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
60
///~...
" 10.4 MPa
> .m =
E 0
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40
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20 I
1
I
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I
I
10 Particle diameter (It m)
I
I
I
I11
100
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
I!
II electron beam II II
[1--
refined zo
portion refined Dy remeltincj
workpiece
\~,~,b~6."~('e" k
.
)
3.4. Principle of rapid solidification at the surface of a block of material following local melting with a traversing heat source (Lux and Hiller 1972).
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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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 . %
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,
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
=
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~2 102 l.Lm 10 I
0.1 I
0.01
10"s
I
I
.
I
!
I
I
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"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
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x
Chadwick
196
9 Cooksey .tal
01
-
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Elliott
+ L i v i n g s t o n et al
_
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. 1967
1968
13 B u r d e n & Jones 1970 -
9 Zimmermann I r ill
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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
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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
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10-6-
eo
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1
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I
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m
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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
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9
20
.
~o
9
0
~: 113 ~
O ~
9O
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/rr
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7075-T651
E lOO
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!
80
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9
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.
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c: ~ I-.
l
/
30 20
60
/
c
V)
J
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40 3O 8009
20 -
10
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----
9 !
70 -
IJJ
CD
2O N E 10
0
so
c
o
~
.~
60
c ID 3O J--.
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-
.~
5O m
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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.
REFERENCES
Abraham, T. (1991) Internat. J. Powder Metall., 27, 161. Alonso, J. A., Gallego, L. J. and Lopez, J. M. (1988) Phil. Mag., A58, 79. Amano, N., Odani, Y., Takeda, Y. and Akechi, K. (1989) Metal Powder Rept., 44, 186. Anderson, I. E. (1991) Adv. Mater. Process., 140(1), 30. Anon (1989) Adv. Mater. Process., 135(6), 12. Ashley, S. (1998) Mech. Eng., 120(6), 72. Boettinger, W. J. (1982) in Rapidly Quenched Metals, Vol. I, eds. Masumoto, T. and Suzuki, K. (The Japan Institute of Metals, Sendai, Japan), p. 85. Boettinger, W. J., Shechtman, D., Schaefer, R. J. and Biancaniello, E S. (1984) Metall. Trans., A15, 55. Boettinger, W. J., Bendersky,L. A., Coriell, S. R., Schaefer, R. J. and Biancaniello, E S. (1987) J. Cryst. Growth, 80, 17. Boswell, E G. and Chadwick, G. A. (1977) Scripta Metall., 11, 459. Boulby, K. A. and Wood, J. V. (1986a) Powder Metall., 29, 33. Boulby, K. A. and Wood, J. V. (1986b) Metal Powder Rept., 41, 299. Davies, H. A., Aucote, J. and Hull, J. B. (1974) Scripta Metall., 8, 1179. Davis, L. A. (1976) in Rapidly Quenched Metals, eds. Grant, N. J. and Giessen, B. C. (MIT Press, Cambridge, MA), p. 369. Duwez, P. (1967) Trans. Amer. Soc. Metals, 60, 607. Edgington, J. (1977) in Fibre-Reinforced Materials (Institute of Civil Engineers, London), p. 129. Feurer, U. and Wunderlin, R. (1977) DGM Fachbereicht, No. 38. Gelinas, C., Angers, R. and Pelletier, S. (1988) Mater. Lett., 6, 359. Gibeling, J. C. and Nix, W. D. (1978) Scripta Metall., 12, 919. Gilgien, E, Zryd, A. and Kurz, W. (1995) Acta Metall. Mater., 43, 3477. Gilman, E (1990) Metals and Materials, 6, 504. Grant, N. J. (1970) Fizika, 2 Suppl. 2, 16.1.
44
H. J o n e s
Greer, A. L. (1994) Nature, 368, 688. Guthrie, M. S. (1993) in Rapidly Solidified Alloys: Processes, Structures, Properties, Applications, ed. Liebermann, H. H. (Marcel Dekker, New York), p. 665. Hashimoto, K. (1997) Mater. Sci. Eng. A, A226/228, 891. Herlach, D. M. (1994) Mater. Sci. Eng. Rept., R12, 177. Herlach, D. M., Gillesen, E, Volkmann, T., Wollgarten, M. and Urban, K. (1992) Phys. Rev. B, 46, 5203. Herlach, D. M., Cochrane, R. E, Egry, I., Fecht, H. J. and Greer, A. L. (1993) Internat. Mater. Rev., 38, 273. Holbrook, A. L. (1986) Internat. J. Powder Metall., 22, 39. Hunt, J. D. (1979) in Solidification and Casting of Metals (The Metals Society, London), p. 3. Hunt, J. D. and Lu, S.-Z. (1996) Metall. Trans. A, 27A, 611. Ikawa, Y., Itami, T., Kumagai, K., Kawashima, Y., Leatham, A. G., Coombs, J. S. and Brooks, R. G. (1990) ISIJ Internat., 30, 756. Inoue, A., Zhang, T. and Masumoto, T. (1993) J. Non-Cryst. Solids, 156-158, 473. Jackson, K. A. and Hunt, J. D. (1966) Trans. Metall. Soc. AIME, 236, 1129. Johnson, W. L. (1996) Curr. Opin. Solid State & Mater. Sci., 1, 383. Johnson, W., Mamalis, A. G. and Hunt, H. (1976) Metallurgia and Metal Forming, 43(3), 68. Jones, H. (1984) J. Mater. Sci., 19, 1043. Jones, H. (1988) Mater. Lett., 6, 181. Jones, H. (1991) in Proceedings of the Second Japan International SAMPE Symposium, eds. Kimpara, I., Kageyama, K. and Kagawa, Y., SAMPE, p. 306. Juarez-Islas, J. A. and Jones, H. (1987) Acta Metall., 35, 499. Kawashima, A., Habazaki, H., Asami, K. and Hashimoto, K. (1996) Sci. Rept. RITU A, A42, 91. Kim, W. T. and Cantor, B. (1994) Acta Metall. Mater., 42, 3045. Kim, Y. H., Inoue, A. and Masumoto, T. (1990) Mater. Trans. JIM, 31,747. Kui, H. W., Greer, A. L. and Tumbull, D. (1984) Appl. Phys. Lett., 45, 615. Kurz, W. and Fisher, D. J. (1981) Acta Metall., 29, 11. Kurz, W., Giovanola, B. and Trivedi, R. (1986) Acta Metall., 34, 823. Kurz, W., Giovanola, B. and Trivedi, R. (1988) J. Cryst. Growth, 91, 123. Lu, S.-Z., Hunt, J. D., Gilgien, P. and Kurz, W. (1994) Acta Metall. Mater., 42, 1653. Lux, B. and Hiller, W. (1972) Praktische Metallographie, 8, 218. Maringer, R. E. and Mobley, C. E. (1974) J. Vac. Sci. Technol., 11, 1067. Massalski, T. B., Kim, Y. W., Vassamillet, L. E and Hopper, R. W. (1981) Mater. Sci. Eng., 47, 1. Midson, S. E and Jones, H. (1982) in Rapidly Quenched Metals, Vol. II, eds. Masumoto, T. and Suzuki, K. (The Japan Institute of Metals, Sendai, Japan), p. 1539. Moir, S. A. and Jones, H. (1991) Mater. Lett., 12, 141. Nagahuma, H., Ohtera, K., Higashi, K., Inoue, A. and Masumoto, T. (1993) Phil. Mag. Lett., 67, 225. Oguchi, M., Inoue, A., Yamaguchi, H. and Masumoto, T. (1990) Mater. Trans. JIM, 31, 1005. Ohnaka, I. (1985) Internat. J. Rapid Solidification, 1, 219. Overshott, K. J. (1979) Electron. Power, 25, 347. Peker, A. and Johnson, W. L. (1993) Appl. Phys. Lett., 63, 2342. Pelletier, S., Gelinas, C. and Angers, R. (1990) lnternat. J. Powder Metall., 26, 51. Perepezko, J. H. (1984) Mater. Sci. Eng., 65, 125. Prasad, Y. V. S. S., Ramachandrarao, E and Chattopadhyay, K. (1984) Acta Metall., 32, 1825. Rabinkin, A. and Liebermann, H. H. (1993) in Rapidly Solidified Alloys: Processes, Structures, Properties, Applications, ed. Liebermann, H. H. (Marcel Dekker, New York), p. 691. Raman, R. V., Patel, A. N. and Carbonara, R. S. (1982) Prog. Powder Metall., 38, 99. Rama Rao, E and Radhakrishnan, V. M. (1984) in Metallic Glasses, ed. Anantharaman, T. R. (TransTech, Aedermannsdorf, Switzerland), p. 225. Schmitt, H. (1979) Powder Met. Internat., 11, 17. Smith, C. H. (1993) in Rapidly Solidified Alloys: Processes, Structures, Properties, Applications, ed. Liebermann, H. H. (Marcel Dekker, New York), p. 617. Spaepen, E (1987) in Undercooled Alloy Phases, eds. Collings, E. W. and Koch, C. C. (TMS, Warrendale, PA), p. 187. Vero, B., Fauszt, A., Tardy, E, Takacs, J., Konczos, G. and Lovas, T. (1992) Mater. Sci. Technol., 8, 645. von Allmen, M., Huber, M., Blatter, A. and Affolter, K. (1984) Internat. J. Rapid Solidification, 1, 15. Wang, K. K., Gustavson, L. J. and Dumbleton, J. H. (1988) Mod. Dev. Powder Metall., 20, 361.
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.
This Page Intentionally Left Blank
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
This Page Intentionally Left Blank
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
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 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.
Mechanical Alloying
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.
REFERENCES
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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
-Iw w
=
300
r9
250
I,.,
o
>
O
350
"o w
O
200
O
9
9 ,
9 1
200
,
9 .
.
0
9 I
400
,
,
,
I
I
I
600
Distance from surface
I
I
I
I
800
•
1000
(llm)
Figure 5.5. Hardness distribution across a laser surface solutionized zone of Ni-base superalloy, Inconel 718 (Hirose et al. 1996).
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
K. F. Kobayashi . . . .
~:
I
. . . .
I
. . . .
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. . . .
1
. . . .
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1.5
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,,=0.5
,e.,m
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0
5
10 15 20 25 30 Traverse speed (mm/s)
35
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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i
10
,,
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20 30 40 50 60 70 Absorbed energy (J/mm 2)
80
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
28
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._+ o" 4)
\
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/
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Ferrite (F)
"
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 the TIG weld, the width of the softened region in the laser-beam weld is less than 1/4 that of the TIG weld. The softened region is caused by reversion (dissolution) of strengthening/~
Laser Processing .
109
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Figure 5.9. Cross-sectionof a laser-beam weld of I mm thick 6061-T6 A1 alloy using 2.5 kW CO2 laser.
(Mg2Si) precipitates due to the heat input of welding. The hardness of the softened region in the laser-beam weld is almost fully recovered to the level of the base alloy after an aging treatment at 448 K for 28.8 ks. In contrast to this, the TIG weld subjected to the aging treatment after welding has an appreciable hardness drop in a region between 4 and 7 mm from the bead center. This region is thought to be the region where the solute (Mg and Si) content in the matrix is insufficient for precipitation of the strengthening fl (MgzSi) phase because of precipitation of nonhardening fif and/or fl1~ phases during weld thermal cycle. Therefore, the TIG weld will require a solution annealing treatment before artificial aging to fully recover the heat-affected zone hardness unlike the laser-beam welds. In laser-beam welding, since the cooling rate in the heat affected zone is rapid enough to prevent the precipitation of fl~ and/or fl~ phases, supersaturated solid solutions can be obtained over the whole region of the heat affected zone, and thereby sufficient age-hardening due to precipitation of fi1~ phase can be achieved after the artificial aging. As expected from the hardness profile in Fig. 5.10, the tensile strength of the specimen subjected to the aging treatment after laser-beam welding is almost equivalent to that of the base plate. Laser-beam welding has been applied to joining an age-hardened Ni base superalloy, Inconel 718, plate (3.2 mm in thickness) using a 2.5 kW CO2 laser (Hirose et al. 1997a, Hirose et al. 1998) and its properties have been compared with those of TIG welds. Inconel 718 is strengthened by intermetallic ~,~ and V ~ phases, which are precipitated after solution annealing treatment followed by a two step aging treatment. Microstructures of the weld fusion zones are shown in Fig. 5.11. The TIG weld fusion zone has a cellular dendritic structure having austenite (Vf)/Nb-rich Laves phase eutectic and Nb-rich MCtype carbides in the interdendritic regions (Fig. 5.11 (a)). The laser-beam weld fusion zone
110
K. F. Kobayashi
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~
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,
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.After aging
,
,
,
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12
Figure 5.10. Hardness profiles obtained from laser-beam and TIG welds of 1 mm thick 6061-T6 AI alloy plates in as-welded condition and after aging at 448 K for 28.8 ks.
has significantly finer solidification structure than that of the TIG weld (Fig. 5.1 l(b)). In the laser-beam welding, the cooling rate during solidification that is estimated from the relationship between dendrite arm spacing and cooling rate is 100 times greater than that in the TIG welding. Thus, the high energy concentration of the laser-beam welding results in rapid solidification and thereby fine solidification structure. Figure 5.12 shows the hardness distributions obtained from the laser-beam and TIG welds in the as-welded condition and after the two-step aging treatment. The fusion zone width of the laser-beam weld is approximately 1/6 of that of the TIG weld. Apparent softening occurs in the weld fusion and heat-affected zones because of a reversion (dissolution) of strengthening y' and y" phases. However, the laser-beam weld has a significantly narrower softened region that is approximately 1/7 of that of the TIG weld. In the laser-beam welds, the hardness of the fusion and heat-affected zones is fully recovered to that of the aged base alloy after the aging treatment. In the TIG welds, the hardness of the weld fusion zones after the aging treatment is not completely recovered to the level of the base metal. Since in the TIG weld fusion zones coarse Nb-rich phases are formed in interdendritic regions, the solute concentrations of the y' matrix may be insufficient for full precipitation of the y' and y" phases. These hardness distributions apparently reveal an advantage of laser-beam welding that its low overall heat input results in fine solidification structure and minimize the weld fusion and heat affected zones. The tensile strength of the laser-beam weld joint subjected to the aging treatment after welding fully recovers to that of the aged base alloy, whereas that of the TIG weld joint does not because of the coarse interdendritic phases and lower hardness in the weld fusion zone.
Laser Processing
111
Figure 5.11. Microstructures of weld fusion zones of Ni-base superalloy, Incone1718. (a) TIG welding and (B) CO2 laser-beam welding.
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112
K. E Kobayashi
Narrow weld-fusion and heat-affected zones, and minimizing distortion in laser-beam welding, are also beneficial to fusion welding, of brittle materials such as intermetallic compounds. Laser-beam welding of a molybdenum-containing TiA1 intermetallic compound (Ti-46at%A1-2at%Mo) has been performed using a 2.5 kW CO2 laser (Hirose et al. 1995a, Hirose et al. 1995b). The cooling rates in the welded zones have been controlled by changing the welding speed and pre-heating temperature. The base metal has the lamellar structure containing y(TiA1) + otz(Ti3A1). However, the microstructure and hardness of the laser fusion zones change depending upon the cooling rates as shown in the schematic CCT diagram (Fig. 5.13). When the cooling rate at 1390 K is above 3000 K/s, the ot ~ y transformation is suppressed and the fusion zone structure results in a single phase of ot2. If the cooling rate is less than 3000 K/s, massive y phase is seen together with the Ot2 phase. In laser-beam welding, where the cooling rates are approximately 80 K/s to 1800 K/s, the fusion-zone microstructure mainly consists of massive ot2, massive y and lamellar (c~2 + y) phases. Thus metastable phases of massive Ot2 and massive y form owing to rapid cooling during laser-beam welding. A fully lamellar structure is formed if the cooling rate is less than 20 K/s. The hardness of the fusion zones increases from 315 Hv of the base metal hardness to above 500 Hv with increasing cooling rate. Cracking in the weld fusion zones is sensitive to the cooling rates and hence to the microstructures of the fusion zones. When the hardness value of the fusion zones is reduced below 410 Hv, cracking dramatically decreases. As a result, crack-free welds can be obtained at welding speeds below 50 mm/s and pre-heating temperatures above 673 K (Fig. 5.14). The cross-section of a crack-free weld bead is shown in Fig. 5.15. During tensile testing, the laser welded specimens without cracking fractured in the base metal. Thus, laser-beam welding with preheating can provide narrow weld fusion and heat-affected zones, and optimal microstructures of the weld fusion zones that are beneficial to welding of brittle materials.
5.7.2 Microjoining
Microjoining in electronics industries also is an important application of laser-beam processing. Laser-beam soldering has been applied to surface-mount technology, namely interconnection between electronic components and printed circuits since 1980's (Chang 1986, Semerad et al. 1993, Messier and Millard 1994). In laser-beam soldering, ease of controlling the shape and location of the heating region brings about reliable solder joints with minimal component heating and is suitable for fine-pitch soldering and high-density packaging. While CO2 laser and Nd:YAG laser are used for laser soldering, Nd: YAG laser is more suitable because of preferential absorption by solders as opposed to board materials (Messier and Millard 1994). Since only the solder joint area is heated during laser-beam soldering process without damaging the boards and components, this process is more beneficial to interconnections using solders with a wide range of melting temperatures in comparison with conventional infrared reflow process (Messier and Millard 1994). Thus, this can overcome the problem of soldering with Pb-free solders with higher melting temperatures than those of eutectic or near eutectic Sn-Pb solders (Yang et al. 1995). Rapid heating and cooling thermal
Laser Processing
113
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A
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=3
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!--
Cf~2 300(K/: OHv
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Figure 5.14. Effects of preheating temperature and welding speed on crack frequency of laser-beam welds of Ti-46at%Al-2at%Mo.
114
K. E Kobayashi
Figure 5.15. Cross-sectionof a crack free laser-beamweld of 2 mm thick Ti-46at%A1-
2at%Mo using 2.5 kW CO2 laser.
cycles in laser-beam soldering result in the microstructures and mechanical properties of the solder joints different from those of the conventional infrared reflowed joints. Laser-beam soldering of eutectic Sn-Ag solder has been performed and the effect of the soldering process on the microstructures and mechanical properties have been evaluated (Yang et al. 1995). In this study eutectic Sn-Ag solder joints between quad-flat pack surface-mounted components and printed wiring boards have been produced using a Nd : YAG laser soldering system. Much lower overall heat input and shorter duration time in laser-beam soldering results in very high cooling rates (~104 K/s) from peak temperature compared with the infrared reflow soldering (~10 K/s). Therefore, in laserbeam solder joints, much finer eutectic Ag3Sn intermetallic in pure tin and thinner Cu-Sn intermetallic layers at copper pad and wire interfaces form compared with the infrared
Laser Processing
115
reflowed joints. The high peak temperature (673 K to 873 K) caused by laser irradiation compared to the infrared reflow (~528 K) dissolves a certain amount of copper from leads and pads. This results in the formation of a considerable number of equiaxed Cu6Sn5 dendrites in the bulk solder. The difference of microstructure between laser-beam soldered and infrared reflowed joints causes higher microhardness of the laser-beam soldered joints (~27 Hv) than that of infrared reflowed joints (~16 Hv), suggesting better mechanical properties under static, creep and fatigue loading. Laser-beam processes have also been applied to bond Tape Automated Bonding (TAB) leads to integrated circuit pads (Spletter and Crowley 1990). Since laser inner lead bonding can be performed using lower temperature and lower pressure compared to conventional gang bonding process, it can minimize the thermal and mechanical stressinduced damages of circuit. TAB laser bonding is performed by positioning a focused laser-beam over a lead and then pulsing the laser once, creating a reliable metallurgical bonding between the plated lead and the bump (Spletter and Crowley 1990). This bonding process takes advantage of a thin tin plating to absorb the laser energy and accelerates melting of the copper lead and gold bump. Since this process achieves rapid alloying of copper and gold, bonds without the brittle intermetallic compounds often found in gold/tin alloys formed with conventional methods can be obtained (Spletter and Crowley 1990).
5.8. C O N C L U S I O N S Lasers are lights that obey straightforward physical laws and make it easy to predict their behavior. Knowing the relevant optical equations and how to apply them will explain most of the features of laser materials processing. For example, the high energy density available from focused sources is utilized to melt only the surface of a material in a very short time. This characteristic can apply to form non-equilibrium materials. In other cases like EXCIMER laser, it is suitable for applications in the field of micromachining. Although laser processing is prominent, the user needs to consider the following problems. What is the best quality that one can achieve and what is the fastest way to get there? What will be new and different are the orders of magnitude differences between the feature sizes and quality. These differences will demand better optics, stable environments, and system design, part handling and inspection systems.
REFERENCES
Akgun, O. V. and Inal, O. T. (1995) J. Mater. Sci., 30, 6097. Aziz, M. J. (1982) J. AppL Phys., 53, 1158. Badan, B., Giordano, L., Ramous, E. and Rudilosso, S. (1985) in Rapidly Quenched Metals V, eds. Steeb, S. and Warlimont, H. (Elsevier Science Pub., Amsterdam),p. 851. Bergmann, H. W. (1985) Surface Eng., 1, 135. Bergmann, H. W. and Mordike, B. L. (1981) J. Mater Sci., 16, 863. Bergmann, H. W., Fritsch, H. U. and Hunger, G. (1981) J. Mater. Sci., 16, 1935. Bimberg, D. (1986) in Laser Treatment of Materials, ed. Mordike, B. L. (DGM Informationsgesellschaft, Bad Nauheim, Germany),p. 3. Bloembergen, N. (1979) in Laser-Solid Interactions and Laser Processing 1978, eds. Ferris, S. D., Leamy, H. J. and Poate, J. M. (Amer. Inst. Phys., New York), p. 1.
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Boettinger, W. J., Biancaniello, E, Kalonji G. M. and Cahn, J. W. (1980) in Rapid Solidification Processing: Principles and Technologies, eds. Mehrabian, R., Kear, B. H. and Cohen, M. (Claitor's Pub., Reston, VA), p. 50. Boettinger, W. J., Coriell, S. R. and Sekerka, R. E (1984) Mater. Sci. Eng., 65, 27. Boettinger, W. J., Coriell, S. R. and Trivedi, R. (1988) in Rapid Solidification Processing: Principles and Technologies, eds. Mehrabian, R. and Paris, P. A. (Claitor's Pub., Reston), p. 13. Bradley, J. R. and Kim, S. (1988) Metall. Trans. A, 19A, 2013. Cahn, J. W. (1956) Acta Metall., 4, 572. Carrad, M., Gremaud, M., Zimmermann, M. and Kurz, W. (1992) Acta Metall. Mater., 40, 983. Chande, T. and Mazumder, J. (1982) Appl. Phys. Lett., 41, 42. Chande, T. and Mazumder, J. (1983) Metall. Trans., 14B, 181. Chang, D. U. (1986) Weld. J., 65, 32. Cline, H. E. and Anthony, T. R. (1977) J. Appl. Phys., 48, 3895. Cullis, A. G., Webber, H. C. and Chew, N. G. (1982) in Laser and Electron-Beam Interraction with Solids, eds. Appleton, B. R. and Celler, G. K. (North-Holland, Amsterdam), p. 131. David, S. A. and Vitek, J. M. (1985) in Rapidly Quenched Metals V, eds. Steeb, S. and Warlimont, H. (Elsevier Science Pub., Amsterdam), p. 847. David, S. A. and Vitek, J. (1989) Internat. Mater. Rev., 34, 23. David, S. A., Vitek, J. M. and Hebble, T. L. (1987) Weld. J., 68, 289-s. Davies, H. A., Aucote, J. and Hull, J. B. (1974) Scripta Metall., 8, 1179. Draper, C. W. and Poate, J. M. (1985) lnternat. Metall. Rev., 30, 85. Egami, T. and Waseda, Y. (1984) J. Non-Cryst. Solids, 64, 113. Elhossary, E M., Fabian, D. J. and Webb, A. P. (1990) Thin Solid Films, 192, 201. Frauenheim, E M., Stephan, U., Bewilogua, K., Jungnivkel, E, Blaudeck, P. and Fromm, E. (1989) Thin Solid Films, 182, 63. Fritsch, H. U. (1986) in Laser Treatment of Materials, ed. Mordike, B. L. (DGM Informationsgesellschaft, Bad Nauheim), p. 19. Garcia-Alonso, M. C., Escudero, M. L., Lopez, V. and Macias, A. (1996) ISIJ lnternat., 36, 172. Gilgien, P., Zryd, A. and Kurz, W. (1995) Acta Metall. Mater., 43, 3477. Gill, S. C. and Kurz, W. (1995) Acta Metall. Mater., 43, 139. Grange, R. A. and Kiefer, J. M. (1941) Trans. ASM, 29, 85. Greenwald, L. E., Breinan, E. M. and Kear, B. H. (1979) in Laser-Solid Interactions and Laser Processing 1978, eds. Ferris, S. D., Leamy, H. J. and Poate, J. M. (Amer. Inst. Phys., New York), p. 189. Herzer, G. (1992) J. Magn. Magn. Mater., 112, 258. Hirose, A. and Kobayashi, K. E (1994) Mater Sci. Eng., A174, 199. Hirose, A. and Kobayashi, K. F. (1995) ISIJ Internat., 35, 757. Hirose, A., Ueda, T. and Kobayashi, K. E (1993) Mater. Sci. Eng.,A160, 143. Hirose, A., Arita, Y. and Kobayashi, K. E (1995a) J. Mater Sci., 30, 970. Hirose, A., Arita, Y. and Kobayashi, K. E (1995b) J. Society Mater Sci. Japan, 44, 1145. Hirose, A., Arita, Y., Nakanishi, Y. and Kobayashi, K. E (1996) Mater Sci. Eng., A219, 71. Hirose, A., Sakata, K. and Kobayashi, K. E (1997a) in Proc. of 4th Decennial Int. Conf. on Solidification Processing, Sheffield, UK, p. 675. Hirose, A., Todaka, H. and Kobayashi, K. E (1997b) Metall. Mater. Trans., 28A, 2657. Hirose, A., Sakata, K. and Kobayashi, K. E (1998) Int. J. Mater. Production Technol. (submitted). Hirose, E, Takagi, M., Moil, H. and Imura, T. (1992) Jpn. J. Appl. Phys., 31, 3940. Hirose, E, Takagi, M., Mori, H. and Imura, T. (1994) Mater Sci. Eng., A181/A182, 1359. Hsu, S. C., Chakravorty, S. and Mehrabian, R. (1978) Metall. Trans., 9B, 221. Hsu, S. C., Kou, S. and Mehrabian, R. (1980) Metall. Trans., liB, 29. Hunziker, O. and Kurz, W. (1997) in Proc. of 4th Decennial Internat. Conf. on Solidification Processing, Sheffield, UK, p. 599. Ion, J. C., Salminen A. S. and Sun, Z. (1996) Weld. J., 75, 225-s. Jackson, K. A., Gilmer, G. H. and Leamy, H. J. (1980) in Laser and Electron Beam Processing of Materials, eds. White, C. W. and Peercy, P. S. (Academic Press, New York), p. 104. Khan, P. A. A., Debroy, T. and David, S. A. (1988) Weld. J., 67, 1-s. Kobayashi, K. E and Shingu, P. H. (1982) in Rapidly Quenched Metals IV, vol. I, eds. Masumoto, T. and Suzuki, K. (Japan Inst. Metals, Sendai), p.103.
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Kobayashi, K. E and Shingu, E H. (1988) J. Mater. Sci., 23, 2157. Kobayashi, K. E, Kumikawa, M. and Shingu, E H. (1984) J. Cryst. Growth, 67, 85. Kou, S., Hsu, S. C. and Mehrabian, R. (1981) Metall. Trans., 12B, 33. Kumagai, N., Asami, K. and Hashimoto, K. (1986) J. Non-Cryst. Solids, 87, 123. Kuriyama, H., Kiyama, S., Noguchi, S., Kuwahara, T., Ishida, S., Nohda, T., Sano, K., Iwata, H., Kawata, H., Osumi, M., Tsuda, S., Nakano, S. and Kuwano, Y. (1991) Jpn. J. Appl. Phys., 30, 3700. Kuriyama, H., Kuwahara, T., Ishida, S., Nohda, T., Sano, K., Iwata, H., Noguchi, S., Kiyama, S., Tsuda, S., Nakano, S., Osumi, M. and Kuwano, Y. (1992) Jpn. J. Appl. Phys., 31, 4550. Kuriyama, H., Nohda, T., Ishida, S., Kuwahara, T., Noguchi, S., Kiyama, S., Tsuda, S. and Nakano, S. (1993) Jpn. J. Appl. Phys., 32, 6190. Kurokawa, H., Yonezawa, T. and Mitani, T. (1987) IEEE Trans. Magnetics, 23, 2410. Kurz, W. (1995) in Mathematical Modelling of Weld Phenomena, ed. Cerjak, H. (The Institute of Metals, London), p. 40. Kurz, W., Giovanola, B. and Trivedi, R. (1986) Acta Metall., 34, 823. Kurz, W. and Trivedi, R. (1994) Mater Sci. Eng., A179/180, 46. Lanotte, L. and Lannotti, V. (1995) J. Appl. Phys., 78, 3531. Lee, H. J., Akiyama, E., Habazaki, H., Kawashima, A., Asami, K. and Hashimoto, K. (1995) Corros. Sci., 37, 321. Lin, C. J. and Spaepen, E (1981) Appl. Phys. Lett., 41, 55. Lin, C. J. and Spaepen, E (1982) Appl. Phys. Lett., 41, 721. Lin, C. J. and Spaepen, E (1983) in Chemistry and Physics of Rapidly Solidified Materials, eds. Berkowitz, B. J. and Scattergood, R. (TMS-AIME, New York), p. 273. Lin, C. J. and Spaepen, E (1984) Mater Res. Soc. Symp. Proc., 28, 75. Lin, C. J. and Spaepen, E (1986) Acta Metall., 34, 1367. Lin, C. J., Spaepen, E and Turnbull, D. (1984) J. Non-Cryst. Solids, 33, 251. Lipton, J., Kurz, W. and Trivedi, R. (1987) Acta Metall., 35, 957. Liu, A. Y. and Cohen, M. L. (1990) Phys. Rev., B41, 10727. Liu, E L., Yen, R., Bloembergen, N. and Hodgson, R. T. (1979) Appl. Phys. Lett., 34, 864. Lowndes, D. H., Pennycook, S. J., Jellison, Jr., G. E., Withrow, S. E and Washburn, D. N. (1987) J. Mater. Res., 2, 648. Lu, E X., Yang, B. X., Cheng, D. G., Ye, R. Z., Yu, W. X. and Sun, J. B. (1992) Thin Solid Films, 212, 220. Matteazzi, E, Lanotte, L. and Tagliaferri, V. (1985) in Rapidly Quenched Metals V, eds. Steeb, S. and Warlimont, H. (Elsevier Science Pub., Amsterdam), p. 1199. Mazumder, J. and Steer, W. M. (1980) J. Appl. Phys., 51, 941. Mazzoldi, E, Della Mea, G., Battaglin, G., Miotello, A., Servidori, M., Bacci, D. and Jannitti, E. (1980) Phys. Rev. Lett. 44, 88. Mehrabian, R. (1982) Internat. Metall. Rev., 27, 185. Mehrabian, R., Kou, S., Hsu, S. C. and Munitz, A. (1979) in Laser-Solid Interactions and Laser Processing 1978, eds. Ferris, S. D., Leamy, H. J., and Poate, J. M. (Amer. Inst. Phys., New York), p. 129. Messler, Jr., R. W. and Millard, D. L. (1994) Weld. J., 73, 43. Metzbower, E. A. (1990) Crystal Properties and Preparation, 22, 656. Metzbower, E. A. and Denney, E E. (1986) Proc. Inernat. Conf. Power Beam Technol., p. 10-1. Mitsuhashi, A., Asami, K., Kawashima, A. and Hashimoto, K. (1987) Corros. Sci., 27, 957. Miyazawa, Y. and Pound, G. M. (1974) J. Cryst. Growth, 23, 45. Moon, D. W., Metzbower, E. A., Phillips, R. H. and Denney, E E. (1987) in Proc. Conf. Laser vs. Electron Beam in Welding, Cutting and Surface Treatment: State of the Art, Part 2, p. 280. Pound, G. M. and Lamer, V. K. (1974) J. Amer. Chem. Soc., 74, 2323. Prasad, S. V., Zabinski, J. S. and Dyhouse, V. J. (1992) J. Mater. Sci. Lett. 11, 1282. Rappaz, M., Carrupt, B., Zimmermann, M. and Kurz, W. (1987) Helvetica Physica Acta, 60, 924. Renaud, L., Fouquet, E, Millet, J. E, Mazille, H. and Crolet, J. L. (1991) Mater Sci. Eng., A134, 1049. Roessler, D. M. (1986) in The Industrial Laser Handbook, eds. Belforte, D. and Levitt, M. (Penny Well Books, Tulsa), p. 16. Rosenthal, D. (1946) Trans. ASME, 68, 849. Rozgonyi, G. A., Baumgart, H., Phillip, E, Uebbing, R. and Oppholzer, H. (1982) in Laser and Electron-Beam Interraction with Solids, eds. Appleton, B. R. and Celler, G. K. (North-Holland, Amsterdam), p. 177. Sadler, L. Y. and Shamsuzzoha, M. (1997) J. Mater Res., 12, 147.
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Chapter 6 Thermal Plasma Processing 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 RF Plasma Torches 6.2.2 Plasmagen Gases 6.2.3 Plasma-Particle Interaction 6.2.4 Plasma Processing Systems 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
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Chapter 6 Thermal Plasma Processing R V. ANANTHAPADMANABHAN and N. VENKATRAMANI
6.1. INTRODUCTION The twentieth century, variously termed at different times the semiconductor age, the silicon age, and the oxide age, has seen some breathtaking achievements in the development of new materials. These achievements have been made possible by matching developments in the material processing technologies. Today's advanced industries need precisely tailored materials to meet stringent operational requirements. However, a close look at the conventional manufacturing technology tells us that it is beset with barriers, both theoretical and man-made. Table 6.1 summarizes the limitations of the current material processing technologies and how they can be overcome in the emerging technologies of the future. It may be seen that many of the shortcomings of conventional technologies can be overcome by the use of electron, plasma and laser beams. In particular, thermal plasma processing has the capability to transcend thermodynamic and kinetic barriers by opening up new reaction paths inconceivable in the conventional sense. This technology, which until recently was confined to the realms of laboratory, is fast emerging as a major technique not only to synthesize and purify materials, but also to modify and shape surfaces (Boulos 1996, Venkatramani 1995, Kong and Lau 1990, Pauloski 1995). Plasmas of technological interest can be classified into two categories: nonequilibrium plasmas and thermal plasmas (Fauchais et al. 1983, Fineman 1987). Nonequilibrium or cold plasmas, more popularly known as glow-discharge plasmas, are low-pressure plasmas characterized by high electron temperatures and low ion and neutral temperatures. They are widely used in lighting, surface cleaning, etching, film deposition and polymerization. Thermal plasmas are characterized by the electron temperature being equal to the gas temperature. Normally, plasmas in the temperature range of 2,00030,000 K and with charged particle density of 1019-1021 m -3 are termed thermal plasmas. Thermal plasma processing has emerged as a major processing route to synthesize many alloys and ceramics of industrial importance. The different categories of plasmas and their applications are listed in Table 6.2.
