runs the jockey- weight, W, which is so adjusted that, when in the zero position, at the left-hand end of the lever, it just counterpoises the weight of the lever itself. As the weight is run out to the right it can thus be made to counter- poise the puU of the test-piece, T, acting through the shackle, S, and the stirrup, R. The lower shackle, S', is attached to an hydrauhc ram, P, which, by the application of water pressure, derived either from the mains or from a special pump, can be caused to move downwards and thus to apply the load to the test-piece. If the test-piece stretches, the ram, P, will 200 STUDY OF PHYSICAL METALLURGY simply descend under constant pressure and take up the stretch before applying additional load, unless the water is supplied to the ram so fast that the stretch of the test-piece cannot keep pace with it ; in that case the pull on the test-piece may increase, even while fairly rapid extension is taking place. The necessity for the use of the lever in this type of testing machine, for the purpose of measuring the pull or load on the test-piece, is due to the fact that this pull caimot be directly deduced from the water pressure acting on the hydraulic ram. This discrepancy arises from the very large amount of friction which exists between the ram and the cup-leather which is used to make a tight joint between the ram and the cyMnder in which it moves. The lever, however, although convenient in many ways, also possesses several serious disadvantages. In the first place, the knife-edges require care and attention at intervals because they are necessarily exposed to severe stresses and even to shocks when a strong test-piece breaks suddenly. Next, the lever — ^Lf a single lever is used — ^is long and heavy and occupies much space. The whole machine, indeed, is bulky and costly. A photograph of a 100-ton tensile testing machine of this type, as installed at the National Physical Laboratory, is shown in Fig. 81, Plate XVII. In this machine the hydraulic cylinder and ram are in the pit below the floor level, while the hydraulic pump and accumulator are in another part of the building, so that the machine as shown in the picture is only a part of the whole installation. A further disadvantage of the lever and jockey- weight is that they possess a considerable amount of inertia, and may thus momentarily exert forces far larger than the scale-reading of the jockey-weight would indicate. Particularly when a test-piece is stretching rapidly this factor may introduce serious errors, especially if measure- ments of a more deUcate tjrpe are to be made. The hydraulic measuring type of tensile testing machine is a much simpler and more compact apparatus than the lever tjrpe ; Mke the other it possesses, as a rule, an independent hydraulic ram for the purpose of exerting the loading pull and of taking up the stretch of the test-piece, but the load on the test-piece is transmitted to a second hydraulic apparatus, PLATE XVII. S3 [To face 2J. 200. THE MECHANICAL TESTING OF METALS 201 which serves solely for the purpose of measuring the amount of pull exerted at any moment. This is accomplished by an appliance developed largely by Martens, and known in German as the " Mess Dose " or " hydraulic measuring box." In this appliance the pull of the specimen is transmitted to one side of a chamber which is partly composed of a flexible dia- phragm. This chamber is completely fiUed with water, which is in communication with a sensitive pressure gauge. When the test-piece exerts the pull on the movable side of this chamber the pressure within the chamber rises until the pull of the test-piece is exactly counterpoised, and the pressure then recorded by the gauge measures the pull of the test-piece, being, of course, proportional to the area of the movable side of the box. In this way the actual pull can be very much magnified, and the scale of the pressure gauge, in terms of tons, can be made very open. This arrangement has the great advantage of compactness and cheapness and the total absence of serious inertia. The results, however, necessarily depend upon the accurate calibration of the measuring box and pressure gauge and upon the constancy of these two. The latter has, however, been amply established by the continued re-calibrations under- taken at the Material Priifungsamt, at Lichterfelde, Berlin, and there can be little doubt that for accuracy and reliability this hydraulic measuring type of testing machine compares favourably with the lever types. Another simple type of testing machine, which also depends upon the hydraulic measurement of the load applied to the test-piece, is that developed by Amsler in Switzerland. In this type of machine the ram which applies the testing load also serves to measure it. For that purpose the ram is made a slightly loose fit in its cylinder and no cup-leather is used. The pressure is applied by means of oil, fed from a high-pressure pump, and a thin film of this oil is continually allowed to leak past the ram and to return to the reservoir. In this way the friction of the ram is reduced to negligible proportions and maintained at a very constant value. The pressure of the oil behind the ram is then used as a measure of the load on the test-piece, and for that purpose a very simple and accurate 202 STUDY OF PHYSICAL METALLUIIGY form of direct gravity pressure gauge is employed, in which the oil pressure is measured by the deflection of a heavy pendulum. This type of machine is found to give very accurate and rehable results, if properly caUbrated and kept in the requisite good order. In use it is decidedly more convenient than the lever type ; it occupies much less space and is free from inertia errors. Fig. 82, Plate XVIII. , shows an example of this type of machine, both the high pressure pump and the pendulum manometer being shown. Quite recently Dalby (*) has perfected a new type of instru- ment for the autographic recording of the indications of a testing machine when employed for tensile work, which is self- contained, in the sense that it obviates the need for any means of stress-measurement on the part of the testing machine proper. This result is achieved by the use of what is termed a " weigh bar," which is simply a piece of steel of high " elastic hmit " through which the puU of the machine is transmitted to the test-piece. Such a " weigh bar " is in reality a very stiff but perfect spring, since its elastic extension is always strictly proportional to the load upon it. The Dalby instrument, which will be referred to in greater detail when the determina- tion of " stress-strain diagrams " is considered, measures the extensions of such a weigh-bar and thereby determines the load on the test-piece. It is, of course, essential that the weigh-bar should be very much larger in cross-section than the test-piec©> so that the latter may be broken before the elastic limit of the former shall be passed. For the calibration of the weigh-bar itself — i.e., for the purpose of determining once for all what load is required to produce various elastic extensions in the weigh-bar — the use of some accurate form of testing machine of the ordinary type, or working by direct dead-weight loading is, of course, required — but one standard machine could readily supply calibrated weigh-bars for any number of simple machines having no separate measuring appliances. Leaving aside the question of the experimental details of testing machines as unsuited to the scope of the present book, we may now turn to a consideration of the nature of the data which are derivable from a tensile test. In order to understand PLATE XViri. Fig. «i'. {To face p. 202. THE MECHANICAL TESTING OF METALS 203 these we must briefly describe the course of events which take place when a specimen of metal is broken in a tensile test. As the load is gradually applied there is first a period in which no visible change occurs in the specimen — assuming, for example, that we are deaUng with a material like mild steel. Delicate measurements show, however, that even in this stage the metal undergoes a measurable amount of stretching. This stretch, however, is temporary or " elastic " and disappears almost entirely when the load is removed. It is found, too, that during this period the amount of extension is strictly proportional to the appUed stress, thus following the well-known law of Hooke, that (elastic) strain is proportional to stress. The actual amount of elastic stretch which occurs under a given stress varies very considerably with different materials. This amount is generally expressed in the form of the elastic constant, known as " Young's modulus," which is calculated, from measured loads and extensions, as the load in pounds per square inch which would suffice to stretch a piece of the material — ^if elastic stretching of such an amount were physically possible — to double its initial length. The same thing may be put rather more rationally by saying that Young's modulus is 1,000 times the stress, in pounds per square inch, which is sufficient to cause a test-piece to stretch by one thousandth of its original length. In mild steel this constant has a value of the order of thirty milHons, or 3 X 10* pounds per square inch. In other materials it varies from such a value as twelve millions for some light aluminium alloys to eighteen miUions for bronze or brass. The value of this constant, curiously enough, is very little affected by such processes as cold rolling or drawing, so that it appears to depend upon the nature of the atoms present in the material rather than on their arrangement. The measurement of the elastic stretching of metals under tensile loading is a matter of some delicacy, as the changes of length to be determined are very minute. Instruments for this purpose, known as extensometers, have been designed in various forms, but only two need be mentioned here. In one of these, designed by Ewing (^), clips are attached to two points 204 STUDY OF PHYSICAL METALLURGY :^ of the test-bar a known and definite distance apart — usually eight inches. From the upper clip a rod is suspended, to which is attached a microscope ; in the field of view of this micro- scope is a piece of glass engraved with a very fine scale, and this piece of glass is attached to the lower clip on the test-bar. The arrangement is diagrammatically shown in Fig. 83. The small glass scale moves with the lower clip, while the microscope moves with the upper, and the relative movement of the two shows the extension of the test-piece between the clips, the movement being, however, magnified in the first place by the system of levers which constitute the supports of the microscope and the scale, while the principal magnification is obtained optically in the microscope. With this instrument changes of length of sjs.hisTS of ^^ ^'^"^ ^^^ ^® measured. Even more delicate, but less con- venient, is the instrument employed by Martens (*). In this type of extenso- meter a clip is again attached to the one end of the test-bar ; from this clip rods pass down the sides of the test-bar and parallel to the face of the bar — shown as RR in the diagram of the arrange- ment given in Pig. 84. At the point P, corresponding to the lower cHp position in other extensometers, small rectangular prisms or double knife-edges are placed between the test-piece and the rods in such a way that the opposite knife-edges, E in the diagram, bear upon a rod and the test-bar respectively. If now there is any movement of the rod relatively to the bar, the little prism will be slightly tilted. In the Martens extenso- meter this prism carries a mirror, whose angular movements are measured by means of a telescope and scale. These give readings proportional to the minute extensions of the test-bar, since the test-bar will move relatively to the " idle " rod when the former undergoes elastic extension. These and other extensometers serve to determine the curve of 3C T-Test Piece PP-Points or Attachment S - Scale Fig. 83. — Diagram of the Ewing Extenso- meter. THE MECHANICAL TESTING OF METALS 206 r> o elastic stretching and the value of Young's modulus, as well as for the determination of what is known as the " elastic limit " or sometimes, perhaps more correctly, the " Umit of propor- tionaUty." Suppose that the load on a test piece is increased step by step and the extension is read with a suitable extenso- meter. If the values of stretch thus obtained are plotted against load as abscissae, a curve is obtained which represents the behaviour of the metal. A curve of this kind, on which the observed points are marked, is shown in Fig. 85, which is typical of mild steel. There is first a considerable range within which the observed points he on a slop- ing straight line. This is the elastic range in which stretch is proportional to load, and the slope of the hne merely indicates the value of Young's modulus — i.e., how much the material stretches for each increment of load. After a certain point, however, which in this example lies at a stress of eight tons per square inch, the observed points visibly fall away from the straight hne. This means that the material has ceased to obey Hooke's Law, the stretch is no longer strictly proportional to the load and the " limit of proportionality " has been passed. This "elastic Mmit" is a very important point in the behaviour of metals, but we shall see later that, as determined by a simple tensile test, its position is much influenced by the treatment which the metal has pre- viously undergone. The " primitive " elastic hmit, as this point is sometimes called when it is desired to point out that we are deahng with a simple determination upon a material in its " primitive " condition, may be displaced in either direction by means of mechanical treatment — ^it is raised by tensile over-strain, cold rolling or drawing, or any other mode of cold working which causes an extension of the material. If, on the other hand, the metal has previously been compressed longi- tudinally, or treated in such a manner as to produce diminution Fig. 84. — Diagram of Martens' Extenso- meter. 206 STUDY OF PHYSICAL METALLURGY of length, then the " primitive " elastic Umit will be displaced downwards from its true position. Unless, therefore, the material has been treated in such a way as to remove the effects of previous mechanical treatment, the determination of an elastic limit by extensometer is not very reliable. Bairstow (') has shown that, however much the apparent or " primitive " elastic limit may have been displaced by mechanical treatment ■008 ■D07 ■006 toos ^■004 %003 \ «: . ^^ .,9 eOJ*^ Mt. «<:'*>' 10 20 ZSyLZinc IS Composition Fig. 96.— Diagram showing Comparative Results of various Tests applied to a series of Alloys. (Zinc- Aluminium.) THE MECHANICAL TESTING OF METALS 239 subjected to a great variety of tests. An example of this kind is offered by the investigation of the alloys of aluminium and zinc contained in the Tenth Report to the Alloys Research Committee of the Institution of Mechanical Engineers by the author and S. L. Archbutt. The results of the various tests, referred wherever possible to material of strictly comparable kinds, are shown in Fig. 95, which is a graphical summary of the results of all the tests applied. The ordinates of the various curves there shown differ, of course, from one curve to another, so that no comparison of absolute values can be made ; the curves are all plotted to the same scale of chemical composition of the alloys, i.e., percentage contents of zinc, so that the shapes of the curves are strictly comparable. It will be seen at once that the curves of alternate bending impact, static elastic limit and ultimate stress all rise steadily with increasing zinc content ; the yield-point curve shows a curious dip before it takes the upward course, and this is balanced — and con- firmed — by the corresponding hump on the curve of elongation in the tensile tests. The alternating stress curve — i.e., the curve of half the elastic range — also rises steadily, although slowly, with increasing zinc content. The ductility curves, on the other hand — i.e., " reduction of area " and, after a pre- liminary " hump," the elongation curve also — ^fall markedly with rising zinc content. It is interesting to note that the curve representing the Arnold alternate bending test practically follows the ductihty curves, also falling steadily towards the right. The most remarkable curve, however, is that of the single-blow impact tests (Izod), which rises to a definite maximum and then falls, thus indicating that this test takes account of the combination of strength with ductihty. These curves are shown here as an Ulustration of the manner in which it is desirable that the various methods of test should be systematically correlated in order to enable us to form a correct judgment of their relative values and modes of interpretation. Eefebences. (1) Larrard. Proc. Inst. CivU Engineers, CLXXIX., 1909— -10, Pt. 1, 240 STUDY OF PHYSICAL METALLURGY (2) Brinell. Teknisk Tidsrift, XXX., p. 59. See also Wahlberg, Journ. Iron and Steel Inst., No. 1, 1901, p. 243. (3) Benedicks. Kecherches Physiques et Physico-Chimiques sur I'Acier au Carbone, Upsala, 1904, p. 75. (4) Saniter. Jonrn. Iron and Steel Inst., 1908, III., p. 73. Internat. Testing Assoc, New York Congress, 1912, Sec. A. (6) Ludwik. Zeitschr. d. Oesterreich. Ingen. U. ArcMtekten Vereins, LIX., p. 191 and p. 205. (6) Turner. Proc. Birmingham PMl. Soc, V., Part 2. (7) Shore. American Machinist, XXX., Pt. II., p. 747. (8) Bauschinger. Mitteilungen aus d. Mech. Techn. Laboratorium in Miinchen, 1886. (9) Wohler. Zeitschr. f. Bauwesen, 1860—1870. Engineering, XI., 1871. (10) Osborne-Keynolds and Smith. Phil. Trans. Eoy. Soc, CXCIX., p. 265, 1902. (11) Stanton and Bairstow. Proc. Inst. Civil Engineers, CLXVI., 1905—6, Pt. IV. (12) Sankey. Engineering, December 20th, 1907. (13) Arnold. Proc. Inst. Civil Engineers, 154, Supplement, 1903. (14) Charpy. Memoirs Soc. Ingenieurs Civ. Ser. 6, CVIII., 1904. (15) Izod. Engineering, September 5th, 1903. (16) Keport of Committee on Notched Bar Impact Tests, Internat. Testing Association, New York Congress, 1912, Section A. (17) Stanton and Bairstow. Proc Inst. Mechanical Engineers, November, 1908. (18) Coker. Proc. Inst. Mechanical Engineers, 1913. CHAPTER XI THE EFFECT OE STRAIN ON THE STRTJCTITEE OF METALS Having in the last two chapters reviewed in general terms the methods of measuring the resistance of metals to various forms of stress, we have now to consider the relationship which exists between the mechanical properties of metals and their internal structure and constitution. We wUl best approach this subject by seeing how the internal structure of a metal is afEected when the metal is subjected to strain — i.e., when its shape or dimensions are altered as a result of the application of stress. We shall follow the same plan as in earUer chapters of this book and begin by considering the simplest variety of metal from the structural point of view, viz., a simple pure metal, and by considering its behaviour under a simple kind of stress, viz., tension. The obvious way of attacking this question is by the examina- tion, under the microscope, of specimens cut from a previously strained piece of a pure metal, such as a broken tensile test- piece. If this is done for any simple metal, such as the purest obtainable iron, the structure is seen to be very markedly changed from what it was before straining. An example is shown in Fig. 96, Plate XIX., which represents strained Swedish iron. The irregular but approximately equi-axed polyhedral crystals of the original material are found to be replaced by elongated grains exhibiting considerable lengthening in the direction of tension. This altered structure could be accounted for by simply supposing each of the original crystals of the metal to have become elongated in much the same manner as the whole bar has become lengthened during the test. But then the question arises, what are these elongated grains — are they still crystals ? The answer is that they exhibit all the essential features upon which the conclusion is based that the grains of the unstrained metal are true crystals ; there is the uniformity P.M. R 242 STUDY OP PHYSICAL METALLURGY of etching, indicated by the boundaries developed on etching, and the oriented lustre which indicates uniform crystalline orientation throughout each of these elongated areas. Etching- figures can also be obtained which lead to the same conclusion. It would seem, then, that the crystals of a simple ductile metal, when that metal undergoes plastic or permanent extension, are themselves elongated while yet retaining the essential character of crystals. How can a crystal become elongated while yet remaining essentially a crystal ? The fundamental difficulty implied in this question lies in the some- what natural association of the term " crystal " with hard and brittle materials, such as salt or sugar or rock crystal ; this association is, however, purely accidental and due solely to the fact that the crystals most commonly and readily observed are those of brittle substances. The manner in which the constituent crystals of a simple metal undergo extension is conveniently observed by preparing a small test-piece, best, perhaps, in the form of a strip of sheet metal, and polishing and etching one of the flat sides of this piece. Under the microscope it then exhibits the familiar polyhedral structure. By means of a special little straining machine, actuated by a screw, such a thin strip of metal can be readily subjected to plastic extension while its surface is kept under constant observation (^). What is observed in these cir- cumstances is as foUows : so long as the metal remains in the elastic stage, or rather until the yield-point is reached, there is no visible change in the appearance of the surface ; elastic extension produces no visible effects under the microscope. As soon, however, as the yield-point is reached and plastic or permanent deformation occurs, definite changes are seen to take place on the surface observed under the microscope. Across the previously plain white faces of the crystals fine black lines appear, at first in isolated groups, but in steadily increasing numbers until the surfaces of the crystals are cross- hatched with fine black lines. In iron or very mUd steel these lines appear to be curved, irregular and branched, but in such metals as lead, tin, silver, gold, etc., the lines are very straight and regular. The appearance of these lines is illustrated in PLATE XX. Fig. 97. Pig, 98. \_Toface p. 242. PLATE XXI. Fig. ion. ..-*. ■ '¥:'^- -•■' X"' •'.■'■■ "^^::s \ \ ^ -^-^^^ \ »■ _-—'' — -" K^^—"^ — * « t '^"^^^Z!^— - — "^ T "— ^__^-^— „\ i^ - * y » _--^ ^ * '^ ■*m * , >ir . \ f * \ vv^^v \ Fig. 99. Fig. 101. [^'Toface p. 24S. STRUCTURE OF METALS UNDER STRAIN 243 Figs. 97 to 101 . Figs. 97 and 98, Plate XX., show the same field of view, in iron, before the commencement of straining and after a considerable amount of plastic strain has taken place. Figs. 99 and 100, Plate XXI., show the black lines in lead (at 150 and 1,000 diameters respectively), and Fig. 101, Plate XXI., in nickel steel. If the process of straining is carried further, the number and apparent width of the black lines increases until they form a very close network, which becomes confused. As fracture is approached this confusion leads to a general roughening of the surface, which prevents sharp focussing with the microscope, but, even close to a tensile fracture, the lines can be seen in very large numbers crossing one another in various directions. The true nature of these lines is now well understood, and is summed up in the name " slip bands " which is given to them. Their origin lies in t he fa ct t£aFj£hen.-aujCEX§t5;l of a ductile^nietaLisiorcibly altered in shapej._it_adapts itself to thenew^configuration iropiMed u.ponit }>y- a process of slidin g or s lipjaduabu^&ccurs certain _oLJfcs- crystaJlogriaEMc planes A. Before Strain inq. B. AFter Straining. Fig. 102. — Diagram of the Forma- tion of Slip Bands. ._^ ^ The manner in which this occurs is well illustrateoBy the behaviour of a pack of cards, or of a pile of books, when the pile is distorted — the shape of each individual card or book remains unchanged but the shape of the pile is changed by the sliding of the individual cards or books over one another. The process may be made clearer by the diagram of Fig. 102, which is intended to indicate in a very approximate manner the condi- tion of a cross-section of two adjacent crystals before and after plastic straining. The upper sketch represents the unstrained crystals, whose smooth upper (polished) surface is indicated by the line ABC — ^the step at B is the slight difference of level B 2 244 STUDY OF PHYSICAL METALLURGY between adjacent crystals formed as the result of etching ; the boundary between the crystals is represented by the full hne BD , while the potential planes of gliding or sHp, differently oriented in each crystal, are indicated by the dotted lines. After straining, sHp has taken place on some of these slip-planes, and minute steps have consequently been formed in the surface at the points marked s, s, s in the lower figure — these steps will, of course, slope in different directions in different crystals. Seen from above, by normal illumination, these short, steep, sloping surfaces will appear simply as narrow black lines. The reason for this appearance is at once explained by reference to the sketch of the paths of rays of light coming from a micro- scope objective on to such a stepped surface, as shown in Fig. 103. The vertical rays which strike the flat surfaces between the crystals are reflected back directly into the objective, and these surfaces consequently ap- pear bright ; those rays, how- ever, which strike the sloping surface of the step — ^the " slip band " — will be reflected away from the microscope, and the steps will, therefore, send no light into the objective, and will appear black in the field of view. The optical characteristics just described afford a very striking demonstration of the fact that " slip bands " really are of the nature of steps, as indicated in the diagram of Fig. 103. While rays of light falling normally on these bands are reflected outwards, conversely, rays falling obHquely on to such sloping surfaces will, if the angle of incidence be properly adjusted, be reflected into the lens of the microscope ; at the same time, obhque rays falhng on the smooth flat areas between the steps will pass outside the microscope objective. When, therefore, a crystal on whose surface slip-bands have been developed is Fig. 103. — Diagram illustrating the Optical behaviour of Slip Bands. I'LATE XXIl. Fig. 10-1. Fig. 10.5. Fig. 10(). [^Tofacep. 245. STRUCTURE OF METALS UNDER STRAIN 245 illuminated by a beam of oblique light at a suitable angle, the elip-bands should appear as bright lines on a dark background. This is actually the case, as is illustrated in Figs. 104 and 105, Plate XXII., which show the same field of view as seen by normal and obUque lighting. Still more conclusive, if possible, than this optical evidence is the demonstration of the nature of shp-bands which has been obtained by the author by means of actual cross-sections of a surface on which slip-bands had been developed. The difficulty in obtaining such a cross-section lay not alone in the minuteness of the surface features looked for, but also in the fact that, when a specimen is cut and polished, the edges of the surface are always more or less damaged and rounded off. Sharp images at high magnifications, therefore, could not be obtained by simply cutting a previously pohshed and strained specimen through at right angles to the original surface and then poUshing the cross-section thus obtained. The author overcame this difficulty by first coating the surface on which slip-bands had been formed with a thick layer of electro-deposited copper. When subsequently a section was cut through the compound mass — in this case of iron and copper — the boundary on this section, between copper and iron, gave the exact outline of the section of the original surface. When observed at a high magnification, such a section shows the slip-bands as true, though minute, steps. The appearance, under a magnification of 1,200 diameters, is shown in Fig. 106, Plate XXII., and it will be seen at once how the outline there seen corresponds with the outline of the lower figure in Fig. 102. It is of interest to note that the approximate depth of the minute slip-band steps seen in Fig. 106 is ^hJtjo o^ ^^ inch, which is about the length of a wave of sodium (yellow) light {^). If we realise the true nature of this phenomenon of deforma- tion by slip, we see at once that it throws a flood of light on the behaviour of metals under stresses sufficient to bring about plastic strain. If such plastic strain occurs solely by slip, then the truly crystaUine nature of the metal should remain un- altered by the straining process. We have already seen that strained, i.e., elongated crystals, stiU exhibit the essential 246 STUDY OF PHYSICAL METALLURGY characteristics of crystals, so that, broadly speaking, this generalisation is correct. J ^t ther e are a whole series of other phenomena connected Avith the proceWTi^pl'astic strain which T;^^ul3r~be^xtreniely_djS^^5Tor~e'xplain _sagiIactoriIy~OTi any theOTywliich.j:eqmrfid.±he structure of nietaHoremamperfectly ..iSEystalliHtej even when very severely strained. The^rarcum- stance that a metal is hardened by strain, in the sense of having a much higher yield-point and Hmit of elasticity, and even a higher breaking stress, is one of these facts. These circum- stances have led the author to adopt a view put forward in the first place by Beilby (^), to the effect that while plastic deforma- tion — and although we have merely discussed its occurrence in connection with tensile strain, its nature and mechanism is the same whatever the system of forces which have brought it about — takes place by shp on the gliding planes of metaUic crystals, yet that the act of slipping is accompanied by some- thing further. Just as the rubbing action of poMshing produces on metal surfaces a thin layer of altered, amorphous material, so we may well expect that the sUding oyer one an othe r of adjacent shp,suja[aces jgill produce^ajocal disturbance of mole- cjaiaJL-anaoggm^nt. If the slip is sligE5r~then-ilrpfobably happens that this derangement is also sKght and temporary and that the disturbed molecules are stiU able to rearrange themselves pretty much in their original system. In such a case the crystaUine orientation is not at all disturbed. If, on the other hand, the slip has been more pronounced, the resulting local disturbance will also be more far-reaching, a greater number of molecules will be disturbed, and they wiU no longer be able to re-arrange themselves in the old crystalline system. A more or less thin layer of amorphous metal wiU thus be formed on each surface of shp. At first — for a short time — these layers will probably possess a certain degree of mobihty, like the surface film which adjusts itself under surface tension. During this period these layers would act as a sort of lubricant, facihtating further slip on the same gliding planes. After a time, however, when the disturbed molecules have had time to "set " in the amorphous condition, we should have on each plane where slip has taken place a layer of hard, non-plastic, STRUCTURE OF METALS UNDER STRAIN 247 amorphous metal. These would effectually prevent further slip on that particular set of gliding planes, and the crystals would be Hmited, in accommodating themselves to further plastic strains, to sUp on other surfaces which had not been " used " in the previous straining. But all the surfaces of easiest slip will have been used at the first straining, so that to effect plastic deformation the second time more force would be required. Not only this, but the hard and brittle amorphous films on the surfaces of previous slip would also act £is a stiffen- ing skeleton for the whole crystal and thus offer additional resistance to the commencement of fresh slip. This conception, therefore, affords a rational and simple explanation of the apparent hardening of metal by plastic deformation. The essence of the whole conception lies in this — that plastic deformation can only occur on the sHp-planes of a true crystal, so that plasticity is, in metals, a property bound up with crystalline structure. As amorphous layers are formed, the capacity of the crystals to yield by slip is diminished, and ultimately, when a certain not very large proportion of the metal has been converted into the amorphous condition, the metal has been rendered incapable of further plastic deforma- tion — ^it has been rendered hard and brittle. As this stage is reached when the metal still consists largely of portions of crystals embedded in amorphous layers, it is obviously im- possible to convert the whole of the metal into the amorphous condition, because the straining process whereby this conversion can be brought about is rendered impossible by the skeleton of hard amorphous matter already present. There is one difficulty about the explanation of the hardening of metal by plastic strain, as given by the conceptions explained above, which requires brief mention. This lies in the fact that the hardening produced by plastic strain is, in certain cases at all events, uni-directional. For instance, a piece of steel which has been "hardened," i.e., has attained a raised peld-point and apparently raised elastic limit, as the result of tensile over- strain, is not really hardened in every way, for, if it is tested in compression, it is found that for stresses of that kind the apparent elastic limit has been lowered, so that the total range 248 STUDY OF PHYSICAL METALLURGY of elasticity, from the limit in compression to the limit in tension, has not been materially altered. It is difficult to see how the formation of hard amorphous layers on surfaces of shp can account for a softening of the metal in relation to compression while producing hardening as against tension. It may be, of course, that there is something directional about the manner in which the crystalline arrangement is broken up on the sUp surfaces, so that a configuration results which allows of sUp in one direction but not in the other. A more probable explanation has, however, been put forward by Humfrey (*), who suggests that amorphous layers present or formed in the crystal boundaries act in such a way that, when distorted by plastic strain, they serve to assist the action of any forces tending to restore the metal to its original shape, but lend added resistance to further change of shape in the same direction as previous strain. This explanation wiU, however, be better understood when the general question of the structures at crystal boundaries has been discussed. The conception of an amorphous and temporarily mobile layer formed on all surfaces of internal sUp affords an interesting explanation of phenomena connected with over-strain, more particularly observable in iron and steel. These are the phenomena of semi-plasticity, which occur immediately after a tensile test-piece has been exposed to a stress beyond the yield-point. If the load is removed from such a test-piece and small loads are gradually applied to it, extensometer observations at once show that the material is in a pecuHar condition. Even under sUght loads the proportionaHty between load and extension is no longer observed ; the test- piece shows a certain gradual creep, the extension slowly increasing with time. If the small load is removed, the metal does not at once return to its original length, but a small more or less permanent " set " is left behind, although this may diminish gradually as the result of backward creeping. The somewhat complex phenomena may be summed up by saying that the material has for the time lost the truly elastic character which it normally exhibits under such small loads, and behaves as if it were in a semi-plastic condition. If left to itself, the original STRUCTURE OF METALS UNDER STRAIN 249 perfect elasticity is recovered after a time which varies from a few hours to several days, according to the nature of the material, while exposure to a temperature of 100° C. for twenty minutes is sufficient to bring about complete " recovery " (®). Such phenomena could only be explained with great difficulty if purely crystalline shp were alone concerned in plastic defor- mation. The conception of amorphous layers possessing tem- porary mobility, however, affords a simple explanation. So long as the amorphous layers formed on the shp surfaces retain their temporary mobility, the metal may well be expected to behave in a semi-plastic manner, since the amorphous layers on the slip surfaces would behave like films of viscous fluid, slowly moving under the appUed stresses, thus giAring the slight permanent " set " and the gradual creeping observed. In the course of time, or more rapidly in consequence of shght warming, these layers become rigid in the permanently hard, amorphous condition, and the phenomena of semi-plasticity accordingly disappear. Having freely availed ourselves of the conception of " amor- phous " metal, it may be well to state precisely what is under- stood, in this connection, by the term " amorphous." In accordance with Beilby's views, and their recent development by the author and his collaborators (®), the amorphous condi- tion is one in which the crystalline arrangement of the molecules is completely broken up, so that the molecules remain in a state of irregular arrangement similar to that which is supposed to exist in the hquid state. In fact, the " amorphous phase " is regarded as being — ^from the point of view of the phase doctrine — ^identical with the hquid phase. Now extremely under-cooled hquids are well known in such substances as glass, vitreous silica, etc., and they are hard, brittle bodies devoid of plasticity and of crystalline structure. Amorphous metal is, therefore, to be regarded as being identical in nature with the Hquid metal if that could be cooled down to the ordinary temperature without undergoing crystalUsation. It is admitted that such extreme under-cooling of metals has never been actually accompHshed, and objections to the whole concept of amorphous metal have been baaed on that ground. 