PERGAMON MATERIALS SERIES VOLUME 2
Non-equilibrium Processing of Materials
PERGAMON
MATERIALS
SERIES
Series Editor: Robert W. Cahn F R S Department of Materials Science and Metallurgy, University of Cambridge, UK
Vol. 1 CALPHAD (Calculation of Phase Diagrams): A Comprehensive Guide
by N. Saunders and A. P. Miodownik Vol. 2 Non-equilibrium Processing of Materials edited by C. Suryanarayana A selection offorthcoming titles in this series:
Phase Transformations in Titanium- and Zirconium-based Alloys
by S. Banerjee and P. Mukhopadhyay Wettability at High Temperatures by N. Eustathopoulos, M. G. Nicholas and B. Drevet Ostwald Ripening by S. Marsh Nucleation by A. L. Greer and K. F. Kelton Underneath the Bragg Peaks: Structural Analysis of Complex Materials by T. Egami and S. J. L. Billinge The Coming of Materials Science by R. W. Cahn Thermally Activated Mechanisms in Crystal Plasticity by D. Caillard and J. L. Martin
PERGAMON MATERIALS SERIES
Non-equilibrium Processing of Materials edited by C. Suryanarayana The George S. Ansell Department of Metallurgical and Materials Engineering Colorado School of Mines Golden, Colorado, USA
1999
PERGAMON An Imprint of Elsevier Science Amsterdam- Lausanne- New Y o r k - Oxford- Shannon- Singapore- Tokyo
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Contents xv xvii xix
Series Preface Preface List of Contributors
CHAPTER 1 INTRODUCTION
1
R. W Cahn
CHAPTER 2 THERMODYNAMICS AND KINETICS OF METASTABLE PHASE FORMATION
5
K. N. Ishihara 2.1. Introduction 2.2. Thermodynamics of Metastable Phase Formation 2.2.1 Free Energy of Elements 2.2.2 Free Energy of Alloy Phases 2.2.3 Determination of Free Energy of Metastable Phases 2.2.4 Lattice Parameter of the Supersaturated Phase 2.3. Kinetics of Metastable Phase Formation 2.3.1 Nucleation of Metastable Phases 2.3.2 Nucleation of Alloy Phases 2.3.3 Crystal Growth Rate of Metastable Phases 2.4. Summary 2.5. List of Symbols References
5 5 5 7 9 11 12 13 15 18 19 19 20
CHAPTER 3 RAPID SOLIDIFICATION
23
H. Jones
3.1. Introduction 3.2. Methods of Rapid Solidification 3.2.1 Droplet Methods 3.2.2 Spinning Methods 3.2.3 Surface-Melting Methods 3.3. Constitution and Microstructure Formation by Rapid Solidification
V
23 24 24 26 27 28
3.3.1 Non-Equilibrium Constitution 3.3.2 Microstructure Formation 3.4. Properties, Performance and Applications of Rapidly Solidified Materials 3.4.1 Properties and Performance 3.4.2 Applications of Rapid Solidification References Selected Bibliography
CHAPTER 4 MECHANICAL ALLOYING
29 32 34 34 41 43 45
49
C. Suryanarayana 4.1. Introduction 4.2. Nomenclature 4.3. The Process of Mechanical Alloying 4.3.1 Raw Materials 4.3.2 Process Control Agents 4.3.3 Type of Mills 4.3.3.1 SPEX Shaker Mills 4.3.3.2 The Planetary Ball-Mill 4.3.3.3 Attritor Mills 4.3.3.4 Commercial Mills 4.3.3.5 New Designs 4.4. Mechanism of Alloying 4.5. Consolidation 4.6. Synthesis of Non-Equilibrium Phases 4.6.1 Solid Solubility Extensions 4.6.2 Synthesis of Intermetallics 4.6.3 Disordering of Ordered Intermetallics 4.6.4 Nanocrystalline Materials 4.6.5 Solid-state Amorphization 4.6.6 Displacement Reactions 4.7. Powder Contamination 4.8. Modeling 4.9. Industrial Applications 4.9.1 Nickel-Base Alloys 4.9.2 Iron-Base Alloys 4.9.3 Aluminum-Base Alloys 4.9.4 Magnesium-Base Alloys 4.10. Concluding Remarks References
vi
49
50 51 51 52 52
53 53 53 53 54 55 56 58 59 62 64 66 67 71 74 75 76 77 79 79 80 80 81
CHAPTER 5 LASER PROCESSING
89
K. F: Kobayashi 5.1. Principles of Lasers 5.2. Classifications of Laser Processing 5.3. Analysis of the Laser Melting and Quenching Process 5.3.1 Heat Transfer 5.3.1.1 Absorption Phenomena of Laser 5.3.1.2 Pulsed Laser 5.3.1.3 Continuous Wave (CW) Laser 5.3.2 Kinetic Conditions for Solidification 5.3.2.1 Kinetic Conditions for Amorphous Phase Formation 5.3.2.2 Solidification Modes in the Laser Processing of Materials 5.3.3 Microstructure Selection Maps (MSM) for Alloys 5.4. Laser-Quenching 5.4.1 Amorphous Phase Formation 5.4.2 Formation of Crystalline Phases 5.5. Laser Surface-Alloying and Cladding 5.5.1 Ferrous Alloys 5.5.2 Non-Ferrous Alloys 5.6. Laser-Annealing 5.7. Laser-Beam Joining 5.7.1 Welding of Structural Materials 5.7.1.1 Ferrous-Based Alloys 5.7.1.2 Non-Ferrous Alloys 5.7.2 Microjoining 5.8. Conclusions References
89 89 92 92 92 93 94 96 96 97 97 98 98 99 100 100 101 102 106 106 106 108 112 115 115
CHAPTER 6 THERMAL PLASMA PROCESSING
121
I? I!Ananthapadmanabhan and N. Venkatramani 6.1. Introduction 6.1.1 Advantages of Plasma Processing 6.2. Thermal Plasmas 6.2.1 Principles of Plasma Generation 6.2.1.1 DC Plasma Torches 6.2.1.2 AC Plasma Torches 6.2.1.3 F W Plasma Torches 6.2.2 Plasmagen Gases 6.2.3 Plasma-Particle Interaction 6.2.4 Plasma Processing Systems vii
121 121 123 123 124 126 126 127 128 129
6.3.
Processing of Materials 6.3.1 Plasma-Spraying 6.3.1.1 Structure of Sprayed Deposits 6.3.1.2 Plasma Spray-Deposited Materials 6.3.1.3 Low-Pressure Plasma-Spraying 6.3.1.4 Reactive Plasma-Spray Forming 6.3.1.5 Plasma Spheroidization 6.3.2 Plasma Reactors and Furnaces 6.3.2.1 Plasma Decomposition 6.3.2.2 Plasma Metallurgy 6.3.2.3 Processing of Ceramics 6.3.2.4 Treatment of Hazardous Wastes 6.3.3 Processing of Metastable Phases 6.3.3.1 Plasma Deposition of Diamonds 6.3.3.2 Thermal Plasma Synthesis of Ultrafine Alumina Powder 6.4. Summary and Conclusions Acknowledgments References
129 129 131 132 134 135 135 137 137 138 140 141 141 141 142 147 148 148
CHAPTER 7 153
SPRAY-FORMING Bing Li and E. J. Lavernia 7.1.
7.2. 7.3.
7.4.
7.5.
Introduction Principles Variations and Distinctions 7.3.1 Variations 7.3.1.1 Crucible: Induction Skull Melting/Spray-Forming 7.3.1.2 Atomizer: Circular vs. Linear Spray-Forming 7.3.1.3 Atomizer: Close-Coupled vs. Free-Fall Spray-Forming 7.3.1.4 Atomization Gas: Reactive Spray-Forming 7.3.1.5 Substrate: Near Net Shape Spray-Forming 7.3.1.6 Spray-Forming and Co-Injection 7.3.2 Nomenclature 7.3.2.1 Spray Atomization 7.3.2.2 Related Processing Techniques Applicability Non-Equilibrium Phenomena in Spray-Forming 7.5.1 Non-Equilibrium Nature 7.5.1.1 Rapid Solidification in Atomized Droplets 7.5.1.2 Transient Semi-Solid Layer in Deposition Stage 7.5.2 Non-Equilibrium Related Features in the Deposit 7.5.2.1 Metastable Phases 7.5.2.2 Extended Solid Solubility 7.5.2.3 Absence of Macrosegregation/Minimized Microsegregation viii
153 154 157 157 158 158 159 160 163 163 166 166 167 168 172 173 173 176 178 178 178 179
7.5.2.4 Refinement Effects of Non-Equilibrium Features on Mechanical/Physical Properties 7.5.3.1 Direct Effects 7.5.3.2 Indirect Effects 7.6. Concluding Remarks Acknowledgments References
180
7.5.3
181 183 185 189 189 189
CHAPTER 8 ION-MIXING
197
B. X. Liu
8.1. 8.2. 8.3.
Introduction Brief Description of Underlying Physics in Ion Mixing Thermodynamics of Alloy Phase Formation 8.3.1 Miedema's Theory and Alonso's Method 8.3.2 Interfacial Free Energy in the Multilayers 8.4. Experimentation of Ion-Mixing 8.4.1 Sample Design 8.4.2 Ion-Mixing Parameters 8.4.3 Characterization Methods 8.5. Amorphous Phase Formation 8.5.1 Glass-Forming Ability of Systems with a Negative Heat of Formation 8.5.2 Glass-Forming Ability of Systems with a Positive Heat of Formation 8.5.2.1 Amorphous Alloys Formed Within Restricted Compositions 8.5.2.2 Amorphous Alloys Formed in a Broad Composition Range 8.5.2.3 Nominal and Intrinsic Glass-Forming Ability 8.6. Formation of Metastable Crystalline Alloys 8.6.1 Structural Classification of the Metastable Crystalline Phases 8.6.1.1 Solid Solutions 8.6.1.2 h.c.p.-I and f.c.c.-I Phases 8.6.1.3 h.c.p.-II and f.c.c.-II Phases Based on b.c.c. Metals 8.6.2 Free Energy Calculation of the MX Phases 8.7. Interface-Generated Solid-State Vitrification in Systems with a Positive Heat of Formation 8.8. Concluding Remarks Acknowledgments References
197 199 200 200 201 203 203 205 205 205 206 207 207 207 211 213 213 213 213 215 216 218 219 220 220
CHAPTER 9 PHYSICAL VAPOR DEPOSITION
225
J. S. Colligon
ix
9.1. 9.2. 9.3.
Introduction Development of PVD Deposition Methods 9.3.1 Evaporation 9.3.2 Molecular Beam Epitaxy (MBE) 9.3.3 Sputtering 9.3.4 Ion-Assisted Deposition 9.3.5 Magnetron Sputtering 9.4. Influence of Energy on Coatings 9.4.1 Film Morphology and Density 9.4.2 Nucleation and Adhesion 9.4.3 Metastable Phases 9.4.4 Microstructure 9.5. Applications of PVD Coatings 9.5.1 Optical Coatings 9.5.2 Corrosion Protective Coatings 9.5.3 Hard Coatings 9.6. Future Trends Acknowledgment References
225 225 228 228 229 229 231 235 238 238 239 240 240 243 243 245 246 248 249 249
C H A P T E R 10 CHEMICAL VAPOR DEPOSITION
257
E Teyssandier and A. Dollet
10.1. Introduction 10.1.1 Presentation 10.2. Gas-Phase Transport and Reactivity 10.2.1 Non-Reactive Fluid Flow 10.2.1.1 Geometrical Effects 10.2.1.2 Influence of Operating Conditions 10.2.2 Reactive Flows 10.2.2.1 Thermodynamic Approach 10.2.2.2 Gas-Phase Mechanisms and Kinetics 10.2.2.3 Simulations of the Gas-Phase Composition 10.2.2.4 Nucleation and Growth of Solid Particles from the Gas Phase 10.3. Solid Phase Formation 10.3.1 Solid-Gas Thermodynamic Equilibrium Approach 10.3.2 The Structure of CVD Layers 10.3.3 The Driving Force for Crystal Growth from the Vapor Phase 10.3.4 Surface Mechanisms and Kinetics 10.3.4.1 Rate of Elementary Surface Processes 10.3.4.2 Estimation of the Kinetics of Surface Processes in CVD 10.3.5 Heterogeneous Nucleation 10.3.5.1 Nucleation Theories X
257 258 259 260 260 261 261 263 264 264 265 267 267 268 270 271 271 272 273 274
10.3.5.2 10.3.5.3 10.3.5.4 10.3.5.5 10.3.5.6 10.3.5.7
Two- and Three-Dimensional Nucleation Epitaxial Nucleation and Growth Multi-Component Nucleation and Chemical Reactions Nucleation Enhancement Selective Vapor Deposition Theoretical and Experimental Studies of Nucleation in CVD 10.3.6 Crystal Growth 10.3.6.1 Mechanisms of Crystal Growth 10.3.6.2 Classification of Crystal Faces 10.3.6.3 Crystalline Morphology 10.3.6.4 Morphology and Normal Growth Rate of Crystal Faces 10.3.6.5 Relation between Preferred Orientation and Morphology of Polycrystalline Films 10.3.6.6 Theoretical Studies of Crystal Growth from the Vapor Phase 10.4. Conclusions References
275 275 276 276 277 277 278 278 278 279 279 281 281 281 282
C H A P T E R 11 COMBUSTION
SYNTHESIS
289
S. B. Bhaduri and S. Bhaduri
11.1. Introduction 11.2. Thermodynamic Considerations 11.3. Kinetic Considerations 11.4. Field-Activated Combustion Synthesis 11.5. The "Azide" Process 11.6. SHS Reactions in Synthesizing Ti3SiC2 11.7. Controlled Reactions in the Ti-B Binary System 11.8. Auto-Ignition Synthesis of Nanocrystalline Oxides 11.9. Non-Equilibrium Effects 11.10. Concluding Remarks Acknowledgments References
289 291 294 295 297 299 301 304 307 307 308 308
C H A P T E R 12 NANOSTRUCTURED
MATERIALS
313
c. Suryanarayana and C. C. Koch
12.1. 12.2. 12.3.
Introduction Classification Preparation 12.3.1 Inert Gas Condensation
313 313 314 314 xi
12.3.2 Rapid Solidification 12.3.3 Electrodeposition 12.3.4 Chemical Reactions 12.3.5 Mechanical Attrition 12.3.6 Devitrification 12.4. Structure 12.4.1 Microstructure 12.4.2 Atomic Structure of the Crystal Lattice 12.4.3 Atomic Structure of the Grain Boundaries 12.4.4 Triple Junctions and Higher-Order Grain Junctions 12.5. Stability 12.5.1 Grain Growth in Nanocrystalline Materials 12.5.2 Grain Growth at Ambient Temperature 12.5.3 Examples of Grain Growth Inhibition 12.5.4 Isothermal Grain Growth Kinetics 12.6. Particulate Consolidation 12.7. Properties 12.7.1 Diffusion and Sinterability 12.7.2 Mechanical Properties 12.7.2.1 Elastic Properties 12.7.2.2 Hardness and Strength 12.7.2.3 Ductility and Toughness 12.7.2.4 Superplastic Behavior 12.7.2.5 Deformation Mechanisms in Nanoscale Materials 12.7.3 Magnetic Properties 12.7.3.1 Fundamental Properties 12.7.3.2 Soft Magnetic Materials 12.7.3.3 Hard Magnetic Materials 12.7.3.4 Other Ferromagnetic Nanocrystalline Materials 12.7.4 Chemical Properties 12.7.4.1 Corrosion Behavior 12.7.4.2 Catalytic Properties 12.8. Applications--Present and Potential 12.8.1 Structural Applications 12.8.1.1 Cutting Tools 12.8.1.2 Nanocomposites 12.8.1.3 Superplastic Materials 12.8.1.4 Coatings 12.8.2 Magnetic Applications 12.8.3 Catalysts and Hydrogen Storage Materials 12.8.4 Functional NanostructuresmElectronic Applications 12.9. Concluding Remarks References
xii
316 317 318 318 319 319 320 321 321 322 323 323 324 324 325 326 327 327 328 328 329 329 331 331 333 333 333 333 335 335 335 335 337 337 337 338 339 339 340 340 341 341 342
CHAPTER 13 POWDER CONSOLIDATION
347
J. R. Groza
13.1. 13.2. 13.3.
Introduction Metastability in Powder Consolidation Consolidation of Metastable Powders 13.3.1 Sintering Mechanisms 13.3.2 Scaling Laws 13.3.3 Powder Contamination 13.4. Consolidation Methods 13.4.1 Conventional Sintering 13.4.1.1 Cold Compaction 13.4.1.2 Warm Compaction 13.4.1.3 Conventional Sintering of Nanopowders 13.4.1.4 Nano-Composite Densification 13.4.1.5 Grain-Size Control 13.4.2 Pressure Consolidation of Metastable Powders 13.4.3 Pressure-Assisted Consolidation Methods 13.4.4 Non-Conventional Sintering Methods 13.4.4.1 Microwave Sintering 13.4.4.2 Field-Assisted Sintering 13.4.4.3 Shockwave Consolidation 13.5. Concluding Remarks Acknowledgments References
347 348 349 352 353 355 356 356 356 358 359 361 361 363 364 366 366 366 368 368 369 369
CHAPTER 14 BULK AMORPHOUS ALLOYS
375
A. Inoue
14.1. 14.2. 14.3. 14.4. 14.5. 14.6.
14.7. 14.8. 14.9.
History of Bulk Amorphous Alloys Dominant Factors for High Glass-Forming Ability Continuous Cooling Transformation of Alloys with High Glass-Forming Ability Preparation Methods and Elemental Effect Structural Relaxation and Glass Transition Physical Properties 14.6.1 Density 14.6.2 Electrical Resistivity 14.6.3 Thermal Expansion Coefficient Mechanical Properties Viscoelasticity Soft Magnetic Properties xiii
375 376 380 383 387 392 392 393 394 395 399 403
14.9.1 Formation and Soft Magnetic Properties of Bulk Amorphous Alloys 14.10. Viscous Flow and Micro-Formability of Supercooled Liquid 14.10.1 Feature of Phase Transition of Bulk Amorphous Alloys 14.10.2 Deformation Behavior of Supercooled Liquid 14.10.3 Micro-Forming of Supercooled Liquid 14.11. Applications and Future Prospects References
403 406 409 409 412 413 413
Author Index
417
Subject Index
419
xiv
Series Preface
My editorial objective in this new series is to present to the scientific public a collection of texts that satisfies one of two criteria: the systematic presentation of a specialised but important topic within materials science or engineering that has not previously (or recently) been the subject of full-length treatment and is in rapid development; or the systematic account of a broad theme in materials science or engineering. The books are not, in general, designed as undergraduate texts, but rather are intended for use at graduate level and by established research workers. However, teaching methods are in such rapid evolution that some of the books may well find use at an earlier stage in university education. I have long editorial experience both in covering the whole of a huge field- physical metallurgy or materials science and t e c h n o l o g y - and in arranging for specialised subsidiary topics to be presented in monographs. My intention is to apply the lessons learnt in more than 35 years of editing to the objectives stated above. Authors (and, in some instances, editors) have been invited for their up-to-date expertise and also for their ability to see their subjects in a wider perspective. I am grateful to Elsevier Science Ltd., who own the Pergamon Press imprint, and equally to my authors, for their confidence.
ROBERT W. CAHN, FRS
(Cambridge University, UK)
XV
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Preface
Metals and materials have been used for a long time. But the rapid technological developments during the later half of the twentieth century demanded materials that are stronger, capable of use at much higher temperature, more corrosion-resistant, and much less expensive than those currently in use. These demands become even more significant when we are on the threshold of the new century and the millenium. Significant improvements in properties can only be achieved by processing the materials under far-from-equilibrium (or non-equilibrium) conditions. Accordingly, several new processing technologies have been developed during the past few decades, including rapid solidification, spray-forming, mechanical alloying, ion-mixing, vapor-deposition, laser processing, plasma-processing, etc. New types of materials such as amorphous alloys and nanostructured materials have also been developed. Professor Robert Cahn has put these developments in a proper perspective in his introductory chapter, which also sets the tone for the chapters that follow. Remarkable advances have been made in recent years in the science and technology of these processes used to synthesize, characterize, and apply materials processed under non-equilibrium conditions. There have been many success stories. Some of these techniques have evolved from laboratory curiosities to commercial-scale manufacturing in a few short years. In other cases, industrial necessity prompted development of the technology, and the science followed later. The very large number of technical papers published in these different areas have been summarized in reviews, monographs, or conference proceedings, but have almost always been devoted to a single processing technique. I felt that a book of this sort was needed to provide a critical and comparative study of the different processes in a single volume. The different chapters in this book have been written by people who are worldrecognized experts in their respective fields. Each chapter describes the principles, processing techniques, special features of the materials produced and their applications. An extensive list of references is provided at the end of each chapter which will facilitate location of additional information on specific aspects of any technique. Even though great care has been taken to provide coherency and uniformity amongst the different chapters, the individual style of the authors has not been completely sacrificed. After all, variety is the spice of life! The book is primarily intended for use by graduate-level students of materials science and engineering. However, scientists beginning their research careers in any of the processes covered in this book can use it to get a fairly good knowledge and overview of the state of the technique and its potential. Scientists who wish to enter this area of research can also use this book to get a comparative idea of the potential and limitations of the different non-equilibrium processing techniques.
xvii
I wish to thank all the contributing authors for their considerable efforts in summarizing the voluminous literature of their specialties, and their dedication, diligence, and patience in working with me (and tolerating the many editorial iterations). The encouragement provided by Robert Cahn (Cambridge, UK) and John Moore (Golden, CO, USA) is greatly appreciated. I would like to express my appreciation to Phil Mestecky and his staff at Elsevier for bringing this book out in a timely and attractive fashion. Finally, I wish to thank my wife Meena for all her support during the preparation of this volume.
Golden, CO, USA August 31, 1998
C. SURYANARAYANA
xviii
List of Contributors
E V. ANANTHAPADMANABHAN Laser and Plasma Technology Division, Bhabha Atomic Research Centre, Mumbai400 085, India S. BHADURI Department of Metallurgical and Mining Engineering, University of Idaho, Moscow, ID 83844-3024, U.S.A. S. B. BHADURI Department of Metallurgical and Mining Engineering, University of Idaho, Moscow, ID 83844-3024, U.S.A. R. W. CAHN Department of Materials Science and Metallurgy, University of Cambridge, Pembroke Street, Cambridge CB2 3QZ, U.K. J. S. COLLIGON Manchester Metropolitan University, Chester Street, Manchester M1 5GD, U.K. A. DOLLET Institut de Science et de Grnie des Matrriaux et Procrdrs, CNRS, UPR8521, Universitr, Avenue de Villeneuve, F-66860 Perpignan Cedex, France J. R. GROZA Department of Chemical Engineering and Materials Science, University of California at Davis, One Shields Ave., Davis, CA 95616, U.S.A. A. INOUE Institute for Materials Research, Tohoku University, Sendai 980-8577, Japan K. N. ISHIHARA Department of Socio-Environmental Energy Science, Kyoto University, Kyoto 6068501, Japan H. JONES Department of Engineering Materials, University of Sheffield, Mappin Street, Sheffield S1 3JD, U.K.
xix
K. E KOBAYASHI Department of Manufacturing Science, Graduate School of Engineering, Osaka University, 2-1 Yamadaoka, Suita, Osaka 565-0871, Japan C. C. KOCH Department of Materials Science and Engineering, North Carolina State University, Raleigh, NC 27695-7907, U.S.A. E. J. LAVERNIA Department of Chemical & Biochemical Engineering and Materials Science, University of California, Irvine, CA 92697-2575, U.S.A. B. LI
Department of Chemical & Biochemical Engineering and Materials Science, University of California, Irvine, CA 92697-2575, U.S.A. B. X. LIU
Department of Materials Science and Engineering, Tsinghua University, Beijing 100084, China C. SURYANARAYANA The George S. Ansell Department of Metallurgical and Materials Engineering, Colorado School of Mines, Golden, CO 80401-1887, USA E TEYSSANDIER Institut de Science et de G6nie des Mat6riaux et Proc6d6s, CNRS, UPR8521, Universit6, Avenue de Villeneuve, F-66860 Perpignan Cedex, France N. VENKATRAMANI Laser and Plasma Technology Division, Bhabha Atomic Research Centre, Mumbai400 085, India
XX
Chapter 1 Introduction ROBERT W. CAHN
Man has used some of his materials in non-equilibrium form for a long time: The Japanese samurai sword of the thirteenth century, with its martensitic edge; age-hardened aircraft alloys, introduced early in this century; vapor-quenched superconducting alloys in the 1930s; melt-spun transformer laminations consisting of metallic glass in the 1970s. It was the chance discovery of age-hardening by Wilm at the start of this century, supplemented by Merica's recognition that the hardening stemmed in some way from a supersaturated solid solution, that really drew metallurgists' attention to the practical importance of non-equilibrium states. From the mid-1930s onwards, crystallographers like Guinier and Preston, physical chemists like Turnbull and physical metallurgists like Nutting, devoted their skills to unravelling the nature of age-hardening, at the same time as other metallurgists such as Kurdyumov in Russia and Cohen in the USA began the long task of understanding the details of the martensitic transformation in steels. Metastability had become a way of life for some metallurgists. The attitude of metallurgists to metastability underwent a sudden sea-change when Pol Duwez in California, in 1960, invented what at first came to be called splat-quenching (a term which, incidentally, he hated) and later, rapid solidification processing. This novel technique, centred on rapid quenching of hot molten alloys, introduced to the materials community things like hugely supersaturated solid solutions and, particularly, metallic glasses. Like all important innovations, this one did not spring fully formed from its progenitor's fertile imagination. Some years earlier, Duwez had studied the cooling rates accessible with a thin solid foil and a fast-moving stream of helium, and found them limited to around 10,000 K/s. A number of engineers had used rapid freezing, for instance of wires, purely as a means of manufacturing convenient shapes cheaply. Duwez realized that, to make progress, he had to quench, not a solid but a melt. What was new about Duwez's researches was that, for him, rapid freezing was a means to enhance accessible cooling rates by a large factor, and thereby to generate new states of matter: it was a scientific study in the first instance, and it led to technological advances as a by-product. What went on at California Institute of Technology at the beginning of the 1960s was a textbook illustration of the principle that technological innovation often (by no means always) springs from new scientific insights, and from experiments designed to create such insights.
2
R.W. Cahn
Once it had become clear that strength, in particular, could be greatly enhanced by supersaturating solid solutions and then heat-treating such non-equilibrium structures, metallurgists began to examine other ways of achieving such structures. Mechanical alloying began in the nickel industry as a way of making useful nickel-base alloys, and was then adopted by academically inclined researchers as a way of producing alloys in non-equilibrium states: in particular, alloys such as Ni3A1 which in equilibrium are fully ordered could be produced in disordered form, and some of these artificially disordered alloys then turned into glasses; this was a kind of two-stage metastability. Physical vapor deposition, which as mentioned above had been studied since the 1930s, was looked at far more systematically, and new devices such as lasers (that notorious family of solutions looking for a problem) were pressed into service both for this and for other processing strategies. The microelectronic industry put great effort into chemical vapor deposition and into ion implantation and ion mixing, and some of this work gave new ideas to classical metallurgists. Some years ago, an influential committee of the National Academies of Science and Engineering in the USA pointed to the overwhelming importance of improvements in processing to the national economy. Non-equilibrium processing of the types mentioned can be said to have responded to this need. Research was not limited to the academic (as in much recent work on mechanical alloying and milling) or to the tiny, as in microelectronics. New major processing techniques were improved: plasma-processing and spray-forming are large-scale techniques which happen to incorporate non-equilibrium states of matter. Another approach, mostly pioneered in Russia, was combustion synthesis (which, like splat-quenching, passed through several names); long ago, this approach was used to weld together streetcar rails in situ, but it has come a long way since then, and recently has turned full circle as a way of improving casting technique. Finally, and in some ways the most exciting, is the development of research on nanostructures, introduced to the world of modem science by Gleiter in Germany. Nanostructural features affect both functional materials (quantum dots, soft magnets) and structural materials. A century earlier, something similar had formed the subject-matter of the science of colloid science, but this never really took off as a wholly respectable science: maybe, in part, this was because early colloid scientists had poured cold water on the hypothesis that giant (macro-) molecules can form from organic monomers, and held back the development of polymer science by a couple of decades at least. It is always hard to forgive willful obscurantists! Now, many years later, colloid scientists (they still exist) have made fascinating discoveries about "colloidal crystals", consisting of equisized colloidal particles rather than of atoms. There have been books devoted to rapid solidification, to nanostructures, to metallic glasses (before they found application in that most cogent of industries, that of golf), to laser-processing and to other techniques toombut this, I do believe, is the first book to cover these variegated techniques in a critical and comparative way. I have long known Prof. Suryanarayana as a masterly editor, and here he has excelled himself.