6.1.1 Advantages of plasma processing The high temperature, high enthalpy, and fast quench rate which are characteristics of the process, offer unlimited scope to the materials scientist to synthesize newer ma121
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P V. Ananthapadmanabhan and N. Venkatramani
Table 6.1. Limitations of current material processing technologies Barrier
Approach to overcome the barrier
Emerging technology
Thermodynamic barrier
Processing in non-equilibrium and thermal plasmas
Spatial barrier
Access to non-equilibrium regimes, use of exotic intermediates and precursors Transfer heat to volume rather than across surface Micro- and nanoscale processing
Temporal barrier
Nano- and picosecond processing
Purity barrier
Absolute chemical control and limit impurity content
Heat transfer barrier
and
specificity
Ecological barrier
One-step eco-friendly technology, conserving materials and recycling wastes
Electromagnetic processing Beam processing; electron and laser beams and low-pressure plasma processing Laser processing, thermal plasma processing with rapid quench Processing by electron, plasma and laser beams and low-pressure plasmas Thermal plasma processing
Table 6.2. Plasmas of technological interest Category of plasma
Applications
Low pressure plasmas (10-4-10 -2 torr)
Sputtering and surface modification processes, plasma source for ion implantation Etching, microelectronic processing Plasma chemistry, plasma polymerization
Medium pressure plasmas (10-2-1 torr) Subatmospheric pressure plasmas (1-100 torr) Atmospheric plasmas (100+ torr)
Plasma-spraying, plasma-melting, material synthesis
terials with improved properties. The temperature encountered in the plasma medium is 10,000-15,000 K, whereas the maximum temperature available in conventional heat sources is 3,000 K. This enables the adoption of novel approaches to accomplish chemical reactions and metallurgical processes, inconceivable in the conventional sense. The high-temperature gas phase chemistry in the plasma medium promotes various chemical reactions at much faster rates than those encountered in conventional processing techniques. The high quench rate, typically of the order of 105-107 K/s, prevents recombination of the products and favors homogeneous nucleation resulting in nano-sized particles. Furthermore, it offers flexibility in the choice of reactants, which can be solids, liquids or gases. Plasma processes are, usually, one-step processes, with the capability of handling large throughputs in small reactor volumes in relatively short processing times. They also do not substantially contaminate the process environment, thereby ensuring high purity of the product.
Thermal Plasma Processing
123
Plasma processing has been applied to synthesize a variety of alloys, intermetallics, and refractory ceramics. Plasma-spray technology is a well-established surface treatment technique to prepare protective coatings for a variety of industrial applications. The focus of this chapter will be on thermal plasma processing of materials. The organization of the chapter is as follows: Introduction to the subject of thermal plasma processing will be followed by a description of the different techniques of plasma generation. Applications of thermal plasmas in material processing are discussed in detail in Section 6.3. Plasma synthesis of ultrafine alumina is discussed in detail towards the end of this section to illustrate the formation of non-equilibrium phases in reactions occurring in plasma medium.
6.2. T H E R M A L
PLASMAS
The classical plasma is a fully ionized gas consisting of electrons and ions that interact with each other only through Coulomb forces (Eliezer and Eliezer 1989). This model is valid only for astrophysical or fusion plasmas; it does not hold good for industrial plasmas, which are partially ionized gases and are dominated by atomic and molecular processes. Thermal plasma is a viscous, electrically and thermally conducting fluid. Plasma retains many of the properties of gases, and behaves in conformance with the physical laws valid for gases. The specific properties of plasma that distinguish it from a gas become apparent in the presence of a magnetic field, when the plasma acquires non-isotropic properties. Thus, the fundamental difference between plasma and a non-ionized gas lies in their responses to electromagnetic forces. The electrically charged particles present in plasma are affected by externally applied electric and magnetic fields, and also interact with one another. The electric field that they set up is so extensive that every particle is affected by a multitude of other particles. Consequently, the collective behavior of the particles is much stronger in plasma than in a non-ionized gas. The unique feature of the thermal plasma jet that distinguishes it from other heat sources is its high power density. The energy density of thermal plasma devices is of the order of several GW/m 2, which is 10-100 times the power density of conventional oxy-fuel flames. Although higher power densities are obtainable with electron and laser beams, the source strengths available with these devices, especially laser devices, are not high enough for large-scale material processing applications.
6.2.1 Principles of plasma generation Thermal plasmas for material processing applications are usually of electric origin. Laboratory thermal plasmas are produced in devices known as plasma torches or plasmatrons (Heberlein 1992). Plasma torches are electrothermal devices, which convert the electrical input energy to thermal energy of the plasma. Depending on the primary source, which can be direct current (DC), alternating current (AC) at mains frequency, or at radio frequency (RF), they are known as DC, AC or RF torches. There are also combinations of these torches known as hybrid devices. Figure 6.1 gives a schematic of the different types of plasma torches.
124
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Figure 6.1. Schematicof different plasma sources: (a) DC plasma torch--rod cathode type (non-transferred arc mode); (b) DC plasma torchmrod cathode type (transferred arc mode); (c) DC plasma torch--hollow cathode type; and (d) RF plasma torch.
6.2.1.1 DC plasma torches The most commonly used plasma device for material processing applications is the DC arc plasma torch. A direct current arc is struck between a cathode rod and a nozzle anode and the energy from the arc is abstracted by the plasma gas injected in the inter-electrode region. Forced gas flow extends the arc in the anode nozzle which is strongly water-cooled (Fig. 6.1(a)). A thermal arc pinch effect is produced by the joint action of the cold wall arc channel and the cold gas sheath around a very high-temperature conducting core (the arc column). The DC arc in a plasma torch needs to be stabilized, i.e., it should remain stationary against fluctuations. This is often done by constricting the arc to a well-adjusted narrow high-temperature, highly conducting arc column. Various torch configurations are possible depending upon the arc stabilization mode: tangential vortex gas input in the arc channel, axial gas input along the cathode, segmented anode arc and magnetic stabilization (Fauchais et al. 1983). The magnetic field can be self-induced (by an arc current greater than 8000 A), or externally generated. The electric arc is stabilized by using a constricted anode nozzle and by the resultant aerodynamic effects in the streaming plasma gas. Stabilizing action of the vortex gas flow provides a cold boundary layer near the anode wall so that heat loss to the wall is reduced. This results in the thermal energy being highly concentrated with improved torch stability and efficiency. The localized heat flux at the nozzle can be as high as 160 W/mm 2. The obvious choice of material for the anode is copper, although molybdenum and graphite are also used. Depending on the nature of the gas and the working parameters, the anode losses range
Thermal Plasma Processing 3
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Figure 6.2. Cross-section of a typical DC plasma torch: (1) cathode; (2) anode; (3) anode segment; (4) insulator; (5) cathode segment; (6) water inlet; (7) plasma gas inlet; and (8) water outlet.
between 80 and 10% of the energy input in the arc. The anode losses are proportional to the current density and are a function of the arc voltage. The cathodes can be of the thermionic type such as tungsten, carbon or molybdenum, which, obviously, must be used in non-oxidizing atmosphere. Under certain conditions of oxidizing atmosphere, one can use zirconium or hafnium cathodes. Heat losses at the cathode are generally quite low (less than 10% of the input power). For a tungsten cathode tip, the erosion rate has been estimated to be of the order of 10 -9 to 10-10 g per Coulomb, so that its lifetime is about 150 h. Non-thermionic cathodes are normally made of copper. The cross-section of a typical DC plasma torch is shown in Fig. 6.2. This design is based on a rod type cathode and nozzle-shaped anode with tangential gas entry through the insulator module. When a gas is injected into the electrode gap and a high intensity current is passed, a DC arc is established between the electrodes. The plasma gas extracts energy from the arc and emerges out of the nozzle as a high-temperature, high velocity jet. The temperature at the core of the plasma jet ranges between 15,000-20,000 K (Gerdeman and Hecht 1972). The typical operating parameters of a low power, laboratory DC plasma torch (non-transferred arc mode) are given in Table 6.3 The plasma torch can be operated in a transferred or non-transferred mode, depending on whether the arc is electrically transferred to the work piece or not, as illustrated in Fig. 6.1 (a) and 6.1 (b). The transferred arc mode is mainly used for cutting metal sheets, refining, and remelting operations. The non-transferred arc mode is used in plasma-spraying and other material processing operations.
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P. V. Ananthapadmanabhan and N. Venkatramani
Table 6.3 Typical operating parameters of a low-power laboratory DC plasma torch (non-transferred arc mode) Arc voltage (V) Arc current (A) Arc gas Torch input power (kW) Water Inlet temperature ((C) Water outlet temperature ((C) Cooling water flow rate (1/min) Power loss at the electrodes (kW) Flow rate of Argon (g/s) Cold gas flow velocity (m/s) Thermal energy of emerging plasma (kW) Enthalpy of plasma (kJ/g) Plasma temperature at the nozzle exit (K) Plasma exit velocity (m/s)
40 250 Argon 10 33.2 36.3 15 3.2 1.19 13.27 6.8 5.71 10,600" 468.9*
* Calculated.
6.2.1.2 AC plasma torches When alternating current at mains frequency is used in a two-electrode device, each electrode alternately acts as the cathode and anode. In each half cycle a DC discharge is established once the breakdown potential is crossed and the discharge is maintained so long as the applied voltage remains above the maintenance potential. When the voltage falls below the maintenance potential, the discharge gets extinguished, and the space charge gets swept away. The continuous alternation of the polarity of the electrodes limits strongly the stability of the discharge and the choice of the electrode materials.
6.2.1.3 RF plasma torches Radio-frequency (RF) inductive or capacitive coupling (between 1-20 MHz) can be used to generate a plasma column, which can be stabilized in a quartz tube. The main advantage of this technique is its electrodeless feature, allowing the use of aggressive gases such as oxygen, chlorine (Mackinnon and Reuben 1975) or UF6, etc. which cannot be used as plasmagen gases in DC plasma torches. The minimum power necessary to sustain the discharge is determined by the nature of the gas, its pressure and the frequency of the electromagnetic field (Rykalin 1978). The gas velocity in an RF torch is several tens of m/s, which is much lower than that encountered in DC torches (hundreds of m/s) (Dembovski 1985). Furthermore, the diameter of the reactor is important (20-200 mm) and the energy density is much lower than in arc plasma generators. RF plasma torches capable of operating up to power levels of 1 MW (Ionarc Tafa torches) and a lifetime of a few thousand hours are presently employed for material processing applications. However, the electrothermal efficiency of these torches is low (40-50%) and are mostly limited to specialized applications.
Thermal Plasma Processing
,•,4
11,S00K
10,500K
9,600 K
8,600 K
7,600 K
127
6,800 K
E2
0
4
8
12
16
20
24
Axial Distance(mm) Figure 6.3. Axial and radial temperature distribution in a nitrogen plasma jet (Fauchais et al. 1983).
The arc column can be extended by using hybrid torches, which consist of a combination of torches connected in tandem. The extended hot zone increases the residence time and ensures that the reaction goes to completion.
6.2.2 Plasmagen gases The most commonly used gases for plasma generation are argon, nitrogen, helium, hydrogen, and air. Plasma gas flow rate and the electric power to the plasma torch must be properly balanced in order to get a stable arc. The choice of plasma gas depends on many factors, such as the design features of the torch, in particular the electrode materials (Venkatramani 1995). In the case of plasma torches employing tungsten cathode, the choice of the plasma gas is limited to inert gases and non-oxidizing gases. Gas enthalpy is another important factor deciding the choice of the gas (Gerdeman and Hecht 1972). The energy content of nitrogen and hydrogen, which are diatomic, is considerably higher than that of argon or helium. This is due to the dissociation reaction in the case of nitrogen and hydrogen prior to ionization. The heat transfer coefficient is also higher, enabling efficient plasma-particle heat transfer. The low cost and high internal energy of nitrogen makes it the most commonly used gas. If a completely inert atmosphere is required, argon is usually preferred. Reactive gases like hydrogen, oxygen, chlorine, methane and ammonia/nitrogen can be used to impart reducing, oxidizing, chloriding, carburizing, or nitriding effects, respectively, to the plasma. The axial and radial temperature fields in the plasma jet issuing out of a nontransferred arc torch are shown in Fig. 6.3. It is seen from the figure that the temperature falls off rapidly downstream the axis and radially away from the centre-line. The maximum temperature and the temperature profile depend on the input power, nature and flow rate of the plasma gas and the torch design. But the significant point to be noted is the steep gradient in the axial and radial temperatures. Particle injection into the plasma stream is not a trivial problem. In fact, it can be very difficult due to the high viscosity of the plasma. Precursor powders are usually entrained
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P V. Ananthapadmanabhan and N. Venkatramani
in a carrier gas for injection into the plasma. The carrier gas can be inert or reactive. There is an increasing tendency to use liquid reactants, which is known as liquid injection plasma synthesis (LIPS) (Kong and Lau 1990).
6.2.3 Plasma-particle interaction The heat transfer process from the plasma to the particles is of fundamental importance in all plasma processing applications. Plasma-particle heat transfer is very complicated and depends on the thermo-physical properties of the material as well as the enthalpy and composition of the plasma. Many theoretical models (Fauchais et al. 1983, Bonet 1976) have been developed for the plasma-particle heat transfer phenomenon. The heattransfer process is decided by the thermal boundary at the particle-plasma interface and the properties of the gases in this boundary layer, which is characterized by steep temperature gradient and strong deviations from chemical and kinetic equilibrium (Pfender 1994). The heat transfer coefficient to the particle may be calculated from a correlation based on the well-known Ranz-Marshall equation (Ranz and Marshall 1952): Nu = 2.0 + 0.6Re~
~
(6.1)
where Nu, Re and Pr are the Nusselt, Reynolds and Prandtl numbers, respectively. Although this equation must be corrected by additional factors to account for the strong variation of gas properties in the boundary layer, it does give a fair estimate of the heat transfer coefficient. The coefficient of heat transfer between the plasma gas and the surface of the particle can now be calculated from the definition of the Nusselt number: Nu - hdp/k
(6.2)
where h is the heat transfer coefficient, k is the thermal conductivity of the plasma, and
dp is the diameter of the particle. The goal pursued while optimizing the process parameters is to ensure that all particles entering the plasma jet are heated to the desired temperature, which may be the melting point in the case of plasma-spraying or the reaction temperature in chemical synthesis or dissociation reactions. If the particles are small or if they conduct heat well, this objective is easier to achieve. On the other hand, in the case of large or poorly conducting particles, it might be necessary to choose processing conditions that lead to intense evaporation from the surface of the particles. The large thermal gradients within the particle may also lead to partially molten particles, which lead to porosity in the coatings. Numerous onedimensional and elaborate two-dimensional models have been worked out for plasmaparticle heat transfer problem. In the case of good heat-conducting particles, the equation describing particle melting in the plasma is given by (Pauloski 1995):
telqd2h(Tg - Tp)
-
3
1/61-IdpppCp{(Tm
- 300) -k- A H f
}.
(6.3)
The left-hand side of Eq. (6.3) describes heat transferred from the flame to the particle,
Tg is the gas temperature and Tp is the particle temperature, which is assumed to be uniform in the entire volume of the particle. The right-hand side describes the heat absorbed
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by the particle, pp and Cp being the density and specific heat of the particle respectively. A Hu is the latent heat of fusion of the material per unit volume and te the time for heating the particle from room temperature to its melting point. It is assumed that the energy transferred by radiation from the plasma is small.
6.2.4 Plasma processing systems Plasma-processing encompasses a wide range of operations including plasma-cutting and welding, plasma-melting and machining, plasma-spraying, chemical synthesis and mineral dissociation. A wide variety of plasma reactors have been developed based on DC non-transferred and transferred arc plasma torches. The various types of plasma reactors used for material processing applications have been described by Fauchais et al. (1983). The exact design of the reactor set-up depends on the specific application, but the general features remain more or less the same. The schematic of a typical plasma reactor developed in our laboratory for material synthesis is shown in Fig. 6.4. The major subsystems are plasma torch, powder-feeder, reaction chamber, torch power supply and control console. The reaction chamber consists of double walled water-cooled stainless steel segments provided with sampling and diagnostic ports. The plasma torch is mounted on one end of the reaction chamber as shown in the figure. There is also provision to collect the product and feed any reactive gas. The power sources generally employed are DC motor generator sets, transductor controlled rectifier units or silicon controlled rectifier (SCR) system. The SCR systems are currently the most extensively used power sources. A high frequency (HF) igniter to initiate the plasma is also included in the power supplies. The HF unit is connected in parallel with the DC power supply, which must be protected from it. The control console has instruments for the measurement and control of the torch inputs (viz. electrical, cooling water, gas and powder feed rates). Powder is injected into the plasma by means of the powder feeder. The turn-table-feed type powder feeder is commonly used. This ensures smooth volumetric flow of powder and has very good control over the feed rate and requires low carrier gas. 6.3. P R O C E S S I N G OF M A T E R I A L S
Plasma processing is an emerging technology that utilizes the exotic properties of the plasma medium to effect physical, chemical or metallurgical reactions to produce new materials or impart new functional properties to conventional materials. Plasma processing has been successfully applied to develop advanced ceramic coatings, synthesis of nanocrystalline materials, processing of minerals and ores, and treatment of hazardous wastes.
6.3.1 Plasma-spraying Plasma-spraying is a surface modification technique that combines particle melting, rapid solidification, and consolidation in a single process. This technique has been widely used to develop protective coatings of ceramics, alloys, and composites to enhance the surface properties of critical components operating in severe environment (Pauloski 1995,
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P. V. Ananthapadmanabhan and N. Venkatramani 5
6
7
8
=
=
QIO
OlO
['~
Figure 6.4. Schematic of thermal plasma reactor for material synthesis: (1) cooling water; (2) control console; (3) power supply; (4) powder feeder; (5) plasma torch; (6) reaction chamber; (7) sampling duct; (8) powder collector; and (9) auxiliary instruments.
Gerdeman and Hecht 1972). In conventional plasma-spraying, the material to be coated is introduced in the plasma jet in powder form by means of a carrier gas. The powder particles, as they enter the plasma jet, are heated and melted and the molten droplets are accelerated to very high velocities, as high as 100-400 m/s. As these molten droplets strike the substrate surface, they flatten and get anchored to the surface irregularities to form an adherent coating. Plasma-spraying has several advantages over other coating processes. Because the high temperature generated by the plasma arc is far beyond the melting point of any known material, any material that melts and remains in the liquid state over a distinct range can be spray-coated. Although the plasma temperature may be in the range of 7,00010,000 K, the substrate temperature can be kept low (below 400 K), avoiding the risk of substrate distortion. The spray technique does not impose any restriction on the workpiece dimensions and large samples can be coated. The adherence of the coating to the substrate surface is predominantly mechanical in nature (Pauloski 1995). When the molten particles impact a roughened surface, they get interlocked to the surface irregularities resulting in a mechanical adherence, which is believed to be the predominant mechanism in plasma-sprayed coatings. However, depending on the nature of the substrate material, its temperature and the impinging particles, a metallurgical bond can result as in the case of nickel aluminide or molybdenum on steel substrates.
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Figure 6.5.
,
~
,
131
-
SEM photograph of a typical sample of plasma-sprayed coating of alumina.
The coating quality is affected by a number of interdependent experimental parameters, which need to be carefully controlled. Thermal interaction of the particles with the plasma jet, the particle trajectory and the dwell time are important factors influencing coating adhesion and other properties. Particle size decides the particle trajectory and degree of melting, which controls the porosity and adhesion of the coatings. As it is practically impossible to have monosized powders, particle size in a narrow range is selected. Although the exact particle size depends on the nature of the material to be spray-coated, the recommended size for ceramic powder is 3 0 - 4 0 / z m and for metal powders 40-60 /~m.
6.3.1.1 Structure of sprayed deposits The sprayed layer has a typical lamellar structure, built up by successive layers on the material (Gerdeman and Hecht 1972). Heat transfer calculations show that freezing of the particles occurs in a few microseconds and that complete solidification occurs in perhaps 100/zs. The zone of thermal effect in the underlying material is, therefore, quite small and temperature gradients reach 105 K/cm. The wetting and flow properties of the liquid droplets are important since they influence the porosity within the coating and at the substrate interface. The surface morphology of a typical sample of plasma-sprayed alumina is shown in Fig. 6.5. The deposit shows grains consisting of melt-quenched particles with interganular and intragranular porosity, which are characteristic of plasma-sprayed coatings. The porosity of the coating is about 15-20%, which can be considerably reduced by optimizing the experimental parameters and carrying out the coating operation under reduced pressure.
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The flow and solidification of molten particles on impact with a cold surface is a difficult problem to treat theoretically because of the interaction of heat transfer and crystal growth. The very high velocity at which the various processes occur also makes direct experimental observation extremely difficult. Madejski (1976) has carried out a theoretical treatment of the impact of a molten droplet on a cold substrate, taking into account surface tension, viscosity and crystallization effects. The results show that the degree of flattening (D/d), defined as the ratio of the diameter of the flattened disc (D) to that of the original drop (d), is related to the density p, viscosity/z and impact velocity of the drop v by the following equation: D / d = 1.29(pvd/#) ~
(6.4)
Substituting appropriate values for A1203, in the velocity range of 100-400 m/s gives a value of 3-6 for the degree of flattening, which seems to be in reasonable agreement with various experimental observations (Zolotowski 1968). This suggests that at least for alumina, flattening is completed before crystallization becomes significant and the two processes can largely be treated separately. Plasma-spraying conditions are very similar to those occurring during splat-quenching of metals which has been widely studied. The cooling rate has been of particular interest and various techniques have been used to estimate the thermal conditions as a molten particle spreads and solidifies (Jones 1973). The main results of this work show that the cooling rate is largely independent of substrate material and is controlled by the interface between substrate and particle if the flattened particle has a thickness of less than a few #m. The estimated heat transfer coefficient under these conditions is of the order of 105 W m -2 K -1 . Applying this result to the case of plasma sprayed A1203 gives a cooling rate of the order of 107 K/s and a freezing time of about 10 #s. This also suggests that there is a large temperature gradient across the interface between the particle and the substrate (or particle and previously deposited layer) and the temperature of the underlying material remains relatively low. This is particularly so if the substrate is cooled during spraying by an air blast to prevent heating by plasma gases and to remove the sensible heat and heat of fusion of the sprayed particle. 6.3.1.2 Plasma spray-deposited materials Applications for plasma-sprayed materials range from microelectronics to large engineering structures. There are excellent reviews and publications (Pauloski 1995, Gerdeman and Hecht 1972, Zaat 1983, Fauchais and Vardelle 1994, Notomi and Hisatome 1996, Boulos 1992, Matejka and Benko 1989) exclusively devoted to the coating process and applications, which can be referred to for further details. Plasma-sprayed coatings of metals, alloys and ceramic materials have been used for various applications. Plasma-spraying has been applied to produce wear-resistant coatings of ceramics and metals on various components. Successful wear-resistant coatings have been developed using cobalt- and nickel-based alloys containing carbides of chromium, tungsten and titanium as abrasion-resistant particles (Rangaswamy and Herman 1986, Mohanty and Smith 1995, Knight et al. 1995). Self-lubrication can be provided by incorporating low-friction materials or dry lubricants such as graphite, molybdenum disul-
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phide, or organic lubricants. Plasma-spraying has been effectively used to protect metallic and non-metallic components from high temperature and oxidation. Plasma-sprayed thermal barrier coatings consisting of multi-layered and graded ceramic and metal coatings are being extensively used for protecting the surfaces of metal components, gas turbines and other critical engine parts exposed to corrosive environment. Another important application of plasma-spraying is reclamation, where it is used to repair worn out or badly machined components. Worn out surfaces of brake drums, grinders, spindles, etc. can be reclaimed by plasma coating. Oxide coatings are used for many high-tech as well as conventional applications. Plasma-sprayed coatings of stabilized zirconia, alumina, and alumina-titania have been extensively studied (Bratton and Lau 1981, Stecura 1977, Ananthapadmanabhan et al. 1991, Ramachandran et al. 1997). Aluminium oxide (often mixed with titania) and chromium oxide are used on many textile machinery components such as thread guides (Matejka and Benko 1989). They are also widely used on pumps and mechanical seals to provide resistance to both abrasion and corrosion. Plasma-sprayed zirconia coatings, as thermal barriers, are effective in improving the efficiency of gas turbine engines. The preferred material is 6-8 wt.% Y203-ZrO2 or 24 wt.% MgO-ZrO2 deposited over Ni-CrA1-Y bond coat applied by vacuum plasma-spraying (Meetham 1988). The latter provides oxidation resistance to the superalloy substrate and is also strain tolerant, which helps to accommodate stresses arising out of thermal expansion mismatch between the base metal and the ceramic coating. Thermal shock resistance of zirconia coatings has been greatly enhanced by using composite coatings of alumina and calcia stabilized zirconia. The improved thermal shock behavior is attributed to destabilization of calcia stabilized zirconia leading to the generation of microcracks which act as thermal stress relief centers (Ananthapadmanabhan et al. 1991). When ceramic coatings are applied on metal substrates, the large mismatch in the thermal expansion coefficients of the metal base and the coating material can result in interface stresses. This can lead to poor adhesion and even coating failure. In order to overcome this, an intermediate coating, referred to as bond coat, is applied on the metal substrate before applying the ceramic coating. The primary function of the bond coat is to reduce the interface stress and improve coating adhesion (Matejka and Benko 1989). Nickel aluminide and its derivatives, Ni-Cr-A1-Y, are frequently used as the bond coat material. The reaction of nickel and aluminum to form the intermetallic is highly exothermic, resulting in high adhesion strength. Nickel aluminide coatings are usually prepared by using special composite powder of nickel and aluminum (Matejka and Benko 1989). This composite powder forms the compound during the spray process. A one-step process to form plasma-sprayed Ni3A1 coatings has been developed (Ananthapadmanabhan et al. 1995) by using a blend of nickel and aluminum powders as the feedstock material. The process parameters have been optimized to effect compound formation as the powder particles deposit on the substrate surface. Coating adhesion was measured by tensile adhesion strength as described in ASTM C 633 (Lin and Berndt 1994). The adhesion strength of the coatings was found to be much greater than 60 MPa.
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The cooling rate for plasma-sprayed particles approaches 106 K]s and this produces a fine-grained microstructure within the splat (Scott and Kingswell 1991). This rapid solidification can produce amorphous deposits from some ferrous alloys (Herman and Bhat 1980), while metastable or non-stoichiometric phases can be retained in some ceramics. Phase composition and stability of yttria stabilized zirconia coatings have been extensively studied. The phase structure of yttria stabilized zirconia coatings have been shown to consist of non-equilibrium phases. The work of Miller et al. (1981) and Bratton and Lau (1981) on plasma-sprayed zirconia thermal barrier coatings showed that the coatings consisted of non-equilibrium phases. Irrespective of the starting material, the as-sprayed material was found to contain a non-transformable tetragonal (high yttria) phase. It was observed that partially stabilized zirconia (ZRO2-8% Y203) as well as fully stabilized zirconia (ZRO2-20% Y203) contain primarily the non-equilibrium tetragonal phase, which transformed to the equilibrium cubic and tetragonal (low yttria) during thermal annealing. Plasma-sprayed coatings of alumina and alumina-titania have been shown to consist of metastable alumina (Heintze and Uematsu 1992, Ramachandran et al. 1997). The work of Heintze and Uematsu (1992) has shown that plasma-sprayed coatings of alumina predominantly consist of y-alumina. They also found that u-alumina can be formed by careful optimization of the process parameters.
6.3.1.3 Low-pressure plasma-spraying Atmospheric plasma-spraying (APS) is seldom satisfactory when used to prepare coatings of metals, alloys, and non-oxide ceramics. When coatings of refractory carbides, nitrides and borides are prepared by APS, the high velocity plasma jet acts as an aspirator and sucks in air from the surrounding atmosphere resulting in oxide inclusions, which can seriously affect the coating performance. Studies on the phase composition of TiB2 coatings prepared by APS showed appreciable amount of boric oxide, formed by partial in-flight oxidation of TiB2 particles (Ananthapadmanabhan et al. 1993). The low melting point of B203 (about 800 K) affects the high temperature performance of TiB2 coatings. This problem can be overcome by low-pressure plasma-spraying (LPPS) or vacuum plasma-spraying (VPS), where the coating operation is performed in a vacuum chamber. Alternately, the chamber can be filled with inert gas to minimize coating contamination. The quality of the coatings produced by LPPS is superior to that of those prepared by APS, in terms of coating properties as well as chemical purity. LPPS reduces porosity and improves coating adhesion. Reactive metals like titanium, zirconium and tungsten can be spray-coated with minimum oxide contamination. LPPS has been extensively used to prepare coatings of refractory carbides, nitrides and other non-oxide materials for various applications. Recent developments include the successful application of LPPS to fabricate components for solid oxide fuel cell (SOFC) components (Notomi and Hisatome 1996). Dense electrolyte coatings of yttria stabilized zirconia have been prepared by low pressure plasma-spraying. A 1 kW module has been developed and successfully operated for 3,000 h.
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6.3.1.4 Reactive plasma-spray f o r m i n g Conventional plasma-spraying consists in feeding metal, alloy, or ceramic powders in the plasma jet, where the particles melt and the molten droplets are deposited on to a substrate surface, forming an adherent coating. It is essentially a physical process involving melting and resolidification of the powder. In reactive plasma-spraying, the high temperature plasma environment is utilized to realize inflight chemical reactions by using suitable reactive gases along with the feed powder (Mutasim et al. 1992). The metal or ceramic powders are introduced in the plasma jet along with a suitable reactive gas. The reactant powder particles melt and the molten droplets react with the reactive gas, resulting in the in situ formation of the desired compound, which is deposited on the substrate. This technique has been used to develop refractory carbide coatings. Reactive plasma-spraying of Ti using methane, propylene, or acetylene as the reactive gas has been carried out (Mutasim et al. 1992, Ananthapadmanabhan et al. 1993) to produce Ti-TiC composite coatings. Dissociation of the hydrocarbon to form finely dispersed aerosol of carbon is the first step in the formation of the carbide. These fine particles of carbon subsequently react with the molten droplets of the metal to form the metal carbide, which is deposited on the substrate. By proper selection of the process parameters, composite coatings containing 60 vol.% of TiC dispersions have been obtained. It was also observed that higher substrate temperatures increased the carbide content. A recent development in this field is a one-step process to form Fe-TiC cermet coatings by reactive plasma-spraying of ilmenite (Ananthapadmanabhan and Taylor 1998). The principle involves the reduction of ilmenite by methane in a thermal plasma environment according to the following chemical reaction:
FeTiO3 + 4CH4 = Fe + TiC + 3CO + 8H2.
(6.5)
Ilmenite powder is introduced into the plasma jet using a mixture of methane and argon as the carrier gas and the substrate is kept at about 1200 K. Ilmenite is reduced to iron and titanium dioxide, which is further reduced to its lower oxides and finally to TiC. Fig. 6.6 shows the surface morphology of a typical Fe-TiC coating, revealing excellent inter-particle cohesion and very little porosity. Reactive plasma-spraying can be extended to develop other refractory carbide and nitride coatings by suitable choice of precursor materials and reactive gases. Reactive plasma-spraying is also used for depositing hydroxy apatite (HA) coatings for dental implants. Conventional atmospheric plasma-spraying of hydroxy apatite, Cal0(PO4)6(OH)2 results in the formation of biodegradable phases such as, fi-tricalcium phospate and tetracalcium phosphate by the thermal dissociation of HA, rendering it unsuitable for biomedical applications (Hench 1991). The decomposition of HA to biodegradable compounds has been eliminated by a reactive plasma-spray technique developed by Patil et al. (1996).
6.3.1.5 Plasma spheroidization Powders of metals, alloys and ceramics can be melted in a plasma jet. Melting of a particle results in the formation of a spherical drop under the action of the surface tension forces and this shape is usually retained after solidification, thus providing the name to
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P. V. A n a n t h a p a d m a n a b h a n and N. Venkatramani
Figure 6.6. SEM photograph of Fe-TiC coating formedby reactive plasma-spraying.
the process. This treatment may be used simply to give a spherical particle shape for particular application, such as Fe304 for photocopying, plasma spray-quality powders for thermal spray applications and UO2 for dispensed nuclear fuels. Spheroidization is particularly useful to prepare spray-quality powders of special materials for thermal spray applications. Spheroidization of a powder blend of Ni-15% A1 has been carried out in a thermal plasma reactor (Ananthapadmanabhan et al. 1997). The as-collected powder showed poor crystallinity as indicated by the broad X-ray diffraction pattern (Ananthapadmanabhan et al. 1995). Annealing at 800 K in an atmosphere of flowing argon resulted in the formation of well-crystalline Ni3A1. The process parameters have been optimized to yield a precursor powder of Ni-A1, which forms Ni3A1 on the substrate surface during plasma-spraying. A typical photomicrograph of plasma spheroidized Ni-A1 powder is shown in Fig. 6.7. Besides possessing spherical morphology, the powder has excellent flow characteristics, making it ideally suited for thermal-spray applications. Spheroidized particles may also be prepared by atomization of rods or wires fed into a plasma. The diameter of the particles produced depends on the diameter of the torch nozzle, the gas flow rate, plasma density and temperature and surface tension of the material in the liquid state (Rykalin 1976). Alumina particles of 40 # m diameter have been produced by feeding alumina rod in a hydrogen plasma. The heat transfer can be improved if the arc is transferred to the wire but this is not possible with non-conducting materials. An example of large scale industrial application is the spheroidization of magnetite for photocopying applications, where magnetite particles, 125/zm diameter, are produced in 600 kW AC air plasma heater with a power consumption of 2 kWh&g (Fey et al. 1975).
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Figure 6.7. SEM photograph of plasma spheroidized Ni-A1 powder.
6.3.2 Plasma reactors and furnaces
Plasma technology can be gainfully utilized to treat minerals and ores to yield value-added materials. Decomposition processes which are already proved both thermodynamically and technically feasible and favorable include dissociation of carbonates, oxides, sulfides, halides and complex polymetallic ores (Dembovski 1985). It has also been used for metallurgical applications and in the synthesis of advanced ceramics. Table 6.4 shows the high power plasma systems used for diverse applications.