250 STUDY OF PHYSICAL METALLURGY These cannot be discussed here, but they have been fully dealt with by the author elsewhere C). At the present moment it may well be claimed that the theory of an amorphous phase in metals, produced as the result of mechanical dis- ruption of the crystalline arrangement of the molecules, holds the field as the best working hypothesis available. There are admittedly outstanding difficulties, but these the author regards as being inevitable in a theory whose wide generalisations cover a field which has as yet been but partially explored (^). The general effect of slowly- appfied or " static " plastic deformation, which has just been discussed, may now be summarised as follows : the metal in the cast or the annealed state is an aggregate of crystals ; when the metal undergoes plastic deformation the crystals are deformed in the same general sense as the mass of metal. This change of shape of the crystals is accomplished by a process of sHp or gliding by which layers of the crystal sfide over one another along certain gliding planes. Such slip occurs on a great multitude of planes in each direction and, as a rule, it must occur on at least three sets of planes, but it may occur on a greater number. This slip is accompanied by a certain disturbance of the crystalline arrangement of the molecules on or near the surfaces of sHp. If the deformation has been slight, the disturbed molecules may be able to rearrange themselves in accordance with the crystalline system by which they are surrounded, and the completely crystalline character of the metal will then be unaffected by the deformation. When the deformation is severe, and if it is intensely localised, as when a single crystal is bent upon itself, the molecular disturbances at the slip surfaces become larger, and permanent layers of amorphous metal are formed on each surface where slip has occurred. Ultimately the metal, as the result of extreme deformation, becomes a mass of crystal fragments embedded in relatively thick layers of amorphous metal. This change is accompanied by an increase in the hardness of the metal and — ^in the majority of metals — ^by a measurable decrease of density. Reference must be made here to another process which is PLATE XXIII. Fig. 107. Fig. lOS. ITofiicep. 2.51. STRUCTURE OF METALS UNDER STRAIN 251 known to occur in the plastic deformation of crystals and which occurs in metals, although it is doubtful whether it is the direct result of strain or only arises during a subsequent anneaHng process. This is the formation of " twin " lameUse, or " twin- ning," as it is generally called. The mutual relationship of twinned crystals is best described by saying that in their internal arrangement and symmetry they are mirror images of one another ; we can conceive of the twinned portion being generated from the normal structure by a simple swinging round of all the crystal elements through a definite angle, which is, however, intimately connected with the natural angles of the crystal. In the mineral calcite this swinging round can be brought about by mechanical means, such as the forcible insertion of a knife blade into the edge of the crystal at a suitable point. In metals the presence of twin lamellse makes itself evident in the etched micro-sections of many metals, such as copper, brass, silver. An example of their appearance, with its typical rectilinear and parallel banding, has already been given in Fig. 55, Plate XI. ; another very beautiful example, taken from pure silver, is shown in Fig. 107, Plate XXIII . ) . The existence of twinned regions is also revealed by the shape assumed by slip-bands in passing through such twinned lamellse ; this is illustrated in Fig. 108, Plate XXIII., which relates to strained copper. The view has been widely held (®) that twinning in metals occurs directly as a purely mechanical action under the influence of strain ; recent experi- ments have, however, shown that in the majority of metals, if not in all, twinned crystals are never found in the cast material, if it has never been strained in any way, and that the apphca- tion of strain to such cast metal does not at once produce twinning ; even the very shght amount of strain involved in gently poKshing the surface of a piece of pure cast silver is sufficient to produce most vigorous twinning when next the metal is annealed {^°). This evidence tends to the conclusion that twinned bands are not formed during straining, but that they are a form which the crystals of many metals assume very readily during the re-crystaUisation which occurs on anneahng after straining. It is interesting to note that typical twinning 252 STUDY OF PHYSICAL METALLURGY is not seen in a iron, but that 7 iron exhibits twinning in a marked degree. We have now to consider the mechanism of fracture, first when it results from the application of a steady load and then in relation to the phenomena of " fatigue " described in con- nection with alternating stress testing in the previous chapter. It wiU be seen at once that the mechanism of fracture wiU be different according as it is brought about by a gradually applied load, by shock, or by alternating stresses. Taking steadily-appMed tension as our example, wftfind that the crystals become increasingly elongated with increasing extension, slip pcoum ng to an increasing extent on jnore and more numerous gliding planes. t)n these planes, too^^amoiphgus_layers are formed, j)Jit the amorphous material is stiU in its temporarily mobila_c©nditiori. ^^TJItimately the proc§g^.,^rsli£[^^reacbes a limit, ^nd the surfaces part and the test;£iece breaks? When a fracture thus.produced js examined, it . .exhibits„a , typical "fibrous" appearance, but the fibres, in a pure metal, are merely the drawn-out ends of the elongated crystals, and can be fecOgnised as such under the microscope. This can be readily done if the fracture is embedded in electro-deposited metal, such as copper, and a longitudinal section is then cut. Such a section, showing the broken ends of the elongated crystals, is given in Fig. 109, Plate XXV. When shock acts upon a pure metal, the result is somewhat different. Although a certain amount of deformation by shp takes place, that is essentially a process requiring time, and shock fracture does not allow such time. Consequently, instead of undergoing deformation by sHp on the gliding planes of the crystals, the metal undergoes fracture along their cleavage planes. The process of fracture under steady load and under shock have this in common, that in both cases fracture results across the body of the crystal — the individual crystals are not torn apart from one another along their boundaries. This is a typical feature of the fracture of all "sound" materials ; where inter-crystalline fracture occurs it is a certain sign of some abnormahty in the composition or condition of the metal. Yet the mode of fracture under shock results in an appearance so STRUCTURE OF METALS UNDER STRAIN 253 different from that under steady load that the suggestion has been made that metallic crystals may exhibit either a " duc- tile " or a " brittle " behaviour. Both are, however, merely manifestations, under differing conditions, of their essentially crystalline constitution. The explanation of the process of failure by " fatigue," i.e., under the repeated alternations of a stress which would not cause fracture if steadily apphed, is also furnished by the con- ceptions described above as to the behaviour of a crystalline aggregate under strain. A stress which is to cause ultimate failure after repeated alternations must be large enough to produce a small amount of local yielding in the metal. This may be so small in amount as to be unobservable, even with a delicate extensometer, and in that case the stress would be regarded as lying within the apparent or " primitive " elastic limit, but the microscopic examination of polished test-pieces under load has shown that the formation of sUp-bands in isolated crystals here and there in the metal may and does occur for stresses of this kind. Some crystals, by their shape and the orientation of their gUding planes are unfavourably situated to resist the particular system of stresses which has been applied, and a sHght local slip takes place. If the load remains in steady action, nothing further occurs. If, however, the stress is reversed — i.e., if the metal is being subjected to alternating stress — ^then this slight amount of slip will also be reversed, particularly as the slip surfaces will stiU be covered with the temporarily mobile layer of amorphous metal. Such reversal will be repeated with each reversal of the applied stress, and at each successive slip the layer of amorphous material will be increased. After a time, however, by virtue of its temporary mobiUty, this film of quasi-hquid metal will be squeezed out between the gliding surfaces, and the site of the initial minute sUps will develop into a fine crack. As this process continues, that particular crystal soon begins to lose its strength, and additional stress is thereby thrown upon its immediate neighbours, which undergo slip and gradual disintegration in the same way. Ultimately, the crack or flaw thus originated works its way across the entire section of 254 STUDY OF PHYSICAL METALLURGY the metal, and failure of the piece results. If the action has been fairly rapid, i.e., if the stress was high enough to produce somewhat rapid disintegration by repeated reversals of slip, the resulting fracture exhibits the crystal faces upon which sKp has taken place as a number of bright facets resembhng those produced in a " brittle " shock fracture, and it is this appearance which has led to the mistaken idea that alternating stresses cause metal to " become crystalline." Actually, as we have seen, the metal is crystalline from the beginning, never really " fibrous," and the " fibrous " or " crystal- line " appearance of the fractures depends on the mechanism of fracture, and not on any change in the crystalline structure of the metal. The explanation of fatigue fracture which has just been given, although it has been deduced from the general character of plastic deformation by slip, was first given by Ewing and Humfrey {"), as the result of direct experimental observations in which they watched the formation of shp-bands in certain crystals, and their gradual growth into cracks on the polished and etched surfaces of pieces of Swedish iron submitted to the Wohler test. One set of photo-micrographs obtained by them in this way is reproduced in Figs. 110 to 113 inclusive, Plate XXI. The manner in which fracture under comparatively low alter- nating stresses takes place affords an insight into the true mean- ing of the parabolic curve obtained by the Wohler or other fatigue test. With the higher stresses, fairly rapid failure results, because slight shp takes place in a considerable number of crystals, and this slip is also greater in amount than it would be under lighter stresses — ^the exact relation between the two is, however, dependent upon certain properties of the crystals with which we are not well acquainted. On the other hand, when the stress has been reduced to such a value that no local slipping occurs anywhere in the metal, then an indefinite number of reversals produces no effect on the metal — i.e., the test yields a point on the horizontal part of the parabolic curve. The limiting stress at which slip ceases to occur within the metal is thus the " safe range " for alternating stresses, and its physical meaning is simply that it is the " true ' ' PLATE XXIV. Fig. 110. Fig. 111. ^■: Fig. 112. Fig. 113. [To fiiiv p. 2oi. STRUCTURE OF METALS UNDER STRAIN 255 elastic limit, since nothing but purely elastic deformation is produced thereby. Actually, the alternating stress test is probably the most sensitive method of determining this true elastic limit, and it has the advantage that it is unaffected by the particular condition of cold work in which the metal may find itself, as the application of alternating stresses rapidly restores the elastic limits in both tension and compression to their " true " or normal positions. While it cannot be claimed that all the intricate phenomena connected with the failure of metals under vibratory or alter- nating stresses have been fully and finally explained, the methods of Physical Metallurgy have so far elucidated the problem, in the manner indicated above, that a clear insight into the mechanism of failure under the simpler forms of alternating stress has been gained. The myth that metals " become crystalline " under the influence of vibration has been finally dissipated by the demonstration of the manner in which the yielding of crystals by slip can lead to failure under alternating loads. Indeed, microscopic study has shown very conclusively that, so far as iron or steel is concerned, no perceptible change of crystalline structure or arrangement ever occurs at the ordinary temperature. Quite recently. Garland (^*) has shown, by means of micro-sections of specimens of non-ferrous metal and alloys taken from ancient Egyptian tombs, that the meta-stable cored structures of cast soHd solutions, as well as the distorted crystals produced by cold work, persist unchanged through periods of thousands of years. The author's examination of some specimens of ancient iron from Adam's Peak, in Ceylon (^^), and similar studies by Sir Robert Hadfield ("), have shown that in pieces of iron probably over 4,000 years old the structure is exactly the same as in iron that has been made yesterday. It is well, therefore, that the modern explanation of "fatigue" pheno- mena rests on something much more secure than the unfounded supposition that vibration can cause changes of crystalline structure in a metal so far removed from its softening tem- perature as iron or steel, or even brass or bronze. In the discussion of the behaviour of metals under strain 266 STUDY OF PHYSICAL METALLURGY we have so far confined our attention, not only to pure or '' simple " metals, composed of an aggregate of crystals of one kind, but we have dealt only with the behaviour of the material forming the mass or interior of a crystal and not at all with the conditions which exist at the crystal boundaries where adjacent crystals meet. Yet the behaviour of the crystal boundaries is of primary importance to the quahties of the metal. The extent to which this is the case appears in a striking manner from some early experiments of Arnold (^^) with an alloy of gold and bismuth. He showed that such an aUoy consists of perfectly ductile crystals of pure gold surrounded — ^for aUoys of low bismuth content — ^with a thin brittle film in which the bismuth is concentrated. The presence of these thin brittle films in the boundaries is sufficient to render the whole mass of metal extremely fragile, so that it can be readily broken up with a hammer. Individual crystals, however, after being thus broken apart can be hammered out into thin sheets. This is a striking case of a condition which is really abnormal, viz., of inter-crystaUine brittleness, due in that case to the presence of a deleterious impurity. In all metals and alloys it is now well known that such inter-crystaUine brittleness only arises when the crystal boundaries are weakened, either by the presence in them of a weak and brittle constituent, or because the inter-crystaUine cohesion has been weakened by the effects of some special treatment to which the material has been subjected. It has been clearly shown that in all normal fractures, both of iron and steel and of other metals and alloys, even including such a brittle metal as bismuth, the path of the fracture, whether produced by static stress, shock, or fatigue, never follows the inter-crystalline boundaries, but always crosses the crystals, sometimes on relatively straight surfaces, as in cleavage fracture produced by shock, sometimes on greatly stepped sHp-surfaces, as in a tensile frac- ture. An illustration of this fact, which observation of numbers of test-pieces has confirmed, is given in Fig. 114, Plate XXV., which is the section of a shock-fracture of a very mild steel, consisting almost entirely of ferrite. This fracture may be con- trasted with that shown in Fig. 115, Plate XXV., which is from PLATE XXV. Fki. 109. Fifi. l\-, Pig. 114. Fig. 115. [Tufacej). 2.i6. STRUCTURE OE* METALS UNDER STRAIN 257 an abnormal steel in which inter-crystalline brittleness is evident. If we regard cohesion within the body of a crystal as being due to the attractions between layers of adjacent mole- cules, we can readily understand the continuity of such cohesion acting throughout the entire mass of any one crystal. The forces which are at work in producing cohesion between adjacent crystals must, however, be of a somewhat different character, for it is obvious that the regular arrangement of molecules in oriented layers cannot be carried on through a crystal boundary, while it appears that the actual cohesion between adjacent crystals is stronger than that between different layers of the same crystal. The view, at one time widely held, that a considerable number of minute cavities exist where adjacent crystals meet in such a way that they do not "fit in " with one another, is obviously contrary to the observed fact that the boundaries are surfaces of extra strength and not of weakness. During the past three years, a view has been independently put forward by a number of workers (^*) according to which the interstitial spaces between adjacent crystals are fiUed up by an amorphous film or layer whose nature is similar to that formed on poUshed surfaces or on surfaces of slip, but whose origin is entirely different, at any rate so far as cast or annealed metals are concerned. The basis of this view is that where adjacent crystals meet there is a kind of " region between the orientations " where the molecules of the crystallising liquid, or of the re-crystalhsing soUd, are left in a state of approximate balance between the opposing tendencies to arrange themselves according to one or other of the adjacent crystalline orientations. In conse- quence of this approximate balance of directive forces they remain in the irregular arrangement characteristic of the liquid state, and persist as thin layers of greatly under-cooled liquid, i.e., of amorphous material, which is readily able to accommodate itself to the irregularities of surface of both of the adjacent crystals, and thus, by bridging the gap in the chain of cohesion, acts as an inter-crystalline cement. This view, although widely supported, is also strongly P.M. s 258 STUDY OF PHYSICAL METALLURGY opposed by those who beheve that the tendency to crystallise is so extremely powerful in metals that any considerable under- cooling of the kind suggested is impossible {^). No satisfactory alternative theory, however, has yet been put forward, and the " amorphous cement theory, " at all events holds the field as a satisfactory working hypothesis. How far the picture of what occurs at the crystal boundaries really represents the truth remains for further research to determine. The theory — or hypothesis — as it stands, however, serves to explain a whole series of interesting phenomena. Some of these have only been discovered because the author and his collaborators were led, by the indications of the theory in question, to look for them experimentally. The most interesting of these from the present point of view relates to the behaviour of metals at high temperatures (*), and the facts connected with it are well worth careful consideration. At the ordinary temperature, when a specimen of metal is strained, having previously been polished and etched on a convenient surface, we have seen that the crystal surfaces become covered with shp-bands, as illustrated in Fig. 98, Plate XX. This change of shape of the individual crystals by means of sHp cannot, however, go on without some move- ment at the crystal boundaries. In addition to changing their own shapes, the crystals must move relatively to one another. The strength of the crystal boundaries is found to resist such movement, and the slip-bands are found to be arranged in such a way as to minimise the amount of displace- ment occurring at the actual boundaries. In other words, the metal takes up the new shape imposed on it, as far as possible, by means of shp within the crystals and with as little disturbance as possible of the inter-crystalline boundaries. That there is some definite movement at the boundaries becomes evident if a pohshed specimen is strained without being previously etched ; the effect of the strain at once causes the crystal boundaries to become visible on the surface. As the temperature of a metal is raised, it is found that in general its strength — i.e., its resistance to shp and to fracture — becomes rapidly diminished ; at the same time, the surface PLATE XXVI. Fl(i. IKJ. Fig, 118. Fig. 119. Fig. 120 [ ?b face p. 2.59. STRUCTURE OP METALS XJNDER STRAIN 259 phenomena which are seen on previously polished surfaces also undergo a change. This — according to the amor- phous cement theory — ^is due to the fact that while the resistance to slip, being of the nature of solid friction, is not very much affected by rise of temperature, the resistance of the boundaries — ^being of the nature of the resistance of a very viscous fluid — ^rapidly diminishes with rising temperature. While at the ordinary temperature the boundaries are much stronger than the bodies of the crystals, a point is reached, with rising temperature, when this relation is reversed, and the boundaries yield more readily than the mass of the crystal, Specimens strained at such a high temperature — ^in nearly pure iron this lies in the neighbourhood of 900° C. — show no great development of slip-bands, although signs of their occurrence are still visible, but there are clear signs of much movement at the crystal boundaries. This is illustrated in Fig. 116, Plate XXVI., which shows the previously pohshed surface of a specimen of very mild steel which has been strained in a high vacuum at a temperature of 1,000° C. That this phenomenon is really due to a viscous yielding of the crystal boundaries is shown by two distinct facts. One of these is that at such temperatures the resistance of the metal is dependent upon the rate" of straining according to a law which is that of viscous resistance ; the second is that, if a specimen of the same material, heated to the same temperature, is strained by the sudden application of the load, not only is the strength apparently much higher, but the surface phenomena and the fracture revert to the low- temperature type — i.e., there is again a large development of slip-bands and little movement at the boundaries. This — according to our theory — simply means that the rate of deformation has been so great that the fluid, but still highly viscous, cement between the crystals has not had time to flow, and has accordingly behaved like a hard soMd, much as does pitch under the action of a blow. This hne of evidence has, however, been pushed considerably further. If the inter-crystalline cement becomes softer as the temperature is raised, at some point not very far below the actual melting-point of the metal, the cement should be s2 260 STUDY OF PHYSICAL METALLURGY practically fluid in the more usual sense of the word — i.e., have only a low degree of viscosity ; at the same temperature the crystals themselves should still possess considerable strength, and, consequently, it should be possible to pull the crystals apart, by the aid of a small load, without appreciably deforming them. This leads to the conclusion that at a temperature not far below its melting-point even the most ductile metal should exhibit a very high degree of inter-crystalline brittleness. This somewhat unexpected conclusion has been very completely verified by the author and D. Ewen (*), who produced perfectly brittle fractures of the purest lead, tin, gold and bismuth by means of a slight pull applied to the metals at temperatures ranging from 3° to 10° C. below their melting-points. This question, with its many interesting points of outstanding difficulty, cannot be further pursued here, particularly as many of the fundamental points are still under active discussion on the part of those interested in the investigation of this subject. On the value of the various arguments employed, the reader — if he is sufficiently interested in the matter — ^will form his own opinion on reading the various papers referred to in connection with the present chapter. At this point we must pass on to other aspects of our subject. While we have so far confined our consideration of the behaviour of a crystalline metalUc aggregate under strain to the case of a simple metal, consisting entirely of crystals of one kind or phase, a very large number of metaUic materials em- ployed in the arts and industries are of duplex constitution, i.e., consist of aggregates of at least two kinds of crystals. We have now to consider how such a duplex aggregate behaves when deformed. It must first of all be understood that in a duplex, or even in more complex metallic aggregates, all the individual phases or constituents exist as distinct crystals. Where we have two solid solutions, such as the a and /8 phases in brass, existing side by side, this conclusion is readily established by examining each of the constituents, when each is found to be character- istically crystalline. Where a typical eutectic or eutectoid is present the case is not quite so simple, as we then have an STRUCTURE OP METALS UNDER STRAIN 261 aggregate consisting of simple homogeneous crystals co- existing with grains which are themselves finely laminated and duplex. Yet these eutectio or eutectoid grains have been shown to possess definite crystalline structure of a certain type (^'). This type is, however, decidedly less regular in its arrangement than a simple metal crystal, so that the occurrence of slip-bands of the ordinary kind is not frequently observed in eutectics. In fact, if a specimen of a pure eutectic alloy is provided with a polished surface and is then plastically strained, it will be found that the effect on its surface appearance is almost identical with that obtained by etching — the laminated structure is very clearly revealed (^). Close examination has shown that what really happens is that slip occurs along, or very close to, the boundaries of the lamellaa of the two con- stituents present. When it is borne in mind that these lamellae are formed by the process of crystallisation, it will be seen that their surfaces must he on or near to some of the principal crystallographic planes, so that it is not really surprising that slip should occur on these surfaces. When the eutectic structure is somewhat coarse and the applied strain is severe, the formation of ordinary slip-bands can be induced, and these always bear some definite relation to the lamellae of the eutectic. Admitting, then, that both constituents of a duplex alloy are essentially crystalline, the duplex metal still differs funda- mentally from the simple metal, owing to the fact that, while in the one the various crystals differ from one another so far as resistance to slip is concerned, only on account of their varying orientation, in the duplex metal we have two kinds of crystals which differ fundamentally in their inherent power of resisting sHp or of allowing it to take place to any appreciable extent without undergoing fracture. Here we may recall the circumstance that in a large number of duplex aUoys one of the constituents either consists of, or contains, a definite inter- metallic compound, and that such compounds are characteristic- ally brittle — i.e., incapable of undergoing slip without suffering fracture. Perhaps this may be expressed by saying that in these compounds all the ghding planes are essentially cleavage 262 STUDY OF PHYSICAL METALLURGY planes, but the important fact is that, as a consequence, one of the constituents of a duplex alloy is generally hard and brittle. The condition of a duplex alloy under strain may thus be regarded as very similar to that of a simple metal into whose structure a certain number of hard, brittle grains have been introduced. The influence of such hard interspersed bodies can be readily understood. In the first place, their presence tends to harden and stiffen the whole alloy to an extent quite disproportionate to the quantity of hard material actually present. This is due to the fact that the ductile crystals of the " matrix " metal are supported by the adjacent harder crystals, and in that way the beginning of the slipping process is retarded until the stiffening effect of the hard constituent is overcome by the application of a larger stress. This is equivalent to saying that the duplex material has a higher yield-point and a higher ultimate strength, since the stiffening support of the harder crystals will be continued to the end of the straining process. Further, if the hard constituent is not present in too large a proportion, and if its individual crystals are small and evenly distributed, they in their turn wiU be supported by the softer material in which they are embedded, and wUl thus be able to undergo a certain amount of deformation — by sUp — without undergoing immediate fracture. That such supporting action is of importance is well known ; thus crystalline Mme- stone, which is ordinarily quite brittle, can be made to undergo considerable flow if it is confined in a tube under great hydro- static pressure ; similarly bismuth can be extruded under pressure in the form of a wire, although isolated crystals are extremely brittle. We have, thus, in a duplex alloy a double system whose constituents react upon one another, each supporting the other and aiding it to resist failure in its own particular way. Such a system, however, is necessarily sensitive to the exact arrange- ment of its parts. Now in an alloy of this kind the arrangement of the two constituents — -the size and shape and relative arrangement of their crystals — ^is dependent upon both me- chanical and thermal treatment, so that it is not surprising to find that the physical properties of duplex alloys, more than those STRUCTURE OF METALS UNDER STRAIN 263 of other types of metal, are largely dependent upon the treat- ment — thermal and mechanical — which they have undergone. One of the most important and typical of these duplex alloys is, of course, carbon steel, and the influence of the duplex structure is readily studied in this case by the examination of cross-sections of fractures obtained in the manner described above (see p. 245). The path of a tensile fracture is illustrated in Fig. 117, Plate XXV. It wiU be seen that apparently the fracture passes indifferently through both ferrite (light) and pearhte (dark), in spite of the fact that the tensile strength of a steel consisting entirely of pearUte is some two and a half times that of a pure ferrite material. The reason for this peculiar feature lies in that close inter-relation of the two constituents which has just been discussed. As the yield-point of such a steel is passed, the ferrite, in spite of the support of the pearhte, begins to yield, and the pearlite itself undergoes a small amount of plastic deforma- tion. This process continues until the limit of endurance of the pearlite is reached, when, even with the support of the surrounding ferrite, it can withstand no further stretching. The elongated pearlite crystals then break, leaving the ferrite to stretch further unsupported, finally breaking in its turn ; this final break, however, is bound to occur in such a way that the resulting fracture will run through the fissures already existing in the pearlite. The result is a fracture which appears to run impartially through both constituents. The signs of what has taken place are, however, to be found in the frequency with which the pearhte grains in the neighbourhood of the fracture are found to be fissured, sometimes in several places. When the same steel is broken in other ways, a different type of fracture results, but in general terms these fractures may be divided into two groups, according as the fracture is or is not preceded by serious plastic deformation. Wherever there is much antecedent deformation the fracture is of the same type as that already described in connection with tension ; on the other hand, if the steel is broken without previous deforma- tion, as by a sudden blow on a notched bar, the fracture appears to occur by the direct cleavage of the weaker and more perfectly 264 STUDY OF PHYSICAL METALLURGY crystalline constituent, i.e., the ferrite, and carefully to avoid passing through the pearlite. Thus alternating stress fractures follow the ferrite, as shown by Stanton and Bairstow ; a shock fracture has been shown in Fig. 114, Plate XXV. Eefebences. (1) Ewing and Eosenhain. Phil. Trans. Roy. Soc, 353a, 1900. (2) Beilby. Numerous papers, including : — Proc. Eoy.Soc, 72, May, 1903; 76, 1905 ; 79, 1907; 82, 1909- Journ. Soc. Chemical Industry, 1903, November. Faraday Soc, June, 1904, Phil. Mag., August, 1904. Smithsonian Eeport, August, 1905. Journ. Inst. Metals, VI., 2, 1911. (3) Eosenhain. Proc. Eoy. Soc, Vol. 74, February, 1905. Journ. Iron and Steel Inst., 1906, II. (4) Humfrey. Journ. Iron and Steel Inst. (Carnegie Memoirs, 1912). (5) Muir. Phil. Trans. Eoy. Soc, 1899. (6) Eosenhain and Ewen. Journ. Inst. Metals, VIII. 2, 1912, and X., 2, 1913. Eosenhain and Humfrey. Journ. Iron and Steel Inst., 1913, I. (7) Eosenhain. Engineering, October, 1913. Internat. Zeitschr. Metallographie, V., 1914. (8) Heyn. Eeport on Progress of Metallography