Chapter 2 Thermodynamics and Kinetics of Metastable Phase Formation 2.1. 2.2.
Introduction Thermodynamics of Metastable Phase Formation 2.2.1 Free Energy of Elements 2.2.2 Free Energy of Alloy Phases 2.2.3 Determination of Free Energy of Metastable Phases 2.2.4 Lattice Parameter of the Supersaturated Phase 2.3. Kinetics of Metastable Phase Formation 2.3.1 Nucleation of Metastable Phases 2.3.2 Nucleation of Alloy Phases 2.3.3 Crystal Growth Rate of Metastable Phases 2.4. Summary 2.5. List of Symbols References
5 5 5 7 9 11 12 13 15 18 19 19 20
This Page Intentionally Left Blank
Chapter 2 Thermodynamics and Kinetics of Metastable Phase Formation KEIICHI N. ISHIHARA
2.1. I N T R O D U C T I O N
Metastable phases are commonly seen in nature. Typical examples are the occurrence of martensite in steel and diamond form of carbon. Recently, non-equilibrium processes such as rapid solidification, mechanical alloying and vapor deposition have been developed enabling formation of a variety of metastable phases. The fundamentals of some of these processing techniques have been reviewed (Turnbull 1981, Shingu 1992, Boettinger and Perepezko 1993, Herlach 1994, Baricco 1995). In this chapter, the thermodynamics and kinetics of metastable phase formation will be treated. The thermodynamics of competitive phases will be first discussed followed by the nucleation problem of metastable phase formation. Thermodynamically, reaction or phase transformation can occur only when the free energy of the system decreases. However, there is the possibility that it may not occur. For example, the free energy of diamond is higher than that of graphite at ambient pressure and room temperature. But diamond is quite "stable" at room temperature and pressure without any noticeable change occurring. Put differently, thermodynamics tells us that a reaction cannot occur when the free energy of the system increases. In the following paragraphs, we will show how the formation of a metastable phase can be expected to occur. Then we will describe how a phase can be formed under a wide range of thermodynamical possibilities. This is kinetics in contrast to thermodynamics, which only tells us whether, or not, a phase can form. In this part, the classical nucleation theory will play an important role. Using this theory, Ostwald's step rule will show that it is the metastable and not the stable phase that is transformed.
2.2. T H E R M O D Y N A M I C S O F M E T A S T A B L E P H A S E F O R M A T I O N
2.2.1 Free energy of elements Figure 2.1 shows schematically the free energies of the or,/~, y and 3 solid phases and the liquid phase of a hypothetical element. In this case, the/3, V and 6 solid phases are metastable with respect to the ot phase since the free energies of these phases are higher
6
K.N. Ishihara
0
6 solid
2 solid solid Liquid
0
Tg
T~m
T~
Tm
Temperature Figure 2.1. Schematic free energy hierarchies.
than that of the u phase at all temperatures. Tm ~ represents the melting point of the ot phase, where the free energy of the u phase is the same as that of the liquid phase. Similarly, Tm ~ and Tmy represent the melting temperatures of the fl and y phases, respectively. The Gibbs free energies of the different competing phases can be compared, at fixed temperature and pressure, to determine the relative stabilities of the phases under equilibrium. At a temperature higher than Tm ~ the free energy of the liquid phase is the lowest, and those of the c~,/3, y and 6 phases are higher in that order. This means that while liquid is the most stable phase under these conditions, the or, /3, y and 6 phases are less stable. We call this the hierarchy of stability. The hierarchy of stability of the different phases changes through a crossover of the free energy curves during a change in temperature and/or pressure. We will, however, ignore the pressure effect in this chapter. In this case, for the range of undercooling from Tm ~ to Tm ~, ot is the stable phase, the free energy of which is less than that of the liquid phase. Only the u phase can be formed from the undercooled liquid in this temperature region. Below Tm ~ it is possible to solidify the fl phase directly from the liquid. Therefore, there are two different ways in which the liquid can solidify. Thermodynamics does not tell us which phase will form preferentially from the liquid state. This will be discussed in the section dealing with the kinetics. In Fig. 2.1, the free energy of the 6 phase is higher than that of the liquid phase at all temperatures. Therefore, the 6 phase cannot form from the undercooled liquid. However, it may form from the vapor phase, depending on the relative position of the free energy of the vapor and 3 phases. The position of the amorphous phase is also indicated in this figure. The amorphous phase is the frozen state of an undercooled liquid. Since the free energy of the amorphous
Thermodynamics and Kinetics of Metastable Phase Formation phase depends on the frozen temperature, called the glass transition temperature Tg, the free energy cannot be precisely determined. The free energy of the amorphous phase is always higher than that of the liquid phase. Experimental determination of the free energy hierarchy has been performed for various elements. When the difference of specific heat between the liquid and solid phases is known, the difference in free energy, A G can be calculated using the equation:
A G - - AHf(Tm Tm - T) - fT Tm ACpdT + T fTTm ACP T dT
(2.1)
where AHu is the enthalpy of fusion, Tm is the melting temperature, and ACp is the difference of specific heat between the liquid and solid phases as a function of temperature. If no experimental data are available, the best approximation is to set A Cp = 0. The various expressions for this can be found in Kelton (1991). The free energies of different metallic elements are summarized by SGTE (Scientific Group Thermodata Europe) and can be found in CALPHAD (1991, 15, 317).
2.2.2 Free energy of alloy phases The minimum degree of undercooling which is thermodynamically necessary for the diffusionless transformation of a liquid alloy into a solid solution for a given alloy composition, is expressed by the To curve in the phase diagram. Figure 2.2 shows part of a hypothetical binary phase diagram and the free energy vs. composition curves at a temperature 7"1. Since the common tangent from A to B is lower than the free energy of either the single phase liquid or solid phases, the most stable state at the concentration from CA to C B is a mixture of the liquid and solid phases. This is indicated as the two-phase region in the phase diagram. However, in the case of partitionless solidification that is often observed during rapid solidification, the position of the free energy curve becomes important. Thus, at 7'1, in the concentration range below Co, the solid phase is more stable than the liquid phase and above Co the liquid phase is more stable than the solid phase. The trace of Co at which the solid and liquid phases have the same free energy can be expressed in the phase diagram; it is called To curve indicated by the dashed line in Fig. 2.2. If the To curve plunges to very low temperatures, single-phase crystals cannot form from the melt. In this case, phase separation from the liquid phase depresses the solidification temperature to the glass transition temperature, Tg, where the liquid transforms to the amorphous phase. The system with the To curves, only slightly depressed below the stable liquidus, is a good candidate for partitionless solidification in the entire composition range (Herlach 1994). It is very important to know the position of the To curves in the phase diagram. Since the equilibrium phase diagram is a way of expressing thermodynamic data, there might be a simpler way of finding the approximate position of the To curve from the general features of the solidus and liquidus lines. The simplest guess, which often works, is to draw the To curve connecting the midpoints between the liquidus and solidus lines at a given temperature.
K. N. Ishihara
uid
~
0
0
i i i i
I
CA
CO
CB
i
(D .4,..a
c~ ~D T1
iquidus Solidus ~
Concentration Figure 2.2. Part of a hypothetical binary phase diagram and the free energy curves of the solid and liquid phases at temperature T 1.
If we know the free energy as a function of temperature and composition, the metastable phase diagrams including To curves and extrapolation of phase boundaries can be drawn. The free energy function of the binary system is now available for the purpose of calculating the phase diagram. This calculation of the phase diagram using the free energy function is called the CALPHAD (CALculating the PHAse Diagram) method (Saunders and Miodownik 1998). Thermocalc program module that contains the database of these functions is useful for this purpose. The recent edition of the ASM handbook on phase diagrams includes reference to the calculation of the phase diagrams (Massalski et al. 1990). The specific journals available for the assessment of free energy functions in connection with phase diagrams are "CALPHAD" and Journal of Phase Equilibria.
Thermodynamics and Kinetics of Metastable Phase Formation
9
2.2.3 Determination o f free energy o f metastable phases
Direct measurement of thermodynamic data of metastable phases is difficult and only a limited data is available. This has been done by measuring either the heat released during transformation of the metastable phase to the stable form or by measuring the electromotive force (emf). Measurement of the heat released on transforming to the stable phase is one of the simplest methods. Usually this measurement has been performed by DSC (Differential Scanning Calorimetry) (Michaelsen et al. 1997), especially to determine the heats of crystallization of the amorphous phases. The specific heat can also be measured by DSC. The specific heat and enthalpy of transformation for a metastable solid phase have been measured and the free energy is calculated using Eq. (2.1) for the pure metal gallium. An example of the experimentally determined free-energy hierarchy for gallium is shown in Fig. 2.3 (Bosio et al. 1973). ot is the stable phase and fl and y are metastable phases in Fig. 2.3. Measurements of emf for the stable and metastable phases have also been reported (Ishihara et al. 1993, G~xtner et al. 1996). In this method, only the partial quantity of free energy, activity, can be directly measured. Such a measurement for the Ag-Cu system is shown in Fig. 2.4. When any thermodynamic value is not available, the enthalpy of mixing based on Miedema's macroscopic atom model can be used (de Boer et al. 1988). This approach gives us not only the enthalpies of mixing for the liquid and solid phases for a binary alloy, but also the interfacial energy between solid phases. Since in the macroscopic atom model, atoms are conceived as 'blocks' of the element, assuming the size of the block as the grain size the enthalpy of the structure including the grain boundaries can be obtained. The enthalpy of nanostructures, frequently obtained by mechanical alloying and other
a 1500
~,
-
~ 10oo 500
0
100
200
300N~
T e m p e r a t u r e (K) Figure 2.3. Experimentallydeterminedfree energyhierarchyfor Gallium (Bosioet aL 1973).
10
K. N. I s h i h a r a
10000
.
.
.
.
.
.
.
.
.
I 9
1
@
1000 I:- ~
't
0 1O0
> r @
10
1
0.0
0.2
0.4
0.6
0.8
1.0
Cu Figure 2.4. Activity of Cu in Ag-Cu solid solution phase measured by the emf method. Solid circles and diamonds show the values measured using the CuSO4 solution and the solid electrolyte RbCulCI, respectively.
(a)
(b)
~xo
9
A
B
Concentration
A
B
Concentration
Figure 2.5. Changes of enthalpy by grain refinement for the case of systems with (a)
positive heat of mixing and (b) negative heat of mixing.
Thermodynamics and Kinetics of Metastable Phase Formation
11
methods, can be evaluated by this method. In this case, the enthalpy of the nanostructure should be close to that of the liquid or amorphous state so that the grain size reaches nanometer dimensions (Bakker 1998). The enthalpy for the simple mixing of powders can be indicated by a horizontal line in Fig. 2.5. When the grain size becomes small, the total enthalpy of grain boundary, AHgb changes to closely resemble the enthalpy of mixing of the liquid phase. In the case of a system which has a positive heat of mixing, A Hgb should raise its value, while in the case of the negative heat of mixing, A Hgb should increase at both low and high concentration ranges but may decrease in the middle concentration range, as shown in Fig. 2.5 (a) and (b). It means that during mechanical alloying, the enthalpy of mixing of the solid phase does not always increase and approach that of the liquid. 2.2.4 Lattice parameter o f the supersaturated phase
The lattice parameter of a solid solution phase varies linearly with composition; this is known as Vegard's law. If the crystal structures of the terminal solid solutions are different, the average atomic volume varies linearly with the concentration. A positive deviation from Vegard's law is often seen in nature (Najafabadi et al. 1993). Extension of solid solubility is commonly observed during non-equilibrium processing and the deviation from Vegard's law has been confirmed for many systems and for wide composition ranges. Several attempts have been made to explain this deviation on the basis of the theory of elasticity (Gschneidner and Vineyard 1962). However, this deviation is attributable to the excess energy of the solid and can be thermodynamically explained using the enthalpy of mixing. Here, a simple idea will be presented. One must first consider how much energy can be stored by the deviation from the ideal volume which is shown by the Vegard's law. This energy E can be expressed as (Ziman 1972): E = (B/v)(AV/V)
(2.2)
where B is the bulk modulus, 7/is the Griineisen constant and A V is the change in the unit cell volume. In the case of a system involving metals with a cubic structure, A V can be expressed in terms of the lattice parameter a, as A V~ V ~- 3 A a / a
(2.3)
where V and a are the ideal unit cell volume and lattice parameter, respectively, according to Vegard's law and Aa is the deviation of the lattice parameter from the ideal value. This energy should be equal to the excess enthalpy of mixing. In the case of a regular solution model, it can be shown that E
= CACB~']
(2.4)
where f2 is the interaction parameter. Then, A a / a = (v/3B)CACB~2
(2.5)
K. N. Ishihara
12 0.006
!
!
!
i
0.004
<~ 0.002
0.000
,
0.0
I
0.2
,
I
I
0.4
0.6
,
I
0.8
1.0
Cu Figure 2.6. Deviation of the lattice parameter form Vegard's law, Aa for the AgCu system. The open and filled circles show the reported results of ball-milled and rapidly-quenched samples, respectively (Najafabadi et al. 1993).
In this model, the deviation from Vegard's law for the regular solution model is parabolic, which is consistent with the reported experimental data. Moreover, the sign of the deviation is the same as the sign of the heat of mixing. This is the reason why a positive deviation from Vegard's law is often seen in nature. This is because, in the case of the system that has a negative heat of mixing, intermetallic compounds are formed rather than a supersaturated solid solution. In the Ag-Cu system, the calculated value of Aa calculated using the heat of mixing by the CALPHAD method (Murray 1984) matches moderately well with the experimental results shown in Fig. 2.6 for y = 2 and B calculated as the average of 1.0 x 10 ll N/m 2 for Ag and 1.5 x 10 ll N/m 2 for Cu.
2.3. K I N E T I C S OF M E T A S T A B L E P H A S E F O R M A T I O N
Many metastable phases are found in nature and have also been produced by the recently developed novel materials processing techniques. According to thermodynamics, as mentioned above, there is a possibility for both stable and metastable phases to form. So, which is more probable? This question can be answered by the kinetics of the process. There are two main reasons why metastable phases form. The first reason is that a metastable phase can nucleate before the stable phase can nucleate. The second reason is that the crystal growth rate of the metastable phase exceeds that of the stable phase. There have been attempts to draw kinetic phase diagrams (Chvoj 1983). However, no systematic study in this field has been performed as compared to the thermodynamic in-
Thermodynamics and Kinetics of Metastable Phase Formation
13
vestigations. In this section, the nucleation and the growth rates of the metastable and stable phases will be discussed.
2.3.1 Nucleation of metastable phases The free energy hierarchy, the relative magnitude of the driving forces for the possible phase transformations, does not tell us which phase is more likely to form under a given set of conditions. It is well known that according to the Ostwald step rule (Ostwald 1897) the first phase to appear will be the metastable phase which has a higher free energy than the stable phase. Here, the proof of the rule using the classical nucleation theory will be presented (Ishihara et al. 1985). The relative magnitude of the homogeneous nucleation frequencies for various phases from undercooled liquid metals and alloys is considered. The volume nucleation frequency, Iv, at a temperature T given by the classical nucleation theory can be expressed as follows:
Iv--Kvexp
-
AG*f(O)) kT
(2.6)
where Kv is a kinetic parameter, 0 is the wetting angle of the nucleus on the surface of a substrate material, and k is the Boltzmann's constant. A G* is the excess free energy of the critical nucleus and is expressed as: 16zr o'3l) 2 AG* -- -~- ( A G 2 )
(2.7)
where o" is the surface free energy between the liquid and solid, v is the molar volume, and A G is the free energy difference between the liquid and solid phases, o" may be empirically expressed as a function of the latent heat and molar volume as: o- -- 0.065AHfv -2/3
(2.8)
where A H f is the enthalpy of fusion. AG may be simply expressed by Eq. (2.1) when
ACp = O ,
as
AG = A H f ( T m - T)/Tm
(2.9)
where Tm is the melting temperature. From these equations, the following expression can be obtained: AG* = m167r(0.065)3AHf
3
T2
(2 10)
(Tin - T) 2"
The magnitude of Kv depends on the atomic transport rate in the liquid across the solidliquid interface. The value of Kv for the nucleation of the metastable and the stable phases may be considered equal since both the phases appear from the same liquid phase. So, in the case of homogeneous or fixed wetting angle nucleation, the nucleation frequency
K. N. Ishihara
14
expressed by Eq. (2.6) may be compared by calculating the differences between the value of A G* and A G .f (f refers to the metastable phase). The critical temperature, where the nucleation frequency of the metastable phase is equal to that of the stable phase, can be given as a solution to the equation AG* = AG .I
(2.11)
This equation has a significant solution under the condition that the ratio of enthalpies of fusion for stable and metastable phases, D, is less than unity:
D-
AH~f/AHf < 1.
(2.12)
When the above condition is satisfied, the solution of Eq. (2.11) is to --
A _ ~/-DA
(2.13)
Tr - x/~A where A is the ratio of the melting temperatures (T~m/Tm) and Tr is the reduced temperature (T/Tm). Since Ttm/Tm < 1, the critical temperature, to, should exist between Tm ~ and OK. A schematic illustration of AG* and A G .I vs. temperature is shown in Fig. 2.7. The condition (B < 1) should usually be satisfied because A Hf approaches A G at low temperatures and A G .I for the metastable phase is always less than that of the stable phase. Therefore, it may generally be expected that the nucleation frequency of the metastable phase is larger than that of the stable phase if the undercooling is large enough. Thus,
]! Meta~ table
o
I
0
to
A
Reduced Temperature (T/I'm) Figure 2.7. Example of the relationship of A G* for stable and metastable phases when A = 0.8 and B = 0.5. For the temperature region below t0, the nucleation of a metastable phase is more likely than that of the stable phase.
Thermodynamics and Kinetics o f Metastable Phase Formation
15
300
250 liquid --~ 3' o
200
[.~ 150 100
0
1
2
time (s) Figure 2.8. The trace of temperature during rapid cooling of Ga droplet. It is clearly shown that the sample is crystallized first into the metastable y phase (melting temperature 238 K). Then, it transforms into the metastable/3 phase (melting temperature 257 K) and finally it transforms to the stable ot phase. The free energies of these phases are shown in Fig. 2.3.
the Ostwald step rule tends to be valid for large undercooling or supersaturation. Experimental confirmation of this rule can be found in several alloy systems (Mackintosh et al. 1983, Levi et al. 1988, Cortella et al. 1993, Abe et al. 1994, Yoshioka et al. 1994, Yang et al. 1996). The sequence of phase changes during rapid cooling of Ga is shown in Fig. 2.8 (Mackintosh et al. 1983), confirming the Ostwald step rule.
2.3.2 Nucleation o f alloy phases
The question to be answered here is what would be the composition of the solid phase that nucleates first from the undercooled liquid of a given initial composition. Figure 2.9 shows schematically the possible nucleus composition at a given undercooling. For the undercooled liquid of composition Co, the formation of a solid phase with a composition between a and b reduces the free energy of the system. The magnitude of reduction in the free energy is the largest for the nucleus with the composition at c. But, the most probable nucleus composition can be explained by the relative value of the nucleation frequency. The theory of nucleation in a binary system has been presented earlier (Kamenetskaya 1967, Kelly and van der Sande 1982, Gressin et al. 1983, Ishihara, et al. 1985, Lin et al. 1988, Evans and Greer 1990, Thompson and Spaepen 1993, Ghosh 1994). According to Kamenetskaya (1967), when a solid nucleus of radius r with a concentration Cs forms
16
K. N. I s h i h a r a
e~t) Solid ~D
O Liquid
"""
i
Co
i
a Concentration
c
Figure 2.9. Schematic illustration of the free energy-concentration diagram explaining the magnitude of the driving force for nucleation of solid phase in a solid-solutionforming binary alloy system.
from a liquid of initial composition Co, the total free energy change of the system is given as"
A G --
1
3Nv
Gt(CI) + \ 3Nv
Gs(Cs) q
N
(2.14)
where N is the number of atoms in a system, Gl and Gs are the free energies of the liquid and solid phases and Ct is the liquid composition after nucleation. The mass conservation law gives the relation"
1
3--lqv- CI +
3Nv
(2.15)
The nucleus should have the composition that gives the smallest free energy increase on nucleation. Such a composition should satisfy the condition when the free energy surface shows a saddle point. Considering a binary system in which both the liquid and solid phases are ideal solutions, the surface energy can be expressed as: cr -- aA(1 -- Cs) -k- ffBCs
(2.16)
Thermodynamics and Kinetics of Metastable Phase Formation
1200
17
Co To
1000
800
e~
M 600 N
~' 400 To 3
200
O
0.0
0.2
I
i
i
0.4
0.6
0.8
P 1.0
Concentration Figure 2.10. Example of a calculated phase diagram with the ideal liquid and solid solutions. The curves SP and SQ show the most probable nucleus concentration for composition-independent surface energy and for composition-dependent surface energy, respectively.
where O"A and cr8 are the surface energies between the pure metals and the liquid. The results of calculation of the composition of the most probable nucleus as a function of temperature are shown in Fig. 2.10. This estimation agrees well with the experimental results (Fecht 1991). The area below MOSN, which is the trace of the compositions a and b in Fig. 2.9, refers to the region where the driving force for solidification from the initial liquid of composition Co is positive. The line SQ shows the trace of the saddle points. The temperature at which partitionless solidification is most probable, the intersection point of SQ and C = 0.2 in Fig. 2.10, is exactly To~3.
18
K. N. Ishihara
2.3.3 Crystal growth rate of metastable phases There have been only a few systematic studies of the growth rate. In general, the growth rate S can be expressed as a function of the undercooling by the equation (Jackson 1984)"
S-
So(AT) f
(2.17)
with So and fl (1 < fl < 2) as constants, and A T is the degree of undercooling. So and fi have no physical meaning. However, when fl of the metastable phase is larger than that of the stable phase, the growth rate of the metastable phase may become larger than that of the stable phase at a large undercooling, as shown in Fig. 2.11. Such a situation can occur when the stable phase has a faceted solid-liquid interface while the metastable phase has a molecularly rough interface (Tuckerman and Bechhoefer 1992). The experimentally investigated growth rates of the stable and metastable phase in the F e - C - X system are schematically shown (Magnin and Kurz 1988).
metastable phase
~D
o
stable phase
Tm(metastable) Tm(stable) Temperature Figure 2.11. Hypothetical relation of the growth rate for both stable and metastable phases. In this figure, fl of stable and metastable phases is 1.4 and 2.0, respectively. For SO,the value for the metastable phase is three times larger than that of the stable phase.
Thermodynamics and Kinetics of Metastable Phase Formation
19
2.4. S U M M A R Y
The thermodynamic nature of the metastable phases was discussed. The possibility of the formation of metastable phases was also discussed from the kinetic point of view-nucleation and growth of metastable phases. Furthermore, emphasis was laid on the relationship between the excess heat of mixing and the deviation of lattice parameter from Vegard's law.
2.5. L I S T O F S Y M B O L S
A a Aa B C
A Cp D E G AG A G*
AHf A Hgb
Iv Kv k N r S So T
Tm Tr V AV v fl y cr 0 f2
ratio of melting temperatures lattice constant difference of lattice constant bulk modulus concentration difference of specific heat ratio of enthalpy of fusion stored energy free energy difference of free energy excess free energy enthalpy of fusion enthalpy of grain boundary nucleation frequency kinetic parameter Boltzmann's constant number of atoms in a system radius of nucleus growth rate constant temperature melting temperature reduced temperature unit cell volume difference of unit cell volume molar volume constant Grtineisen constant surface energy wetting angle interaction parameter
20
K. N. I s h i h a r a
REFERENCES
Abe, T., Akiyama, A. and Onodera, H. (1994) ISIJ International, 34, 429. Bakker, H. (1998) Enthalpies in Alloys (TransTech Publications, Ztirich, Switzerland), p. 31. Baricco, M. (1995) Key Eng. Mater, 103, 1. Boettinger, W. J. and Perepezko, J. H. (1993) Rapidly Solidified Alloys: Processes, Structures, Properties, Applications, ed. Liebermann, H. H. (Marcel Dekker, New York), p. 17. Bosio, L., Cortes, R. and Defrain, A. (1973) J. Chem. Phys., 70, 357. Chvoj, Z. (1983) J. Non-Equilib. Thermodyn., 18, 201. Cortella, L., Vinet, B., Desr6, E J., Pasturel, A., Paxton, A. T. and van Schifgaarde, M. (1993) Phys. Rev. Lett., 70, 1469. de Boer, E R., Boom, R., Mattens, W. C. M., Miedema, A. R. and Niessen, A. K. (1988) Cohesion in Metals: Transition Metal Alloys (North-Holland, Amsterdam). Evans, E V. and Greer, A. L. (1990) in Principles of Solidification and Materials Processing, Vol. 2, eds. Trivedi, R., Sekhar, J. A. and Mazumder, J. (TransTech Publications, ZUrich, Switzerland), p. 741. Fecht, H. J. (1991) Acta Metall. Mater, 39, 1003. Gartner, E, Busch, R., Haasen, E and Bormann, R. (1996) in Metastable Phases and Microstructures, Vol. 400, eds. Bormann, R., Mazzone, G., Shull, R. D., Averback, R. S. and Ziolo, R. E (Mater. Res. Soc., Pittsburgh, PA), p. 119. Ghosh, G. (1994) Mater Sci. Eng., A189, 277. Gressin, E, Eustathopoulos, N. and Desr6, E (1983) Scripta Metall., 17, 711. Gschneidner Jr., K. A. and Vineyard, G. H. (1962) J. Appl. Phys., 33, 3444. Herlach, D. M. (1994) Mater Sci. Eng. Rept., R 12, 177. Ishihara, K. N., Maeda, M. and Shingu, P. H. (1985) Acta Metall., 33, 2113. Ishihara, K. N., Diaz de la Torre, S., Hosoe, T., Yamamoto, A. and Shingu, E H. (1993) in Processing Materials for Properties, eds. Heinen, H. and Oki, T. (TMS, Warrendale, PA), p. 667. Jackson, K. A. (1984) Mater Sci. Eng., 65, 7. Kamenetskaya, S. S. (1967) Sov. Phys. Crystallogr., 12, 71. Kelly, T. E and van der Sande, J.B. (1982) in Chemistry and Physics of Rapidly Solidified Materials, eds. Berkowitz, B. J. and Scattergood, R. O. (TMS, Warrendale, PA), p. 35. Kelton, K. E (1991 ) Solid State Phys., 45, 75. Levi, C. G., Jayaram, V., Valencia, J. J. and Mehrabian, R. (1988)J. Mater. Res., 3, 969. Lin, H. M., Kim, Y. W. and Kelly, T. E (1988) Acta Metall., 36, 2537. Mackintosh, A. J., Ishihara, K. N., and Shingu, P. H. (1983) Scripta Metall., 17, 1441. Magnin, E and Kurz, W. (1988) Metall. Trans., 19A, 1955, 1965. Massalski, T. B., Okamoto, H., Subramanian, E R. and Kacprzak, L. (1990) Binary Alloy Phase Diagrams (ASM Intemational, Materials Park, OH). Michaelsen, C., Gente, C. and Bormann, R. (1997) J. Mater. Res., 12, 1463. Murray, J. L. (1984) Metall. Trans., A15, 261. Najafabadi, R., Srolovitz, D. J., Ma, E. and Atzmon, M. (1993) J. Appl. Phys., 74, 3144. Ostwald, W. (1897) Z. Phys. Chem., 22, 22. Saunders, N. and Miodownik, A. P. (1998) CALPHAD: A Comprehensive Guide (Elsevier, Oxford). Shingu, P. H. (ed) (1992) Mechanical Alloying (TransTech Publications, Ztirich, Switzerland). Thompson, C. V. and Spaepen, E (1993) Acta Metall., 41, 2021. Tuckerman, L. S. and Bechhoefer J., (1992) Phys. Rev., A46, 3178. Tumbull, D. (1981) Metall. Trans., 12A, 695. Yang, G., Zhuang, H. and Biswas, P. (1996) NanoStructured Mater., 7, 675. Yoshioka, M., Hancock, B. C. and Zografi, G. (1994) J. Pharmac. Sci., 83, 1700. Ziman, J. M. (1972) Principles of the Theory of Solids, second edition (Cambridge University Press, Cambridge), p. 67.