6.3.2.1 Plasma decomposition Decomposition of zircon sand in thermal plasma reactors to produce zirconia is well established. The principle involved is the thermal dissociation of zircon into its component oxides, followed by quenching the reaction products to yield a mixture of ZrO2 and SIO2. Zircon is thermodynamically stable up to 1950 K, above which it spontaneously dissociates into zirconium dioxide and silicon dioxide. The reverse reaction is prevented by freezing the high temperature equilibrium by rapid quenching of the products. Zircon sand is fed into the plasma by means of a powder feeder. As the zircon particles pass through the plasma, they are heated to temperatures above 2,000 K, when dissociation takes place and the product is rapidly quenched. Zirconia is separated from the oxide mixture by leaching away silica with 30-50% solution of sodium hydroxide. Extensive studies on the influence of operating parameters on the dissociation reaction have been carried out by Ananthapadmanabhan et al. (1994a). Results of the study showed that the extent of dissociation of zircon was affected by the power input to the
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P. V. Ananthapadmanabhan and N. Venkatramani
Table 6.4 High-powerplasma torches for material processing applications Torch make Huls Linde-Retech Tioxide Aero-Spatiale SKF Ionarc Metco Daido Tetronics Voest-Alpine Chemische Werke Huls AG Voest-Alpine
Power ( M W ) 8.5 0.7 5.0 5.0 7.0 0.35 0.1 1.0 3.0 7.5 8.0 4.0
Applications Chemical synthesis Plasma melting Titanium dioxide Ferro-manganese Recovery from dust Zirconia production Plasma-spraying Scrap melting Tundish heating Steel remelting Synthetic gas for direct reduction of iron ore Ferromanganese
plasma, plasma gas flow rate, carrier gas flow rate, feed rate of zircon, etc. Ionarc Tafa has developed a 300 kW multi-torch furnace for processing zircon sands at 300 kg/h feed rate. The process has been shown to be economically superior to the conventional process. The product purity is superior with an electricity consumption of 1.32 kWh/kg of the product compared to 9.9 kWh/kg for the conventional process (Dembovski 1985). A plasma reactor has also been used to convert ilmenite to titanium dioxide. The process involves reductive dissociation of ilmenite to TiO2 and iron in a thermal plasma medium using controlled amount of methane as the reductant. The process has been extended to prepare Fe-TiC powder (Taylor et al. 1995).
6.3.2.2 Plasma metallurgy The use of RF and DC plasmas for metallurgical applications is getting established in many countries throughout the world. The introduction of plasma furnaces has revolutionized the metallurgical industry by improving the operational efficiency and quality of the products. The conventional iron and steel industry is highly capital-intensive. The blast furnace operation utilizes dwindling resources of expensive high-grade coking coal, which is used for heating and reducing operations. Consumption of this expensive coke can be considerably reduced by augmenting the conventional blast furnace operation with the plasma process (Dembovski 1985). The process consists in partially replacing coking coal by superhot reducing gas synthesized by plasma-heating of a mixture of natural gas and air or carbon dioxide. Superhot reducing gas can reduce the use of coke by as much as 75%, resulting in substantial cost savings. This process is referred to as the direct reduction process and the direct-reduced iron (DRI) is used for producing low-carbon steels. Mintek, South Africa has developed pilot-plant-scale furnaces for producing low-carbon steel by smelting DRI and iron ore (McRae et al. 1985).
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The application of thermal plasma technology to steel-plant dust is a recent development. Dust and metal-bearing wastes from steel plants consist of oxides of various metals and their disposal in landfill sites poses environmental hazards, since toxic metals may be leached into drinking water supplies. Thermal plasma processing of steel-plant dust not only offers an effective way to destroy the hazardous waste, but also to recover valuable metals such as zinc, nickel and chromium. SKF of Sweden has developed the Plasmared processes for recovering metals from metal-bearing wastes in a reducing atmosphere generated by the decomposition of fossil fuels in a plasma stream. The Enviroplas process is a DC plasma arc process, developed by Mintek, South Africa, which can treat arc-furnace dust, alloy-steel dust and other metal-bearing wastes. Production of ferrochrome, ferromanganese and nitrogen alloying of steel are other areas where plasma technology has been extensively used. DC transferred arc furnaces have replaced the conventional submerged arc furnaces for producing ferrochromium from chromite (FeCrO3) ore. The DC arc process ensures higher chromium recovery with reduced levels of phosphorous and sulfur (Jones et al. 1997). Nitrogen alloying of steel by plasma treatment is illustrative of the unique advantages of the process over conventional metallurgical alloying processes. Interstitially dissolved nitrogen improves the mechanical properties and corrosion behavior of steel significantly. The strength and corrosion resistance of some steel grades are commonly enhanced by addition of nitrogen-enriched ferrochrome, ferromanganese or calcium nitride. In conventional induction furnaces with nitrogen partial pressure of 105 Pa, the amount of nitrogen dissolved in steel at 2,500 K is hardly 0.06%, which can be increased to about 0.2% by weight under very high partial pressures of nitrogen. In plasma furnaces, however, a substantial fraction of nitrogen is present in atomic, excited and ionic state. Under identical conditions of temperature and pressure, plasma processing can increase the dissolved nitrogen content in iron or steel by as much as an order of magnitude over the levels attained in induction furnaces. The tensile strength, yield point and creep strength of 18Cr-8Ni steel are found to increase with the amount of dissolved nitrogen. A nitrogen content of 0.8% raises the yield strength from about 200-700 MPa (Dembovski 1985). Substantial improvement in the long-term creep strength and corrosion resistance in various corrosive media are also observed. Reduction of metal halides and oxides to metal in hydrogen plasma offers many technological advantages. The atomic hydrogen and other active species present in the plasma medium considerably enhance the kinetics of the reduction reaction. The high residual hydrogen contents of metals gained by this technique can be brought down to acceptable levels by subsequent vacuum annealing. Zirconium tetrachloride vapor has been reduced with hydrogen to produce metallic zirconium in a hybrid DC transferred arc plasma reactor (Jones et al. 1997). This plasma reduction technique has been successfully applied to reduction of WO3 and MOO3. Arc plasma devices operating in the range of 150-250 kW have been used to produce tungsten and molybdenum. Similar oxide reduction processes have been extended to produce metals of Ni, Cr and other elements using an argon plasma medium containing additions of hydrogen and hydrocarbons.
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6.3.2.3 Processing o f ceramics High-temperature vapor phase chemistry in the thermal plasma medium offers an attractive route for powder synthesis. The process offers flexibility in the choice of the reactants, which can be solids, liquids or gases. The starting material is introduced in the form of a gas, liquid or solid. The reaction proceeds in several steps: vaporization, vapor phase reaction, nucleation and condensation. Carbides, nitrides, borides and oxides of refractory metals have been successfully produced in RF and DC torch plasma-based reactors. 6.3.2.3.1 Oxide powders. Plasma processing has been widely used to produce ultrafine powders of alumina, magnesia, titania and other oxide powders. A plasma-based technology has been developed for large-scale manufacture of pigment-grade titanium dioxide by the Tioxide Group in UK under the name ICON process. The process essentially makes use of a DC plasma torch to oxidize TIC14. The principal advantages of the plasma process over the conventional process are simplification of the process steps, narrow particle size range and precise control over product quality. Ultrafine silica is used as a filler material in composites and as a thixotropic additive material in the paint industry. The active silica product is prepared by vaporizing SiO2 in an oxidizing atmosphere followed by rapid quenching. The resulting powder is highly active and amorphous. Ultrafine aluminum oxide, iron oxide and zirconium oxide have been synthesized in an inductively coupled plasma (ICP) reactor (Kagawa et al. 1983). Nitrate solutions of the respective metals were introduced into the plasma jet, where decomposition of the fine droplets of the nitrate into the oxide occurred. The fast quench rate favored the formation of metastable ),-phase of A1203 from aluminum nitrate. The dissociation of zirconium nitrate in the ICP reactor resulted in the formation of ZrO2, which was predominantly in the metastable tetragonal phase. 6.3.2.3.2 Non-oxide ceramics. The carbides of silicon, titanium, boron, and tungsten are characterized by high hardness and excellent wear and corrosion resistant properties. These materials find application in the cutting tool industry and also as abrasives. The high energy density of the transferred arc has been used to produce metallic vapors that on condensation produce submicron aerosols. These vapors are also reacted in a suitable environment to produce a variety of products. Submicron size silicon carbide powders have also been produced with SIC14 and methane in a DC plasma torch (Allaire et al. 1991). Ultrafine powders of silicon nitride and aluminum nitride have also been produced using arc plasmas using ammonia as the nitriding gas (Ageorges et al. 1994). Processing of ultrafine powders of refractory carbides and nitrides is one of the widely studied areas of plasma processing. Most of the research efforts have been directed to the study of gas-phase reactions in plasma medium involving metal halides and suitable reactive gas to give ultrafine powders of refractory carbides and nitrides (Yoshida et al. 1979). Titanium carbide and tungsten carbide have been synthesized in DC and RF plasma reactors from titanium tetrachloride in a methane atmosphere. Ti and W metal powders have also been used to synthesize TiC and WC powders (Ananthapadmanabhan et al. 1994b). Plasma processing in RF plasma torches has also been used to prepare titanium
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nitride. The process involves the use of titanium halide or titanium metal powder with ammonia or nitrogen as the reactive gas. The liquid injection plasma synthesis (LIPS), where a liquid precursor is injected into the plasma jet, is widely used to process nanosized ceramic powders (Kong and Lau 1990). The liquid injection can be either parallel to the flow or counter current to the flow. Ultrafine oxide and non-oxide ceramic powders and mixed oxide compounds have been synthesized using the LIPS and counter current plasma synthesis (CCPS). Refractory carbides, nitrides, high temperature super-conducting oxide and mixed oxide powders have been produced in RF and DC based plasma reactor systems. 6.3.2.4 Treatment of hazardous wastes Conventional incinerators are, often, unsatisfactory for effective treatment and disposal of hazardous wastes, which are often mixtures of metals, organic and inorganic compounds. Production of toxic gases such as dioxin, insufficient thermal destruction of inorganic constituents, production of polluted flue gases, etc., are some of the drawbacks of conventional waste management systems. Thermal plasma technology has been shown to be potentially useful in the management of industrial waste (Boulos 1996), hospital waste, etc. Multi-torch plasma furnaces have been developed for destruction of highly toxic waste materials such as polychlorinated biphenyls (PCBs), pesticides, herbicides and organic solvents. These systems are also suitable for nuclear waste treatment (Ageorges et al. 1994).
6.3.3 Processing of metastable phases The unique feature of plasma processing is its ability to overcome thermodynamic and kinetic barriers and gain access to non-equilibrium regimes. The high temperature and the high quench rate (105-107 K/s) help in processing metastable and amorphous phases. As the reactants enter the plasma jet, they are heated to their melting temperatures and very often, an appreciable fraction of the particles vaporize and react with each other. The atomic and molecular clusters of the product are carried to the cooler regions of the reactor, where they are quenched. The entire process is completed in a few milliseconds. This time frame is too small for atomic rearrangement to form crystalline structures. It has been observed that thermal plasma processing of silica and other silicate minerals, such as zircon invariably produces highly reactive silica in an amorphous state as confirmed by X-ray diffraction of the product. Plasma process is capable of producing metallic glasses, amorphous silicon, diamond and other metastable phases of ceramic materials.
6.3.3.1 Plasma deposition of diamonds Conventionally, diamond is synthesized by recrystallization of metal-solvated carbon under high temperature (2,000~ and extreme pressures (50-100 kbar). Vapor-phase synthesis of diamond by the hot-filament chemical vapor deposition (CVD) technique has been known for many years. However, the technique is characterized by extremely low deposition rates. The use of a thermal plasma jet for diamond deposition is a recent development. Thermal dissociation of methane and other hydrocarbons in DC and RF plasma jets have been successfully applied to develop diamond coatings (Yoshida et al.
142
P. V. A n a n t h a p a d m a n a b h a n a n d N. Venkatramani
Figure 6.8 Surfacemorphologyof plasma deposited diamond
1979, Meyer et al. 1989, Metaxas 1996). In contrast to typical growth rates of 0.05/xm/h in hot filament CVD methods, the plasma process is characterized by high growth rates of the order of 900 #m/h (Ohtake et al. 1991). Ramachandran et al. (1996) have injected a mixture of methane and hydrogen gas into an argon plasma jet issuing out of a DC arc plasma torch operated in the non-transferred arc mode. Figure 6.8 shows the surface morphology of diamond deposited on silicon substrate. The cauliflower-like structure is characteristic of high growth rate. The Raman spectrum, indicating the characteristic Raman Shift of the diamond deposit, is shown in Fig. 6.9. The precursor gases employed consisted of a mixture of methane and hydrogen in the ratio 1:99. The experimental conditions strongly influence the diamond growth rate, both qualitatively and quantitatively (Baldwin et al. 1994). In particular, the substrate temperature and hydrogen/methane ratio are significant. Although the exact mechanism of diamond formation is still not understood, high temperature gas phase reactions under plasma environments involving methyl radicals is proposed to be the main operating mechanism.
6.3.3.2 Thermal plasma synthesis of ultrafine alumina powder A classic example illustrating the non-equilibrium nature of reactions occurring in the plasma medium is exemplified by plasma oxidation of aluminum powder (Ananthapadmanbhan et al. 1996). The experimental set-up used is shown in Fig. 6.4. Aluminum metal powder, stored in a powder feeder, was fed into the plasma using argon as carrier gas. Compressed air was injected into the reactor. The powder particles melt and vaporize and the molten droplets and vapors of aluminum react with oxygen, resulting in the
Thermal Plasma Processing
143
800
4330
cr~ t
600
400 ,,
A
15IXl
Wovenumbers ( cn~ I )
|
1000
Figure 6.9 Raman spectrum of diamond deposit.
Table 6.5 Typicalexperimental parameters Power input (kW) Arc voltage (V) Plasma gas Ar (LPM) Secondary gas N2 (LPM) Carrier gas Ar (LPM) Reactive gas air (LPM) Powder feed rate (g/min)
16 40 20 02 10 40 05
formation of aluminum oxide that is deposited on the reactor walls. Typical operating parameters are given in Table 6.5. X-ray diffraction results indicate that the powder consists of the metastable ),aluminum oxide. Formation of the ),-phase is consistent with the highly non-equilibrium nature of the process. Vapors and molten droplets of aluminum react with oxygen to form aluminum oxide, which crystallizes in the metastable ),-phase by virtue of the rapid quench rate associated with the process. Figure 6.10 is a scanning electron micrograph (SEM) of a typical sample of alumina powder collected from the main section of the reactor. Individual particles are not discernible, but agglomerates consisting of nanosized particles can be seen. A transmission electron micrograph (TEM) of a typical sample is shown in Fig. 6.11. The clusters are resolved into individual particles with a size ranging
144
P. V. Ananthapadmanabhan and N. Venkatramani
Figure 6.10. SEM photographof plasma synthesized alumina
from a few nanometers to about 30 nm. The spherical morphology of the particles is also evident from the figure.
6.3.3.2.1 Formation of },-alumina. The free energy minimization plot for the system A1O-N has been developed using the CSIRO thermopackage (Turnbull and Wadsley 1988). Figure 6.12 shows the thermodynamic equilibrium diagram for the system aluminumoxygen-nitrogen in the molar ratio 1:2:8 (close to the experimental conditions) under 1 atmosphere pressure. The Y-axis shows the concentration of the species in number of mols and the X-axis shows the temperature. Argon as the plasma gas is also included in the thermodynamic calculations. However, only the oxides of aluminum, atomic species of oxygen and nitrogen and aluminum vapor are shown in the figure. It can be seen from Fig. 6.12 that at temperatures greater than 5,000 K, which are commonly encountered in plasma reactors, the stable species present are aluminum vapour, atomic oxygen and atomic nitrogen. In the temperature region 3,500-4,500 K, A10(g) becomes increasingly stable. The other oxide species of aluminum such as A120, A102 and A1202 are also formed in this temperature zone, although at lower concentration levels, possibly by the following reactions: Al(g) + O(g) = A10(g)
(6.6)
145
Thermal Plasma Processing w
~
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. . ~.. .
.:,:
*
'
.:
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.
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~--
.
.
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Figure 6.11. Transmission electron micrograph of plasma synthesized alumina.
2Al(g) + O(g) = A120(g)
(6.7)
2Al(g) + 20(g) = A1202(g)
(6.8)
Al(g) + 20(g) = A102(g)
(6.9)
2A10(g) = A120(g) + O(g).
(6.10)
As the temperature falls below 4,500 K, oxygen molecules and nitrogen molecules (not shown in the figure due to their high concentration) start forming by the recombination of the respective atomic species. At temperatures below 3,500 K and above 2,500 K, liquid aluminum oxide is formed. Solid ot-A1203, which is the only thermodynamically stable form of alumina, results as the temperature drops to below 2,300 K, which is the
146
P. V. Ananthapadmanabhan and N. Venkatramani
1 0.9 0.8 0.7 0.6 o.s
~s)
N(g)
0.4 0.3 0.2 0.1 0 0 0 0 '~--
0 0 it) ~--
0 0 0 tN
0 0 rid r
0 0 0 03
0 0 ur) ~
0 0 0 ~I*
0 0 it') ~
0 0 0 IZ)
0 0 it') it')
0 0 0 t,,ID
T ~ ~ Figure 6.12. Free energy minimization plot for the system AI-O-N (1:2:8).
melting point of alumina. Below 2,300 K, the stable species are solid ot-A1203, nitrogen and oxygen molecules (not shown due to their high concentration). Thermodynamic considerations would lead us to conclude that ct-Al203 is formed. However, the conditions existing in the plasma reactor, augmented by thermochemical effects associated with the oxidation of aluminum, favor the formation of non-equilibrium phases. The first step leading to the formation of alumina is the melting and vaporization of aluminum powder injected into the plasma. Molten droplets and aluminum vapor react with oxygen downstream of the reactor to produce aluminum oxide according to the following equation: 2A1 + 3 0 2 -- A1203, AH ~ -- - 1675.7 kJ/mol.
(6.11)
Oxidation of aluminum to alumina is highly exothermic, the standard enthalpy of formation of aluminum oxide from the elements being -1675.7 4- 1.2 kJ/mol (Chase et al. 1985). The enthalpy of formation increases with the reaction temperature and is -2243.73 kJ/mol at 3,000 K. The heat released during oxidation coupled with the high temperature in the plasma medium vaporizes aluminum oxide, leading to the formation of suboxides of aluminum and if the temperature is higher than 4,500 K aluminum vapor and oxygen atoms are formed. As aluminum vapor and oxygen atoms move down to regions where the temperature is between 3,500-4,500 K, they react to form A10 gas and other oxides of aluminum. This is followed by a rapid quench of the vapors as they move to the cooler zones of the reactor tube. The quench rate in plasma processing has been estimated to be
Thermal Plasma Processing
147
of the order of 106-108 K/s and the extent of supercooling achievable is much greater than 0.2Tm, favoring homogeneous nucleation resulting in nanosized particles of ),-alumina. In principle, the selection of a metastable phase during solidification is associated with kinetic factors and requires suppression of the nucleation of the stable phase. (See also Chapter 2 on "Thermodynamics and Kinetics of Metastable Phase Formation" by K. N. Ishihara in this book). The nucleation rate of a solid condensing from liquid or gaseous phase is decided by its critical nucleation free energy, A G*, which is the energy barrier for nucleation. The relative nucleation rates of a and y phases of alumina depend on the ratio of their respective critical nucleation free energies, A G * / A G•*. McPherson (1973) has estimated the critical free energies of nucleation for a- and v-alumina as a function of temperature. The estimated ratio of the relative critical free energy for nucleation of a- and y-alumina indicates that v-alumina would nucleate in preference to a-alumina at temperatures lower than 1,740~ The estimated solidification temperature for alumina particles falls in the range of 1,500-1,600~ under plasma conditions, suggesting the formation of v-alumina in preference to a-alumina. Formation of v-alumina in preference to the thermodynamically stable a-phase is typical of rapid quench techniques such as plasma-spraying (Heintze and Uematsu 1992), arc discharge (Hirayama 1987) and electro-hydrodynamic atomization (Jayaram and Levi 1989). Spherical particles of alumina prepared by electro-hydrodynamic atomization have been shown to consist of metastable phases like y- and 6-alumina. Aluminum oxide prepared by vapor condensation from arc, flame oxidation of aluminum salts also consists of metastable phases rather than the thermodynamically stable form. However, the final phase obtained depends on the thermal excursion the particles experience during recalescence and subsequent cooling of the solid powder. This depends on the particle size and the cooling rate. The transformation of y-A1203 to a-A1203 for plasma-sprayed alumina droplets has been analyzed by McPherson (1981), who estimated the achievable supercooling and the thermal history of droplets of different sizes. The analysis shows that particles less than 1 # m in diamater consist entirely of y-A1203, whereas particles >50 # m are completely transformed to the a-phase. Levi et al. (1988) have developed a more rigorous model to examine microstructure evolution continuously through the liquid cooling stage, recalescence and subsequent cooling of the solid powder. Their analysis shows that decreasing the droplet size results in higher supercoolings and more efficient cooling leading to the formation of the metastable phase. The particle size range normally observed in plasma processed alumina is of the order of a few tens of nanometers suggesting that high supercoolings favor the formation of the v-phase. The above analysis helps to illustrate the physicochemical principles involved in the formation of non-equilibrium phases under plasma environment. The generalities of the model can be extended to any other specific case. 6.4. S U M M A R Y AND C O N C L U S I O N S
An overview of the recent developments in the field of thermal plasma processing has been presented. Thermal plasma processes are characterized by high temperature, high enthalpy and power density, 10-100 times higher than those encountered in conventional
148
P.V. A n a n t h a p a d m a n a b h a n
a n d N. V e n k a t r a m a n i
furnaces. Different plasma devices ranging from a few hundreds of k W to a few M W p o w e r levels have been designed and developed, over the past decade, to meet a variety of requirements ranging from p o w d e r synthesis to primary and secondary metallurgical operations and treatment of hazardous wastes. The use of plasma-sprayed ceramic and alloy coatings for industrial applications is a well established technique. But, there have been newer developments such as the reactive plasma-spray technique, which combine plasma processes and chemical reactions in one single operation. A further development in this area is the direct conversion method using cheap raw materials. The gas-to-solid conversion in plasma process to fabricate diamond coatings is a variant of this method. Contrary to conventional processing techniques, which are limited by thermodynamic and kinetic barriers, plasma processing, by virtue of the high temperature, high power density and rapid quench, offers unlimited scope for preparation of metastable phases and amorphous structures. The presence of reactive species provides new reaction paths to synthesize a variety of novel materials.
ACKNOWLEDGMENTS
The authors thank Mr. U. K. Chatterjee, Head, Laser & Plasma Technology Division for his constant support and encouragement. The cooperation and help received from Mr. K. P. Sreekumar and Dr. D. S. Patil are also gratefully acknowledged.
REFERENCES
Ageorges, H., Megy, S., Chang, K., Baronnet, J. M., Williams, J. K. and Chapman, C. (1994) Plasma Chem. Plasma Process., 13, 613. Allaire, E, Parent, L. and Dalleuie, S. (1991) J. Mater. Sci., 26, 4160. Ananthapadmanabhan, P. V. and Taylor, P. R. (1998)J. Alloys and Compounds (in press). Ananthapadmanabhan, P. V., Sreekumar, K. P., Muraleedharan, K. V. and Venkatramani, N. (1991) Surf. Coat. Technol., 49, 62. Ananthapadmanabhan, P. V., Sreekumar, K. P., Ravindran, P. V. and Venkatramani, N. (1993) Thin Solid Films, 224, 148. Ananthapadmanabhan, P. V., Sreekumar, K. P., Iyer, K. V. and Venkatramani, N. (1994a) Mater Chem. Phys., 38, 15.
Ananthapadmanabhan, P. V., Sreekumar, K. P. and Venkatramani, N. (1994b) Plasma Chemical Processing of Refractory Carbides, Plasma Science and Technology, ed. Das, I. M. L. (Allied Publ., New Delhi, India), p. 627. Ananthapadmanabhan, P. V., Sreekumar, K. P., Ravindran, P. V. and Venkatramani, N. (1995) in Thermal Spraying--Current Status and Future Trends, Vol. 2, ed. Ohmori, A. (High Temperature Soc. of Japan, Tokyo), p. 1127. Ananthapadmanabhan, P. V., Sreekumar, K. P., Venkatramani, N. and Taylor, P. R. (1996) J. Alloys and Compounds, 244, 70. Ananthapadmanabhan, P. V., Sreekumar, K. P., Venkatramani, N. and Taylor, P. R. (1997) in EPD Congress 1997, ed. Mishra, B. (TMS, Warrandale, PA), p. 209. Baldwin, Jr., S. K., Owano, T. G. and Kruger, C. H. (1994) Plasma Chem. Plasma Process., 14, 383. Bonet, C. (1976) Chem. Eng. Prog., 72, 63. Boulos, M. (1992) J. Thermal Spray Technol., 1, 33. Boulos, M. I. (1996) Pure Appl. Chem., 68, 1007. Bratton, R. J. and Lau, S. K. (1981) inAdvances in Ceramics, Vol. III, Science and Technology ofZirconia, eds. Heuer, A. H. and Hobbs, L. W. (American Ceramic Society, Columbus, OH), p. 226.
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Chase, Jr., M. W., Davies, C. A., Downey, Jr., J. R., Frurip, D. J., McDonald, R. A. and Syverud, A. N. (eds.), (1985) JANAF Thermochemical Tables, third edition (The American Chemical Society and The American Institute of Physics for the NIST, New York). Dembovski, V. (1985) Plasma Metallurgy: The Principles, Materials Science Monographs, No. 23, (Elsevier, Amsterdam). Eliezer, Y. and Eliezer, S. (1989) The Fourth State of Matter: An Introduction to the Physics of Plasma (Adam Hilger, Bristol, UK). Fauchais, E and Vardelle, M. (1994) Pure Appl. Chem., 66, 1247. Fauchais, P., Bourdin, E., Coudert, J. E and McPherson, R. (1983) in Topics in Current Chemistry, Vol. 107, eds. Veprek, S. and Venugopalan, M. (Springer-Verlag, Berlin/Heidelberg), p. 59. Fey, M. G., Wolf, C. B. and Harvey, J. E (1975) in Proceedings of the International Round Table: Transport Phenomena in Thermal Plasmas (Odeillo, France). Fineman, J. (1987) Plasma Technology in Metallurgical Industry (Iron and Steel Society, Warrendale, PA). Gerdeman, D. A. and Hecht, N. L. (1972) Arc Plasma Technology in Materials Science (Springer-Verlag, Wien). Heberlein, J. V. R. (1992) Pure Appl. Chem., 64, 629. Heintze, G. N. and Uematsu, S. (1992) Surf. Coat. Technol., 50, 213. Hench, L. L. (1991) J. Amer. Ceram. Soc., 74, 1487. Herman, H. and Bhat, H. (1980) in Synthesis of Metastable Phases, eds. Machlin, E. S. and Rowland, T. J. (TMS- AIME, Warrendale, PA), p. 115. Hirayama, T. (1987) J. Amer. Ceram. Soc., 70, C-122. Jayaram, V. and Levi, C. G. (1989) Acta. Metall., 37, 569. Jones, H. (1973) Rep. Prog. Phys., 36, 1425. Jones, R. T., Barcza, N. A. and Curr, T. R. (1997) Plasma Developments in Africa (Internal Report, Mintek, Randburg, South Africa). Kagawa, M., Honda, E, Onodera, H. and Nagae, T. (1983) Mater Res. Bull., 18, 1081. Knight, R., Smith, R. W. and Lawley, A. (1995) Inter J. Powder Metall., 31, 205. Kong, E C. and Lau, Y. C. (1990) Pure Appl. Chem., 62, 1809. Levi, C. G., Jayaram, V., Valencia, J. J. and Mehrabian, R. (1988) J. Mater Res., 3, 969. Lin, C. K. and Berndt, C. C. (1994) J. Thermal Spray Technol., 3, 75. Mackinnon, I. M. and Reuben, B. G. (1975) J. Electrochem. Soc., 122, 806. Madejski, J. (1976) Inter J. Heat and Mass Trans., 19, 1009. Matejka, D. and Benko, B. (1989) Plasma Spraying of Metallic and Ceramic Materials (John Wiley, New York). McPherson, R. (1973) J. Mater Sci., 8, 851. McPherson, R. (1981) J. Mater Sci., 16, 3141. McRae, L. B., Barcza, N. A. and Curr, T. R. (1985) in Proceedings of the International Conference on Mineral Science and Technology, Vol. 2, ed. Haughton (Council for Mineral Technology, Randburg, South Africa), p. 827. Meetham, G. W. (1988) Mater Design, 9, 308. Metaxas, A. C. (1996) Foundations ofElectroheat: A Unified Approach (John Wiley, Chichester, UK). Meyer, D. E., Dillon, R. O. and Woollam, J. A. (1989) J. Vac. Sci. TechnoI., AT, 2325. Miller, R. A., Smialek, J. L. and Garlick, R. G. (1981) inAdvances in Ceramics, Vol. III, Science and Technology ofZirconia, eds. Huer, A. H. and Hobbs, L. W. (American Ceramic Society, Columbus, OH), p. 241. Mohanty, M. and Smith, R. W. (1995) J. Therm. Spray Technol., 4, 384. Mutasim, Z. Z., Smith, R. W. and Mohanty, M. (1992) in Thermal Spray--International Advances in Coating Technology, ed. Berndt, C. C. (ASM International, Materials Park, OH), p. 1019. Notomi, A. and Hisatome, N. (1996) Pure Appl. Chem., 68, 1101. Ohtake, N., Mashimo, Y. and Yoshikawa, M. (1991) New Diamond Sci. Technol., 173, 178. Patil, D. S., Sreekumar, K. E, Venkatramani, N., Iyer, R. K., Ram Prasad, Koppikar, R. S. and Munim, K. R. (1996) Bull. Mater Sci. (India), 19, 115. Pauloski, L. (1995) The Science and Engineering of Thermal Spray Coatings (John Wiley, New York). Pfender, E. (1994) in Proceedings of the International Symposium on Heat and Mass Transfer, ed. Fauchais, E (Bell House, Turkey), p. 223. Ramachandran, K., Patil, D. S., Venkatramani, N., Biswas, A. R., Venkateswaran, S. and D' cunha, R. (1996) Indian Vac. Soc. Bull., 27, 23. Ramachandran., K., Selvarajan, V., Ananthapadmanabhan, P. V. and Sreekumar, K. P. (1997) Plasma Devices and Operations, 5, 1.
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Rangaswamy, S. and Herman, H. (1986) Advances in Thermal Spraying. (Pergamon Press, Oxford, UK), p. 101. Ranz, W. E. and Marshall, W. R. (1952) Chem. Eng. Prog., 48, 141. Rykalin, N. N. (1976) Pure Appl. Chem., 48, 179. Rykalin, N. N. and Kudinov, V. V. (1976) Pure Appl. Chem., 48, 229. Scott, K. T. and Kingswell, R. (1991) in Advanced Surface Coatings, eds. Rickerby, D. S. and Matthews, A. (Blackie, London). Stecura, S. (1977) Amer. Ceram. Soc. Bull., 56, 1082. Taylor, E R., Manrique, M., Pirzada, S. A. and Abdel-Latif, M. (1995) Plasma Chem. Plasma Proc., 15, 545. Turnbull, A. G. and Wadsley, M. W. (1988) The CSIRO Thermochemistry System, version 5 (IMEC, Australia). Venkatramani, N. (1995) Bull. Mater. Sci. (India), 18, 741. Yoshida, T., Kawasaki, A., Nakagawa, K. and Akashi, K. (1979) J. Mater. Sci. 14, 1624. Zaat, J. H. (1983) Ann. Rev. Mater. Sci., 13, 9. Zoltowski, P. (1968) Rev. Int. Htes. Temp. et. Refract., 6, 65.
Chapter 7 Spray-Forming 7.1. 7.2. 7.3.
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 7.4. Applicability 7.5. 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 7.5.2.4 Refinement 7.5.3 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
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Chapter 7 Spray-Forming B ING LI and ENRIQUE J. LAVERNIA
7.1. I N T R O D U C T I O N
Materials science and engineering is a very broad, interdisciplinary field which encompasses several sciences, primarily chemistry, physics, and metallurgy. The materials under investigation include metals, semiconductors, ceramics, polymers, glasses, and composites. The areas of study range from the nature of atomic arrangement to the phenomenology of the macroscopic material properties; and from elegant fundamental theories to exciting practical applications. Among the manifold aspects, an everlasting important philosophy in materials science and engineering is to produce a material having the required physical and mechanical properties at the lowest cost. To fulfill this goal, it is essential to explore advanced materials with superior properties while simultaneously maintaining a competitive cost. This is made possible by development/utilization of novel processing techniques. Each of the above mentioned categories of materials has its own unique processing requirements, either conventional or novel. For metallic materials, there are mainly two most frequently utilized conventional processing techniques, i.e., casting (ingot casting, die casting, investment casting, etc.) and powder metallurgy. Meanwhile, many novel processing techniques are being currently explored/investigated, such as mechanical alloying (Koch 1989), laser surface modification (Steen 1985), ion implantation (Picraux 1984), self-sustained high-temperature processing (Mabuchi et al. 1995, Alman et al. 1995, Rawers and Wrzesinski 1992), and spray-forming, to name a few. Each of these novel processing techniques has its own advantages and disadvantages. Among these many techniques, spray-forming is especially attractive in the sense that it is potentially capable of fabricating advanced materials with improved properties at costs that are industrially competitive. The improved properties are directly associated with the non-equilibrium nature of spray-forming, which is the focus of this chapter and will be discussed in detail in the following sections. The cost effectiveness is derived from the inherent simplicity of this technique, which, in turn, is partially attributable to the non-equilibrium phenomena. It is instrumental to illustrate its inherent simplicity by comparing spray-forming with other processing techniques, especially the two well-established processing techniques: powder metallurgy and ingot casting. Compared with powder metallurgy, spray-forming combines powder production, sieving, canning, cold pressing, degassing and hot press153
154
Bing Li and E. J. Lavernia
ing into a single step. This is inherent in the concept of spray-forming which incorporates consolidation into the process of powder production. Compared with the conventional ingot casting, spray-forming eliminates macrosegregation and minimizes microsegregation, hence avoiding many intermediate forming steps, such as long term homogenization treatment and extensive thermomechanical processing, which are typically required in conventional ingot casting techniques to eliminate/minimize segregation. The concept of spray-forming was first introduced by Singer in the early 1970s (Singer 1970, 1972). His initial intent was to develop a processing technique which was capable of reducing the overall cost in manufacturing structural materials. It was soon realized that spray-forming is also unique in obtaining superior microstructures and hence improved mechanical properties. Because of the many potential advantages of this technology over conventional processing techniques, spray-forming aroused more and more interest in the academic field, as well as in the industrial community worldwide over the last three decades. There are now more than 20 research organizations and 20 industrial companies located in the United States, Europe, and Eastern Asia (China, Japan, Korea, etc.) actively involved in the development/utilization of this technology (Leatham 1993). These developmental activities are coupled with a large volume of publications in the field of spray-forming, including journal articles, conference papers, book sections (Yule and Dunkley 1994), and books (Lavernia and Wu 1996). These efforts may be categorized as: (a) understanding of the fundamental phenomena associated with this technology, such as heat transfer, undercooling, and solidification; (b) application of this technology to new alloy systems rather than the widely studied aluminum alloys, such as Cu, Fe, Ni, Ti, Zr, and Mo based alloys; (c) optimization of the processing parameters to achieve better microstructural characteristics and mechanical properties; (d) modification of this technology to extend its application domain from monolithic alloys to metal matrix composites; (e) novel design of this technology to enhance its capability in the control of product geometry (near net shape, such as flat sheet, circular tube, etc.); and (f) commercialization of this technology.