from 1909 to 1913. Internat. Testing. Assoc, New York Congress, 1912. Tammann. Zeitschr. f. Elektrochemie, XVIII., July, 1912. Guertler. Internat. Zeitschr. f. Metallographie, V., 1914. (9) Eosenhain and Ewen. Journ. Inst. Metals, VIII., 2, 1912. Heyn. Discussion on Eosenhain, Journ. Iron and Steel Inst., 1904, I. (10) Eosenhain and Ewen. See (6), above. (11) Ewing and Humfrey. Phil. Trans. Eoy. Soc, 200a, 1902. (12) Garland. Journ. Inst. Metals, X., 2, 1913. (13) Hadfleld. Journ. Iron and Steel Inst., 1912, II. Eosenhain. Journ. Faraday Soc, 1913. (14) Hadfleld. Journ. Iron and Steel Inst., 1912, II. (15) Arnold. Discussion on Fourth Eeport Alloys Eesearch Cte., Proc. Mech. Engs., February, 1897. (16) Brillouin. Annales de Chimie etde Physique, 1898 (VII.), XII., p. 377 ; 1896, VI., p. 540. Sears. Trans. Cambridge Phil. Soc, XXI., p. 105. Bengough. Journ. Inst. Metals, No. 1, 1912, VII., p. 176. Eosenhain and Ewen. See (6), above. Eosenhain and Humfrey. See (6), above. (17) Eosenhain and Tucker. Phil. Trans. Eoy. Soc, 209a, 1908. Vogel. Zeitschr. Anorg. Chemie., 1912, LXXVI., p. 425. Guertler. Internat. Zeitschr. f. Metallogr., II., 90, 1912, CHAPTER XII THE THERMAL TREATMENT OF METALS In the previous chapter the manner in which the structure of a metal behaves under strain has been discussed, and we have seen how, as the result of the process of internal slip accom- panied by the formation of amorphous metal in increasing quantities with increasing strain, the original structure of a cast or annealed metal is modified by the application of " cold work," i.e., by changes of shape produced by the application of stresses to the metal in the cold state. The corresponding effects when metal is subjected to " hot work," i.e., when changes of shape are produced in the metal while it is hot, will be discussed in the later portions of the present chapter, but, before we can deal satisfactorily with that branch of our subject, we must first consider the process known as " anneal- ing." Essentially the anneahng process is the converse of the process of straining, since it results in undoing, to a certain extent, the effects produced by the latter. The changes of shape which have been produced by strain are not, of course, undone when the metal is subsequently annealed, but, so far as the internal condition of the metal is concerned, the result of annealing, if properly completed, is to restore the metal to the perfectly crystalline condition which we have learnt to regard as normal, and thus to remove the hardening effects of strain. The annealing process consists, in all cases, in exposing the metal to a sufficiently high temperature and following this by comparatively slow cooling. In most cases the exposure to a high temperature need not be more than sufficient to aUow the whole mass of the metal to attain the same temperature, but if it is desired to obtain an annealing effect at moderately low temperatures, much longer exposure is required. The details 266 STUDY OF PHYSICAL METALLURGY in regard to this point vary so much that reference must be made in each case to works dealing ^vith the individual metal in question. Here we can only consider one or two typical cases. Before considering the properties of any particular metal in regard to annealing, we must first define somewhat more precisely what is meant by such a term as " anneaUng tem- perature." When a cold-worked piece of metal, such as a hard-drawn wire or a rolled or hammered sheet, is heated there seems to be no doubt that the first notable effect is that of mechanical softening. Whether we measure this softening by some form of hardness test, such as a ball or cone impression, or by the " scleroscope," or whether we determine the reduction in tensile strength of a wire, it is always found that, for a given period of heating, there is some fairly definite minimum tem- perature at which softening occurs. If the micro-structure of the strain-hardened material be examined after various intervals of such heating, however, it is usually found that the effects of annealing upon the structure do not become visible for some little time after the first definite softening effect has been produced. This is quite in accordance with the " amor- phous " theory, since the first effect of raising the temperature would be to cause the re-crystalhsation of the amorphous material present in the metal — so far as such re-crystallisation is possible without a re-arrangement of the whole crystaUine structure. If strain has not been too severe, the various layers of the original crystals will still be correctly oriented relatively to one another and, when the intervening amorphous layers are caused to resume the crystalline arrangement, this will result in the practical re-constitution of the original crystals, but these will retain the elongated or otherwise distorted shapes which they assumed under plastic deformation. If the strain has been severe, however, the original orientation of the various layers of a crystal will no longer be preserved, and this reconstitution of the former crystals will only occur to a very small extent, the resulting structure being a mass of extremely minute crystals separated by their boundary films of amorphous metal. In its mechanical properties such a material will differ THE THERMAL TREATMENT OF METALS 267 from the fully re-crystallised material only in certain special directions ; as compared with the strength or hardness of the material before annealing had begun, its strength and hardness will be small and practically identical with that of the fully- annealed metal. Even though the amorphous boundary layers are numerous in this state of the material, their total cross-sectional area is still small, and their influence on the hardness or tensile strength is small compared with that which the amorphous layers in the slip-surfaces of the crystals had produced in the strain-hardened metal. It thus arises that the mechanical softening effect of anneahng appears to precede the visible effect on the micro-structure. In practice, also, the action is considerably complicated by the fact that the various stages of the annealing process do not occur uniformly throughout any mass of metal. Somewhat more prolonged heating, after the first mechanical softening effects have been obtained, results in the rapid formation of a large number of minute crystals, which follow the universal tendency of the constituents of a crystaUine aggregate to increase in size and to decrease in number. Accordingly some of these crystals grow at the expense of their neighbours. On the hypothesis of an amorphous inter- crystaUine cement, this process simply means that the structure of the metal changes in such a manner that the total quantity of amorphous metal — which is essentially in an unstable condi- tion and, therefore, tends to revert to the stable crystalline form — ^is diminished in consequence of the diminution of the total area of inter-crystalline boundaries in the piece of metal. This tendency to grow is fairly rapid at first, but diminishes with increasing size of crystals ; \ it has been found b^ Sauveur (^), however, that in mild steel, at aU events, there is a " critical deformation," i.e., a certain degree of deformation by cold work, which particularly favours rapid growth of new crystals on annealing, both greater and lesser degrees of distortion resulting in much smaller rates of subsequent crystal growth on re-heating. This is a phenomenon for which it is by no means easy to suggest a satisfactory explanation. Whatever the theoretical explanation which may ultimately 268 STUDY OF PHYSICAL METALLURGY be accepted, there can be no doubt that the process of annealing in most metals occurs in two stages, which sometimes overlap in different parts of the same piece of metal. We have first the simple effect of more or less complete mechanical softening leaving the strained structure apparently unaltered, followed by the process of gradual re-crystaUisation by which new crystals of equi-axed shape, and thus possessing boundaries of minimum area, are produced and gradually developed in size. This process of re-crystallisation and of gradual crystal growth is also accompanied by a small change in the physical pro- perties of the metal — there is a further softening effect, accom- panied by sUght changes in density, thermo-electric power, electric conductivity, etc. One very important point must, however, be borne in mind. If metal is annealed after cold working with the object of fitting it to stand further cold working, it becomes important to allow the second stage of the annealing process to continue for a certain time, so that a crystalhne structure of moderate size may be developed. The reason for this necessity lies in the fact that there appears to be a definite limit to the amount of distortion which can be safely apphed to a given metal without producing internal fracture. Thus, in wire-drawing, if it is attempted to harden a wire too much by successive cold drawing, a point ia reached where the wire begins to " draw hollow," — i.e., the inner layers of the metal fracture in consequence of excessive distortion. Simi- larly in cold-rolHng, metal may tend to flake and spht in conse- quence of " over work." There is some reason to believe that if a metal is only just softened by slight annealing without permitting the growth of new crystals of moderate size, its power of safely undergoing further distortion is more hmited than it would be in a more thoroughly annealed condition. Perhaps the re-arrangement of the amorphous layers within the old crystals is not sufficiently perfect to restore complete ductility, but it is found far safer to allow annealing to proceed far enough to obUterate, as far as possible, the traces of the previous appUcation of " cold work " before fresh deformation is apphed. On the other hand, it is equally important to avoid unduly prolonged annealing, with its consequent coarsening of THE THERMAL TREATMENT OE METALS 269 crystal size. Quite apart from the fact that in many metals prolonged heating causes damage owing to oxidation, the coarser crystals do not appear to possess as great a degree of strength and ductility as an aggregate of smaller ones. The difference is not very marked on a tensile test, particularly in the case of a pure metal, but shock tests show a decided difference in favour of the finer structure. In duplex alloys, which depend for their valuable qualities upon the co-opera- tion of two constituents of widely different mechanical proper- ties, the time of anneaUng is often of vital importance, since it governs the relative arrangement of the constituents ; with this aspect of annealing, however, we shall deal in connection with the subject of " heat treatment " generally. The annealing process has been studied with great care in several pure metals, notably in lead (Ewing and Rosenhain {^) ), copper (Beilby (^) ), and in gold (Beilby (*), Rose (*) ). In regard to duplex alloys, the matter is complicated by the presence of the second constituent, and, for the sake of sim- plicity, we shall confine our attention to the pure metals just named. In the case of lead, it has been found that very rapid re- crystallisation occurs at a temperature of 200° C. ; at 100° C. re-crystallisation occurs to a marked extent after an exposure of thirty minutes, while it is found that definite signs of re- crystallisation can be detected after a few months in very severely cold-worked lead, even if stored at the temperature of an ordinary room. Ancient lead, as found on the roofs of old buildings, always exhibits a coarsely crystalline structure, which has, no doubt, grown slowly during many years. As regards the softening effect, the author has recently made some experiments by means of a ball-hardness test, and has found that definite softening of severely strained lead occurs, without the application of artificial heat, after a few weeks, while slight effects may even be observed after a few hours at the ordinary temperature. Some of the other soft metals — i.e., metals with a low melting-point — also undergo anneaUng at the ordinary temperature. Rose (') states that tin, cadmium and zinc all undergo softening at the ordinary temperature. 270 STUDY OP PHYSICAL METALLURGY In regard to copper, Beilby has shown that exposure to a temperature of 200° C. is sufficient to produce definite effects on the mechanical properties of a hard-drawn wire. Here, of course, the question arises how " anneahng temperature " is to be defined. The influence of a given temperature depends upon the time for which it is allowed to act, and the degree of straining or hardening to which the metal had previously been subjected also plays a part in the resulting data. It is therefore necessary to adopt some arbitrary definition of " annealing temperature." Different workers have adopted different standards, suitable to the particular purpose for which the data were required in each case. Probably the best course is to reaHse that there is no definite and absolute " annealing temperature," although for the purpose of any practical process there is always a definite temperature and duration of heating which gives the best results. Since, however, such a temperature could not be calculated from any arbitrarily defined and experimentally determined " annealing temperature," it scarcely seems worth while to attempt to lay down any general definition or convention on this subject. The annealing of gold has been very fully studied, more particularly by Beilby and Rose. Beilby finds that an effect on the mechanical properties of a hard-drawn gold wire only makes itself felt when a temperature of 250° C. has been attained. Rose, on the other hand, taking the " temperature of anneahng " as the " temperature at which a piece of metal becomes almost completely softened and re-crystallised in half an hour," finds that the very purest gold, not containing more than one part of impurity in 10,000, has an " annealing temperature " as low as 150° C, while the corresponding temperature for gold containing 0-05 per cent of copper is 250° C. This striking influence of shght impurities on the temperature of annealing has also been traced in copper, and no doubt serves to account for the higher temperature found by Beilby. In the case of wrought iron and steel, there is evidence to show that even a temperature of 100° C. is sufficient to bring about a notable change in freshly over-strained material. THE THERMAL TREATMENT OP METALS 271 Muir (^) has shown that the elasticity of iron or mild steel which has been temporarily destroyed by straining beyond the yield-point can be rapidly restored by exposure to boiling water. The common practice of exposing a razor to steam immediately after stropping probably also acts by removing the traces of over-strain from the thin edge of the steel which has been over-strained by bending backwards and forwards during stropping. Steel fully hardened by quenching in cold water also appears to be very slightly altered by exposure to 100° C. Strain-hardened iron or steel, however, although some slight and gradual changes of properties become noticeable at temperatures near 300° C, does not undergo any serious or rapid annealing until a temperature of 500° C. is reached, aM for all practical purposes this temperature may be regarded as the lowest possible annealing temperature. This fact appears very clearly in the curves published by Goerens (*). It is perhaps desirable to point out here that the " annealing temperatures," in so far as they have been mentioned in the preceding paragraphs, refer to the minimum temperatures at which certain effects have been obtained. For practical pur- poses very much higher temperatures are universally adopted ; primarily the reason for this practice hes in the fact that these higher temperatures produce the desired effects of softening and re-crystalHsation very much more rapidly and, under proper control, without damaging the metal. The duration of exposure and the temperature attained must be carefully adjusted to each requirement, and it is particularly important to note in this connection that the various metals differ very widely in regard to the rate at which re-crystallisation takes place. Thus aluminium exhibits extreme sluggishness in this respect, so that prolonged exposure to a relatively high annealing temperature is required to bring about complete re-crystallisation. In brass and copper, on the other hand, the rate of crystal growth is relatively rapid, and a few seconds exposure to a temperature of 600° or 700° C. is sufficient to bring about fairly complete anneahng in these metals. We have already referred to the evils associated with 272 STUDY OF PHYSICAL METALLURGY excessive annealing, but the matter is of such considerable importance that rather more detailed consideration is required. The effects of prolonged annealing, more especially in the case of those metals which are usually annealed at temperatures above 600° C, are two-fold. In the first place, there are the effects of mere exposure to a high temperature upon the internal structure of the metal, and, in the second place, there are in some cases equally serious effects due to the chemical and physical action of the atmosphere in which the metal is heated. Taldng the latter first, we have effects of oxidation, which are particularly felt in iron and steel where removal of carbon and formation of oxide in the surface layers always occurs in industrial annealing, even though efforts are made, by packing the steel in iron boxes filled with Ume or other powdered material, to exclude the air as far as possible. In the case of very thin pieces of metal this oxidising action may penetrate far enough to impair the value of the piece entirely. A case of the opposite kind is presented by many varieties of copper, which normally contain a certain propor- tion of oxygen. If such copper is heated in a reducing atmo- sphere, or if while hot it is even temporarily exposed to reducing gases, the metal is so seriously injured as to be ren- dered useless for most practical purposes. This action, known as the " gassing " of copper, arises from the fact that hydrogen is able to diffuse rapidly through red-hot copper ; in the interior of the copper this gas meets with small grains of cuprous oxide, and a chemical reaction follows by which the cuprous oxide is reduced to copper while the hydrogen is burnt with the formation of an equivalent quantity of steam, thus : — Cu^O -}- Hg = 2Cu -j- HgO. The water vapour produced is, however, unable to diffuse away through the copper, and a smaU blow-hole or cavity is formed in the place previously occupied by the grain of cuprous oxide. This action often suffices to render the copper extremely weak and porous. Injury to metal as the result of an annealing process is, however, much more frequently caused by the effect of THE THERMAL TREATMENT OF METALS 273 prolonged exposure to a high temperature upon its internal structure. We may consider first the case of a simple metal — i.e., a pure metal or an alloy of the " simple " solid solution type in which crystals of one kind only are present. In such a material the effect of prolonged annealing upon the resulting structure can only be of one kind, viz., the coarsening of the structure by an increase in the size and a reduction in the number of the crystals which compose the mass. We have already seen that the crystals which constitute a mass of metal exhibit a tendency to increase in size and to diminish in numbers — ■ a change which is usually described as the " growth " of the constituent crystals of a metal. At the ordinary temperature such growth is held in check by the internal resistance which the hard and solid metal offers to any molecular re-arrangement. When the temperature is raised this resistance is diminished, and a stage is reached when the tendency to growth is sufficient to overcome the resistance. If the metal has been strained, the tendency for some re-arrangement to occur is probably much stronger than in the unstrained state, and the temperature at which re-arrangement will begin — ^what we have previously called, roughly, the " anneahng temperature " — ^is probably lower than that at which crystal growth will occur in an unstrained sample of the same metal. Indeed, there is some doubt whether some degree of strain is not essential to start the process of crystal growth at all ; if this is the case, however, only very little of such impetus can be required. The real reason for crystal growth lies in the tendency, which every physical system exhibits, to dissipate any stores of potential energy which it may contain. Now there can be no doubt that the crystal boundaries contain stored energy ; whether we regard it, in accordance with the amorphous cement theory, as being identical with the latent heat of fusion, or whether we confine ourselves to the more general statement that the crystal boundaries are the seat of " surface energy " or of " surface tension," we reach the conclusion that the final state of physical equihbrium of a crystalline aggregate — ^towards which it tends to change gradually whenever the conditions are favourable to change — ^is that of a single crystal forming the entire mass, P.M. T 274 STUDY OF PHYSICAL METALLURGY with a total absence of inter-crystalline boundaries. We do not often meet with metal in this state — the attainment of this ultimate physical equilibrium is even rarer than the attainment of complete physico-chemical equilibrium in accordance with the requirements of the phase rule. None the less, these final states indicate the directions in which changes are likely to occur on prolonged heating. In the present case the tendency results in the growth and re- arrangement of crystals always in the direction of reducing the area of their inter-suriaces. And this tendency extends not only to aggregates of crystals of one kind, but also to duplex and more complex aggregates, with results — ^in the case of duplex aUoys — which we shall have to consider below. The question then arises whether the increasing size of crystals produced in a simple metal by prolonged heating is injurious or otherwise so far as the useful properties and, more especially, the mechanical properties of the metal are concerned. There can be little doubt that, within reasonable Umits, the mechanical properties of a simple metal are better the smaller the constituent crystals of which it is built up. Under the tensile test, coarseness of structure usually results only in a slightly lowered yield-point, while the ultimate stress and the elongation are little impaired, although the reduction of area at fracture is sometimes markedly less. This, of course, refers to ductile metals and alloys, the mechanical properties of brittle materials being relatively unimportant. On the other hand, under both shock and fatigue tests a coarse structure, even in a simple metal, gives unsatis- factory results. The reason is easily understood, for the crystal boundaries act as a species of strengthening skeleton ; at each boundary a crystal is supported by its neighbours, since a slip-band, for example, must change its direction — often in two planes — where it passes from one crystal to another, and if the stresses at work are such as to favour slipping in one of these directions, they will not be so favourable for slip in the other, so that the weakness of one crystal will be balanced by the strength of another. Where the crystals are large, single surfaces of slip or cleavage extend unbroken through THE THERMAL TREATMENT OP METALS 275 relatively large areas, and fracture— particularly under shock or fatigue — may be brought about under conditions which a finer grained structure would have resisted. The whole state of affairs is somewhat like the effect of a " break of joint " in brickwork or masonry ; within an individual crystal we may compare the structure to that of bricks laid without break of joint, but at the boundaries there is not only a complete breaking of joint, but also a change in the direction of the courses. In connection with the effect of prolonged heating or annealing on the general scale of the crystal structure, a further remarkable feature is sometimes met with. This is a tendency which whole groups of adjacent crystals appear to possess of assuming a similar, although not strictly identical, orientation. Etching reveals boundaries between the members of such a group, but the boundaries only appear after somewhat deep etching, and are then very fine, indicating that the difference of surface level produced by the etching process is very small. On straining or breaking such a piece of metal it is often found that slips or cleavages run with very little change of direction through the whole group, which thus behaves almost like a large single crystal. An example of a cleavage shock fracture in wrought iron which had attained this condition is shown in Fig. 118, Plate XXVI., where straight cleavage edges running almost unbroken through whole groups of crystals are clearly shown. Such cleavages running almost unbroken through a number of crystals have also been described by Stead (') as occurring in certain mild steel sheets after rolling, although their origin in that case was probably not connected with over- annealing. While examples of the evil effects of over-annealing are most frequently met with in connection with iron and steel, they occur equally in the case of copper and brass. There is, however, in these metals a feature of the structure which tends to mitigate somewhat the effects of a coarse crystal structure. This is the occurrence of the twinned lamellte which are so con- stantly found in metal which has been wrought and annealed. Although essentially of the same crystalUne system as the T 2 276 STUDY OF PHYSICAL METALLURGY parent crystals, yet the twinned regions do impose a change of direction on planes of slip or of cleavage, so that these do not have such a free and uninterrupted path as in a crystal of similar size free from twinning. The _boundar iea__of_jfcffiian-ed regio ns do not contain any amorphous cement and, if one can judge from their behaviour on anneaUng, they do not serve for the storage of latent energy, since they do not tend to coalesce on heating, so that their value from the mechanical point of view is not so great as that of true inter-crystal boundaries. The influence of prolonged heating — i.e., of over-annealing or " over-heating " — on duplex alloys is considerably more marked than in the case of simple metals. The tendency towards alterations of structure, which result in the diminution of inter-crystal surfaces, is quite as strong in the duplex alloy as in the simple metal, so that the mere growth of crystal size, quite apart from the arrangement of the two constituents, takes place exactly as in a simple metal, and with much the same results. The mechanical properties of such duplex alloys, however, usually derive their importance, in the manner already explained in the previous chapter, from the proper juxtaposition and co-operation of two constituents of unequal hardness and ductUity. In such an alloy we have, in fact, a ductile and relatively weak constituent supported and " held up " to its work by the interspersed crystals of a harder material, while this harder material is in turn prevented from undergoing brittle fracture by being firmly embedded in the softer constituent. Ideally, these effects would be at their maximum if the two constituents were mixed in the most intimate manner, so that the most desirable structure is one in which each of the two constituents is present in the smallest possible crystals whUe the two constituents are evenly mingled. When, however, a duplex aggregate is exposed to prolonged heating, there is not only the change in average crystal size to be looked for, but also a more or less rapid segregation of the two constituents. As the crystals of the predominant constituent grow they necessarily push before them the crystals of the second phase, until two or more of these meet and coalesce into a larger crystal. Not only this ; when a duplex structure THE THERMAL TREATMENT OP METALS 277 is formed, whether by sohdification from the liquid or by Beparation from a sohd solution, the second — and usually the harder — phase is left occupying the points where the crystal boundaries of the primary phase meet, and frequently thin layers of the hard phase extend some Uttle way along the inter-crystal boundaries, thus forming a network which is closely inter-connected with the crystals of the softer con- stituent. The tendency to reduce boundary areas which comes into play on heating, however, rapidly alters this arrange- ment ; the crystals of the harder phase tend to assume rounded or polygonal outlines and to retract any outlying arms, thus very materially lessening the interlocking between the two constituents. Changes of this kind, with their accompanying deterioration in powers of resistance, particularly to shock, are readily observable in all duplex alloys, including such materials as the brasses containing the a and jS phases, certain aluminium- copper alloys, and in many types of bronze. The most im- portant and, in many ways, the most interesting case is that presented by steel. The thermal treatment of steel, however, is governed by several considerations of a somewhat special kind, so that a section of this chapter will be completely devoted to it. The special features in connection with steel which affect the whole question of heat-treatment are : first, the duplex character of the " hard " constituent, i.e., of pearlite, and, second, the critical points and the corresponding constitutional and structural changes through which steel passes on heating and cooling. In view of these latter, it is desirable to discuss the behaviour of steel under two separate heads, according as the temperatures involved lie below or above the critical range — i.e., the hnes EG, GI, HI and ID of the diagram (Fig. 64, p. 161), which represents the constitution of the iron- carbon alloys. At temperatures considerably below Ar^, represented on the diagram by the line HIJ, the annealing of steel is slow, but if the temperature of Ar^ is closely approached — i.e., in the neighbourhood of 700° C. — a material amount of structural 278 STUDY OP PHYSICAL METALLURGY change is produced. This is exactly of the general type indicated above, resulting in an increa se in the size of the ferrite crystals and a bjJIin g up of the ^pgarhte i nto larger, rounger masses with more sharply defined edg^. This is the first stage of the process," but, owing to"the peculiar nature of pearlite, the process can go a step further. Pe arlite consists of la minae or grariule_s_aLce mentite em bedded in ferrite ; this laminated or granulated structure, however, is solely due to the manner in which this constituent separates from its matrix and, as a matter of fact, when very slowly cooled, steel shows no laminated pearhte. The l aminated or granular structure can also be destroyed IT^the steel is exposed to .prolonged heat ing at a ^Eemperature approaching 700° C, provided that the steel has been raised to that temperature without first passing through the critical point. This last condition simply means that the steel must be exposed to prolonged heating while in the pearlite condition. If this is done, the laminae or granules of c^entite gradually retract and " ball up '' intojajger granules, while the ferrite with which they were formerly interlardedTbeconies incor- poruted with the adjacent primary ferrite crystals,. The result- i ng structure consists of ferrite with some cementite scattered among it, e ither in isolated rounded lumps or in the form of layers OT films in the cryitaT boundaries^ An example of the ~ laEteFstructure is shown in Fig. 119, Plate XXVI., taken from a steel boiler-plate which cracked in use. Both these structures are extremely undesirable. The cementite, when present as isolated balls or lumps, is of very little use to the steel, which then possesses merely the properties of pure ferrite having a coarse structure — these, of course, being far from equal to those of a mild steel having a correct structure. Where the cementite lies in filaments in the crystal boundaries it is apt to play the part of a mere weak and brittle cement, much as the bismuth did in Arnold's experiment with gold (see p. 256). When in the condition of pearlite or Sorbite, where cementite is intimately associated with layers of ductile ferrite, the inherent brittleness of cementite is largely neutralised, and the hardness of the duplex constituent — ^pearlite — ^is well adjusted to stiffen and support the ferrite matrix of a mild steel. Cementite by THE THERMAL TREATMENT OF METALS 279 itself, however, ia not only excessively hard and brittle as compared with ferrite, but it is also too small in quantity and bulk to serve as an efficient stiffener. Its presence in the isolated state thus becomes a source of weakness rather than of strength. The behaviour, both under test and in service, of steels having this " free cementite " structure is now well recognised to be most unsatisfactory. Fortunately the pro- longed heating of steel to a temperature near 700° C. is not often Hkely to occur, and, if steel has accidentally been treated in that way, its normal structure can be restored by the use of the method of " heat refining " to be described below. It should be added that in the case of very mild steel, containing only very small amounts of carbon (0-10 per cent, and less) the influence of Ar^ and Ac^ is small, and the effect of pearlite or cementite on their strength is also less important. Consequently, up to the temperature of the line EG, of the diagram Fig. 64, these steels behave very much like simple metals and undergo gradual coarsening when heated for a prolonged period. As Sauveur has shown, when the steel has been exposed to one particular degree of strain, the subsequent growth of its crystals at these temperatures is exceptionally rapid. On the other hand, the very pure electroljrtic iron recently studied by Stead and Carpenter (*) does not appear to undergo any marked crystal growth within this range of tem- perature. These curious facts serve to show that our under- standing of the causes at work in crystal growth is by no means complete. We have now to consider the structural changes which result when steel is heated above the temperature of Acj {i.e., above the Une EGI of Fig. 64). We shall confine ourselves to steels containing less than 0-90 per cent, of carbon, i.e., to the hypo - eutectoid^ eteels. In their condit ion befo re heating t hese steels will generally be in the condition of ferrite vlus pear hte. although the normal arrangement of these constituents may be seriously distorted as the result of cold work or strain. On heating such steel above the temperature of Acg which, according to the carbon content will vary from 900° C. to 720° C. approxi- mately, the ferrite at once undergoes the allotropic change to 280 STUDY OF PHYSICAL METALLURGY the condition of y iron, and, since ce mentite is soluble in 7 iron, the cementite pjeseiit^in^th^ lajmeUae .of..M9:rJJi ^-Q ^^ "'^"^ b egins to diffuse into.ih,SLyiron._ At first this diffusion results in the amalgamation of the ferrite and cementite lamellae or granules of the pearlite itself, but this stage is usually accomplished during the process of heating-up, since the ferrite in contact with cementite undergoes thg allotrQpi s_jE]iaJiaaJa,^-iroa at much lower temperatures — i.e., as^oon^as jbhe_steeMbas_been raised above Ac^. After a time, whose exact length is not accurately known, since it necessarily varies according to the scale of the initial ferrite-cementite structure and also in accordance with the effects upon the rate of diffusion which are undoubtedly exerted by some of the common impurities of steel, the diffusion process is completed, and the steel wUl consist of an aggregate of homogeneous crystals of the y iron solid solution. The transformation of the ferrite from the a tjo_ the y state. However, quite apart from the influence of adjacent carbon in lowering the transformation temperature, does not occur suddenly or uniformly throughout the mass even of a single ferrite crystal. What really takes place is the growthof crystals ofjthe jiew.jjhase — y iron--atJ:he expens^of_theJi^riie_ or ferrite 'plus cementite. Now the growth of such new crystals commences at a considerable number of points, and, as a rule, these points will lie in the boundaries of the existing ferrite crystals, principally because these boundaries are hkely to contain particles of cementite or of pearlite here and there whose presence facilitates the beginning of the transformation. The result is, however, of vital importance, for it foUows that while the new y iron structure has some relation to the pre- viously existing ferrite, this relation is not one of numerical equality — i.e., there will, as a rule, be a number of new crystals of 7 iron in place of each pre-existing ferrite crystal — and, if the latter have been large, then the corresponding crystals are likely to be fairly numerous. At all events, the new 7 crystals^ mUjjwhen first formed-. at ajtemperature just above Ar^, b ^ more numerous and smaller than the pre-existing ferrite^ Xike any other variety of crystals existing in a metal at a THE THERMAL TREATMENT OP METALS 281 relatively high temperature, these newly-formed crystals, if maintained at a temperature above that of Ar^, will begin to grow, and the growth will be more rapid the higher the tempera- ture. If therefore, steel of this kind is kept at temperatures above the critical range for any considerable time, or even if for a short time its temperature is raised above 900° C, a coarse structure is rapidly developed in the 7 iron solid solution. If the steel is then slowly cooled, j^e^transfnrmaitJQa. nf .thfi.