Chapter 3 Rapid Solidification 3.1. 3.2.
Introduction Methods of Rapid Solidification 3.2.1 Droplet Methods 3.2.2 Spinning Methods 3.2.3 Surface-Melting Methods 3.3. Constitution and Microstructure Formation by Rapid Solidification 3.3.1 Non-Equilibrium Constitution 3.3.2 Microstructure Formation 3.4. Properties, Performance and Applications of Rapidly Solidified Materials 3.4.1 Properties and Performance 3.4.2 Applications of Rapid Solidification References Selected Bibliography
23 24 24 26 27 28 29 32 34 34 41 43 45
This Page Intentionally Left Blank
Chapter 3 Rapid Solidification H. JONES
3.1. I N T R O D U C T I O N
The term solidification simply means formation of solid so the parent material(s) can be gaseous (vapor) or liquid (e.g., melt). Thus, solidification can involve continuous or discontinuous deposition with or without a chemical reaction or passage of a solidification front through a volume of melt. The term rapid solidification is normally applied to solidification from the melt. The description "rapid" can be taken to imply a short-time interval between initiation and completion of solidification and a high velocity of propagation of the advancing solidification front. Such rapid solidification is most readily achieved by imposing a high cooling rate during solidification. This requires a thin layer, filament or droplet(s) of melt to be generated in good contact with an effective heat sink. Rapid solidification then occurs either directly, as a result of coupling between external heat extraction and the transport of latent and specific heat required to propagate the solidification front, or indirectly during the recalescence that follows nucleation of solidification at large undercooling. Rapid solidification processes such as melt-atomization, spray-deposition, splat-quenching, melt-spinning and planar flow-casting and their derivatives all generate high cooling rates during solidification and their products contain structural features formed at different levels of prior undercooling. More control over the nucleation stages of rapid solidification can be obtained by slow cooling of a volume of melt (which can be quite large) that has been freed of agents than would otherwise catalyze nucleation at low undercooling. Nucleation of rapid solidification can then be triggered by contact with a nucleating phase at predetermined levels of undercooling, though in this case the slow cool that follows completion of solidification may be insufficient to preserve to ambient temperature some of the non-equilibrium features generated by the rapid solidification step. Correspondingly, growth velocity during rapid solidification can be coupled to the velocity of a heat source (such as a laser or electron beam) traversing the surface of a block of material. In this case the heat source generates a stable melt-pool which melts new material at its leading edge as fast as previously melted material refreezes at its trailing edge. Melting and freezing front velocities increase from zero at the bottom of the melt-pool to a value related to the traverse velocity at the sample surface. The present short survey is in three sections. The first section reviews methods of rapid solidification from the melt and their characteristics. The second section reviews current knowledge of the effects of 23
H. Jones
24
rapid solidification on constitution and microstructure formation. The third section gives examples of resulting property changes and of current or potential applications.
3.2. M E T H O D S
OF R A P I D S O L I D I F I C A T I O N
These are most conveniently discussed under the headings droplet, spinning and surface melting technologies. Droplet technologies all involve generation of multiple or single droplets of melt which solidify either individually or in a state of partial or complete coalescence on a deposition surface. Spinning technologies involve stabilization of a flowing melt-stream so that it freezes as a continuous filament, ribbon or sheet in contact with a moving chill surface or heat extracting fluid. Surface melting technologies involve rapid melting at a surface followed by rapid freezing sustained by rapid heat extraction into the unmelted bulk. Cooling rates during solidification can reach 101~ K/s in cross sections as small as ~ 0 . 1 / z m with undercoolings that can attain hundreds of K, front velocities as high as 100 m/s and solidification times as short as nanoseconds.
3.2.1 Droplet methods These are all variants of the long-established technology of making lead shot by teeming molten lead through a preheated steel dish perforated by an array of equal sized holes (Johnson et al. 1976). The fundamental principle is that a pendant drop or free falling melt-stream tends to break up into droplets as a result of surface tension. This process can be intensified, for example, by impingement of high velocity jets of a second fluid, by centrifugal force at the tip of a rotating cup or disc or by an applied electric field. The standard method, known as atomization, uses high velocity gas or water jets to rapidly fragment a volume of melt into large numbers of small droplets (Fig. 3.1) (Schmitt 1979). The jets form, propel and cool the droplets that may freeze completely in flight to form powder or Metal
Metal
Connifed~ II1 Figure 3.1. Principle of spray droplet (atomization)by impingementof high velocity gas jets on to a free falling or emergent melt-stream (Schmitt 1979).
Rapid Solidification
25
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Figure 3.2. Particle size distribution for Sn-5wt%Pb alloy as a function of atomizing gas pressure (Anderson 1991).
may form individual splats or a thick deposit by impact on to a suitable substrate. Powder or shot particulate generated by such techniques displays a range of particle sizes and shapes even for a single set of operating conditions. Increasing atomizing gas pressure, for example, then increases the yield of finer particle sizes (Fig. 3.2) (Anderson 1991) while the more effective quench of water atomization leads to more irregular shapes. Although particles more or less identical in size and shape from the same powder sample can show quite different microstructures (because of different local conditions), smaller particles tend to cool more rapidly and/or undercool more prior to solidification so tend to solidify more rapidly (i.e., at higher front velocity). Splats formed from droplets of given size tend to solidify even more rapidly because of more effective heat extraction from the larger surface area they offer, especially when at least one of their surfaces is in good contact with an efficient heat sink, such as a water-cooled rotating copper drum. A spray deposit can maintain the same microstructure as the equivalent splats provided that their solidification time is sufficiently less than the time interval between deposition of successive splats at a given location on the substrate (Oguchi et al. 1990). Atomizers range in size from laboratory units with capacities of less than 1 kg per run to commercial scale facilities with capacities as large as 50,000 tonnes per year (Anon 1989). Powders destined for high performance applications tend to be atomized with inert gases or in vacuum to minimize formation of oxides or other potentially damaging inclusions. The variability of powder particle sizes, shapes and microstructures within bulk atomized powder samples, splats and spray deposits (which is acceptable for many purposes), and the need to consolidate the materials for typical engineering applications, has stimulated work on droplet methods that employ lower cooling rates and allow measurement or control of the temperature at which nucleation occurs. Such methods include emulsification techniques in which a distribution of droplets is produced within a carrier liquid which can be inert, which can act as a flux for impurities in the droplets that otherwise would activate nucleation at low undercoolings or which can trigger nucleation at a spe-
26
H. Jones
cific undercooling (Perepezko 1984). A variant uses rapid solidification to produce a fine secondary network of eutectic in a single- phase matrix. This network is then spheroidized by heat treatment above the solidus to produce a distribution of liquid droplets in the solid matrix (Prasad et al. 1984). Information regarding melting and nucleation temperatures of the droplets can then be obtained by controlled heating and cooling coupled with differential calorimetry, and the results related to microstructures formed in the solidified droplets (Kim and Cantor 1994). Undercoolings of hundreds of degrees can be generated in relatively large volumes of melt by prior removal of impurities e.g., by fluxing (Kui et al. 1984), by levitation (containerless) melting (Weber et al. 1991), or by induction remelting in contact with a water-cooled copper hearth. A recent triumph of this latter approach has been to form substantial pieces of zirconium-based metallic glasses that display ultra-high strengths and good fatigue characteristics along with extensive plasticity in the supercooled liquid range (Johnson 1996). Benefits of the related levitation melting approach have included much fundamental information on the formation of solidification microstructures at very high undercoolings and growth velocities that cannot be accessed by other rapid solidification techniques (Herlach 1994, Herlach et al. 1993). 3.2.2 Spinning methods
These derive from the simplest system in which a single melt-stream emerging from an orifice is stabilized by surface film formation or solidification before it can break up into droplets (so-called melt-extrusion or melt-spinning), to the most sophisticated in which
gas pressure
eit
heoting winding
./~
~ /orifice
/ribbon
rotot/ng drum Figure 3.3. Schematicof free-jet chill-block melt-spinning (Overshott 1979).
Rapid Solidification
27
one or more melt-streams are used to make wide or composite ribbons by impingement on single or twin chill roll surfaces. In standard free-jet chill-block melt-spinning (Fig. 3.3) (Overshott 1979) the melt-stream forms a single ribbon typically 10 to 50 # m thick and a few millimeters wide by impingement on to a single chill roll rotating at a surface speed of some tens of meters per second. In planar flow-casting the distance from the nozzle orifice to the chill surface is reduced to less than ~0.5 mm to eliminate formation of a melt-pool and associated instabilities. Thin strips up to 500 mm wide have been cast in tonnage quantities by this method, either for direct application or pulverization into platelet particulate for consolidation into other product forms. Melt-extraction (Maringer and Mobley 1974), in which the rotating chill roll forms the products by direct contact with the surface of a crucible of melt or the melted extremity of the solid electrode, and its derivative melt-overflow (Boulby and Wood 1986a, b), do not involve generating a melt-stream. Continuous rapidly solidified filament with round section can be produced by the rotating water bath process in which a melt-stream emerging from a cylindrical orifice is solidified directly on entry into a rotating annulus of water contained within the lip of a rotating water bath (Ohnaka 1985). This process can also be used to generate powder particulate when the conditions result in the break-up of the melt-stream on entry into the water bath (Raman et al. 1982). Free-jet chill-block melt-spinning and meltextraction can be used to generate a flake product directly if continuity is interrupted by a series of regularly-spaced notches on the chill-block. Melt-spun ribbon or planar flow cast sheet can be pulverised, for example, by means of a blade-cutter mill (Gelinas et al. 1988, Pelletier et al. 1990) into platelets (_~0.3 mm across). These provide a more suitable feedstock for consolidation by powder metallurgy than continuous ribbon, for example, which exhibits a much lower packing density.
3.2.3 Surface-melting methods These derive from spot or traverse welding techniques and differ from them only in that the depth melted is limited to ensure that the ensuing solidification will be sufficiently rapid. In its simplest form, a single pulse or continuous traversing heat source is used to rapidly melt the surface of a block material, the unmelted bulk acting as the heat sink during the subsequent rapid solidification (Fig. 3.4) (Lux and Hiller 1972). The resulting rapidly solidified material has the same composition as the underlying parent material, although the rapid solidification may produce a different microstructure and much improved properties. A second possibility is to preplace or inject alloy or dispersoid additions at its surface so they are incorporated into the melt-zone to form a surface region of composition different from the underlying bulk. The third possibility is to melt a different material preplaced on the surface so that mixing with the underlying material is limited to the minimum required for effective bonding. All three variants offer the practical possibility of generating a more durable surface on an underlying material that is in all other respects entirely adequate for the application in view. Both nanosecond and picosecond power laser sources have been used to generate some quite spectacular non-equilibrium effects in surface melt-zones as shallow as 0.1 /xm in which cooling rates during solidification have been estimated to reach 10 l~ K/s (von Allmen et al. 1984) or more and solidification times to be as short as 10 -9 s (Spaepen 1987). Both traversing laser and electron beams
28
H. J o n e s
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Figure
have been used to treat entire surfaces via repeated incremental lateral displacement of the beam by the width of the melt-zone at the start of each new traverse (Lux and Hiller 1972). Samples with crack-free, relatively smooth treated surfaces can now be produced by appropriate control of the process parameters. The technique can also be used to develop coupling between the traversing beam and the solidification front so that the effects on resulting microstructure of systematic variations in front velocity up to -~ 1 m/s can be determined experimentally (Boettinger et al. 1984, Zimmermann et al. 1989). (See also Chapter 5 on "Laser Processing" by K.E Kobayashi in this book.)
3.3. C O N S T I T U T I O N AND M I C R O S T R U C T U R E F O R M A T I O N BY RAPID SOLIDIFICATION
Rapid solidification produces changes in constitution because the large undercoolings and front velocities involved promote formation of non-equilibrium phases and extensions in composition range of surviving equilibrium phases. Microstructural differences include changes in mode of growth and size refinement as a result of the short diffusion distances imposed. These aspects will be dealt with separately.
29
Rapid Solidification
3.3.1 Non-equilibrium constitution Non-equilibrium phases formed by rapid solidification include glasses, quasicrystalline phases, new crystalline phases and equilibrium phases with extended composition ranges. Glass formation requires that the entire or residual melt is undercooled below the glass-forming temperature before the advance of crystal formation can complete the solidification process. Glass formation is highly sensitive to composition, with pure metallic melts in particular requiring very high cooling rates during solidification to suppress the rapid advance of crystal formation. Alloy systems showing deep eutectics, however, form glasses rather more readily in the composition ranges near to these eutectics. Critical cooling rates for glass formation can be predicted by modelling the nucleation and growth of the competing crystalline (or quasicrystalline) phases. The melt will become a glass if cooling is rapid enough not to intersect the knee of the time-temperature-transformation (TTT) curve of the competing crystal forming process (Fig. 3.5) (Davies et al. 1974). In the absence of such diagrams, easy or bulk glass formation has been variously associated with ratio of glass formation to melting temperatures, large intervals between glass formation and crystallization temperatures (the so-called "supercooled liquid" range), low thermodynamic driving force for crystallization, large negative heats of mixing, simultaneous precipitation of more than one phase, large atomic size differences between the components, as well as deep eutectic formation (Peker and Johnson 1993, Massalski et al. 1981, Inoue et al. 1993). (See also Chapter 14 on "Bulk Amorphous Alloys" by A. Inoue in 2000
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30
H. Jones 9
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this book.) Figure 3.6 (Alonso et al. 1988) shows a glass formability map for binary solute additions to nickel in which high enthalpy of formation arising from atomic size difference and large negative enthalpy of mixing are seen to correlate well with glass formation under melt-spinning conditions (cooling rate ~ 106 K/s). Formation of non-equilibrium crystalline phases by rapid solidification occurs when their nucleation and growth kinetics results in a higher rate of formation than the competing equilibrium phase. Formation of metastable (cementite) cast irons rather than graphitic cast irons at temperatures below the austenite-cementite eutectic temperature is a well-known example for relatively modest cooling rates. Possibilities for such metastable phase formation multiply manyfold at the high undercoolings and cooling rates that typify processes such as melt-spinning and splat quenching. The TTT diagram in Fig. 3.7 (Herlach et al. 1992) illustrates the competitive kinetics of formation of three contending intermetallic phases and or-A1 in undercooled Al-12 at %Mn alloy. Based on assumptions made, equilibrium A16Mn should form at cooling rates > 5 • 103 K/s, displaced by T-
Rapid Solidification
31
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Figure 3.8. Maximum reported extensions of solid solubility of various solutes in aluminum. Black bars indicate the increases beyond equilibrium solid solubilities achieved by rapid solidification (Midson and Jones 1982).
32
H. Jones
phase at 5 x 103 to 2 x 105 K/s and an icosahedral phase at 2 x 105 to 2 x 106 K/s, in accord with experimental results. Extension of solid solubility by rapid solidification, again, has been widely observed, and can amount to orders of magnitude in some cases where equilibrium solid solubilities are very low (Fig. 3.8) (Midson and Jones 1982). While limited atomic size and electrochemical differences between solvent and solute appear to be necessary conditions for such solubility extension, the actual outcome is expected to be determined by conditions at the advancing solidification front. Fig. 3.9 (Boettinger et al. 1987, Kurz et al. 1988) shows the continuous increase in solute concentration with increasing solidification front velocity for ot-Ag solid solution in the Ag-15 wt%Cu alloy. The experimental data match predictions of dendrite growth modeling and approach 15 wt%Cu (segregation-free solidification, SFS) at 0.4 m/s above which plane front solidification is expected to be stabilized. The phenomenon of solute trapping is also expected to take effect in this velocity range in which the time for incremental advance of the solidification front is less than the time required for diffusional jumping of solute atoms so that instead they are engulfed by the advancing solid. Fig. 3.10 (Boettinger et al. 1984) shows the corresponding effect of varying alloy composition as well as growth velocity. The increase in velocity required for SFS with increasing Cu alloying content up to 15 wt% can be associated with incidence of plane front solidification while the decrease in this velocity above 15 wt%Cu can be attributed to control by solute trapping (Jones 1988).
3.3.2 Microstructure formation Figure 3.10 (Boettinger et al. 1984) is also an example of a solidification microstructure selection diagram (SMSD) in which dominant solidification mode is mapped on a plot which has a material parameter, such as alloy composition, as one axis and a processing parameter, such as front velocity, as the other axis. Figure 3.11 (Boettinger 1982) shows a further example in which the competition is between glass formation, formation of primary Pd3Si, y and x phases and formation of the Pd3Si-K eutectic. Modelling of the competitive growth situation implicit in such diagrams presents a challenge for solidification theory. The fundamental premise is that the solidification mode with the highest growth temperature will dominate under given conditions. Figure 3.12 (Gilgien et al. 1995) shows a comparison between experimental fact and predictions of such competitive growth modelling for Al-rich A1-Fe alloys, which form the basis of several high performance rapidly solidified aluminium alloys. The level of agreement is encouraging. While growth temperature can be crucial in determining which constituent and which mode of solidification prevails, an equally important consideration for performance and properties is the size-scale of the resulting microstructural constituent. Measurements of grain size ~.0 vs. cooling rate T (Fig. 3.13) (Jones 1991), suggest a relationship of the form ~.0 - K l J ' - n with n -- 1 and K1 ~ 107 # m K/s for Ni-5wt%A1 and Ti6wt%A1-4wt%V and ~ 106 # m K/s for Fe-6.3wt%Si and type 316 stainless steel. A prediction for pure aluminium assuming homogeneous nucleation and linear growth kinetics gave n = 0.9 and K ~ 2 x 107 (K/s) n (Boswell and Chadwick 1977). The recent observations of nanoscale dispersions of extended and non-equilibrium phases often adjacent or related to glass-forming composition ranges has also been attributed to
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operation of homogeneous nucleation (Greer 1994). Spacings of cellular or-A1 in A1Fe and A1-Cu alloys grown at 0.03 to 1 m/s show good agreement with predictions of the tip stability model of Kurz et al. (1986) and a newer array model of Hunt and Lu (1996). An example is shown in Fig. 3.14 (Lu et al. 1994). The tip stability model predicts ,k = K 2 / V 1/2, where K2 is 4Jr(DF/ATo) 1/2, in which D is the solute diffusivity in the melt, 1-" is the ratio of interfacial energy to entropy of fusion and AT0 is the alloy freezing range. Approximate analytical models predict primary dendrite spacing )~1 = K3G1/2V 1/4 with K3 = K4(AToDI') 1/4, where G is temperature gradient, K4 is 2.83k 1/4 (Hunt 1979) or 4.3/k 1/4 (Kurz and Fisher 1981) with k as a solute partition coefficient, and experimental data lend some support to this relation (Moir and Jones 1991). Secondary dendrite arm spacing )~2 is predicted to follow )~2 -- K5 i?-n with K5 "~ 50/xm (K/s) n and n = 1/3, as exemplified for or-A1 in dilute A1-Cu and AI-Si alloys in Fig. 3.15, although the experimental data for other systems show 0.25 < n < 0.5 and 16 < K5 < 160 # m K n s -n (Jones 1984). Predictions for the interphase spacing of a regular eutectic give )~ = K 6 / V 1/2 (Jackson and Hunt 1996), where 7 < K6 < 50 # m 3/2 s 1/2 for Al-based eutectics (Juarez-Islas and Jones 1987). Figure 3.16 shows eutectic spacing A vs. V for the A1-Cu eutectic, for which K6 ~ 9.4 # m 2/3 s-1/2 right up to the velocity for extinction of stable eutectic separation at V ~ 0.5 rrds (Jones 1984, Zimmermann et al. 1989).
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!
25
~l 30
COMPOSITION (wt % Cu)
Figure 3.10. Conditions for microsegregation free solidification of a-Ag as a function of growth velocity and alloy composition for Ag-rich Ag-Cu alloys (Boettinger et al. 1984).
3.4. P R O P E R T I E S ~ P E R F O R M A N C E
AND A P P L I C A T I O N S O F R A P I D L Y S O L I D I F I E D
MATERIALS
These will be discussed under the separate headings of Properties and Performance, and Applications (Sections 3.4.1 and 3.4.2).
3.4.1 Properties and performance Metallic glasses can exhibit high flow stresses and hardness with fracture stress approaching E/50, the theoretical limit, where E is elastic modulus. This can be combined with ductile behavior in bending (Fig. 3.17) (Davis 1976), shear and compression. Young's and shear moduli are typically 20-40% lower than in the crystalline state because the atoms in the glass are not constrained by a lattice to make self-similar displacements. Bulk modulus is however only a few percent lower in the glassy state reflecting the dense random packing which gives a density only 1-2% less than for the crystalline state. Further increments in hardness and strength along with increased stiffness can be obtained by precipitation of nanocrystals within a metallic glass matrix so as to suppress the inhomogeneous shear process (Fig. 3.17) that characterizes flow and fracture of metallic glass.
35
Rapid Solidification ,
-
25 A
E
E
I0
>~-
5
m
1
I
J
~
I
o ~, ~ ~ \ 0
'
9 Gloss
\\
/
~o~._ "\ - ~ ~ \
N
~
I
t
i
i
'i
I
l
I 15
!
a
/
\',
%~
~
,/X
\~
-
_
i '
/
g
/
/
Gio,=+x z~
9
A
/k
~2.5
> ::E I'--
tw (..9
_
D Pd 3 s i + X + ~
"
/
eut~tic
///
0.5
0.25 -
"~' J
~ I
I
l
0
r 1 5
\~
\~
o //
\ \
l
/
K + euttctie
o
/
1 ~ ~ AT. % Cu
I l0
~
I
A
I
Figure 3.11. Constituent selection diagram for Pd-Cu-Si alloys for solidification front velocities of 0.25-50 mm/s (Boettinger 1982).
1000
AI plane front '
. . . . . . .
'
- ~,~,-,~ ,-,i ,-,2\",'~ . . . . . . . .
100
.=.
10
AI cell / dend. + AI-AIsFe eul.
E u
,~TITITIT11]Trr[~
L---,
>
I-l~Tr'-'-
r
/! IIIIW%
,/111 L V
.01
.'"
-- "
AI3F e dend- +
AI-AI,Fe eat
AI den~. +" - - ~ - J / ~ " "-s "-A/F~ d e n d ~ 1 2 1 ; e ;ut" ~l-Al,,Feeut,l ~ t ~ 3 -, " 3 " ". / , ~ P ~ ' I . AI-AI,'}FeeutI,,, ,
.001 0
1
2
3
" 4
Fe-content [at%] Figure 3.12. Comparison between experimental (full lines) and predicted (broken lines) solidification microstmcture selection map for Al-rich A1-Fe alloys (Gilgien et al. 1995).
H. Jones
36
lOO
104
10 ~
106
10"
"i" K/s Figure 3.13. Grain size vs. cooling rate during solidification of [[5]] Ti-6wt%Al4wt%V, (A) Ni-5wt%A1, (~) Fe-6.3wt%Si, and ( 9 type 316 stainless steel (Jones 1991). Full line represents a prediction (Boswell and Chadwick 1977) for aluminium.
This is illustrated in Fig. 3.18 (Kim et al. 1990) for an Al-based A1-Y-Ni-M (M = Mn or Fe) metallic glass. Metallic glasses exhibit a well defined fatigue limit corresponding to ,~0.35 of the flow stress. Fatigue life above this limit tends to be shorter than for crystalline materials because absence of work hardening allows cracks to initiate at relatively low stresses. Their wear resistance can be superior to that of crystalline materials of the same hardness under sliding conditions, but tends to be inferior under abrasive wear conditions. Metallic glasses undergo considerable creep deformation at temperatures close to but below their glass transition temperature Tg compared with crystalline metals and ceramics, Newtonian viscous flow is observed at much lower temperatures and up to higher stresses. Correlations between minimum creep rate and rupture time and between applied stress, minimum creep rate and temperature are shown in Fig. 3.19 (Rama Rao and Radhakrishnan 1984, Gibeling and Nix 1978) for an F e - N i - P - B glass. The homogeneous flow behavior of metallic glasses at temperatures just below Tg can be applied to hot work them into useful shapes. The increased scattering of electrons by the disordered structure and mixed population of atomic species in a metallic glass gives rise to high electrical resistivity, two or three times that of the same composition in crystalline form. In addition metallic glasses based on ferromagnetics such as iron, cobalt or nickel exhibit excellent soft ferromagnetic behavior. Nanocrystalline counterparts can exhibit superior soft or hard ferromagnetic
Rapid Solidification
37 0.5 w t . %
Dendritic
10 ~
Cellular 10-' 10 -2
~L
2wt. % 10 ~
,~
10"
t., 10 2 ~=~ e,~
10~
$=I
101
~ o~t ~g 02
4wt.
%
!
10"2
l . . . . . . . . . . . . . . . . . . ]
10 ~
..... ~ ~, . .
6wt.%
.
10"1
"~
10"2
"..................
1 8wt.
10 ~
%
10-1
10 .2 -
10 ~
10"1
,
10 .2
Velocity (cm/sec) Figure 3 . 1 4 . a - A 1 cell spacing vs. velocity of solidification front for five A1-Fe alloy compositions compared with predictions of tip stability (dashed lines) and array model (shaded regions) (Lu e t al. 1 9 9 4 ) .
behavior depending on their constitution and alloy compositions. Excellent corrosion resistance can also be developed by appropriate levels of alloying with chromium along with additions such as Mo, W, P or B. Catalytic effectiveness can also be enhanced for the glassy state compared with the crystalline condition. Crystalline materials made by rapid solidification can also show high hardness and strength by combinations of fine grain size, presence of fine disperoids/precipitation and increased solid solution hardening, without the reduced stiffness associated with the glassy state. The Hall-Petch plot shown in Fig. 3.20 (Nagahuma et al. 1993) for rapidly solidified A1-Ni-Mn(-Zr) alloys compared with wrought unalloyed aluminium illustrates both the increased direct contribution of particle and solid solution hardening ('~500 MPa) and the indirect effect of stabilizing small grain size with an increased Hall-Petch slope
38
H. J o n e s
104
I
i V
103
i
w
i
I
i
i
,
~
I
I
I
I
i
i
=
I
I
1,
i
~'.
~2 102 l.Lm 10 I
0.1 I
0.01
10"s
I
I
.
I
!
I
I
100
"1" K/s
10 s
t
101~
Figure 3.15. Dendrite cell size ~-2 vs. cooling rate J" during solidification for A1-4 to 5wt%Cu (open points) and A1-7 to 11 wt%Si (filled points). Full line represents ),2 j"l/3 = 50 # m (K/s) 1/3 (Jones 1984).
10
'"~1
o . Eote0t,0 ~m
x
Chadwick
196
9 Cooksey .tal
01
-
A
Elliott
+ L i v i n g s t o n et al
_
o
Clark&Elliott
0.011 0.001
0.01
r~__
. 1967
1968
13 B u r d e n & Jones 1970 -
9 Zimmermann I r ill
J
1964
Ibaraki&Okamoto M9ore&
~
9
1970 1976
et al 1988 I ~ I,f
0.1
,
,
,
,! 1
V mm/s
10
100
1000
Figure 3.16. Eutectic interphase spacing ,~ vs. V for the regular (lamellar) or a-A1A12Cu eutectic (Jones 1984, Zimmermann et al. 1989).
39
Rapid Solidification
..ii.iiii.i.!.ii!iiii!i!i!!!i!!ii!i!!!ii!;ii!i!!!i!!!i!i!iiii!i!i!i!!ii!i!i!i:!!i i~ii~iiii~i!i!ii!ii~i~ii~i~ii!~i!i!i~!ii~ii~i!i~i~i~i~i~ii~ii~ii~i ;i~!i~i~!ii!ii~i!~i~ii~i~ii~ii i .....................................................................................................................................
~!i!iii!i!iiiiiiiiiiiiiiiiiiii!iiiii!iiiii!iiiiiiiiiiiiiiii iiiiliiiiiiiii!iiiiiiiiiiiii!iiiiiiiiiiiiiiiiiiiiiiiiiii iiiiii!i!iiiiiiiiiiiiiiiii!iiiiiiiiiiiiiiiiiiiiii!i
%!~i (,,.~~ ,;
5:.i% 9,,,,,r
Figure 3.17. Shear band offsets on the outer surface of a bent metallic glass ribbon (Davis 1976).
6 0 0 _ AleeY 2 N i 9 M 1 "r" 4 0 0
~
1~
ol21 Mn
@11 Fe * brittle
200 EL (..9 uJ
60, 50 I
,
I
I
I
I
1400
d3
1200
EL
-
1000
.....