7.2. P R I N C I P L E S
The principle of spray-forming is schematically shown in Fig. 7.1. It consists of two stages: spray atomization and spray deposition. Accordingly, spray-forming is also termed as spray atomization and deposition in the academic field as well as the industrial community. The spray atomization stage further contains the following detailed physical changes: (a) A stream of molten metallic material is impinged upon by highly energetic gas jets in different, normally symmetrically opposite directions, as shown in Fig. 7.1. The result is disintegration of the molten metallic stream into small, irregular ligaments. (b) The small, irregular ligaments transform into spherical droplets shortly after their formation (Lavernia et al. 1992). This change is normally termed spheroidization in
Spray-Forming
155
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Deposition substrate
Figure 7.1. Two stages of spray-forming: sprayatomization and spray deposition.
spray-forming. The driving force for this change comes from the decrease of the surface energy associated with the ligament (minimum energy principle). The surface energy is the product of surface tension and surface area. The surface tension is a constant for the same material, while the surface area depends on the shape. For a given volume, a sphere has the minimum surface area. Accordingly, there is always a tendency for any shape to change into spherical to minimize the total surface energy. In other words, formation of spherical droplets from irregular ligaments is thermodynamically favorable. In a practical situation, whether spheroidization really occurs depends on the associated kinetics. It has been demonstrated, both experimentally and theoretically, that spheroidization can be completed almost immediately after the formation of the irregular ligaments during spray-forming, mainly because of the high surface tension of most molten metallic materials and the size of the ligaments. (c) The sizes of the spherical droplets are normally on the order of microns, but typically fall into a wide range. Ample experimental evidence suggest that, in any spray-forming experiment, the size distribution of these droplets obeys a specific distribution called log-normal distribution (Lubanska 1970, Liang et al. 1992a, Marinkovich et al. 1989, Irani and Callis 1963, Allen 1981, Grant et al. 1993a, Delplanque et al. 1997). Most of the droplets have sizes close to a certain finite value (median size). When the size deviates further from this value, the probability of finding a droplet having this size decreases. (d) The atomization gas flow field forms a specific geometrical shape, normally a spray cone. The micron-sized droplets are dispersed in the spray cone. (e) The spherical droplets experience momentum transfer with the atomization gas flow field. This momentum transfer governs the movement of the droplets in the spray
156
Bing Li and E. J. Lavernia
cone, including trajectory, flight time, and subsequent deformation behavior during the deposition stage. (f) The spherical droplets experience thermal exchange with the atomization gas flow field. This thermal exchange causes the spherical droplets to cool down and solidify during their flight towards the substrate. The thermal exchange is affected, to a certain extent, by the momentum transfer between the droplet and the gas field. The spray deposition stage involves collection of the droplets in the spray cone onto a water-cooled substrate, forming a dense bulk material that is normally referred to as a deposit: (a) By the time when these droplets reach the substrate, they are in the solid, liquid or semi-solid states. (b) At the beginning of the deposition, the droplets impinge directly on the top surface of the substrate. Once these droplets are collected on the substrate, they act as the substrate themselves for the subsequent impinging droplets. (c) Impingement of the droplets with the substrate leads to extensive modification of the droplets, including their morphology and the associated microstructural features, such as dendrites (see Fig. 7.2), which are developed during solidification in flight. This deformation plays a critical role in the evolution of the microstructure of the deposit. (d) The deposited bulk material experiences thermal exchange with the water-cooled substrate and the surrounding environment, leading to the dissipation of its thermal energy. Simultaneously, during spray deposition, the arrival of droplets from the spray cone onto the substrate continuously incorporates thermal energy into the deposited bulk material. These two effects compete with each other, determining the overall thermal behavior of the deposit. This further affects some aspects of the microstructure, such as grain size. The microstructure in turn determines the mechanical properties. Spray-forming is normally carried out in an environmental chamber. The chamber is evacuated and back-filled with a protective gas such as N2, He, and Ar. The purpose of this arrangement is to enhance the capability to either suppress or exploit the possible reactions, such as oxidation, between the atomization/environmental gas and the metallic materials (the droplets and the deposit). The products of spray-forming include the bulk deposit and powders that are usually termed as oversprayed powders. During the deposition stage, some powders (solidified droplets) cannot be collected into the substrate to form the bulk deposit. Instead, they are carried by the gas flow field into a collecting apparatus (cyclone). These powders are termed as oversprayed powders in spray-forming technology. In commercial applications, the bulk deposit is the main product of practical value, while the oversprayed powders must be minimized in quantity. In scientific research, both the bulk deposit and the oversprayed powders are useful in providing valuable information on the understanding of the fundamentals in spray-forming.
Spray-Forming
Figure 7.2. SEM
157
micrograph showingdendrites formed in spray atomized A1-Mn-LiZr powders (Lavernia et al. 1992).
7.3. VARIATIONS AND D I S T I N C T I O N S
Advances in our understanding of the science and technology of spray-forming have led to the development of novel designs, as well as modifications to the original technology, although the fundamental principle remains the same. A comprehensive discussion should encompass a description of these advanced designs. Since spray-forming is currently a very active field, an exhaustive description of all these variations is difficult. Accordingly, in the following sections, several representative variations will be described. It is also worth noting that there is some confusion in the literature regarding the term sprayforming, as applied to related manufacturing technologies. This issue is addressed in the sections that follow. 7.3.1 Variations
Most of the variations to spray-forming involve modification of one or more of the key components of a spray-forming facility, which include the following: crucible, atomizer, atomization gas, and substrate, as will be described in sequence below.
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Bing Li and E. J. Lavernia
7.3.1.1 Crucible: induction skull melting~spray-forming In spray-forming, there is generally a crucible used to melt the master alloy and then to hold the molten metal. During spray-forming of non-reactive alloys with low to intermediate melting temperatures, such as aluminum and magnesium based alloys, ferritic steels, and nickel based intermetallics, the crucible is generally made of graphite or ceramic materials such as zirconia and alumina. For reactive/refractory alloys, these conventional crucibles are no longer usable. In spray-forming, the molten metal is generally heated to a temperature about 50 ~ 100 K higher than its melting temperature and then maintained there for a significant time duration (> 10 min). Such a processing procedure is essential to ensure smooth flow of the molten metal through a delivery tube into the atomizer or, in the worst scenario, to avoid freeze-up of the molten metal at the delivery tube. For reactive alloys, such as Ti (Morris et al. 1992, Li et al. 1996a, Li and Lavernia 1997) and Zr (Li et al. 1997) based alloys, their highly reactive nature in the molten state will lead to significant reaction between the molten alloys and the crucible material during such a long period, making selection/utilization of conventional crucibles extremely difficult if not completely impossible. For refractory alloys, such as Zr (Li et al. 1997), Mo, Nb, V, Ta (Delplanque et al. 1997), and W based alloys, their melting temperatures (>1800~ are normally close to or above those of the commonly used crucible materials (zirconia or graphite). Accordingly, the crucible will be destroyed before the alloys can be melted. To circumvent these problems, the facility has to be significantly modified, and an alternate melting approach, namely skull melting, has to be incorporated. The novelty in the concept of skull melting/spray-forming may be best illustrated by a comparison between the usage of the conventional crucible and the skull crucible (refer to Fig. 7.3). In spray-forming experiments that employ a conventional crucible, the crucible is charged with a master alloy and then heated up using an induction furnace. In the process of heating, the crucible is thermally insulated from its environment, preventing unnecessary loss of thermal energy and enhancing melting efficiency. Consequently, both the crucible and the charge remain at an elevated temperature as they are heated up. In skull melting/spray-forming, the crucible is water-cooled during the entire process of heating of the charged material. When the center of the charge is melted, its exterior regions remain solid because of their intimate contact with the cold crucible. These solid regions of the charged material form a crucible of the shape akin to the water-cooled crucible. This crucible is called a "skull crucible", which leads to the name of this technique. The most intriguing feature of skull melting technique is that the molten metal will be contained in a crucible (the skull crucible) made of the metal itself; hence contamination resulting from the reaction between the molten metal and the crucible is minimized, if not totally eliminated. 7.3.1.2 Atomizer: circular vs. linear spray-forming
In terms of the geometrical shape of the atomizer, there are two typical types of sprayforming techniques: circular and linear. In a circular atomizer, the gas jet nozzles are arranged concentrically (Fig. 7.4a). The number, cross-section area, and shape of the nozzles are critical parameters that control the efficiency of the atomizer. During sprayforming, the atomization gas streams exit these nozzles, forming a circular cone (Fig. 7.5b). The deposit produced this way normally exhibits a bell shape (Fig. 7.5b). In a
Spray-Forming
159
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linear atomizer, the gas jets are located along two parallel lines (Fig. 7.4b). During sprayforming the atomization gas forms a prism-like geometry, giving rise to a deposit with a high aspect ratio (Fig. 7.5a). Hence, a linear nozzle, in combination with a moving substrate, may be readily used to attain a sheet-like geometry (Fig. 7.5a).
7.3.1.3 Atomizer: close-coupled vs. free-fall spray-forming Schematic diagrams illustrating these two types of spray-forming are shown in Fig. 7.6. In close-coupled spray-forming, the molten metal is released from the crucible through a delivery tube. When it exits the tube, the high-speed atomization gas jets impinge on the melt stream at the tip of the tube (Fig. 7.6a). In free-fall spray-forming, the molten metal is released from the bottom of the crucible, forming a stream. The molten metal stream travels downwards, through the atomizer unit, until it is atomized at a certain point below the atomizer. From the point it exits the bottom of the crucible to the point it is atomized, the molten stream is not constrained by any tubes as in close-coupled atomizer (Fig. 7.6b). Instead, it falls freely, which leads to its name: free-fall sprayforming. Close-coupled spray-forming and free-fall spray-forming are complementary to each other. Close-coupled spray-forming is more susceptible to freeze-up problems. Since the delivery tube is in close contact with the cold atomization gas, there is a higher probability for the melt stream to freeze at the tip of the tube. When this happens, further delivery of the molten metal through the tube is impossible, and the entire experiment fails and has to be restarted. Free-fall spray-forming can avert freeze-up problem because of the absence of the delivery tube. A concomitant problem is the instability of the molten stream in the absence of the delivery tube. The molten metal stream may be easily deflected by the turbulence in the surrounding gas envelope. This instability may lead to oscillation and asymmetry of the atomization cone (Lavernia and Wu 1996).
160
Bing Li and E. J. Lavernia
Gas jet nozzles Orl ac(
delivery tube
(a)
Space to accommodate delivery tube
Linear atomizer twins
~as jet slits
(b) Figure 7.4. Arrangementof atomization gas jet nozzles (slits) in circular and linear atomizers: (a) circular, and (b) linear.
7.3.1.4 Atomization gas: reactive spray-forming There are two types of atomization gases: protective and reactive. Helium and argon are common protective gases, while 02, CO, CO2 and H2 are normally categorized as reactive gases. Nitrogen (N2) is either protective or reactive, depending on the nature of the material being synthesized. In many situations, the atomization gas employed is protective. In a recently developed variation, i.e., reactive spray-forming (Zeng et al. 1992, Zeng et al. 1995), the atomization gas is deliberately selected to contain certain amount of reactive gases. Reactive spray-forming is mainly utilized to synthesize a special category of high strength materials: dispersion-strengthened alloys. The most representative example in this category is the oxide dispersion-strengthened (ODS) material. Its earliest application may be dated back to 1910 when Coolidge incorporated thorium oxide into tungsten wire to improve its elevated temperature mechanical properties (Coolidge 1910). The first
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350~ 500~ 600~ Ni25Mo75~ > Amor.~ > hcp-II~ > fcc-II 60 min 30 min 30 min Apparently, the phase evolution sequences of these two reactions are in good agreement with that deduced from the calculated free energy diagram, thus confirming the applicability of the Miedema's model in this respect (Zhang and Liu 1994a). The author's group has also performed similar annealing experiments for several positive A H f systems and these results are also in support of the calculations (Chen and Liu 1996). The use of calculated free energy diagrams can explain the alloy phase formation upon IM, as discussed in the previous sections by some representative examples.
8.7. I N T E R F A C E - G E N E R A T E D
S O L I D - S T A T E V I T R I F I C A T I O N IN S Y S T E M S W I T H A
P O S I T I V E H E A T OF F O R M A T I O N
In 1983, Schwarz and Johnson reported for the first time that solid-state reaction (SSR) of multilayers consisting of two different crystalline metals (La and Au) can lead to the formation of an amorphous alloy. They argued that spontaneous vitrification was expected to take place in metal systems with a negative A H i and one metal diffusing anomalously fast in the other. However, the author's group has found very recently that in some binary metal systems with positive A H f , vitrification could be achieved by SSR, e.g., in the HfTa (Liu et al. 1996), Y-refractory metals (Chen and Liu 1997) and Cu-Ta (Pan et al. 1998) systems. Consider the Y-Nb system as an example, Figure 8.8 shows the X-ray diffraction patterns of the Yv2Nb28 multilayered films before and after thermal annealing. It should be emphasized that similar solid-state vitrification was also observed by author's group very recently in some other systems with a positive A H f (Chen and Liu 1997a). To test whether a large size difference is a decisive prerequisite in solid-state amorphization, the Au-Ta system with a negative A H f o f - 4 9 kJ/mol (Pan et al. 1995) was selected as the atomic radius ratio and the volume ratio between Au and Ta are 1.008 and 0.936, respectively. When the Auz3Ta77 multilayers were heated up continuously at a rate of 5 K/min to 200~ partial amorphization was observed. After maintaining at 200~ for 30 min, the annealing temperature was raised to 350~ and kept for 30 min. The structure of the multilayers turned entirely into an amorphous state. When the temperature was further raised to about 650~ and kept for 20 min, diffraction lines from metals Au and Ta emerged again, indicating crystallization of the Auz3Ta77 amorphous alloy. These results suggest that solid-state vitrification can also be achieved in a system with a small atomic size difference.
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Figure 8.8. X-ray diffraction patterns of the Y72Nb28 multilayered films with 121 layers and a total thickness of 350 nm before and after thermal annealing: (a) as-deposited; (b) amorphized by annealing at 300~ for 1.5 hours.
8.8. C O N C L U D I N G R E M A R K S
(1) In the past 15 years, study of IM of multilayered films scheme has provided extensive data, which promoted significantly the non-equilibrium processing of materials. IM has so far synthesized a great number of new amorphous as well as new MX alloys. It is anticipated that this technique is capable of synthesizing more new metallic glasses in those not yet studied systems known as non-glass-forming ones. (2) Miedema's theory was employed for free energy calculation and steady-state thermal annealing results showed that the calculated energetic sequence of the related phases was relevant, at least in its outline, suggesting the applicability of Miedema's theory in this respect. (3) Taking into account the interfacial free energy, a unified thermodynamic model was proposed for predicting the possibility of metallic glass formation by IM in binary metal systems with either negative or positive AHz, and for a reasonable explanation for the formation of MX phases. (4) The interface concept was also tested by solid-state reaction experiments in some positive A Hf systems and it turned out that metallic glasses were indeed obtained by thermal annealing in the properly designed multilayered films with enough fraction of interfacial atoms. These results offer a neat evidence demonstrating the decisive role of the interfacial free energy in amorphization taking place in the multilayered films.
220
B. X. Liu
(5) Kinetically, due to the very restricted conditions available in IM for a crystalline phase to nucleate and grow, only simple-structured phases can be formed and they feature a close crystallographic correlation with their matrix. (6) Based on the IM and SSR studies, a new glass-producing technique is developed, i.e., the multilayer technique, in which the multilayered films are firstly designed according to the interfacial free energy calculation and then either by thermal annealing or ion irradiation. (7) Further studies are still needed in this area, especially in the systems with a positive A Hf. For instance, detailed growth of the amorphous phase, local microstructures and the abnormal properties of the new metallic glasses formed in such systems are highly interesting topics to be investigated by IM as well as SSR technique.
ACKNOWLEDGMENTS
The author is grateful to his current and previous students for their continuous contributions to this research project, especially to Mr. Y.G. Chen for his substantial assistance in writing and finalizing this chapter. Special thanks go to the Editor of this book, Professor C. Suryanarayana, for his critical reading of the manuscript of this chapter and for many helpful suggestions. The continuous financial support from National Natural Science Foundation of China under several projects and by the Administration of Tsinghua University are greatly appreciated.
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B. X. L i u
Pan, E et al., (1998) unpublished data. Pan, E, Chen, Y. G. and Liu, B. X. (1995) AppL Phys. Lett., 67, 780. Pan, E, Chen, Y. G., Zhang, Z. J., and Liu, B. X. (1996) J. Non-Cryst. Sol., 194, 305. Peiner, E. and Kopitzki, K. (1988) Nucl. Instrum. Methods Phys. Res., B34, 173. Rauschenbach, B. (1986) Appl. Phys., A40, 47. Rauschenbach, B. and Hohmuth, K. (1987) Nucl. Instr. Methods Phys. Res., B23, 323. Schmid, A. and Ziemann, P. (1985) Nucl. Instr. Methods Phys. Res., B7/8, 581. Schwarz, R. B. and Johnson, W. L. (1983) Phys. Rev. Lett., 51, 415. Shin, S. M., Ray, M. S., Rigsbee, J. M. and Greene, J. E. (1983) Appl. Phys. Lett., 43, 249. Takayama, S. (1976) J. Mater. Sci., 11, 164. Thompson, M. W. (1969) Defects and Radiation Damage in Metals. (Cambridge University Press, Cambridge). Townsend, E D., Kelly, J. C. and Hartley, N. E. W. (1976) Ion Implantation, Sputtering and their Applications. (Academic Press, London), p. 31. Tsaur, B. Y. (1980) Thesis, California Institute of Technology, Pasadena, CA. Tsaur, B. Y., Lau, S. S., Hung, L. S. and Mayer, J. W. (1981) Nucl. Instr. Methods, 182/183, 67. Turnbull, D. (1955) Impurities and Imperfections. (Amer. Soc. Met., Cleveland), p. 121. van Rossum, M., Shreter, U., Johnson, W. L. and Nicolet, M.-A. (1983) in Proceedings of MRS Fall Meeting, Boston, MA. von Allmen, M. (1983) in Laser-Beam Interaction with Materials, ed. Carl, J. (Washington International Society of Optical Engineering, Washington), p. 140. Yan, Z. H., Liu, B. X., and Li, H. D. (1986) Phys. Stat. Sol., (a) 94, 483. Yan, Z. H., Liu, B. X., and Li, H. D. (1987) Nucl. Instr. Methods Phys. Res., B19/20, 700. Zhang, Z. J. (1995) Ph.D. Thesis, Tsinghua University, Beijing, China. Zhang, Z. J. and Liu, B. X. (1994a) J. Appl. Phys. 76, 3351. Zhang, Z. J. and Liu, B. X. (1994b) J. Appl. Phys., 75, 4948. Zhang, Z. J. and Liu, B. X. (1994c) J. Phys.: Condens. Matter, 6, 2647. Zhang, Z. J. and Liu, B. X. (1994d) J. Phys.: Condens. Matter, 6, 9065. Zhang, Z.J. and Liu, B. X. (1995) Phys. Rev., B51, 16475. Zhang, Z. J., Bai, H. Y., Qiu, Q. L., Yang, T., Tao, K. and Liu, B. X. (1993) J. Appl. Phys., 73, 1702. Zhang, Z. J., Jin, O. and Liu, B. X. (1995) Phys. Rev., B 51, 8076.
Chapter 9 Physical Vapor Deposition 9.1. Introduction 9.2. Development of PVD 9.3. 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
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Chapter 9 Physical Vapor Deposition JOHN S. COLLIGON
9.1. INTRODUCTION Physical Vapor Deposition (PVD) is the general name given to coating processes where the transport of material to the substrate is effected by a physical driving mechanism. Such mechanisms include evaporation, sputtering, ion-plating, and ion-assisted sputtering. Often chemical reactions do take place to form coatings of oxides, nitrides and carbides, but the transport process does not generally rely on chemical reactions. The coated samples can usually be maintained at a temperature between ambient and 100~ during the deposition and the working pressure is below 10 -2 Pa. PVD thus differs quite markedly from Chemical Vapor Deposition (CVD) which occurs at atmospheric pressure, or just below, and at temperatures normally in excess of 500~ CVD also requires the use of complex gaseous species (such as fluorides, organometallics) which need to be handled very carefully whereas PVD requires either no gas, for evaporation, or, relatively simple species (Oxygen, Nitrogen) for most of the other PVD processes. Recent improvements in PVD techniques, in particular the use of plasmas and the unbalanced magnetron for the deposition process, have led to its wide application in industry. Such coatings are used for corrosion protection, to provide hard surfaces for engineering applications, to form decorative coatings on a range of household and industrial products, for energy-saving window glass, for media storage systems such as compact discs and for a range of microelectronics applications. Table 9.1 lists some of the conference series that cover this huge field. The present chapter starts with a brief historical review of the subject. It then describes the main PVD methods and continues with a discussion of the relation between deposition conditions and film properties. Some of the main applications of PVD coatings are then described and the chapter concludes with an outline of some of the possible future trends in this important industrial field.
9.2. D E V E L O P M E N T
OF P V D
The first recorded mention of a PVD process was by Grove (1852) in 1852 who observed the formation of a deposit in a glass tube in which a discharge had been running and identified this as being the same material as one of the electrodes; this was the first reported evidence of the ejection of atoms from an ion-bombarded material, now known as sput225
226
J. S. Colligon
Table 9.1. Established conference series containing papers on PVD. 9 9 9 9 9 9 9 9 9 9 9 9 9 9
InternationalConference on Metallurgical Coatings and Thin Films (ICMCTF) AnnualTechnical Conference of the Society of Vacuum Coaters E-MRSSpring Meeting InternationalConference on Crystal Growth and International Conference on Vapor Growth and Epitaxy InternationalVacuum Congress and International Conference on Solid Surfaces InternationalConference on Thin Films InternationalConference on Nanometer-scale Science and Technology (Nano) InternationalConference on Ion Beam Modification of Materials (IBMM) InternationalConference on Surface Modification of Metals by Ion Beams (SMMIB) InternationalConference on Advances in Surface Engineering EuropeanConference on Diamond, Diamond-like Materials, Nitrides and Silicon Carbide (Diamond) InternationalConference on Plasma Surface Engineering (PSE) NationalSymposiumof the American Vacuum Society MRS Fall Meeting
tering and one of the main PVD processes. A few years later Faraday (1857) exploded metal wires in vacuum and thereby produced an evaporated film. Many years then passed before the concept of ionising some of the depositing species was proposed by Berghaus (1938) and, subsequently exploited by Mattox (1963) in the process known as ion-plating. Ion-plated coatings were seen to have superior adhesion properties which were identified with the additional energy imparted to some of the evaporated atoms which had been ionized and to a general increase in energy of support gas and other species bombarding the growing film. The process operated at a relatively high pressure (0.5 Pa) however, so that many collisions occurred in the gas phase which tended to reduce the energies of ionized species and increase those of neutral particles. The many collisions allowed depositing species to be bounced around shaped substrates to the rear side, giving the technique so-called high throwing power, but the parameters were not easy to control. Weissmantel (1976) proposed the use of ion beams in the deposition process where material was either evaporated onto a surface being simultaneously bombarded by ions, or was sputtered by one ion beam and the deposited film bombarded by a second beam. Here the precise control of the added energy and ion to atom arrival ratio could be achieved. A review of the early work was given by Weissmantel et al. (1979). This type of system, further exploited by Harper et al. (1982a, 1982b, 1983), Sainty et al. (1984), Martin et al. (1983, 1984), Colligon (1995), Kheyrandish et al. (1994) and many others (Rossnagel and Cuomo 1988), has allowed controlled studies of the influence of added energy on film composition, stress, density, microstructure, hardness, wear and durability to be made. An alternative approach has been proposed by Yamada and Takagi (1995) who developed an ion cluster-beam method. In this method, illustrated schematically in Fig. 9.1, a beam of clusters of atoms from 2 - 2 0 in each cluster is created by expanding a vapor of the metal through a nozzle. Some atoms in the clusters are then ionized by an electron beam so that the whole cluster may be accelerated toward the negative substrate. How-
Physical Vapor Deposition t
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ce material Figure 9.1. Schematic diagram of ionized cluster-beam system
ever, all the above ion beam systems have a common limitation in that they do not allow the production of coatings over large areas (normally a maximum of a few square centimeters) nor do they allow high growth rates to be attained (normally maximum rates are 0.2-1 nm/s). There are some developments in the cluster-beam method by Gspann (1993) who has recently produced cluster-beam sources with much improved deposition rates but the treated areas are still of order of a few square centimeters. The major breakthrough in industrial processing has however come from the new generation of unbalanced magnetron systems (Window and Savvides 1986a, b, Savvides and Window 1986) which have, in the 1990s, replaced many simple evaporation methods and ion-plating techniques which had been the main coating methods during the period from the 1960's. Film deposition rates of 10-500 nm/s are now possible and the sample area can be several square meters. Similar rates of deposition, but not such large area capability, have been found for recent plasma-based systems, including the pulsed plasma methods (Conrad et al. 1990). All these systems can be developed to produce reliable coatings and, in some cases, can produce metastable phases of alloy materials because the energetic atoms are suddenly trapped into rapidly cooling surroundings. It is not always possible to relate the film properties to the deposition parameters because there are complex interactions between magnetron power, system pressure, ion/atom arrival ratio and deposition energies. Fundamental studies on new coatings are, therefore, often best undertaken on ion-beam systems.
J. S. Colligon
228 \ rH.:..
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9.3. D E P O S I T I O N M E T H O D S
This review will outline the main features of the various PVD methods but will concentrate more on the more recent ion-assisted deposition (IAD) techniques where additional energy is provided to the growing film by ion-bombardment. Earlier studies of traditional evaporation and sputtering methods have been fully reviewed in several textbooks by Holland (1956), Chopra (1969), Vossen and Kern (1978) and Bunshah (1982).
9.3.1 Evaporation In this process materials are vaporised by heating using direct resistance-heating, radiation, eddy current heating, electron-beam heating, or running an arc discharge. One advantage of the technique is that the evaporation can be carried out in a UHV chamber so that contamination of the substrate by residual gases is minimal. For best purity an electron-beam heating system should be used as the material can then be evaporated from the center of the crucible in which it is held and reactions with the crucible walls are minimised. A schematic diagram of a typical e-beam coating system is shown in Fig. 9.2. Fundamental equations for evaporation rates and tables of evaporation temperatures and procedures are given in Chapter 4 of the book edited by Bunshah (1982). Some decomposition of chemical compounds always occurs and special procedures have to be followed if fully stoichiometric films are to be formed. The energies of species arriving at the substrate are only of the order of 0.05-0.1 eV which means they have limited energy for surface migration to find good nucleation sites. Substrates must be free of contaminants to allow reasonable levels of film adhesion to be attained so considerable attention has to be paid to cleaning procedure. Because of the lack of surface mobility
Physical Vapor Deposition
229
there is a high probability that the coating will contain voids and will be less dense than the bulk material. This leads to the formation of stress in the coating, usually tensile, which often results in the film delaminating and curling up.
9.3.2 Molecular beam epitaxy (MBE) A special technique of evaporation has been developed, primarily for the semiconductor industry, in which component species are injected into a UHV chamber from gas effusion cells and form the desired film on a substrate. This is known as Molecular Beam Epitaxy. In some systems organometallic gases are used which react at the high temperature substrate to form the desired deposit and a residual volatile species. Following the work of Giinther (1958) an alternative three-temperature method was developed which allowed III-V semiconductor compounds to be deposited from the pure constituent elements; usually the type-III element cell was run at a higher temperature than that for the more volatile type-V element and the substrate temperature was chosen to be high enough to re-evaporate excess type-V species. A review of fundamental aspects of MBE has been published by Foxon and Joyce (1981). In later work an important correlation between the oscillations in intensity of RHEED signals and layer coverage was discovered, allowing the number of layers deposited to be calibrated in situ (Neave et al. 1983). The whole field is well reviewed by Herman and Sitter (1989) and Chow (1991) and a comprehensive bibliography of early data relating to MBE of III-V compounds (from 1958-1983) has been published by Ploog and Graf (1984). Later improvements have allowed the deposition conditions to be improved by applying reactive gas in pulses (Goodman and Pessa 1986) which allows growth at lower temperatures and hence reduces diffusion to produce sharper boundaries. The choice of gas species and effusion cell temperature is very important as the structure of the molecule forming the film (e.g., dimer or trimer) is of major importance. The MBE process has been linked to other stages in the semiconductor production line as described, for example, by Migauchi and Hashimoto (1986) and by Melngailis (1987).
9.3.3 Sputtering When an ion strikes a material its energy is transferred by a series of collisions to the atoms of the first few layers of that material. Some of these atoms will move toward the surface and, if their energy is greater than the binding energy they will escape. The ratio of number of ejected atoms per incident ion is known as the sputtering coefficient and, for an incident ion of energy about 1 keV, it has a value for most materials ranging between 0.1 and 2 atoms per ion. A full review of the mechanisms and major trends in the sputtering phenomena was published by the author in 1968 (Carter and Colligon 1968, Chapter 7). The sputtering coefficient increases, after an initial quadratic behavior below 100 eV, almost linearly with ion energy as the ion range increases and more and more atoms participate in the energy exchange process. However, when the ion energy increases such that the ion range exceeds several atomic layers, the energy is too deep to be transferred back to the surface and the sputtering coefficient begins to decrease again.
230
J. S. Colligon
The fundamental understanding of the process has not changed significantly since the late 1960s but the theory of the process has been fully developed, the model by Sigmund (1969, 1971) proving to be very successful in predicting sputtering coefficients. Of particular importance to thin film deposition is the energy of the sputtered species which has a peaked distribution with a maximum at about 4-7 eV and a long tail down to over 50 eV. These energies are more than an order of magnitude greater than those of evaporated species and allow the depositing species more mobility on the surface to form denser, well-bonded films. One problem with the method is the change in composition of the surface of a sputtered compound material which arises from preferential sputtering of components. This results primarily as a result of the different energy transfer to different mass species. Fortunately, after a period of bombardment, the system is selfcompensating since the number of atoms of the species sputtering most readily is reduced and the rate of further sputtering of this component returns to the stoichiometric level. It is always important therefore to sputter a compound target for a period of time and to cover the substrate during this stabilizing period. Another problem from compound targets is that the spatial distributions of component ejected species can be different. Thus a substrate at a specific position in space facing the target may not collect component atoms in the correct stoichiometric proportions, even though the global total sputtering of species is stoichiometric. Clearly it is preferable, where possible, to move substrates so that they sweep through these spatial distribution variations. Apart from this problem, the process of sputtering is very versatile and control of the deposition rate is relatively simple. The first commercial systems were based on so-called planar diode sputtering and had the form illustrated in Fig. 9.3. Electrons are accelerated towards the anode and have ionizing collisions with gas atoms. These ions then bombard the cathode, causing sputtering and the liberation of further electrons which cause the discharge to be sustained. There are also many ion-gas collisions leading to charge exchange so that the target is bombarded by a mixture of ions and neutrals with energies below the total applied voltage. Typical operating conditions for dc metal sputtering are; cathode current density: 1 mA/cm 2, discharge voltage: 3 kV, argon pressure: 10 Pa, cathode-to-substrate distance: 40 mm, deposition rate: 40 nrn/min. For sputtering of insulating materials a radio-frequency (rf) supply has to be used, usually at the allowed 13.56 MHz frequency. These rf systems run at lower pressures, in the range 1-2 Pa. Although materials can be deposited using planar sputtering systems, the deposition rate is rather low, the ionization efficiency in the plasma is also low and the substrate becomes heated. These limitations have been overcome by the introduction of magnetron sputtering systems which are described in the next section. Deposition rates in these systems are as high as several microns per minute and can be controlled by varying the power density supplied to the target region, the target area, coating pressure and target-substrate distance. Although the substrates are near to the ionized plasma region the degree of ion bombardment of the growing film is low in these systems. The film quality is thus similar to that obtained for other sputtering methods but with a slight improvement in film density arising from the lower operating pressure; atoms sputtered from the target do not suffer many collisions in transit and therefore maintain, on average, a higher energy.
231
Physical Vapor Deposition
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9.3.4 Ion-assisted deposition
Evidence that additional ion bombardment of a surface being coated improved the adhesion and film properties was demonstrated in the early 1960s by Mattox (1963, 1965) using a system known as ion-plating, which was patented much earlier by Berghaus (1938). A typical ion-plating system is illustrated schematically in Fig. 9.4. The coating material is evaporated using an electron-beam and a discharge is struck in the upper region using a support gas such as argon or, for nitrides, nitrogen. Operating pressures are generally in the region 0.1-1 Pa which means that ions and atoms in the discharge make many collisions in passing to the negatively biased substrate. The substrate is thus bombarded by a variety of atomic, ionized and molecular species, including ions and atoms of the coating material with energies from thermal values up to the applied voltage of several keV. A further effect is the irradiation of the substrate and growing film by photons from the discharge. These energy-assisted deposition conditions have been shown to produce coatings with improved adhesion (Mattox 1965, Rigsbee et al. 1986, McLeod and Mah 1974, Enomoto and Matsubara 1975, Takagi et al. 1974, Teer 1983, Ahmed and Teer 1984), increased density (Mattox and Kominiak 1972), improved abrasion resistance (Lardon et al. 1978, Broszeit et al. 1985), and enhanced optical reflectivity (Henderson 1977, Varasi et al. 1985). However the film-growth parameters cannot be controlled independently; for example, a change in pressure also effects gas collisions and therefore energies of arriving particles.
232
J. S. Colligon
working gas
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Figure 9.4. Ion-platingsystem.