-y,- irpn .-SaUd_ sohition into...thfL normaLJigmte:*e£dikL.££ni£ture takes place duri ng the passage through the critical range : th^even ts which took place on heating are now to some exte nt reversed ^ and the resulti ng ferrite and pearhte crystals^are soinewhat_niota nunieirQUS--than~tha,3(Lin:tH,.oryatiala frQHUEhich they have been formed. On this reverse transformation, Btiwever, a proce"^ of rejection takes place, the dissolved iron carbide bemg pu sne a out ot solution in the y iron to form the c ementite l amellae of the pearlite . In such a process the presence of nuclei is of vital importance. Such nuclei may possibly exist in the form of traces of undissolved cementite remaining from the original structure, but this is only hkely to occur if the steel has not been overheated, since prolonged exposure to a high temperature should result in the complete solution of the cementite. In that case nuc lei may be furnished b2i;jTfl.];tir,1fis "f impmity-menhaTiina lly enclosed uTtEFste d] It has in fact been shown by Ziegler (^) that such particles can and do act as nuclei and play ah important part in determining the structure of a steel. If this is the case, we may meet with instances where a very similar tjrpe of structure is formed repeatedly after successive heatings and coolings to tem- peratures beyond Acg. Kroll (^°) describes some experi- ments in which this took place. As a rule, however, the action of such foreign nuclei is not very prominent, and the new structure formed on cooling is directly dependent as to scale and arrangement on the y iron structure from which it springs. The processes which have just been outlined are of practical importance in two ways. First, we find that by " overheating " steel, i.e., by exposing it to unduly high 282 STUDY OF PHYSICAL METALLURGY temperatures, or for too long a time to any temperature above Acj, the growth of a very coarse y iron structure results, and this, on coohng down, gives rise to a correspondingly coarse ferrite-pearlite structure. Not only this, but the arrangement and forms assumed by the pearlite which is formed from such steel is characteristic ; thereisji_strong^^tendency_for the ferrite to take the form of straightJ)ands_with,filQ33SSt-ed^JJjL4 sharply angular jgatches of pearlite between Jhem, the ferrite Easels' frequgjitlaL c£QSS^ing„ane anoth er at a ngles of 60'^'^'Tms is a structure very similar to that found in steel ingots as cast, or in steel castings which have been annealed. A typical example is shown in Fig. 120, Plate XXVI. Such a coarse, sharply angular structure is, of course, extremely undesirable ; there is a mini- mum of interlocking between the ferrite and the pearlite, and the straightness of the arrangement f acihtates the propagation of slip or cleavage through the crystals. Such structures are, in fact, frequently met with in steel objects which have failed in service. Under test they generally exhibit some degree of weakness as regards shock and alternating stresses, but their tensile strength and elongation are frequently quite satisfactory. The most typical feature, however, is a decided drop in the jdeld-point as compared with that of the same material in a more normal condition. The use of steel showing this type of structure is, of course, to be strongly deprecated wherever strength and reliability are of importance. The second practical application of our knowledge of the processes of re-crystallisation which occur when steel is heated and cooled through the critical range lies in the process known as " heat refining." Steel possessing an undesirably coarse structure may have to be dealt with, whether it be found in the form of castings or of overheated material, or as a product of special processes, such as case-hardening, where prolonged exposure to high temperatures is unavoidable. Even ordinary forged or rolled material — particularly if finished too hot — may stand in need of having its structure refined. This can be done by making use of the re-crystalHeation processes which occur as described above. By rapidly heating the steel to a temperature just above Acg, maintaining it there for a time THE THERMAL TREATMENT OP METALS 283 only just long enough to allow the whole mass of material to attain the temperature, and then cooling as quickly as practic- able, the maximum refining effect is obtained. From the de- scription given above, the mechanism of this process can be readily followed ; only in rare cases does the presence of foreign bodies acting as nuclei serve to interfere materially with such a refining operation. While rapid cooling is desirable in order to secure that the pearlite shall be formed as far as possible without any " balling up," actual quenching is neither necessary nor desirable for the purpose of heat-refining itself — the temperature of Arg being decidedly too high for safe quenching of any but small pieces of steel. If there is a case to be hardened, as in case-hardened objects, or if, for other reasons, it is desired to convert the peailite of^the steel into Martensite_ or Troostite, this is best done by re - heating the steel to a temperature just above Acj^ . coo ling, down to ju st above A r,, an d t hen quenching^__, There is one condition of steel, more or less akin to simple over-heating, which does not permit of restoration to a satis- factory condition by any form of heat treatment. This is what is known as " burnt " steel, and is the result of heating to an excessively high temperature. The explanation of " burning " now usually accepted is that it occurs when the steel is heated to a temperature above that of the solidus curve AD of the constitutional diagram, Fig. 64. It is supposed that when incipient fusion occurs in the boundaries of the crystals an opportunity is given for furnace gases, and more especially oxygen, to invade the steel and to oxidise the crystal boundaries, and thus to render the steel permanently weak and brittle. Whether this explanation be correct or not, the micro- scope shows in burnt steel the presence of foreign matter, probably oxide, at points corresponding to the location of the boundaries of the large iron crystals which existed at the time when the steel was excessively hot. These traces cannot be removed by heat treatment, and, although the general structure may be refined, the brittleness along these old boundaries is never removed and the steel is only fit for re-melting. 284 STUDY OF PHYSICAL METALLURGY In the discussion of heat treatment which has here been given, attention has been confined to pure carbon steels, and, even in regard to these, no attempt has been made to follow the exact details of the genesis of the various features to be met with in the micro-structure of steel as variously treated. Some reference must, however, be made to the thermal treatment which is of vital importance to the newer varieties of steel, in which other elements, such as nickel, chromium manganese, vanadium, titanium, and, finally, tungsten and molybdenum are introduced in notable and, sometimes, in very considerable pro- portions. It is impossible within the limits of an introductory volume to enter upon the wide field of the metallography of these alloy steels, but it is not too much to say that their practical value is very largely, if not entirely, dependent upon a knowledge of the proper thermal treatment to be appHed to them. The use of these alloying elements has given us the power of influencing the position and relative importance of the critical points in low-carbon steels, and, in consequence, we are able to obtain materials which are still in the range of stable existence of 7 iron at the ordinary temperature ; others again are in an intermediate or " Martensitic " stage in which they are capable of attaining great hardness, while in others the fundamental properties of the ordinary constituents of carbon steel are so profoundly modified that we are practically dealing with a new material, although microscopically it may still be " ferrite " or " Austenite." The study of the correct thermal treatment of these complex steels is thus obviously a matter of considerable difficulty ; in many cases the true causes at work are not yet understood, and results are obtained on empirical lines — a certain form of heat treat- ment being found to yield the desired results. The study of these ternary and quaternary alloys by the methods of Physical Metallurgy is being steadily taken up, and in time it may be confidently expected that we shall have as sound and complete a knowledge of their transformations, and a correspondingly certain command over their properties and treatment, as we now enjoy in the case of pure carbon steels. THE THERMAL TREATMENT OF METALS 285 Kefeeences. (1) Sauveur. Intemat. Testing Assoc, New York Congress, 1912. (2) Ewing and Eosenhain. Phil. Trans. Roy. See, 195a, pp. 279 —301, 1900. (3) Beilby. May Lecture, Journ. Inst. Metals, VI., 2, 1911. (4) Eose. Journ. Inst. Metals, IX., 2, 1912 and X., 2, 1913. (5) Muir. Phil. Trans. Eoy. Soc, 1899. (6) Goerens. Journ. Iron and Steel Inst., 1911, III., pp. 320 — 400. (7) Stead. Journ. Iron and Steel Inst., 1898, II. (8) Stead and Carpenter. Journ. Iron and Steel Inst., 1913, II. (9) Ziegler. Eev. de M6tallurgie, 1909. (10) Kroll. Journ. Iron and Steel Inst., 1910, 1., p. 304. CHAPTER XIII THE MECHANICAL TREATMENT OF METALS, INCLUDING CASTING In this chapter, under the heading of " Mechanical Treat- ment of Metals," we shall deal briefly with those processes which are employed in order to bring metal into the shape or form in which it is required. The majority of these processes are strictly spealdng " mechanical," in that they depend upon the direct application of mechanical forces or agents to the metal. For the sake of convenience we propose, however, to include the process of " casting " under this head, for although in that process the mechanical — i.e., the hydrostatic — pressure of molten metal is employed in order to bring the material to the shape of the mould, yet casting is not usually regarded as a " mechanical " process. Castings, in the form of ingots, how- ever, form the starting-point for all rolling and forging processes, and many of the qualities of the ingot are of vital importance to the final result. While, in general, the primary object of all the processes to be considered here is the production of a definite shape or form, these processes exert a great influence upon the internal structure and mechanical properties of the metals to which they are applied, and it is largely from that point of view that they will be considered here, as bearing upon the correlation of constitution and structure with strength. In regard to the process of casting, we shall leave on one side the whole art of the foundry in regard to the design and construction of moulds and all matters connected with the furnaces for melting the metals. Although these matters do not lie outside the scope of Physical Metallurgy, their treatment is best relegated to a special treatise dealing with that branch of the subject. Here we are concerned principally with the MECHANICAL TREATMENT OF METALS 287 effects of the casting process upon the constitution, structure and properties of the resulting material. Two totally different types of " casting " must be distin- guished at the outset. In one, which is the essential province of the founder, the casting itself is the final product, and — subject to some thermal treatment, such as subsequent anneal- ing — the properties of the material are those which it receives as it solidifies from fusion. In the second type of " casting " the object which is directly produced from the molten metal is merely a step in a long process ; the ingot is to serve as the starting point for the production of rolled, forged, drawn or otherwise wrought material. In consequence of this difference, the methods and aims of the processes employed vary very widely between the two types of " casting." It must be pointed out, however, that in many respects the long and costly processes of " working " metal are rendered necessary by the imperfection of the cast material itself — ^its internal structure requires thorough refining before it reaches the requisite standard, whether of strength or ductility. Modern practice, always tending to seek economy by reducing the number of manipulations, is constantly seeking to produce by direct casting, products which could formerly be obtained only by more complicated and costly methods. The great growth of the " malleable castings " industry is a striking example of this kind, while recently in Sweden the production of large steel castings has been pushed so far that it was thought possible to produce guns of fairly large calibre by casting direct ('). The production of very satisfactory crank-shafts for marine engines by casting alone is claimed as a commercial success in certain Continental works. For a given material, the constitution and structure of any casting will depend almost entirely upon the rate of cooling, and principally upon the rate of solidification which the metal has undergone in the mould. In castings of our first type, which are to be used as such, it is obviously important to secure the best possible structure by securing the correct rate of coohng. For the majority of materials the best rate of cool- ing is simply the most rapid rate that can be employed without 288 STUDY OP PHYSICAL METALLURGY producing cracking or warping of the casting. Thus it is almost universally found that metal from the same pot cast in " chill " moulds yields a far finer structure and better mechanical properties than the same material cast in a mould made of sand. A " chill " mould, being meide of metal, conducts the heat away rapidly and leads to rapid sohdification, while the sand mould is a poor conductor of heat and keeps the metal hot for a long time. Exceptions to this rule occur in metals which undergo one or more critical transformations during cooMng, and in such cases it may be necessary to avoid that suppression of such changes which too rapid a rate of coohng might bring about. A typical example of this kind is found in cast-iron, whose structure and constitution varies very markedly according to the rate of coohng. If the iron is cooled rapidly, particularly if it contains little siUcon, the formation of graphite during the first stages of sohdification is prevented by casting in a chill mould, with the result that the iron when cold consists of cementite and pearhte, or even Martensite-Austenite. This constitutes what is known as " white iron," which, but for the usual impurities, is practically a very hard high-carbon steel. Such chiUed iron has its important uses, and in other cases its great hardness is not objectionable, but where the castings have to be cut or machined, this hardness would render them useless. Even apart from the effects of chiU moulds, or of chills locally introduced in order to produce special effects, the rate of coohng or of sohdification of a casting can be regulated to a considerable extent by controlhng the temperature of the metal at the moment of pouring. A high casting temperature, quite apart from other disadvantages attaching to it, reduces the rate of sohdification, because the excess heat of the metal ia communicated to the mould before the freezing-point of the metal is reached, so that the thermal capacity of the mould is largely exhausted before solidification begins. The resulting very gradual solidification causes not only a very coarse structure, but allows of the occurrence of serious segregation, the more fusible portions of the alloy tending to rise into that part of the casting which is the last to sohdify. It is, therefore MECHANICAL TREATMENT OF METALS 289 an almost universal rule that casting temperatures should be kept as low as possible, consistent with adequate fluidity of the metal to allow of the proper filling of the mould. Not only is it desirable to keep down the temperature at which metal is poured into the mould, but any overheating of the molten metal should be avoided, although, if it is allowed to cool down to the proper temperature before pouring, the greater part of the harm is avoided. None the less, there seems to be some ground for thinking that metals which have been thus over-heated tend to develop a coarser structure than would otherwise be the case. There is also a considerable likelihood that the metal will be injuriously affected as regards chemical composition during such over-heating. It must be borne in mind that there is a constant tendency to estabUsh a state of chemical equilibrium between the molten metal and any substances with which it is in contact. When very hot, the chemical actions due to this cause are much intensified, and many metals absorb or reduce some of the constituents of the crucible, stirring or skimming rods, etc., with considerable avidity. The absorption of silicon by aluminium or its alloys, when heated in a crucible containing any form of sihcates, is a typical example. The metal is, moreover, in contact, not only with the walls of the crucible or furnace, but also with the furnace atmosphere. The absorption of gases by hquid metal appears to increase with increase of temperature — ^this depar- ture from the laws which govern the solution of gases in liquids such as water, is probably due to the fact that the metals form loose compounds' with such gases as hydrogen and nitrogen. The gases thus absorbed are frequently retained by the metal until crystalhsation sets in, when the gases are hberated and render the casting unsound. While rapid cooUng of castings through the freezing range is eminently desirable from many points of view, a Mmitation is imposed in many cases by the circumstance that contraction stresses will produce fracture or warping if too great a rate of cooling is adopted. These stresses arise if the inner layers of a casting are still fluid while the outer layer has already become sohd and rigid. AH metals undergo changes of volume at the P.M. V 290 STUDY OP PHYSICAL METALLURGY moment of solidification, and if a fluid core is contained in a solid envelope, severe stresses must be set up as soon as the core begins to solidify and endeavours to change its volume in accordance with its normal behaviour. The outer shell wiU resist any such change of volume, whether it be expansion or contraction, and severe stresses wiQ be set up, which may easily restdt in fracture. The manner in which a casting is allowed to cool will also affect its behaviour in this respect ; if the relative rates of cooling of thick and of thin portions are equalised, much more rapid rates can be adopted than would otherwise be permissible. The shape of a casting, quite apart from either the character of the metal or the other circum- stances of casting, exerts a con- siderable influence on the resulting structure, since the shape frequently regulates the manner in which cool- ing occurs, and thus determines the directions in which the most rapid flow of heat takes place during solidi- flcation. Now, it is a universal fact that crystals always grow along the Mnes of heat flow, the growth taking place in the direction opposite to that of the flow. The reason for this behaviour on the part of the crystals is readily understood when we realise that in a fluid metal whose temperature is at the freezing-point each crystal will continue to grow in aU directions until it meets with the growing arms, or " dendrites," from an adjacent crystal. Now consider a mass of cooling metal whose cross-section is circular, repre- sented as the outer circle in the diagram of Fig. 121. In such a mass the centre will be at the highest temperature, and we may draw as isothermal lines — i.e., lines of equal temperature at any moment — the circles 1 — 2 — 3 4, Fig. 121 each representing a fall of temperature of — say 5° C. Now take the instant when the outer layer, represented by the outer circle, just reaches the freezing-point. At a number of separate Fig. 121. — Diagram of Iso- thermals in a Cylindrical Kod while Cooling. MECHANICAL TREATMENT OF METALS 291 points in this circle, crystal nuclei will spring into being more or less simultaneously, and each will begin to grow outward in all directions. In the direction radially outwards this growth is arrested by the walls of the mould, and parallel to the circumference of the circle the growing arms of adjacent crystals soon meet one another and put a stop to further extension of individual crystals in that direction. Only in the direction radially inwards is there room for unimpeded growth. There the metal is, at any instant, too hot to allow of the formation of fresh nuclei and, as each successive layer cools to the freezing-point, the advancing dendrites of the circle of existing crystals occupy the area. This process of growth in directions at right-angles to the isothermal surfaces, or contrary to the lines of flow of heat, governs the crystalline arrangement of all castings so far as a fringe around their external boundaries is concerned. The depth of this fringe will depend upon the rate of cooling, since the process, as described above, continues so long as the temperature gradient is steep enough to ensure that the layer which is at any instant at the freezing-point is never thick enough to allow of the formation of any considerable number of fresh crystal nuclei. In castings of moderate thickness this state of affairs usually persists until the whole of the metal is solid, and the casting is seen to consist of crystals growing inwards from the sides. When the casting is thicker, however, a period is reached when the temperature-gradient becomes very slight in the interior of the metal, which is still molten. In these circumstances a large number of independent nuclei spring up in the liquid, and each of these grows in all directions, thus producing a mass of approximately equi-axed crystals. Thick castings, therefore, usually consist of fringes of crystals growing inwards from the outer boundaries of the casting, with a large mass of equi-axed crystals forming the centre or core. An example of this kind, taken from a casting of lead, is shown in Fig. 122, Plate XXVII. An important feature in the crystalline structure of castings is that which is met with where sharp angles, and, more particularly, re-entrant angles, occur. The arrangement 172 292 STUDY OF PHYSICAL METALLURGY assumed by the crystals in such a case is that illustrated in Fig. 123, Plate XXVII., where we seethe two systems of "fringe" crystals meeting at right-angles. Such a structure is known to be weak, and castings frequently crack at such sharp corners. This is often ascribed to the existence at such points, of surfaces where a number of crystal boundaries all he in one plane. We know, however, that crystal boundaries are not in themselves sources of weakness, but rather the reverse, and the true explanation of weakness at such points must, therefore, be sought elsewhere. Actually, several sets of circumstances conspire to render such re-entrant corners weak. In the first place, if we draw the isothermal surfaces, as sketched in Fig. 124, for such a portion of a casting, we see at once that, close to the corner, the heat con- ducted into the mould from one of the cooling surfaces which forms one branch of the angle must retard the cooUng of the adjacent sm:- face and vice versd ; there- fore the isothermals must curve inwards towards such a corner. It follows that the metal actually at the corner will, for every successive layer, be the last to soKdify. The metal at these points will, therefore, still be liquid when the remainder of the corresponding layers is already solid, and, as these already solid parts tend to contract as they cool, there is an obvious tendency to produce shrinkage cracks at these points where the metal is weakest, because hottest. There is also a tendency for gases and other impuri- ties to be forced into these positions, and thus to assist the other forces in bringing about local injury. The simple cases which have been discussed above serve as sufficient examples of the principles which apply to the behaviour of castings during solidification. If we can Fig. 124. — Diagram of Isothermals at a Ke-entrant Angle ia a Cool- ing Casting. PLATE XXVII. Fic. 12'.'. Fig. 125. Fig. 123. \To far-eiJ.2m. MECHANICAL TREATMENT OE METALS 293 draw the isothermals representing the cooling process — and their general outline can usually be inferred from the shape and conditions of cooling of the casting in each particular case — we know at once the direction in which fringe crystals will grow. The isothermals themselves practically indicate the successive stages of solidification, since the edge of the solidified portion at any moment is necessarily an isothermal — viz., that of the freezing-point. Strictly speaking, the arguments here set forth apply only to the solidification of a pure metal, every portion of which undergoes solidification at one tempera- ture ; in practice, however, the majority of alloys behave in a very similar manner. In the case of most solid solutions the freezing-range is short as compared with the temperature gradients met with in solidifying metal. In eutectiferous alloys the comparatively tardy solidification of the eutectic tends to accentuate most of the effects due to unequal cooUng and contraction. Incidentally, this method of considering the isothermals of a cooUng casting will also enable those interested to work out the effects of special features, such as the insertion of local " chills " in sand moulds or other devices for producing local rapid or slow cooling. It is only necessary to bear in mind that the isothermals are close together where cooling is rapid, and far apart where it is slow, in order to be able to sketch with some accuracy the shape of the isothermals for any given case. The various points discussed and, in a few instances, illus- trated here can be demonstrated in a very simple manner by moans of an ordinary piece of good tin-plate having a fairly thick coating of tin. By heating a portion of such a plate over a flame, the tin can be melted, and it will then solidify very much as if it were a section of an ingot ; if left to itseff crystals of the " fringe " type will grow inwards, giving a structure typical of a small ingot cast in a chill mould. Other types of crystallisation, illustrating the relation between crystal growth and isothermal lines, can be readily produced by local chilling or heating, and the effect on the resulting structure watched. This structure is readily revealed by etching the tin surface with strong hydrochloric acid or some 2&4 STUDY OF PHYSICAL METALLURGY other suitable reagent, the crystals thus revealed being amply large enough to be visible without magnification. An example of the structures produced in this way on tin-plate is given in Fig. 125, Plate XXVII. The structure and constitution of a casting cannot always be rendered satisfactory, since in large masses of metal the rate of cooUng is not under ready control. In the particular case of steel, the structure can be refined by the process of " heat refining " described in the previous chapter, and such treatment has a further advantageous effect in altering the maimer in which impurities, notably sulphides and phosphorus, are dis- tributed in steel castings, thin inter-crystalline films becoming coagulated into small and relatively harmless globules. It must be noted, however, that the process of heat refining is almost entirely confined to steel. It is entirely inapplicable to metals or alloys which do not undergo a critical transformation result- ing in the crystallisation of a new phase. In the case both of brass and bronzes of certain compositions such critical changes do exist, but their power of producing a refining effect on the structure is smaller than in the corresponding case of steel (*). In the case of ingots, whose shape is subsequently to be altered by v/orking, the exact character of the internal structure appears to be of less immediate importance than in ordinary castings which are to retain the shape in which they solidified. Yet, even in ingots, for forging or rolhng, certain features resulting from the casting and the solidifying process are of great importance. The actual size and arrangement of the crystals, although it does affect the ultimate structure of the wrought material, is of secondary importance, but the mechanical soundness of the ingot is vitally important. We have already seen that metals, on freezing, are apt to become porous, or to contain cavities of various sizes from two distinct causes, one being the liberation of gases and the other the shrinkage or contraction of the metal. The liberation of gases can be minimised by keeping the temperature of the molten metal as low as possible, but, even then, some gas appears always to be liberated on freezing. On the other hand, the shrinkage of the metal during freezing cannot be prevented, MECHAITICAL TREATMENT OF METALS 295 and the only course which can be adopted is either to secure that all the shrinkage shall take effect in one part of the ingot, or in an attached " header " or " riser " from which hquid metal is allowed to run into the mould, thus " feeding " the ingot as it solidifies. The exact nature of these actions has, perhaps, been most fully studied in steel, partly because the very large ingots which are habitually cast in that metal exhibit these phenomena on a correspondingly large scale. The phenomena occur, however, equally in all alloys. In the case of steel, the question of gases can be largely reduced by allowing the metal to remain in the furnace until it is in perfectly quiet fusion or " dead melted." If this is not done, it imphes that the chemical actions of the refining furnace, which lead to the generation of gases, have not been completed, or at least that the gases thus generated have not been allowed time to escape from the metal. It is probable that the action of aluminium in rendering steel " quiet in the mould " consists in stopping these chemical actions ; the metallic aluminium reduces any iron oxides still present, and the alumina thus formed is incapable of reacting with the carbon still present in the steel ; the generation of carbon monoxide or dioxide is thus at once arrested, and the steel becomes quiet. The aluminium may also exert a purely physical effect in assisting the internal circulation or " feeding " of the ingot ; at all events, the net effect is that steel which has been treated in that way forms a perfectly sohd ingot so far as small scattered cavities or " blow-holes " are con- cerned, but, on the other hand, a deep central cavity or " pipe " is formed. When an ingot is subjected to roUing, if it contains numerous small cavities, these are closed up, and at worst appear as fine longitudinal tubes in the finished steel — ^in many cases the surfaces of such blow-holes are sufficiently clean to become welded together during rolling or forging. In other cases, however, they leave discontinuities in the steel, and these are dangerous in the finished material. On the other hand, the deep central " pipe " of an otherwise solid ingot can never be removed by rolling — ^nor even covered up — and the steel-maker is reduced to cutting off as scrap the whole of the piped portion 296 STUDY OF PHYSICAL METALLURGY of the ingot. The maker of steel ingots is thus placed between two serious difficulties, and many efforts have been made to develop a system whereby ingots could be obtained which, while solid, shall yet be free from any deep pipe. Among these the Whitworth process of " fluid compression," in which the steel is exposed to intense hydrauUc pressure during sohdifica- tion, is well known. Such treatment wiU undoubtedly close all small shrinkage or gas cavities, and to some extent it also prevents the formation of a pipe as a result of contraction and internal " feeding." It is, however, a very costly process, particularly as regards the plant required. Almost equally costly is the " Harmet " process of ingot compression ; in this process the steel is poured into a conical ingot mould whose thinner end is upwards. By means of suitable hydrauKc presses, the steel, as it solidifies, is slowly driven upwards into this mould, so that it is slightly " wire-drawn." The result is a compression of the sohdified shell of the ingot and a gradual squeezing out of the still liquid core. The resultiag ingots are very perfect in every respect, but the cost of the plant and the time occupied make it impossible to use this process except for special work. Recently, Talbot (^), has suggested a process in which the compression of the outer solid shell of an ingot, while the interior is still fluid, is accomphshed by passing the ingot in this state through the roUs. There is, of course, the obvious risk that the shell may burst and the hquid contents squirted out into the miU. Apart from this, the process appears to be a promising one, although proofs of its practical value on the large scale have yet to be obtained. Another method for obtaining sound ingots is that described by Hadfield {*), in which the upper portion of the ingot-mould is developed into a very small furnace, burning charcoal fuel in a blast of air, thus keeping the top of the ingot very hot and fluid, with the object of providing a " head " of liquid steel from which the interior of the shrinking ingot below may be freely fed. For the non-ferrous metals such processes have been worked out in a few special cases, but, as a rule, in these industries the ingots are used as they are formed by simple cooUng. Owing MECHANICAL TREATMENT OE METALS 291 to the smaller size of the ingots, and the fact that the value of the metal itself is so much higher in comparison with the cost of re-melting it than is the case with steel, the whole question has never attained that degree of importance which it has assumed in the steel industry. A study of the methods of the steel industry in this respect would, however, be of undoubted value to those who have to deal with other metals, and it is unfortunate that the somewhat arbitrary