3~ 800
w
2
6
~b
2'o V~
3'o
4'0
(%)
Figure 3.18. Hardness Hv, Young's modulus E, fracture strength af and elongation to fracture Ef of A1-2at%Y-9at%Ni-lat%M (M = Mn or Fe) metallic glass ribbon vs. volume fraction Vf of fcc o~-A1 nanocrystals introduced into the metallic glass matrix (Kim et al. 1990).
40
H. Jones
10%
a)
~i~i I
I
I
"i
I
O
10-5 -
10-6-
eo
10-7 102
1
103
200
b)
I
-
c,, 1 0 0
-
I
10/+ t,ls)
m
E E
' "
105
,
106
9
b
50_.22
,l
'23
,I
,
I
-2z~ -25 P= tog[~s]- O03t+T
1
-26
Figure 3.19. (a) Minimum creep rate ks vs. rupture time tr for Fe-40at%Ni-14at%P6at%B metallic glass for stress range 600-1600 MPa (temperatures: ( 9 523 K, (O) 573 K); (b) applied stress o- vs. parameter P (ks is the strain rate in s -1 , T is the test temperature in K) for the same glass. From (Rama Rao and Radhakrishnan 1984) with data from (Gibeling and Nix 1978).
(3.7 MPa mm 1/2 compared with 2.3 MPa mm 1/2 for the wrought unalloyed condition). The new alloy chemistries afforded by rapid solidification have also led to improved thermal stability, such as in the A1-Fe-V-Si (8009) alloy developed by Allied Signal. Figure 3.21 (Gilman 1990) shows that 8009 retains relatively high strength in the range 150450~ where conventional wrought A1-Cu (2219) and A1-Zn-Mg-Cu (7075) lose most of their high ambient temperature strength as a result of precipitate coarsening. The re-
Rapid Solidification 1200
1000
(3.
8O0
:
O
1
41
i
AI-Ni-Mm-Zr
(~
AI-Ni-Mm
[-]
Fujila and Taba/a
/~
Wyrzykowski and Grabski
1
Hansen
-
t-"
600
-o
400
200
0
20
I
I
I
40
60
80
d 1/2
t
100
,,,
[
120
140
/ mm
Figure 3.20. Proof strength o-0.2 vs. inverse square root of grain size d for rapidly solidified A1-Ni-Mn(-Zr) alloys ((30) compared with results for wrought unalloyed aluminium ([]Am) (Nagahuma et al. 1993).
markable resistance to coarsening of the high volume fraction of the nanoscale silicide dispersoid in 8009 thus raises maximum service temperature by some 200 K. Creep resistance and corrosion resistance are also notably improved relative to established wrought ingot alloys of comparable room temperature strength. The limit on alloying set by incidence of unacceptable ingot macrosegregation in wrought ingot processing has been relaxed by rapid solidification processing for a number of other classes of material such as silicon irons for transformers, tool and bearing steels, nickel based superalloys and copper alloys for conductor applications. It is also a route of choice for processing the new class of high performance Nd-Fe-B permanent magnet materials based on the remarkable properties of the Nd2Fe14B compound.
3.4.2 Applications of rapid solidification It is worth recalling that modem interest in rapid solidification processing stems from curiosity-driven research in the 1960s aimed at establishing its capabilities for generating new alloy phases (Duwez 1967) and size-refined microstructures (Grant 1970). There is no doubt that it has contributed much to our present understanding of what governs formation of alloy phases, their mode of occurrence in solidification microstructures and their effect on properties and performance. It has also demonstrated the development potential of a number of previously underemployed production processes such as melt-atomization and spray forming, containerless processing, melt-spinning and thin-strip casting, and surface melt-traversing, all of which are undergoing intensive appraisal. A number of new
H. Jones
42
Tem~rature (*F) 0 o,. 5OO v
100
200
300
400
500
600
700
800
900
"
70
~,~oo
50
q,)
ta,,,. W
80O9
cO 3OO
,,.'K
c
co
40 N tll
3O
N. 200 2219-T851 ~ A
7075-T651
0
9
20
.
~o
9
0
~: 113 ~
O ~
9O
6O0
8O "~
/rr
soo
~' 4oo w
o~
t~
2219-T851
7075-T651
E lOO
o.,.., ...,.
!
80
r
"-
1
I
i
"
!
'
9
9
"
1
'"I
9
"
!
9
1
I
"'
' I
.
!
' " '"
.
I
"
i
I --
|
9
1_
"
I[
"
0 80
7O 7075-T651
-~ 6 o 9
j
C:~ " c ~0 -
c: ~ I-.
l
/
30 20
60
/
c
V)
J
5O
2219-T851
40 3O 8009
20 -
10
0
----
9 !
70 -
IJJ
CD
2O N E 10
0
so
c
o
~
.~
60
c ID 3O J--.
cll I..- 2OO
-
.~
5O m
8OO9
30O
70
50
10
100 150 200 250 300 350 400 450
Temperature (*C) Figure 3.21. Tensile properties of rapidly solidified 8009 (A1-Fe-V-Si) alloy vs. test temperature compared with high strength wrought alloys 2219-T851 and 7075-T651 (Gilman 1990).
R a p i d Solidification
43
materials have been developed. These include a range of new soft (Smith 1993) and hard (Guthrie 1993) magnetic materials for application in power distribution as well as in magnetic and electronic devices. New high performance light alloys have been established for both wear resistance (Amano et al. 1989) and high temperature (Gilman 1990) performance. Materials with enhanced catalytic performance (Hashimoto 1997) and for fuel cells (Kawashima et al. 1996) have been developed. Powder metallurgy tool steels and superalloys made from prealloyed atomized powder have an established market for inserts, engineering parts and hard facings (Abraham 1991). Direct spray deposition has been employed to produce mill rolls with two or three times the life of conventionally produced rolls (Ikawa et al. 1990). Rapid solidification has been used to develop new alloys for medical implants (Wang et al. 1988) and to make improved Ag-based feedstock particulate for producing dental amalgams (Vero et al. 1992). In addition rapid solidification has provided feedstocks for other developmental technologies such as mechanical alloying, plasma and high velocity spraying, and manufacture of metal matrix composites. It has also provided cost-effective routes to make steel fibres for reinforcement of concrete and castable refractories (Edgington 1977) and aluminium flake for imparting conductivity to plastics (Holbrook 1986) as well as to form ductile brazing foils for a wide variety of precision joining applications (Rabinkin and Liebermann 1993). The newest application, that of a metallic glass golf club head (Ashley 1998) with strength and hardness twice that of cast stainless steel or titanium but of lower modulus and intermediate density, made by bulk undercooling prior to solidification, would have been inconceivable 30 years ago.
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44
H. J o n e s
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R a p i d Solidification
45
Weber, J. K. R., Krishnan, S. and Nordine, R C. (1991) J. Metals, 43(7), 8. Zimmermann, M., Carrard, M. and Kurz, W. (1989) Acta Metall., 32, 3305.
SELECTED BIBLIOGRAPHY
Ananthararman, T. R. (ed.) (1984) Metallic Glasses, Production, Properties and Applications (TransTech, Aedermannsdorf, Switzerland), 300 pp. Anantharaman, T. R. and Suryanarayana, C. (1987) Rapidly Solidified Metals (TransTech, Aedermannsdorf, Switzerland), 260 pp. Cahn, R. W. and Greer, A. L. (1996) in Physical Metallurgy, eds. Cahn, R.W. and Haasen, P., fourth edition (Elsevier Science, Oxford), p. 1723. Gilman, J. J. and Leamy, H. J. (1978) Metallic Glasses (American Society for Metals, Metals Park, OH), 350 PP. G~intherodt, H. J. and Beck, H. (1981) Glassy Metals I (Springer-Verlag, Berlin), 270 pp. Gtintherodt, H. J. and Beck, H. (1983) Glassy Metals II (Springer, Berlin, Germany), 400 pp. Herman, H. (ed.) (1981) Ultrarapid Quenching of Metals and Alloys (Academic Press, New York), 448 pp. Jacobson, L. A. and McKittnick, J. (1994) Mater. Sci. Eng. R, R l l , 355. Jones, H. (1982) Rapid Solidification of Metals and Alloys (Institution of Metallurgists, London), 83 pp. Jones, H. (1984) J. Mater. Sci., 19, 1043. Jones, H. (1991) Metals and Materials, 7, 486. Liebermann, H. H. (ed.) (1993) Rapidly Solidified Alloys: Processes, Structures, Properties, Applications (Marcel Dekker, New York), 808 pp. Luborsky, E E. (ed.) (1983) Amorphous Metallic Alloys (Butterworths, Stoneham, MA), 548 pp. Otooni, M. A. (ed.) (1997) Elements of Rapid Solidification (Springer-Verlag, Berlin), 320 pp. Sahm, P. R., Jones, H. and Adam, C. M. (eds.) (1986) Science and Technology ofthe Undercooled Melt (Martinus Nijhoff, Dordrecht, The Netherlands), 304 pp. Srivatsan, T. S. and Sudarshan, T. S. (1993) Rapid Solidification Technology: An Engineering Guide (Technomatic., Lancaster, PA), 737 pp.
These volumes provide extensive reviews of their subject areas that were comprehensive at the time of their compilation. More specialized review articles also continue to appear but are too numerous to list here. In addition, the series of international conferences that began in Brela, Yugoslavia, in 1970 and reached its ninth at Bratislava in 1996 and the tenth in Bangalore, India in 1999 contain numerous useful review articles within their Proceedings (for the ninth conference see Mater. Sci. Eng. A, 15th June 1997 issue, Vols. A 226-228, 1110 pp). The proceedings of two recent NATO workshops are also useful: Undercooled Metallic Melts: Properties, Solidification and Metastable Phases, eds. D. M Herlach et al., Mater. Sci. Eng. A, 1994, A178(1,2), 320 pp, and Science and Technology of Rapid Solidification and Processing, ed. M. A Otooni, Kluwer, Dordrecht, 1995, 390 pp). International Journal of Rapid Solidification, Vol. 1, No. 1 (1984) to Vol. 9, No. 4 (1996), is the only journal dedicated to publication of research papers in the field. It continues with its new name International Journal of Nonequilibrium Processing with Vol. 10, No. 1 in 1997.
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Chapter 4 Mechanical Alloying 4.1.
4.2. 4.3.
Introduction Nomenclature The Process of Mechanical Alloying 4.3.1 Raw Materials 4.3.2 Process Control Agents 4.3.3 Type of Mills
4.3.3.1 SPEX Shaker Mills 4.3.3.2 The Planetary Ball-Mill 4.3.3.3 Attritor Mills 4.3.3.4 Commercial Mills 4.3.3.5 New Designs 4.4. Mechanism of Alloying 4.5. Consolidation 4.6. Synthesis of Non-Equilibrium Phases 4.6.1 Solid Solubility Extensions 4.6.2 Synthesis of Intermetallics 4.6.3 Disordering of Ordered Intermetallics 4.6.4 Nanocrystalline Materials 4.6.5 Solid-State Amorphization 4.6.6 Displacement Reactions 4.7. Powder Contamination 4.8. Modeling 4.9. Industrial Applications 4.9.1 Nickel-Base Alloys 4.9.2 Iron-Base Alloys 4.9.3 Aluminum-Base Alloys 4.9.4 Magnesium-Base Alloys 4.10. Concluding Remarks References
49 5O 51 51 52 52 53 53 53 53 54 55 56 58 59 62 64 66 67 71 74 75 76 77 79 79 80 80 81
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Chapter 4 Mechanical Alloying C. SURYANARAYANA
4.1. I N T R O D U C T I O N
Mechanical alloying is normally a dry, high-energy ball-milling technique and has been employed to produce a variety of commercially useful and scientifically interesting materials. This technique was developed around 1966 by John S. Benjamin and his colleagues at the Paul D. Merica Research Laboratory of the International Nickel Company (INCO). The main goal was essentially to combine the advantages of precipitation-hardening and oxide-dispersion strengthening in a nickel-based superalloy intended for gas turbine applications. Benjamin (1976, 1990)reviewed the historic origins of the process. This process was referred to as "milling/mixing", but a patent attorney for the INCO coined the term m e c h a n i c a l alloying.
The formation of an amorphous phase by mechanical grinding of an Y-Co intermetallic compound in 1981 (Yermakov et al. 1981) and in the Ni-Nb system by ball-milling of blended elemental powder mixtures in 1983 (Koch et al. 1983) brought about the recognition that mechanical alloying is a potential non-equilibrium processing technique. Beginning from the mid-1980s, a number of investigations have been carried out to synthesize a variety of stable and metastable phases including supersaturated solid solutions, crystalline and quasicrystalline intermediate phases, and amorphous alloys (Suryanarayana 1995a). Efforts were also under way since the early 1990s to understand the process fundamentals of mechanical alloying (Maurice and Courtney 1990). Additionally, it has been recognized that this technique can be used to induce chemical (displacement) reactions in powder mixtures at room temperature or at least at much lower temperatures than normally required (Heinicke 1984, McCormick 1995). Because of all these special attributes, this simple, but effective, processing technique has been applied to metals, ceramics, polymers, and composite materials. The important attributes of mechanical alloying are listed in Table 4.1. The different facets of this technology have been reviewed periodically (Gilman and Benjamin 1983, Weeber and Bakker 1988, Koch 1989, Koch 1991, Herr 1995, Suryanarayana 1996, Lu and Lai 1998, Murty and Ranganathan 1998). The literature available till 1994 has been collected together in an annotated Bibliography (Suryanarayana 1995a). Additionally, a number of conferences are devoted to the science and technology of mechanically alloyed materials (Arzt and Schultz 1989, Yavari and Desr6 1990, Froes and de Barbadillo 1990, de Barbadillo et al. 1993). Furthermore, proceedings of the annual International Symposia on 'Metastable, Mechanically 49
50
C. Suryanarayana
Table 4.1. Attributes of mechanical alloying Production of a fine dispersion of second-phase particles Extension of solid solubility limits Refinement of grain sizes down to the nanometer range Synthesis of novel crystalline and quasicrystalline phases Development of amorphous (glassy)phases Disordering of ordered intermetallics Possibility of alloying of difficult to alloy elements Inducement of chemical (displacement) reactions at low temperatures Scaleable process
Alloyed and Nanocrystalline Materials' (ISMANAM) are regularly published in Materials Science Forum published by TransTech Publications, Switzerland (Shingu 1992, Yavari 1995, Schulz 1996, Fiorani and Magini 1997).
4.2. N O M E N C L A T U R E
Two different terms are commonly used in the literature to denote the processing of powder particles in high-energy ball-mills. Mechanical Alloying (MA) describes the process when mixtures of powders (of different metals or alloys/compounds) are milled together. Material transfer is involved in this process to obtain a homogeneous alloy. On the other hand, milling of uniform (often stoichiometric) composition powders, such as pure metals, intermetallics, or prealloyed powders, where material transfer is not required for homogenization, has been termed Mechanical Milling (MM). Some investigators have referred to MM as Mechanical Grinding (MG). Since "grinding" is normally thought of as an abrasive machining process that involves mainly shear stresses and chip formation, the term "milling" is preferred to include the more complex triaxial, perhaps partly hydrostatic, stress states that can occur during ball-milling of powders (Koch 1991). It should be realized that MA is a generic term, and some investigators use this term to include mechanical milling/grinding as well. However, we will distinguish between these two terms by using MA or MM, depending on whether material transfer is involved or not during processing. Reaction-milling (RM) is the process developed by Jangg et al. (1975). In this process the powder is milled without the aid of any process control agent (see later for its function during milling) to produce fine dispersions of oxides and carbides in aluminum (Jangg 1989). This proprietary process of Eckart-Werke in Germany achieves the dispersion of carbides by adding lamp-black or graphite during milling of aluminum. Adjusting the oxygen content via close control of the milling atmosphere (oxygen, argon, nitrogen, air, etc.) produces the oxides. Thus, the final product of milling contains the aluminum phase along with A14C3 and A1203 and these alloys are given the trade name DISPAL. Even without intentional additions of graphite, fine dispersions of A14C3 and A1203 could be produced when a process control agent is added during milling of metal powders. Thus, Singer et al. (1980) observed the presence of 10-nm size dispersoids of A14C3 and A1203
Mechanical Alloying
51
in a hot-extruded sample after milling the aluminum powder in the presence of Nopcowax 22-DSP. Milling of metal powders in the presence of reactive solids/liquids/gases (enabling a chemical reaction to take place) is now regularly employed to synthesize metal oxides, nitrides, and carbides (Avvakumov 1986, Calka et al. 1996). Thus, milling of titanium in a nitrogen atmosphere has produced titanium nitride (Calka 1991, Suryanarayana 1995b); several other compounds have also been produced in a similar way. Milling of tungsten with carbon (graphite) has produced tungsten carbide (Radlinski and Calka 1991). Milling of metal powders with boron has produced borides, e.g., TiB2 (Calka and Radlinski 1990). Another variation of milling adopted is cryomilling (Luton et al. 1989), in which the milling operation is carried out in the presence of liquid nitrogen. This process produces 2-10 nm sized aluminum oxy-nitride particles which strengthen the aluminum powder. 4.3. T H E P R O C E S S OF M E C H A N I C A L
ALLOYING
The actual process of MA starts with mixing of the powders in the right proportion and loading the powder into the mill along with the grinding medium (generally steel balls). This mix is then milled for the desired length of time until a steady state is reached when the composition of every powder particle is the same as the proportion of the elements in the starting powder mix. Sometimes the powder is milled to an intermediate state either to form metastable phases or to achieve certain desired properties. (During MM, however, the objective would be to produce certain desired properties in the material, e.g., formation of nanostructured or amorphous phases, or disordering of ordered intermetallics since no alloying/homogenization is required to occur). The milled powder is then consolidated into a bulk shape and heat treated to obtain the desired microstructure and properties. Let us now look at the individual components mentioned above. 4.3.1 Raw materials
The raw materials used for MA are widely available commercially pure powders that have particle sizes in the range of 1-200/zm. But the powder particle size is not very critical, except that it should be smaller than the grinding ball size. This is because the powder particle size decreases exponentially with time and reaches a small value only after a few minutes of milling. These powders fall into the broad categories of pure metals, master alloys, prealloyed powders and refractory compounds. The oxygen content of the commercially pure metal powders ranges from 0.05-0.2 wt.%. Therefore, it becomes important to choose high-purity powders, especially if one is interested in studying phase formation, since more often than not, the nature and amount of impurities in the system decide the nature of the final phase formed and the chemical constitution of the alloy. Dispersion strengthened materials usually contain additions of carbides, nitrides, and oxides. Oxides are the most common and these alloys are known as oxide-dispersion strengthened (ODS) materials. In the early days the powder charge for MA consisted of at least 15 vol% of a ductile compressibly deformable metal powder to act as a host or a binder. However, in recent years, mixtures of fully brittle materials have been milled successfully, resulting in alloy formation (Koch 1991). Thus, the requirement of having a ductile metal
52
C. Suryanarayana
powder during milling is no longer necessary. Accordingly, ductile-ductile, ductile-brittle, and brittle-brittle powder mixtures are milled now a days to produce novel alloys. Mixtures of solid powder particles and liquids have also been milled in recent times (Ivanov 1992).
4.3.2 Process control agents
A process control agent (PCA) is added to the powder mixture during milling, especially when the powder mix involves a substantial fraction of a ductile component. The PCAs are mostly organic compounds, which act as surface-active agents. The PCA adsorbs on the surface of the powder particles and minimizes cold welding between powder particles and thereby inhibits agglomeration. The surface-active agents adsorbed on particle surfaces interfere with cold welding and lower the surface tension of the solid material. Because the energy required for milling is equal to the product of new surface area generated times the surface tension, a reduction in surface tension results in the use of shorter milling times and/or finer powders. A wide range of PCAs has been used in practice. These are mostly organic compounds used at a level of about 1-4 wt.% of the total powder charge and include stearic acid, hexane, oxalic acid, methanol, ethanol, acetone, heptane, Nopcowax-22DSP, octane, toluene, etc. (Suryanarayana 1998). Majority of these compounds decompose during milling, interact with the powder and form compounds, and these get incorporated in the form of dispersoids inside the powder particles during milling. Thus hydrocarbons containing hydrogen and carbon, and carbohydrates containing hydrogen, carbon, and oxygen are likely to introduce carbon and/or oxygen into the powder particles, resulting in the formation of carbides and oxides which are uniformly dispersed in the matrix. These are not necessarily harmful to the alloy system since they can contribute to dispersion strengthening of the material. The hydrogen subsequently escapes as a gas or is absorbed into the metal lattice on heating or sintering. The presence of air in the milling container or milling of the powders at very low temperatures also has been shown to minimize welding, most probably due to the increased brittleness of the powder particles at such low temperatures. Metal powders (with an fcc structure) milled in a hydrogen atmosphere have been found to become brittle and not stick to themselves or the container. The choice of a PCA for milling depends on the nature of the powder being milled and the purity of the final product desired. The nature and amount of PCA used during milling determine the final powder particle size and powder yield. It should also be realized that there is no universal PCA. One has to decide on a PCA by looking at the possible interactions between the metal and the components in the PCA.
4.3.3 Type of mills Different types of milling equipment are used to produce mechanically alloyed powders. They differ in their capacity, efficiency of milling and additional arrangements for cooling, heating, etc.
Mechanical Alloying
53
4.3.3.1 SPEX shaker mills Shaker mills such as SPEX mills, which mill about 10 g of the powder at a time, are most commonly used to investigate the fundamental aspects of MA and for alloy screening purposes. In this mill, the vial containing the powder to be milled and the grinding balls, is secured in the clamp and swung energetically back and forth several thousand times a minute. The back-and-forth shaking motion is combined with lateral movements of the ends of the vial, so that the vial appears to be describing a figure-8 or infinity sign as it moves. With each swing of the vial the balls impact against the sample and the end of the vial, both milling and mixing the sample. Because of the amplitude (about 5 cm) and speed (about 1,200 rpm) of the clamp motion, the ball velocities are high (on the order of 5 m/s) and consequently the force of the bali's impact is unusually great. Therefore, these mills can be considered as high-energy variety. 4.3.3.2 The planetary ball-mill Another popular mill for conducting MA/MM experiments is the planetary ball-mill (referred to as Pulverisette) in which a few hundred grams of the powder can be milled at a time. The planetary ball-mill owes its name to the planet-like movement of its vials. These are arranged on a rotating support disk and a special drive mechanism causes them to rotate around their own axes. The centrifugal force produced by the vials rotating around their own axes and that produced by the rotating support disk both act on the vial contents, consisting of material to be ground and the grinding balls. Since the vials and the supporting disk rotate in opposite directions, the centrifugal forces alternately act in like and opposite directions. This causes the grinding balls to run down the inside wall of the vialmthe friction effect--followed by the material being ground and grinding balls lifting off and traveling freely through the inner chamber of the vial and colliding against the opposing inside wall--the impact effect. Even though the linear velocity of the balls in this type of mill is higher than that in the SPEX mills, the frequency of impacts is much more in the SPEX mills. Hence, in comparison to SPEX mills, Fritsch Pulverisette can be considered lower energy mills. 4.3.3.3 Attritor mills Attritors are the mills in which large quantities of powder (from a few pounds to 100 lb) can be milled at a time. The velocity of the grinding medium is much lower (about 0.5 m/s) than in Fritsch or SPEX mills and consequently the energy of the attritors is low. Attritors of different sizes and capacities are available. The operation of an attritor is simple. The powder to be milled is placed in a stationary tank with the grinding media. This mixture is then agitated by a shaft with arms, rotating at a high speed of about 250 rpm (Fig. 4.1). This causes the media to exert both shearing and impact forces on the material. Although much less efficient than either the SPEX or Fritsch mills, the laboratory attritor works up to 10 times faster than conventional jar mills. 4.3.3.4 Commercial mills Commercial mills for MA are much larger in size than the mills described above and can process several hundred pounds of powder material at a time. Mechanical alloying for
C. Suryanarayana
54
Water-cooled stationary tank
Gas seal
Steel ball bearings Ball mill Figure ,1.1.
Rotating impeller
Arrangement of rotating arms on a shaft in the attrition ball-mill.
commercial production is carried out in ball-mills of up to 3,000 lb capacity. The milling time decreases with an increase in the energy of the mill. For example, a process that takes only a few minutes in the SPEX mill may take hours in an attritor and a few days in a commercial mill.
4.3.3.5 New designs Several new designs of mills have been developed in recent years for specialized purposes. These include the rod mills, vibrating frame mills, and the equipment available from Dymatron, Inc., Cincinnati, OH; Super Misuni NEV-MA-8 from Nisshin Giken, Co., Ltd., Tokyo, Japan (with the ability to control the temperature of milling from very low temperatures by spraying liquid nitrogen and up to a high temperature of 300~ by electrical heating); Uni-Ball-Mill from Australian Scientific Instruments, Canberra, Australia (it is possible to control the nature and magnitude of impact of the balls in this machine with the help of magnets); HE Mill from M.B.N. srl, Rome, Italy; and Zoz Maschinenbau GmbH,
Mechanical Alloying
55
Kreuztal, Germany. Some special equipment is also designed for specific laboratory applications. The operating details and special features of these mills can be found in the respective brochures available from the manufacturers. 4.4. MECHANISM OF ALLOYING During high-energy milling the powder particles are repeatedly flattened, fractured and rewelded. Whenever two steel balls collide, some amount of powder is trapped in between them. Typically, around 1,000 particles with an aggregate weight of about 0.2 mg are trapped during each collision (Fig. 4.2). The force of the impact plastically deforms the powder particles, creates new surfaces, and enables the particles to weld together. This leads to an increase in particle size. Since in the early stages of milling, the particles are soft (if we are using either ductile-ductile or ductile-brittle material combination), their tendency to weld together and form large particles is high. A broad range of particle sizes develops, with some as large as three times bigger than the starting particles. The composite particles at this stage have a characteristic layered structure consisting of various combinations of the starting constituents. With continued deformation, the particles get work hardened and fracture by a fatigue failure mechanism and/or by the fragmentation of fragile flakes. Fragments generated by this mechanism may continue to reduce in size in the absence of strong agglomerating forces. At this stage, the tendency to fracture predominates over cold welding. Owing to the continued impact of grinding balls, the structure of the particles is steadily refined, but the particle size continues to be the same. Consequently, the interlayer spacing decreases and the number of layers in a particle increases. As mentioned earlier, the rate of refinement of the internal structure (particle size, crystallite size, lamellar spacing, etc.) is roughly logarithmic with processing time and therefore the size of the stating particles is relatively unimportant. In a few minutes to an hour, the lamellar spacing usually becomes small and the crystallite (or grain) size is refined to nanometer (1 n m - 10 -9 m or 10/k) dimensions (Fig. 4.3). The ease with which this can be achieved is one reason why MA has been extensively employed to produce nanocrystalline materials (Koch 1993, Suryanarayana 1995c, Koch 1997). After milling for a certain length of time, steady-state equilibrium is attained when a balance is achieved between the rate of welding, which tends to increase the average particle size, and the rate of fracturing, which tends to decrease the average composite particle size. At this stage each particle contains substantially all of the starting ingredients, in the proportion they were mixed together and the particles reach saturation hardness due to the accumulation of strain energy. The particle size distribution at this stage is narrow, because particles larger than average are reduced in size at the same rate that fragments smaller than average grow through agglomeration of smaller particles. The average particle size obtained at this stage depends on the relative ease with which agglomerates can be formed by cold welding, fatigue and fracture strength of composite particles, and resistance of particles to deformation. From the foregoing it is clear that during MA, heavy deformation is introduced into the particles. This is manifested by the presence of a variety of crystal defects such as dislocations, vacancies, stacking faults, and increased number of grain boundaries. The presence of this defect structure increases the diffusivity of solute elements into the matrix. Further-
C. Suryanarayana
56
Figure 4.2. Ball-powder-ball collision of powder mixture during MA (Gilman and Benjamin 1983).
more, the refined microstructural features decrease the diffusion distances. Additionally, the slight rise in temperature during milling further aids the diffusion behavior, and consequently, true alloying takes place amongst the constituent elements. While this alloying generally takes place nominally at room temperature, sometimes it may be necessary to heat treat the mechanically alloyed powder for alloying to be achieved. This is particularly true when formation of intermetallics is desired.