An altemative plasma-based ion-assisted deposition system is the arc-plating system. Here a discharge is run in the abnormal glow region of the current-voltage characteristic for a diode system. Extremely high (many amperes) currents flow in the discharge which forms an arc with the current entering the target at a small (few mm dia) spot. Target material is vaporized and, in the intense discharge, a high fraction of this material is then ionized. A biased substrate will therefore collect both atoms and ions of the film species (together with neutral species). The arc is struck with an igniter (used to produce fieldemission of electrons) and starts with additional gas in the chamber (such as argon) but, the arc once struck, the argon can be removed and the discharge maintained in the vapor. In operation the arc moves randomly over the target surface but this can be controlled by applying a movable magnetic field (steered arc system). Film quality is as good as for ion-plating with the exception that there is a tendency for large droplets of target material to erupt from the surface and fall on the substrate, thus spoiling film uniformity. This problem can be minimized by running at lower arc currents (but with associated lower deposition rates) or by moving a magnet below the cathode to "steer" the arc (Ramalingam et al. 1987, Erttirk and Heuvel 1989). Later versions have however introduced an arc deflecting column to entrain the arc around a 90 degree bend so that the substrate may be
233
Physical Vapor Deposition
coit'~
/
arc spot
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c afhode Figure 9.5. Filteredarc system.
mounted out of the line of sight from the target. This is known as the filtered arc system and is shown in Fig. 9.5 (Vyskocil and Musil 1990). In this situation only the ionized fraction of the ejected material will reach the substrate but this can be high (approaching 100% in some systems) so even this geometry provides high deposition rates. To provide controlled energy-assisted deposition for fundamental studies of the process, ion beam systems were developed in the 1980s. As mentioned earlier, the first of these was the apparatus devised by Weissmantel and his colleagues (Weissmantel 1976, Weissmantel et al. 1979) but many other groups produced similar systems (Harper and Cuomo 1982, Harper et al. 1982a, b, 1983, Sainty et al. 1984, Martin et al. 1983, 1984, Colligon 1995, Kheyrandish et al. 1994). Figure 9.6 shows some of the various possibilities for ion-assisted deposition and includes an e-beam system for evaporation plus partial ionization (PIB, Fig. 9.6(a)) (Bai 1990)), ion-beam deposition (IBD, Fig. 9.6(b)) (Colligon et al. 1976, Lawson et al. 1975, Armour et al. 1986, Sharples et al. 1987), dual ion-beam deposition (DIB, Fig. 9.6(c)) (Harper and Cuomo 1982, Harper et al. 1982a, 1982b, 1983, Sainty et al. 1984, Martin et al. 1983, 1984, Colligon 1995, Kheyrandish et al. 1994), single-beam plus evaporation (SBE, Fig. 9.6(d)) (Weissmantel 1976, Weissmantel et al. 1979), ion-beam-induced chemical deposition (IBC, Fig. 9.6(e)) (Harriott and Vasile 1988) and ionized cluster-beam deposition (ICB, see Fig. 9.1) (Gspann 1993, Takagi et al. 1977, Yamada and Takagi 1995, Yamada 1994). In many of the ion-beam systems the Kaufman source was used (Kaufman et al. 1982) which produces broad beam gaseous ion species with a relatively low energy spread over an energy range 50-2000 eV and with intensity of order 1-5 mA/cm 2. Most of the above techniques are well known and the principle of operation is clear from the diagrams. In IBD the ions are usually extracted from an ion-implanter source at energies of 10-20 keV and, after magnetic separation, are slowed down to about 10 eV using carefully designed ion-optical lenses to prevent beam divergence and to steer the ions
234
J. S. C o l l i g o n substrate at-V _
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Figure 9.9.
Variation in refractive index of CeO2 a a function of oxygen ion-assist energy. Films marked (x) are not stable on exposure to air. (From work of Netterfield et al. (1985)).
controlled studies of this type of new optical coating should lead to many new composite films with unique properties.
9.5.2 Corrosion protective coatings Another area where film density is important is that of corrosion protection. Ion-beam methods for producing such coatings have been reviewed by Wolf (1992) who notes that stress is an important factor in this type of application. Tensile stress can be reduced by increasing the ion/atom ratio in many cases, for example in Cr films using Ar ion assisted deposition (Roy et al. 1989a, b). Again the novelty of energy-assisted PVD is that new combinations of materials can be studied and their corrosion properties evaluated. Work reported by Hsieh et al. (1992), evaporating Nb and Cr with 250 eV Ar + ion-assistance at an ion/atom ratio of 0.68, has shown that Nb-50%Cr films are the most dense and column free but, because of its ability to form a Nb-rich oxide, Nb-30%Cr films had the best corrosion resistance. The present author has produced a range of Fe-Cr alloys for experiments which showed that the Cr concentration had to exceed 16% before complete passivation occurred (Fujimoto et al. 1993). The ability to "seed-in" additional elements has been exploited by Donahue et al. (1997) who showed that the addition of yttrium to magnetron-sputtered films of Ti-A1Cr-N drastically reduced the oxidation rate at elevated temperature.
246
J. S. Colligon
-
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9.5.3 Hard coatings Some of the main coatings used to produce durable hard layers include the nitrides of Ti, Zr, Hf and Cr, more complex nitrides which include Ti-A1-N, Ti-Zr-N, cubic BN, Ti-Nb-N, Ti-A1-O-N, diamond and diamond-like carbon, carbides of Ti, W and Ta and oxides. Hard coatings based on carbon, often referred to as diamond-like carbon (DLC) coatings, have been investigated using ion-assist methods since the early work of Aisenberg and Chabot (1971) who noted that diamond-like films were formed by directing a 40 eV beam of positive carbon ions at a substrate. Weissmantel (1981) later showed that similar films could be formed using ion-assisted techniques. More recently He et al. (1994) have shown that hard films with a high sp 3 content can be formed by sputtering graphite with Ar + ions onto silicon or AISI 52100 bearing steel. The growing film is bombarded with ions formed by introducing CH4 into a second source and running at energies from 0.2 to 25 keV. The highest sp 3 content occurred at the lowest (200 eV) energy where a Knoop hardness of 5100 kgf/mm 2 was registered using a 20 g load. A similar energy requirement was seen by Yugo (1994) in a microwave plasma running in a CH4/H2 mixture but here a two-step process was required, first a high level (40%) of CH4 was used to develop a carbon cluster and then, after 30 minutes, the CH4 level was reduced to only 0.5% allowing production of sp 3 bonds. This extensive field of DLC coatings has been reviewed by Ensinger (1996) who, in other work (Ensinger and Wolf 1989), has shown that these films provide good corrosion protection when applied to alloy steels. Silicon-carbide coatings have also attracted con-
247
Physical Vapor Deposition
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siderable interest (Learn and Haq 1970, Seaward et al. 1986, Yamada 1989, Berezhinskii et al. 1990, Yen and Murray 1990, He et al. 1995). Figure 9.11 shows data by He et al (1995) for dual beam ion-assisted growth using CH4 gas in the ion-assist source and Ar + ions to sputter silicon onto the substrate. It can be seen that the ion-assist energy must exceed 300 eV before the SiC compound forms. At this energy there was considerable re-sputtering of material and hence a low net rate of film growth. Cubic boron-nitride (c-BN) is an extremely useful coating for industrial applications owing to its high hardness, chemical and thermal stability at extremely high temperatures and low reactivity with ferrous alloys. It also has a large band-gap (>6 eV) and can be doped p-, or, n-type. Unfortunately, to form c-BN using conventional methods, pressures of over 11.5 GPa and temperatures in excess of 2000 K are required, which would destroy the properties of most substrates. However, ion-assist methods allow these films to be formed in vacuum at 460~ (see, for example Park et al. (1997)). In all cases a mixture of argon and nitrogen ions is required with an energy of about 300 eV and an ion/atom ratio of about 1.5. According to a model of the process by McKenzie (1993) and simulation experiments by Dworschak et al. (1995) the compressive stress, induced by the ions colliding with the film atoms, and the formation of interstitials is assumed to trigger the conversion of hexagonal-BN to c-BN. One of the most interesting aspects of current PVD methods is their ability to produce coatings with grain sizes smaller than those previously possible. This offers the possibility of producing nanoscale composite metastable alloys and pure materials which have unique properties including high hardness and high ductility. When the dimensions of components of the film become less than 100 nm the term/3 in the Hall-Petch equation
H = Ho + fld -1/2
248
J. S. Colligon
where H is the hardness, H0 a material constant and d the grain size, no longer applies since, according to Carsley et al. (1995) the material now has a significant fraction of its atoms in a grain boundary phase and so must be considered as a two-phase material. By considering typical grain boundary widths and grain sizes the fraction of each phase can be estimated and the overall flow stress of the composite material calculated, using bulk values of flow stress for the grains and amorphous material values for the boundary regions. The predicted values for hardness were compared by Carsley et al. with experimental data of Hughes et al. (1986) and E1 Sherik et al. (1992) for nickel and show good agreement. Similar agreement was also seen for iron and also copper. The maximum hardness is predicted to occur for grain sizes of about 4 nm for these materials. Vep~ek et al. (1995, 1996), Vep~ek and Reiprich (1996) and Christiansen et al. (1998) have predicted that mixed phase nitride materials should produce hard coatings and have demonstrated significant increases in hardness for crystalline TiN in amorphous Si3N4 formed by plasma-assisted CVD. Maximum hardness of 50-70 GPa was found for grain sizes of order 6-7 nm. Other combinations of nitrides are predicted to exhibit similar high values of hardness and Musil and Vlcek (1998) have noted that there are two types of mixed-nitride coatings: nanocrystalline metal nitride/hard matrix (such as c-TiN/a-Si3N4) and nanocrystalline metal nitride/soft matrix (such as c-TiN/Ni, Cu, Co, Ag, Au, rare earth elements). The metal nitride can be of several forms, including TiN, A1N, ZrN, NbN and WN. The former types of mixed nitride coatings are very hard but are brittle whereas the latter are of similar hardness but are ductile (Musil and Regent 1998, Musil and Zeman 1998, Nakayama et al. 1994). Musil et al. have shown that these mixed nanocomposite coatings can be formed using magnetron sputtering if the substrate is heated sufficiently to allow the second, amorphous phase, to migrate to the grain boundaries of the first phase (Musil and Vlcek 1998, Nakayama et al. 1994). It is this segregation that is believed to be responsible for limiting the growth in grain size. The value of substrate temperature required for formation of nanocomposite ZrN/Cu was found by Musil and Vlcek (1998), Musil and Regent (1998) to be of order 600~ and it is hoped that this can be reduced by improving the technique, for example, by incorporating a small quantity of a third element such as yttrium, which is known to reduce the grain size in metals (Liu et al. 1997). Another method for producing hard coatings is to use multilayer films A-BA-B . . . . , where the Young's modulus of one material is lower than that of the other. Seitz and Koehler (1956) predicted that the hardness would also increase for these "superlattice" structures and this prediction has been confirmed by Helmersson et al. (1987), for TiN/VN multilayers, by Kitagawa et al. (1998) for TiN/Ti multilayers and by Wu et al. (1997) for CNx/ZrN multilayers. 9.6. F U T U R E
TRENDS
The energy-assisted deposition processes are thus now producing a range of novel coatings which are well-bonded and can increase the lifetime of the coated product by significant factors. Coatings are formed mainly by magnetron sputtering, arc-ion-plating or other plasma-based techniques and it is expected that these methods will further improve to provide more efficient operation (for example, better utilisation of targets in magnetron systems) and better control of deposition parameters (for example addition of
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electromagnetic coils to control the degree of ionization in magnetron sputtering). However the main further development appears to be in the production of new metastable alloy coatings (Bates and Arnell 1994, Monaghan and Arnell 1991) and nanocomposite coatings (Hughes et al. 1986, E1-Sherik et al. 1992, Vep~ek et al. 1995, 1996, Vep~ek and Reiprich 1996, Christiansen 1998, Musil and Vlcek 1998, Musil and Regent 1998, Musil and Zeman 1998, Nakayama et al. 1994) According to Musil and Vlcek (1998), ions bombarding a surface during coating impart sufficient energy by momentum transfer to allow high temperature phases to form within the collision cascade. The surrounding atoms however provide fast cooling at rates of order 1014 K/s so that the high temperature phase is frozen into the solid. Thus the presence of (a) an additional small quantity of several additional elements in the film, (b) additional energy to provide the effective temperature and pressure conditions within the film to form the high temperature phase, and (c) high cooling rates, will allow many phases to be formed at relatively low temperatures. Ion-assisted coatings can provide all the above parameters and the immediate work needs to study what types of coating can be formed and how stable they are in working conditions. The possible control of composition, crystal structure and stress in the coatings will lead to some exciting new possibilities in the near future, perhaps even the formation of the theoretical hard beta-phase cubic C3N4, predicted by Liu and Cohen (1989), Cohen (1994) and Yu et al. (1994) and currently occupying the attention of many research groups (Sj6str6m et al. 1996, Vep~ek et al. 1995, Riviere et al. 1995, Boyd et al. 1995).
ACKNOWLEDGMENT
The author wishes to thank many colleagues for providing information for this review.
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Chapter 10 Chemical Vapor Deposition 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 10.3.5.2 Two- and Three-Dimensional Nucleation 10.3.5.3 Epitaxial Nucleation and Growth 10.3.5.4 Multi-component nucleation and chemical reactions 10.3.5.5 Nucleation Enhancement 10.3.5.6 Selective Vapor Deposition 10.3.5.7 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
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Chapter 10 Chemical Vapor Deposition E TEYSSANDIER and A. DOLLET
10.1. I N T R O D U C T I O N
In vapor deposition, a solid material is deposited from gaseous molecules onto a substrate to form a coating. The process is nowadays classified into two basic families according to the initial proposition of Blocher (1967): physical vapor deposition (PVD) when coatings are prepared from pure condensation processes, and chemical vapor deposition (CVD) when deposits are formed by chemical reactions only. These two routes were developed separately and many modifications have since been introduced, so that this distinction is no longer strictly valid. A great variety of processes related to CVD now exist. Some were developed to lower the deposition temperature by introducing (i) a glow discharge (plasma-activated CVD, PACVD, or plasma-enhanced CVD, PECVD) or (ii) less thermally stable organometallic precursors (organometallic CVD, OMCVD). Some were developed on the basis of a specific devicewa laser beam can be used, either for gas phase excitation or substrate heating in the laser-enhanced CVD process (LECVD). Certain denominations refer to specific types of coatings and related processes, such as epitaxially grown films for electronic or optoelectronic applications by the metalorganic vapor phase epitaxy process (MOVPE), or infiltration of ceramic materials in fibrous preforms by the chemical vapor infiltration (CVI) process. Distinction is usually made between two basic types of reactors: hot-wall and coldwall reactors (Fig. 10.1). In cold-wall reactors, the parts to be coated are heated by various means: electromagnetic induction or joule effect for metallic parts; tungsten-halogen lamps (RTPCVD, rapid thermal processing CVD) for non-conductive materials. In the hot-wall case, the whole reactor, usually tube-shaped, is heated by means of a classical furnace. Though the same phenomena take place in both cases, they present characteristic features: deposition can occur at hot walls thus leading to gas phase depletion, and the hot part in a cold-wall reactor is responsible for large thermal gradients and significant buoyancy-driven convection. These points will be further detailed. A lot of industrial coatings are currently produced by the CVD process, in the field of electronic, optical, optoelectronic or mechanical applications, and a considerable body of literature as well as scientific journals are devoted to CVD. The reader can also refer to the non-exhaustive review articles already published on the subject (Bryant 1977, Hess et al. 1985, Jasinski et al. 1987, Jensen 1987, Besmann et al. 1989, Jensen 1989b, Wahl 1989, Kleijn 1991, Hampden-Smith and Kodas 1995). The present review will report on 257
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the thermally-activated process that excludes any kind of physical activation either in the gas phase or on the surface. The basic mechanisms involved in the CVD process will be reviewed and illustrated.
10.1.1 Presentation Difficulties in obtaining a comprehensive view of the growth process arise from the number of interdependent parameters such as temperature, pressure, composition of the initial gas phase, flow rate, type of reacting chamber, etc. and the complexity of the various steps. The steps of the individual mechanisms involved, which are summarized in Fig. 10.2, can be divided into two highly interdependent parts: (a) phenomena occurring in the gas phase, such as bulk transport of the species in a hydrodynamic flow with diffusion mechanisms due to concentration gradients and possibly homogeneous reactions or gravity effects; in a cold wall reactor one must furthermore take into account thermal diffusion (Soret effect), and (b) the surface phenomena including adsorption, diffusion, reaction and desorption. In this case a distinction has to be made between the nucleation period when interactions with the substrate take place, and the steady-state growth of the coating from the gas phase.
259
Chemical Vapor Deposition
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The range of variation of the different parameters is as follows. In the thermally activated process, the temperature domain usually ranges from about 100 up to 2000~ The lowest deposition temperatures can be reached only by the use of organometallic precursors, since inorganic precursors, which are much more thermally stable, require at least 800~ to be decomposed. The total pressure in the reactor extends from atmospheric pressure down to a few tenths of a Torr (1 Tort = 133 Pa), so, according to the large characteristic dimensions of the reactors a continuum description is generally appropriate. Furthermore, since the gas velocities are low (Reynolds number < 100) and the Mach number is small (Ma < 1) the flows are laminar and the gas is assumed to be incompressible. The wide variety of relevant scientific domains is one of the difficulties of the CVD process and the various above-mentioned phenomena will be detailed in this review. The presentation order corresponds to the occurrence of the phenomena, from the introduction of the reactive mixture at the reactor entrance to the growth on the surface to be coated.
10.2. G A S - P H A S E T R A N S P O R T AND R E A C T I V I T Y
In CVD, the various chemical elements required to grow the film are transported by molecules in a forced flow from the entrance of the reactor towards the hot surface. The multicomponent, diffusive and reactive fluid flow can be an extremely complex system, and both flow phenomena and homogeneous chemical reactions affect the film growth. The fundamental equations describing the transport properties in continuous media are well known (Bird et al. 1960) and they will not be reviewed here. For a comprehensive review of transport phenomena in CVD see Holstein (1992), Jensen (1994). We will first focus on the fluid flow, ignoring its potential chemical reactivity which will be subsequently discussed.
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F. Teyssandier and A. Dollet
10.2.1 Non-reactive fluid flow 10.2.1.1 Geometrical effects The reactor has a critical influence on the uniformity of the fluid flow. The direction of the flow with respect to the gravitational field as well as the reactor geometry induce specific flow patterns such as recirculation cells. Under mass-transfer-controlled conditions the local flux of active species and hence the thickness uniformity and structural homogeneity of the coating is highly dependent on the flow pattern. Transport phenomena were described for a long time by simple analytical models based on boundary-layer similarity and film theory. The validity of the boundary layer hypothesis is questionable in many CVD reactors (Rosenberger 1987) and the absence of such a layer has been demonstrated in several cases by 3D numerical modeling since its thickness eventually exceeds the dimensions of the reactor for typical operating conditions. A stagnant layer theory was also encountered in the CVD field. This non-physical theory, which allowed decoupling convective and diffusive phenomena in several CVD analytical models, originates from flow visualization experiments. Eversteyn et al. (1970) interpreted a TiO2 particle free layer adjacent to the susceptor as evidence for a stagnant layer. Laser Doppler velocimetry experiments (Talbot et al. 1980), as well as modeling (Fotiadis and Jensen 1990) have revealed that thermophoretic forces were responsible for driving the seed particles away from the substrate. None of the above-mentioned approaches gives a satisfactory description of the physical processes involved in CVD, but they have provided simple models which have successfully predicted qualitative trends. The same film theories have also been developed for heat transfer (thermal boundary-layer), and for mass transfer (diffusion boundarylayer) (Hitchman 1980). All these theories are now encompassed in the approach which solves the coupled partial differential equations by numerical integration using finitedifference, finite-volume or finite-element methods. 1D to 3D simulation models are required depending on the type of reactor encountered. Both vertical and horizontal reactors are used. In the absence of gravitational effects, flow patterns are equivalent. This is not the case in CVD reactors on Earth, where natural convection resulting from essentially thermally induced density gradients is superposed on forced convection. The influence of natural convection on the process may be quite different according to the disposition of the reactor orientation. Vertical axisymmetric reactors such as impinging jet, stagnation point flow or rotating disk reactors have been widely investigated on the laboratory scale. The ideal stagnation point and rotating disk flows have uniform vertical gas velocities producing a boundary layer of constant thickness in which the gases move radially outwards. Due to the axisymmetric reactor geometry, 1D simulation provides a reliable description of the flow phenomena which agrees very well with flow visualization experiments (Wahl 1977, Fotiadis et al. 1987, Wang et al. 1995). Forced and natural convection in vertical reactors are coaxial phenomena. They have the same direction in the upward forced-flow case which stabilizes the axial density gradient. Recirculations may however be present in this case because of radial temperature gradients. Downward non-isothermal forced-flow results in counter-current influences that usually generate strong recirculation cells (Wahl
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1977). These cells are often eliminated by increasing the inlet flow-rate thus improving the thickness uniformity. Flow in horizontal reactors is much more complex, and any horizontal density gradient in a gravitational field results in buoyancy-driven free convection without threshold (Rosenberger 1987, Gokoglu et al. 1989). Both longitudinal and transverse rolls are then generated and 3D modeling is in this case required (Ouazzani and Rosenberger 1990). Under mass-transfer-controlled conditions the thickness uniformity is highly dependent on the flow pattern. The deposition profiles calculated in Fig. 10.3 (Jensen 1989a, Moffat and Jensen 1986) for two different flow patterns show that growth rate is enhanced or reduced respectively in the domains facing downward or upward secondary flow directions. The same phenomena can take place at the edge of substrates, frequently leading to their rounding off. Mass-transfer inside a reactor depends on both the reactor enclosure and the operating conditions. The latter can be modified from one experiment to another, but the reactor geometry is selected once for all. The design of reactors can be optimized by modeling a priori the flow and heat transfer. These calculations are invaluable in obtaining a better understanding of their behavior for various operating conditions. Several basic options can thus be checked: the shape, distance between inlet and susceptor, inlet tube diameter, introduction of baffles. The influence of these reactor features has been extensively illustrated in the case of a cold-wall vertical reactor (Fotiadis and Jensen 1990).
10.2.1.2 Influence of operating conditions Once the reactor has been selected, the experimenter has to determine the most appropriate operating factors. Several parameters can be chosen and these include temperature, total pressure, nature and proportion of precursors, and total flow rate. The nature and shape of the parts to be coated are the imposed parameters. Our understanding of the basic phenomena involved in the CVD process provides us guidance for selecting or optimizing the operating parameters. Analytical models were first used for that purpose. The influence on the growth rate of the basic parameters such as total pressure or temperature can be easily schematized and therefore memorized thanks to fairly simple models, such as the one proposed by Van den Brekel (1977) which successfully predicts qualitative trends. To quantify the relative influence of gas phase diffusion resistance transfer against the surface reactions kinetic limitation, he introduced the so-called CVD number, which is in fact a simplification of the Damk6hler (surface) number when only first-order surface reactions are considered (Hess et al. 1985). The general behavior of a CVD reactor can also be deduced from dimensionless groups that arise from the scaling of the governing transport equations (Jensen 1989). These groups are also helpful in determining the influence of a specific parameter, such as temperature, on the physical properties of gases and thus the transport properties in CVD reactors (Rosenberger 1987). 10.2.2 Reactive flows In many situations, the residence time and gas temperature are high enough to allow reactions to occur in the gas phase prior to reaching the surface. The new species so
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produced are transported towards the substrate at variable rates according to their physical properties. Therefore, a large variety of species having different reactivities can be present in the vicinity of the deposition surface. The main CVD processing domains are traditionally reported in terms of growth rates and their dependence on temperature. Surface kinetics or gas phase mass-transfer control operating domains are thus deduced from the overall activation energy. To determine the apparent activation energy, it is usual practice to plot the logarithm of the growth rate against 1/T. Three temperature domains can thus be observed: (a) in the low-temperature domain, the deposition rate increases rapidly with temperature. The steep slope which is indicative of a high overall activation energy is associated with a growth rate governed by the surface phenomena. For very short residence times and/or very low temperature of the gas mixture, almost no intermediate species are formed, and the flow is assumed to be chemically frozen. (b) at intermediate temperatures, surface kinetics become fast and the ability of the process to supply reactive species to the surface becomes the growth-limiting factor. In this range of temperature, convection in the bulk and diffusion in the vicinity of the surface become the rate-limiting transport phenomena. Their weak dependence on temperature is responsible for the low overall activation energy. Higher gas phase temperatures allow homogeneous chemical reactions to develop under kinetic control. Near-equilibrium conditions are reached in the gas phase for long residence times and/or high temperatures. (c) at high temperatures, a decrease in the growth rate is often encountered. Several phenomena may account for such a behavior: the nucleation and growth of particles in the gas phase that deplete the amount of reactive species on the surface, modifications of the fluid flow dynamics or desorption of species that exceed the number of molecules impacting on the surface. In the intermediate and high temperature domains (cases (b) and (c)) gas phase phenomena and especially mass transport phenomena are the factors that control growth rate uniformity. The transition from the surface kinetics rate limiting domain to the mass transport limiting domain is of prime importance in the process as it leads to very specific features of the growth mechanisms. Gas-phase transport-limiting process is often responsible for surface roughening, whereas surface kinetics limitation tends to smooth out the surface (Van den Brekel 1978, Palmer and Gordon 1988, Palmer and Gordon 1989) (see also crystal growth section in part 3).
10.2.2.1 Thermodynamic approach Because the detailed reaction mechanisms and corresponding kinetic data are rarely available, chemical equilibrium has been the first traditional approach to CVD process modeling. Thermodynamic calculation has been applied to a great variety of chemical systems and for various purposes which will be detailed in the "Solid-gas thermodynamic equilibrium approach" section in part 3. Application of thermodynamic calculations to gas phase in the absence of solid state materials leads to pseudo-thermodynamic equilibrium involving a supersaturated gas phase. This allows determining bulk compositions and identifying key species, as well as conditions under which they might form (Allendorf 1993). Though this equilibrium situation, which is representative of a gas phase composition that would be reached at infinite residence time by a reactive mixture flowing at constant temperature (as long as homogeneous nucleation does not take place), is never
264
F. Teyssandier a n d A. D o l l e t
encountered in CVD reactors, near- equilibrium states are often approached in hot-wall reactors. Accurate modeling and, ultimately, control of the CVD process requires knowledge of both the identity of the species formed during gas phase transport and their rates of formation. When these are known, the flux of reactive species toward the surface can be predicted. For that purpose, gas phase reaction mechanisms have to be identified and their corresponding rate constants must be measured or estimated.
10.2.2.2 Gas-phase mechanisms and kinetics In the dilute medium encountered in the CVD process, it is assumed that there are only two types of kinetic processes: unimolecular bond dissociation reactions and bimolecular reactions. For further details on experimental and theoretical development of gas-phase kinetics the reader can refer to classic texts (Benson 1976, Forst 1973, Gilbert and Smith 1990, Robinson and Holbrook 1996). Only a brief review of the basic principles will be given here. In a unimolecular reaction, an isolated molecule undergoes a chemical change. Reaching an internal energy sufficient for the reaction to occur requires excitation, which most frequently is obtained from collision with a bath gas at thermal equilibrium. Unimolecular reaction rates can be obtained experimentally from static (Neudorfl et al. 1980) and flow pyrolysis (Ring et al. 1970), shock tubes (Mick et al. 1996) and laser powered (Jasinski and Estes 1985) homogeneous pyrolysis techniques. They may also be estimated from theoretical models. The energy-dependent rate constant k ( E ) can be computed from the RRKM model (Rice-Ramsperger-Kassel-Marcus) which is the application of the transition state theory to a microcanonical ensemble of excited reactant molecules. This model gives an exact upper bound to the rate coefficient. A simpler model (QRRK) (Westmoreland et al. 1986) in which a molecule is treated as if all its vibrations can be represented by a simple average frequency (usually the geometric mean) can be used instead of the RRKM model. Three categories of bimolecular reactions (A + B --+ products) may be distinguished: displacement, association and metathesis reactions (in which an atom is transferred from A to B). The bimolecular adaptation of the unimolecular QRRK model (Dean 1985, Westmoreland et al. 1986) can be used to predict the pressure and temperature dependence of addition or recombination reactions that initially form excited intermediates. Once the various gas phase reaction mechanisms are identified and their rate measured or estimated, the kinetic data can be introduced in full-scale models or simpler (plug-flow or well-stirred) time-dependent models such as the Sandia CHEMKIN/SENKIN package (Kee et al. 1989, Lutz et al. 1995). These tools are useful in interpreting the evolution of gas mixtures (Teyssandier and Allendorf 1998).
10.2.2.3 Simulations of the gas-phase composition Even the simplest chemical systems may involve numerous gaseous species. The model proposed by Coltrin et al. (1989) for silicon deposition from silane, which includes only two elements (Si and H), is composed of 27 chemical reactions involving 17 gas phase chemical species. Adding carbon to allow SiC deposition, increases the gas phase chemical model up to 83 reactions involving 41 species (Allendorf and Kee 1991). Some of these species are not important in the deposition process and such systems can be simpli-
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265
fled for specific total pressure and temperature conditions (Wang et al. 1994). However the rejection of some species or chemical reactions according to concentration criteria can be misleading. As a fact, the selected subsystem must constitute a closed system with respect to the chemical reactions considered, i.e., the important reaction pathways for the selected species must not be ignored. Furthermore, due to the large difference of surface reactivity between closed-shell molecules and radicals, the latter species, even when present at low concentration, may have a significant influence in the surface process. For instance, though Sill2 forms only at high temperatures and is thus usually present at the interface with the surface substrate in far lower concentrations than the precursors introduced at the reactor entrance (silane or disilane), its unity reactive sticking coefficient (RSC) at the surface is at least three orders of magnitude larger at 1000 K than those of the saturated silanes (Sill4, Si2H6, Si3Hs, etc.). The contribution of silylene to silicon growth rate (at least at high temperatures) can be for that reason significant. Full-scale modeling including gas-phase reactions allows determining the contribution of the various gaseous species to the deposition rate. The partial deposition rates of Sicarrier species at the surface facing a circular impinging jet (Fig. 10.4) reveals that (i) the deposition rate increases rapidly towards the edge of the impingement surface as a result of the increasing amount of highly reactive intermediate species formed, (ii) the contribution of the initial precursor Si2H6 is constant on the surface that reflects its supersaturation which is a consequence of its low surface reactivity, and (iii) Si3H8 is responsible for the step variation of the deposit thickness observed at the edge of the surface.
10.2.2.4 Nucleation and growth of solid particles from the gas phase The non-controlled introduction of solid particles in a CVD reactor is detrimental to the properties of the films. These particles are usually introduced into the reactor as impurities or can be formed by agglomeration of smaller particles nucleated in the gas phase. They act as extrinsic seeds for heterogeneous nucleation, deplete the species necessary for the heterogeneous reactions and create undesirable microstructures including porosity that lower the mechanical properties of the deposits. The elementary processes involved in the formation and growth of particles from gaseous precursors include gas phase chemical reactions, homogeneous nucleation, heterogeneous condensation, coagulation and coalescence or fusion (Fig. 10.5). Both high temperatures and long residence times are favorable to homogeneous nucleation of such particles. This situation is especially encountered at the center of recirculation cells (infinite residence time) in hot-wall reactors. In contrast, the risk of solid particle formation is considerably reduced under low total pressure and/or in diluted reactive mixture. Because of the large temperature gradients found in cold-wall CVD reactors, the thermophoresis forces may cause the solid particles to deviate significantly from the fluid particle trajectories. This phenomenon is the preponderant mass-transfer mechanism used in the modified chemical vapor deposition process (MCVD) to manufacture optical waveguide preforms (Kim and Pratsinis 1988, Lin et al. 1990). Since the thermophoretic coefficient increases with decreasing pressure, its influence is expected to be enhanced with decreasing pressure (Fotiadis and Jensen 1990). Accordingly, low pressures, which reduce buoyancy-driven flows, favor rapid transport of particles away from the deposition zone and reduce particle contamination. The carrier gas has also a significant influence on
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10.3. S O L I D P H A S E F O R M A T I O N
Once introduced into the CVD reactor, the reactants may experience some chemical transformations during their transport towards the substrate. When reactants and reaction products reach the solid-gas interface, the gaseous chemical species undergo new complex physical and chemical interactions that may lead to film growth provided that some thermodynamic and kinetic criteria are satisfied. We will first discuss the energetics of the solid-gas transformation, and subsequently, the kinetics of the elementary heterogeneous processes. The nucleation and crystal growth phenomena will be finally addressed.
10.3.1 Solid-gas thermodynamic equilibrium approach The final equilibrium state of a thermally-activated CVD system, in which various chemical reactions are involved in the formation of a solid material from a gas phase, is imposed by thermodynamics whether or not this state is reached for kinetic reasons. In other words, when quantitative predictions are difficult to obtain because of the lack of mechanistic and kinetic information, thermodynamic calculations may provide reliable qualitative trends. Several review articles have already been devoted to this subject (Bernard and Madar 1990, Bernard and Madar 1991, Spear and Dirkx 1990). When all gaseous and solid species are taken into account, the "closed box" (nonflowing closed system) thermodynamic calculation at equilibrium provides valuable qualitative information on the nature of the different phases formed, on their amounts,
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F. Teyssandier and A. Dollet
as well as on the influence of the various parameters. Two methods have been developed for the calculation of chemical equilibria in complex chemical systems: one based on mass action relationships, and the other based on minimization of the Gibbs free energy of the system. The latter is the method currently used in commercial computational programs. In the calculations it is implicitly assumed that (i) thermodynamic equilibrium is reached (no kinetic control), (ii) transport phenomena can be neglected, (iii) all the species that may form at equilibrium are taken into account in the calculation, and (iv) the Gibbs free energy of the species is known with good precision. When the residence time of the gas mixture in the hot zone of the reactor is short and/or the temperature of the gas mixture is low, kinetic limitations cannot be ignored and a detailed mechanistic investigation of CVD reactions is required.
10.3.2 The structure of CVD layers One of the major advantages of the CVD technique is that each of the operating parameters can be selected within a wide domain of variation, leading to a great variety of film structures. Strictly speaking, the possible structures of CVD layers are so numerous that it is almost impossible to discuss all of them in detail; nevertheless, they can be classified into several categories and sub-categories, so that their essential features can be captured. The first level of classification consists in dividing materials into amorphous materials and crystalline materials. Amorphous materials (AM) consist of metastable solid phases, generally highly non-organized, exhibiting only short-range atomic order. Their composition may differ significantly from the stoichiometric composition of the corresponding thermodynamically stable phases. Moderate temperatures are required in order to obtain this kind of material using the CVD technique, and PECVD, which allows near room temperature deposition, is often preferred to prepare AM on substrates which are very temperaturesensitive. One of the main features of deposited CVD AM is that, despite their disordered structure at the atomic scale, they may exhibit some structural organization, exemplified by columnar microstructures which have been observed in numerous cases (Thornton 1977). Crystalline materials (CM) can be divided into epitaxially grown monocrystalline materials and polycrystalline materials. The term epitaxy is used to indicate that the growing film structure is controlled by the (mono)crystalline substrate structure. When the mismatch between the respective lattice parameters of the substrate and the desired film is not too large, and also when particular growing conditions are satisfied (see further), the substrate acts as a seed crystal, imposing its lattice parameters to the growing material. When the film and the substrate are of the same nature and the film grows exactly in the same crystalline structure as the substrate, the growth process is called homoepitaxy. When the film and the substrate are made of different materials (lattice mismatch is highly probable in this case), the film may still grow with a crystalline structure imposed by the substrate lattice, but mechanical strains appear inside the film, the intensity of which depends mainly on the degree of mismatch. If the initial mismatch is non negligible, structural defects such as dislocations will nec-
Chemical Vapor Deposition
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essarily be created in order to relax the strains. This second epitaxial growth process is referred to as heteroepitaxy. A large variety of materials such as Si, GaAs, GaN, SiC, etc. are epitaxially grown using the CVD technique, mainly for applications in the field of electronics. Polycrystalline materials (PCM) are easily obtained using standard CVD conditions. They have generally highly anisotropic microstructures (depending on the operating conditions but also on the film thickness), generating anisotropic physical or chemical properties. The various surface morphologies of CVD PCM have been classified in many categories such as: plate-like, lenticular, star-shaped, pyramidal, spherulitic, cauliflower, dendritic, etc. Processing/morphology maps can be drawn to provide a general idea of the influence of operating parameters on surface morphology. As an example, Figure 10.6a shows such a map for CVD TiCN (Cheng et al. 1987). Another approach used to schematize the microstructure evolution is to draw a structure zone model (Movchan and Demchishin 1969, Thornton 1977). Three zones have been distinguished by Movchan and Demchishin (1969) for the coating microstructures (Fig. 10.6b) which appear to be a function of T~ Tm, where T is the substrate temperature and Tm the PCM melting point. The first coating structure is observed at T/Tm < 0.3. At such moderate temperatures, surface diffusion of the adsorbates is low: tapered crystals with domed tops are observed, separated by voided boundaries. The second structure corresponds to 0.3 < T~ Tm < 0.5. Surface diffusion is favored in this case, columnar faceted crystals with dense boundaries are observed (Thornton 1977), the size of which tends to increase with coating thickness. The last structure (0.5 < T~ Tm < 1) is controlled by bulk diffusion: the deposit is composed of equiaxed grains and exhibits a bright surface. In order to better account for the substrate roughness, a fourth zone (not presented in Fig. 10.6b), zone T, has been added by Thornton (1977) between zone 1 and zone 2, which consists of fibrous grains with dense boundaries. The above zone models were established for PVD, however, the general trends remain the same for CVD. The values of T~ Tm separating the various zones may vary significantly from one material to another, and should only be considered as approximate values in the best cases. For any zone, it is observed that the grain size increases with T~ Tin.