distinction often drawn between " ferrous " and " non-ferrous " metals and their treatment tends to keep the two branches of metal- lurgy distinctly out of touch with one another in such matters. Apart from the question of mechanical soundness in an ingot, the casting and solidification processes exert a powerful influence on the character of the final product in regard to the distribution of the constituents and impurities. It may be well to point out here that metallographic examination has fully proved that the changes of form which are imposed upon an ingot by subsequent working are participated in by each small section of the ingot, so that the relative positions of different parts remain as undisturbed as possible in the resulting product — i.e., the metal which formed the core of the ingot forms the core of the finished bar, or rail or forging, and similarly for the skin and intermediate parts. If, therefore, there is a marked difference of composition between different parts of an ingot, there wiU still be similarly marked differences between the corresponding parts of the finished piece — there is no " mixing " action in the hot or cold working of a metal, except in regard to a certain amount of diffusion, which, how- ever, generally merely affects the minute structural constituents, but not the composition of any considerable regions of metal. Our insight into the process of sohdification is sufficient to explain why and how differences in the chemical composition in different parts of an ingot may arise. The principal factors are that the ingot undergoes solidification in a certain succession, the peripheral and, usually, the lower portions freezing first while the inner and upper portions are still hquid. But solidi- fication occurs in stages or succession, not only in regard to position in the ingot, but also in regard to chemical composi- 298 STUDY OF PHYSICAL METALLURGY tion. The materials usually cast in ingots are not pure metals, and, therefore, consist of portions having different freezing- points. Solidification, therefore, begins by the separation of the least fusible constituents — those whose crystals are deposited when the " liquidus " of the particular alloy present is first passed. These crystals, as they grow, push before them the more fusible constituents. This separation, however, will be very far from complete, yet a certain definite concentration of the more fusible constituents must and does occur. It follows that the portions of an ingot which solidify first contain an excess of the primary metal, while the last portions to solidify contain a larger proportion of eutectic and other relatively fusible constituents. In this way the phenomenon known as " segrega- tion " is produced. This makes itseK very strongly felt in such a complex body as steel, in which relatively fusible impurities are always present ; in a comparatively simple metal like brass, on the other hand, where the whole freezing-range is short and fusible impurities are not usually present, segregation is not an important feature. Each material must, however, be con- sidered and examined on its merits in this connection. The prevention of segregation in alloys where it occurs to a serious extent is always a difficult matter. Rapid cooUng, which does not give the materials time to separate appreciably, is one remedy, but this is frequently inapplicable. The surest remedy is the alteration of the composition of the metal in such a way as to lessen the total freezing-range. As a rule this simply amounts to the careful removal of impurities, particularly of those which produce a relatively fusible con- stituent and thereby increase the total freezing-range of the alloy — since, from this point of view, impurities must be con- sidered as component members of the alloy system in question. In some cases, as, for instance, by the operation of the Harmet press referred to above, it is possible to squeeze out the most fusible and most impure portions of the metal, and in other cases the drastic remedy of severely " cropping " the ingot and thus cutting to waste the most segregated portions of the metal can, and must, be adopted. The difficulty of elimi- nating segregation and its evils is, however, very considerable, MECHANICAL TREATMENT OP METALS 299 and increases rapidly with increasing size of the masses of metal to be dealt with. We now pass on to consider some of the mechanical operations by which the ingot is brought into the desired shape or form. These operations may be applied to the metal either while it is hot, and then constitute "hot working," or while it is cold. A mere distinction by temperature, however, is unsatisfactory for this purpose, since a temperature of 400° C. is such as to constitute " hot " working for an aluminium alloy, while it would imply cold working for steel. The real basis for a rational distinction between hot and cold working lies in the effects which the operation produces. We know now that at all temperatures the apphcation of strain — i.e., of forcible changes of shape — results in certain processes within the metal, such as deformation by slip accompanied by more or less movement at the crystal boundaries and a more or less consider- able disturbance of the molecular arrangement at all surfaces where displacement has taken place. The vital distinction between hot and cold work, however, lies in the fact that at high temperatures a compensating annealing action is con- tinually taking place, allowing the disturbed molecules to re-assume a crystaUine arrangement, and thus constantly tending to re-establish the normal condition of an aggregate of equi-axed crystals in the metal, even while deformation is going on. The rate at which this annealing action takes place will, of course, depend, in a given metal, upon the temperature. At a sufficiently high temperature the rate at which the crystals are re-formed and grow will be as rapid, or even more rapid, than that at which they are disturbed by mechanical deforma- tion. If the working operation is stopped at any instant and the metal is allowed to cool down from such a high temperature, there will be no direct signs of the application of work, i.e., the metal will be completely annealed and will consist of an aggregate of equi-axed crystals. These crystals will be very small compared with those of the original ingot, for example, because the crystals have been deeply disturbed, and those finally present have only been allowed a very short time for their formation and growth, but there will be no signs of SOO STUDY Of IPiBiYSICAL METALLUflGY distortion and no mechanical hardening effect which could be removed by subsequent anneaUng. If now we lower the working temperature, a point will be reached where the rate of annealing is no longer rapid enough to obliterate the effects of progressive deformation or working. If such metal is cooled down, there wiU be signs of distortion and of strain-hardness. Such metal wiU be capable of softening by subsequent annealing. In such a case we see the beginnings of " cold work." In the great majority of industrial rolling and forging operations the process of working is continued down to a temperature well below the hmit at which this stage is reached, so that all ordinary wrought metals show signs of " cold work " and are more or less strain-hardened. This condition, due to what is known as a " low finishing tem- perature," is, however, still far removed from the extreme of cold working which is met with in hard-drawn wire or cold- roUed sheets, where very large amounts of deformation have been appUed in the cold deliberately with the intention of producing strain-hardening. The considerations which have just been indicated serve at once as a guide to the nature of forging and rolling operations. At very high temperatures large deformations can be applied without hardening the metal, simply because continuous re-crystallisation undoes the work of straining. On the other hand, this same process of re-crystaUisation wiU set in unchecked the moment that the deforming process is stopped, so that a rapid coarsening of the structure is likely to result, even during the mere gradual cooling of any large mass of metal. For that reason it is generally desirable to continue the working operations until a moderately low temperature is reached. This will result in slight strain-hardening of the metal, but will, at the same time, prevent the growth of coarse crystals. The slight strain-hardening thus introduced is not in any way objectionable, since the apparently raised yield-point or elastic limit which it causes is readily recognised and allowed for when the material is tested — a coarse crystalline structure is a much more serious disadvantage in a material. It should be noticed here, also, that the temperature at which strain- MECHANICAL TREATMENT OF METALS 301 hardening will set in depends upon the rate of deformation to wliich the metal is subjected. To a given temperature belongs a certain rate of annealing or re-crystallisation ; if the amount of deformation applied per second is so large that re-crystallization cannot reorganise the crystal structure before deformation is repeated or continued, then strain- hardening sets in. The influence of the modern rapid methods of roUing and working metal thus generally tend to the intro- duction of strain-hardening at finishing temperatures which would be amply high enough to leave the metal completely soft if slower methods of working had been employed. The lower limit of temperature to which ostensible " hot work " can be pushed is thus defined by the amount of strain- hardening which is permissible in a given case. An upper limit also exists. In the first place, the majority of metals cannot be raised above a certain well-defined temperature without " burning " them. We have already referred to this matter in relation to steel, but a similar limitation applies to most metals and alloys — viz., that it is not safe to heat them above the temperature at which the " solidus " curve of the constitutional diagram is crossed — and in impure commercial alloys this often lies as much as 100° C. below the temperature of the solidus in the pure binary system to which the alloys ostensibly belong. In pure metals the solidus, of course, coincides with the liquidus, and there is thus no risk of " burn- ing " a really pure metal in the usual sense of the word, although excessive heating will bring about great oxidation— in some metals — and disintegration will follow. Even apart from this, however, all metals become extremely weak and brittle close to their melting-points, so that much lower temperatures should always be adopted for purposes of hot working. Our existing knowledge of the behaviour of metals under deformation, when applied slowly and quickly respectively, affords some insight into the effect of rapid hot working, as by hammering or rolhng, and the slower process of forging in the press. Under the hammer there is a sudden application of stress and, unless the metal is at such a temperature that the viscous resistance at the crysta,! bQundaries is exceptionally 302 STUDY OF PHYSICAL METALLURGY low, the deformation will take effect chiefly within the crystals, producing a maximum of slip and a minimum of displacement of the crystals as a whole. With steady, slow deformation as in a press, the reverse is likely to be the case if the metal is in a really soft condition. Rolling occupies a somewhat inter- mediate position, since the deformation, although rapidly applied, is sometimes applied in very numerous and gently- graded stages. Most industrial roUing processes must, how- ever, be regarded as applying deformation by something much akin to shock . How far the differences in the internal mechanism of deformation affect the resulting structure is not yet known, and it is not a matter upon which we have anything to guide an anticipation. That there are differences between materials forged under a hammer or under a press is well known, but the internal mechanism by which they are set up, once the original deformations have been applied, yet remains to be discovered. The whole subject of the behaviour of metals when deformed at high temperatures has only been taken up quite recently (^), so that further important results may weU be anticipated from continued research on those lines. Although we have said above that the annealing action which occurs in hot metals " undoes " the effects of the straining actions of hot working, this is true only with certain limitations, as there are some effects which ordinary annealing does not readily remove. These effects are in reality nothing but internal changes of form, just as the general change of shape is an external change of form. If, for instance, an ingot contains a core of different composition which lies as an ellipsoidal mass in the original ingot, this ellipsoid will be elongated into a rod with slightly tapering ends when the ingot is rolled out into a long bar or rod. Now if this is done at a high temperature, or if the roUed bar is afterwards annealed, although the crystal structure will be restored to that of an equi-axed aggregate, the core of different composition will not be restored to its original position or shape. If now a longitudinal section of such a bar is examined and traces of the elongated rod-like core are found on the sectipn, these traces will be symptoms of the MECHANICAL TREATMENT OF METALS 303 work to which the metal has been subjected. Rightly regarded, however, such a trace of deformation is neither more nor less a trace of working than is the elongated shape of the bar itself. What we have just said of a single central core or layer of different composition applies necessarily also to a considerable number of distinct layers of different composition, so that ultimately we may expect to find even certain of the micro- constituents arranged in long lines or rods in a piece of metal which has been hot-rolled or forged. In such cases each individual crystal of these constituents is restored to its equi- axed condition by annealing, but ordinary annealing does not, and cannot, undo the longitudinal distribution of the various particles of the different constituents. Such longitudinal distribution must not, in itself, be taken as a sign of " cold work " — the only really sure sign of cold work, i.e., of deformation which can be in a sense "undone" by annealing, is that of elongated individual crystals. The disposition of the constituents in elongated lines or threads is most frequently met with in duplex alloys, such as mild steel, or in the a -f /3 brasses ; in a pure metal there can, of course, be no such disposition of constituents, and the same apphes to any simple solid solution. The case of steel is peculiar, since that material, at high temperatures, is a homo- geneous solid solution in which two separate constituents do not exist. The hot working of steel is, however, usually continued down to temperatures well below Ar^, with the result that, during the later part of the operation, the material already has a duplex structure, and the pearlite can and does arrange itself in longitudinal Jines, which are not obhterated by ordinary annealing, particularly if the steel is not taken into the y-iron range. We may thus have, and indeed frequently find, steels in which all traces of cold work have been removed by anneahng at moderate temperatures, but in which the pearHte is still arranged in long lines . An example of this kind is seen in Fig. 126, Plate XXVIII., where the etching has been carried far enough to show the outlines of the ferrite crystals whose equi- axed forms indicate the absence of serious strain-hardening. These longitudinal distributions in rolled materials are not 304 STUDY OF PHYSICAL METALLURGY confined to structurally distinct constituents, but may be traced in regard to substances which are present in soMd solution, but which were not uniformly disseminated in the original ingot. In the case of steel, phosphorus occurs typically in this way, being present in solid solution in the ferrite of the ingot, but in the form of solid-solution cores, so that the phosphorus content of each crystal increases from its centre to its periphery. When rolled out, these crystal cores assume the form of elongated masses, and, although the ferrite itself undergoes complete re-crystalhsation, possibly repeatedly, there is nothing to cause the phosphorus to migrate except the process of diffusion, which is particularly slow in that case. The result is that in the finished material the phosphorus-rich ferrite still remains in long bands or streaks, and these bands pass indifferently through numbers of individual crystals — indeed, an individual crystal may lie partly within and partly outside one of these bands— the growing ferrite crystal has simply used the material it found at hand, whether rich in phosphorus or not. The existence of these bands has been clearly shown by Stead (^), who develops them on the poHshed surface of steel by the process of heat-tinting. The phosphorus- rich bands do not oxidise at the same rate as the rest of the steel, and the bands appear as regions of different colour. The author and Haughton C) have recently shown that the same effect may be produced by the electro-chemical deposition of copper from a solution of ferric chloride and hydrochloric acid containing a small amount of copper chloride. This reagent develops the banded structure of steel containing phosphorus in a striking manner, as illustrated in Fig. 127, Plate XXVIII. The striking persistence of these bands, not as a matter of cold working, but as an indication of actual displacement, apart from crystalline structure, is clearly shown by an experiment made by the author. A typically banded steel was subjected to a Brinell ball impression ; a section was then cut and polished, and the manner in which the impress of the ball has displaced the bands is clearly seen in Fig. 128, Plate XXVIII. The piece of steel was then very thoroughly annealed so as to remove all strain-hardness and to allow of complete re- PLATE xxviir. Fk; 120. Fid. 1L>7. Fig. 12.S. Fig. 129. [To face p. 304. MECHANICAL TREATMENT OF METALS 305 crystallisation ; a second photograph (Fig. 129, Plate XXVIII.) shows the distorted bands still in existence. Such bands can, in fact, only be removed by a sufl&ciently prolonged heating, allowing diffusion to take place, which would render the distribution of the phosphorus uniform and thus remove the banding completely. A somewhat striking fact in connection with the hot working of metals, which deserves mention at this point, is the circum- stance that while, of course, there are many metals and alloys which do not permit of working, either hot or cold, there are also some which, while readily workable in the cold, are quite brittle when hot, and, therefore, cannot be subjected to hot working. 1 The best-known examples of this kind are the zinc-copper aUoys (brasses) which contain less than 30 per cent, of zinc. This fact is the more remarkable since other alloys containing a soUd solution of the a tjrpe, mainly consisting of copper, are readily worked hot ; the aluminium-copper alloys are a good example of this land. In the zinc-copper series the possibility of hot work again arises as soon as the yS soUd solution appears in the alloys, in spite of the fact that, in itself, the ^ body is a much harder material than the a sohd solution. This anomalous behaviour of the brasses suggests that there may be a transformation in these alloys occurring at some temperature between that at which they are ductile and that at which they become brittle, which accounts for the change, but there is no thermal evidence for such a change, and no miscroscopic evidence has as yet been obtained. The matter is one, however, which has not been exhaustively investigated, although it is a subject which would well repay careful study. In some cases where alloys which are ductile in the cold prove to be brittle when hot — and such cases are really rare— a transformation is always known to occur which results in the appearance, at high temperatures, of a brittle phase not present at low temperatures. The alloys of nickel, zinc and copper, known as " German sUver," also become brittle when hot and are always rolled and drawn in the cold. • See also p. 145 and footnotp. P.M. X 306 STUDY or PHYSICAL METALLURGY The general mechanism of cold working has already been considered, both in our general treatment of the plastic deformation of metals in general and in connection with our discussion of the respective meanings of the terms "hot " and "cold "work. But little need be added here. The cold work- ing of the more ductile metals, owing to its convenience and cleanness, and the possibility which it presents of producing finished articles in large numbers by such processes as drawing, stamping, spinning, etc., possesses very great practical import- ance. Not only is the process employed for the purpose of bringing metals into the desired shape, but the hardening effect of the operation is also relied upon to give to otherwise soft metals the hardness and stiffness requisite for many purposes. Wh'ere the strain-hardness of cold-worked metal is merely required to afford stiffness for small articles, and where — generally — ^it is not a question of resistance to serious and continued stresses, the utilisation of strain-hardness is perfectly legitimate and rational. In some cases — particularly in the case of cold-drawn wires — ^the artificially induced strength appears to be more or less permanently reliable, although even there, in the case of wire ropes subjected to repeated bending, fatigue failures occur owing to the circumstance that as regards alternating stresses strain-hardness is of no avail. More serious are those cases where rods of hard-drawn or cold- rolled alloys are employed for such purposes as bolts or in other positions where they are called upon to carry important loads. Practical experience in such cases confirms the conclu- sions to be drawn from research data, that the extra " strength " due to strain-hardness cannot be safely relied upon for continued resistance, particularly where stresses are alternating or intermittent. The best recent practice shows a strong and highly rational tendency to avoid the use of any material which has been severely cold-worked, unless it has been subsequently annealed in such a way as to remove strain-hardening more or less completely. The behaviour of various metals and alloys under the action of cold working varies very widely, the amount of plastic MECHANICAL TREATMENT OF METALS 307 deformation which can be safely employed without inter- mediate annealing depending partly upon the actual ductility of the annealed metal and also upon some other properties which are not yet well understood. It is not always the material which shows the greatest " ductility " under a tensile test, for example, which will " stand " the greatest amount of cold working. The actual size of the crystal structure appears to exert an important influence, extremely small and very large crystals being apparently aUke unfavourable to the appli- cation of a large amount of cold work. Excessive cold working makes itself felt in various characteristic ways, according to the shape and nature of the metal object which is being formed. A typical case is that of a rod or wire which is being reduced in diameter and increased in length by " drawing " through a draw-plate. In this process the metal is drawn through a hole, which tapers in the direction of drawing in such a way that the metal is caused to flow as it passes through the hole. In this operation, xi ,oa -o ^ ,.,..,.,,,. ,. „ Fig- 130.— Rod passing which IS indicated diaggramatically through a Drawing-die. in Fig. 130, the outer skin is not only extended longitudinally, but is at the same time subjected to considerable radial pressure. This pressure serves to " hold up " the outer skin and to enable it to endure a larger amount of deformation than the material could withstand in the absence of such support. When such a rod or wire is too severely treated it follows that fracture takes place in the interior relays, and it usually occurs at the centre where the metal is furthest removed from the supporting effect of the external pressure exerted by the sides of the hole. The result of these internal fractures is that the rod or wire " draws hollow " the internal portions breaking up into short lengths, although the outer skin may remain completely intact, as indicated in the sketch. Fig. 131. Such "over-drawn" or "hollow-drawn" material is, of course, useless and cannot be restored by annealing or by mechanical treatment. Somewhat similar x2 s 308 STUDY OF PHYSICAL METALLURGY although less perfectly defined, effects are also produced when the cold-rolling of sheet metal is pushed too far, the effect in that case generally taking the form of an exfoliation of the sheet, thin layers peehng off or crumbling away. In many cases, however, in the case of sheet metal, cracking at the edges occurs before this stage is reached. A question which cannot as yet be answered with entire certainty is, whether metals which have been extremely severely strain-hardened as the result of cold work can or do undergo any degree of spontaneous anneaUng or re-crystaUisa- tion in the course of time, even without any elevation of temperature. There can be no doubt that in the case of the softer metals, such as lead, tin, cadmium, etc., spontaneous annealing at the ordinary temperature takes place with con- siderable rapidity. On the other ^^ hand, so far as iron and steel are {""^p^^ '}> ^ ^ )^ '? concerned, the evidence is entirely in the opposite sense ; no signs of Fig. I31.-Diagram of the ^^y spontaneous anneaUng or re- Longitudinal Section of crystallisation in these metals have HoUow-drawn Rod or ^^^j, ^een observed, and, in view of the fact that definite annealing does not occur in these metals until a temperature of 500° C. is reached, it is hardly to be expected that any notable rate of re-crystallisation or of annealing can exist at the ordinary temperature. In regard to metals and alloys of an intermediate type, the state of affairs is somewhat doubtful. The most important case is that of brass, such as that used for the manufacture of cartridge cases. These, if excessively hardened during the stamping process, and particu- larly if stored in hot climates, such as that of India, exhibit a tendency to what is known as " season cracking," the brass becoming fissured in a way which suggests that a change of volume, i.e., a contraction, has taken place. This phenomenon is explained by some observers as a result of spontaneous anneahng or re-crystallisation taking place slowly in the course of months or even years, and accelerated by the slightly elevated temperature. Such " season cracking " is also met MECHANICAL TREATMENT OF METALS 309 with in other articles made of cold- worked brass, but very often these are articles which have been placed in positions where a decided rise of temperature has been experienced — as, for instance, in the chains of a gasolier — while similar articles kept in stock in a cool place have remained sound. Cohen (*) accepts the view that spontaneous annealing does occur in extremely strain-hardened brass, and that rapid re-crystalhsa- tion can be induced in such metal at the ordinary temperature by " innoculation " with a piece of the same metal in the fully crystalline condition. In some experiments described by that author an etched surface of annealed brass was pressed against the surface of a piece of cold-rolled brass, and local re-crystalhsation of the latter resulted at the points of contact. Cohen calls this— somewhat fancifully — the " strain disease " of metals, but his views have not yet attained general accept- ance, and the whole matter requires much further investigation. Meanwhile it is well, however, to bear in mind that excessively strain-hardened metals are apt to develop cracks and faults, and that, at best, strain-hardness is not a particularly reliable factor in constructive or other work. Under the heading of the mechanical treatment of metals, and more especially under that of the " cold working of metals," must be included yet another class of operations which are applied to metals on the most extensive scale, viz., the cutting of metals by means of tools of every description, but usually by what are known as machine tools. The operations of cutting tools have for their purpose the shaping of metal objects, and frequently also the production of a finished surface required for various purposes. In most respects the operations of cutting are matters which concern the engineer who has to deal both with the finished products which are turned out by the machine tools and with these machine tools themselves. The manner in which the tools actually operate on the metal is, however, a question of Physical Metallurgy. The action of all cutting tools may fairly be described as depending upon the production, over very small areas at a time, of stresses severe enough to produce local failure or breakage — generally by crushing — of the metal which is being 310 STUDY OF PHYSICAL METALLURGY worked. In the ordinary type of edge-tool this concentration of stress is brought about by the sharp edge of the tool. This really constitutes a very small area of contact, at any instant, between the tool and the work, and, since the whole pressure of the tool is concentrated on this minute area, the locally- developed stress attains a very high intensity, resulting in local fracture by crushing or shear, or both. That such crushing or shearing action really exists may be demonstrated by preparing a micro-section through the " root " of a cut taken by a machine tool, such as that used in a lathe. This can be done by suddenly withdrawing the tool while in action and then cutting a section through the " root " of the turning which was being removed. A photo-micrograph taken in this way is reproduced in Fig. 132, Plate XXIX., where the defor- mation of the structure of the steel in the vicinity of the cut can be readily traced. It is easily understood that the sharper the tool and the lighter the cut, the more completely locahsed will be the straining effects of the applied stress. If, on the other hand, very heavy " cuts " are taken with heavy tools whose edges are comparatively large in area, very large forces must be appHed to the tool and to the work, and a much wider region of dis- turbed and strained structure will result. Thus, the opera- tions of shearing and punching are really extreme cases of the operative process in all machine-tool cutting, but, owing to the fact that in these operations large masses of metal are removed simultaneously, the applied stresses are not so minutely concentrated, and relatively large regions in the vicinity of the " cut " are seriously strain-hardened. This is the reason why the material in the neighbourhood of punched holes or of sheared edges is always brittle — a circumstance long known to engineers ('). The presence of this strain-hardened material is the more injurious because it exists in the middle of, or at all events closely associated with, the general mass of normal material, while its properties are so materially different that the two cannot act together in resisting stresses, and cracks are hable to be developed in the strain-hardened regions. In the case of light cuts taken with sharp tools there is Uttle PLATE XXIX. y, ■-"A-'^v^. : Via. V42. 