4.5. CONSOLIDATION
Widespread application of mechanically alloyed powders requires production of the powder in tonnage quantities and also efficient methods of consolidating the powders into bulk shapes. All the consolidation methods generally used in powder metallurgy processes can also be used for mechanically alloyed powders. However, since the powder particles in the mechanically alloyed condition are stronger and smaller (typically a few microns, even though the grain size may be only a few nanometers) in size than those used in conventional powder metallurgy operations, some special precautions need to be taken to minimize their activity and high level of interparticle friction. Conventional consolidation of powder to full density through processes such as hot extrusion and hot isostatic pressing normally requires use of high pressures and high tem-
Mechanical Alloying
57
E r N ......, t,,..
~
N
r ,,.=.=, t~
Decreasing ball-to-powder ratio
t.,.=
13.
Milling time, h Figure 4.3. Schematic showing refinement of particle and grain sizes with milling time. The rate of refinement increases and the minimum particle (grain) size reached is lower at higher ball-to-powder weight ratios.
peratures for extended periods of time. Unfortunately, this results in loss of the benefits achieved due to the metastable effects or nanostructures obtained by MA. Therefore, novel and innovative methods of consolidating the mechanically alloyed powders are required. (See also Chapter 13 on "Powder Consolidation" by Joanna R. Groza in this book.) Mechanically alloyed powders have been consolidated into useful bulk shapes by several processes. Since the mechanically alloyed powders have a high hardness, cold compaction is not an option. Furthermore, ODS alloys have been found not to densify during simple sintering. The most common method of consolidation is hot compaction followed by hot extrusion, or by direct hot extrusion at temperatures greater than half the melting point. This process could be used if one is not concerned with the loss of metastable effects, such as the crystallization of the amorphous phase in the powder. Some of the consolidation methods used for mechanically alloyed powders, both on laboratory and industrial scales, include hot isostatic pressing, powder rolling, Ceracon processing, plasma activated sintering, electrodischarge compaction, and explosive forming (shock or dynamic consolidation methods). Combination of these techniques is also sometimes used (He and Ma 1996). Some investigations were also carried out to compare the efficiencies of the different consolidation techniques in retaining the material in a nanostructured condition (Suryanarayana et al. 1997). Retention of nanostructures seems to be most efficient when consolidation is carried out at very high pressures and when the powder is exposed to these high pressures (and temperatures) for only a short period of time, e.g.,
58
C. Suryanarayana
Table 4.2 Important milestones in the development of mechanical alloying/milling 1966 1981 1983 1988 1989 1990
Development of ODS nickel-base superalloys Amorphization of intermetallics Amorphization of blended elemental powder mixtures Synthesis of nanostructures Occurrence of displacement reactions Observation of disordering of intermetallics
in shock consolidation methods. These operations are usually carried out at temperatures considerably lower than those used for conventional powders. For example, while conventional y-TiA1 powders are hot isostatically pressed at about 1,100~ (Beddoes et al. 1992), mechanically alloyed y-TiA1 powders could be consolidated to full density by hot isostatically pressing them at about 725~ (Suryanarayana et al. 1997).
4.6. S Y N T H E S I S OF N O N - E Q U I L I B R I U M P H A S E S
MA of suitable alloy compositions can result in the formation of a variety of equilibrium and non-equilibrium alloy phases. These include equilibrium and supersaturated solid solutions, stable and metastable crystalline and quasicrystalline intermediate phases, and glassy (amorphous) alloys (Koch 1991, Suryanarayana 1996). This technique is also capable of producing true alloys starting from pure elements that are not either easy to form by conventional means or sometimes even impossible to prepare, e.g., elements which are immiscible under equilibrium conditions. Rapid solidification processing (RSP) (rapid quenching of metallic melts at rates of ~106 K/s) also produces very similar microstructural and constitutional effects as MA (see Chapter 3 on "Rapid Solidification" by H. Jones in this book) and, therefore, comparisons are frequently made of the product phases obtained by these two techniques. However, MA processing is carried out entirely in the solid state. Thus, limitations imposed by phase diagrams, such as immiscibility in the liquid and solid states, do not apply to the MA process and therefore the alloying capabilities are likely to be much more extensive than by conventional ingot techniques or RSP methods. Furthermore, it should also be noted that whereas the technique of rapid solidification started as an academic curiosity in 1960 (Duwez et al. 1960, Duwez 1967) and progressed into an industrially mature technology about 20 years after the demonstration of its capabilities (Anantharaman and Suryanarayana 1987, Liebermann 1993), MA started as an industrial necessity in the late 1960s and only recently the scientific basis of MA is beginning to be understood. Some of the important milestones in the development of MMMM are listed in Table 4.2. Recent advances in understanding the "science" behind MMMM processing will be discussed first under the following headings: (1) Solid Solubility Extensions; (2) Synthesis of Intermetallics;
Mechanical Alloying (3) (4) (5) (6)
59
Disordering of Ordered Intermetallics; Nanostructures; Solid-State Amorphization; and Displacement Reactions.
The industrial applications of mechanically alloyed materials, such as oxidedispersion strengthened (ODS) alloys and others, will then be discussed towards the end of the Chapter.
4.6.1 Solid solubility extensions Solid solubility extensions have been achieved in many alloy systems by non-equilibrium processing methods such as RSP and vapor deposition. Similar results have also been reported in metal powder mixtures processed by MA. However, variation of solid solubility limits with the process parameters has not been systematically investigated. On mechanically alloying the blended elemental powder mixtures, interdiffusion between the two components occurs and solid solutions can form. This solid solubility is expected to increase with milling time as diffusion progresses and reach a saturation level, beyond which no further extension of solubility occurs. These can be considered as the solid solubility levels achieved by MA. Solid solubility levels have been generally determined from the changes in the lattice parameter values calculated from the shifts in peak positions in the X-ray diffraction patterns. There are certain difficulties associated with the solid solubility limits determined this way and, therefore, they may not be accurate for reasons discussed in an earlier article (Suryanarayana 1996). The absence of a second phase in the X-ray diffraction pattern is usually taken as an indication of formation of a homogeneous solid solution. But this may not always be true, since the appearance of the second phase in the X-ray diffraction pattern depends on the amount of the second phase, its particle size, the strain and defect content present in the material (this may shift the peak positions, sometimes asymmetrically). Therefore, caution must be exercised in interpreting the X-ray diffraction patterns from mechanically alloyed materials. Even though many of these factors have not been taken into account in many of the cases, solid solubility extensions, beyond the equilibrium values, have been reported in a number of alloy systems and some of these values are listed in Table 4.3. It should also be noted that in many cases equilibrium solid solutions (e.g., Cu-Ni, A1-Mg, Cu-A1) were synthesized by MA starting from blended elemental powders. Solid solutions (both equilibrium and metastable) form during MA because diffusion distances between particles are considerably reduced. The increased defect density and a slight rise in temperature during MA further aid diffusion. In fact, the temperature during MA seems to have a predominant effect. Increasing the temperature during MA tends to encourage decomposition of the solid solutions formed to reach thermodynamic equilibrium and oppose the effect of alloying that is possible by MA. The balance between these two opposing forces decides the final constitution of the alloy. For example, MA of a Cu-50 at. % Ag powder mixture at room temperature resulted in the formation of a homogeneous fcc solid solution. But MA at 473 K resulted in a fully decomposed two-phase
C. S u r y a n a r a y a n a
60
Table 4.3 Selected list of solid solubility extensions obtained by mechanical alloying Solvent
Solute
Maximum room temperature solid solubility (at.%) (equilibrium
by MA
(Massalski et al. 1990)) Ag A1
Cu
~0.0
100.0
Gd
~0.0
5.0
Cr
~0.0
Fe
Cd Co
Cu
Fe
0.025
5.0 4.5
Mg
1.2
23.0
Mn
0.3
18.5
Nb
--~0.0
15.0
Ru
~0.0
14.0
Ti
~0.0
Zn
3.5
3.0 ~50 6.0
C
~0.0
Cr
4.0
40
Cu
~0.0
100
Ag
~0.0
100
Hg
~0.0
70
Sb
1.2
3.7
Sn
~0.0
9.8
Ti
~0.0
AI
20.8
50
9.4
Cu
0.3
40
Mn
Co
4.0
50
Nb
AI
6.0
60
Ge
3.4
Ni
1.2
10
AI
7.2
27
C
"~0.0
12
Cr
22.5
50
Ga
10.8
50
In
~0.0
13
Nb
3.2
10
Ti
9.7
28
Ni
8.6
Ru
A1
0.7
14
Ti
A1
11.2
60
Cu
~0.0
10
Mg
~0.0
5O
Si
~0.0
8
W
Fe
2.0
40
Zr
A1
1.0
17.5
Ni
~0.0
7.0
Mechanical Alloying
61
mixture. Partially decomposed structures were obtained during milling at intermediate temperatures (Klassen et al. 1997). A similar conclusion was drawn by Ma et al. (1997) to explain the differences in the alloying behavior between two liquid immiscible systems Ag-Fe and Cu-Fe. Another interesting observation that has been made in mechanically alloyed materials is that the solid solubility achieved increases with increasing solute content in the starting blended elemental powder mixture, making it difficult to report a specific solid solubility value (Clark et al. 1995). This has been tentatively explained on the basis that the concentration gradients between the solvent and the solute are steeper at higher solute contents and that these are expected to result in increased diffusion and consequently in higher solubility levels. Neither the actual mechanism(s) for formation nor the limits for the solid solubility extensions in alloy systems obtained by MA have been well investigated. It was suggested (Schwarz et al. 1985) that the increased solid solubility of Ti in Ni in mechanically alloyed Ni-Ti powder mixtures was due to the metastable equilibrium between the ot-Ni solid solution and the Ni-Ti amorphous phase as opposed to the stable equilibrium between the ot-Ni and Ni3Ti phases. That is, the extent of supersaturation is limited to the composition where the amorphous phase starts forming. A similar conclusion was also arrived at by Murty et al. (1993). However, it has been noted in recent years that significant solid solubility extensions can be achieved in alloy systems (e.g., Ti-Mg) even when an amorphous phase did not form in them (Zhou et al. 1995). Thus, alternative explanations need to be sought to explain the formation of solid solutions and their limits of solid solubility. For example, it was suggested (Suryanarayana and Froes 1990) that formation of supersaturated solid solutions is related to the occurrence of nanocrystalline phases. The large volume fraction of grain boundaries in these materials is expected to enhance the diffusion and consequently extend the solid solubility levels in these types of systems. In support of this argument, it has been shown that the solid solubility of Mg in Ti is zero when the grain size of the titanium phase is on the order of microns, while significant solubility is reported when the titanium grain size is on the order of nanometers. Formation of (supersaturated) solid solutions by MA is more noteworthy in liquid immiscible systems. Accordingly, it has been reported that extensive solid solutions form in Cu-Co, Cu-Fe, Cu-Ta, and Cu-V systems. Gente et al. (1993) and Huang et al. (1994, 1997) suggested that the increased diffusivity due to the presence of structural defects and local stresses in nanocrystals is responsible for the formation of supersaturated solid solutions in these systems. Formation of homogeneous solid solutions is favored when the crystallite size of the constituents is reduced to below about 1-2 nm. The absolute limits of solid solubility extension in RSP alloys have been determined from thermodynamic considerations using the concept of To temperature at which, for a given composition, the solid and liquid phases have the same free energy. Supersaturated solid solutions can be obtained only when the liquid could be undercooled to a temperature below To (Boettinger and Perepezko 1993). Even though MA also is a nonequilibrium processing technique, there have not been any attempts so far to rationalize even the limited experimental data of solid solubility extensions; furthermore, the possible limits have not been defined theoretically. However, since no liquid phase is involved
62
C. S u r y a n a r a y a n a
during MA, it is doubtful if the same criterion as in RSP can be used to rationalize the solid solubility extensions achieved during MA. Even though the mechanisms by which supersaturated solid solutions form by MA and RSP techniques are different, comparisons have been frequently made, with no specific conclusions. For example, solid solutions have been formed in the whole composition range in the Cu-Ag system by both the techniques. A similar situation was obtained in several other systems by rapid solidification methods, but not by MA. Solid solutions have been obtained in the full composition range in the Cu-Fe, A1Sb-InSb, and Cu-Co systems by MA but not by RSP. The levels of solid solubility achieved are also different in different systems by the two different techniques (Suryanarayana 1999).
4. 6.2 Synthesis of intermetallics Because of the intimate mixing of the constituent metal powders during MA, it has been possible to synthesize a variety of intermetallic phases at appropriate compositions in a number of alloy systems. The intermetallics synthesized include both equilibrium and metastable crystalline phases and non-equilibrium quasicrystalline phases. Only a few reports exist on the synthesis and transformation behavior of quasicrystalline phases. Icosahedral phases have been synthesized by MA in Mg-Zn-A1 (Ivanov et al. 1989a), Mg-Cu-A1, A1-Cu-Mn (Eckert et al. 1989, 1990a), A1-Cu-MnGe (Bahadur et al. 1992), A1-Cu-Fe (Nasu et al. 1992, Asahi et al. 1994), A1-Cu-Ru (Politis et al. 1989), Ti-Ni-Fe-Si, and Mg-A1-Pd systems (Takeuchi et al. 1993, Suryanarayana 1996). These icosahedral phases have been found to be similar to those produced by RSP. It has been reported (Eckert et al. 1990b) that the nature of the phase formed was different depending on the milling intensity. At a very high milling intensity, a crystalline intermetallic phase formed while at a very low intensity, an amorphous phase formed. A quasicrystalline phase formed at intermediate intensities. The formation of the crystalline intermediate phase at the highest milling intensity was attributed to the increased temperature rise during MA resulting in the crystallization of the quasicrystalline phase. In support of this, it may be noted (Bahadur et al. 1996) that an as-milled Ti-Ni-Fe-Si amorphous alloy transformed to an icosahedral quasicrystalline phase on annealing at a temperature of 1025 K. This trend is somewhat contrary to what has been observed in RSP alloys. In RSP, a stable crystalline phase formed when the melt was cooled at low cooling rates and an amorphous phase formed at the highest cooling rates. A quasicrystalline phase formed at intermediate cooling rates (Suryanarayana and Jones 1988). A similar observation of formation of metastable phases at low milling intensities and stable phases at higher milling intensities has been reported in some other alloy systems also (Gerasimov et al., 1991, Guo et al. 1994, Suryanarayana et al. 1999). The equilibrium Hume-Rothery electron compounds--fl-, y- and E-brass phasesm were synthesized by MA of pure Cu and Zn powders mixed in the proper proportions (McDermott and Koch 1986, Martelli et al. 1988, Koch 1991). A number of intermetallics in other alloy systems have also been synthesized (Koch 1991, Suryanarayana 1995a). These include aluminides, silicides, and other intermetallics. Ceramic materials such as borides, carbides, nitrides, and oxides have also been synthesized by reaction milling. In
Mechanical Alloying
63
some cases, formation of the intermetallics seems to have taken place by a combustion synthesis-type reaction. Amongst these cases, there are also examples where the combustion took place only after "interrupted milling", i.e., milling of the powders for a given length of time, aging the powder at room temperature after stopping the milling, and then resuming the milling operation. An NiA1 compound was synthesized this way (Atzmon 1990). Similar results were also reported in other systems (Mukhopadhyay et al. 1996) and in fact these results compare very favorably with those obtained in the regular combustion synthesis process (Moore and Feng 1995). The synthesis of only a few metastable intermetallic phases has been reported by MA. A metastable deformation-induced martensite phase has been reported to form in mechanically alloyed Cu-Zn alloy powders (McDermott and Koch 1986). Similarly, synthesis of a bct-Fe phase on prolonged processing of iron powder in a nitrogen atmosphere has also been reported (Rawers et al. 1995). It is surprising that not many such deformationinduced metastable phases have been reported in mechanically alloyed powder mixtures in view of the fact that there are several instances where metastable phases are formed by cold working; a good example being the formation of deformation-induced metastable E (hcp) martensite in austenitic stainless steels. Recently, a metastable crystalline phase with rhombohedral crystal structure was reported to form in mechanically alloyed A1-Ge powders (Yvon and Schwarz 1993, Chattopadhyay et al. 1996). Formation of a metastable phase with a simple cubic structure was also reported in mechanically alloyed Te-Ag alloys (Chitralekha et al. 1995). Extreme caution should be exercised in identifying metastable crystalline phases produced by MA. It is possible that some of the phases formed are due to interstitial contamination, and these may be misinterpreted as new intermetallic phases. An accurate chemical analysis of the final resultant powder can, in many cases, clear up this confusion (Suryanarayana 1995b). Both disordered and ordered intermetallics (some of them metastable) have been synthesized by MA. In some cases, the intermetallics were synthesized directly by MA; but, in others an additional heat treatment was required after MA to form the intermetallic. For example, ordered intermetallics have been found to form directly on MA in Al-rich Al-transition metal systems. Some of the examples include A13Nb (Suryanarayana et al. 1994a), A15Fe2 (Mukhopadhyay et al. 1994), A13Zr (Suryanarayana et al. 1994b, Li et al. 1995a, Desch et al. 1991, 1996), (A1,X)3Zr, where X = Fe or Ni (Suryanarayana et al. 1994b, Li et al. 1995a), A13Hf, and (A1,X)3Hf where X = Fe or Ni (Li et al. 1995b). Reasons for the formation of ordered intermetallics have not been investigated in detail. It may be assumed that a phase will exist either in the ordered or disordered condition depending upon the balance between atomic disordering introduced by MA and the thermally activated reordering. The reordering is caused by the difference in energy between the ordered and disordered states. Thus, if this difference in energy is small, the alloy will exist in the disordered state whereas if it is large the alloy will be in the ordered state. It has been shown that both MA and MM produced the NiA1 phase in the ordered state while the FeA1 phase was found to be in the disordered state (Schropf et al. 1994). The ordering energy is related and scales up with the enthalpy of formation ( A H f ) and A H f values
64
C. Suryanarayana
for NiA1 and FeA1 are 72 and 25 kJ/mol, respectively (Schropf et al. 1994), confirming the above argument. A number of intermediate phases have been synthesized in the disordered condition. This is not surprising because MA involves heavy deformation and this destroys ordering in the lattice (Stoloff and Davies 1966). The number of disordered intermetallic phases synthesized is too numerous to list here and can be found in (Koch 1991, Suryanarayana 1995a, Lu and Lai 1998, Murty and Ranganathan 1998).
4.6.3 Disordering o f ordered intermetallics It has been long known that partially ordered phases are stronger than those wholly disordered or fully ordered (because at a certain value of the long-range order parameter, S, superdislocations separate into unlinked singles). Thus, it is of interest to study the mechanical behavior of materials in various states of partial order. Disordering phenomena of ordered alloys have also been studied to understand the mechanism of disordering and also to produce the disordered material that has a higher deformability than the ordered alloy. Alloys can be disordered by irradiation (Russel 1985), rapid solidification (Koch 1988), or heavy plastic deformation (Stoloff and Davies 1966). The present section will deal with the phenomenon of disordering of ordered intermetallics by MM. An exhaustive review on this aspect has been recently published (Bakker et al. 1995a). Mechanical milling introduces high energy into the material being processed. This energy can be stored in the material as atomic disorder and/or grain boundaries, i.e., A G (milling) = A G (disorder) + A G (grain boundaries). The atomic disorder in an intermetallic can be manifested in three different ways (Bakker et al. 1995a). First, the two atomic species involved can occupy the "wrong" sublattices and this is referred to as anti-site disorder. This type of disorder was observed in a number of mechanically milled compounds with the A15 structure, e.g., Ni3A1, Ni3Si, Fe3Ge and those with the B2 structure, e.g., CoGa. Second, triple defects can be generated. In CoGa, for example, vacancies on the Co-sublattice in combination with Co anti-site atoms in a ratio of 2:1 constitute the triple defects. Third, there could also be redistribution of interstitials wherein the interstitial atoms in the octahedral sites are transferred to the tetrahedral sites, e.g., Mn3Sn2, Fe3Ge2. Additionally, grain refinement increases the grain boundary area and this also raises the free energy of the system. The sum of the energy of these two effects (disordering and creation of grain boundaries) will be the total energy introduced into the material during milling. Mechanical milling of ordered intermetallics has been shown to result in one of the three following types of transformation (Bakker et al. 1995a, 1995b): (1) formation of a solid solution of one component in the other, i.e., the terminal solid solution based on the major component, e.g., Nb3A1, V3Ga, Ni3A1, Fe3Ge, Ni2V, and NbAu2; (2) formation of an amorphous phase, e.g., Nb3Sn, NiZr, NiV2, CoZr; and (3) formation of a different phase with a complex crystal structure, e.g., Ni3Sn2, TiSi2.
65
Mechanical Alloying CoZr
1.0
tv W IW
0.5CuTi
d c/ i3AI 0.0
M E C H A N I C A L A L L O Y I N G TIME
Figure 4.4. Schematic representation of variation of long-range order parameter S with milling time.
It has been noted that upon milling, the long-range order parameter (S) in the intermetallic is gradually reduced and, in many cases, the material may become totally disordered (S = 0), e.g., Ni3A1 (Jang and Koch 1990). In other cases, S is reduced with milling time but does not reach S = 0, i.e., partial order and partial disorder co-exist, e.g., CuTi (Seki and Johnson 1990). In other cases, the S value does not decrease at all and is maintained at S = 1; but, with continued milling the material becomes amorphous, e.g., CoZr (Cho and Koch 1993). These three situations are schematically represented in Fig. 4.4. Thus, upon milling, an ordered intermetallic can transform into a disordered crystalline phase (solid solution) or an amorphous phase either with or without complete loss of long-range order. If the product is a crystalline phase, the material has an extremely fine grain size, usually in the nanometer range. Disordering of intermetallics has been studied primarily by X-ray diffraction techniques. By measuring the intensity of the superlattice reflections relative to that of the fundamental reflections, the long-range order parameter S is evaluated. Additionally, it has also been noted that the lattice parameter increases slightly (0.3-0.8%) in the disordered state (Gialanella et al. 1992). This measurement of change in lattice parameter, of course, is not possible if there is a change in the crystal structure due to disordering. Mossbauer spectroscopy techniques, superconducting properties and magnetization methods have also been employed to study the disordering phenomena (Bakker et al. 1995a). Recently, Zhou and Bakker (1996) have demonstrated that magnetic measurements are very powerful in determining the nature of disordering if one of the atoms involved has magnetic moments.
66
C. Suryanarayana
During disordering of CoZr, Cho and Koch (1993) did not observe any decrease in the intensity of superlattice reflections (relative to fundamental lines) with milling time before amorphization occurred. It was thus concluded that the grain boundary energy in this system is high enough to drive the crystal-to-amorphous transition without any disordering occurring. Zhou and Bakker (1996) observed that magnetization in the mechanically milled Co-Zr alloys increased continuously with milling time up to about 40 h (no reduction in the relative intensity of superlattice reflections was observed) and also that the lattice parameter of the disordered nanocrystalline phase increased with milling time. From these two observations, they concluded that disordering of CoZr occurs by atomic disorder and grain refinement (i.e., creation of additional grain boundaries). Thus, a combination of techniques can provide confirmatory evidence and also an accurate account of the mechanism of disordering. Whether an intermetallic transforms to a solid solution or an amorphous phase on milling is determined by the relative free energy values of the amorphous and crystalline phases with respect to the energy stored in the intermetallic by mechanical milling. For example, the intermetallic will transform to the solid solution if the enthalpy of the amorphous state as estimated by the Miedema analysis (Bakker 1998) is higher than that of the solid solution. Thus, if AG (milling)> AG a-c, where AG a-c represents the difference in free energy between the ordered crystalline and amorphous phases, complete amorphization occurs. On the other hand, if AG (milling)< AG a-c, no amorphization occurs; instead, a solid solution would form. Similar conclusions can also be reached by observing the nature of the phase diagrams. If the phase diagram shows that the intermetallic transforms to a (disordered) solid solution before it melts, then milling of the intermetallic will result in the formation of a (disordered) solid solution. Continued milling may then produce the amorphous phase. But if the intermetallic melts congruently (irreversibly ordered or permanently ordered intermetallic), then milling of the intermetallic will produce an amorphous phase directly. This has been shown to be true in a number of instances (Suryanarayana 1995a, Bakker et al. 1995a).
4. 6.4 Nanocrystalline materials Nanocrystalline materials are single-phase or multiphase materials, the crystal size of which is of the order of a few (typically 1-100) nanometers. Because of the extremely small size of the grains, a large fraction of the atoms is located in the grain boundaries and thus the material exhibits enhanced combinations of physical, mechanical, and magnetic properties (compared to material with a more conventional grain size, i.e., > 1 #m). Nanocrystalline materials show increased strength, high hardness, extremely high diffusion rates, and consequently reduced sintering times for powder compaction. The reader is referred to some recent comprehensive reviews for full details of the processing, properties, and applications of these materials (Gleiter 1989, Siegel 1991, Koch 1993, Suryanarayana 1995c and Chapter 12, this volume). Nanocrystalline materials can be synthesized by a number of techniques starting from the vapor phase (e.g., inert gas-condensation), liquid phase (e.g., electrodeposition, rapid solidification), and solid state (e.g., mechanical attrition). The advantage of using MA
Mechanical Alloying
67
for synthesis of nanocrystalline materials lies in its ability to produce bulk quantities of material in the solid state using simple equipment and at room temperature. Plastic deformation generally occurs by slip and twinning at low and moderate strain rates, while at high strain rates it occurs by the formation of shear bands, consisting of dense networks of dislocations. Transmission electron microscopy investigations have shown that plastic deformation by ball-milling, as anticipated, occurs via shear bands which are approximately 0.5 # m wide (Hellstern et al. 1989). The plastic strain in the material increases due to increasing dislocation density in the early stages of ball-milling. At a threshold dislocation density, even at only moderately elevated temperatures, the material relaxes into subgrains separated by low-angle boundaries, leading to a decrease of atomic level strain. During subsequent milling the process of high deformation/subgrain formation is repeated resulting in the subgrains becoming finer and finer, and the relative orientation of the subgrains with respect to each other ultimately becoming completely random. Once the subgrains reach a critical level of refinement, further refinement becomes virtually impossible since the stresses required for dislocation movement are enormously high due to the Hall-Petch strengthening. Thus, nanostructures with a minimum grain size are produced. In the case of fcc metals, this minimum grain size was found to scale with the melting temperature, or bulk modulus, of the given element (Eckert et al. 1992). However, in the case of bcc and hcp metals the minimum grain size attained was essentially independent of the melting temperature of the metals (Koch 1993). Nanostructure processing leads to the formation of supersaturated solid solutions and intermetallic compounds, which are otherwise difficult to obtain. For example, a supersaturated solid solution can be obtained in a system, if the grain size is in the nanometer regime; this is irrespective of the method by which the material is produced in the nanocrystalline state. Accordingly, the solid solubility of Bi in Cu is increased from the equilibrium value of <10 -4 to 4 at.% in the nanocrystalline state (Birringer et al. 1988). Similarly, formation of intermetallics such as Pd3Bi and CuEr has been achieved by nanostructure processing. As mentioned above, the enhanced solid solubility limits in mechanically alloyed materials is mostly due to the increased volume fraction of atoms residing at the grain boundaries in nanostructured materials. Since all the aspects of nanostructures are covered in Chapter 12 on "Nanocrystalline Materials" in this book, this topic will not be further covered in this chapter. 4.6.5 Solid-state amorphization
Amorphous phases are formed in suitable alloy systems by processing techniques such as rapid solidification from the liquid state, vapor-deposition, and laser-processing (Cahn 1991). In these methods, there is a change in the state of matter, i.e., a solid phase is formed from either the liquid or the vapor phase, and the high effective "quenching" rate and associated undercooling have been shown to be responsible for the amorphization process. There are also methods of amorphizing a solid without passing through the liquid (or vapor) state at any stage. Known as solid-state amorphization reactions (SSAR), these include irradiation, interdiffusion of metallic elements, pressure-induced vitrification, and mechanical deformation. Two reports of amorphization have triggered an increased research activity in this area. These are the synthesis of an Ni-Nb amorphous
68
C. Suryanarayana
phase starting from blended elemental powders of Ni and Nb by MA (Koch et al. 1983) and amorphization of intermetallic compounds (e.g., YCo3, Y2Co7, YCo5) by MM (Yermakov et al. 1981, 1982). Since then, amorphous phases have been reported in a number of alloy systems by both MA and MM. Amorphous phases have also been obtained by MM of quasicrystalline phases produced by RSP (Zhang and Wang 1996). Amorphous phases have been synthesized by MA starting from either elemental (of two or more types) or by MM of intermetallic compound powders. A reasonably exhaustive listing of the alloy systems in which amorphous phases have been found to form by MA or MM can be found in Gaffet et al. (1993) and Suryanarayana (1999). The semimetals Si and Ge have also been partially amorphized by MM (Gaffet et al. 1993). A further point to be considered is whether, or not, there exists a difference in the way an amorphous phase forms by MA from a mixture of blended elemental powders (A and B) and by MM from the intermetallic compound (Am Bn). Two reaction routes are possible: (i) A + B --+ (AmBn)crystalline ~ (AmBn)amorphous (ii) (Am Bn)crystalline ~ (Am Bn)amorphous Instances of occurrence of amorphization by both the routes can be found in the literature (Koch 1991, Suryanarayana 1995a, 1999). In general, the end product is the same irrespective of whether the starting material is a mixture of the elements, a single intermetallic phase, or a mixture of two or more such phases. However, it has been reported that MgToZn30 can be amorphized only by MM of the intermetallic, and not by MA of the blended elemental powder mixture (Calka and Radlinski 1989). Amorphization in ordered alloys seems to follow the sequence (Jang and Koch 1990): Ordered phase ~ disordered phase (loss of long-range order) -+ fine-grained (nanocrystalline) phase ~ amorphous phase. As mentioned earlier, it appears that an intermetallic (either as the starting phase or that formed on MA of blended elemental powders) becomes amorphous when its free energy is raised above that of the amorphous phase. The number of alloy systems that show amorphization by MA or MM is too numerous to list here. Even though it appears that any alloy system can be made amorphous under appropriate conditions of milling, it may not be true. This is because a number of variables, the more important of them being milling energy, milling temperature, and impurity contamination, control the constitution of the final product. Increased milling energy (achieved by a higher ball-to-powder weight ratio, increased speed of rotation, etc.) is normally expected to introduce more strain and increase the defect concentration in the powder and thus lead to easier amorphization. However, higher milling energies also produce more heat (and higher temperatures) and this can result in the crystallization of the amorphous phase. Therefore, a balance between these two effects will determine the nature of the final product phase. Examples of the occurrence of both crystalline and amorphous phases at high milling intensities (energies) are available in the literature (Koch 1996).