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F. Teyssandier and A. Dollet
It should be kept in mind that although they seem to be very attractive, the above approaches remain empirical since they are completely specific to the reactor, and also to the material for processing/morphology maps. 10.3.3 The driving f o r c e f o r crystal growth f r o m the v a p o r p h a s e
The gas-solid thermodynamic equilibrium assumption, discussed above, is obviously never satisfied in practice in the CVD process since it implies that there is no net mass exchange between the two phases, that is, no condensation or crystallization of the vapor, nor solid evaporation. The equilibrium condition is: # g -- ~s
(10.1)
w h e r e # g and/Zs are the chemical potentials of the gas phase and solid phase, respectively.
Nevertheless, when #g ~: /Zs, the system tends to evolve towards its minimum energy, i.e. thermodynamic equilibrium. The condition #g > /Zs is then required in order for deposition to be observed. The chemical potential excess in the vapor phase, A/z, is the chemical driving force for crystal growth from the vapor. For quantification purposes, the relative supersaturation a defined as: cr -- e Alz/RT
(10.2)
is generally used instead of A#. The absolute supersaturation: aa = a - 1, is also used in some cases. From the values of a or aa, the trends in the evolution of a solid-gas systems can be predicted: 9 0 < a < 1 (or aa < 0) dissolution of solid into gas phase 9 cr = 1 (or aa = 0) solid-gas equilibrium 9a > 1 (or O"a > 0) solid deposition from vapor phase a can be related to the partial pressure of the various species present in the gas phase by using the classical thermodynamic relations for the chemical potentials. However, rigorous analytical expressions of a are not practically easy to obtain for real CVD systems, as many chemical species and sometimes several complex solid phases have to be taken into account. Authors often illustrate the evolution of the microstructure of CVD films by the correlation diagram proposed by Blocher (1974) in the PVD case. However, the microstructure of CVD films depends on supersaturation in a more complex manner, because of the chemical reactions. Figure 10.7 shows the experimentally observed evolution of the microstructure of CVD films as a function of temperature (which influences the surface mobility of adsorbed species) and flux of active species towards the surface. The microstructural order of the deposit is mainly the result of a competition between these two factors. Several attempts have been made to evaluate the supersaturation in CVD processes, in order to correlate supersaturation with film microstructure. The validity of the calculations carried out for that purpose is often questionable, either because equilibrium
271
Chemical Vapor Deposition
Effect of increased surfacetemperature
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Figure 10.7. Effect of temperature and flux of reactive species on CVD film microstructure.
approximations have been used (Langlais et al. 1982), or because simple but meaningless definitions of supersaturation have been chosen (Goto et al. 1992). Vandenbulcke and Vuillard (1977) have proposed a more rigorous approach by evaluating supersaturation from the concentration of species at the gas-solid interface, taking into account surface reactions. They have plotted the observed microstructure of boron deposits as a function of the supersaturation and flow rates. It has been found that formation of amorphous deposits is favored at low temperatures and high supersaturation. Later, Vandenbulcke et al. (1991) have also proposed a schematic plot of the various microstructures of carbon (graphite and diamond) prepared by CVD as a function of the calculated supersaturation. However, in this study like in most of the others, only very simplified chemical schemes have been considered.
10.3.4 Surface mechanisms and kinetics Surface kinetics are completely ignored in the thermodynamic approach which predicts only the final state of a solid film assuming its thermodynamic equilibrium with the saturated vapor. However, the formation of structural defects in solid films and metastable solid phases (which is evidence of the non-equilibrium state of a CVD system) can only be explained by kinetic limitations induced by the heterogeneous processes. In this section, a very brief survey of elementary surface processes will be first given, then the methods for obtaining the related kinetic parameters will be presented and illustrated on several examples dedicated tO the CVD process. 10.3.4.1 Rate of elementary surface processes The most important heterogeneous phenomena occurring at the film surface during deposition are (i) adsorption (physisorption or chemisorption), (ii) surface diffusion, (iii) reaction on the surface, and (iv) desorption (Fig. 10.2). The amount of species adsorbed at steady-state depends on the nature of both the surface and the adsorbates, on the pressure and on the temperature. Surface processes are thermally activated but it is worth to note that the energy barrier for a given heterogeneous reaction is almost always significantly lower than that of its homogeneous counterpart (Laidler 1987). As will be seen in Section 10.3.6.1, a solid surface is not smooth at the atomic level: various kinds of surface site exist at which adsorbed species may be attached more or less strongly (Laidler 1987, Tsong 1990). At low temperature, most of the adsorbates are not able to overcome the energy barriers of the various surface processes. They cannot move over distances large
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enough to find the energetically most favorable surface sites for incorporation, a condition which is required for the film to grow as a crystalline material. Nevertheless, some local minima of potential energy exist at the immediate vicinity of the adsorption sites, in which they can be trapped; in this case, only an amorphous solid phase can be grown. Among the most important factors determining the kinetics of the above-mentioned heterogeneous processes are the binding energies of surface atoms. The rate constants of the elementary surface processes are generally assumed to obey an Arrhenius equation. However, the measurement or the evaluation of Arrhenius parameters of surface processes is not a simple task since they depend on the surface coverage by adsorbates (Zhdanov 1991), and also, when occurring, on adsorbate-induced surface reconstruction (Zhdanov 1991, Somorjai and Van Hove 1995). 10.3.4.2 Estimation o f the kinetics of surface processes in CVD Kinetic parameters or binding energies of surface atoms can be obtained experimentally by means of various techniques such as temperature-programmed desorption (TPD), isothermal desorption (ID), field-ion microscopy (FIM), scanning tunneling microscopy (STM), laser-induced thermal desorption (LITD) (Zhdanov 1991, Tsong 1990). Among the most efficient techniques, FIM and STM are used to obtain direct information on surface diffusion, whereas kinetic information on desorption and surface reactions are generally obtained from TPD and ID analyses. Molecular beam reactive scattering (MBRS) is also an efficient method for providing information on the surface kinetics in CVD according to Gates (1996). It should be pointed out that such experiments are usually expensive and difficult to perform. Moreover, since chemical vapor deposition is generally carried out at relatively high temperatures, there are so many possible reaction pathways that a complete experimental determination of the kinetics of all possible surface processes in CVD appears rather unrealistic. Because of the complexity of the heterogeneous processes, the kinetic equations which are most often included in CVD reactive transport models are empirical overall kinetic equations. Some reaction mechanisms are assumed a priori and the related kinetic parameters are subsequently adjusted until an acceptable agreement between theoretical and experimental results (mainly the deposition rate profiles) is obtained (Roenigk and Jensen 1987). The simplest procedure for describing the surface kinetics in CVD is probably the introduction of a reactive sticking coefficient (s), i.e., the probability of surface reaction. During thermal accommodation (Herlin et al. 1991), the increase of the internal energy may be either transformed into kinetic energy favoring desorption or into rotovibrational energy inducing breaking of bonds. Sticking coefficients can be obtained from specific measurements such as molecular beam scattering at very low pressure (Buss et al. 1988, Fisher et al. 1992). They can also be deduced from indirect experimental ways, for instance, by examining the conformality of films deposited inside micrometric trenches. A model simulating deposition in such geometries, i.e., Monte-Carlo (Dew et al. 1992) or analytical (Cale and Raupp 1990, IslamRaja et al. 1991) models, provides the s value (Fig. 10.8). It is possible to find the value of s for a particular chemical species if its incident flux at the surface can be monitored independently (Matsuda et al. 1990, Fisher et al. 1992). Langmuir kinetic equations, taking into account the surface coverage by adsorbates, are also often chosen a priori to describe the surface kinetics of the CVD process (Roenigk and Jensen 1987).
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Figure 10.8. Simulationof the deposition profile in an infinite trench for various values of the sticking coefficient (1, 0.1, 0.01) (IslamRajaet al. 1991).
Several theoretical methods, that allow calculating the kinetics of the various surface processes (Molecular Dynamics (MD) and Monte-Carlo (MC) simulations, transitionstate theories (TST), electronic structure calculations) can be used to analyze the kinetics of the surface processes involved in CVD (Koseki and Ishitani 1992, de Sainte Claire et al. 1996); however, a severe drawback of these sophisticated methods lies in their complexity. A paper by Wang and Pollard (1995) has been recently dedicated to the presentation of a simpler theoretical approach for modeling surface reaction kinetics in CVD, which combines simple statistical mechanics, TST, and bond dissociation energies. The method requires four steps: (i) identification of the nature of the various sites and the corresponding bonding configurations of the adsorbates, (ii) evaluation of the thermodynamic properties of the reacting species, (iii) description of the nature of the elementary processes, and (iv) assessment of the structure and properties of the transition states. The method has been applied to tungsten deposition from WF6 and Sill4, and Fig. 10.9 shows some typical adsorbates on the (100)W surface. Because of its relative simplicity, this approach can only provide rough estimations of the kinetic parameters of the various surface reactions but calculations are not time consuming and, as a consequence, a very large number of elementary steps can be investigated. In conclusion of this section, it must be recalled that surface chemistry plays a key role in determining the structure of films grown by CVD. This subject has been further discussed, for instance, in a recent review article by Gates (1996). Unfortunately, it should be also emphasized that in spite of the importance of the elementary surface processes in CVD, reliable kinetic data remain generally difficult to obtain.
10.3.5 Heterogeneous nucleation Since it is the initial stage of crystal growth from a vapor phase, heterogeneous nucleation is one of the most dominating factors in the microstructure formation. In this section, a brief insight into the fundamentals of heterogeneous nucleation will be first given and some of its implications in the field of CVD will be illustrated. Nucleation consists in the formation of crystalline embryos (clusters or nuclei) at the substrate surface, due to localized density fluctuations in the supersaturated vapor. Some gaseous species are first adsorbed on the substrate surface, which either move along this surface or are temporarily trapped in potential wells. During their movement on the sur-
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ooeO H F Si
W
Figure 10.9. Schematic diagram of the (100)W surface with typical adsorbates on the surface. A tungsten atom is present at the apex of each pyramid (Wang and Pollard 1995).
face the adatoms may interact with each other to build-up clusters, but the adsorbates which have not been incorporated in nuclei of critical size before the end of their life time re-evaporate. If a sufficient number of stable nuclei can be formed, the moving adatoms increasingly attach to the stable nuclei, which become bigger, and a redistribution of their sizes is observed. When they become big enough, they come into contact with each other, leading to their size increase at the expense of their number. 10.3.5.1 Nucleation theories Two distinct theories are generally used for analyzing the nucleation process. The first one, mainly based on thermodynamic considerations, is referred to as classical nucleation theory. Fundamental data such as nucleation barrier height, critical size of nuclei, nucleation frequency, equilibrium shape of nuclei, can be obtained using this theory (see Chakraverty 1973, Bennema 1973, Mutaftschiev 1993, among other articles). When the supersaturation is high, the classical nucleation theory becomes difficult to use and a second theory, the atomistic nucleation theory, must be applied. Nucleation is then regarded as a kinetic polymerization process: instead of a critical size and a critical energy, a critical number of atoms and cluster dissociation energies are defined (Rhodin and Walton 1963). Convenient analytical expressions of the nucleation rate, critical size and cluster density are given by several authors, such as Lewis and Campbell (1967), in the case of atomistic theory. Most often in CVD, the critical nucleus size is of the order of only a few atoms and the atomistic nucleation theory is used instead of the classical theory (Liu and Dandy 1996). An example of the use of the non-classical (atomistic) theory in the field of CVD can be found in a paper by Liu and Dandy (1996) who have studied the influence of substrate materials, surface diffusion and adsorption state on the initial stage of diamond nucleation on carbide-forming substrates. The theoretical results obtained have
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been found to agree qualitatively well with previous experimental observations, in spite of the hypotheses formulated (not recalled here). The major findings of this study are that: (i) surface diffusion of adatoms is the most determining factor influencing the nucleation kinetics of diamond (it depends on the substrate and temperature); (ii) when increasing the substrate temperature, chemisorption is favored at the expense of physisorption and surface diffusivity is also increased. As a result, the mean residence time of adsorbed species on the surface is increased and nucleation is enhanced. 10.3.5.2 Two- and three-dimensional nucleation When an embryo does not perfectly wet the substrate, the value of the free energy of interface formation is high and a significant supersaturation is required to overcome the nucleation barrier. Three-dimensional stable clusters are formed in this case, the orientation of which with respect to the substrate does not depend on supersaturation. On the other hand, when the substrate is perfectly wetted, two-dimensional nucleation may take place, provided that the gas phase is supersaturated with respect to the 2D solid phase. In this case, however, the orientation of the 2D phase may depend on the ledge free energy and on the supersaturation (Mutaftschiev 1993). A temperature increase favors the transition from 3D nucleation to 2D nucleation, since the enhancement of their mobility allows the adatoms to surmount local energy barriers, such as barriers for jumping down steps (Levi and Kotrla 1997).
10.3.5.3 Epitaxial nucleation and growth At equilibrium, an adsorbate layer undergoes two kinds of opposing interactions: the firstones are the interactions between the adatoms, and the second-ones are the interactions between the adatoms and the substrate atoms. If the substrate-adsorbate interactions are strong as compared to the adsorbateadsorbate interactions, and if the distance between neighbor-substrate atoms is comparable to the distance between atoms in the normal bulk structure of the growing film, the nucleation of two-dimensional clusters on the substrate and layer-by-layer growth may be favored. In this case, the substrate imposes its structure on the two-dimensional growing material, even if the distance between the atoms of the growing crystal planes (hkl) parallel to the substrate surface is slightly different from the distance between the atoms of the corresponding isolated three-dimensional bulk material. Because the influence of the substrate on the film structure decreases with increasing number of new monolayers, the chemical potential of the 2D phase, which is minimum for the first monolayer, gradually increases for the upper monolayers. For layer-by-layer growth to be observed, the gas phase may be undersaturated with respect to the infinite three-dimensional solid phase, but must always remain supersaturated with respect to the growing two-dimensional phase. This process, known as layer-by-layer or Frank-Van der Merwe mechanism (Frank and Van der Merwe 1949a,b), can only be observed below the so-called roughening temperature (RT). If the substrate-film lattice mismatch is more significant, the substrate may still impose its structure on the first monolayer (the gas phase is then supersaturated with respect to the 2D film structure), but the internal strains in this layer become large. However, as mentioned above, the influence of the substrate on the upper ad-layers decreases and, up
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Frank Van der Merwe (FVM)
Stranski-Krastanov ( S K )
Volmer-Weber (VW)
Figure 10.10. Nucleation/growth modes in epitaxy.
to a critical film thickness, the substrate can no longer impose an abnormal structure to the film and layer-by-layer growth mode stops. If the system is supersaturated with respect to the 3D solid phase, three-dimensional nucleation follows and the film continues to grow with its normal bulk structure. This particular process is known as the Stranski-Krastanov mechanism (Stranski and Krastanov 1938). If the substrate-film lattice mismatch is large and the substrate-film interactions are not strong compared to the adsorbate-adsorbate interactions, three-dimensional nucleation occurs on the bare substrate surface, at supersaturation with respect to the 3D film structure. This simple three-dimensional nucleation process is called the Volmer-Weber mechanism (Volmer and Weber 1926). The above mechanisms are schematically summarized in Fig. 10.10.
10.3.5.4 Multi-component nucleation and chemical reactions Up to now, the discussion has been restricted to the case of single-component nuclei formed from a single-component vapor without taking into account the kinetics of the heterogeneous chemical reactions, which are of primary importance in CVD. Although most of the fundamental features of the nucleation process have already been exposed, formation of multi-component nuclei with chemical reactions presents some specific aspects which deserve particular discussion. A general analysis of the problem has been presented by Kukushkin and Osipov (1995) who have classified the elementary processes which are likely to occur when heterogeneous nucleation from a multi-component vapor takes place. They have taken into account the possible formation of solid solutions as well as eutectics and chemical compounds and have proposed the following processes: (i) direct formation of multi-component nuclei, (ii) formation of nuclei after a chemical reaction within an adsorbed layer, (iii) chemical reaction within the nuclei, (iv) eutectic separation in a solid phase, (v) evaporation of nuclei, and (vi) further growth of nuclei. A particular sequence of some of these processes can be formulated to describe nucleation from a multi-component vapor under conditions of interest. 10.3.5.5 Nucleation enhancement Increasing the substrate temperature generally enhances nucleation. This is mainly the result of the increased surface mobility of adatoms which have thus a better chance to be incorporated in existing embryos and subsequently form critical nuclei. The use of scratched substrates as well as ion bombardment are known to be efficient solutions for nucleation enhancement in CVD. For example, a bias pretreatment can be
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used for the growth of heteroepitaxial diamond films on monocrystalline silicon substrates by plasma CVD (Strckel et al. 1998). A dc bias voltage is applied to the substrate for a few minutes prior to deposition that results in a bombardment of the substrate by ions from the plasma. This special procedure can increase the nucleation density up to about 1012 cm -2, whereas the diamond nucleation density is only about 104 cm -2 on untreated substrates; however, only a very small fraction of the nuclei are epitaxially oriented. Surface active species (surfactants) can be intentionally incorporated during epitaxial growth in order to change the growth mode, and particularly, to promote the layer-by-layer mode (Markov 1997). Predeposition of one monolayer of such species (which are not incorporated into the film) may suppress the undesirable islanding observed in some cases and may also reduce the defects resulting from strain accommodation. 10.3.5. 6 Selective vapor deposition Selective deposition (Carlsson 1990, Gladfelter 1993) is one of the most interesting features of the CVD process. The deposition selectivity has most interesting potential for advanced submicron devices on patterned substrates, but other fields are emerging such as growth of artificial 2-D and 3-D materials. Selectivity usually refers to the process ability to control the deposition mechanisms in such a way that the film growth occurs only on one surface in the presence of others (area-selective deposition). However, several categories of selective growth systems exist, and CVD allows depositing simultaneously and selectively on different materials:
- different phases: phase-selective deposition different microstructures: microstructure-selective deposition - different chemical composition: chemical-composition-selective deposition -
Selective growth is based on differences in the initial interfacial reactions between the vapor and the various surfaces of the substrate and particularly on the differences in incubation times required for the nucleation. The selective growth period is thus up to the longest incubation time for nucleation. As seen above, epitaxial growth is carried out at supersaturation below that required for 3D heterogeneous nucleation, i.e., at near equilibrium conditions. Thermodynamic analysis is for that reason used to predict trends in selectivity (Cho et al. 1987, Madar and Bernard 1990, Engqvist et al. 1992). A new understanding of selectivity has been proposed for the area-selective deposition of epitaxial silicon in the presence of SiO2 blanket areas (Fitch 1994). At low pressures, hydrogen atoms passivate and block adsorption of silicon species, whereas, at high pressures, nuclei etch away faster than they grow. 10.3.5. 7 Theoretical and experimental studies o f nucleation in CVD For a long time, only analytical approaches based on empirical or semi-empirical relations have been used for the theoretical analysis of nucleation, but owing to the increasing power of computers, numerical simulations become more and more the rule for that purpose. Several numerical models have already been proposed for studying the nucleation process in CVD (Zachariah et al. 1996, Mahalingham et al. 1997), which provide key information on the reaction kinetics and coalescence mechanisms.
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However, it must be kept in mind that some basic experimental information is required in order to check the accuracy of the predictions. Many experimental studies on nucleation kinetics have been conducted in CVD. Atomic Force Microscopy (AFM) and Scanning Electron Microscopy (SEM) have often been used to determine the cluster density and size as a function of time and CVD process parameters. Since AFM and SEM measurements are usually performed ex-situ, several deposition experiments of variable duration are necessary in order to obtain the time evolution of the cluster characteristics. Nevertheless, some experimental analyses of nucleation kinetics have also been carried out in-situ using techniques such as ellipsometry (Drevillon 1993), or light-scattering (Sheldon et al. 1992). In this case, however, a schematic nucleation model must be used in order to interpret the experimental measurements.
10.3.6 Crystal growth Heterogeneous nucleation is the first stage, let us say the birth of a crystalline solid. The following stage, the ripening of this solid, is crystal growth. Like for nucleation, we have thought that it could be helpful to remind very briefly some of the most basic ideas of crystal growth before discussing some particular examples related to CVD. Far more complete and rigorous analyses of the fundamentals of crystal growth will be found in many excellent books and review articles (Hartman 1973a, Hurle 1993, Villain and Pimpinelli 1995, Levi and Kotrla 1997).
10.3.6.1 Mechanisms of crystal growth At an atomic scale, the surface of a growing crystal can be represented by terraces, ledges (or steps) and kinks as first suggested by Kossel (1927). The species to be incorporated are referred to as growth units. The most favorable sites (with the lowest energy) for incorporating the growth units are the kink sites (so-called half-crystal sites). Accordingly, the growth units first adsorb on a terrace, then diffuse until they either reach a ledge and then finally a kink site where their incorporation takes place, or desorb (TLK model). Nucleation should normally provide the source of repeatable steps on flat terraces, assuming that supersaturation is high enough for the nucleation barrier to be overcome; however, surprisingly, crystal growth can be observed even at very low supersaturation. The source of repeatable steps has remained unidentified until Frank (1949) proposed a spiral growth model in which the steps are continuously generated at the emergence of a screw dislocation.
10.3.6.2 Classification of crystal faces The various faces of a given crystal have different density of steps and kinks. Some faces present theoretically no steps (like terraces) and have a flat appearance: they are referred to as singular faces. Other faces, which have a rough aspect at an atomic scale present a more or less important density of steps (and kinks) and are referred to as non-singular faces. A more precise classification has been proposed from the concept of periodic bond chain (PBC) (Hartman and Perdok 1955). A given (hkl) crystal face is referred to as a flat face (F), stepped face (S), kinked face (K). F faces correspond to singular faces, whereas
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S and K faces correspond to non-singular faces. It must be mentioned that F faces have a fiat aspect only below the so-called roughening transition.
10.3.6.3 Crystalline morphology Some experimental observations have suggested that a crystalline form {hkl } has a greater morphological importance when the spacing between the (hkl) planes, dhkl, is large. This result, known as the Bravais-Friedel law, has been further extended by Donnay-Harker (Hartman 1973a) and Donnay-Donnay (Hartman 1973a). The concept of attachment energy Eatt, the energy (per molecule) released when a new slice is attached to a crystal face, is most often used to obtain the crystalline morphologies (habits) of large crystals. In the Hartman-Perdok theory (Hartman and Perdok 1955), it is proposed that (i) the higher Eatt for a given (hkl) face, the less the morphological importance of this face in the growth form, (ii) the growth rate of F faces, R• is proportional to Eatt. This law has been found to be satisfying in many cases (Hartman and Bennema 1980), although the relation: R• (x Ea2tt, should be preferred for needle-like crystals (Van der Sluis and Kroon 1990). Using the above relations, it is possible to estimate the growth rate of the various slow-growing faces, and then to construct the crystal habit (see below). 10.3.6.4 Morphology and normal growth rate of crystal faces Since incorporation of growth units is energetically more difficult on terraces than at kink or step sites, F-faces should theoretically be the slowest-growing crystal faces. This is experimentally verified unless the reactivity of non-singular faces is lowered due to strong surface reconstruction and/or adsorption of impurities (elements which are not incorporated in the solid structure). Crystal habit is controlled by the slowest-growing faces, that is immediately obvious from an examination of Fig. 10.11. Starting with a morphology in which all possible crystal faces exist, and keeping the growth rate of these faces constant, it can be inferred that: (i) the fastest growing faces disappear quickly, (ii) two adjacent faces grow independently when the angle 0 between them is 90 deg (Alexandru 1969). Stable growth of CVD layers. Aspar et al. (1993) and Rodriguez-Clemente et al. (1993) have studied the surface morphology of polycrystalline A1N films prepared by CVD. They have first used the Donnay-Harker (Hartman 1973a) approach in order to draw a theoretical equilibrium morphology assuming that the growth rate of a given (hkl) face is inversely proportional to its lattice spacing dhkl. From the PBC of A1N structure deduced from literature analysis, they have concluded that the slow-growing F faces are the non-polar face { 100 } and the polar faces { 101 }, { 10T }, {001 }, {001 }. In this case, the two approaches Donnay-Harker (Hartman 1973a) and Hartman-Perdok (1955) have given similar results, but the real morphologies observed did not match exactly the theoretical-one, and varied with the operating parameters. To account for these discrepancies, Rodriguez-Clemente et al. (1993) have discussed the reactivity of A1N crystal faces using a broken bond model. They have reached the conclusion that the negative polar faces { 101 } and {001 } should be more reactive than their opposite positive polar faces, and that they should be inhibited by hydrogen adsorption. Some other authors (Dollet et
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F. Teyssandier a n d A. Dollet
R1
~
l
RFI
'~RF2
FX~'l IRr3 F lll a)
b)
Figure 10.11. Left: time evolution of crystalline morphology from the relative growth rate of possible faces. (a) F faces with comparable growth rates (all faces survive), (b) S face between two F faces (the fast-growing S face disappear). Right: illustration of the competitive growth process (2D). 2 identical cubic crystals are considered. The growth rates perpendicular to the substrate surface are R1 et R2 for crystal 1 and 2 respectively. Each line represents a time step.
(oo-l)~ {lOl} 33 ~
a)
I subs rate b)
Figure 10.12. Surface morphology of (112) oriented AIN, prepared by CVD (Dollet et al. 1997). (a) SEM picture (left). (b) Theoretical morphology, using the SHAPE program (Dowty 1991) (right).
al. 1997) have also studied the growth of polycrystalline A1N films by CVD, and have observed surface morphologies essentially composed of pyramidal crystals in which the { 101 } and {001 } forms dominate (Fig. 10.12), in agreement with the theoretical findings of Rodriguez-Clemente et al. (1993). Owing to reactor modeling results, it has also been suggested that the surface coverage by ammonia, or perhaps the deposition rate, may play a key role in the preferred orientation. Unstable growth. During crystal growth from the vapor, some fluctuations (thermal, chemical, etc.) may disturb the moving solid-gas interface. If these fluctuations cannot be smoothened, they may become amplified and morphological instabilities will result. Dendrites, skeletal crystals (Nanev 1994), etc. are good examples of crystal morphologies
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which form under unstable growth conditions. It has been shown in previous atomistic simulations (Xiao et al. 1991) that a growing crystal may exhibit compact faceted or open dendritic morphologies depending on various factors such as supersaturation, mean free path or temperature. The apparition of protrusions has been observed to result from volume diffusion whereas the anisotropic surface kinetics has been found to exert a stabilizing effect. Although atomistic simulation seems to be the most powerful method to tackle the problem of morphological instability, it should be mentioned that continuum equations with moving boundary conditions have also been solved for the same purpose in CVD (Palmer and Gordon 1989) but this approach has generally been restricted to the growth of non-faceted crystals or amorphous materials.
10.3.6.5 Relation between preferred orientation and morphology of polycrystalline films Polycrystalline films are composed of a large number of individual crystals which are generally not oriented in the same direction. A mechanism known as evolutionary selection or competitive growth has been proposed by van der Drift (1967) to explain the preferential orientation of deposited polycrystalline layers. The nuclei are assumed to be randomly oriented on the substrate surface and during their growth, some crystals come into contact and a geometrical selection operates. Only the crystals having their fastest direction of growth nearly perpendicular to the substrate surface will finally emerge at the film surface. The crystals which are less favorably oriented are gradually buried by the others, and, as a result, the grain size and orientation sharpness increase with increasing film thickness. This is illustrated in Fig. 10.11, where it can be seen that the crystal on the fight will be soon buried by the one on the left. An evolutionary selection process has been proposed by Wild et al. (1990), to explain the growth of polycrystalline diamond films by CVD, which agrees well with the observations from X-ray texture analyses, showing that the initial orientation of the tiny crystals is random, and that the sharpness of the preferred (110) orientation increases with increasing film thickness. A 2D geometrical model has been formulated by the authors in order to simulate the growth of the diamond films.
10.3.6.6 Theoretical studies of crystal growth from the vaporphase Evolution of crystalline morphology, grain size and texture, formation of defects inside the film (voids, vacancies, etc.) during film growth results from a large number of complex phenomena at the atomic scale which are difficult to apprehend without numerical models. Beside continuous geometrical models (Wild et al. 1990, Tang et al. 1990, Barrat et al. 1996), atomistic models such as ballistic aggregation models (Henderson et al. 1974, Dew et al. 1992), molecular dynamics (MD) and Monte Carlo simulation models (Dawnkaski et al. 1996, Smith 1997) become more and more used for understanding the exact mechanisms of solid phase formation in CVD. 10.4. C O N C L U S I O N S
Modern manufacturing methods are placing more emphasis on high performance techniques. An accurate control of these techniques requires an improved knowledge of the physico-chemical phenomena that are involved in the growth of materials. We have tried
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to summarize in this review the current status of mechanistic research in the field of materials grown from a vapor phase in the presence of a purely thermally activated process. In particular, it has been emphasized that numerical modeling supported by in situ diagnostics is a key tool for the understanding and improvement of the CVD process. However, things become more problematic for the heterogeneous phenomena since they are even more complex than their homogeneous counterparts. Calculations assuming thermodynamic equilibrium can be run in order to observe, at least qualitatively, the influence of the operating parameters on the composition of the solid-gas system. Nevertheless, since CVD remains in essence a non-equilibrium process, whenever possible, a kinetic analysis of both the homogeneous and the heterogeneous processes should be preferred. It is worth to note that an increasing amount of work is dedicated to the numerical simulation of nucleation and crystal growth in CVD. Several interesting simplified approaches have been proposed for that purpose. However, sophisticated atomistic simulations are increasingly used and already provide invaluable information though their performances are still limited by the computing power available. Unfortunately, in spite of the progress accomplished in this field, evaluating the kinetics of the elementary surface processes will undoubtedly remain, for a rather long time, the most difficult task. The understanding and full-scale modeling of the whole process, including the formation of film microstructure, appears to be the most important challenge of the next years for research in chemical vapor deposition.
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Sydney, Toronto) Bernard, C. and Madar, R. (1990) Mater. Res. Soc. Symp. Proc., 168, 3. Bernard, C. and Madar, R. (1991) Surf. Coat. Technol., 49, 208. Besmann, T. M., Stinton, D. E and Lowden, R. A. (1989) MRS Bulletin, 14, 45. Bird, R. B., Stewart, W. E. and Lightfoot, E. N., (1960) Transport Phenomena. (John Wiley & Sons, New York,
Chichester, Brisbane, Toronto, Singapore). Blocher, J. M. (1967) in Vapor Deposition, eds. Powell, C. E, Oxley, J. H. and Blocher, J. M., (John Wiley & Sons, New York), p. 3 Blocher, J. M. (1974) J. Vac. Sci. Technol., A, 11,680. Bryant, W. A. (1977) J. Mater. Sci., 12, 1285. Buss, R. J., Ho, E, Breiland, W. G. and Coltrin, M. E. (1988) J. Appl. Phys., 63, 2808. Cale, T. S. and Raupp, G. B. (1990) J. Vac. Sci. Technol., B, 8, 1242. Carlsson, J. O. (1990) Crit. Rev. Solid State Mater. Sci., 16, 161. Chakraverty, B. K. (1973) in Crystal Growth: An Introduction, ed. Hartman, E (North-Holland, Amsterdam), p. 50
Cheng, D. J., Sun, W. E and Hon, M. H. (1987) Thin Solid Films, 146, 45. Cho, K. I., Yang, J. W., Park, C. S. and Park, S. C. (1987) in Proceedings of the X International Conference on CVD, Honolulu USA, eds. Cullen G. W. and Blocher J. M. (The Electrochemical Society, Pennington), p. 379.