1 • r ( -"■^ -■ r gjn** • 1 * / t * Fig. 133. Fig. 134. [To face p. 310. MECHANICAL TREATMENT OF METALS 311 or no strain-hardening of adjacent layers — or, rather, the strain-hardened layer under the machined surface is extremely thin and its mechanical effects may be neglected. The modern •development of high-speed steels, however, has brought about the practice of taking exceedingly heavy cuts at high speeds, and where this is done the thickness of the strain-hardened layer becomes much greater. Whether, even in cases of ex- treme practice, this layer ever becomes thick enough to be of mechanical importance has yet to be ascertained, but it is well worth noticing that the difference between machining and shearing or punching is one of degree rather than of kind, and that the acceleration of the machining processes tends to make them approximate in their effects to those of shearing and punching. Should the development of high-speed cutting be carried considerably further than is at present the case, this aspect may become a serious one. The behaviour of various metals and alloys under the action of cutting tools is well known to differ widely ; some materials are easy to machine, a good surface can be produced upon them, and the cutting of screws or other more or less delicate work is readily executed. In other materials, machining is difficult ; smooth surfaces cannot be readily produced, owing to the tendency of the metal to tear up or to clog the tool. " Machinery brass," which usually contains a notable propor- tion of lead, is an example of an easily machined metal, while pure aluminium and pure copper are both examples of difficult metals. The difference is, in reality, very simply explained by differences in ductility. In an extremely ductile metal the very severe local stresses produced by the edge of the tool are not sufficient at once to cause local fracture by crushing or shear — the soft and ductile metal flows even under this extreme local pressure, and the effect is to clog the tool, which tends to " dig " into the metal or to tear it up in flakes. In order to machine well a metal must possess a certain degree of brittleness, which will allow it to fracture readily and cleanly under the stresses applied by the cutting edge, thus leaving a freely-curKng turning to flow away from under the tool. That this explanation is correct may be 312 STUDY OP PHYSICAL METALLURGY inferred from the fact that a small addition of an embrittling element immediately improves the machining qualities of an otherwise difficult metal. Ordinary soft brass is rendered slightly brittle, and enormously easier to machine, by the addition of a few per cent, of lead ; aluminium is rendered distinctly less ductile by the addition of 15 or 20 per cent, of zinc, but such an addition suffices to convert one of the most difficult metals, from the machining point of view, into one of the best and easiest. Even in the case of mUd steel, the normal ductility of the pure, or approximately pure, iron-carbon alloys is too great for easy and rapid machining in automatic machines for the production of screws and bolts, and for that purpose a " special " steel is sometimes employed which is intentionally kept high in phosphorus content ; this, by slightly increasing brittleness, greatly improves the behaviour of the steel under the machine tool. References. (1) Cubillo. Journ. Iron and Steel Inst., 1912, 1., p. 314. (2) Stead and Steadmaii. Journ. Inst. Metals, No. 1, 1914, XI. Bengough and Hudson. Journ. Soc. Chem. Ind., 1908, XXVII., 43 and 654. (3) Talbot. Journ. Iron and Steel Inst., 1913, I. (4) Hadfield. Journ. Iron and Steel Inst., 1912, II. (5) Rosenhain and Humfrey. Journ. Iron and Steel Inst., 1913, 1. (6) Stead. Journ. Iron and Steel Inst., 1900, II., p. 60. (7) Rosenhain and Haughton. Journ. Iron and Steel Inst., 1914, I. (8) Cohen. Elektrochemische Zeitschr., XVII., 181, 1910. Zeitschr. Phys. Chem., LXXI., 301, 1910. (9) Howe. Journ. Amer. Inst. Mining Engineers, 1914. CHAPTER XIV DEFECTS AND FAILURES IN METALS AND ALLOYS In the course of the general survey of the more important aspects of Physical Metallurgy contained in the preceding chapters, frequent opportunities have occurred for referring to injurious results which are liable to follow upon certain forms of treatment, or, rather, of maltreatment of metals. Since the more detailed study of the physics of metals is intended, so far as the practical point of view is concerned, to aid in lessening the liabiUty to defects and failures, this treatment of the subject throughout the present book is obviously justified. On the other hand, there are several important classes of defects whose nature and origin has either not been discussed or has not been dealt with to the extent which their importance justi- fies, so that a summary of the whole subject from this point of view is required. Throughout our treatment of the constitution and structure of alloys we have intentionally confined our attention to the pure alloys, and have, for the sake of simplicity, left impurities out of consideration entirely. In practice, however, impurities are of very great importance, since the art of the metallurgist who is concerned with the purification of metals is not suffi- ciently advanced to afford a commercial supply of really " pure " metals, although the materials at present available are far superior to those which had to be employed some twenty years ago. Curiously enough, it seems that this increase in the purity of metals, employed either as such or in alloys, has not proved an entirely unmixed blessing ; in certain industries diffi- culties have arisen which are commonly ascribed to the " excessive " purity of the modern metal. There may be a sound basis for such a view in some cases, but this simply means that the material formerly employed was really an " alloy," i.e., that the particular impurity present exerted a 314 STUDY OF PHYSICAL METALLURGY avourable influence on the resulting metal, and that the purer metal from which this alloying influence is removed has ceased to be equally satisfactory. In such cases the diffi- culty can generally be overcome by seeking the exact nature of the difference between the modern and the old " pure " metal and making a suitable alloying addition, if that be found necessary. In general, however, the term " impurity " is employed to denote a chemically " foreign " element which may be present in a metal or an alloy, and which exerts a more or less injurious influence on the resulting material. Such impurities may be either metallic or non-metallic elements present as such, or they may be compounds. The nature of their action depends upon whether they exist in one or other of the three possible states, viz. : — (a) As solid solutions ; (6) As separate constituents crystallising from the metal or alloy ; or (c) As mechanical enclosures or suspensions. In the case of the two former classes of substances, which are generally other metals or metalloid bodies (such as phos- phorus, silicon, etc.), the impurities are in reality additional constituents of the alloys, and their effects are of the same nature as those of the intentionally added elements — they really make the material an alloy of a more complex order, whose constitution and properties require careful investigation. Within the limits in which they occur in practice the effects of such impurities are in most cases well known, and, although there may well be considerable differences of opinion as to the proportions of various impurities which are admissible in materials required for specific purposes, there is general agree- ment that a definite limit to the allowable proportions should be set. In some instances this is accomplished by the specifica- tion of certain mechanical properties in the final material, while in other cases direct chemical analysis is required. There seems to be little doubt that the latter is much the safer course, and in those cases where the imposition of a chemical analysis in a specification has been shirked this is probably due to the DEFECTS IN METALS AND ALLOYS 315 influence of manufacturers who find it difficult to maintain the standard of purity which analytical results would demand. Mechanical tests depend upon many factors besides the chemical composition in regard to impurities, and direct determination of every important factor is surely the only sound plan. To enumerate the various impurities which are ordinarily met with in the whole range of industrial materials would be impos- sible in this book, so that reference will only be made to a few typical cases. Of impurities which exist in a state of solid solution, in iron and steel, phosphorus and silicon are the most typical examples. The injurious influence of phosphorus on mUd steel is well recognised, and speciflcations demand that the proportions present shall not exceed a Hmit which is flxed in the case of the most important class of articles, such as axles, tyres, springs, etc., at the low figure of 0-035 per cent., while for steel rails as much as 0-080 per cent., and in some cases even 0-10 per cent., is allowed. It would be difficult to believe that such a small proportion of phosphorus, if uniformly disseminated in a state of solid solution in the f errite of a mild steel, could produce any seriously injurious effect. Two factors must, however, be borne in mind. In the first place the phosphorus is present, not as free phosphorus, but as iron phosphide, FeaP, so that every 0-10 per cent, of phosphorus involves the presence in the steel of 0-48 per cent, of iron phosphide. Further, this iron phosphide is not uniformly disseminated throughout the steel. When the solid solution first crystallises from fusion the typical process of core-formation takes place, and the phosphorus is concentrated in the peripheral parts of the first large crystals which are formed. As has already been pointed out, when such material is rolled or forged, these phosphoric portions are elongated and the phosphorus retains this distribution, thus forming the typical banded structure which has already been illustrated (see Fig. 127, Plate XXVIII.). This condition implies that the phosphorus concentration in certain portions of the steel may be from four up to as much as ten or twelve times as great as that indicated by the average value furnished by analysis, and it is these higher local concentrations which lead to failure under test or in service, 316 STUDY OP PHYSICAL METALLURGY In the non-ferrous metals, impurities which exist in solid solution are not frequently met with, although in copper there is reason to believe that a small proportion of arsenic remains in that condition. Opinions as to the value, or otherwise, of arsenic in copper are very much divided ; some EngUsh metallurgists claim great advantages for arsenical copper, and require an arsenic content of not less than 0-35 per cent, in such materials as copper for fire-box plates in locomotives, etc. Others in this country, and the majority of foreign metallurgists, regard the presence of arsenic as definitely undesirable and prefer the use of pure copper. It is fairly evident, however, that arsenic in copper cannot be regarded with anything hke the suspicion which rightly attaches to phosphorus in steel. Impurities which separate from solution in the molten alloys and form distinct micrographic constituents are a much more numerous class, particularly in the non-ferrous metals. In iron and steel they are found only in cast-iron, where the iron-phosphide eutectic is an important example of this type of impurity. This constituent is characteristic of practically all varieties of cast-iron, where it is seen generally in the form of a granular eutectic differing slightly in colour from the pearlite with which it is often associated. An example is shown in Fig. 133, Plate XXIX. At times difficulty may be experi- enced in distinguishing between granular pearUte and the phosphide eutectic ; in such cases heat-tinting at once shows the difference very clearly. In the non-ferrous metals, impurities of this kind are met with in almost every kind of alloy. Typical examples are the com- pounds of iron and of silicon which occur in this way in even the purest of commercial aluminium. An example of these is shown in Fig. 134, Plate XXIX. In the case of copper, cuprous oxide is soluble in molten copper and forms a well-marked eutectic on freezing. The most injurious forms of impurity are frequently to be found in this class, owing to the tendency which often exists for such separate constituents to assume the form of thin brittle walls or membranes surrounding the crystals of the primary metal or alloy. The case of bismuth and gold has already been mentioned (p. 256), while bismuth or antimony DEFECTS IN METALS AND ALLOYS 317 in copper produce a similar effect. Cuprous oxide, on the other hand, does not tend to form membranes, and is consequently- harmless if not present in undue proportion — ^in fact, in ordinary- commercial copper the presence of a certain proportion of copper oxide is found essential to the good quality of the metal. The class of impurities just discussed — i.e., those -which dissolve in the molten metal, but separate as micro-con- stituents on freezing — shade off, o-wing to the incompleteness of our knowledge, into the remaining class, viz., those -which are present in mechanical suspension both in the hquid and the sohd state. The so-called " slag enclosures " of steel (sulphides and silicates, principally of manganese) are of this kind, and so are the majority of metalUc oxides and other compounds (other than inter-metallic) present in metals. In many cases it is impossible to state definitely at the present time -whether these materials possess any real solubility in the molten metals -within the range of ordinary -working temperatures. Yet the point is of considerable importance, since if these impurities are al-ways present in mechanical suspension, it should be possible to secure their removal by some process of mechanical separation, the simplest of -which is that of allowing them to float to the top or settle to the bottom of the molten metal, according to the relative densities of the metals and the impurities. In the case of steel, the separation of " slag enclosures " by this method has been successfully tried, but it entails serious cost, and is, in practice, only effected in ■ the case of steels kept in quiet fusion in an electric furnace. The mode of origin of the enclosures of steel (^) is of some interest, as it is typical of that of the enclosures in other alloys. Steel as it ordinarily leaves the open-hearth furnace or the converter contains both iron oxide and iron sulphide, either in solution or in suspension. Both these are extremely injurious ingredients, and a reducing agent — ferro-manganese — is generally added to the steel in the final stages in order to remove these substances. The result is the formation of man- ganese sulphide or, more probably, of a mutual solution of manganese and iron sulphides, and also manganese oxide, which either takes up silica from the slag or the refractory lining, or 318 STUDY OF PHYSICAL METALLURGY combines with silicon, which may be present in the steel, to form manganese silicate. If time is allowed, the globules of manganese sulphide and silicate rise to the top and join the slag to which they properly belong, but the finer particles only rise very slowly, and a great number remain entangled in the steel. Opinions still differ considerably as to the effect which these enclosures produce on the mechanical properties of steel, but in the author's experience a large number of cases of failure in service have been directly traceable to the influence of these enclosures. They are particularly harmful in steel which is to be subjected to quenching operations, as hardening-cracks frequently start from the larger enclosures. Typical examples of such enclosures, taken from cases where failure has been traced to their influence, are shown in Figs. 136 and 136, Plate XXX. The distribution of sulphide enclosures in steel is readily studied by means of a process developed by Baumann C*) from an earlier process of Heyn (^). This is known as "sulphur printing " and simply consists in pressing against a roughly polished surface of the steel in question a sheet of silver bromide paper such as that used for photographic purposes. This paper is previously soaked in a ten per cent, solution of sulphuric acid in water and the steel is thus exposed to the action of the dilute acid held in the damp paper. Under the action of this acid every particle of sulphide which is exposed in the surface of the steel is attacked and minute streams of hydrogen sulphide issue from every such particle. The hydrogen sulphide thus evolved acts upon the silver bromide, producing a dark speck of silver sulphide opposite every particle of sulphide present in the steel surface, and a direct contact print, generally known as a " sulphur print " is thus obtained which clearly indicates the distribution of sulphur in the steel. Since the various impurities of steel generally segregate more or less together, the appearance of a sulphur print serves as a guide to the presence or absence of general segregation in the steel as well as to the distribution of sulphur itself. In non-ferrous metals the conditions with regard to suspended impurities are somewhat different. Alloys such as brass or PLATE XXX. Fig. 135. Fig. 136. \_To face p. 3 IS DEFECTS IN METALS AND ALLOYS 319 bronze are not, in modern practice, prepared by any process of direct refining, such as that employed in the case of steel — the production of non-ferrous alloys is analogous rather to the production of high-class steel in crucibles in which pure materials are melted together in the desired proportions, and little or no opportunity for the formation of enclosures on any considerable scale can occur. Slag enclosures comparable to those found in steel are, therefore, unknown in non-ferrrous alloys. The presence of oxides, always in fine particles or crystals, and sometimes in an exceedingly fine state of division, is, however, a common feature of many alloys. The presence of zinc oxide in brass, and of oxide of tin in bronze (*), is known to render these metals " thick," so that they do not run freely when poured into a mould. This is due to the fact that the oxides which are present in a state of very fine division show no tendency to rise to the surface or to flux off with materials derived from the crucible or protective slag. Their removal can, however, be effected by the introduction of some reducing agent whose oxide is fusible, so that it will agglomerate into globules large enough to float to the surface. In regard to alloys of copper this question has already been discussed in Chapter VII., pp. 152, 153. Finally, before leaving the subject of impurities not dissolved in the alloy, mention must be made of " accidental " impurities which are sometimes found. Thus the writer has in his posses- sion a rod of drawn brass in the centre of which an ordinary three-cornered steel flle lies securely embedded. More serious, because more frequent, are the impurities arising from the attempted introduction into an alloy of substances which are either incapable of entering into the metal at all or have not been given long enough time or high enough temperature to pass into solution. Sometimes such " additions " simply fall to the bottom of the crucible and spoil the lower part of the melt, but at other times they become disseminated through the mass and affect it injuriously. Such cases generally occur when attempts are made to introduce into ordinary alloys any of the rarer or more refractory metals. The danger of such an occurrence can, however, be guarded against if the 320 STUDY OF PHYSICAL METALLURGY principle is adhered to that in preparing an alloy in its final condition materials of very widely different melting-points should never be directly melted together. In such cases it is most desirable to prepare intermediate " rich " alloys whose melting-points approximate to one another, and, by mixing these according to simple calculations based on the composi- tion of each, to reach the desired final composition. The defects of metals and alloys so far considered, viz., those arising from the presence of " impurities," form one class of a still wider group which embraces all defects whose nature is that of an error of chemical composition, either local or general. If the general average composition of an alloy is wrong, this can hardly be described as a defect, since we are then in reality deaHng with a totally different alloy. Yet such errors of com- position are by no means unknown, and, apart from actual errors of mixing, they are apt to occur to a greater or less degree in consequence of actions to which alloys are exposed during manufacture. Thus, during melting, losses of certain metals may occur either by direct volatihsation or by oxida- tion, or by both. The loss of zinc from brass at every melting is a well-known example, while the elimination of carbon from steel, both during oxidising melting and anneahng, is well known. Much more frequent than serious errors in general average composition, are local differences of composition from one portion of the mass to another, particularly in ingots and in the wrought products derived from them. The origin of these differences has already been considered, particularly in regard to steel, in our discussion of the solidification of an ingot and of castings in general. In non-ferrous alloys segregation may occur in a totally different manner, viz., by the separation of a constituent which appears at a relatively high temperature in the form of isolated crystals or dendrites. These, if either much denser or much Hghter than the mother-liquor in which they are formed, show a strong tendency either to sink or to float to the top. The separation of graphite from molten carburised iron in the form of " kish " is a well-known example of the latter kind, which is paralleled in non-ferrous materials DEFECTS IN METALS AND ALLOYS 321 by the tendency of the crystals of the antimony-tin compound present in certain white bearing metals to rise to the top when solidification begins. In certain lead alloys, on the other hand, the primary crystallites of lead fall rapidly downwards through the mother-liquor, forming a layer of almost pure lead at the bottom of any slowly-cooled ingot. In some aluminium alloys also, where relatively heavy crystalline compounds of aluminium with the alloying metal are formed, such segregation takes place. In these cases the only method of prevention Mes in chilUng the molten metal so as to bring about rapid solidifica- tion ; there is then no time for such gravity-separation to occur. The sources of defects in metals which have been discussed in the preceding pages may be grouped together as being " congenital," in the sense that they are present in the molten metal before its first sohdification and cannot be removed by subsequent treatment — in many cases not even by simple re-melting. The majority of other causes of defects, such as those due to errors in thermal or mechanical treatment, may be grouped together as " acquired," since they are not inherent in the metal. These can always be remedied by re-melting, and in many cases even by less drastic processes. In the course of the discussion, in the preceding chapters, of some of the principal processes to which metals are subjected we have had frequent occasion to refer to the results which follow on the more important errors of treatment ; consequently, any detailed discussion of these causes of defects is hardly required here. In order, however, to afford the reader some rather more systematic insight into the sources of defects in metals, an attempt has been made to tabulate these. This tabulation does not claim to be exhaustive ; the field of Physical Metallurgy is so large that it is impossible for any one person to be completely acquainted with the details of aU the materials and processes involved, and such an acquaintance would be required in order to draw up a thoroughly exhaustive tabulation of this kind. In some cases, therefore, it may well be found that particular sources of defects which arise in some particular material or process are not indicated in the table or P.M. Y 322 STUDY OF PHYSICAL METALLURGY are included in a somewhat vague general group. All the more important causes of defects are, however, at least represented in the table. Where the previous treatment of the subject referred to in any particular section of the table does not appear to be sufficient to explain the nature of the actions referred to, explanatory remarks have been incorporated in the table itself, while in other cases references have been made to actual examples which serve as typical illustrations of the more strik- ing types of defects. In other cases, references to the pages of this book in which the matter has already been dealt with are given. The table may thus serve as a special index to the whole subject of defects in metals in so far as these have been dealt with in the present work. Tabulation of the Causes of Defects and Failukes in Metals. I. — Chemical Composition. Errors in chemical composition may arise either from errors in the original mixture or from actions occurring during melting. (1) Larger errors in chemical composition may lead to : — (a) Unsuitable mechanical properties. This really amounts to the use of a wrong material for a given purpose. (6) Too low a softening or melting temperature, a defect which only makes itseU felt in cases where the metal is exposed to high temperatures or in special circumstances. (c) Undue corrodibility (see pp. 328—333). (2) Smaller errors of chemical composition may lead to : — (d) The formation of brittle impurities which may be — (i.) Suspended or mechanically enclosed (see p. 318). (ii.) Structurally formed, sometimes as brittle inter-crystalline envelopes (see p. 256). (e) Variations in critical temperatures leading to subsequent abnormal behaviour under thermal or mechanical treatment. This type of defect appUes particularly to the class of " special " or alloy steels. II. — Treatment. (A) Oasting. (a) Overheating the molten metal, leading to change of chemical composition either by loss through volatilisation, oxidation or reaction with the containing vessel, etc. Absorption of gases, with subsequent liberation during freezing, may also result. PLATE XXXI. Tig. 1,37. Fig. l:!S. Fig. 139. Fig. 140. yVi) face p 323. DEFECTS IN METALS AND ALLOYS 323 (6) Excessively high casting temperature, leading to unduly slow cooling, coarse structure and liability to segregation (see p. 287). (c) Too low a casting temperature, leading to incomplete filling of the mould, " cold shuts " and similar defects. (d) Inadequate cleaning or skimming of the molten metal, leading to the admixture of slag or dross. (e) Admixture of foreign matter from the mould, such as particles of sand, gas bubbles, etc. These defects are due to bad moulds, insufficient drying, inadequate " venting," or to unduly vigorous pouring of the metal resulting in the disintegration of the surface of the mould. The correct design of the pattern to suit the metal employed also plays a part in this connection. (/) Errors in the design of the pattern or the mould leading to " draw- ing," unsoundness and shrinkage cracks (see p. 292). Here, again, the adaptation of the pattern to the nature and pecuhari- ties of the metal employed plays an important part, since patterns which can be readily cast in one metal sometimes yield unsatisfactory castings with other metals. (g) In ingots the corresponding causes operate to produce piping, blow -holes and segregation (see pp. 295 — 297). (B) Hot Working. (Forging, pressing, rolling, hot-stamping, etc.) (a) Commencing at too high a temperature. The metal is then too weak and brittle, and the formation of cracks and fissures results ; these are sometimes partially closed or covered over by subse- quent operations, but reveal their presence when fracture occurs. (b) Finishing at too low a temperature. This results in more or less severe strain-hardening, which may or may not be objectionable (seepp. 299— 301). (c) Excessively rapid deformation or " reduction." This may lead to cracking of the metal. (d) Too gradual an application of work. It is not certain whether this leads to ultimate defects, but in modern industrial practice it is not likely to occur, since the tendency is to accelerate all operations to the utmost limit. (e) Insufficient total amount of " work " or reduction. This is apt to result from a desire to effect economy by the use of an ingot which requires less forging to bring it into the desired final shape. This results in a merely superficial working of the metal, leaving the interior very coarse and weak. A typical case is illustrated in Figs. 137 and 138, Plate XXXI., which show the micro- structure of the outer shell and of the interior portions of a large steel shaftwhich failed in service. Had the coarse internal structure been due to over-heating it would have extended to the surface and signs of serious decarburisation (see p. 326 (e), below) would have been found. y2 324 STUDY OF PHYSICAL METALLURGY (/) Defects arising from mechanical causes during hot working, leading to such features as " laps," " rokes," etc., which result from the partial welding up of fissures or of portions of metal which have become accidentally overlapped. These defects should be detected by inspection before the metal is taken into use. {g) Defects due to working into the surface of the metal of scale or other matter left on the surface during rolling or forging. WhUe it is comparatively easy to brush away scale, etc., from the upper surface of metal passing through rolls or under the press, scale may adhere to the under side, which is not in view, and this may become roUed or pressed into the surface. (C) Cold Working. The operations of cold working may produce defects by being : — (a) Too violent. The use of either too high a speed or of too great an amount of reduction, either at any one pass or by a number of successive passes appUed without intermediate annealing, may be grouped together as unduly " violent " cold work. The resulting defects are either cracks on the surface or at the edges in the case of sheet metals, or in the case of rods or wires they take the form of " hollow drawing." Sometimes an alloy is employed for cold working which is reaUy unsuited to the pro- cess, such as a brass containing any considerable quantity of lead. If such a material is cold-drawn or roUed only very minute cracks are formed, not visible at the surface, but readily detected in other ways. As a general rule it may be stated that cold working may be continued, without permanent damage to the metal, so long as the tensile strength continues to rise. (b) Too slight. This is not only a cause of defect in those cases where the additional stiffness due to strain-hardening is rehed upon (see p. 300), but may cause excessive superficial haxdening and severe internal stresses. (c) Unequal, portions of the same piece of metal being much more severely treated than others. This leads to a state of severe internal tension, and may lead to the development of " season cracks " (see p. 308), or it may cause the metal to crack on anneaUng, owing to the unequal rates of re-crystaUization in the differently strained portions. (d) Unintentional, inadvertent or accidental cold working. This very frequently occurs in the fitting and erection of engineer- ing objects, such as large boilers, bridges, roofs and other struc- tures. Where things do not fit sufficiently well, they are " set " by forcible means until a fit is obtained. In other cases, minor irregularities of shape or surface are abolished by such processes as "hammer dressing." While such processes are employed, as a rule, in good faith and in complete reliance on the known DEFECTS IN METALS AND ALLOYS 325 ductility of the material in question, the result is often a highly dangerous condition of severe local cold-working and consequent strain-hardening. The dangerous nature of such defects arises from the inability of the material at those points to withstand further deformation or working to Which it may be subjected, and its inability to co-operate with the ductile material around it. A striking example is illustrated in Fig. 139, Plate XXXI., which represents the structure of a locally severely strain-har- dened portion of a steel plate which formed the domed end of a vessel which exploded disastrously under steam pressure. The fracture took place at the point where the severely strain- hardened steel joined the surrounding ductile material. The local strain -hardening in this case was due to an unduly severe " fitting " operation in the workshop. The ill effects of local strain -hardening also occur in connection With such operations as punching and shearing (see p. 310). (D) Thermal Treatment. Annealing or Ee-heating may cause defects owing to being carried out — (a) At too high a temperature, leading to — (1) " Burning," and (2) "Overheating." An example of " burning " in steel is shown in Fig. 140, Plate XXXI., which represents the structure of a portion of a steel shaft which had been "burnt," the typical enclosure of oxide at the boundaries of large grains or cells being well marked. An over- heated structure taken from a steel forging which failed in service is shown in Fig. 120, Plate XXVI. (6) For too prolonged a period, even at a moderate temperature. This treatment tends to cause an aggregation or " balling up " of some of the most important micro -constituents, and especially of euteotoid bodies such as the pearhte of mUd steel. A case of failure resulting from such treatment is illustrated in Fig. 115, Plate XXV., which refers to a boiler plate which developed typically inter-crystaUine cracks. When etched with sodium picrate, which stains cementite black, this steel is seen to con- tain little or no pearhte, the cementite having become " balled up " into inter- cry staUine masses. These appear as dark fila- ments in Fig. 119, Plate XXVI. (c) In an injurious atmosphere. Copper is injured by " gassing," i.e., by exposure at high temperatures to a reducing atmosphere (see p. 272). In other metals an oxidising atmosphere is frequently injurious, the oxidation and decarburisation of steel being well known. 326 STUDY OF PHYSICAL METALLURGY Quenching may cause defects if it is carried out — (o) From too liigli a temperature. This results, in the case of steel, in the development of a coarse martensitic structure with which much brittleness is associated. Too high a quenching tempera- ture also involves considerable risk of cracking and warping. (6) From too low a temperature. In that case the entire operation fails to attain its object, and the metal remains soft or otherwise unaltered. (c) Unequally, so that the rate of cooling is different for different parts of the same object. This generally results in serious warping. Even such large objects as heavy guns are quenched by vertical immersion in a deep oil-bath in order to avoid cooUng one side before the other. (d) In an unsuitable bath. The quenching effect depends upon the cooling power of the liquid as well as upon the temperature of quenching, although the most powerfully cooling liquid cannot compensate for too low a quenching temperature. In a liquid, high specific heat, low viscosity and low vapour pressure tend towards great coohng power. Water is therefore particularly powerful, while mercury, in spite of a higher thermal conduc- tivity, is far less effective. Tempering may lead to defects if carried out — (a) At too high a temperature. The result is to remove the effects of quenching too completely. (ft) At too low a temperature, whereby the effects of quenching are insufficiently removed, and the metal (in the case of steel) remains too brittle for its purpose. (c) Unequally. This rarely occurs in small objects, such as tools. Where, however, large masses of steel are subjected to special heat treatment by quenching followed by tempering, the attain- ment of a uniform condition throughout the mass is of the highest importance, particularly if the object is a gun or an armour plate or a part of a motor-car or aeroplane subjected to severe shocks. Want of uniformity in such cases leads to a concentration of stresses with resulting failure. The ills to which metals are heir, which have been very briefly discussed and illustrated in connection with the above tabulation, do not by any means constitute a complete list of the causes of failure in metal objects employed in engineering construction, and brief mention of some of the other factors should be made here. The causes of failure, which have been treated above as fully as the introductory character of this work will permit, all reside in the metal itself or in the treat- DEFECTS IN METALS AND ALLOYS 327 ment which it has received during manufacture. At one point only (under II. (C), {d) ) have we referred to a class of defects due rather to the abuse of the material by the constructor than to the production of the metal itself. There are, however, numerous other cases in which " abuse " of the material may fairly be put down as the cause of failure. The use of a wrong material or insistence on a wrongly-drawn specification may equally be regarded as abuse of the metal ; errors of design and of construction come under the same head, and these constitute the most difficult cases which have to be dealt with in the investigation, by the methods of Physical Metallurgy, of cases of failure in practice. The investigator can in such cases only endeavour to exhaust the possibilities with regard to defects in the metal itself, in order to show that the fault must lie elsewhere. Another class of failure in metals arises from want of per- manence of the material under the conditions to which it is exposed. We have already considered such matters as the gradual breakdown of metals under " fatigue " or alternating stresses which exceed the true elastic limits. Another type of gradual disintegrating use is that where the metal is subjected to abrasion. Such objects as the rails and tyres of a railway, the bearings of engines, etc., constitute examples of this class ; the conditions of wear existing in such circumstances are, however, fully recognised in engineering practice, and " failure " is only regarded as arising where the rate of wear has been unduly rapid. Par more serious and difficult in every way is another type of " wear " to which many metals are subjected, viz., the whole class of actions known as " corrosion." This subject is so large that it requires treatment in a separate volume, but a brief reference to it is required here. Since we find practically all the industrially useful metals in nature in the state of their oxides, or compounds of their oxides or sulphides, it is not surprising to find that there is a strong tendency, which requires to be constantly guarded against, for all metals when in contact with the agents of nature, viz., oxygen, water and carbon dioxide, to revert to their 328 STUDY OF PHYSICAL METALLURGY condition of chemically stable equilibrium and, therefore, to undergo one or other of the processes of corrosion, all of which tend to the formation of oxides or of salts. The metaUic state, in the conditions which prevail at the surface of the earth, is for the majority of metals a state of chemical meta-stabihty, and there is room for astonishment that we should be able to defeat the action of the atmosphere to such an extent as to maintain in prolonged existence the great quantities of metal on which modern engineering depends. The general " cause " of corrosion, therefore, is not far to seek, but in regard to its immediate mechanism and the detailed conditions which govern its incidence and rate of progress, a great deal of study has been expended, although definite and finally conclusive results have only rarely been reached. The possibihty of the continued existence of oxidisable metals in the metallic state depends upon that universal property of " inertia " which manifests itself in a reluctance to undergo any change, and more particularly in the reluctance with which the commencement of any change is initiated. In many cases, therefore, the study of the " causes " of corrosion resolves itself into a study of the special conditions which facilitate the