M e c h a n i c a l Alloying
69
A few reports also exist which describe the crystallization of amorphous alloys by MA. Eckert et al. (1988) reported that if the milling intensity were too high, Zr-Ni alloys which were amorphized at a lower milling intensity would exhibit crystalline structures. This was explained as due to the temperature rise of the powder caused by the increased kinetic energy of the milling media. Similar results were also reported in other systems (Mizutani and Lee 1990, Trudeau et al. 1990, Trudeau 1994). Mechanical milling has also resulted in the crystallization of amorphous alloys in some cases. Crystallization of amorphous alloys by MA/MM, whether formed during milling or as the starting phase, has been referred to as mechanical crystallization; this behavior is not well understood. It has also been reported that a crystalline F-TiA1 phase formed on milling beyond the amorphous phase in Ti- A1 alloys (Oleszak et al. 1993). This has been, however, suggested to be due to interstitial contamination of the powder (Suryanarayana 1995b). There have also been conflicting results on the effect of milling temperature on the nature of the phase formed. Generally, a lower milling temperature was found to accelerate the amorphization process (Pathak et al. 1993, Yamada and Koch 1993). See Koch (1994) for a detailed review on the temperature effects during MA/MM. The mechanism of amorphization by MA and/or MM is not clearly understood. The early investigators (Yermakov et al. 1981, 1982) assumed that the powder particles melted because of the very high rate of plastic deformation, and consequent rise in the powder temperature. Subsequent rapid quenching of the liquid by heat conduction into the less deformed, and hence cooler, interior regions of the particles results in the formation of the amorphous phase (like in RSP). However, energy input calculations and temperature measurements suggest that the temperature rise is not large enough for the powder particles to melt. Researchers now believe that amorphization during MA is not purely a mechanical process and that a solid-state reaction similar to that observed in thin films (Schwarz and Johnson 1983) occurs during MA. During MM, destabilization of the crystalline phase is thought to occur by the accumulation of point and lattice defects, and anti-phase boundaries. Schwarz and Johnson (1983) proposed that formation of amorphous phases from two elements in the solid state requires that (i) the two elements have a large negative heat of mixing in the amorphous (liquid) state (alloy systems with a positive heat of mixing have also been amorphized by MA; but, negative heat of mixing is an essential condition for glass formation by RSP) and (ii) the two elements have very different atomic sizes, allowing one of the components to diffuse more rapidly than the other. The negative heat of mixing provides the thermodynamic driving force for the reaction and the large difference in the diffusivities is necessary for favoring the kinetics of the reaction. Thus, owing to these factors, the free energy of the amorphous phase is lower than that of the mixture of the pure crystalline compounds, thus explaining the formation of the amorphous phase by MA of a mixture of pure components. Beke et al. (1991a, b) estimated the elastic mismatch energy stored in an ordered solid solution when its long-range chemical order is destroyed. They showed that if Tc > Tm and the ratio of the elastic mismatch energy to the ordering energy is high enough, amorphization is expected to occur. This has been proved to be a valid criterion in many cases of amorphization observed by irradiation.
70
C. S u r y a n a r a y a n a
Recently, attempts have been made to predict the ability of an alloy to become amorphous under MA/MM conditions. Since both atomic sizes and heats of formation are important in the formation of amorphous phases, Zhang (1993) plotted (RA -- R s ) / R 8 (where RA and R8 are the atomic radii of the components A and B, respectively) against A H, heat of formation of the amorphous phase, and observed that the regions where an amorphous phase forms or does not form can be separated by a straight line given by the equation:
(RA -- RB = A R ) / R B
= 0.068 A H + 0.716.
The proportion of correct predictions for the glass-forming alloys was 89.2% while that for the non-glass-forming alloys was only 71.4%. A theoretical model was also proposed in which amorphization was assumed to be realized through interstitial impurity formation during MA (Chakk et al. 1994). It was assumed that amorphization occurred when impurity atoms penetrated into interstitial sites and distorted the lattice locally. When the local distortions reached some critical value, the long-range order of the lattice was destroyed (destabilization of the crystalline phase) and an amorphous phase formed. Even though amorphization has been observed by both MA and MM, there is no quantitative data available to define the "critical" energy input (or other parameters) which can be equated to the critical cooling rate which should be exceeded to form the amorphous phase in RSP studies. Modeling of MA/MM (currently in progress in different laboratories) can provide additional information in this direction. Significant differences have been observed in the formation of metastable intermediate phases or amorphous alloys by MA/MM and RSP techniques. Formation of an amorphous phase has been reported in the Ti-A1 system by MA (due to contamination?); not possible by RSP methods. An important difference noted between these two methods is that while amorphous phases form in the vicinity of eutectic compositions by rapid solidification, they form near the equiatomic composition by MA. Furthermore, the glass formation range (the compositions range over which glasses can form) is wider in the mechanically alloyed condition than in rapid solidification. Formation of nanostructures has been reported in many alloy systems by mechanical alloying methods, but not by rapid solidification methods. But, devitrification of amorphous phases obtained by rapid solidification has led to the synthesis of nanostructure composites useful for magnetic applications (Suryanarayana 1995c, Lu 1996). Another important difference between mechanically alloyed and rapidly solidified alloys containing dispersoids appears to be in the size and distribution of the dispersoids. Figure 4.5 shows a pair of transmission electron micrographs comparing the matrix grain sizes, and size and distribution of dispersoids, in HIP compacts of Ti3Al-based alloys containing Er203 dispersoids. A Ti3A1 alloy to which 2 wt.% Er was added was rapidly solidified and then HIPed at 850~ (Fig. 4.5a). The mechanically alloyed material was Ti-25Al-10Nb-3V- 1Mo (at.%) to which 2 wt.% Er had been added and the alloy powder was HIPed at 1,000~ (Fig. 4.5b). Even after HIPing at a higher temperature, in comparison to the rapidly solidified alloy, the mechanically alloyed material showed a finer matrix
Mechanical Alloying
71
grain size, more uniform distribution of the dispersoids, absence of large dispersoids at the grain boundaries, and absence of dispersoid-free zones near the grain boundaries. Majority of the above-mentioned and other differences between materials processed by these two techniques could be traced to the fact that rapid solidification is a liquid-tosolid transformation while MA is a fully solid-state transformation.
4. 6.6 Displacement reactions It was recognized in 1989 that MA may be used as the basis of a chemical refining process (Schaffer and McCormick 1989) by demonstrating that the pure metal Cu can be synthesized when CuO and Ca were ball-milled together at room temperature. Simultaneous reduction of CuO and ZnO by Ca has also been shown to result in the formation of flbrass. This method of metal synthesis directly from their oxides or chlorides has now been extended to the synthesis of a number of metals, e.g., Zn, Ti, Zr, Ta, Gd, Er, Sm, V, W and some intermetallics, e.g., SmCos. Results of the Australian group have recently been summarized in a review article (McCormick 1995). Other groups also have now taken up this type of work using MA. Most of the reactions studied to date have been displacement reactions of the type: MO + R ---~ M + R O where a metal oxide (MO) is reduced by a reductant (R) to the metal (M). Metal chlorides also have been reduced to metals this way. The displacement reactions studied by MA are characterized by a large negative free energy change at room temperature and are therefore thermodynamically feasible at room temperature. However, commercial operations by the pyrometallurgical techniques are conducted at elevated temperatures to overcome the kinetic barriers and achieve sufficiently high reaction rates. Mechanical alloying can provide the means to substantially increase the reaction kinetics of the displacement reactions. This is because during MA, the repeated fracturing and welding of powder particles increases the area of contact between the particles due to a reduction in particle size and also allows fresh surfaces to come into contact repeatedly. This allows the reaction to proceed without diffusion through the product layer. Additionally, the high defect densities induced by MA/MM accelerate the diffusion process. As a consequence, these reactions will now occur at room temperature. If a reaction cannot occur at room temperature, the particle refinement and consequent reduction in diffusion distances (due to microstructural refinement) can at least significantly reduce the reaction temperatures. The product of the displacement reactions normally consists of two phasesmthe metal (or a compound) and the oxide or chloride associated with the reductant. The removal of the unwanted reaction by-product can be difficult due to the high reactivity of the metal phase associated with nanocrystalline grain sizes and intermixing of the phases induced by the MA process. The by-product phase can be easily removed if the metal particles are embedded in a continuous matrix of the byproduct phase. Removal of the byproduct is achieved by leaching the product mixture in a dilute acid or hot water, or by vacuum
C. Suryanarayana
72
RS+HIP'ed
850~
,,/!~; ~~
" ,
[ .....
' MA + HIP'ed
1000~
B
b
Figure 4 5 Transmission electron micrographs showing the difference in the m a trix grain size, and size and distribution of dispersoids in rapidly solidified (RS) and mechanically alloyed (MA) materials
t
Mechanical Alloying
TIC14-
Kroll Process +Mg (!) @ 1000" C
73
Ti sponge Melt
_., %.
Ingot Casting PREP
Ti powder Figure 4.6. Schematic illustration showing the route by which titanium alloy powder can be synthesized directly from its tetrachloride by mechanical alloying with a reducing agent.
distillation. The use of carbon or hydrogen as the reductant produces gaseous C0 2 or water vapor as the byproduct and obviates the need for leaching/distillation. The displacement reactions have now been studied to synthesize metals, alloys, intermetallics, and rare-earth permanent magnet alloys. The rare-earth alloys synthesized include Nd2Fel4B, Sm2Fel7N2.6, Sm2Fel4Ga3C2, SmCo5, Sm2Co17, SmFeTi and SmCoFe. The product of MA in these cases is usually an amorphous phase that needs to be heat treated to crystallize the hard-magnetic phase. An important advantage of the MA process is that nanocrystalline microstructures are developed and these result in remanence-enhanced magnetic materials. Since with remanence enhancement there is no need for an alignment process, significant potential exists for achieving cost savings and simplification of the manufacturing process. The mechanically driven displacement reaction offers a number of advantages over conventional metal processing techniques. First, it enables the reduction of a number of oxides and halides to pure metals at room temperature, thus effecting energy savings. Second, if a number of components are reduced simultaneously, then it is possible to produce an alloy without first having to convert the oxides to pure metals and then to the desired alloy. Third, for powder metallurgy applications it allows the direct formation of powder product without first having to manufacture the bulk alloy and then convert it to powder form. Thus, a number of high-temperature processes can be combined into one single room-temperature process with the potential for significant cost savings (Fig. 4.6) (Schaffer and McCormick 1992).
74
C. S u r y a n a r a y a n a
4.7. P O W D E R C O N T A M I N A T I O N
A major concern in the MA process is the impurities that get into the powder and contaminate it. The small size of the powder particles, availability of large surface area, and formation of fresh surfaces during milling all contribute to the contamination of the powder. Thus, it appears as though powder contamination is an inherent drawback of the technique, unless special measures are taken to avoid/minimize it. As mentioned earlier, in the early stages of MA, the metal powder coats the surface of the grinding medium and the inner walls of the container. This was expected to have prevented contamination of the powder and so no attention was paid to the problem of powder contamination. However, when different results were reported by different groups of researchers on the same alloy system, it was recognized that contamination could be a serious problem. This problem appears to be ubiquitous and is now encountered in many investigations, especially when reactive metals such as titanium and zirconium are milled. The magnitude of contamination appears to depend on the time of milling, the intensity of milling, the atmosphere in which the powder is milled, and the difference in strength~ardness of the powder and the milling medium. Whereas, 1-4 wt.% Fe has been found to be normally present in most of the powders milled with the steel grinding medium, amounts as large as 7 wt.% (20 at.%) Fe in a W-C mixture milled for 310 h (Wang et al. 1997), 13 wt.% (33 at.%) Fe in pure W milled for 50 h in a SPEX mill (Yang et al. 1993), and 60 ac% Fe in a W-5 at.% Ni alloy milled for 60 h in a SPEX mill (Courtney and Wang 1992) were also reported. These are very high levels of contamination. Iron contamination has also been observed in other alloy systems. For example, the presence of 16 at.% Fe was reported in a Ta-30 at.% A1 alloy powder milled for 400 h in a ball-mill (E1-Eskandarany et al. 1991). Similarly, large amounts of oxygen (up to 36.5 at.%) and nitrogen (up to 22.6 at.%) have also been reported to be present in AI-Ti powders milled for 400 h in a low energy ball-mill (Saji et al. 1992). Contamination of metal powders can be traced to (i) chemical purity of the starting powders; (ii) milling atmosphere; (iii) milling equipment; and (iv) the process control agents added to the powders. Contamination from source (i) can be either substitutional or interstitial in nature, while contamination from source (ii) is essentially interstitial in nature and that from (iii) is mainly substitutional, even though carbon from the steel equipment can be an interstitial impurity. Contamination from the PCA leads to interstitial contamination. Several attempts have been made in recent years to minimize the powder contamination during MA. One way of minimizing the contamination from the grinding medium and container is to use the same material for the container and grinding medium as the powder being milled. In this case also there will be wear of the grinding medium and the debris gets incorporated into the powder. Thus, even though there is no contamination, the chemistry of the final powder will be quite different from the starting powder; the metallic content (from the container and balls) would be higher than in the initial powder. This can be compensated for if we know how much of the metallic content is increased. The above solution may be possible in some cases; but it is difficult in many cases due to the non-availability of the special grinding medium and containers.
Mechanical Alloying
75
The problem of milling atmosphere is serious and has been found to be the major cause of contamination in many cases. In fact, it has been observed that if the container is not properly sealed, the atmosphere surrounding the container, usually air (containing nitrogen and oxygen), leaks into the container and contaminates the powder. Thus, when reactive metals like titanium and zirconium are milled in improperly sealed containers, the powders are contaminated with nitrogen and oxygen. Attempts have been made to improve the container seal integrity to prevent the outside atmosphere leaking inside. Use of high-purity argon (99.998%) atmosphere and improvements of the seal quality resulted in the processing of high-quality titanium alloy powder with as little as 100 ppm oxygen and 15 ppm nitrogen (Goodwin and Ward-Close 1995). This process, however, may not be economically viable and hence may not be feasible on an industrial scale.
4.8. M O D E L I N G
From the description of the MA process it is easy to realize that MA is a process involving a number of both independent and interdependent variables. Like in any other process, modeling of MA is also carried out to identify the salient factors affecting the process and establish process control instrumentation. By modeling the process effectively, it is possible to bring down the number of actual experiments to be conducted to optimize the process and achieve a particular application. The number of variables involved in the MA process is very large. For a particular alloy system the variables include the type of mill, intensity of milling, type of milling media, ball-to-powder weight ratio, the atmosphere under which the powder is milled, purity of the powders, milling time, milling temperature, and nature and amount of the PCA used. All these have a significant effect on the constitution of the powder. Even on a local scale, the nature of impacts between two balls, the frequency of impacts, the amount of powder trapped between two balls during a collision could vary from point to point. Thus, modeling the MA process is a difficult task. In spite of this, some attempts have been made in recent times (Courtney 1994, Courtney and Maurice 1996, Magini and Iasonna 1995, Magini et al. 1996, Abdellaoui and Gaffet 1996) and moderate success has been achieved in modeling the mechanics of the process. From the actual experiments conducted, attempts have also been made to correlate the phases formed with the process parameters during milling. But, we are far from a situation where we can predict the final chemical constitution of the powder (type and description of phases formed). As Courtney and Maurice (1996) succinctly put it: "It is important to appreciate what can be expected of even the most successful efforts of this kind. 'Absolute' predictions are unlikely; this is exacerbated in the case of MA by its inherently stochastic nature. In addition, models require input 'data'; material properties (often not known to precision, particularly at the high strain levels to which powder is often subjected during MA) and process characteristics (collision frequency and velocity, also seldom known a priori). Thus, realistic goals of process modeling are to correctly predict general trends, and perhaps even to predict resulting properties/dimensions within an order of magnitude. The benefits of successful models lie not in their abilities to predict outcomes, but in that they
76
C. Suryanarayana
Table 4.4 Room temperature and elevated temperature mechanical properties of commercial ODS nickel- and iron-base superalloys Alloy
MA 6000
MA 754
MA956
Test
0.2% YS
UTS
temperature o f (~
MPa
MPa
% El.
% RA
RT
1220
1253
7.2
6.5
871 (1600)
675
701
2.2
4.6
982 (1800)
445
460
2.8
1.9
RT
586
965
21
33
871 (1600)
214
248
31
58
982 (1800)
169
190
18
34
RT
517
655
20
35
97 69
100 72
12
30
1000 (1832) 1100 (2192)
MA 6000:15 Cr, 4.5 A1, 0.5 Ti, 2.0 Mo, 4.0 W, 2.0 Ta, 1.1 Y203, balance Ni. MA 754:20 Cr, 0.3 A1, 0.5 Ti, 0.6 Y203, balance Ni. MA 956:20 Cr, 4.5 AI, 0.5 Ti, 0.5 Y203, balance Fe.
help identify critical process and material variables, and reduce the amount of testing needed for process optimization".
4.9. I N D U S T R I A L A P P L I C A T I O N S
The major industrial applications of mechanically alloyed materials have been in the areas of thermal processing, glass processing, energy production, aerospace, and other industries. These applications are based on the oxide-dispersion strengthening effect achieved in mechanically alloyed nickel-, iron- and aluminum-based alloys. Mechanically alloyed materials are strong both at room and elevated temperatures (Table 4.4). The elevated temperature strength of these materials is derived from more than one mechanism. First, the uniform dispersion (with a spacing of the order of 100 nm) of very fine (5-50 nm) oxide particles (commonly used are Y203 (yttria), ThO2 (thoria), and La203 (lanthana)), which are stable at high temperatures, inhibit dislocation motion in the metal matrix and increase the resistance of the alloy to creep deformation. Another function of the dispersoid particles is to inhibit the recovery and recrystallization processes, because of which a very stable large grain size is obtained; these large grains resist grain rotation during high temperature deformation. A stable large grain size can also be obtained by secondary recrystallization mechanisms. Second, the very homogeneous distribution of alloying elements during MA gives both the solid-solution strengthened and precipitation-hardened alloys more stability at elevated temperatures and overall improvement in properties. Mechanically alloyed materials also have excellent oxidation and hot corrosion resistance. The increased resistance to oxidation-sulfidation attack is due to the homogeneous distribution of the alloying elements and the improved scale adherence due to the dispersoid itself (Hack 1987, Fischer et al. 1991).
Mechanical Alloying
77
Temperature, OF 1600
90 ~E
&
-~
50
--~
40
"-~
30
~
20
:3
0 0
o
i
" DS Mar-M200 + Hf
A SC PWA 454 o TD-Nickel
70 60
2200
9MA 6000
8O '
x:"
1900
=,,.,..
10 0 700
800
900
1000
1100
1200
Temperature, ~
Figure 4.7. Comparison of 1000-h specific rupture strength of INCONEL MA6000 with dispersion strengthened Mar-M200 + Hf, TD-Ni, and single-crystal PWA 454. Courtesy of Inco Alloys International.
4.9.1 Nickel-base alloys The most significant advantage of oxide-dispersion strengthened superalloys is the improvement in the stress rupture properties. Figure 4.7 compares the specific rupture strength (strength/density) for a 1,000-h life as a function of temperature for several nickel-base superalloys used for turbine blade applications. Mar-M200 is a nickel-base alloy containing by weight percent 9.0 Cr, 5.0 A1, 2.0 Ti, 12.0 W, 10.0 Co, 1.0 Nb and 1.8 Hf, while PWA454 is a nickel-base alloy containing 10.0 Cr, 5.0 A1, 1.5 Ti, 12.0 Ta, 4.0 W and 5.0 Co, and TD Ni is nickel containing 2.0 wt.% ThO2. It is clear from this figure that the MA 6000 alloy can maintain a given stress for a much longer time than a conventional alloy for similar vane applications. This is mainly due to the benefits of the combined strengthening modes in the mechanically alloyed material. The mechanically alloyed superalloys are considered mainly for three groups of applications--gas turbine vanes, turbine blades, and sheets for use in oxidizing/corrosive atmospheres. The largest use of MA 754 is as vanes and bands for aircraft gas turbine en-
78
C. Suryanarayana
~
R
Figure 4.8. Some typical gas turbine engine components made out of INCONEL alloy MA 754. Courtesy of Inco Alloys International.
gines (Fig. 4.8). For applications requiring good resistance to thermal fatigue, such as gas turbine vanes, Inconel MA 754 is given a strong texture. The majority of the grains are aligned so that their (100) axes are parallel to the principal working direction and along the length of the bar. Such texture results in low modulus of elasticity (149 GPa) in the longitudinal direction. The low modulus improves resistance to thermal fatigue by lowering stresses for given thermal strains. The alloy MA 758 is used in the glass industry for high-temperature components requiring both elevated temperature strength and resistance to extremely corrosive molten glass. It is also used for internal combustion engine components, mainly in critical fuel injection parts. One novel industrial application of alloy MA 754 is a high-temperature atmosphere-circulation fan in a "floating" furnace design being commercialized in Japan. Large rounds are used for the hub and plate material is used for blades of the fan, which operates at temperatures over 1,100~ The MA 6000 alloy is a more complex alloy developed as a blade material for advanced gas turbines. It is used for first and second stage turbine vanes and blades machined form solid bar. Unlike cast alloys, MA 6000 exhibits nearly flat rupture-life curves at high temperatures due to the combination of oxide dispersion strengthening and high grainto-width ratios (typically > 10 to 1). Because of its composition, MA 6000 has excellent resistance to oxidation and sulfidation. The characteristics of this alloy allow blade cooling to be reduced or eliminated as the metal temperature can be increased by 100~ or more in engines where the stresses are medium or low.
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79
4.9.2 Iron-base alloys The mechanically alloyed iron-base superalloys combine the high-temperature strength and stability of oxide dispersion strengthening with excellent resistance to oxidation, carburization, and hot corrosion. These alloys are suitable for use in gas turbine combustion chambers. Incoloy alloy MA 956 (20 Cr, 4.5 A1, 0.5 Ti, 0.5 Y203, balance Fe) is particularly well suited for use in heat processing applications. For example, vacuum furnace fixtures made of MA 956 have shown excellent durability and are able to compete with wrought molybdenum, which is also used in these applications. In comparison to molybdenum, MA 956 is about 30% lower in density, providing weight savings and cost advantages. Furthermore, since MA 956 has a lower vapor pressure than molybdenum, it will not coat the inside of the vacuum chamber or the parts being heat treated. Thus, MA 956 rods, flats, and sheets are used in numerous atmosphere and vacuum furnace applications including muffles, baskets, trays, and thermowells. Alloy MA 956 in tubing form has also been used for high temperature, severe service applications such as coke injection lance pipes in steel making. The alloy MA 956 is also being used in glass processing industry because of its resistance to attack by molten glass. Because of this corrosion resistance, the alloy is being evaluated for applications such as firing-kiln rollers, muffle tubes, and furnace racks. Other applications include molten-glass resistance heaters, thermocouple protection tubes, glass-processing components used in nuclear waste disposal and the bushings used to make single and multistrand fibers. More recently, MA 957 (14 Cr, 1.0 Ti, 0.3 Mo, 0.25 Y203, balance Fe) has been evaluated for use as the fuel cladding in fast neutron, breeder reactors. Conventional austenitic alloys are unsuitable for this application due to the dimensional swelling phenomenon caused by the high neutron fluxes. The mechanically alloyed materials are also being evaluated for heat exchanger components in high-temperature gas-cooled reactors.
4.9.3 Aluminum-base alloys The success of mechanically alloyed superalloys led to the development of dispersionstrengthened aluminum alloys. Since an aluminum oxide layer is always present either on the surface of the powder particles at the start of processing or during milling, its incorporation into the alloy contributes to significant improvements in the properties of the alloy. Furthermore, since aluminum is a ductile metal, PCAs are added to assist in minimizing cold welding during processing. Aluminum carbides are formed during MA by the decomposition of the PCA. Both the oxide or carbide type dispersions are about 30-50 nm in size and stabilize the ultrafine grain size. This results in 50% increase in strength, higher fracture toughness and improved resistance to stress corrosion cracking and fatigue crack growth of the mechanically alloyed materials. IncoMAP alloy AL-9052 (4.0 Mg, 1.1 C, 0.6 O, balance A1) has a density 5% less than that of conventional agehardenable aluminum alloy of comparable strengths such as 2024. With its combination of lightweight, high strength, and corrosion resistance, IncoMAP alloy AL-9052 is evaluated for aerospace applications where marine corrosion is also a factor. Addition of lithium to mechanically alloyed aluminum alloys has produced an ultra lightweight alloy IncoMAP alloy AL-905XL (4.0 Mg, 1.3 Li, 1.1 C, 0.6 O, balance
80
C. Suryanarayana
A1). Its density is 8% lower and stiffness 10% greater than the age-hardenable conventional alloy 7075-T73 of comparable strength. The excellent combination of the properties makes this alloy very attractive for airframe applications. In particular, the freedom from age-hardening treatments makes it possible to produce forgings and heavy sections with homogeneous metallurgical structures.
4.9.4 Magnesium-base alloys A useful application of the MA technique was in the production of supercorroding magnesium alloys that operate as short-circuited galvanic cells to corrode (react) rapidly and predictably with an electrolyte such as seawater to produce heat and hydrogen gas (Black 1979). Such an alloy system is suitable as a heat source for warming deep-sea divers, as a gas generator to provide gas for buoyancy, or as a fuel in hydrogen engines or fuel cells. Magnesium-base alloys containing Fe, Cu, Co, Cr, or Ti have been evaluated for such applications. The Mg-5 to 20 at.% Fe alloy is ideal because of its extremely fast reaction rate, high power output, and the high percentage of theoretical completion of the actual reaction. For corrodable release links an alloy with a slower reaction rate, such as Mg-5 at.% Ti is useful. There have also been a number of investigations in recent years to use the MA processing technique to produce metal hydrides. This is because metal hydrides are materials for safe storage of hydrogen and they can store hydrogen with a higher volume density than liquid hydrogen. However, these are sensitive to surface oxidation and hence can be a limiting factor in their commercial utilization. Nanocrystalline hydrides have a high density of defects and interfaces that could enhance diffusion, and therefore nanocrystalline intermetallics would not require activation treatments at high temperatures and pressures after exposure to air (Tessier et al. 1996). Several magnesium-base and iron-base intermetallics are being evaluated for this application. Figure 4.9 summarizes the applications of mechanically alloyed materials. The range of these applications is continuing to increase.