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Coltrin, M. E., Kee, R. J. and Evans, G. H. (1989) J. Electrochem. Soc., 136, 819. Dawnkaski, E. J., Srivastava, D. and Garrison, B. J. (1996) J. Chem. Phys., 104, 5997. De Sainte Claire, E, Song, K., Hase, W. L. and Brenner, D. W. (1996) J. Phys. Chem., 100, 1761. Dean, A. M. (1985) J. Phys. Chem., 89, 4600. Dew, S. K., Smy, T. and Brett, M. J. (1992) J. Vac. Sci. Technol., B, 10, 618. Dollet, A., Casaux, Y. and Rodriguez-Clemente, R. (1997) in Proceedings of the XIV International Conference and XI European Conference on CVD, Paris France, eds. Allendorf, M. D. and Bernard, C. (The Electrochemical Society, Pennington), p. 286. Dowty, E. (1991) Shape Computer Program, ESM Software (Hamilton, OH, USA). Drevillon, B. (1993) Prog. Crystal Growth Charact. Mater., 27, 1. Engqvist, J., Myers, C. and Carlsson, J. O. (1992) J. Electrochem. Soc., 139, 3197. Eversteyn, E C., Severin, E J. W., Van den Brekel, C. H. J., and Peek, H. L. (1970) J. Electrochem. Soc., 117, 925. Fisher, E. R., Ho, E, Breiland, W. G. and Buss, R. J. (1992) J. Phys. Chem., 24, 9855. Fitch, J. T. (1994) J. Electrochem. Soc., 141, 1046. Forst, W. (1973) Theory of Unimolecular Reactions (Academic Press, New York). Fotiadis, D. I. and Jensen, K. E (1990) J. Cryst. Growth, 102, 743. Fotiadis, D. I., Kremer, A. M., McKenna, D. R. and Jensen, K. E (1987) J. Cryst. Growth, 85, 154. Fotiadis, D. I., Kieda, S. and Jensen, K. E (1990) J. Cryst. Growth, 102, 441. Frank, E C. (1949) Discuss. Faraday Soc., 5, 48. Frank, E C. and Van der Merwe, J. H. (1949a) Proc. Roy. Soc. Lond. A., 198, 205. Frank, E C. and Van der Merwe, J. H. (1949b) Proc. Roy .Soc. Lond. A., 200, 125. Gates, S. M. (1996) Chem. Rev., 96, 1519. Gilbert, R. G. and Smith, S. C. (1990) Theory of Unimolecular and Recombination Reactions (Blackwell Scientific Publications, Oxford). Gladfelter, W. L. (1993) Chem. Mater., 5, 1372. Gokoglu, S. A., Kuczmarski, M., Tsui, E and Chait, A. (1989) in Proceedings of the VII European Conference on CVD, Perpignan France, eds. Ducarroir, M., Bernard, C. and Vandenbulcke, L. (Colloque C5 Journal de Physique Supp6ment No. 1-2, Les Editions de Physique, Les Ulis France), p. 17. Goto, T., Tsuneyoshi, J., Kaya, K. and Hirai, T. (1992) J. Mater. Sci., 27, 247. Hampden-Smith, M. J. and Kodas, T. T. (1995) Adv. Mater. CVD, 1, 40. Hartman, E (1973a) Crystal Growth: An Introduction, ed. Hartman, E (North-Holland, Amsterdam). Hartman, E (1973b) in Crystal Growth: An Introduction, ed. Hartman, E (North-Holland, Amsterdam), p. 367 Hartman, E and Bennema, E (1980) J. Cryst. Growth, 49, 145. Hartman, E and Perdok, W. G. (1955) Acta Cryst., 8, 521. Henderson, D., Brodsky, M. H. and Chaudhari, P. (1974) Appl. Phys. Lett., 25, 641. Herlin, N., Pealat, M., Lefebvre, M., Alnot, E and Perrin, J. (1991) Surf. Sci., 258, 381. Hess, D. W., Jensen, K. E and Anderson, T. J. (1985) Rev. Chem. Eng., 3, 97. Hitchman, M. L. (1980) J. Cryst. Growth, 48, 394. Holstein, W. L. (1992) Prog. Crystal Growth and Charact., 24, 111. Hurle, D. T. J. (1993) Handbook of Crystal Growth, ed. Hurle, D. T. J. (Elsevier Science, Amsterdam). IslamRaja, M. M., Capelli, M. A., McVittie, J. E and Saraswat, K. C. (1991) J. Appl. Phys., 70, 7137. Jasinski, J. M. and Estes, R. D. (1985) Chem. Phys. Lett., 117, 495. Jasinski, J. M., Meyerson, B. S. and Scott, B. A. (1987) Ann. Rev. Phys. Chem., 38, 109. Jensen, K. E (1987) Chem. Eng. Sci., 42, 923. Jensen, K. E (1989a) J. Cryst. Growth, 98, 148. Jensen, K. E (1989b) Adv. Chem. Ser., 221, 199. Jensen, K. E (1994) in Handbook of Crystal Growth, Vol.3b, ed. Hurle, D. T. J. (Elsevier Science, Amsterdam), p. 541 Kee, R. J., Rupley, E M. and Miller, J. A. (1989) Chemkin-II: a fortran chemical kinetics package for the analysis of gas-phase chemical kinetics, SAND89-8009B.UC-706, 1. Kim, K. S. and Pratsinis, S. E. (1988) AIChE, 34, 912. Kleijn, C. R. (1991) in Proceedings of the VIII European Conference on CVD, Glasgow U.K., eds. Hitchman, M. L. and Archer, N. J. (Colloque C2 Journal de Physique Suppl6ment No. 7), p. 19. Koseki, S. and Ishitani, A. (1992) J. Appl. Phys., 72, 5808. Kossel, W. (1927) Nachr. Ges. Wiss. GOttingen, math-phys. Klasse, 135.
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Chapter 11 Combustion Synthesis 11.1. 11.2. 11.3. 11.4. 11.5. 11.6. 11.7. 11.8. 11.9.
Introduction Thermodynamic Considerations Kinetic Considerations Field-Activated Combustion Synthesis The "Azide" Process SHS Reactions in Synthesizing Ti3SiC2 Controlled Reactions in the Ti-B Binary System Auto-Ignition Synthesis of Nanocrystalline Oxides Non-Equilibrium Effects 11.10. Concluding Remarks Acknowledgments References
289 291 294 295 297 299 301 304 307 307 308 308
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Chapter 11 Combustion Synthesis S.B. BHADURI and S. BHADURI
11.1. I N T R O D U C T I O N
The technique of self-propagating high-temperature synthesis (SHS) is being used all over the world, not only in synthesizing many advanced materials but also in various industrial applications. Other terminologies used to describe the process include combustion synthesis (CS), gasless combustion, etc. The current interest in the topic can be credited to the pioneering work of Merzhanov and his group in Russia, starting in the early seventies (Merzhanov and Borovinskaya 1975). However, contrary to the popular belief, the process itself has a much longer history. The earliest mention of a combustion phenomenon in materials synthesis can be found in 1825. Berzelius (1825) reported that Zr metal powders burn to form the corresponding oxide. The thermodynamics of the process was explained by Goldschmidt (1898). In fact, Goldschmidt can be rightly called the originator of the aluminothermic process, a subgroup of SHS reactions. During the next 100 years, there were some sporadic industrial discoveries utilizing the combustion phenomena (Alexander 1941, Krapf 1964). Despite those discoveries, there was no widespread scientific interest in the process until vigorous research was initiated by Merzhanov and his colleagues. In the USA, it is generally believed that the work in this field started with support from the Advanced Research Project Agency (ARPA) in the mid- 1980s. Long before that, however, Walton and'Poulos (1959) started utilizing the aluminothermic processes for synthesizing cermets. This activity went on for several decades. As the SHS related activities took firm root, excellent reviews started appearing in the literature. Frankhouser et al. (1986) published an extensive monograph on the subject. This monograph was interesting in that it had an industrial bias and was a balanced review of the work done in the Soviet Union and elsewhere in the West. Soon after that, two of Munir's reviews appeared in 1988 and 1989 (Munir 1988, Munir and Anselmi-Tamburini 1989). Munir's reviews together describe the important aspects of the then Soviet work, which by then was already a decade and a half old. Yi and Moore (1990), as well as Merzhanov (1990), reviewed somewhat similar topics. The paper by Koizumi is an important contribution in reviewing the Japanese work (Koizumi 1990). A short paper by Hlavacek (1991) points out the historic aspects of the process and mentions the earliest work by Berzelius, Goldschmidt and others. Other reviews also appeared in the literature (Subrahmanyam and Vijaykumar 1992) including the most recent ones by Moore and Feng (1995a, b). However, in spite 289
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of almost ten reviews in the field within the last decade, the contents of the reviews have not progressed much further since the publication of the review by Munir and AnselmiTamburini (1989). Since some of the information included in this review was actually quoted from the earlier Russian references, it is believed that the data need confirmation in view of the more recent reports. Furthermore, there have been some interesting developments in the field, which need proper attention and dissemination. Finally, in spite of a great deal of applications in industry, worldwide skepticism towards the process has prevailed. Hence, there is a need for reviewing some of the recent and interesting developments in the field. The present attempt will not be comprehensive. Instead, it will focus on the synthesis of newer materials and developments of new techniques with a view towards ultimate application. Unless large-scale applications of SHS are realized, skepticism will remain. We would like to point out that the industrial applications of SHS are by no means insignificant in this country. But elsewhere in the world, particularly in Russia, other countries of the CIS, Europe (notably Spain), and Asia (specifically Japan, China, Korea, Taiwan and India), industrial applications have been substantial. Combustion synthesis, or Self-propagating High temperature Synthesis (SHS), relies on the ability of highly exothermic reactions to be self-sustaining. In a typical combustion synthesis reaction, the mixed reactant powders are pressed into a pellet of a certain green density and subsequently ignited either locally at one point (propagating mode) or by heating the whole pellet (simultaneous combustion mode). The reaction can be ignited in a number of ways including laser irradiation, radiant flux, resistance heating coil, spark, chemical oven, etc. The exothermic reaction is initiated at the ignition temperature T/g and generates heat, which is manifested in a maximum or combustion temperature, Tc. The maximum temperature obtained can be controlled by adding diluents, the nature and amount of which can be determined depending on the maximum temperature desired. The products of combustion synthesis are usually extremely porous, typically 50% of the theoretical density. Techniques such as hot pressing, hot isostatic pressing, and shock waves have been employed to densify the products. Compared with conventional processing methods, combustion synthesis has the following advantages: (1) the high temperatures generated can volatilize low-boiling-point impurities and therefore produce high-purity products, (2) the simple exothermic nature of the process avoids the need for expensive processing facilities and equipment, (3) the short processing times result in low operating and processing costs, (4) the high thermal gradients and rapid cooling rates can result in the formation of non-equilibrium phases, (5) inorganic materials can be synthesized and consolidated into a final product in one step by utilizing the chemical energy of the reactants. Because of these advantages of the process, it has been possible to produce novel and improved materials with improved mechanical, electrical, optical, and chemical properties. It has been estimated (Yi and Moore 1990) that over 500 compounds have been synthesized by the combustion synthesis process. These include resistive heating elements, abrasives, cutting tools, and polishing powders, shape memory alloys, high-temperature intermetallics, electrodes for electrolysis of corrosive media, thin films and coatings, functionally-graded materials, and composite materials.
Combustion Synthesis
291
This review will start off with a discussion of the thermodynamic aspects of the process. This will be followed by some specific examples of materials development such as nitrides, silicides, metal matrix composites, and oxides.
11.2. T H E R M O D Y N A M I C
CONSIDERATIONS
Once ignited, high temperatures can be achieved in a very short time owing to the highly exothermic nature of the reactions. The enthalpy of the synthesis reaction for the propagating case, AHreaction, is equal to AHf,298, the standard enthalpy of formation of the product. For an adiabatic reaction with the reaction temperature denoted by Tad, the heating of the product from 298 K to Tact takes place in response to the input of A Hreaction. Depending upon whether Tad is lower than, equal to, or higher than the product melting temperature Tm, the following three cases arise: AHf,298 --
f T~d
Cp(product) dT
if Tad Tm
(11.3)
where, v is the fraction of the liquid phase in the mixture of liquid and solid phases, since both of them are expected to exist at the melting point. For many compounds, the thermodynamic data required to calculate Tad are available in the literature. However, agreement between the experimentally obtained data and theoretically calculated values is lacking. This is because the reactions are not truly adiabatic. At best, the calculations predict the upper limit of Tad. Nevertheless, such calculations provide important guidelines for maintaining the exothermic state of the reactions. Based on the observations of a number of reactions (Munir and AnselmiTamburini 1989, Moore and Feng 1995a, b), a linear correlation has been found between AHf,298/Cp,298 and Tad. It has been suggested in the literature that the reaction will not be self-sustained if AHf, z98/Cp,298 < 2000~ (or Tad values are less than 1800~ (Munir and Anselmi-Tamburini 1989). Table 11.1 shows the results of Tad calculations for various silicides. AH/,298 for the materials under consideration was obtained from (Barin et al. 1977, Chart 1973). Cp (in U/mole K) is written as A + B T 4-C T 2, where the coefficients A, B and C are taken from (Barin et al. 1977) and are listed in Table 11.1. If substantial melting takes place, (as in the case of TisSi3 (Rupp et al. 1992)), the calculations may require the value for Cp (liquid), which is taken as 33.74.n J/mole K, where n is the number of atoms in a molecule and is equal to 8 for TisSi3. Table 11.2 compares the as-predicted values of Tad for various silicides with those reported (Subrahmanyam and Vijaykumar 1992), which are based on the Russian data. Even though the Russian values are available for only some materials,
292
S.B. Bhaduri and S. Bhaduri
Table 11.1. Thermodynamic data and adiabatic temperatures, Tad, for various silicides.
Compound
A
B
C
A Hf
A Hm
(kJ/mol)
(kJ/mol)
M.P. (K)
Tad A H f / Cp
(K)
ZrSi
44.9
8.74
2.95
153.92
58.73
2368
3.84
2909
ZrSi2
62.8
15.27
-2.78
158.50
44.75
1790
2.47
2228
188.0 47.84
30.58 11.36
-14.9 -5.41
572.42 128.96
59.59 40.18
2483 1843
3.17 2.86
2739 2349
TiSi2
70.01
17.47
8.99
133.54
43.44
1773
2.05
1808
Ti5Si 3 CrSi
46.75 51.71
10.70 8.69
-4.8 -8.36
138.40 54.60
178.88 35.12
2408 1730
2.90 1.21
2528 1215
Zr5Si 3 TiSi
CrSi2 Cr5Si 3
65.23 197.5
22.38 49.00
-7.74 -25.5
79.62
42.39
1730
1.26
1252
221.73
33.02
1920
1.21
1233
ReSi
52.29
9.57
-3.74
52.42
50.16
2153
1.03
997
ReSi2
67.39
10.98
-6.07
89.86
73.22
2253
1.41
1420
Re5Si 3 MnSi
189.7 49.00
44.93 12.69
-13.9 -6.36
156.42 78.29
46.89 29.41
2233 1548
0.83 1.71
901 1553
Mn5Si 3
200.2
53.87
-19.5
199.68
21.71
1573
1.03
1100
12.57
-2.78
137.28
69.17
2203
2.14
1863
449.28
73.51
2753
2.49
2222
NbSi2 Nb5Si 3 TaSi2 Ta5Si 3
62.82 188.1 72.84 178.6
30.61
-14.9
7.67
-9.01
118.39
89.03
2473
1.82
1671
38.90
-8.86
332.80
72.10
2773
1.85
1772
MoSi2
67.44
11.90
-6.53
130.96
74.83
2303
2.06
1872
Mo5Si3 W5Si 2
182.3 67.43
34.84 10.98
-11.9 -6.06
307.84 92.19
56.65 85.16
2463 2433
1.72 1.44
1672 1471
W5Si 3
178.6
38.98
-8.82
133.79
63.53
2593
0.74
945
11.61
-9.36
124.8
51.75
1953
2.00
1792
459.26
49.08
2283
2.26
1703
VS2 V5Si 3
71.05 187.3
118.1
-17.2
A, B, and C in J mol-1 K - l , 10 -3 J mol-1 K - l , 105 J mol-1 K - l , respectively.
the agreement is quite sound for most of the cases provided, while there are substantial differences in the predicted values of Tad in the case of TiSi and CrSi2. Tables 11.3 and 11.4 show similarly the calculated Tad values for various carbides as well as borides. Again, there are agreements as well as discrepancies between our data and the Russian predictions. Notable differences exist in the carbides of ZrC and HfC and the borides CrB2 and MoB. Similar comments can also be made in the case of various borides. At this point several comments can be made. First, since these predictions serve as important guidelines, the new values should be carefully considered before initiating any reaction. Second, combustion of stoichiometric TiB composition has a peculiar result. Even though the predictions favor reactions of higher exothermicity than that of the TiB2 reaction, in practice such reactions are quite sluggish. This is due to the formation of a significant amount of the liquid phase, which can be attributed to the large quantity of unreacted Ti remaining within the system. This occurrence can be exploited in fabricat-
Combustion Synthesis
293
Table 11.2. Comparison of predicted Tad values of selected silicides.
Material
Tad (K) Present work
Tad (K) Russian work
TiSi TiSi2 Ti5Si 3 ZrSi ZrSi 2 Zr5Si 3 VSi2 V5Si 3 NbSi2 Nb5Si 3 TaSi2 Ta5Si3 CrSi CrSi2 Cr5Si3 MoSi2 Mo5Si 3 WSi 2 WsSi 3
2349 1808 2528 2909 2228 2739 1792 1703 1863 2222 1671 1772 1215 1252 1233 1872 1672 1471 945
2000 2500
2800
1900 1800
1800 1900 1500
ing m e t a l m a t r i x c o m p o s i t e s . Third, it is s e e n that, the f o r m a t i o n o f several c o m p o u n d s is quite s l u g g i s h as w e l l ( c o r r e s p o n d i n g to Tact t e m p e r a t u r e s in the r a n g e o f 1000 K). T h e r e are m a n y strategies that are f o l l o w e d in initiating r e a c t i o n s in such cases. C o m m o n m e t h ods to ignite such r e a c t i o n s are the " E x p l o s i o n M o d e " , the " C h e m i c a l O v e n " , and u s i n g
Table 11.3. Calculated adiabatic temperatures for selected carbides.
Material TiC ZrC HfC VC NbC TaC Mo2C WC
Tad (K) Present work
Tad (K) Russian work
3290 3758 4512 2228 2821 2901 1017 1163
3210 3400 3900 2400 2800 2700 1000 1000
294
S.B. Bhaduri and S. Bhaduri Table 11.4. Calculated adiabatic temperatures for selected borides. Tad (K) Present work
Tad (K) Russian work
TiB
3388
3350
TiB2
3198
3190
ZrB 2
3322
3310
HfB 2
3662
3520
VB 2
2546
2670
VB
2490
NbB 2
2715
2400
TaB 2
2728
2700
CrB 2
1475
2470
CrB
1646
Material
MoB 2
1480
MoB
2659
1800
WB
1617
1700
FeB
1605
Fe2B
1172
CoB
1939
Co 2B
1721
NiB
2026
Ni2B
1106
ternary reactions with several different phases taking part in the reaction. A recent advancement in this field is the use of Field Assisted Combustion Synthesis (FACS), which will be described later. Fourth, Tad calculations in the case of nitrides have not been discussed since it is felt that the "Azide" process has better capabilities, which again will be discussed later. Finally, the thermodynamics of the "Auto Ignition" or "Combustion Synthesis in the Auto Ignition" mode will also be described.
11.3. K I N E T I C C O N S I D E R A T I O N S
In order to understand the kinetics of the process, a cylindrical compact is placed vertically in a reaction chamber and ignited from the top. A self-supporting reaction wave is initiated under suitable conditions and its steady speed of combustion is measured. Of interest is the dependence of the reaction upon the variables of the experiment (such as specimen diameter, specimen composition, size of the initial pressed particles, and the extent of dilution with an inert component such as the reaction product itself). These are well documented in Munir (1988), Munir and Anselmi-Tamburini (1989), and Moore and Feng (1995a, b). Azatyan et al. (1979) have derived an expression for the rate of propagation of a planar combustion front based upon an elementary model of the process. In this model, the
Combustion Synthesis
295
calculations are performed assuming a uniform radial distribution (a planar front). With the additional assumptions of there existing a single heat source and that the only heat loss from the body is by conduction and radiation, the generalized form of Fourier's one dimensional heat equation takes the following form:
OT _ kO2T
Cpp Ot
ot
a T4
O~ 4- Qpdp(T, rl) - 2-(Tr - To) - 2e--(r
O(T, o) -
Orl
-
Td)
(11.4)
(11.5)
Ot
where Cp is the heat capacity, p, the density, T, the temperature, To, the ambient temperature, 77, the fraction reacted, t, the time, k, the thermal conductivity, x, the axial distance, Q, the heat of reaction, or, the heat transfer coefficient, r, the sample radius, E, the emissivity coefficient, a, the Stefan-Boltzman constant, and 4~, the kinetic function. Neglecting the transient effects from the ignition surfaces and the heat losses, the above equation can be rewritten as:
OT kO2T Cpp 0---;~ 4- Qp~(T, 7) _
(11.6)
The solution to the above equation can be written as
u2 - a n Q p
E Koexp
RTad
(11.7)
where an is a constant depending upon the order of the reaction, u is the velocity, E is the activation energy and K0 is a constant. The apparent activation energy of the reaction can be found from a graph of ln(u / Tad) vs. 1/ Tad. 11.4. F I E L D - A C T I V A T E D C O M B U S T I O N
SYNTHESIS
During the discussion on thermodynamic calculations, it was mentioned that the reactions are frequently quite sluggish and do not propagate in a self-sustaining manner. Several alternate methods for propagating such reactions were also suggested. Unfortunately, these alternative methods lead to other undesirable phases. A specific example of this phenomenon is the combustion of MoSi2 in the Mo-Si binary system. As noted earlier, the Tad value for MoSi2 formation is at the threshold of self- propagation. In order to enhance the propagating reaction, the explosion mode can be tried, which in turn leads to the formation MosSi3, a compound that generally does not have the high temperature oxidation resistance of MoSi2. To alleviate this problem, Munir and his co-workers ignited the combustion synthesis reactions in the presence of electric fields (Gedevanishvili and Munir 1994, Feng and Munir 1994, Shon and Munir 1995, Munir et al. 1995, Feng and Munir 1995a, b, Gedevanishvili and Munir 1995, Xue and Munir 1996, Munir 1996). This process is now referred to as field-activated combustion synthesis (FACS). The idea here
296
S.B. Bhaduri and S. Bhaduri Table 11.5 Materials synthesized using FACS. Materials
Reference
SiC MoSi2-SiC MoSi2-Nb, MoSi2-ZrO 2 B4C-TiB 2 W5Si 3, WSi2
Feng and Munir (1994, 1995a, b) Gedevanishvili and Munir (1994) Shon and Munir (1995) Xue and Munir (1996) Gedevanishvili and Munir (1995)
is to create localized Joule heating in the reaction zone. Indeed, this occurred when monitoring the voltage, current, and the resistance of the sample, which essentially remained constant during propagation. A second discovery was the apparent dependence of wave propagation on the field itself, as clearly demonstrated by the propagation ceasing whenever the field was shut off. Thirdly, the propagation velocity was directly related to the magnitude of the field. A fourth observation noted a threshold value of the field, below which combustion did not take place, while at relatively high values of the field, combustion took place spontaneously. Finally, the actual mechanisms for such reactions could be studied by turning off the field and stopping the propagating reaction at any time. Such reactions are easy to model by adding Joule heating term of m E 2 to the right hand side of Eq. (11.6). Feng and Munir (1995a) modeled such a reaction using a twodimensional version of the modified Eq. (11.6) and the results are shown in Figure 11.1. The graphs show profiles of current densities near the top, center and bottom of the sample. The profiles show that a planar front is moving through the sample. The modeled experimental configuration is shown in Figure 11.2. The system consists of spring-loaded graphite electrodes between which the powder compact is placed. The compact is then ignited from the side. Table 11.5 shows some of the materials produced by the FACS process. Gedevanishvili and Munir (1994) studied the SHS reactions in the MoSi2-SiC system. In such a case, the formation of both constituents is sluggish. However, beyond the threshold voltage, a desirable composition of MoSi2-SiC composite was produced. The threshold field level decreased as the percentage of Si and C in the system was increased. Feng and Munir (1994, 1995a and b) used the process to react Si and C to form SiC. The threshold voltage was 9.5 V and simultaneous combustion occurred for voltages greater than 30 V. Shon and Munir (1995) reported FACS processing of MoSiz-ZrO2 and MoSiz-Nb composites. The secondary phases are added to enhance the mechanical properties with the application of field, without the formation of undesirable phases. In the absence of the field, however, the reactions do not self-sustain and undesirable phases form. The same result occurs in the case of B4C-TiB2 composites. While the formation of B4C is sluggish, the reaction to form TiB2 is highly energetic. Under normal circumstances, the thermodynamic equations predict the propagating reactions only in restrictive compositions. However, with FACS, propagating reactions can take place easily (Xue and Munir 1996). FACS is also useful in producing WsSi3 with a Tad of 954 K. Since a propagating reaction is not normally expected here, a high voltage is necessary.
Combustion Synthesis
297
top
~--.-4~..$s
I000 |
500
.
:"
?.Us "~Ss "less
1000
~"
500
.-. 1000 E 500
0
0.0
0.5
1.0
1.5
2.0
x (cm) Figure 11.1. Currentdensity profiles from side ignition (Feng and Munir 1995a).
11.5. THE "AZIDE" PROCESS Nitrides are typically formed by igniting a reaction between a metal and nitrogen gas. There are two competitive processes that take place during such reactions: the kinetics of reaction and the permeation of the gas. Typically, permeation of the gas is the rate-controlling step. The conversion efficiency can be increased by making use of two strategies: increasing permeation by increasing the porosity of the compact or by increasing the gas pressure. Frequently, the formation of a molten phase adversely affects the permeation of the gas. Inclusion of diluents helps to reduce Taa values and without allow-
S.B. Bhaduri and S. Bhaduri
298 Graphite
Laser beam Igniter!
ON2N
()
,a~.... ~:,.;~,J, PrO luct Rea tants
power (
O
~
?
Wave Fro ~ ' K~actants 'Product
Figure 11.2. Schematic representations of field activated combustion synthesis (Feng and Munir 1995a).
ing a molten metal to form. Despite these strategies, conversion to the final nitride may not be significant. For example, if the initial starting materials have low melting temperatures, then controlling their melting is extremely difficult. A case in point is the synthesis of A1N by SHS, during which substantial melting of A1 is a problem. Solid-state nitrogen sources can solve this problem. One such source of nitrogen is sodium azide, NAN3. The material is cheap, readily available, and is unstable. Holt (1984a, b) pioneered the "Azide" process for making nitrides by the SHS process. Originally, the process was developed for making nuclear fuels. Typical reactions with the respective Tad values are written as: 3U + NaN3 = 3UN + Na 3UO2 + 6Ca + NaN3 = 3UN + 6CaO + Na
Tad = 3669 K Tad -- 2824 K
During the process, the nitrides form and metallic sodium vaporizes. In the case of the second reaction, CaO is extracted from the product and is expected to offer much better control. Similar ideas can be used in making important ceramic materials as well, which are described by the following equations: A1 + NaN3 = A1N + N2 + Na 3Si + 4NAN3 = Si3N4 + 4N2 + 4Na Ti + NaN3 = TiN + N2 + Na
Tad -- 3010 K Tad = 2537 K Tad = 3022 K
Table 11.6 shows the Tad values calculated in the case of the "Azide" process vs. the conventional gaseous N2 reactions. In some cases the calculated Tad values for the "Azide" process are lower (e.g., in the cases of TiN and Si3N4) than the conventional situation, while in other cases the Tad values by the "Azide" process are higher. Nevertheless, generating N2 in-situ is an important advantage, so much so that excess pressure of N2 is not required (Holt 1984a, b, Amosov et al. 1992). In spite of these potential advantages, the growth in this field has not been significant. Recently, Amosov et al. (1992) used this process for synthesizing several nitrides. They reported that NH4C1, when incorporated into the powder mixture, sometimes catalyzes the reaction. Lee and Chung (1997) explained the catalytic effect of NH4C1 on the basis
299
Combustion Synthesis
Table 11.6 Comparisonof calculated Tad values of nitride formation. Material TiN ZrN Si3N4 A1N UN
Tad (K) (Azide process)
Tad (K) (Conventional)
3022 2864 2537 3010 3669, 2824
4900 4900 4300 2900 3000
of vapor phase decomposition of NH4C1, formation of SiClx, followed by the nitridation of SiClx to form Si3N4. The maximum temperature of combustion was 2123 K, which is much lower than the Tad value calculated. Since vapor phase reactions took place, the product contained loose powders, whiskers, and platelets. Lee et al. (1995) also used the "Azide" process to synthesize A1N powder. As predicted by the thermodynamic calculations, Tad for such reactions are higher than those for the formation of Si3N4. In order to counteract the fragmentation of the processed pellet, a casket of Ti and C powder was built around the A1 and NaN3 mix. The reaction between A1 and NaN3 was actually ignited by the combustion reaction between Ti and C. The conversion was reasonably high. Similar to the Si3N4 reaction, gas phase reactions also took part in A1N formation. 11.6. SHS REACTIONS IN SYNTHESIZING T i 3 S i C 2
Considerable attention has recently been focused on Ti3SiC2 due to its mechanical behavior (Goto and Hirai 1987, Pampuch et al. 1989, Lis et al. 1993, Racault et al. 1994, Tong et al. 1995, Arunajatesam and Carim 1995, Radhakrishnan et al. 1996, Barsoum and E1-Raghby 1996). A complete understanding of the mechanical properties can be obtained only if bulk samples of Ti3SiC2 are synthesized. Such samples have been obtained by sintering or hot-pressing Ti3SiC2 powders obtained through combustion synthesis. Combustion synthesis yields powders that contain up to 10-30 vol% TiC impurities. Consequently, these powders must be pulverized and hot pressed to obtain bulk samples for testing. As opposed to this two-stage process, we have hot pressed Ti, Si, and C powders at 1400~ and 40 MPa for 90 min to obtain bulk samples (Radhakrishnan et al. 1998). X-ray diffraction (XRD) pattems indicate the presence of Ti3SiC2 with about 10 vol% of TiC. The hardness of Ti3SiC2 at room temperature is 4.1 GPa and decreases to 1 GPa beyond 1000~ This suggests significant yielding of the compound beyond 1000~ If this is the case, it should be possible to fabricate complex shapes by extrusion, forging, injection molding, etc. A fracture toughness of 6.15 4-0.4 MPa m 1/2 was obtained from chevron-notched specimens. The fracture surface revealed a layered profile and cleavage planes (Fig. 11.3). Despite possessing superb mechanical properties, Ti3SiC2 lacks sufficient resistance to oxidation. A composite with SiC was investigated to improve the resistance to oxida-
300
S.B. Bhaduri and S. Bhaduri
lO~m Figure 11.3. SEM micrograph of combustion synthesized Ti3SiC2.
Figure 11.4. SEM Micrograph of fracture surface of displacement reaction processed Ti3SiC2.
tion. The solid state displacement reaction between TiC and Si was utilized to fabricate Ti3SiC2/SiC composites. The composite microstructure revealed a uniform distribution of SiC in a TizSiC2 matrix (Fig. 11.4). Two morphologies of SiC were identified: blocky and needle-shaped. The composite, when subjected to oxidation in air at 1000~ registered a weight gain of the order of a few milligrams after 10 hours. A very thin oxide layer was formed. There-
301
Combustion Synthesis i: ~: ; ~ ~ ~
:.
9
Figure
:~ : ~ : ~ :!~i : ~
,
~ ,
:
~
:~
. . . . .
:,
~
; ~ ,,
I
~
~
:::,i~;:i~!~ I~
:i,',,,;,,,~
,~- ~,
~
~
11.5. A Vickers indentation on Ti3SiC2. The bar is 5 0 / z m ,
fore, our objective to improve the oxidation resistance was successful (Radhakrishnan et al. 1998).
The effect of SiC on the mechanical properties was also investigated. A hardness of 4.74 GPa was obtained. The increase in hardness can be attributed to the addition of SiC. Cracks from Vickers indentation, when examined using SEM, provided evidence of crack bridging and crack deflection (Fig. 11.5). The bend strength of the composite was 405 MPa. The fracture surface showed fracture around SiC particles as well as a layered Ti3SiC2 matrix.
11.7. C O N T R O L L E D
REACTIONS
IN THE Ti-B
BINARY
SYSTEM
Another example of the utility of thermodynamic predictions in controlling vigorous combustion reactions is in the Ti-B binary system. It is generally believed that the formation of single-phase TiB2 is highly exothermic (Frankhouser et al. 1986, Sheppard 1986, Munir 1988, Munir and Anselmi-Tamburini 1989, Yi and Moore 1990). A low melting ternary element such as Fe or Cu may be incorporated to control the reaction between Ti and B (Li et al. 1992). Addition of extra elements, however, gives rise to other undesirable phases, rendering the interpretation of mechanisms and the evaluation of properties difficult. Therefore, control of the reaction can be sought by simply shifting to a composition away from the TiB2 stoichiometry (Peng et al. 1996, Peng and Bhaduri 1997). In the following paragraphs, the results of this work are presented. First, a thermodynamic study was carried out to investigate how the exothermic nature changed. Pertinent thermodynamic properties of TiB and TiB2 are shown in Table 11.7. Figure 11.6 shows the variation of Tad with respect to the free Ti acting as a diluent. Phase changes affect the temperature of the system and consequently the reaction rate. The reaction (1 + x)Ti + 2B --+ TiB2 + xTi, for instance, absorbs the latent heat of fusion of 100.4 kJ/mol, while maintaining Tad close to the melting temperature. This trend con-
302
S.B. Bhaduri and S. Bhaduri
Table 11.7 Thermochemical properties of the compounds in the Ti-B system.
Tad
Melting point Tm (~
B
C
(oC)
Product
Ti
B
Comments
56.379
25.867
- 17.464
2919
2913
1670
2092
Tad = Tm
54.066
-0.033
-21.631
3115
2200
1670
2092
Tad > Tm
AH 0 Product
Heat capacity
kJ/mole
A
TiB 2
323.8
TiB
160.2
3100 ~El
2800
-
25~
-
22~
-
1900
-
\+
C
A
O v "O m
+ +
I.--
+ ~+ _+_1.
(1 + x ) T i +
_../X._2.
(1 + x ) ' l i + 2 B = T i B 2
16~
- --O--
13~
.......
B=TiB+xTi
~+
+xTi
~
3. (2 + x)'l'i + 3 . 5 B = 1 . 5 T i B 2 + 0 . 5 T i B + x r i
+
4. (2 + x)'ri + 2 . 5 B = 0 . 5 " f i B 2 + 1.5"fiB + x'fi
0
Figure 11.6.
I .......
0.2
I .......
0.4
I .......
0.6
I .......
0.8
Free Ti (mole)
I .......
1
I .......
1.2
!,,,
1.4
Thermodynamic predictions of how Tad varies with departure from stoichiometry.
tinues up to an addition of 0.65 mole of Ti, as depicted in Fig. 11.6. Tad in the reaction (1 + x)Ti + B ~ TiB + xTi drops drastically with the addition of Ti. Second, the combustion behavior was closely monitored in the chosen compositions. Elemental titanium ( - 3 2 5 mesh, >99.4 pct pure and - 1 0 0 mesh, >99.5 pct pure) and boron ( - 3 2 5 mesh, >99 pct pure) powders were used as starting materials. The three compositions chosen for reaction contained 18, 15, and 12% B, the rest being Ti. The compositions investigated were either stoichiometric or off-stoichiometric composition of TiB. Six grams of each composition were measured and then mechanically pressed at pressures of 135,340, and 500 MPa in a uniaxial press with a die of 13 mm in diameter to form cylindrical pellets, giving two groups of pellets with an identical density of 50-70% of theoretical. Figure 11.7 shows a series of photographs of the combustion front of the above compositions ignited under the same conditions. The white zone is the reaction product, while the black zone is the reactant mixture. The specimens of proper compositions in the Ti-B system retained their original form of cylindrical shape and dimension in cross-section
Combustion Synthesis
{a }
s
(
s
~s
303
!~
Figure 11.7. Propagation of the controlled combustion wave fronts in compositions" (a) Ti-18 wt.% B (b) Ti-15 wt.% B, and (c) Ti-12 wt.% B.
Table 11.8 Reaction rate versus composition and size of Ti powder.
Compositions
Density (% of theoretical)
Ti-12 wt.% B Ti-15 wt.% B Ti- 19 wt.% B
50-60 50-60 50-60
Velocity (mm/s) degassed at 550~ for 10 hrs -325 mesh Ti -100 mesh Ti 13 10 2.5
2 1.4 1.2
while shrunk in the longitudinal direction by about 10%. It is also noted that the combustion front was a planar surface propagating through the compact. The steady front is attributed to an optimum combination of the reaction parameters (such as composition, green density, and particle size). For compositions with Ti content of 88 wt.% the combustion of reactant mixture did not take place, whereas combustion reactions in compositions with Ti of 82 and 85 wt.% were completed. The effects of green density and composition on reaction rate are summarized in Table 11.8 based on visual observations. For compositions of 15-18 wt.% B with green density less than 70% of theoretical, the combustion of the reactants was smooth. Compositions with higher Ti content (e.g., 88 wt.%) propagated half way down the sample and then were self-quenched. When the green density of the compact reached up to about 70% of theoretical, none of the compacts of the three compositions were ignited. It is believed that the higher green density of the compact resulted in higher thermal conductivity, resulting in more heat loss from the combustion front to its surroundings.