commencement of chemical action between the metal and the natural agencies to which it is more or less exposed. In most questions of corrosion two different views may be taken, which are typified by the two schools of thought in regard to the corrosion of iron and steel. One of these schools, headed by Cushman (^) and Walker {^) of America, takes the view that the corrosion of iron and steel is essentially an electrolytic process, the source of the electrolytic action lying in the minute galvanic cells which are formed by the impurities present in the steel. This school, therefore, seeks to render iron and steel more or less incorrodible by making it as pure as possible, and there seems to be some evidence that extremely pure iron possesses very great " inertia " in regard to corrosion. Opposed to this view is the purely chemical theory, according to which iron undergoes direct oxidation in the presence of water, without the necessary intervention of any electrolytic actipn. In support of this view are the DEFECTS IN METALS AND ALLOYS 329 experiments of J. N. Friend (') and the work of Heyn («), who have shown that, although the absence of electrolytic cells in a very pure steel does retard the commencement of corrosion, it does not affect the rate of corrosion once the process has begun. Whichever of these rival views should ultimately prove correct, the prospect of producing a cheap form of iron or steel which shall be practically incorrodible is extremely remote. The ancient iron of India and Ceylon is sometimes quoted as an example of practically incorrodible metal, but the permanence of the Delhi column and of the iron chains fastened, thousands of years ago, on the steep pilgrim's paths on Adam's Peak in Ceylon, must be due to some special cause, since samples of this iron, which have come into the author's hands, have rusted freely on the surfaces exposed by cutting them up. The infer- ence is that these irons are not really incorrodible, but were originally covered with a really effective protective coating. Although a definite proof has not been given so far, it seems probable, from the author's observations, that this protective coating was simply a coating of cinder or slag derived from the crude manufacturing process employed (^). In any case it is evident that, even if the purest possible iron really proved to be incorrodible, its use would be limited by its softness and weak- ness as compared with the better grades of steel. The principal question of practical importance in regard to the prevention of corrosion in iron and steel thus resolves itself into a question of finding a suitable protective coating which shall effectively protect the metal from corroding actions. Oil paint, as ordinarily made and used, is known to be far from perfectly effective, and not only requires constant renewal, but also allows corrosion to take place beneath it. This is principally due to the fact that there is a certain degree of solubility for moisture in the vehicles of oil paints. Many other protective coatings have been proposed and tried, with varying success, ranging from such substances as bituminous varnishes to coatings of other metals applied in various ways. Of these latter a coating of tin, applied in the molten state, is much used, and is very efEective where the coating is perfectly intact. Its expense, however, precludes a very wide extension 330 STUDY OF PHYSICAL METALLURGY for engineering purposes. Coatings of zinc — under the name of " galvanised " iron or steel — are much used, and are effective up to a certain point, particularly as the zinc acts as an electro- lytic protector for any portions of the steel which may become exposed. In time, however, the zinc coating is entirely corroded away, and then attack upon the exposed iron is rapid. A particularly interesting form of " protective " coating is that formed by a layer of Portland cement or con- crete. Steel embedded in this material, which is now so extensively used in ferro-concrete constructions, appears to be perfectly protected from corrosion so long as the concrete retains an alkaline reaction ; any water percolating the concrete finds sufficient free lime to be rendered alkaline, and, therefore, harmless to the steel. In running water, however, the free lime tends to be washed out of the concrete, and it appears probable that ultimately corrosion may set in. A similar action may also occur in the sea— a circumstance which is not without its grave significance in view of the use of ferro-con- crete for marine and hydraulic work in many places. It may be hoped that in such places means have been used to render the concrete as nearly watertight as possible in order to mini- mise the circulation. An extremely serious source of corrosion for all metals is found in stray electric currents which are apt to seek an easy path through the earth by passing into any metallic bodies which may be present. If moisture is present, electrolysis, with rapid destruction of the metal, is set up if the current happens to pass in the unfavourable direction. Such electric currents should be particularly guarded against in the vicinity of ferro-concrete which is liable to be moist. Electrolytic corrosion in ferro-concrete not only results in the destruction of the embedded steel, but, owing to the large bulk of iron oxide which is formed, rapidly disrupts the whole' mass. The corrosion of non-ferrous metals and alloys has been principally studied in regard to brass, although corrosion is, of course, encountered in practically all metals, with the exception of the noble metals (silver, gold, platinum, etc.), and the newly-developed " rare " metals, such as tungsten, which are apparently incorrodible at the ordinary temperature. The DEFECTS IN METALS AND ALLOYS 331 corrosion of brass, however, has assumed exceptional import- ance owing to the large number of brass tubes employed in the construction of condensers for steam engines, more especially for marine purposes. It has been found that in many cases, although not in all, the corrosion of brass takes place by a process which consists in the selective solution and removal of the zinc, leaving the copper behind in the form of a spongy mass containing a considerable proportion of oxide. This process is generally known as " dezincifica- tion," and its mechanism has been recently explained (}"). The action takes place, in the case of brass exposed to contact with sea water, particularly at temperatures above 40° C, by the formation of a basic chloride of zinc, which adheres to the sides of the brass tubes. On the one side this basic chloride constantly reacts with the sea-water, parting with some of its zinc in a soluble form, either as chloride or sulphate, and, on the other side, constantly renewing its own zinc-content at the expense of the brass. The adherent patches of this basic chloride thus act as a species of catalytic agent. Fortunately it has been found that this action is largely inhibited by the presence in the brass of 1 per cent, of tin or, better still, of 2 per cent, of lead. The protection of brass and of other copper alloys from corrosion in condenser tubes and other situations where protec- tive coatings are out of the question, by the aid of specially generated electric current, has recently been suggested and adopted with a certain measure of success. The simplest method of applying this principle consists in generating the requisite electric current in situ, by attaching to the article to be protected pieces of a more readily attacked metal. The combination of the two metals then acts as a primary cell, in which the " protector " is steadily dissolved while the protected metal remains intact. Thus plates of iron, zinc or aluminium have been attached to brass tubes and have served as efficient protectors where there has been sufficiently good electric contact between the two metals and for a moderate distance from the point of contact. In some cases, as Bengough (^'') has shown, the protection is effective over a longer distance, 332 STUDY OF PHYSICAL METALLURGY owing to a secondary effect whereby a thin but strongly adherent skin of calcium carbonate is deposited on the surface of the brass, thus protecting it from the sea-water. The principle of electrolytic protection has also been applied in another way, by employing a definite and independent source of current and thus converting the whole apparatus to be protected— boiler, condenser or other apparatus — into an electrolytic cell, in which specially introduced electrodes of some such substance as carbon serve as anodes, while the whole apparatus acts as cathode. This arrangement, if carried out with sufficient attention to the electrical contacts and resist- ances, gives promise of working very efficiently, although the maintenance of an outside source of electric current is some- what of an inconvenience. On a ship, however, where electric current is now universally available, this inconvenience is extremely sKght as compared with the saving in condenser tubes and the troubles of frequent repairs resulting from per- foration of tubes by corrosion. Finally, reference must be made to the question of the corrosion of aluminium and of its light alloys ; at the present time the use of these materials is still very restricted, but this is due to special causes, one of which is the fear of want of permanence, on the ground both of corrosion by external agencies and of internal disintegration. In regard to corrosion, aluminium itself has long suffered from the excessive zeal of its early friends, who claimed for it practical incorrodibility. This claim probably arose from the fact that brightly-poUshed aluminium can be exposed to the air for long periods without showing signs of material corrosion, but under the action of sea-water, and particularly in the presence of other metals, aluminium corrodes with decided rapidity. The action of the sea in this respect is largely a mechanical one. Aluminium when exposed to the air rapidly becomes covered with a very thin film, which either consists of oxide or contains a consider- able proportion of oxide ; this film serves as a protective coating which, if undisturbed, serves to prevent further corrosion almost indefinitely. When, however, as friction acts on this film it is, after a time, worn away, and, if the DEFECTS IN METALS AND ALLOYS 333 friction continues, the fresh formation of a protective film is prevented and corrosion proceeds rapidly. Pure aluminium, however, owing to its relatively small strength and stiffness, is of little value for structural purposes, and recourse must be had to one of the strong and light alloys which have been produced, as the fruits of Physical Metallurgy researches, in recent years, both in this country and in Ger- many (^^). Since these materials rival the strength of steel while only one-third the weight, some considerable importance attaches to their behaviour in regard to corrosion. Those alloys which contain manganese have been found to be rather less liable to rapid corrosion than pure aluminium, but the fact must none the less be faced that, if exposed to the action of sea-water, these materials require protection approximately to the same extent as ordinary iron or steel ; for mere exposure to a moist atmosphere, however, the aluminium alloys are far superior to polished steel in regard to their power of main- taining a bright, untarnished surface. The second factor which still stands in the way of the wider usefulness of aluminium alloys is the prevalent fear that they may not be permanent in their properties, quite apart from corrosion. This belief has no doubt sprung from the behaviour of certain alloys of aluminium with iron, tin, nickel, and some other elements. These exhibit the peculiar property of spontaneous disintegration ; a small ingot of such an alloy will, if left to itself, fall into a heap of fine powder in the course of a few hours {^^). This behaviour, however, is a property of certain definite compounds of aluminium with other metals, and these compounds only exhibit this behaviour when they constitute the whole, or nearly the whole, of the alloy. This only occurs in alloys in which the element other than aluminium is present in comparatively large quantities, so that the whole process is probably unknown in true " light " alloys which consist principally of aluminium. For certain purposes, especially for a process known as "die casting," some alloys containing a large preponderance of another metal, such as zinc or iron, have been used, and the product, which has a white ap- pearance, has been sent out as " aluminium castings." These 334 STUDY OP PHYSICAL METALLURGY have shown frequent cases of warping and disintegration, and the unfortunate experiences of those who have had to deal with them have reflected on aluminium alloys as a whole. In order to settle this question as definitely as possible, an investi- gation on a very large scale has been undertaken at the National Physical Laboratory. Over a thousand different specimens of aluminium alloys, both pure and containing intentionally added impurities, have been prepared in all conditions of treatment, and are being kept under various conditions of storage as regards temperature and moisture. These specimens have been prepared for accurate measurement, both as to dimensions and electrical resistance, and, by periodical observa- tions, are being watched with a view to detecting the slightest signs of change. A preliminary set of castings, which have been watched for over a year, have so far shown no signs whatever of any change, and to this extent the popular belief in the disintegration or warping of certain aluminium alloys is already discredited by the results of accurate investigation. That this is likely to be the case with whole groups of important alloys is indicated by the fact that samples of some of these materials have been tested at the National Physical Laboratory at intervals over a period of eight years without showing any signs of deterioration. Although the principal interest in this connection attaches to the light alloys of aluminium, some reference must also be made to the very great powers of resisting corrosion and oxida- tion which are displayed by the alloys of aluminium with copper in which the latter metal preponderates. These aluminium coppers, formerly, but quite erroneously, called " aluminium bronzes " i}^), have been tested by prolonged exposure to the sea, first for a period of three years by complete immersion, and then for a similar period " between wind and water." These tests show that the alloy containing approximately 10 per cent, of aluminium is extremely resistant to sea-water corrosion, the test-pieces having lost in weight to a minute extent, which could be detected only by the aid of accurate weighing on a balance, at the end of that period. Alloys of this type can also be exposed to a red heat in an oxidising atmosphere for several DEFECTS IN METALS AND ALLOYS 335 days without becoming appreciably tarnished by oxidation. No doubt the formation of a pellicle of aluminium oxide, supported by the great mechanical strength of these alloys, is responsible for this satisfactory behaviour, which places them above any other alloys of copper in these respects. Here, as elsewhere, however, a compensating disadvantage is to be found. The pelUcle of aluminium oxide which protects these materials from corrosion or oxidation also protects them from alloying with tin, so that fchey cannot be readily soldered like other alloys of the brass or bronze type. Refekenoes, (1) Roseuhain. Internat. Testing Assoc, Copenhagen Congress, 1909, 1., 4. Internat. Testing Assoc, New York Congress, 1912, II., 2. Le Chatelier. Bull. Soc d' Encouragement, September, 1902. Arnold and Waterliouse. Journ. Iron and Steel Inst., 1903, I. Andrews. Engineering, July, 1906. Stead. Iron and Steel Magazine, IX., No. 2, February, 1905. Howorth. Journ. Iron and Steel Inst., 1905, II. Hoyn and Bauer. Bericht d. Materialprufungsamt, 1906, XXIV., p. 233. Fay. Amer. Soc. Testing Materials, III., 1908, and VIII., 1908. Law. Journ. Iron and Steel Inst., 1907, II. Ziegler. Rev. de M6taUurgie, 1909, and 1911. MatweiefE. Rev. de M6tallurgie, 1910, p. 447. Levy. Journ. Iron and Steel Inst., 1911, III. Rohl. Journ. Iron and Steel Inst., 1912, IV., p. 28. (2) Baumann, Metallurgie III., p. 416. (3) Heyn and Bauer, " MetallograpMe " II., p. 130(G6sclien, Leipzig, 1909). (4) Heyn and Bauer. Zeitschr. Anorg. Chem., 1905, XLV., p. 52. (5) Cushman. Journ. Iron and Steel Inst., 1909, 1., p. 13, and many other papers. (6) Walker. Journ. Iron and Steel Inst., 1909, I., p. 13, and many other papers. (7) Friend. Journ. Iron and Steel Inst., 1908, II ; 1909, II. ; 1911, III ; 1912, IV. ; 1913, IV., and other papers. (8) Heyn and Bauer. Journ. Iron and Steel Inst., 1909, f., p. 109. Mitteilungen a. d. KSnigl. Materialpriifungsamt, Berlin, 1908, XXVI., p. 2, and 1910, XXVIII., p. 62. (9) Rosenhain. Faraday Soc Trans., July, 1913. (10) Bengough. Journ. Inst. Metals, 1913, No. 2, X. 336 STUDY OF PHYSICAL METALLURGY (11) Rosenhain and Archbutt. Proc. Inst. Mechanical Engineers, 1912. Wilm. Metallurgie, Ser. VIII., No. 8, AprU, 1911. (12) Law. Faraday Soc. Journ., June, 1911. Rosenhain and Lantsberry. Proc. Inst. Meclianical Engineers, 1909. (13) Report of Nomenclature Committee, Journ. Inst. Metals, 1914, 1. INDEX OF NAMES Abel, 119 Andrews, 13, 15 ;, 159, 335 Arohbutt, 15, 106, 136, 158, 239, 335 Arnold, 13, 175, 192, 231, 239, 240, 256, 264, 278 Arnold and Waterhouse, 335 Batrstow, 212, 240, 264 Baker, 14 Barrett, 129 Bauer, 159, 329 Baumann, 318, 335 Bauschinger, 212, 224, 240 Baykoft, 174, 192 Beilby, 24, 25, 37, 246, 249, 264, 269, 270, 285 Benedicks, 60, 170, 183, 192, 219, 240 Bengough, 158, 264, 331, 335 Bijl, 129 Blough, 146, 159 Bramah, 14 Brillouin, 264 BrineU, 218, 240 Brown, 129 Brunton and Brown, 14 Burgess, 106, 168, 192 Cabpentbk, 106, 141, 164, 156, 158, 159, 192, 285 Cartand, 77 Chamberlain, 108 Charpy, 12, 108, 126, 128, 130, 158, 233, 240 Cohen, 309, 312 Coker, 237, 240 Crowe, 192 P.M. Cubillo, 312 Curie, S., 129 Curry, 154, 159 Cushman, 328, 335 Dalbt, 202, 208, 212 Degens, 132, 158 Dumas, 129 Dumont, 129 Edwards, 154, 156, 157, 158, 159 Emery, 212 Ewen, 108, 129, 260, 264 Ewing, 14, 77, 110, 129, 203, 212, 254, 264, 269, 285 Fairbairn, 14 Fay, 335 Friend, 329, 335 Galileo, 13 Garland, 255, 264 Gibbs, 129 GioUitti, 35, 37, 146, 159 Goerens, 271, 285 Greenhough, 59 Grenet, 12, 108, 128 Guertler, 108, 110, 128, 129, 264 GuUlaume, 129 GulHver, 130 Gutowsky, 163, 192 Gwyer, 159 Hadfield, 129, 255, 264, 296, 312 Hannover, 134, 158 z 338 INDEX OF NAMES Haughton, 35, 37, 108, 129 Haupt, 129 Hensler, 109 Herrschkowitz, 129 Heussler, 128, 129 Heycock, 37, 146, 148, 159 Heycock and Neville, 17 Heyn, 12, 35, 37, 60, 159, 264, 318 Heyn and Bauer, 335 Hodgkinson, 14 Hopkinson, J., 129 Howe, 13, 192, 312 Howorth, 335 Hoyt, 146, 149, 150, 159 Hudson, 158, 159 Humfrey. 170, 180, 181, 192, 198, 212, 248, 254, 264, 312 IzoD, 233, 239, 240 jijptnee von jornstokff, 129 Keeling, 106 Kennedy, 14 Kirkaldy, 14 KroU, 281, 285 Kurnakoff, 106 Lababdie, 14 Lantsberry, 130, 157, 159, 335 Larrard, 213, 239 Laschtschenko, 129 Laurie, 129 Law, 159, 335 Lebasteur, 212 Le Chatelier, 12, 37, 58, 60, 129, 335 Levy, 336 Liidekin, 128 Ludwik, 222, 240 Maet, 128 Martens, 12, 16, 37, 60, 198, 201, 204, 212, 220, 222 Matweieff, 335 Mazotto, 132, 158 Morris, 129 Muir, 212, 264, 271, 285 Murray, 108, 128 Neville, 37, 146, 148, 159 Omodei, 128 Osborne-Eeynolds, 228, 240 Osmond, 12, 33, 37, 77, 84, 106 Pashkt, 129 Perronet, 14 Primrose, 159 Pushin, 129 Eeindees, 129 Retgers, 128 Roberts-Austin, 13, 84, 106 Rohl, 335 Eoozeboom, 130 Rose, 269, 285 Eosenhain, 35, 37, 43, 60, 77, 106, 128, 130, 158, 159, 192, 212, 264, 269, 285, 312, 335 Eoudelet, 14 Ruff, 192 Sanitee, 175, 192, 240 Sankey, 231, 240 Sauveur, 13, 192, 267, 279, 285 Sears, 264 Shepherd, 141, 148, 158, 159 Sheppard, 146 Shore, 222, 240 Smith, 240 Sorby, H. C, 12, 17 Stansfield, 106 Stanton, 228, 236, 240, 264 Stark, 129 Stead, 13, 34, 37, 60, 66, 77, 158, 275, 285. 304, 312, 335 Steadman, 158, 312 INDEX OF NAMES 339 Take, 129 Talbot, 296, 312 Tammann, 13, 103, 106, 129 Tavaati, 159 Tomlimon, 129 Topler, 128 Towne, 212 Tucker, 77, 158, 264 Turner, 108, 128, 129, 222, 240 Unwin, 14, 15, 212 ViNCENTIKI, 128 Vogel, 264 Walker, 328, 335 Weintraub, 159 Weiss, 110, 129 Werth, 12 White, 106 WUm, 159, 335 Wohler, 224, 240, 254 Wratten and Wainwright, 46 Wiist, 12 Zeiss, 39, 46 Ziegler, 281, 285, 335 Z 2 SUBJECT INDEX Aberrations of objectives, 46 Abrasion resistance of aluminium- manganese-copper alloys, 157 Abuse of metals, 327 Aci, Aca, Aca, defined, 172 Acj, annealing steel above, 279 Aci28, defined, 175 Acas. defined, 174 Accidental impurities, 319 Acetylene blowpipe — avoidance in cutting specimen, 20, 21 Acicular constituent in quenched steel, 178 Acid, hydrochloric, 31 hydrofluoric, 31 nitric, 31 picric, as etching reagent, 31, 173 sulphuric, 32 Acquired defects, 321 Action of cutting-tools, 309 Additions insoluble in alloys, 319 Adhesion of adjacent crystals, 68 Alcohol, amyl, 31 ethyl, 31 use in drying specimens, 36 Alkali metals, 108 AUotropio modifications of iron, 166 theory of hardness of steel, 180 transformation of iron affected by carbon, 177 Allotropy of iron and decomposi- tion of iron solid solution, 166 Alloy, eutectic, 76 steels, 192 Alloy — continued. steels — continued. in y-iron state, 284 heat treatment of, 284 steel, Martensitic, 284 systems, typical, 131 Alloys, binary, 72 chemical study of, 118 containing a compound, 100, 101 electrical properties of, 110 forming compounds, con- ductivity curves of, 114 Heussler magnetic, 109 of iron and carbon, 131 magnetic — Heussler, 128 magnetic properties of, 109 mechanical properties of, 101 micro -structure of, 74 molten, as solutions, 73 partly eutectiferous, 99 constitutional dia- gram, 99 physical properties of, 107 Research Committee, 239 ternary, 123 transformations in solid, 102 " Alpha " hardness tester, 222 Alpha iron, 167 Alternate bending test, Arnold's, 231 Alternating shear test, 229 stress curves, 227 safe range of, 226 342 SUBJECT INDEX Alternating — continued. stress fractures, 264 test, Osborne-Rey- nold's, 228 Stanton's direct, 228—9 Wohler's, 225 stresses, 224 torsion test, 230 Alumina, 29 in aUoys, 153 Aluminium, 81, 110 absorption of silicon by, 289 action of sea on, 332 alleged incorrodibility of, 332 alloys and spontaneous dis- integration, 158, 333 etching of, 31 light, 156—157 research on, 334 segregation in, 321 and zinc, machining of, 312 annealing of, 271 as, deoxidiser, 153 " bronzes " 334 castings, 333 compounds, 333 -copper alloys, 154 -copper alloys (heavy), and corrosion, 334 resistance to oxidation, 334 soldering of, 335 -copper alloys, micro-structure of, 156 phase-fields of, 155 -copper-manganese alloys, 157 -copper, overheating of, 277 -copper-tin alloys, 157 corrosion of, 332 etching of, 31 in copper, 153 in steel, 295 iron and silicon in, 316 -manganese-copper, liquidus model of. 126 Aluminium — continued. oxidation with mercury, 29 polishing of, 24 solubility of copper in, 156 -tin-copper alloys, 157 -zinc alloys, 136 comparative tests on, 238 cooling curves of, 140 phase-fields of, 138 thermal analysis, 139 Alundum, 81 Amalgams, specific volumes of, 108 Amorphous cement and annealing, 267 intercrystal- line, 69, 257 condition of metal, 25 defined, 249 iron-carbide solution in hardened steel, 182 layers in strained metal, 246 phase in hardened steel, 181 theory of, 250 surface film, 34 theory, 246 and strain at high temperatures, 259 of hardening steel, 180 Amsler tensile testing machine, 201 Analogy between metals and ig- neous rocks, 17 Analysis, specification of impuri- ties by, 314 Ancient metal, 255 meta-stability of, 122 Angle, re-entrant, in castings, 292 Angular pearlite, 282 Annealing, 265 duplex alloys, 269 ' duration of, 271 effect on micro -structure, 273 of aluminium, 271 SUBJECT INDEX 343 Annealing continued. and amorphous theory, 266 cadmium, 269 copper, 269, 270 defects arising in, 325 in determination of constitu- tional diagram, 105 gold, 269, 270 and growth of crystals, 267 insufficient, 268 iron and steel, 270, 271 lead, 269 and micro-structure, 266 and oxidation, 272 prolonged, 268, 272 and physical properties, 268 and re-crystallisation, 266 soft metals, 269 spontaneous, 308 steel above Acs, 279 below Ari, 277 temperature, 266, 270 tin, 269 very mUd steel, 279 zinc, 269 Antimony, 72, 112 -copper, 114 in copper, 317 -lead alloys, 134 and lead, alloys of, 131 -tin, 321 Aperture of objectives, 46, 51 Arnold on gold-bismuth alloy, 256 Arnold's alternate bending test, 23 1 determination of eutec- toid point in steel, 175 test, 239 Ari, annealing steel below, 277 Ara, Ars, defined, 172 Ari28 defined, 175 Aras defined, 174 Arrest-points, heat evolved or absorbed at, 103 of steel, names of, 172 with 0-6 per cent of carbon, 174 Arsenic in copper, 316 Atmospheric corrosion, 333 Atomic volume, 110 Austenite defined, 178 Austenitic alloy steel, 284 Automatic machines, steel for, 312 Axles, phosphorus in, 315 Bairstow and Stanton on alter- nating stress fractures, 264 BaUed-up cementite in steel, 325 BaUing-up in duplex alloys, 276 of pearlite, 278 Banding, phosphorus in steel, 237, 304, 316 Base metal thermocouples, 80 Bath for quenching, 326 Baumann on sulphur-prints, 318 Baykofi on hot -etching of steel, 174 Bearing metals, copper phosphide in, 152 white, 135 BeUby on annealing copper, 269, 270 gold, 269, 270 BeUby's amorphous theory, 246 definition of " amor- phous," 249 Bending as ductility test, 217 machine, hand, Sankey's, 231 test, 216 alternating, Arnold's 231 Benedicks' formula for hardness number, 219 on Arj, 170 on Troostite, 183 Bengough on corrosion of brass, 331 "Bet^" iron, 167 hardness of, 180 phase in brass, 305 Binary alloy, cooling curves, 87 344 SUBJECT INDEX Binary alloy system, 72 simple, constitutional dia- gram of, 91 Binary alloy systems, brittleness of middle phases in, 145 conductivity curves of, 111 Binary aUoys, modes of solidiflca- tion, 73 Bismuth, 72 and gold, 316 -gold, Arnold's experi- ments, 256 in copper, 317 intercrystaUine brittle- ness of, near melting- point, 260 -lead -tin, liquidus sur- face of, 126 Bituminous varnish, 329 Black heart malleable castings, 190 Blow holes in ingots, 295 Bolts, cold rolled, 306 hard drawn, 306 Boron sub-oxide as deoxidiser, 152 Boundaries, crystal, movement at, 258, 259 intercrystalline, and strength, 274 of crystals, cementite films in, 278 under strain, 256 Brass, 74 brittle when hot, 305 cold working, 145 corrosion of, 331, 332 hot and cold forging, 305 /3 phase in, 305 hot working, 145 lead, cold worktag of, 324 lead in, for machining, 312 loss of zinc from, 320 "machinery," 311 over-annealing of, 225 Qver -heating of, 277 Brass — continued. oxide in, 319 polishing of, 24 protection against corrosion, 331 rolled, stress-strain diagram of, 211 " season cracking " of, 308 Brinell's bail test, 218 Brittle impurities, 322 membranes, 316 Brittleness, abnornal intercrystal- line, 256 and machining, 311, 312 near punched holes, 310 near sheared edges, 310 of burnt steel, 283 of compounds, 261 of metallic compounds 102 of middle phases in binary systems, 145 Bronze, over -heating of, 277 oxide in, 319 phosphor, 152 polishing of, 24 " Bronzes," aluminium, 334 Bronzes, cold worked, 148 Burgess's thermal curves of iron, 168 Burning, 325 " Burning " metals, 301 Burnisliing, 18, 26 " Burnt " steel, 283 Burnt steel shaft, 325 Cadmium, 72, 112 annealing, 269 crystalline cavities in, 66 Calcium carbonate coating on brass, 332 Calibration of thermo-couples, 81 Carbide of iron, 162 SUBJECT INDEX 345 Carbide of iron — continued. amorphous solution of, in hardened steel, 182 retained in " forced solution," 180 in solid solution, effect of, 177 movement of, 178 state of, in Troostite, 183 Carbide, solubility in y, /3, and a - iron, 167 Carbon, combined, 188 in iron, 74 in iron, limit of, 162 -iron alloys, 160 liquidua of, 162 solidus of, 163 retards transformation of fi-iron, 178 steel as duplex alloy, 263 Carbon and iron, alloys of, 131 Carbonate of lime, coating on brass, 332 Carburising reagents, 191 Carpenter and Edwards on alu- minium alloys, 154 Stead on crystal- lisation of elec- trolytic iron, 279 on zinc-copper alloys, 141 Case-hardening of steel, 190 Casting, 285 circular, isothermals of, 290 defects due to, 322 die, 333 temperature, 70,288,323 Castings, aluminium, 333 and isothermals, 293 best rate of cooling, 288 chill, 288 contraction stresses in, 290 " drawing " of, 323 large steel, 287 malleable, 189 shape of and structure, 290 Castings — continued. eteel, heat refining of 294 steel, sulphides in, 294 unsoundness in, 323 with re-entrant angle, 292 isothermals in, 292 Cast-iron and constitutional dia- gram, 186 and steel distinction be- tween, 186 chilled, 288 oxidising, annealing of, 189 phosphide eutectic, 316 Cast metal, imperfections of, 287 Cavities, gas, 294 shrinkage, 294 Cement, amorphous, 69 and annealing, 267 as protective coating, 330 Cement, interorystalline, amor- phous, 257 Cementation of steel, 191 Cementite, 162 balled-up, in steel, 325 decomposition of, 163 discovery of, 119 films in crystal boundaries, 278 free, 176 in high carbon alloys, 186 in micro-structure of steel, 176 solubility of in a -iron, 164 structurally free, 278 Ceylon, ancient u-on from, 329 metal from, 255 Chamberlain, 129 Changes of form, internal, 302 . Charpy's test, 323 testing machine, 234 Chemical activity of molten metal, 289 composition, errors in, 320, 322 346 SUBJECT INDEX Chemical — continued. metallurgy, definition, 1 study of alloys, 118 theory of corrosion, 328 Chill castings, 288 Chilled cast-iron, 288 Choice of specimens for micro- scope, 18 Chromatic aberration, 46 Chromium oxide, 29 Circular casting, isothermals of, 290 Cleavage and twinning, 276 fracture in steel, 263, 264 in steel sheets, 275 in wrought iron, 275 "Clogging" of tools, 311 Coalescence of lead-tin eutectic, 134 Coating, cement as protective, 330 protective against corro- sion, 329 Coatings, metallic, 329 Cobalt, 112 -copper alloys, conductivity curve of, 112 Cohen on " strain disease," 309 Cohesion, inter-crystalline, 256, 257 Coker's optical method, 237 Cold-forging brass, 305 junction, 81 rolled bolts, 300 "Cold shuts." 323 Cold work, 265 defined, 300 limits of, 307 Cold-worked bronzes, 148 Cold-working, defects arising in, 324 during erection. 324 excessive, 307, 308, 324 lead-brass, 324 of brass, 145 Colloidal iron carbide, 183 Combined carbon and structure of iron, 188 Commencement of corrosion, 329 Comparative tests of zinc — ^alu- minum alloys, 238 Comparison of testing methods, 237, 238 Complexity of iron-carbon system, 160 ternary alloys, 123 Composition, chemical errors of, 320 Compound Ala Zns, 137 Al, Mn, 157 antimony-tin, 321 composition of, by thermal analysis, 104 Sn-Sb, importance of, 135 Compounds and electro-chemical potential, 117 brittleness of, 261 conductivities of, 114 determined residue analysis, 118 fictitious, 118 inter-metallic, 73 of aluminium and dis- integration, 333 Compression, elastic limit in, 247, 248 fluid, of steel, 295, 296 test, 211 Concentration potential curve, 117 Concrete, reinforced, 330 Condenser tubes, corrosion of, 331 Conductivities of compounds, 114 Conductivity curve of — cobalt-copper alloys, 112 copper-antimony alloys, 114 gold-copper alloys, 112 Conductivity curves of alloys forming com- pounds, 114 electric, 110 of euteotiferous alloys. 