4.10. C O N C L U D I N G R E M A R K S
Mechanical alloying is a simple, elegant, and useful processing technique that continues to attract the serious attention of researchers. Even though the technique was originally developed to produce ODS superalloys, the synthesis of novel crystalline and amorphous phases has spurred lots of research investigations in recent years. It is estimated that so far about 4,000 research/review papers have been published in this area, with nearly 500 per year during the last 2-3 years. Both equilibrium and non-equilibrium phases and commercially useful materials could be synthesized starting from elemental powders. This also appears to be an economical process with vast potential. One of the greatest advantages of MA is in the synthesis of novel alloys that cannot be made by any other technique, such as alloying of normally immiscible elements. This is because MA is a completely solid-state processing technique and therefore limitations imposed by the phase diagrams do not apply here.
Mechanical Alloying
81
I~ STRENGTHENEDALLOYS ) ~_Ni-, Fe-, AND AI-BASE . . ~
Figure 4.9. Typical applications of the mechanical alloying process.
On an industrial scale, this is an accepted process. The ODS alloys have a higher temperature operating capability at the same stress or increased load-bearing capability at the same temperature than alloys without an oxide dispersion. The ODS alloys continue to find applications in a wide variety of industries. The two production facilities of Inco Alloys International have a combined annual powder capacity approaching 300,000 kg. The yield of the final product varies greatly with the product form and size, but the final wrought capacity is over 200,000 kg. Mechanical alloying involves powders with very small sizes and these should be handled with caution and care. Because of the large surface area, they are highly reactive and can be pyrophoric and can cause health problems when inhaled. Precautions should be taken not to open the powder to atmosphere immediately after milling since this can lead to oxidation of the powders and in some situations they can even catch fire.
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Gleiter, H. (1989) Prog. Mater Sci., 33, 223. Goodwin, R S. and Ward-Close, C. M. (1995) Mater Sci. Forum, 17%181, 411. Guo, W., Iasonna, A., Magini, M., Martelli, S. and Padella, E (1994) J. Mater Sci., 29, 2436. Hack, G. A. J. (1987) Metals and Materials, 3, 457. He, L. and Ma, E. (1996) J. Mater Res., 11, 72. Heinicke, G. (1984) Tribochemistry (Akademie Verlag, Berlin). Hellstern, E., Fecht, H. J., Garland, C. and Johnson, W. L. (1989) in Multicomponent Ultrafine Micrtostructures, Vol. 132, eds. McCandlish, L. E., Polk, D. E., Siegel, R. W. and Kear, B. H. (Materials Research Society, Pittsburgh, PA), p. 137. Herr, U. (1995) Key Eng. Mater., 103, 113. Huang, J. Y., Wu, Y. K., He, A. Q. and Ye, H. Q. (1994) NanoStructured Mater., 4, 293. Huang, J. Y., Yu, Y. D., Wu, Y. K., Li, D. X. and Ye, H. Q. (1997) J. Mater. Res., 12, 936. Ivanov, E. (1992) Mater Sci. Forum, 88-90, 475. Ivanov, E., Konstanchuk, I. G., Bokhonov, B. D. and Boldyrev, V. V. (1989a) Doklady Phys. Chem., 304(1-3), 82. Jang, J. S. C. and Koch, C. C. (1990) J. Mater Res., 5, 498. Jangg, G. (1989) in New Materials by Mechanical Alloying Techniques, eds. Arzt, E. and Schultz, L. (Deutsche Gesellschaft ftir Metallkunde, Oberursel, Germany), p. 39. Jangg, G., Kuttner, R and Korb, G. (1975) Aluminium, 51, 641. Klassen, T., Herr, U. and Averback, R. S. (1997) Acta Mater., 45, 2921. Koch, C. C. (1988) Inter Mater Rev., 33, 201. Koch, C. C. (1989) Ann. Rev. Mater Sci., 19, 121. Koch, C. C. (1991) in Processing of Metals and Alloys, Vol. 15 of Materials Science and Technology--A Comprehensive Treatment, ed. Cahn, R. W. (VCH Verlagsgesellschaft mbH, Weinheim, Germany), p. 193. Koch, C. C. (1993) NanoStructured Mater., 2, 109. Koch, C. C. (1994) Inter J. Mechanochem. Mech. Alloy., 1, 56. Koch, C. C. (1996) Scripta Mater., 34, 21. Koch, C. C. (1997) NanoStructured Mater., 9, 13. Koch, C. C., Cavin, O. B., McKamey, C. G. and Scarbrough, J. O. (1983) Appl. Phys. Lett., 43, 1017. Li, W., Suryanarayana, C. and Froes, F. H. (1995a) in Synthesis and Processing of Lightweight Metallic Materials, eds. Froes, E H., Suryanarayana, C. and Ward-Close, C. M. (TMS, Warrendale, PA), p. 203. Li, W., Suryanarayana, C. and Froes, E H. (1995b) in Advances in Powder Metallurgy and Particulate Materials, Part 1, compiled by Phillips, M. A. and Porter, J. (Metal Powder Industries Federation, Princeton, NJ), p. 145. Liebermann, H. H. (1993) Rapidly Solidified Alloys: Processes, Structures, Properties, Applications (Marcel Dekker, New York). Lu, K. (1996) Mater. Sci. Eng. Rept., R16, 161. Lu, L. and Lai, M. O. (1998) Mechanical Alloying (Kluwer, Norwell, MA). Luton, M. J., Jayanth, C. S., Disko, M. M., Matras, S. and Vallone, J. (1989) in Multicomponent Ultrafine Micrtostructures, Vol. 132, eds. McCandlish, L. E., Polk, D. E., Siegel, R. W. and Kear, B. H. (Materials Research Society, Pittsburgh, PA), p. 79. Ma, E., He, J.-H. and Schilling, R J. (1997) Phys. Rev., B55, 5542. Magini, M. and Iasonna, A. (1995) Mater Trans. JIM, 36, 123. Magini, M., Iasonna, A. and Padella, E (1996) Scripta Mater., 34, 13. Martelli, S., Mazzone, G., Scaglione, S. and Vittori, M. (1988) J. Less-Common Metals, 145, 261. Massalski, T. B., Okamoto, H., Subramanian, R R. and Kacprzak, L., eds. (1990) Binary Alloy Phase Diagrams, second edition, (ASM International, Materials Park, OH). Maurice, D. R. and Courtney, T. H. (1990) Metall. Trans., 21A, 289. McCormick, R G. (1995) Mater Trans. JIM, 36, 161. McDermott, B. T. and Koch, C. C. (1986) Scripta Metall., 20, 669. Mizutani, U. and Lee, C. H. (1990) J. Mater Sci., 25, 399. Moore, J. J. and Feng, H. J. (1995) Prog. Mater Sci., 39, 243. Mukhopadhyay, D. K., Suryanarayana, C. and Froes, E H. (1994) Scripta Metall. Mater., 31, 333. Mukhopadhyay, D. K., Prisbrey, K. A., Suryanarayana, C. and Froes, E H. (1996) in Proceedings of the International Conferene on Tungsten and Refractory Metals 3, eds. Bose, A. and Dowding, R. J. (Metal Powder Industries Federation, Princeton, NJ), p. 239.
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C. S u r y a n a r a y a n a
Murty, B. S. and Ranganathan, S. (1998) Inter Mater Rev., 43 (in press). Murty, B. S., Mohan Rao, M. and Ranganathan, S. (1993) NanoStructured Mater., 3, 459. Nasu, S., Miglierini, M., Ishihara, K. N. and Shingu, P. H. (1992) J. Phys. Soc. Japan, 61, 3766. Oleszak, D., Burzynska-Szyszko, M. and Matyja, H. (1993) J. Mater Sci. Lett., 12, 3. Pathak, D., Yamada, K. and Koch, C. C. (1993) in Mechanical Alloying for Structural Applications eds. deBarbadillo, J. J. et al. (ASM International, Materials Park, OH), p. 205. Politis, C., Krauss, W., Leitz, H. and Schommers, W. (1989) Modern Phys. Lett., 3, 615. Radlinski, A. P. and Calka, A. (1991) Mater Sci. Eng., A134, 1376. Rawers, J., Govier, D. and Cook, D. (1995) Scripta Metall. Mater., 32, 1319. Russell, K. C. (1985) Prog. Mater Sci., 28, 229. Saji, S., Abe, S. and Matsumoto, K. (1992) Mater. Sci. Forum, 88-90, 367. Schaffer, G. B. and McCormick, P. G. (1989) Appl. Phys. Lett., 55, 45. Schaffer, G. B. and McCormick, P. G. (1992) Mater Forum, 16, 91. Schropf, H., Kuhrt, C., Arzt, E. and Schultz, L. (1994) Scripta Metall. Mater., 30, 1569. Schulz, R. (1996) Materials Science Forum, 225-227 (Proceedings of the International Symposium on Metastable, Mechanically Alloyed, and Nanocrystalline Materials, Quebec, Canada, 1995). Schwarz, R. B. and Johnson, W. L. (1983) Phys. Rev. Lett., 51, 415. Schwarz, R. B., Petrich, R. R. and Saw, C. K. (1985) J. Non-Cryst. Solids, 76, 281. Seki, Y. and Johnson, W. L. (1990) in Solid State Powder Processing, eds. Clauer, A. H. and de Barbadillo, J. J. (TMS, Warrendale, PA), p. 287. Shingu, P. H. (1992) Materials Science Forum, 88--90 (Proceedings of the International Symposium on Mechanical Alloying, Kyoto, Japan, May 1991). Siegel, R. W. (1991) in Processing of Metals and Alloys, Vol. 15 of Materials Science and Technology--A Comprehensive Treatment, ed. Cahn, R.W. (VCH Verlagsgesellschaft mbH, Weinheim, Germany), p. 583. Singer, R. E, Oliver, W. C. and Nix, W. D. (1980) Metall. Trans., llA, 1895. Stoloff, N. S. and Davies, R. G. (1966) Prog. Mater Sci., 13, 77. Suryanarayana, C. (1995a) Bibliography on Mechanical Alloying and Milling (Cambridge International Science Publishing, Cambridge, UK). Suryanarayana, C. (1995b)Intermetallics, 3, 153. Suryanarayana, C. (1995c) Inter Mater Rev., 40, 41. Suryanarayana, C. (1996) Metals & Materials, 2, 195. Suryanarayana, C. (1998) in Powder Metal Technologies and Applications, ASM Handbook, Vol. 7 (ASM International, Materials Park, OH) (in press). Suryanarayana, C. (1999) Prog. Mater Sci. (in press). Suryanarayana, C. and Froes, E H. (1990) J. Mater Res., 5, 1880. Suryanarayana, C. and Jones, H. (1988) Inter J. Rapid Solidification, 3, 253. Suryanarayana, C., Zhou, E., Peng, Z. and Froes, E H. (1994a) Scripta Metall. Mater., 30, 781. Suryanarayana, C., Li, W. and Froes, E H. (1994b) Scripta Metall. Mater., 31, 1465. Suryanarayana, C., Korth, G. E. and Froes, E H. (1997) Metall. Mater Trans., 28A, 293. Suryanarayana, C., Ivanov, E., Noufi, R., Contreras, M. A. and Moore, J.J. (1999) J. Mater Res., 14 (in press). Takeuchi, T., Yamada, Y., Fukunaga, T. and Mizutani, U. (1993) J. Japan Soc. Powder Powder Metall., 40, 951. Tessier, P., Zaluski, L., Zaluska, A., Str6m-Olsen, J. O. and Schulz, R. (1996) Mater Sci. Forum, 225-227, 869. Trudeau, M. L. (1994) Appl. Phys. Lett., 64, 3661. Trudeau, M. L., Schulz, R., Dussault, D. and Van Neste, A. (1990) Phys. Rev. Lett., 64, 99. Wang, G. M., Campbell, S. J., Calka, A. and Kaczmarek, W. A. (1997) J. Mater Sci., 32,1461. Weeber, A. W. and Bakker, H. (1988) Physica B, 153, 93. Yamada, K. and Koch, C. C. (1993) J. Mater Res., 8, 1317. Yang, E., Wagner, C. N. J. and Boldrick, M. S. (1993) Key Eng. Mater., 81-83, 663. Yavari, A. R. (1995) Materials Science Forum, 179-181 (Proceedings of the International Symposium on Metastable, Mechanically Alloyed, and Nanocrystalline Materials, Grenoble, France, June 27-July 1, 1994). Yavari, A. R. and Desr6, R (1990) Journal de Physique 51, Colloque C4, Supplement No. 14 (Proceedings of the International Symposium on Multilayer Amorphisation by Solid-State Reaction and Mechanical Alloying, Grenoble, France, February 21-23 1990). Yermakov, A. E., Yurchikov, E. E. and Barinov, V. A. (1981) Phys. Met. Metallogr., 52(6), 50. Yermakov, A. E., Barinov, V. A. and Yurchikov, E. E. (1982) Phys. Met. Metallogr., 54(5), 90.
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Yvon, E J. and Schwarz, R. B. (1993) in Structural Applications of Mechanical Alloying, eds. deBarbadillo, J.J. et al. (ASM International, Materials Park, OH), p. 393. Zhang, H. (1993) J. Phys.: Condens. Matter, 5, L337. Zhang, E X. and Wang, W. K. (1996) J. Alloys & Compounds, 240, 256. Zhou, E., Suryanarayana, C. and Froes, E H. (1995) Mater. Lett., 23, 27. Zhou, G. E and Bakker, H. (1996) Scripta Mater., 34, 29.
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Chapter 5 Laser Processing 5.1. Principles of Lasers 5.2. Classifications of Laser Processing 5.3. Analysis of the Laser Melting and Quenching Process 5.3.1 Heat Transfer 5.3.1.1 Absorption Phenomena of Laser 5.3.1.2 Pulsed Laser 5.3.1.3 Continuous Wave (CW) Laser 5.3.2 Kinetic Conditions for Solidification 5.3.2.1 Kinetic Conditions for Amorphous Phase Formation 5.3.2.2 Solidification Modes in the Laser Processing of Materials 5.3.3 Microstructure Selection Maps (MSM) for Alloys 5.4. Laser-Quenching 5.4.1 Amorphous Phase Formation 5.4.2 Formation of Crystalline Phases 5.5. Laser Surface-Alloying and Cladding 5.5.1 Ferrous Alloys 5.5.2 Non-Ferrous Alloys 5.6. Laser-Annealing 5.7. Laser-Beam Joining 5.7.1 Welding of Structural Materials 5.7.1.1 Ferrous-Based Alloys 5.7.1.2 Non-Ferrous Alloys 5.7.2 Microjoining 5.8. Conclusions References
89 89 92 92 92 93 94 96 96 97 97 98 98 99 100 100 101 102 106 106 106 108 112 115 115
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Chapter 5 Laser Processing KOJIRO E KOBAYASHI
5.1. P R I N C I P L E S O F L A S E R S
The laser (light amplification by stimulated emission of radiation) is a unique source of electromagnetic radiation. Lasers are distinguished from other light sources by their property to emit a coherent radiation of high spectral purity (Bimberg 1986). The total power output of a laser can be focused down to a spot of an extremely small diameter and so large intensities can be obtained. The use of focused laser energy allows one to accomplish the heat treatment of materials in a controlled spatial and temporal domain. There are at present a large number of commercially available laser sources that could be used to induce the melting needed for surface alloying. Of these a relatively small group, including the carbon dioxide (CO2), neodymium-doped yttrium-aluminum garnet (Nd:YAG), and EXCIMER (excited dimer), account for nearly all the practically used sources (Draper and Poate 1985, von Allmen 1983). The laser characteristics such as wavelength, output power, beam diameter, output mode, pulse length, and repetition rate are important processing factors. Table 5.1 lists the powers and wavelengths of different lasers. The laser processing system for a pulsed or Q-switched source offers cooling rates exceeding those of rapid solidification from the melt by several orders of magnitude. Qswitched mode operation can provide extremely higher peak power pulses and higher pulse rates than normal pulsed mode. Each individual focused laser event produces a melt spot and those melt spots must be raster-scanned on the surface if area coverage larger than the spot itself is desired. In the case of the continuous wave laser the area per unit time scanned is dependent on the effective spot size, the degree of overlap, and the relative transverse velocity.
5.2. C L A S S I F I C A T I O N S O F L A S E R P R O C E S S I N G
Figure 5.1 shows schematically the major laser processes that relate to the formation of non-equilibrium phases. The definition and characteristics of each of these processes is outlined below (Roessler 1986). Laser surface melting and quenching (Fig. 5.1(a)): A laser-beam is scanned over the surface of a substrate to be melted. Behind the laser-beam the molten pool cools very rapidly because of high heat conduction from the molten pool into the still cold substrate. 89
K. F. Kobayashi
90
Laser beam Molten / pool ~ / .............
~ ~
Laser beam /\Molten Alloying / ~pool powder / ~/
Rapidly solidified layer
(a) Laser surface melting and quenching
(b) Laser surface alloying Laser beam
Laser beam Cladding / powder [ \ [ ~,~,~~,~; 9
Alloyed layer
/ Molten ~ pool ~,/
Clad layer
Region heated above annealing / temperature ~
(c) Laser surface cladding Focused IA~mrhmam
Annealed and
/ rapidlycooled layer
(d) Laser annealing
Weld bead
(f) Laser beam soldering
(e) Laser beam welding l-- ......substrate .....
Laser beam ~ A ~ ~
p
I
~Clusters Plasma plume
(g) Laser ablation
Figure5.1. Schematicrepresentationsofvariouslaserprocessesformaterials.
Laser Processing
91
Table 5.1. Powersand wavelengths of different lasers. Source c02 laser (gas) YAG laser (solid: Nd) EXCIMER (gas: ArF) (gas: KrF)
Wavelength (microns)
Power
10.6 1.06
~50 kW CW: ~6 kW Pulse: ~20 kW ~50 W ~150 W
0.193 0.248
This produces a rapidly solidified surface layer and thereby results in the formation of amorphous or metastable phases on the surface. This process is known as laser-glazing. Laser surface-alloying (Fig. 5.1(b)): In surface alloying with a laser, an alloying material, which is delivered to the surface of a substrate by either preplacing a powder on the substrate or blowing the powder into the laser-generated melt pool, is fused to and mixed with the substrate so that the surface results in an alloyed layer after solidification. The alloyed layer shows a fine and homogeneous microstructure with minimal segregation because of the rapid solidification. The thickness of the alloyed layer can be from a few microns to about 1 mm. Laser surface-cladding (Fig. 5.1(c)): Laser surface cladding is similar to laser surface alloying except that the surface clad layer is fused to the substrate with the minimum of dilution with the substrate. Thus, the objective in this process is to provide the characteristics of a cladding material itself to the surface of a substrate. Clad layers having a thickness of up to several millimeters can be formed. Laser-annealing (Fig. 5.1(d)): A laser-beam is scanned over the surface of a substrate and heats it above an annealing temperature before rapid cooling. The objective of laser annealing is to achieve homogenization, solutionizing, recrystallization and/or removing defects locally in the surface layer. Laser annealing can allow very selective annealing of the substrate and extremely high cooling rates in comparison with conventional furnace annealing. Laser-beam welding (Fig. 5.1(e)): Fusion welding with a laser is a deep-penetration welding process. This process is characterized by the formation of a "keyhole". The laser energy causes local vaporization of material and generates a keyhole filled with vaporized metal through the thickness of the workpiece (Fig. 5.1(e)). As the focused laser-beam is scanned over the weld line, the keyhole surrounded by a molten pool moves. Continuous melting and vaporization occur at the leading edge of the keyhole. The molten metal flows back around the keyhole and resolidifies at the rear edge of the molten pool. Thus, a fusion weld joint with deep-penetration can be obtained. The major advantage of laser welding is selective heat input with narrow fusion and heat-affected zones. On the other hand, rapid heating and cooling cycles may result in the formation of unusual metastable phases in the weld fusion and heat affected zones. So, the effect of microstructural changes on the properties of the weld joint should be considered. Laser-beam soldering (Fig. 5.1(f)): Laser-beam joining can be applied to microjoining in electronics industries. Figure 5.1 (f) shows schematically laser-beam soldering between
K. E Kobayashi
92
the lead of a quad-flat pack surface-mounted component and a pad of a printed wiring board. A laser is irradiated onto the lead and the absorbed laser energy melts a solder. The solder solidifies and the solder joint is made after the laser irradiation is stopped. Laser power, beam diameter, and duration time are precisely controlled. Laser-ablation (Fig. 5.1(g)): Laser ablation is a process for the deposition of thin films. When a high energy laser-beam is irradiated onto a target surface, atoms and ions are emitted from the target to form a plasma plume. The plasma expands, cools and is condensed into clusters. The atoms, ions, clusters and/or molecules are collected on a substrate, resulting in a thin film deposition on the substrate surface. The deposited film usually has a non-equilibrium amorphous structure. The applications and characteristics of the above laser processes will be described in detail below in Sections 5.4 to 5.7.
5.3. A N A L Y S I S OF T H E L A S E R M E L T I N G AND Q U E N C H I N G P R O C E S S
5.3.1 Heat transfer The recent availability of lasers with high-power directed energy sources has helped in the development of rapid melting and solidification techniques in which a semi-infinite bulk substrate in intimate contact with a molten layer acts as the quenching medium. All mathematical treatments of classical and laser heat treatments are based on the fundamental equation for heat conduction:
OT = OT { 8t
)~ OT
)~ OT ~
O ~ C p p Ox ! Cpp 8y
)~ OT~
Cpp Oz)
qi
Cpp
(5.1)
Here T stands for the absolute temperature, t for the time, x, y and z are the coordinates, ~. the thermal conductivity, C p the specific heat, p the density and q i is the rate at which heat is supplied per unit time and volume. If all the parameters involved are known, Eq. (5.1) can be solved numerically by either finite difference or finite element methods (Mehrabian et al. 1979, Bergmann et al. 1981, Fritsch 1986).
5.3.1.1 Absorption phenomena of laser Absorption of laser power by the substrate is related to the laser-solid interaction. When an electromagnetic wave interacts with the atoms in the surface of a substrate, the field causes an electrical current. The Joule effect increases the surface temperature of the substrate due to the finite electrical conductivity of the material. Generally, the power of the laser beam is subdivided into the refected part, the absorbed part and the transmitted part. For metals that are not transparent, the last part can be neglected. Assuming that the Wiedemann-Franz law, i.e., the ratio between the electrical and thermal conductivity is constant, is valid for metals, metals with good conductivity exhibit poor absorption of laser. Similarly, the absorption ratio (fraction of the incident laser energy absorbed by the metal surface) increases with increasing temperature because the conductivity of metals decreases. When the surface is heated continuously by a laser, the surface layers start melting which makes the electrical conductivity lower and makes the absorption ratio higher.
Laser Processing
o 0.8
93
Ag
,,m
(u L.
s 0
0.6
~-
0.4
Cu
i n
k,.
0 .Q <
Fe
0.2
0.4
1
2
4
8
W a v e length (l~nn) Figure 5.2. Changein absorption ratio as a function of wavelengthfor metal surfaces.
In electromagnetic fields, the frequency affects the electrical conductivity and hence the laser light with different wavelengths shows different absorption values for a given material. Figure 5.2 shows the change in absorption ratio as a function of laser wavelength for metal surfaces. It is recognized that the CO2 laser with a 10.6 # m wavelength has an inferior absorption ratio to Nd: YAG laser with shorter wavelength (1.06/zm). Additionallyl reactions with the molten metal and the surrounding atmosphere may also affect the absorption ratio.
5.3.1.2 Pulsed laser The nature of interaction between a laser-beam and a material surface is controlled by two variables. One is the absorbed power density from the laser and the other is the interaction time. In the pulsed mode, interaction time means the time of laser irradiation. In the rapid solidification surface treatment process, glazing requires a high power density and substantially shorter interaction times. These requirements are met by the application of an Nd: YAG laser with Q switch. As discussed in detail by Bloembergen (1979) and others (Spaepen and Turnbull 1982, Linet al. (1984), picosecond pulse lasers can be used to induce very high quench rates, up to 1012 K/s, in thin molten overlays on a metallic substrate. During laser-glazing, the modified surface is considered to go through the following stages: (i) energy deposition in a layer on the order of the absorption depth (~50 nm), (ii) very fast melting ( ~ 1000 m/s) of a ~ 100 nm thick layer, and (iii) cooling and solidification.
94
K. F. Kobayashi
The quenching rate, u, is determined as a function of temperature gradient at the interface from the thermal balance.
u --
(
~.V 6 Hf
0T) Ox interface
9
(5.2)
o
Here V and A H f represent the molar volume and molar heat of fusion, respectively. On the other hand from the kinetics of interface motion u = u
Tm-T/ ~ Ti
0
(5.3)
where Tm stands for melting temperature, T/ for interface temperature and u0 is the velocity prefactor. For collision-limited motion u0 almost equals to the speed of sound, while for diffusion-limited motion uo = D~ I, where D is the diffusivity in the liquid and 1 is the interatomic distance. Generally u0 has a value between these two extremes because of a mixed mechanism. Since the temperature gradients in this process are very steep (1011 K/m during melting and 101~ K/m during cooling), u from Eq. (5.2) becomes very large and at the same time, Ti becomes considerably different from Tm from Eq. (5.3). The lifetime of the melt, determined experimentally (Lin and Spaepen 1981, Lin and Spaepen 1983) or estimated from heat flow calculations, is reported to be about 1 ns. Since the temperature of the material rises several thousand degrees in that period, the cooling rate can be estimated to be at least d T / d t = 103 K/10 -9 s = 1012 K/s (Lin and Spaepen 1984). This rapid solidification can lead to the formation of new metastable crystalline phases, or, if the crystal growth kinetics are slow enough, glass formation results. 5.3.1.3 Continuous wave (CW) laser In the pulsed mode, very high cooling rates are achieved due to the short interaction time. At the same time, however, this makes the modified region extremely small. Normally for industrial purposes, very high cooling rates are not required and the continuous mode is applied to modify the surface effectively 9 The absorbed heat-flux distribution, which is a function of the position and time, is treated moving in a CW laser while it is treated as stationary in a pulsed laser. In the basic investigation of laser-beam processing, an important aspect is the use of appropriate heat flow models. If all the parameters in Eq. (5.1) are known, it can be solved numerically by either finite difference or finite element methods 9 However, the material parameters are not always constant and their temperature dependence is not always known. Internal heat sources or sinks during transformation, melting or evaporation are not also always negligible. Cline and Anthony (1977) showed that the moving point source model, which involves an analytical solution of the thermal diffusion equation using Green's functions, could be used to estimate the temperature distribution during CW laser irradiation. During exposure and removal of a laser with a uniform energy flux onto the surface, the heating and cooling of a semi-infinite plate during the glazing process can be modeled conveniently and simply (Greenwald et al. 1979).
95
Laser Processing
During heating, the temperature at a point x below the surface at time t is given by (5.4)
T ( x , t) -- To + qJ(x, t)
and 2qi[~t 7*(x,t)----~
(
-exp
where To is the ambient temperature and tc During cooling
-
-
-~
x2 )
-
x err ( 2 ~ x t ) ]
(5.5)
)~/(CplO) is the thermal diffusivity.
T ( x , t) -- To + qJ(x, t) - ~P(x, t - r)
(5.6)
where r is the interaction time, defined in continuous mode as (beam diameter)/(beam traveling speed). The times for initial melting, r* and final solidification r** are given by Eqs. (5.7) and (5.8). r* =
Tm-T0
ab
(5.7)
(r + r*) 2 r** =
4r*
(5.8)
Here a and b are defined as follows: a -- v/rr/4K
(5.9)
b - qi/)~.
(5.1 O)
From these relations, the trailing liquid zone length, L, may be determined, since L is simply equal to v(r** - r*), in which v is the beam traveling speed. The compatibility of the model is checked by comparing this calculated length with the experimental value. Under conditions in which the energy inputs range from 1 to 100 J/cm, the heating effect is concentrated in a very thin region at the material surface. This gives rise to extremely high thermal gradients, which promote rapid solidification of the melt. At the same time, the melt and the substrate are in intimate contact, which makes the heattransfer coefficient between them infinity. These effects enable the solidification to occur faster in the case of atomization and splat-cooling types of processes. As a consequence of the extremely rapid solidification, a variety of interesting metallurgical microstructures can be developed, some of which are unique to the laser glazing process.
96
K. E Kobayashi
5.3.2 Kinetic conditions f o r solidification Crystallization kinetics, cooling rate, and laser parameters are closely connected for the intelligent processing of materials. 5.3.2.1 Kinetic conditions for amorphous phase formation The conditions for amorphous phase formation are to be derived from the nucleation and growth theory (Uhlman 1972, Davies et al. 1974, Kobayashi and Shingu 1988). Classical nucleation theory gives the following expression for the nucleation frequency J (T) (Turnbull 1950, Pound and Lamer 1974, Miyazawa and Pound 1974, Spaepen 1975, Spaepen and Turnbull 1976):
100 [
J ( T ) -- ~ e x p rl
-
3
1 - Tr
(5.11)
where r/= viscosity of liquid, Tr = reduced temperature (=T/Tm) and Tm = equilibrium melting point; c~ is a factor which depends on the atomic arrangement at the interface with a value close to unity, and/3 is the entropy change due to solidification in units of the gas constant. The growth rate of crystal nuclei is given by molecular kinetic considerations assuming that the atomic transport across the interface region, which usually is assumed as several atoms diameters in thickness, takes place by a diffusive process according to the equation: Ri(T) :
K r/(1 -- Tr)
(5.12)
where K = d 3L f/32rq6 and d = atomic diameter, L f = heat of fusion, and q~ = interface thickness. The temperature dependence of viscosity in the undercooled liquid region is represented assuming a Tamrnann and Hesse type of relation, as log 7? = A +
B
T - T~
(5.13)
where A, B and T~ are constants which are determined assuming r/ -- 1013 poise (10 le Pa sec) at T -- Tg (glass temperature); r/ is the value for the Arrhenius relation at T = Tm and d(log r/)/d(1/T) is the slope of the Arrhenius relation at T = Tm. The general form of the fraction of phase transformed X may be given as - In(1 -- X) -- -4n. ~ fot J (T) [fo t Ri (co) do)]3 dO
(5.14)
where J ( T ) is the nucleation frequency, Ri(w) is the linear growth velocity, X the crystallized fraction and t the time after passing the melting point during cooling. The temperature-time-transformation (TTT) diagram and the continuous-coolingtransformation (CCT) diagram can be obtained from Eq. (5.14) (Cahn 1956, Grange and Kiefer 1941). By the use of Eq. (5.14) we can estimate the condition for amorphous
Laser Processing
97
phase formation by requiring that the crystallized volume fraction X in the case of laser quenching be very low. 5.3.2.2 Solidification modes in the laser processing o f materials
The surface treatment of materials by laser is of great interest in many technical applications. By moving at high speed, a high power beam, such as that produced by a CO2 laser, over a surface heat treatment or rapid melting-solidification of the base material can be obtained at depths varying from a few microns to several millimeters. For the case of CW laser sources and metals, heat flow has been modeled by Mehrabian and co-workers (Mehrabian 1982, Hsu et al. 1978, Hsu et al. 1980, Kou et al. 1981, Sekhar and Mehrabian 1981) and others (Mazumder and Steen 1980, Chande and Mazumder 1982, Chande and Mazumder 1983, Lin and Spaepen 1984). During laser remelting, the rate of solidification governs the possible formation of metastable phases, their composition, and the scale of the microstructure (Boettinger et al. 1980, Kurz et al. 1986, Trivedi et al. 1987). Laser resolidification starts epitaxially on a crystalline substrate and nucleation plays a less important role. At the high rates of absolute stability, one cannot assume equilibrium partitioning any more. Such departure from equilibrium partitioning was discussed by Boettinger et al. (1984) who used two response functions to describe the temperature and the solid composition at the interface, where the variation of the partition coefficient with growth velocity was given by Aziz (1982) and Jackson et al. (1980). An amorphous phase should form in two different cases. In the first case, the cooling rate is too fast for partitioning of atoms to occur. In the second, partitioning and/or rearrangement of atoms for crystallization is too slow for a given cooling rate (Kobayashi and Shingu 1982, Kobayashi et al. 1984). 5.3.3 Microstructure selection maps ( M S M ) f o r alloys
Recently a series of solidification models have been developed which are valid at low and high growth velocities for different microstructures (Trivedi et al. 1987, Kurz et al. 1986, Boettinger et al. 1988, Lipton et al. 1987, Carrad et al. 1992). These models predict phase selection and the solidification microstructure of an alloy as a function of solidification parameters such as the growth rate and temperature gradient. These are discussed for dendrites (Kurz et al. 1986, Gill and Kurz 1995, Gilgien et al. 1995, Kurz 1995) or eutectic (including peritectic) structures (Rappaz et al. 1987, Kurz and Trivedi 1994, Gill and Kurz 1995, Gilgien et al. 1995) in rapid solidification processing such as in welding and laser treatment. MSMs contain information on the phase selection, solidification morphology and sometimes size of the microstructure as a function of composition and growth rate (or growth temperature, often combined with the equilibrium phase diagram). The modification of laser treatment conditions by controlling the atmosphere and the initial temperature of the sample allowed the reduction of instabilities in the liquid bath at low beam velocities. This made it possible to observe the metastable eutectic A1-A16Fe and the transition from the metastable eutectic to Al-dendrites for A1-4at%Fe alloy during laser resolidification (Gilgien et al. 1995). In the case of stainless steel, metastable phases such as y may form in a composition range where under normal welding conditions
98
K. F. Kobayashi
still forms. This becomes clear when one considers the fact that a phase with less solute rejection is preferred at high growth velocities even if its liquidus temperature is lower than that of the equilibrium phase (David and Vitek 1989). In high power laser treatment (>5 kW) or in welding, the melt pool might become so large that buoyancy driven convection becomes important (Kurz and Trivedi 1994). Preheating lowers the surface temperature gradient that decreases the driving force for Marangoni convection. Preheating is therefore useful for low velocity laser resolidification.
5.4. LASER-QUENCHING
5.4.1 Amorphousphaseformation Laser quenching yields an amorphous phase over most of the phase diagram in Pd-Si and Pt-Si systems (von Allmen et al. 1980). For the formation of an amorphous phase by laser-alloying methods, it must not only be produced as a result of an individual laser event, but the amorphous region thus produced must remain amorphous throughout subsequent heat cycling needed for overlap in multi track treatment. Fe-B alloys between 5 and 24 at%B were found to be amorphous following laser-alloying (Lin and Spaepen 1982). In this work, Nd : YAG laser melted the alloy for 30 picoseconds (ps) and it was quenched at the high cooling rate of about 101~ K/s. The formation of amorphous phases from pure metallic melts by short laser pulses has also been reported. Frequency-quadrupled Nd-laser pulses (wavelength: 0.265 #m) of 10 ns duration incident on single-crystal Si formed amorphous zones a few tens of nm thick (Tsu et al. 1979). Liu et al. (1979) formed amorphous Si by using 30 ps pulses of both 0.53 and 0.265 #m radiation and the thickness of the zone was estimated to be about 30 nm. Amorphization kinetics of Si prepared by short laser pulses were discussed by Cullis et al. (1982) and by Rozgonyi et al. (1982). In the case of pure A1, the reported amorphous region contained small grains of crystalline A1N (Mazzoldi et al. 1980). Recently, amorphization of Co-Ti thin films was achieved after quenching following melting by a 5 ns Nd: YAG laser pulse (Vitta et al. 1994). The Co-Ti system is suitable for the study of amorphous phase formation at compound-forming compositions since the compounds exhibit a variety of homogeneity ranges and structures. The competition between the formation of an amorphous phase and the terminal solid solution was investigated in this work. The solidification was partitionless and complete amorphous phase formation was found for Ti contents greater than about 23 at% in laser quenched sampies. This compositional limit of the amorphous phase formation agrees well with the prediction based on the influence of atomic size difference on the topographical stability of terminal solid solution that was proposed by Egami and Waseda (1984). Amorphous alloy films were deposited by EXCIMER laser ablation using sintered Ta-Ni targets (Yano et al. 1996). By irradiating a solid target with a pulsed laser-beam of high energy density, atoms and ions forming plumes are emitted from the target, and deposited on a substrate as an amorphous phase. These Ta-Ni amorphous alloy films are known to have extremely high corrosion resistance in both acid and alkaline solutions (Mitsuhashi et al. 1987, Lee et al. 1995). The pulsed laser deposition was performed with
Laser Processing
99
a KrF EXCIMER laser (Lamda Physics LPX305i) having a wavelength and beam dimensions 248 nm and 15 mm x 35 mm, respectively. In another case, the effect of nitrogen on the pulsed laser deposition of amorphous silicon carbide films was investigated to evaluate the properties and structure (Yee et al. 1996). The mechanical, optical, and electrical properties of amorphous carbon-based films (Kurokawa et al. 1987, Fraunheim et al. 1989, Elhossary et al. 1990, Liu and Cohen 1990, Lu et al. 1992, Schwan et al. 1994) can be controlled by focusing laser pulse from ArF EXCIMER laser (Questek 2740SC) operating at 193 nm onto a target. CO2 laser surface treatment consisting of laser surface melting and subsequent selfquenching is also useful to form amorphous phases (Bergmann and Mordike 1981, Kumagai et al. 1986, Hirose et al. 1992, Hirose et al. 1994). An amorphous surface has a variety of useful characteristics such as extremely high corrosion resistance and unique electrocatalytic activity. Hirose et al. (1994) fabricated an amorphous magnetic alloy layer by CO2 laser quenching. The laser-quenched Fe78B 13Si9 alloy layer consisted of an amorphous phase and was utilized to construct a non-contact-type static torque sensor with good sensitivity and linearity.
5. 4.2 Formation o f crystalline p h a s e s
Laser melting and quenching is used not only for the formation of amorphous phases but also for understanding the solidification process. During CO2 laser surface resolidification of Ni3 (A1, Ti) alloys, a sharp transition was observed with increasing growth velocity from the growth of an ordered intermetallic to metastable structures (Hunziker and Kurz 1997). This transition is attributed to disorder trapped in the ordered phase and the corresponding undercooling. High-intensity laser irradiation was used for melting silicon carbide in a high-pressure inert gas atmosphere (Sadler and Shamsuzzoha 1997). By this technique, an entirely new class of materials such as cast silicon carbide shapes and silicon carbide fibers spun or blown directly from the melt might be possible. Badan et al. (1985) formed cellular or dendritic layers, with carbon and chromium supersaturation and metastable austenite as the main constituent on the steel surface by CW-CO2 laser. In the case of YAG laser processing, non-crystalline Nb40Ni60 was formed (Lin and Spaepen 1986) and the morphology of nanostructures on the scale of 1-3 nm could be altered by the laser quenching process (Scandella et al. 1991). Using KrF (248 nm) EXCIMER laser of relatively long (45 ns full width half maximum (FWHM)) pulse duration, melting and subsequent solidification of thick (190410 nm) amorphous-silicon (a-Si) layers produced by ion implantation was investigated (Lowndes et al. 1987). "Fine grained polycrystalline silicon" region produced by explosive crystallization of a-Si contained large numbers of disk-shaped silicon flakes, and the onset of epitaxial regrowth was marked by a slight decrease in surface melt duration. KrF EXCIMER pulsed laser deposition (PLD) has been employed successfully to grow high-quality films that are fully dense, long lasting, and have low friction surfaces. Film properties such as crystallinity, stoichiometry, morphology and dopant concentration can easily be adjusted by varying PLD parameters, since it is exactly these properties that are important for determining the tribological behavior. The feasibility of using PLD to grow
100
K. F. Kobayashi
WS2 films has been evaluated (Prasad et al. 1992). PLD was used to grow adherent films of WS2 on 440C stainless steel substrates (Zabinski et al. 1994). EXCIMER laser is quite a useful tool for the synthesis of electronic materials. A pulsed laser deposition technique formed strained dielectric superlattices of BaTiO3 (BTO) and SrTiO3 (STO) (Tabata et al. 1994). A large stress of 400-500 MPa is introduced at the interface between the BTO and STO layers. A large dielectric constant of 900 was observed with a stacking periodicity of 2 unit cells/2 unit cells. The superlattices show drastically different electrical behavior from that of the solid solution (Sr, Ba)TiO3 films. Dielectric superlattices of (Sr, Ca)TiO3/(Ba, Sr)TiO3 also formed by a pulsed EXCIMER laser ablation (Tabata and Kawai 1997a). The crystal structure is controlled with atomic order accuracy and a large lattice stress of 0.5-1 GPa can be introduced periodically at the interfaces owing to the lattice mismatch of the constituent layers. For the multi-layered films, a laser ablation technique produced a new concept device with a combination of ferroelectric and ferromagnetic materials (Tabata et al. 1997b). An ideal hetero-epitaxy can be obtained owing to the similar crystal structure of perovskite type ferroelectric Pb(Zr, Ti)O3; (so called PZT) and ferromagnetic (La, Sr)MnO3, which is useful for application in DRAM (dynamic random access memory) and SRAM (static random access memory) devices. For the superconducting materials, a laser ablation technique using ArF EXCIMER laser formed Bi-based layered perovskite oxide films (SrBi2Ta207) on the Nb-SrTiO3 substrate or Bi2Sr2CuO6 superconducting bottom electrode with c-axis perpendicular to the substrate surface (Tabata et al. 1995). A multi-target, laser-ablation technique inserted layers of undoped, semiconducting La2CuO4 or SmzCuO4 in a controlled manner in the Lal.ssSr0.15CuO(LSCO) compound superlattices (Tabata et al. 1993). 5.5. L A S E R S U R F A C E - A L L O Y I N G
AND CLADDING
Laser surface treatment is a material-processing method that utilizes the high power density (CO2 laser) available from focused laser sources to melt metal coatings and a portion of the underlying substrate. Laser cladding uses a laser-beam to melt and alloy powders of different composition placed on top of the substrate as well as a thin layer of the substrate. The main purpose of laser cladding is to form a thin interfacial layer of an alloy on a given substrate with minimum dilution of the clad layer and to impart good surface wear and corrosion resistance to the material. The advantages of this technique are production of noble alloys, minimized clad dilution, reduced alloy material loss, reduced machining costs, and reduced distortion. 5.5.1 Ferrous-alloys
A typical use of laser surface treatment is to modify the surface structure of gray cast iron, with the aim of producing a white ledeburitic microstructure while retaining the gray core (Bergmann 1985). The advantage of such a treatment is the combination of good core ductility, a result of stable Fe-C solidification producing an iron-graphite microstructure, and surface wear resistance resulting from the Fe-Fe3C layer obtained in a metastable state.
Laser Processing
101
The phase analysis of the metastable structure was investigated in a laser melted Fe-C alloy (Zarubova et al. 1996). As the laser transformation hardening is characterized by extremely fast heating and self-quenching rates, the presence of alloying elements such as Cr and Mn in steels, that promote hardenability and tend to partition to the carbides in the initial microstructures, is likely to prevent complete austenitization and development of a nearly fully martensitic hardened layer (Bradley and Kim 1988). To improve wear resistance of the AISI 304L austenitic stainless steel, without degrading the inherent corrosion resistance, laser surface treatment was utilized. For this purpose, molybdenum and tantalum were used as alloying elements and they formed a substitutional solid solution in the austenite phase (Akgun and Inal 1995). A 10 kW CO2 laser clad Fe-Cr-Mn-C alloy on an AISI 1016 steel produced a very fine grained microstructure of ferrite and complex carbide precipitation (Singh and Mazumder 1987a). The clad exhibited far superior wear properties compared to Stellite 6 during block-on-cylinder tests. The improved wear resistance was attributed to the fine distribution of metastable M6C carbides in the ferrite matrix. Laser clad Ni-Fe-Cr-A1Hf alloys showed a high degree of grain refinement, considerable increase in solubility of Hf in matrix and the occurrence of Hf-rich precipitates and new metastable phases which contributed towards an increase in the oxidation resistance of the alloy (Singh and Mazumder 1987b). A 10 kW CO2 laser-clad Ni-Cr-A1-Y alloy with either niobium carbide (NbC) or chromium carbides (Cr3C2, Cr7C3, Cr23C6) had higher wear resistance than that of the Ni-Cr-A1-Y alloy cladding without the carbides (Hirose and Kobayashi 1995). Laser cladding of Ni-Cr and Ni-P alloys on steel enhanced the corrosion resistance (Renaud et al. 1991, Garcia-Alonso et al. 1996).
5.5.2 Non-ferrous alloys
Aluminum, titanium, and copper alloys are the most popular materials, but their low wear resistance prevents their widespread applications. A titanium aluminide layer was formed on an aluminum surface by laser processing. By applying a laser-beam to a powder mixture of aluminum and titanium previously placed on the aluminum base metal, the mixture was superficially melted and titanium aluminide alloy layers with a thickness of 400 # m were formed at a relatively high speed. By forming this titanium aluminide alloy layer on an aluminum surface, the wear properties of aluminum have been considerably improved without losing the advantage of lightness of this material (Uenishi et al. 1992a, Uenishi et al. 1992b). A13Ti-based intermetallic compound matrix composite (IMC) layers were formed by using a mixture of aluminum, titanium and ceramic powders as preplaced powders (Uenishi and Kobayashi 1993, Uenishi and Kobayashi 1994, Uenishi and Kobayashi 1996a, Uenishi and Kobayashi 1996b). These IMC layers exhibited good wear resistance. The same technique can be applied to modify the surface of titanium (Hirose et al. 1993). The intermetallic compound clad layers were successfully formed on titanium substrate. A typical example of the clad layers is shown in Fig. 5.3. The clad layers showed good wear and oxidation resistance.
K. E Kobayashi
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Surface modification of copper using a laser is suggested to improve the strength and wear resistance of pure copper at elevated temperatures without sacrificing its electrical properties. In particular, Cu-Cr alloyed layers were produced on a pure copper substrate by irradiating a laser-beam onto chromium powder beds or plasma spray-deposited on the substrate surface (Hirose and Kobayashi 1994).
5.6. L A S E R - A N N E A L I N G
Laser-annealing is characterized by the following advantages in comparison with conventional processes such as furnace annealing" 9 Chemical cleanliness,
Laser Processing
103
9 Extremely short duration of heating cycle and, therefore, high heating and cooling rates, 9 Less thermal damage and distortion, and 9 Controlled shape and location of the treated region. Because of its selective heating of the workpiece, laser-annealing is promising for heat treatment of semiconductor materials as an alternative to furnace-annealing in the microelectronics industry. In conventional IC (integrated circuit) fabrication, the lattice damage caused by ion implantation is removed and the implanted dopants are activated by furnace annealing for at least 30 min at temperatures of 820 K to 1470 K. Laser annealing can more rapidly and effectively remove the lattice damage and activate the dopants. Although significant damage usually remains in the form of dislocation loops after furnace annealing, pulsed-laser annealed specimens have no damage down to a resolution of approximately 1 nm (Young and Wood 1982, Young et al. 1982). Figure 5.4 shows the carrier concentration as a function of implanted dose for pulsed laser and thermally annealed boron-implanted silicon (Young and Wood 1982). In laser-annealed silicon the carrier concentration increases linearly beyond the dose of 1.5 • 1016/cm2 that corresponds to the equilibrium solubility limit of boron in silicon in this case, whereas it saturates for thermal annealing. It is clearly shown in this figure that the non-equilibrium process of laser annealing can allow active carrier concentration in excess of the equilibrium solubility limit. This is achieved due to the extremely high heating and cooling rates in the laser processing. In laser-annealing, the very short duration time and the ease of controlling the shape and the location of the treating region are also beneficial to form shallow junctions, which are needed for highly integrated devices such as megabit SRAMs (Tsukamoto et al. 1992). In conventional furnace annealing or rapid thermal annealing systems, shallow p + n junction formation using boron ion implantation is difficult, because the implanted boron diffuses very rapidly into silicon on subsequent annealing (Tsukamoto et al. 1992). An ultrashallow p + n junction, less than 50 nm deep, has been formed by low power BF2 ion implantation followed by EXCIMER laser-annealing with single pulse irradiation (Tsukamoto et al. 1992, Tsukamoto et al. 1996). Laser-annealing has also been applied to recrystallize amorphous layers of materials to improve their electrical properties. Polycrystalline silicon thin-film transistors (polySi TFTs) are used for large-area electronics applications such as liquid crystal displays and image sensors, because poly-Si TFTs have higher carrier mobility than amorphous silicon (a-Si) TFTs. EXCIMER laser annealing has been applied to the fabrication of highperformance poly-Si TFTs on non-heat-tolerant glass substrates, because its extremely short duration and shallow melt-regrowth avoids thermal damage to the glass substrate (Sameshima et al. 1986, Sera et al. 1989, Kuriyama et al. 1991). Recently large-grain poly-Si TFTs have been obtained by controlling the solidification velocity during dualEXCIMER laser irradiation both from the front side and from the back side (Shimizu et al. 1993) and by low temperature (<673 K) substrate heating during EXCIMER laser annealing (Kuriyama et al. 1992, Kuriyama et al. 1993). Besides the microelectronics applications, annealing of amorphous materials by laserbeams has been applied to improve the magnetic properties of materials (Matteazzi et al.
K. F. Kobayashi
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1985, Sato et al. 1985, Lanotte and Lannotti 1995). Due to the very short duration time and rapid temperature change, laser-heating can produce uniform crystallization at a temperature that is much higher than that reached by other methods (Lanotte and Lannotti 1995). This is very useful for the synthesis of nanocrystalline materials through crystallization of amorphous materials. It has been known that nanocrystalline ferromagnetic materials have improved soft magnetic properties in comparison with crystalline or amorphous materials (Herzer 1992). Pulsed YAG laser irradiation of thick Fe80.5(BSiC)19.5 amorphous ribbons (>50 mm), which is a transformer core material, has decreased the core losses because of refinement of the domain widths (Matteazzi et al. 1985). CW CO2 laser annealing of Fe73.sCulNb3Si13.sB9 amorphous ribbons produced nanocrystalline structure of 50 nm grain size (Lanotte and Lannotti 1995). The laser annealed materials had good soft magnetic properties (Lanotte and Lannotti 1995). It has also been reported that continuous laser-annealing improved the soft magnetic properties of Fe40Ni40P14B6 amorphous ribbons (Sato et al. 1985). Selective and local annealing by laser-beam can be applied to structural materials. Laser irradiation for solution-annealing (laser solution-annealing) has been performed on the surface of a Ni-base superalloy, Inconel 718, to reduce its sensitivity to hydrogen embrittlement (Hirose et al. 1996). Inconel 718, which is one of the most used Ni-base superalloys, is age-hardened by the homogeneous precipitation of yt and y~f phases, which allows the alloy to have a high strength at moderate temperatures, and good creep and fa-
105
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tigue resistance. However, the high sensitivity of Inconel 718 to hydrogen embrittlement is a serious problem in using the alloy, for example, in a liquid hydrogen fueled rocket engine. Hydrogen-induced cracking tends to occur at the surface of stress concentrated regions. Reducing the hardness of such surface regions locally is known to be effective in preventing such cracking. To this end, laser surface solutionizing has been applied to Inconel 718 using a 2.5 kW CO2 laser. The surface-solutionized zones can be obtained without melting the treated zones by optimizing the laser conditions. While the base metal has a hardness of approximately 450 Hv, the hardness of the laser solutionized zones is reduced to below 250 Hv as shown in Fig. 5.5. No precipitates of yl and y1~ phases are identified in the laser solutionized zones by TEM observations. The surface regions are heated beyond the solvus temperatures of y~ and y~ phases and these phases are dissolved in the matrix. The subsequent rapid cooling can prevent the precipitation of the age hardenable phases. Surface-solutionizing can be achieved by laser-scanning. The desired depth of the solutionized zones can be obtained by controlling the laser parameters, traverse speed and defocus distance, as shown in Fig. 5.6. The ductility of the surface solutionized specimen is almost twice that of the base metal in a tensile test under a 29.4 MPa hydrogen atmosphere at room temperature with hydrogen pre-charging. Thus, the laser surface softening is effective in reducing the sensitivity of Inconel 718 to hydrogen embrittlement with little sacrifice to the strength.
106
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Figure 5.6. Effectof traverse speed on depth of the laser surface solutionized zone for different defocus distances (Hirose et al. 1996).
5.7. LASER-BEAM JOINING Joining is one of the most important applications of lasers including both solid-state and gas lasers, ranging from joining and bonding devices in microelectronics to welding structural materials. Laser-beam joining is characterized by its high energy density and flexibility, and has obvious advantages, such as high speed, narrow fusion and heat affected zones (minimizing distortion and damage of adjacent areas), ability of remote noncontact processing, and ease of automation. 5. 7.1 Welding o f structural materials 5. 7.1.1 Ferrous-based alloys Laser-beam welding is an attractive process for joining steels due to its high productivity and possibility of automation. In particular, automotive and shipbuilding industries have an interest in the applications of laser-beam welding. Laser-beam weld fusion and heat affected zones are subjected to significantly more rapid heating and cooling cycles compared to conventional arc welds. In the case of laser-beam welding of C-Mn steels, HY (high yield strength) steels and HSLA (high strength low alloy) steels, metastable phases such as martensite and bainite having high hardness are produced in the weld fusion and heat affected zones (Metzbower and Denney 1986, Moon et al. 1987, Metzbower 1990). Thus, these laser-beam welded zones have higher hardness in comparison with arc welded zones. Therefore, predictions of hardness of laser-beam welds are desired. Ion et al. have developed simplified analytical models of heat flow and phase transformations, and applied them to laser-beam welding of C-Mn steels to predict the maximum heat affected zone hardness (Ion et al. 1996). In this model the heat cycles in the heat affected
107
Laser Processing
Cooling time from 1073K to 773K o~" 0.30 W W r
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Figure 5.7. Diagram of carbon equivalent against absorbed energy showing theoretical contours for maximum heat affected zone hardness in carbon manganese steels (Ion et al. 1996).
zones have been simulated with a formula based on the thin plate analysis of Rosenthal (1946). Microstructural constituents in the heat affected zones have been estimated from the simple carbon equivalent, Ceq = C + Mn/12 + Si/24 (where the concentrations of the different elements are given in wt.%), and cooling time between 1073 K and 773 K using empirical formulas obtained by analyzing experimental data. The hardness of individual phases, namely martensite, bainite and ferrite/pearlite, has been empirically obtained from the carbon equivalent. Finally the maximum heat affected zone hardness has been determined by summing the contributions from the individual phases using a simple rule of mixtures. Figure 5.7 shows the diagram for predicting the maximum heat affected zone hardness from the carbon equivalent and absorbed energy using this model. In laser-beam welding of stainless steels the microstructural changes in weld fusion and heat affected zones due to rapid cooling, which affect both the mechanical and corrosion properties, should be considered. In conventional arc welding, the room-temperature microstructures of weld-bead are estimated using Schaeffler diagram (Schaeffler 1949). However, the solidification microstructures of laser welds subjected to more rapid cooling and higher solidification rate than arc welds are significantly different from those estimated by Schaeffler diagram (Vitek et al. 1982, David and Vitek 1985, David et al. 1987, Khan et al. 1988, Zambon and Bonollo 1994). It has been shown that a significant narrowing of the duplex ferrite-austenite region in the diagram occurs with cooling rates typical of laser-beam welding as schematically shown in Fig. 5.8 (David et al. 1987). Thus, while
K. F. Kobayashi
108 A
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Ferrite (F)
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32 36 40 (mass%)
Figure 5.8. Schaeffler diagram showing schematically the convergence of the 0 and 100% ferrite lines in laser-beam welding (Davidet al. 1987).
austenitic stainless steels such as type 304, 308 and 316 have two-phase ferrite (3) plus austenite (y) microstructures that are found after primary ferrite solidification under normal welding conditions, these steels often have a fully austenitic microstructure under rapid solidification conditions such as laser-beam welding. This is caused by a change in the solidification mode from the equilibrium primary ferrite solidification mode to a non-equilibrium primary austenite solidification mode due to rapid solidification (Vitek and David 1994). 5. 7.1.2 Non-ferrous alloys The low overall heat input and, thereby, the small fusion and heat affected zones makes laser-beam joining suitable for welding age-hardenable alloys, which have a problem of softening in welded zones, in comparison with conventional arc welding processes. The characteristics of laser-beam welds of age-hardened A1 and Ni base alloys will be hereinafter presented in comparison with their tungsten inert gas (TIG) welds. Age-hardened A1-Mg-Si alloy, 606 l-T6, thin plates (1 mm in thickness) have been welded with full penetration using a 2.5 kW CO2 laser (Hirose et al. 1997b). In laser-beam welding a full penetration weld can be achieved even at a welding speed of 167 mm/s, whereas in TIG welding a welding speed of 5 mm/s or less is required for the full penetration weld. The cross section of a typical laser-beam weld is shown in Fig. 5.9. The bead width of the laser-beam weld is approximately 1/3 that of the TIG weld. Figure 5.10 shows the hardness profiles obtained from the laser-beam and TIG welds in the as-welded condition and after an artificial aging treatment at 448 K for 28.8 ks. In the as-welded condition, although a softened region is found in both the laser-beam weld and