111 SUBJECT INDEX 347 Conductivity — continued. of solid Bolutions, 111 thermal, 117 Cone indentation test ,Ludwik's,222 Congenital defects, 321 Constantan, 111, 115 -copper, 80 Constanta, physical, of metals, 3 Constituent, 119 Constituents, impurities as, 314 mutual support of, under strain, 262 which float, 320 sink, 320 Constitutional diagram, 78, 94 aluminium-zinc, 136 and cast iron, 186 and cooling-curves, 91 and electrical properties, 110 and iron, 167 and micro-structure, 105 and physical properties, 107 and practical steel hardening, 184 and specific volume, 108 and temperature co-efiicient, 115 annealing in determination of, 105 diagrams, horizontal lines on, 102 iron-graphite system, 162 model of ternary alloys, 125 of aluminium-copper aUoys, 154 of copper-tin alloys, 147 of iron-carbon alloys, 161 of lead -tin aUoys, 133 of magnesium-tin, 101 of partly eutectiferous alloys, 99 of simple binary system, 91 of solid solution system, 98 of zinc-copper alloys, 141 practical bearing of, 146 Contraction, local, in tensile tests, 207 Contraction, stresses in castings, 289 Cooling, best rate of, castings, 288 curve of binary alloys, 87 of pure iron, 168 of pure metal, 85 curves, 78 of aluminium-zino alloys, 140 for thermal curves, 79 rate of, and structure of castings, 287 Copper, 70, 74, 81, 111, 112, 114 alloys, etching of, 31 manganese in, 163 alloys of, 140 -aluminium alloys, 154 micro-structure of, 156 heavy, resistance to — corrosion, 334 oxidation, 334 (heavy), soldering of, 335 aluminium in, 153 -aluminium-manganese, 167 liquidus of, 126 -ammonium acetate, 46 annealing, 269, 270 -antimony, 114 antimony in, 317 arsenic in, 316 bismuth in, 317 deoxidation of, 162 electro -chemical deposition in etching, 35 chemically deposited, 304 for embedding sections, 246 etching of, 31 " gassing " of, 272 -manganese alloys, 153 -nickel, 74, 97, 111 over-annealing of, 275 oxide in, 152, 316 phosphide in bearing metals, 153 phosphorus alloys, 152 348 SUBJECT INDEX Copper — continued. polishing of, 24 -potassium chloride, 32 salt solutions for etching, 35 -silver, 100 solubility of, in solid alu- minium, 156 -tin, 103 -tin alloys, 146 anomalous behaviour of f> phase, 149 constitutional diagram of 147 oxide in, 152^ strength of, 148 twinning in, 251 -zinc alloys, 141 constitutional diagram, 141 heat treatment of, 144 oxide in, 152 -tin alloys, 150 " Cores " in crystals and banded structure, 304 Cores of solid solutions, 96 Correlation between tests and ser- vice results, 8 Corrodibility, 322 Corrosion, 327 and basic zinc chloride, 331 and heavy aluminium-copper alloys, 334 and oil paint, 329 and stray electric currents, 330 chemical theory, 328 commencement of, 329 electrolytic theory of, 328 general cause of, 328 " inertia " of pure iron, 328 in non-ferrous metals, 330 of aluminium, 332 of brass, 331 protection against, 331 with 2 per cent, lead, 331 with 1 per cent, tin, 331 of condenser tubes, 331 Corrosion — continued. of light alloys, 333 protective coating against, 329 resistance to, of manganese- aluminium - copper alloys, 157 Couple, thermo-electric, 80 Cracks, from quenching, 326 in hardening steel, 185 quenching, and enclosures, 318 Critical deformation, 267 illumination, 64 -points of 0-2 per cent. carbon steel, 171 -points of iron, 169 range and heat treatment of steel, 277 temperatures, 322 " Cropping " ingots, 298 Cross-sections of slip-bands, 245 Crushing action of cutting tool, 310 Crystal boundaries and strength, 274 cementite films in, 278 in burnt steel, 283 movement at, 258, 259 under strain, 256 viscous flow at, 259 cores and banded structure 304 facets, 66 groups, 275 growth and strain, 273 size and slip, 274 Crystalline aggregate, formation of, 62 fracture, 254 habit, 71 nature of duplex alloys, 260 SUBJECT INDEX 349 Crystalline — continued. structure, 26, 63 of eutectics, 134 permanence of in iron and steel, 256 system, 71 Crystallisation and flow of heat, 68 dendritic, 67 " Crystallisation " of metal under vibration, 264 Crystallisation (supposed) of metals, 266 Crystals, 33 adhesion of, 68 and flow of heat, 290 at re-entrant angle, 292 elongated, 241 " fringe," 293 growth of in annealing, 267 interlocking of, 71 liquid, 67 of y iron, formation of, 280 growth of, 281 plastic, 67 plasticity of, 247 radial, in tin-plate, 293 Crystal size and physical pro- perties, 274 Currents, electric, and corrosion, 330 Curvature of field, 49 Curve, concentration - potential, 117 Curved surfaces produced in grind- ing, 22 Curves, heating and cooling, of pure iron, 168 " recalesoence," 175 Cushman and Walker on corrosion, 328 Cutting operations, 24, 309 specimens for micro -exam- ination, 20 tools, action of, 309 tool, straining effects of, 310 Dalbt's optical stress -strain appa- ratus, 202, 208 Dark-etching constituents, 75 " Dead molted " steel, 295 Decarburisation of steel, 320 surface, of steel, 272 Decomposition of cementite, 163 iron solid solu- tion, 164 and allo- tropy of iron, 166 Defects, acquired, 321 arising from — casting, 322 moulds, 323 over-heating, 322 treatment, 322 arising in — ■ annealing or re-heat- ing, 326 cold working, 324 hot -working, 323 quenching, 326 tempering, 326 thermal treatment, 325 " congenital," 321 in metals and alloys, 313 tabulation of, 322 Deflection method, 82 Deformation, critical, 267 of duplex alloys, 260 Degens on lead -tin alloys, 132 Degree of freedom, 120 Delhi column and corrosion, 329 Dendrites in solid solutions, 74 Dendritic crystallisation, 67 Density of strain-hardened metal, 250 Deoxidation of copper, 152 Deposition of copper, in etching, 35 Depth of focus, 48, 49 Derived differential curve, 85 Determination of fundamental properties, 196 350 SUBJECT INDEX Determination of soUdus in iron- carbon system by quenching, 163 Dezinciflcation, 331 Diagram, constitutional, 78 equilibrium, 78 triangular, for ternary alloys, 124 Diamond, 166 Die casting, 333 Differential curve, 84 Diffraction, 50 Diffusion in solid solutions, 96 Dimensions, effect of, in testing, 194 of slip-bands, 245 Disintegration, spontaneous, of aluminium alloys, 158 Disturbance of structure by polish- ing, 18 " Drawing " of castings, 323 Drawing hoUow, 268, 307 Draw-plate, 307 Dryiag specimens, 36 Ductility test by bending, 217 Duplex alloy, carbon steel as, 263 ultimate strength of, 262 yield-point ia, 262 alloys — annealuig of, 269 baUiag-up in, 276 brittle constituent in, 261, 262 crystalline nature of, 260 deformation of, 260 effect of treatment on, 262, 263 over-annealing and over -heat- ing, 276 Duralumin, 154, 158 Duration of annealing, 271 Dynamic tests, 223 need for, 197 Edges, sheared, brittleness near, 310 Effect of dimensions in testing, 194 Egyptian metal, ancient, 255 Elastic limit, 205 in compression after tensile strain, 247 primitive, 205 " true," 255 limits and strain-hardening, 247, 248 modulus, 203 range after strain-harden- ing, 247 stretching of metal, 203 Electric conductivity, temperature co-efflcient of, 110, 115 currents and corrosion, 330 furnace and steel, 317 lamp for microscope illu- mination, 56 Electrical conductivity, 110 properties of alloys, 110 measurement of, 116 Electro -chemical potential and compounds, 117 study of alloys, 117 -chemically deposited cop- per, 304 -deposited copper for em- bedding sections, 245 Electrolytic iron, recrystallisation of, 279 protection, 332 theory of corrosion, 328 Electron theory, 110 Elongated crystals, 241 structure, 303 Embedding sections in electro- copper, 245 Emery papers, 23 testing machines, 199 Enclosures and mechanical pro perties of steel, 318 SUBJECT INDEX 351 Enclosures — continued. and quenching cracks, 318 mechanical, impurities as, 314 slag, as nuclei, 281 slag, in steel, 317 English type of tensile machine, 109 Equilibrium, attainment of, 122 diagram, 78, 93, 122 pressure factor, 120 (constitutional) dia- gram of simple binary system, 91 meta-stable, 94 stable, 94 ultimate physical, 274 Errors of chemical composition, 320 of design, 327 due to insufftcient observa- tion, 113 of treatment, 321 Etched surfaces, illumination of, 65 Etching, 30 electro-chemical, 304 macroscopic, 35 nature of, 32 by deposition of copper, 35 depths of, 31 figures, 66 pattern, 33 of pure metals, 33 reagents, 31 action of, 64 Eutectic, 119 phosphide, in cast iron, 316 ^Entectics, 133 crystalline structure of, 134 nature of, 261 plastic straining of, 261 predominant partner in, 134 spherulites of, 134 Eutectic alloy, 76 freezing of, 89 Eutectic aUoys, 86 of aluminium and copper, 157 copper oxide, 316 in iron-carbon alloys, 163 point, determined by thermal analysis, 104 Eutectiferous alloys, curve of, con- ductivity of. 111 freezing of, 88, 89 Eutectoid of iron-carbon alloys, 165 point, influence of man- ganese on, 175 pearlite, 173 steel, recalescence of , 175 Evolutions of heat in iron, 167 Ewen and Bosenhain on fracture of hot metals, 260 Ewing and Humphrey on fatigue, 254 Ewing and Eosenhain on anneal- ing lead, 269 Ewing's extensometer, 203 Excessive cold-working, 324 " Excessive " purity, difficulties of, 313 Extension, plastic, 207 Extensometer, Ewing's, 203 Marten's, 204 Turner's, 108 Extensometers, 203 Eye-piece, 45 magnification, 51 Factors of safety, 5 Failures in metals and alloys, 6 investigation of, 6 Fatigue and crystal size, 275 and strain-hardness, 306 defined, 224 fracture under, 253 test, Osborne- Keynolds", 228 352 SUBJECT INDEX Fatigue — continued. Stanton's, 228, 229 Wohler's, 225 Favourable examples, study of, 8 Feeding ingots, 295 Ferric chloride, 31 as etching reagent, 35 Ferrite, 173 before and after straining, 243 Ferro-ooncrete, 330 manganese, 317 Fibrous fracture, 252 Field, flatness of, 48 Fields, phase, 121 Filter, Wratten and Wainwright's, 46 Fine-grinding, 23 Finishing temperature and micro- structure, 300 Fire-clay, 81 Fissures, partial welding of, 324 Fitting by cold -working, 324 Flatness of field, 48 Floating constituents, 320 Fluid compression of steel, 296 Focus, depth of, 48, 49 " Forced solution," of iron-carbide 180 Foreign matter from mould, 323 Forging, 300 insufficient work in, 323 Formation of slip bands, 243 Fracture, cleavage, in steel, 263, 264 crystalline, 254 " fibrous," 252 intercrystaUine, abnormal, 252 of metals near melt- ing point, 260 mechanism of, 252 of iron at high tempera- tures, 259 of steel, shock, 256 types of, 263 Fractured steel shaft, 323 Fractures, alternating stress, 264 Fractures — continued. section of, 252 tensile, 252, 263 microscopic study of, 16 under fatigue, 253 shock, 252 Free cementite, 278 Freedom, degree of , 120 Freezing of eutectic aUoy, 86, 89 -points of metals, 81 Friend on corrosion, 329 Fringe crystals, 293 Fundamental properties, 196 Furnace, electric, and steel, 317 Fusion spots, 106 Galvanising, 330 Gralvanometer, 82 Gamma iron, 163 and constitutional diagram, 167 crystals, formation of, 280 growth of, 281 decomposition of, 164 in alloy steels, 284 solid solution, 164 transformation affected by carbide, 177 twinning in, 252 phase of tin-copper alloys, 149 Garland on ancient Egyptian metal, 255 Gas cavities, 294 Gasses and molten metals, 289 influence on thermal curves, 80 Gas lamp for microscope illumina- tion, 55 liberation on freezing, 294 " Gassing " of copper, 272 G«rman silver, 306 Goerens on annealing iron and steel, 271 Gold, 112 and bismuth, 316 annealing, 269, 270 SUBJECT INDEX 353 Gold — continued. -bismuth, Arnold's experi- ments, 266 -copper alloys, conductivity- curve of, 112 purest, brittleness near melt- ing-point, 260 -silver, 74 Grains, polygonal, 67 Graphical representation of ter- nary alloys, 124 Graphite, formation by decom- position of oementite, 163 in cast iron, 288 separation from iron, 320 Graphitic iron, 187 Grey iron as "steel plus graphite," 188 structure of, 187 Grinding wheels, 21 Groups of crystals, 275 Growth of crystals in annealing, 267 S iron crystals, 281 Gulliver, 123 Gutowsky on iron-carbon soUdue, 163 Hadfield on ancient metal, 265 sound ingots, 296 Hammer-dressing, 324 Hammering, deformation by, 302 effect of, 301 Hannover's porous metal, 134 Hard-drawn bolts, 306 Hardened steel, amorphous iron- carbide solution in, 182 Hardening by strain and elastic limits, 247, 248 of Duralumin, 158 metal by straining, 247 steel, " amorphous" theory, 180 P.M. Hardening — continued. of steel — continued. incipient oc crystals in, 181 and allotropic iron, 167 and quenching strains, 181 cracking in, 185 twinning in, 181 and constitu- tional dia- gram, 184 rules for, 184 " Hardenite," defined, 179 Hardness, defined, 218 number, Bendicks' for- mula, 219 of iron, 180 of Martensite, cause of, 180 of steel, allotropic theory of, 180 test, 218 tester, Martens', 220 testing by the sclero- scope, 222 Harmet process, 296 Haughton and Bosenhain on phos- phorus in steel, 304 " Header," 295 Heat evolutions in iron, 167 flow of, and crystallisation, 68 flow of, and crystals, 290 -refining of steel, 282 castings, 294 specific, 117 tinting, 34, 316 -treatment of — alloy steels, 284 Duralumin, 158 Muntz metal, 144 steel, 277 Heating curve of pure iron, 168 curves, 78 for thermal curves, 78, 79 A A 354 SUBJECT INDEX Hensler alloys, 109 Heussler magnetic alloys, 128 Heyn on corrosion, 329 High-carbon alloys, cementite in, 186 temperatures — behaviour of metals at, 257 fracture of iron at, 259 straining steel at, 259 Holes, punched, brittleness near, 310 Hollow drawing, 268, 307 Hooke's law, 203 Horizontal hnes in constitutional diagrams, 102 Hot-etching of steel, 17 -forging brass, 305 -junction, 81 -rolled metal, longitudinal structure in, 303 -work, 265 -working, defects arising in, 323 defined, 299 of brass, 145 rapid, 301 Hoyt on tin-copper alloys, 149 Hoyt's diagram of zinc-tin-copper alloys, 150 Hum&ey and Ewing on fatigue, 264 Eosenhain on Arj, 170 relative hard- ness of $ and 8 iron, 180 straining iron at high tem- peratures, 258, 259 Humfrey on hardening of steel, 181 Humfrey's theory of strain-raised elastic hmit, 248 Hydrochloric acid, 31 Hydrofluoric acid, 31 Hydrogen, 80 sulphide use in testing, 35 Hypo-eutectoid steel, 183 Hysteresis, magnetic, 109 Ice-box for thermo-couples, 81 lUuminants for microscopes, 53 Illumination— axial or central, 40 critical, 54 for photomicrography, 66 obhque, 38 of metal surfaces, 38 vertical or normal, 39 Illuminator, adjustable, of Eosen- hain microscope, 43, 44 Image produced by objectives, 45 Images, how produced, 45 projected, 57 weak or " mUky," 42 Imitative tests, limitations of, 194 Impact, resistance to, nature of, 232 test, shearing, 235 single blow, 233 repeated blow, 236 torsion, 235 tests, 232 Imperfections of cast metal, 287 Impurities, 313 accidental, 319 classified, 314 as constituents, 314 as mechanical enclosures, 314 brittle, 322 in ingots, 297 in sohd solution, 314 metalloid, 314 specification by analysis, 314 Incipient a-crystals in quenched steel, 181 IncorrodibiUty, alleged, of alu- minium, 332 Ingot, blow holes in, 295, 323 pipe in, 295, 323 relation to forging, 297 SUBJECT INDEX 355 Ingots, 294 cropping, 298 feeding, 295 impurities in, 297 liquid core rolling, 296 non-ferrous, 296 segregation in, 297, 323 sound, Hadfleld's process, 296 soundness of, 294 Insoluble additions to alloys, 319 Instrument making, metals used in, 2 Insufficient annealing, 268 observations, errors due to, 113 work in forging, 323 IntercrystaUine brittleness, 256 of pure metals near melting- point, 260 Interorystalline cement, 257 cohesion, 256, 257 fracture near melt- ing-point, 260 interstices, 68 Interference (of Ught), 50 Interlocking of crystals, 71 Intermetallio compounds, 100, 101 Internal changes of form, 302 strains in hardening of steel, 181 tension, 324 International testing association, 10, 161, 240, 264 Interpenetration (forced) of polish- ing powder and metal, 27 of crystals, 70 Interpretation of equilibrium dia- grams, 122 images, 52 Inverse rate curve, 83 Iron, 70, 72 aUotropio, and hardening of steel, 167 allotropio modifications of, 166 Iron — continued. aUotropy in, and decomposi- tion of y iron solid solution, 166 alloys, etching of, 31 u, 167 and steel, annealing, 270, 271 u, absence of twinning, 252 and carbon, alloys of, 131 and steel specimens kept " passive," 37 0, 167 and y, relative hardness of, 180 p hardness of, 180 carbide, 162 colloidal, 183 forced solution of, 180 in solution, effect of, 177 movement of during transformation of y soUd solution, 178 solution, amorphous, in hardened steel, 182 state of in Troostite, 183 -carbon aUoys, 160 constitutional diagram of, 161 eutectoid in, 165 Kquidus of, 162 solidus of, 163 carbon in, 75 limit of, 162 carbon, meta-stable diagram, 161 carbon system, complexity of, 160 cast, oxidising, annealing of, 189 phosphide euteotio in, 316 critical points of, 169 electrolyi;ic, recrystallisation of, 279 etching of, 31 fracture of at high tempera- tures, 259 A A 2 356 SUBJECT INDEX Iron — continued. from Ceylon, 329 graphite, constitutional dia- gram, 162 graphitic, 187 grey, as " steel plus graphite," 188 grey micro structure of, 187 heat, evolutions in, 167 in aluminium, 316 -manganese, 74 micro-structure of, and critical points, 170 molten, graphite from, 320 nuclei in transformations of, 169 phosphide, 315 pure, 168 and corrosion, 328, 329 effects of quenching, 171 heating- and cooling- curves of, 168 semi - plasticity in over- strained, 248 slip-bands in, 243 Swedish, 61 strained, 241 structure of, and combined carbon, 188 transformer sheets, 109 white, micro - structure of, 187 wrought, cleavage in, 275 " yield-point " in, 207 IsothermaJs and castings, 293 and radial structure, 291 at re-entrant angle, 292 of circular casting, 290 Izod's testing machine, 233 KiSH, 320 KroU on persistence of structure of steel, 281 Laboratory, National Physical, 334 Lag in critical points of iron, 169 LamellsB, twin, 251 Laps, 324 Large areas, need of examining microscopically, 20 Larrard on torsion testing, 213 Law of similarity in testing, 194 Lead, 81, 112 and antimony alloys of, 131 and tin, alloys of, 131 anneaUng, 269 antimony alloys, 134 -brass, cold working of, 324 corrosion of, 331 in brass for machining, 312 polishing of, 24 pure, brittleness of, near melt- ing-point, 260 segregation of, 321 shp bands in, 243 solubihty of tin in, 132 -tin, 75 -tin alloys — constitutional diagram of, 133 Degen's view, 132 Mazotto's views, 132 ' micro-structures of, 133 ternaries derived from, 135 use as solders, 134 -tin-bismuth, liquichis surface of, 126 -tin-euteotic, coalescence of, 134 Le ChateUer, 29 Level, diSerences of, recognised, 39 Levelling devices, 69 Lever testing machines, 200 Liberation of gas, 294 Light alloys — atmospheric corrosion, 333 corrosion of, 333 manganese in, 333 SUBJECT INDEX 357 Light, monochromatic, 46, 47 Limitations of phase rule, 121 Limit, elastic, 205 primitive, 205 of proportionaUty, 205 soUd solubility, 73 solutions, 105 Liquid core ingots, rolled, 296 " Liquidus" 94 Liquidus — maximum in, 101 model of copper-aluminium- manganese, 126 of aluminium-copper alloys, 155 of iron-carbon aUoys, 162 of manganese-copper alloys, 153 of partly eutectiferous system, 100 of solid solutions, 98 surface of lead-tin-bismuth, 126 ternary alloys, 126 Longitudinal structure in hot- rolled metals, 303 Losses from volatihsation, 322 meeting, 320 Ludwik's cone indentation test, 222 " Machinery " brass, 311 Machines, automatic, steel for, 312 Machine, Charpy's testin'g, 234 Izod's testing, 233 Sankey's, 231 straining, for microscope, 242 Machines, tensile testing, 198 Machining and brittleness, 311, 312 properties of alloys, 311 Macroscopic etching with copper salt solutions, 35 Magnesia, 29 Magnesium -tin, 100 Magnetic alloys, Heussler, 128 hysteresis, 109 properties, 109 Magneton theory, 110 Magnification, useful, 51 Malleable castings, 189 Manganese-aluminium-copper, 157 liquidus model of, 126 as deoxidiser, 153 copper alloys, 153 influence on eutectoid point, 175 in hght alloys, 157, 333 -iron, 74 sulphide of, in steel, 317 Martens' extensometer, 204 hardness tester, 220 Mess-Dose, 201 sclerometer, 222 testing machines, 198 Martensite, 283 and twinning, 181 coarse, 326 defined, 178 nature of, 180 Martensitic alloy steel, 284 Materialpriifungsamt, 201 Materials, abuse of, 327 Matter, constitution of sohd, 3 Maximum in liquidus, 101 Mazotto on lead -tin, 132 Measurement of electrical proper- ties, 116 Mechanical enclosures, impurities as, 314 properties of — alloys, 101 steel and enclo- sures, 318 testing of metals, 193 simple princi- ple in, 195 treatment, 286 Mechanism of fracture, 252 Mess-Dose, Martens', 201 358 SUBJECT INDEX Metallic compoundB, brittleness of, 102 Metallography, history of, 12 Metalloid impurities, 314 Metallurgical microscopes, 41 Metals at high temperatures, 257 freezing-points of, 81 hexagonal, 72 mechanical testing of, 193 pure, etching of, 33 temperature coef&- cient of, 115 simple, 33 Meta-stability of ancient metal, 122 metaUic state, 328 stable conditions, 121 diagram for iron-car- bon alloys, 161 equilibrium, 94 Method of residue analysis, 118 Melting losses, 320 temperature, unduly low, 322 Membranes brittle, 316 MetaUic protective coatings, 329 state, meta-stabUity of, 328 Metallurgy, Physical, definition, 1 Microscope by Le ChateUer, 58 Greenhough binocular, 59 iUuminants for, 53 images, interpretation of, 52 metallurgical, 38 optical system of, 44 Rosenhain's metallur- gical, 43 rules for use of, 53 simple metallurgical, 41, 42 tube length, 45 Microscopes, metallurgical, 41 Micro -structure, affected by anneahng, 273 and annealing, 266 Micro -structure — continued. and critical points of iron, 171 and finishing temperature, 300 and shape of castings, 290 as evidence for constitutional diagram, 105 elongated, 303 of alloys, 74 of aluminium -copper alloys, 156 -zinc alloys, 139 of burnt steel, 283 of eutectoid steel, 175 of grey iron, 187 of over-heated steel, 282 of pearhte, 173 of 0'2 per cent, carbon steel, 172 of 0'6 per cent, carbon steel, 174 of 0'9 per cent, carbon steel, 175 of lead-tin alloys, 133 of white iron, 187 Micro -volts, 82 Mild steel, " yield-point " in, 207 Mixed crystals, 73 Mobile layers, temporary, 246, 247 Model, constitutional, of ternary system, 125 Modulus, elastic, 203 Molten alloys as solution, 73 iron, graphite from, 320 metal, chemical activity of, 289 overheating of, 289 metals and gases, 289 Monochromatic light, 47 Mould, foreign matter from, 323 Moulds, defects arising from, 323 venting of, 323 Movement of crystal boundaries, 258, 259 iron carbide, 178 Muir on recovery of iron and steel, 271 Muntz metal, 144 SUBJECT INDEX 359 National Physical Laboratory, 334 Need for dynamic tests, 197 Nickel-copper, 74, 97, 111 -steel, slip-bands in, 243 zinc-copper aUoys, brittle when hot, 305 Nitric acid, 31 Nitrogen, 80 Nomenclature Committee, Inst. Metals, 336 Non-ferrous alloys, segregation in, 320 ingots, 296 metals, corrosion in, 330 Normal and oblique illumination of etched surface, 65 Normal illumination, 39 Notched bar impact tests, 233, 234 test-bars, 236 Nuclei in pearhte formation, 281 in transformations of iron, 169 slag enclosures as, 281 Objective, aberrations of, 46 aperture of, 46 image produced by, 46 Objectives, aperture of, 51 for metallurgy, 42 (microscope), 46 zones of, 47 Obhque Ulumination, 38 light, slip-bands under, 245 Oil paint and corrosion, 329 -quenching steel, 186 Opaque reflector, influence on resolving power, 62, 53 Optical character of slip-bands, 244 system of microscope, 44 Ordinary " pohsh," 18 Orientation, crystalline, 33 of crystals, 63 Oriented lustre, 63 explanation of, 66 Osborne-Eeynold's direct alter- nating stress test, 228 Over-anneahng duplex aUoys, 276 in brass, 276 copper, 275 Overheated steel, 325 angular structure of, 282 Overheating, 325 duplex alloys, 276 molten metal, 289, 322 of brass, bronze, aluminium« copper, and steel, 277 steel, 281 Over-strain and semi-plasticity, 248 recovery from, 249 Overwork, 268 Oxidation in annealing, 272 resistance to, of heavy aluminium - copper alloys, 334 Oxide in brass and bronze, 319 copper, 152, 316 tin-copper alloys, 152 zinc-copper aUoys, 162 Oxidising anneahng of cast iron, 189 Paint and corrosion, 329 Pattern, etching, 33 Patterns, 323 Pearhte, 167, 173 angular, 282 balling-up of, 278 formation, nuclei in, 281 in tensile fracture, 263 laminated, 278 longitudinal hues of, 303 micro -structure of, 173 Permanence of aluminium aUoys, 332 metals, 327 Personal names in metallography of steel, 172 360 SUBJECT INDEX Phase, amorphous, defined, 249 in hardened steel, 181 definition of, 119 Phase-fields in copper-tin alloys, 147 in Hoyt's ternary diagram ,151 in zinc-copper alloys, 141, 142 of aluminium-copper alloys, 155 of aluminium-zino alloys, 138 Phase rule, 119 limitations of, 121 Phases, redundant, 121 Phosphide euteotic in cast iron, 316 of copper, 152 of iron, 315 Phosphor bronze, 152 Phosphoric banding in steel, 304 Phosphorus as deoxidiser, 152 banding in steel, 237, 315 -copper alloys, 152 in axles, 315 in ferrite, 173 in iron and steel,34,315 in rails, 315 in steel, for machining, 312 in tyres, 316 reagent (Eosenhain and Haughton), 35 springs, 315 Photomicrography, 58 illumination for, 57 Physical equihbrium, ultimate, 274 meaning of tension tests, 197 metallurgy, definition, 1 metallurgy, practical im- portance of, 3 properties and anneahng, 268 and crystal size, 274 and strength, 195 Physical — continued,. properties of alloys, 102 Picrate, sodium, etchiug with, 325 Picric acid, 31 as etching reagent, 173 Pipe in an ingot, 295, 323 Plastic crystals, 67 extension of metals, 207 Plasticity of crystals, 247 Platinum, 70, 80 PoHsh attack, 33 " PoKshed " surfaces in hquids, 27 Polishing, 18, 23 cloths, 28, 29 powders, action of, 26, 27 test of, 30 Polygonal grains, 67 Polyhedral structure, 62 of y-iron boM solution, 174 Porous metals, Hanover, 134 Potential, electro -chemical, 117 Potentiometer, 82 Practical bearing of constitutional diagram, 146 " Predominant partner " in eutec- tic, 134 Preservation of specimens, 37 Press-forging, 302 Pressure factor in equihbrium dia- grams, 120 Prevention of segregation, 298 " Primitive " elastic limit, 203 Principle, simple, in mechanical testing, 195 Prism reflector, 41 Process, metallurgy, definition of, 1 Projection apparatus, 58 eye-piece, 58 Projections of ternary models, 126 Prolonged annealing, 268, 272 Properties, machining, of aUoys, 311 Prop rtionaUty, limit of, 205 Protection against corrosion, 329, 331 SUBJECT INDEX 361 Protective coating, cement as, 330 Protectors for thermo-couples, 81 Punched holes, brittleness near, 310 Punching tests, 216 Pure metal, cooling-curve of, 85 structure of, 61 Pure metals, temperature co-effi- cient of, 115 under oblique light, 63 Purity, " excessive " difficulties of, 313 Quantity of heat evolved, arrest points, 103 Quenched steel, 177 Martensite in, 179 transition products in, 178 Quenching, 102 bath, 326 cracks, 326 and enclosures, 318 defects arising in, 326 for determination of iron- carbon aolidus, 163 pure iron, defects of, 171 tin-copper alloys, 149 unequal, and warping, 326 Eadial crystals in tin plate, 293 structure and isothermals, 291 Kails, phosphorus in, 315 Sange of alternating stress, safe, 226 Rapid list working, 301 Bate of cooling and structure in castings, 287 straining and resistance, 259 Bazors, steaming of, 271 Be-appearance of scratching on etching, 28 B6aumur malleable castings, 190 " Becalescence " curves, 175 Recalescence of eutectoid steel, 175 Becovery from over-strain, 249 Becrystallisation and annealing, 266 spontaneous, 308 Be-entrant angle in castings, 292 isothermals at, 292 Beflector, opaque, 41 plain glass, 40 Eefractory metals in alloys, 319 Be-heating, defects arising in, 325 of white iron, 189 Belief polishing, 30 Bepeated blow impact tests, 236 impact, Stanton's test, 236 Besearch on aluminium alloys, 334 Besidue analysis in compounds, 118 Besistance and rate of straining, 259 to abrasion of man- ganese - aluminium - copper alloys, 157 to corrosion of man- ganese - aluminium • copper alloys, 157 to impact, nature of, 232 Besolving power, 52 Bhodium, 80 " Riser," 295 Bokes, 324 Boiling, 300 deformation by, 302 liquid core ingots, 296 Bontgen rays, use of, 17 Rose, on annealing gold, 270 Bosenhain and Ewen on annealing lead, 269 on fracture of metals at high temperatures, 260 362 SUBJECT INDEX Bosenhain and Haughton on phos- phorus in steel, 304 Bosenhain and Humfrey on Arj, 170 on relative hardness of 3 and ■y-iron, 180 on straining iron at high temperatures, 258, 259 Bosenhain and Tucker on lead-tin alloys, 133 microscope, 43 Bosenhain' 8 apparatus for thermal curves, 79 levelling device, 59 quenching apparatus, 102 Botation effect, 63 Eouge, 29 Bules for hardening steel, 185 Rupture, work of, 234 Safe range of alternating stress, 226 Samter on hot-etching of steel, 174 Sankey's testing machine, 231 Sauveur on crystal growth, 267 Scale of structure, 69 rolled into surface, 324 Sclerometer, Martens', 222 Turner's, 222 Scleroscope, Shaw's, 222 Scratches, re-appearance on etch- ing, 28 Sea, action of, on aluminium, 332 Season cracking, 325 " Season cracking " of hrass, 308 and temperature, 309 Section of fracture, 252 Sections embedded in electro - copper, 245 of ternary models, 126 Segregation in aluminium alloys, 321 in ingots, 297, 323 in non-ferrous alloys, 320 prevention of, 298 Semi-plasticity and over-strain, 248 Service, tests results, correlation of, 196 Severely strained metal, structure of, 246 Shaft, steel, burnt, 325 fractured, 323 Shaw's scleroscope, 222 Shear test, alternating, 229 Sheared edges, brittleness near, 310 Shearing impact tests, 235 test, 215 Shepherd and Blugh on tin -copper alloys, 148 Shepherd's copper-zinc diagram, 141 Shock and crystal size, 275 fracture, 252 of steel, 256 Shrinkage cavities, 294 Silica, 81 Silicate of manganese in steel, 317 Silicon, 110 absorption by aluminium, 289 in aluminium, 316 in ferrite, 173 Silver, 70, 81 -copper, 99 -gold, 74 twinning in, 251 Similarity, law of, in testing, 194 Simple metallurgical microscope, 41, 42 principle in mechanical testing, 195 Single blow impact tests, 233 Sinking constituents, 320 Size of crystals and physical properties, 274 Skimming, 323 Slag enclosures in steel, 317 Slip and crystal size, 274 and twinning, 276 -bands at high temperatures, 259 SUBJECT INDEX 363 Slip-bands — continued. defined, 243 dimensions of, 245 formation of, 243 in cross-section, 245 in eutectic alloys, 261 in iron, 243 in lead, 243 in nickel steel, 243 in twinned metal, 251 optical character of, 244 under oblique light, 245 Sodium picrate, etching with, 325 Softening temperature, 266 Soldering heavy copper-aluminium alloys, 335 Solders, lead-tin alloys, 134 Solid matter, metals simplest form, 3 solubility, limit of, 73 solution, cooling-curve of, 97 freezing of, 95 y-iron, 164 impurities in, 314 saturated, 99 solutions, 73 conductivity of. 111 cores in, 96 dendritic cores in, 74 diffusion in, 96 , limit of, 105 liquidus of, 97, 98 solidus of, 97 Solidification, rate of, 287 " Solidus," 94 Solidus determined by quenching, 105 of iron-carbon alloys, 163 of partly euteotiferous sys- tem, 100 Solubility of carbide in a, and y-iron, 167 of cementite in y-iron, 164 Solubility — continued. of copper in solid aluminium, 156 of tin in lead, 132 Sorbite, defined, 184 Sound ingots, Hadfield's process, 296 Soundness of ingots, 294 Special steels, 192 Specification of impurities by analysis, 314 Specific heat, 117 tenacity, 13 volume of alloys, 108 amalgams, 108 Specimens, drying, 36 for microscopic exami- nation, methods of cutting, 20, 21 preservation of, 37 washing of, 36 Spherulites of eutectic, 134 Spontaneous annealing and re- crystaUisation, 308 disintegration in alu- minium alloys, 158, 333 Springs, phosphorus in, 315 Stable equilibrium, 94 Stage of microscope, movable, 42 Stanton and Bairstow on alternat- ing stress fractures, 264 Stanton's direct alternating stress test, 228, 229 repeated impact test, 236 State, metallic meta-stabUity of, 328 Statical tests, defined, 213 Stead and Carpenter on crystal- lisation of electrolytic iron, 279 Stead on cleavage in steel sheets, 275 heat-tinting, 304 Steel, 75, 80, 121 alloy, Austenitic, 284 aluminium in, 295 364 SUBJECT INDEX Steel — contmued. and cast iron, distinction between, 186 and iron, annealing of, 271 annealing above Acs, 279 annealing below An, 277 arrest-points of, names of, 172 baUed-up cementite in, 325 behaviour under tensile test, 207 " burnt," 283 carbon, as duplex alloy, 203 case-hardening of, 190 castings, heat refining of, 294 sulphides in, 294 cementation of, 191 cementite and micro-structure of, 176 cleavage fracture in, 263, 264 containing 0-2 per cent, car- bon, 171 0*2 per cent, car- bon, micro-struc- ture of, 172 0-6 per cent, car- bon, 174 0'9 per cent, car- bon.micro -struc- ture of, 175 critical range and heat treat- ment, 277 " dead melted," 295 deoarburisation of, 320 eutectoid, micro -structure of, 175 recalescence of, 175 fluid compression of, 296 for. automatic machines, 312 fractures, ts^pes of, 256, 263 hardened, amorphous iron- carbide solution in, 182 hardening of and twinning, 181 practice and the constitutional diagram, 184 Steel — continued. hardening, rules for, 185 hardness of, and allotropic iron, 167 heat refining of, 282 heat-treatment of, 277 hot-etching of, 174 hypereuteetoid, quenched, 176 hypo-eutectoid, 183 in electric furnace, 317 large castings in, 287 mechanical properties of and enclosures, 318 mild, stress-strain diagram of, 210 yield-point in, 207 nickel, slip-bands in, 243 of eutectoid composition, 175 oil quenching of, 186 over-heated, 281, 325 angular structure of, 282 phosphide bands in, 237, 304, 312, 315 quenched, 173, 177 Martensite in, 179 transition products in, 178 semi - plasticity in over- strained, 248 sheets, cleavage in, 275 silicate of manganese in, 317 soft, temperature tenacity curve of, 179 strained at 1,000° C, 259 tempered, troostite in, 183 surface deoarburisation of, 272 Steels, alloy, 192 heat treatment of, 284 special, 192, 322 sulphide of manganese in, 317 Stopping down, 47, 49, 51 Strain and crystal growth, 273 effect on crystal boundaries, 256 SUBJECT INDEX 365 Strain — continued. effect on structure, 241 -hardened metal, 260, 300 -hardening and elastic limits, 247, 248 local, 247, 325 -hardness and fatigue, 306 not reliable, 309 uses of, 306 plastic and semi-plasticity, 248 Strained metal, amorphous layers in, 246 structure of, 246 Strained Swedish iron, 241 Straining at high temperatures, 259 duplex alloys, 260 eflects of tool out, 310 ferrite before and after, 243 of eutectics, 261 machine (Ewing and iKosenhain's), 242 steel at 1,000° C, 259 sudden, at high tempera- tures, 269 Strength and crystal boundaries, 274 " Strength " and physical proper- ties, 196 Strength of tin-copper alloys, 148 ultimate of duplex alloys, 262 Stresses, alternating, 224 Stress-strain diagram, 202, 207, 210,211 torsional, 214 indicator, Dalby's, 202 Stretching of metal, elastic, 203 Structure and annealing, 266 and rate of cooling a casting, 287 and shape of castings, 290 Structure — continued. angular of over-heated ste^l, 282 banded, and crystal cores, 304 efiect of strain on, 241 longitudinal, in hot- rolled metal, 303 of severely strained metal, 246 radial, and isothermals, 291 scale of, 69 Sulphides in steel castings, 294 Sulphuric acid, 32 Sulphur-prints, 318 Surface decarburisation of steel, 272 " Surface energy," 273 Surface flow in metals, 26 liquidus, of ternary alloys, 126 " Surface tension," 273 Surfaces in constitutional model of ternary system, 125 Swedish iron, 61 strained, 241 System, crystallographic, 71 Tabulation of defects, 322 Talbot on sound ingots, 296 Temperature and season cracking, 309 annealing, 266, 270 casting, 288, 323 CO -efficient and constitutional diagram, 115 electrical conduc- tivity, 110, 115 finishing, and micro -structure, 300 measurement, 80 softening, 266 tenacity curve for soft steel, 179 Temperatures, critical, 322 366 SUBJECT INDEX Temper-carbon, 189 Tempering, defects arising in, 326 of steel, 183 unequal, 326 Temporary mobile layers, 246, 247 condition of metal, 25 Tensile machine, English type, 199 testing, local contraction in, 207 machine, Amsler's, 201 pressure gauge type, 200 machines, 198 tests, data derived from, 203 Tension, fracture under, 252 internal, 324 surface, 273 tests, 197 comparison with tor- sion, 215 physical meaning of, 197 Ternary alloys, 123 as modifications of binary, 127 derived from lead-tin, 135 graphical representation of, 124 of tin, zinc and copper, 150 simplification of, 127 Ternary investigations, difficulty of, 128 models, projections of, 126 sections of, 126 Test, Arnold's, 231, 239 -bars, notched, 236 Charpy's, 233 compression, 211 Testing by bending, 217 by punching, 216 hardness, 218 in shear, 215 Testing — continued. in torsion, 213 machine, Emery's, 198, 199 Izod's, 233 lever, 200 Martens', 198 mechanical, simple prin- ciple in, 195 methods, compared, 237, 238 of materials, history of, 13 of metals, mechanical, 193 Tests, alternating shear and tor- sion, 229, 230 comparative on zinc-alumi- nium aUoys, 238 direct alternating stress, Stan- ton's, 228, 229 dynamic, 223 imitative, conditions for, 194 impact, 232 repeated blow impact, 236 results in service, correlation of, 196 shearing impact, 236 single blow impact, 233 torsion impact, 235 Wohler's, 226, 254 Thallium, polishing of, 24 Theory, allotropic, of hardness of steel, 180 amorphous, 180, 246, 266 chemical, or corrosion, 328 of amorphous cement, 257 phase, 250 of corrosion, electrolytic, 328 Thermal analysis, 104 of aluminium - zinc alloys, 139 arrest at freezing point, 85 conductivity, 117 SUBJECT INDEX 367 TheTvaal— continued. curve and constitutional diagram, 92 curves of 0'6 per cent. carbon steel, 174 curves, Eosenhain's ap- paratus, 79 treatment, 265 defects arising in, 325 Thermo-couples, calibration of, 81 platinum, 80 protection of, 81 electric couple, 80 properties, 117 Thin sections of metals, 17 Time -temperature curve, 83 Tin, 81 -aluminium - copper alloys, 157 and lead alloys of, 131 annealing, 269 antimony, 321 brass, corrosion of, 321 copper, 103, 146 alloys, anomalous be- haviour of phase, 149 constitutional dia- gram of, 147 oxide in, 152 strength of, 148 twinning in, 148 lead, 75 alloys, constitutional dia- gram of, 133 micro-structure of, 133 -bismuth, liquidus surface of, 126 -magnesium, 101 plate, radial crystals in, 293 pure, brittleness of, near melt- ing-point, 260 solubility of, in lead, 132 -zinc-copper alloys, 150 Tool cut, straining effects of, 310 Tools, clogging of, 311 cutting, action of, 309 Torsion impact test, 235 test, alternating, 230 comparison with ten- sion, 215 tests, 213 Torsional stress-strain diagrams, 214 Transformations in solid alloys, 102 tin-copper alloys, 103, 148 Transformer sheet iron, 109 Transition products in quenched steel, 178 Treatment, defects arising from, 322 effect of, on duplex aUoys, 262, 263 errors of, 321 mechanical, 285 Triangular diagram for ternary alloys, 124 Troostite, 283 defined, 183 Tube length of microscope, 45 Tubes, condenser, corrosion of, 331 Tungsten, 330 Turner's sclerometer, 222 Twinning, 251 and cleavage, 276 and slip-bands, 251 in copper and brass, 275 in copper-tin alloys, 148 in y-iron, 252 in hardened steel, 181 Type-metal, 135 Types of fracture of steel, 263 Tyres, phosphorus in, 315 Ultimate physical equilibrium, 274 Under-cooling of metals, 249 Unequal cold -working, 324 tempering, 326 Unsoundness in castings, 323 368 SUBJECT INDEX Vanadium, 71 Vamisli, bituminous, 329 Varnishing, specimens, 37 Venting of moulds, 323 Vertical illumination, 39, 42 Vibration and supposed " crystal- lisation," 254 Viscous flow at crystal boundaries, 259 Volatilisation, losses from, 322 Volume, specific, of alloys, 108 Walker and Cushman on corro- sion, 328 Warping due to unequal quench- ing, 326 of " aluminium " cast- ings, 334 Washing specimens, 36 Weakness of re-entrant angles, 292 Wear, 327 Welding, partial, of fissures, 324 White bearing metals, 135 iron, re-heating of, 189 structure of, 187 Whitworth " fluid compression," 296 Wohler on fatigue of metals, 224 test, 225, 254 interpretation of, 226 Wohler's alternating stress test, 224 Work, insufficient, in forging, 323 of rupture, 234 Wrought iron, cleavage in, 275 Yield-point in duplex alloys, 262 in iron and steel, 207 Young's modulus, 203 Zeigler on slag enclosures as nuclei, 281 Zinc, 72, 74, 81, 87, 112 -aluminium alloys, 136 comparative tests on, 238 cooUng-curves of, 140 phase -fields of, 138 thermal analysis, 139 and aluminium, machining of, 312 annealing, 269 basic chloride of, 331 -copper alloys, 141 fi phase in, 144, 145 constitutional diagram, 141 y phase in, 145 heat treatment of, 144 oxide in, 152 loss of from brass, 320 -tin-copper alloys, 150 Zones of lenses, 47 BEADBURY, AONKW, & CO. 1 Arnold, E. 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