304
S.B. Bhaduri and S. Bhaduri 14.
J.
- o - - fine Ti
12-"
.-e--- c o a r s e Ti
A
E E 0 tU
10-"
8
="
6
="
4
= =
2-
-
~
O-
9
80
82
9
9
9
9
9
84 wt%
Figure 11.8.
41
- . . . m _ . . O . ~ . ~ . " 9
|
.
86
.
|
.
L
,
88
.
.
.
90
Ti
Effect of titanium content with different particle size on propagation rate of combustion reaction.
The rate of combustion wave propagation was measured normal to the direction of the combustion based on the photograph of the reaction. The photographs were taken with a video camera at a film speed of about 30 frames per second. Both the Ti particle size and its composition have a large effect on the reaction rate, as shown in Table 11.8. The rate is seen to be proportional to particle size and the content of Ti, as depicted in Fig. 11.8. This implies that combustion in the Ti-B system may be optimized by selecting the proper size of powders and compositions for controlling the reaction. Under otherwise identical conditions, the velocity of the propagation through the fine powder mixture is considerably faster than that of the coarser particle mixture, which is quite consistent with conventional wisdom. Most importantly, the combustion velocity ranges are much lower than a typical highly exothermic reaction. Yet, the propagation of the combustion front was still smooth, indicating a controlled combustion. Figure 11.9 contains XRD patterns of SHS-produced samples, showing three phases of TiB, TiB2 and Ti. This figure clearly evidences the effectiveness of the control of reactions attempted in this work. In order to appreciate this, close attention must be paid to the TiB and Ti peaks located at the 20 angles between 38 ~ and 40 ~ From the top to the bottom patterns, the ratio of the peak intensity of TiB over Ti is reduced. Since the bottom pattern is for the TiB stoichiometric composition, free Ti retained in the stoichiometric composition seems to be less. It is also observed that the formation of TiB2 is very insignificant in off-stoichiometric compositions as compared to the stoichiometric composition.
11.8. A U T O - I G N I T I O N S Y N T H E S I S O F N A N O C R Y S T A L L I N E O X I D E S
In the auto-ignition process, atomistic mixing is achieved, resulting in homogeneous complex chemical products. The process is similar to the "Pechini" process (Pechini 1963) or "Glycine-Nitrate" process (Pederson et al. 1992) or has been termed as "Combustion Synthesis" (Aksay et al. 1991). While the term "Combustion Synthesis" tends to get
Combustion Synthesis 9T i B 2 o TiB
305
T rl
eTi
o
(c) 77- 12v~%B
o II m
b)_ Ti- 15 Ht%B
(9
=_
_
_
_
(~) Ti-18Ht%B .
25
.
.
.
I
30
.
.
.
.
.
.
35
.
.
.
.
40
.
.
.
.
45
.
.
.
.
.
50
"
;
- - "
55
-
-
-
i
60
-
-
-
-
65
Figure 11.9. X-ray diffraction patterns of the CS-synthesized products of compositions. (a) Ti-18 wt.% B, (b) Ti-15 wt.% B, (c) Ti-12 wt.% B.
confused with SHS, the "Glycine-Nitrate" and the "Pechini" processes are very special cases of such auto-ignition processes. Yet another variation is called the "Spray Pyrolysis" process (Kingsley and Patil 1988). The auto-ignition process had its origin in rocket propellant technology and utilizes oxidizers (e.g., metal salts) and fuels (typically organic compounds). When properly controlled, high temperatures can be generated by the exothermic properties of reduction-oxidation (redox) reactions between decomposition products of the oxidizer and the fuel, which drives the reaction. The process is controlled by such parameters as heating rate, stoichiometry, choice of fuel, etc. The burning mixtures generate a large amount of gas in a very short period of time leading to very fine particulates of friable agglomerates. The following are important features of the auto-ignition process: 9 It is a versatile process leading to the synthesis of solid solution composites as well as complex compound oxide phases in homogeneous form. 9 It can synthesize ultra-fine powders. 9 It uses cheap raw materials. 9 Rapid reaction leads to nucleation of crystallites without much growth. 9 It is a scaleable and a high production rate process. 9 The products are of high purity and are loosely agglomerated. 9 The products are easily sintered. While several ceramics have been reportedly synthesized by the auto-ignition technique, the effort in this area of research is not well coordinated. A great deal of work is being carried out by chemists with little appreciation for the materials themselves, while materials scientists have been performing this type of work without much understanding of fuels and their combustion behavior. To the best of our knowledge, there has been
306
S.B. B h a d u r i a n d S. B h a d u r i
no review on this subject. In the following, some examples of materials synthesized are presented. Examples of synthesis of simple single-phase oxides include A1203 (Aksay et al. 1991) and CeO2 (Verma et al. 1990). Examples of solid solutions include yttria doped zirconia and ceria doped zirconia (Aksay et al. 1991, Dhas and Patil 1994). However, the most important use of this technique, so far, is in synthesizing complex oxides. Because of the atomistic level mixing of raw materials, the final products are very homogeneous. Examples of complex oxides produced by this process include YAG (Aksay et al. 1991), lanthanum aluminate (Aksay et al. 1991), mullite (Hong et al. 1994), cordierite (Gopichandran and Patil 1993), 1-2-3 superconductors (Rambabu 1990), lanthanum chromite (Chick et al. 1990), barium titanate (Zhong and Gallagher 1995), and various ferrites (Ravindranathan and Patil 1987). It has been realized that heating and cooling rates are fast in such reactions. Therefore, they can be utilized in synthesizing nanocrystalline materials. A group at Alfred University has made use of this process for synthesizing nanocrystalline YzO3-doped ZrO2, and pure ZrO2 (Venkatachari et al. 1995, Huang et al. 1995). We have used the auto-ignition process to directly synthesize nanocrystalline oralumina (Bhaduri et al. 1996a). It is believed that the high temperatures generated for a short time crystallized the c~-phase. Since the cool down period was fast, the crystallites did not have the time to grow. Both XRD and transmission electron microscopy (TEM) studies confirmed the crystallite sizes to be less than 100 nm. This study was extended to the AlzO3-ZrO2 binary system to synthesize nanocomposites (Bhaduri et al. 1994, 1995, 1997b). It is well known that grain growth is considerably reduced whenever ZrO2 is added as a second phase to A1203. Nanocomposite powders containing A1203-14 wt.%ZrO2 were synthesized using the auto-ignition process. Both XRD and TEM studies showed that A1203 was in the c~ form and that ZrO2 has the tetragonal structure. The average crystallite size for A1203 was approximately 40 nm and that of ZrO2 was 13 nm. Examples of solid solution were CaO and CeO2 doped nanocrystalline ZrO2 powders (Zhou et al. 1996, 1997). The compositions contained 2 to 10 mol% of CaO and 4 to 20 mol% of CeO2 in ZrO2. The powders were characterized by XRD, Raman spectroscopy, and TEM. The crystallite sizes were about 5 nm in most of the compositions. The results indicated that the as-synthesized powders were in the tetragonal phase and not in the cubic phase. There is also evidence of extension of solid solubility because of the non-equilibrium nature of the process. Nanocrystalline spinel (or spinel-alumina nanocomposite) powders were also synthesized (Bhaduri et al. 1997a). Al-nitrate and Mg-nitrate were used as the starting materials and urea was used as the fuel. The crystallite sizes and lattice parameters were calculated from the XRD results. Lattice parameters of non-stoichiometric spinel solid solutions satisfied the Vegard's law. Spinel grains were about 13-20 nm range and or-alumina 30-40 nm, as determined by XRD and TEM.
Combustion Synthesis
307
11.9. N O N - E Q U I L I B R I U M E F F E C T S
As mentioned in the Introduction, the rapid heating and cooling effects involved in combustion synthesis can produce non-equilibrium phases. The non-equilibrium phases produced by this technique are neither as extensive nor as large in number as those obtained by rapid solidification or mechanical alloying techniques. Another important difference between combustion synthesis and other non-equilibrium processing techniques is that while combustion synthesis is used more for ceramic systems, rapid solidification and mechanical alloying are used more for metallic systems. Therefore, a direct comparison of the non-equilibrium effects produced by these different techniques is not fair. However, by combustion synthesis also supersaturated solid solutions and glass formation have been reported as also the formation of intermetallic phases away from the equilibrium stoichiometry. Kunrath (1998) reported that supersaturated solid solutions could be obtained in the TiC-Cr3C2 system. While the room temperature solid solubility of Cr3C2 in TiC under equilibrium conditions is very small ( 10 nm (most materials produced to-date are in this size range) and therefore they may not exert much influence on the properties of nanocrystalline materials.
12.5. S T A B I L I T Y
Knowledge of the thermal stability of nanocrystalline materials is important for both technological and scientific reasons. From a technological point of view, the thermal stability is important for consolidation of nanocrystalline powders without coarsening the microstructure. That is, many methods for synthesis of nanocrystalline materials result in powder products. These powders then require consolidation by techniques such as hot forging to obtain useful bulk material. The goal of such powder consolidation is to attain essentially 100% theoretical density while preventing or minimizing grain growth of the nanometer-sized grains. Another processing method that requires information on grain growth is the devitrification of amorphous precursors (Lu 1996). In order to design an appropriate heating schedule to provide the desired nanometer grain size, knowledge of the nucleation rate and grain growth rate are both needed. The basic scientific question regarding grain growth of nanocrystalline materials, as for all phenomena for nanocrystalline materials, is whether the behavior involves "new physics" or is simply the expected grain-size-dependent behavior extrapolated to nanometer grain sizes.
1 2 . 5 . 1 G r a i n g r o w t h in n a n o c r y s t a l l i n e m a t e r i a l s
Grain growth occurs in polycrystalline materials to decrease the system energy by decreasing the total grain boundary energy. Under ideal conditions grain growth kinetics should obey parabolic kinetics (Atkinson 1988, Vandermeer and Hu 1994), namely D 2 -
D 2 -- kt
(12.2)
where D is the length scale, e.g., mean grain diameter, after annealing for time t, k is a temperature-dependent rate parameter and Do is the length scale or the grain size at t = 0. However, most experimental data for isothermal grain growth do not obey this ideal behavior and experimenters often analyze grain growth data with the expression:
D1/n
1/n
- Do
-- k t
(12.3)
where n is an empirical constant typically 2). By applying an Arrhenius-type equation of the form
0exp{
324
C. Suryanarayana and C. C. Koch
to the rate parameter k one gets access to the activation energy Q, which is frequently used to deduce the grain growth mechanism. Another method of obtaining Q for nanocrystalline materials is the so-called Kissinger method in which differential scanning calorimetry (DSC) is used to analyze isochronal grain growth at different heating rates. We shall describe the isothermal grain growth kinetics data by the use of n and Q. However, first we will review several aspects of thermal stability which have been observed in nanocrystalline materials.
12.5.2 Grain growth at ambient temperature Significant grain growth, that is doubling of the grain size in 24h, has been observed in a number of pure, relatively low melting elements such as Sn, Pb, A1, and Mg (Birringer 1989). Because of the large amount of grain boundary enthalpy stored in nanocrystalline materials a high driving force for grain growth is expected. GUnther et al. (1992) and Kumpmann et al. (1993) studied grain growth in pure Cu, Ag, Pd and Ag alloys and found grain growth occurring at much lower temperatures than those observed for recrystallization of the elements after heavy cold deformation. In fact grain growth in Cu and Pd was observed even at room temperature. In some of these cases grain growth was observed to be abnormal. Gertsman and Birringer (1994) observed abnormal grain growth in nanocrystalline Cu after holding the samples for times of more than a month at room temperature. The abnormally coarse grains exhibited a range of sizes, most were 2 #m. The nanometer grains surrounding the coarse grains were 10-50 nm. The general explanation given for abnormal grain growth is an inhomogeneous structure for the as-processed samples such that grain growth inhibitors such as pores, impurities, or even grain boundary structure are not evenly distributed and rapid grain growth can occur where such inhibitors are absent due to the large driving force and high grain-boundary mobility.
12.5.3 Examples of grain growth inhibition Significant stabilization of nanometer grain structures has been observed in many multicomponent materials. It has been observed in several systems that on heating a metastable nanocrystalline solid solution alloy, limited grain growth occurs while solute atoms also segregate to the grain boundaries. The grain boundary segregation then may lower the specific grain boundary energy sufficiently to reduce the driving force and limit further grain growth. This effect was illustrated clearly in Pdl-xZrx alloys (Krill et al. 1995) where x = 0.1 is an equilibrium solid solution and x = 0.2 is a metastable supersaturated solid solution. The 10 at.% Zr nanocrystalline alloy showed significant grain growth with an onset temperature of 325~ while the 20 at.% Zr nanocrystalline alloy had little or no grain growth at temperatures up to 500~ While it is not possible to separate possible kinetic effects, such as solute drag, from this thermodynamic stabilization, the much slower grain growth in the 20 at.% Zr supersaturated alloy strongly suggests that the thermodynamic stabilization is dominant here. Solute segregation to nanocrystalline grain boundaries being responsible for grain size stabilization in supersaturated metastable nanocrystalline
Nanostructured Materials
~
-~
IE,_ r
.
RT(Pd .
.
.
.
.
.
.
.
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.
t
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*-" o
.
325
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............................... 5 ....................i. . . . . . . . . . . . . . . . . . . . . . . . . []
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........................... ........................... j .
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.
.
.
.
.
.
.
.
.
.
.
.
.
.
.
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..........................I-(-Fe;Mn)~-S.: i~~----~ 0
....
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0.2
.
.
.
.
.
.
.
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.
.
.
.
.
.
.
i
...........................j........................... i.....................
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.
t ....
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(. ..... .F~St...i
...........................
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i ....
I
I
0.3
0.4
JI . . . .
0.5
Normalized Annealing T e m p e r a t u r e
t ....
I
I
0.6
0.7
(T/TM)
Figure 12.4. Time exponent, n, for isothermal grain growth of various nanocrystalline materials as a function of the normalized annealing temperature (after Malow and Koch 1996).
solid solutions has also been observed in Nb-Cu (Abe et al. 1992) and Fe-Cu (Eckert et
al. 1993) alloys. A number of factors can influence grain boundary mobility in nanocrystalline alloys. These include solute drag, pore drag, second phase (Zener) drag, and chemical ordering.
12.5.4 Isothermal grain growth kinetics The isothermal grain growth kinetics will be considered using Eqs. (12.3) and (12.4) to define the grain growth exponent n (_ 775 K in nanocrystalline Fe was found to be 248 kJ/mol which is comparable to that for grain growth in zone refined pure Fe with conventional grain sizes (> 10/zm) (Vamdermeer and Hu 1994) and with Q1 for Fe. But at T < 775 K, the activation energy was 125 kJ/mol which is lower than both Qgb and Q1 for Fe. This suggests that grain growth in nanocrystalline Fe at T > 775 K is similar to that in conventional grain size pure Fe, but grain growth in nanocrystalline Fe at T < 775 K is due to another mechanism (Malow and Koch 1997). In conclusion it can be said that in order to learn about the mechanisms that govern the thermal stability/grain growth in nanocrystalline materials it is important to keep factors in mind that might interfere with grain growth. Specifically it would be desirable to avoid porosity, contamination and multiple phases (or precipitation) in future investigations. Furthermore it would be of advantage to compare the resultant activation energies to diffusion data for nanocrystalline materials, if available, as it has been done by H6fler and Averback (1990).
12.6. PARTICULATE CONSOLIDATION Widespread application of nanocrystalline materials 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 nanocrystalline materials. However, because of the small size of the powder particles (typically a few microns, even though the grain size is only a few nanometers), some special precautions need to be taken to minimize their activity and the high level of interparticle friction. Conventional consolidation of powders to full density through processes such as hot extrusion and hot isostatic pressing normally requires use of high pressures and high temperatures for extended periods of time. Unfortunately, however, this results in significant coarsening of the nanometer-sized grains and consequently the benefits of nanostructure processing are lost. Therefore, novel and innovative methods of consolidating nanocrystalline powders are required. Consolidation of nanocrystalline powders has been achieved by electro-discharge compaction (Okazaki 1993), plasma-activated sintering (Groza 1993, 1994), shock (explosive) consolidation (Korth and Williamson 1995, Suryanarayana et al. 1997), hotisostatic pressing (Suryanarayana et al. 1997, Haji-Mahmood and Chumbley 1996, He and Ma 1996a), Ceracon processing (Suryanarayana et al. 1997), hydrostatic extrusion (Liang et al. 1996), strained powder-rolling (Liang et al. 1996), and sinter-forging (He
Nanostructured Materials
327
and Ma 1996b). Optimization of the consolidation parameters is important because retention of the nanostructures requires use of low consolidation temperatures while achieving full (theoretical) density requires use of high temperatures. However, it should be noted that because of the increased diffusivity in nanocrystalline materials, sintering (densification) takes place at a temperature much lower than in coarse-grained materials. This is likely to reduce the grain growth. (See also the chapter on "Powder Consolidation" by J.R. Groza in this book). Full consolidation of nanocrystalline materials has also been achieved by modifying/improving some of the other existing processes. Although it is typically difficult to obtain full density while retaining the nano-scale microstructure, the high driving force and enhanced kinetics due to large curvature effects facilitate full densification of nanocrystalline materials at temperatures consistently below those of coarse-grained materials of the same composition by a few hundred degrees. Retention of fine grain sizes and elimination of sintering aids (especially for ceramic materials) are the specific advantages of the lower processing temperatures. 12.7. P R O P E R T I E S
Because of the very fine grain sizes, nanocrystalline materials exhibit a variety of properties that are different and often considerably improved in comparison with those of conventional coarse-grained polycrystalline materials. These include increased strength~ardness, enhanced diffusivity, reduced density, higher electrical resistivity, increased specific heat, higher thermal expansion coefficient, lower thermal conductivity, and superior soft magnetic properties. 12.7.1 Diffusion and sinterability
Since nanocrystalline materials contain a very large fraction of atoms at the grain boundaries, the numerous interfaces provide a high density of short-circuit diffusion paths. Consequently, they are expected to exhibit an enhanced diffusivity in comparison to single crystals or conventional coarse-grained polycrystalline materials with the same chemical composition (Horvath 1989, H6fler et al. 1993). This enhanced diffusivity can have a significant effect on mechanical properties such as creep and superplasticity, ability to efficiently dope nanocrystalline materials with impurities at relatively low temperatures, and synthesis of alloy phases in immiscible metals and at temperatures much lower than those usually required in other systems. The measured diffusivities in nanocrystalline copper are about 14 to 20 orders of magnitude higher than lattice diffusion and about 2 to 4 orders of magnitude larger than grain boundary diffusion. For example, the measured diffusivity at room temperature is 2.6 x 10 -20 m2/s for 8 nm-grain-sized copper samples compared to 4.8 x 10 -24 for grain boundary diffusion and 4 x 10 -40 for lattice diffusion (Schumacher et al. 1989). The increased diffusivity (and consequently the reactivity) leads to increased solid solubility limits, formation of intermetallic phases (at temperatures much lower than those required for coarse-grained materials and sometimes new phases) and increased sinterability of nanocrystalline powders.
328
C. Suryanarayana and C. C. Koch
Solid solubility limits are usually enhanced when the material is in the nanocrystalline state. In extreme cases, solid solutions can also form in alloy systems that exhibit miscibility gaps both in the liquid and solid states. Typical examples of this phenomenon are the Ag-Fe, Ti-Mg, and Cu-Fe systems (Suryanarayana 1995). While the room temperature solid solubility of Mg in Ti under equilibrium conditions is
3
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0
;
.
01.2
016 Displacement, d/ mm 0|.4
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~="Pmax )/~
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300
200
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-6
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-10
0
0.2
0.4
0.6
Displacement, d
0.8 /mm
Figure 14.15. Charpy impact fracture load displacement curves of the cast amorphous Zr55 A1!0Cu30 Ni5 and Pd40 Cu30 Ni !0 P20 specimens.
pattern nor featureless cleavage-like pattern typical for embrittled amorphous alloys was seen over the whole fracture surface. The feature of the fracture surface indicates that the cast bulk amorphous alloy has good ductility leading to severe deformation even in the impact fracture mode. Besides, secondary failure was observed and the change in the fracture surface appearance from equiaxed dimple-like pattern to vein pattern was recognized near the failure. The secondary failure appears to occur in close relation to casting defects. That is, the casting defect corresponds to the interface between the amorphous regions formed at two different stages. The amorphous phase region near the interface is subjected to subsequent heating. The heating decreases the ductility as compared with the as-cast amorphous phase without subsequent heating, leading to the generation of secondary failures. Although an interface between the formerly and latterly solidified amorphous alloys cannot be completely avoided in the casting process, a decrease in the number of the interfaces is expected to further raise the impact fracture energy. Based on the above-described data, in comparison with the corresponding crystalline alloys, bulk amorphous alloys have (1) higher of and lower E, (2) much larger elastic strain of about 2%, exceeding a yield limit of 0.2% for crystalline alloys, (3) much higher elastic energy up to yielding, (4) absence of distinct plastic elongation at room temperature due to inhomogeneous deformation mode, and (5) relatively high impact fracture energy of about 70 kJ/m 2. When the stress-strain curves of Zr-based bulk amorphous alloys are compared with those for Ti-based crystalline alloys, the elastic energy required up to yielding for the bulk amorphous alloy is larger by about 25 times. The relations between E and crf or Hv for the bulk amorphous alloys are shown in Fig. 14.16, where the data of conventional crystalline alloys are also plotted for comparison (Inoue 1997). When the E values of the bulk amorphous alloys are compared with those of the crystalline alloys with the same o'f, the moduli are smaller by about 60% than those for the crystalline alloys. The relation between E and O'f or By can be classified into two different groups. The distinct classification indicates that the features of mechanical properties of bulk amorphous alloys are different from those for the crystalline alloys. The distinct
399
Bulk Amorphous Alloys 5000
,
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/
4500
4000 _
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/
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Relation between Young's modulus (E) and tensile strength ( o f ) or Vickers hardness (He) for bulk amorphous and crystalline alloys.
14.16.
difference in the fundamental mechanical properties is very important for applications. Some engineering applications by the use of the difference are in progress because of the unique characteristics that cannot be obtained for crystalline alloys in each application field. 14.8. V I S C O E L A S T I C I T Y
Dynamic mechanical properties of storage, E ~, and loss, E", moduli were measured with a dynamic mechanical analyzer by a forced oscillation method. The complex modulus, E*, is given as the ratio of an oscillatory stress to the oscillatory strain. The real and imaginary parts of E* are defined as E ' and E" and described by Eq. (14.8) (Nielsen 1962) E* -- E' + i E " -- Eco2r2/(1 + co2z'2) -Jr- iEcor/(1 + co2.g2)
(14.8)
Here, E is the elastic modulus, co is the measurement frequency and z is the relaxation time. At coz = 1, E' and E" show an inflection point of a reduction curve and a loss peak, respectively. The temperature-dependent measurement was carried out every less than two degrees at heating rates from 0.0083 to 0.25 K/s at frequencies of 0.628, 6.28 and 62.8 rad/s. Figure 14.17 shows the temperature dependence of E ~ and E" of the La55AlzsNi2o amorphous alloy with heating rate (or) at 6.28 rad/s (Okumura et al. 1993). At 0.017 K/s, E ' starts to reduce at 466 K due to the lower relaxation system. The effect of the higher temperature relaxation system appears in U ( T ) at 481 K as a plateau which indicates the transition of the main relaxation system in the measurement from the lower temperature relaxation to the higher temperature relaxation, leading to a second rapid reduction
400
A. Inoue
of E t starting at 485 K with increasing temperature. From 503 K, E t increases due to crystallization. At a higher ot of 0.083 K/s, E ~ follows the same E t(T) curve of the first reduction at 0.017 K/s but E ~ shows more reduction superimposed on the first reduction curve at 0.017 K/s until 487 K. After the transition of the main relaxation system, E ~reduces again due to the higher temperature relaxation and rejoins the E t (T) curve at 0.017 K/s above 493 K. The appearance of the second relaxation system does not depend on the stress level but temperature suggesting the ot dependence, that is, shifting to higher temperature with increasing or. E ~shows more second reduction to a lower value as compared with the U(T) curve at 0.017 K/s because the Tx shifts to a higher temperature of 517 K at 0.083 K/s. With further increasing c~ to 0.13 K/s, the change in the E~(T) curve has the same tendencies. The first and the second reduction E ~(T) curves show the same and the superimposed curves at lower ot but the appearance shifts to the higher temperature of 493 K. The Tx also shifts to 521 K. In the temperature dependence of E" shown in Fig. 14.17(b), two relaxation peaks for the lower and the higher temperature relaxation and the crystallization peak are clearly observed at 473,490 and 511 K, respectively, at 0.017 K/s. The valley temperature of 479 K between the two relaxation peaks corresponds to the temperature for the appearance of the effect of the second relaxation in the E ~(T) curve as shown in Fig. 14.17(a). The peak temperatures, Tpl and Tp2, for the first and the second relaxation, respectively, show contrary results by the increase of or. Tpl slightly shifts to lower temperatures from 473 to 471 K and the height of the peak slightly decreases while Tp2 clearly shifts to higher temperatures from 490 to 501 K and the height rapidly decreases with increasing c~ from 0.017 to 0.13 K/s. At the higher c~, the appearance of the effect of the second relaxation shifts to higher temperature marking the shift of Tp2, and the overlap of the two relaxation peaks becomes less. Thus, the first peak distinctly appears with a slight shift to lower temperature and a slight decrease in height. Because the initial E" value for the second relaxation is lower, caused by the progress of the first peak, the height of the second peak decreases rapidly with increasing or. The E t~(T) curve merges with that of a lower ot at temperatures above Tp2 showing the same tendency as seen in the E ~(T) curve. The onset temperatures, To l and To2, for the first and the second reduction of U , are defined as the points of intersection of two tangent lines of the Et(T) curve. The ot dependence of To1 and To2 at three frequencies of 62.8, 6.28 and 0.628 rad/s is shown in Fig. 14.18(a). No different tendency is observed in the results of different frequencies. With increasing or, Tol negligibly increases to a saturated value while To2 distinctly increases. For example, in the case of 0.628 rad/s, Tol increases from 466 to 467 K and To2 increases from 481 to 493 K with increasing ot from 0.0083 to 0.083 K/s. Because the appearance of the second relaxation in E t (T) has the oe dependence, To2 shows the distinct ot dependence. The second relaxation starts to appear before the first relaxation shows a steep reduction in the E~(T) curve at lower ot (see Fig. 14.17(a)). Thus, the tangent line of the reduction side for To1 becomes sluggish, leading to the slight shift to a lower temperature for To1. At sufficiently higher ot as the first relaxation shows steep reduction, To1 shows very little ot dependence and saturated value. The shifting behavior of Tpl and Tp2 against ot at three frequencies is summarized in Fig. 14.18(b). As the two relaxation peaks become more distinct at the lower frequency
Bulk Amorphous Alloys
401
102 I
a~AI2~Ni..,,
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9
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i
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T/K
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402
A. Inoue
Las.~Ai2~Ni2~,
510
0 62.8rad/s
zx 6.~
To2
El 0.628
5OO
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460
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Heating rate, a / K s-'
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5OO
490
480
Tpl
470
460
~ 10-2
i 10-~
_
,
10o
Heating rate, a / K s" F i g u r e 14.18. Heating rate dependence of Tol and To2 at three frequencies of 0.628, 6.28 and 62.8 rad/s. (b) Heating rate dependence of Tpl and Tp2 at three frequencies.
Bulk A m o r p h o u s Alloys
403
(Okumura et al. 1992), Tpl and Tp2 at the lower frequency shift more drastically than those at the higher frequency. The endothermic reaction for the glass transition in the DSC curve also shifts to higher temperature with increasing or. To correlate with the change in Ef'(T) with or, the apparent activation energies, Q, are evaluated by the Kissinger method (Kissinger 1957). The value of Q obtained from Tp2 at three frequencies varies slightly from 300 to 340 kJ/mol which agrees with the value of 340 kJ/mol from the DSC measurements. It should be mentioned that Q of about 340 kJ/mol represents the value not only for the relaxation but also the value including the effects of the phase separation and relaxation. The real activation energy for the higher temperature relaxation has been obtained as 550 kJ/mol (Okumura et al. 1992a). From this result, it is clear that the shift of the DSC curve is mainly affected by the higher temperature relaxation. Making a comparison of the ratio of E" to E', tan ~ (= E " / E ' ) , between the two relaxations, tan 6 for the higher temperature relaxation exceeds 10 ~ which is one order of magnitude higher than that for the lower temperature relaxation (Okumura et al. 1991). It suggests that the higher temperature relaxation occurs more drastically than the lower temperature relaxation. The appearance of the ot dependence of Tg is mainly caused by the phase separation as evidenced by two E ~ relaxation peaks. In order to examine the ot dependence of the normal glass transition without the effect of the phase separation, poly-methyl methacrylate, PMMA (Iwayanagi and Hideshima 1953) is used. As amorphous alloys with one relaxation peak, to our knowledge, do not show stable supercooled liquid region, the crystallization peak of E ~ overlaps with the relaxation peak. The ot dependence is hardly observed in mechanical measurements considering the margin of errors. Almost all amorphous alloys with a wide supercooled liquid region exceeding 40 K in the DSC curve at 0.67 K/s, such as Zr60Al15Ni25, Mg65Cu25Y10 and Pd48Ni32P20 amorphous alloys, have two relaxation systems suggesting the phase separation in the glass transition region (Okumura et al. 1991/92). These alloys show distinct ot dependence of Tg like the La55AlzsNi20 alloy. On the other hand, the amorphous alloys with one relaxation system and smaller supercooled liquid region, such as Zrv0Cu30, La55A145 and Pd80Si20, show little ot dependence of Tg. A similar two-stage change in the dynamic mechanical properties has been noted for the Zr60AllsNi25 amorphous alloy (Okumura et al. 1992b). 14.9. S O F T M A G N E T I C P R O P E R T I E S
Based on the three empirical rules for the achievement of high GFA, ferromagnetism at room temperature has been found in Fe-(A1, Ga)-(R C, B, Si) (Inoue and Gook 1995, Inoue et al. 1995b, Inoue et al. 1996b), Co-Cr-(A1, Ga)-(P, B, Si) (Inoue and Katsuya 1996), Fe-(Co, Ni)-(Zr, Nb, Ta)-B (Inoue et al. 1997e, Inoue et al. 1998b, Inoue et al. in press) and Co-Fe-(Zr, Nb)-B (Inoue et al. 1998b) systems. This section describes the soft magnetic properties of the Fe- and Co-based bulk amorphous alloys.
14.9.1 Formation and soft magnetic properties of bulk amorphous alloys The Fe-based amorphous alloys with large A Tx (>60 K) are expected to have a high GFA which enables the production of bulk amorphous alloys with diameters larger than
404
A. I n o u e
1.5
I
as-Q.
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,,,,,,,,
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.
_
_ FeTaAI5Ga2P11CsB4
(lmm ~)
............. Fe72AIsGa2P10CeB4Si1 (2mrn~)
0.5 o.o
,,,,,,m
,i,,,,i
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-1.5
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-16oo
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....
o
Applied Field, H(A/m)
I
800
,
16oo
Figure 14.19. Hysteresis B-H loops of the cast amorphous Fe73AI5Ga2PllC5B4 and Fe72A15Ga2PIoC6B4Si 1 cylinders.
1 mm by the casting processes. The cast Fe72A15Ga2PllC6B4 (Inoue et al. 1995b) and FevzA15GazP10C6B4Sil (Inoue et al. 1996b, Inoue et al. 1997e) amorphous cylinders with diameters of 1 to 2 mm had smooth surface and metallic luster. Besides, good castability of these Fe-based alloys is indicated by noting that the outer shape of the cast samples mirrors the inner cavity of the copper mold. Even after an appropriate etching treatment, no appreciable contrast corresponding to a crystalline phase is seen over the whole transverse cross section of the cast Fe72A15Ga2P10C6B4Si] cylinders with diameters of 1 and 2 mm. In addition, neither cavities nor shrinkage holes are observed, indicating that the amorphous alloy has a good castability owing to the lack of the discontinuous change in the specific volume-temperature relation which is ordinarily recognized for crystalline alloys obtained from the liquid. The cast FevzA15Ga2P]]C6B4 amorphous cylinders with diameters of 0.5 to 1.5 mm (Inoue et al. 1995b) show the sequential changes of the glass transition, supercooled liquid and then single-stage exothermic peak which agree with those of the corresponding melt-spun amorphous ribbon. No distinct changes in the thermal stability and crystallization mode of the supercooled liquid are seen between the cast amorphous cylinders and the melt-spun amorphous ribbon.
Bulk Amorphous Alloys
405
The hysteresis B - H loop of the cast Fe72A15Ga2P11C6B4 amorphous cylinder with a diameter of 1 mm was examined in the as-cast and annealed (723 K, 600 s) states (Inoue et al. 1995b, Inoue et al. 1996b). The Bs, He and Br/Bs of the annealed sample are 1.07 T, 5.1 A/m and 0.37, respectively, indicating that the cast amorphous cylinder has good soft magnetic properties. Besides, the ~e at 1 kHz for the annealed amorphous cylinder also shows a high value of 7000. The soft magnetic properties were further improved for the 1 at% Si-containing alloy (Inoue et al. 1996b, Inoue et al. 1997b). Figure 14.19 shows the hysteresis B - H loop of the cast FevzA15Ga2P10C6B4Si1 amorphous cylinder with a diameter of 2 mm, together with the data of the cast Fev3A15Ga2PllCsB4 amorphous cylinder with a diameter of 1 mm. The Bs, He and Br/Bs for the Si-containing cylinder are 1.14 T, 0.5 A/m and 0.38, respectively. In comparison with those (Inoue et al. 1995b) for the Fe-A1-Ga-P-C-B cylinder, Hc decreases and B~/Bs increases for the Si-containing amorphous cylinder. From the thermomagnetic data, it is confirmed that the cast Fev3A15GazP11CsB4 amorphous cylinder has a Curie temperature (Tc) of 600 K which is lower by 185 K than the Tx. The Tc value agrees with that determined from the endothermic peak on the DSC curve. Thus, the cast Fe-based amorphous cylinders with diameters up to 2 mm exhibit good soft magnetic properties of 1.1 T for Bs, 2 to 6 A/m for He and 7000 for #e at 1 kHz. When these soft magnetic properties are compared with those for the corresponding melt-spun amorphous ribbons, no distinct changes in Bs and He are seen. However, the #e value is degraded for the cast cylinder presumably because of the increase in the influence of the demagnetization resulting from the significant change in the sample morphology. Figure 14.20 shows the outer morphology of the bulk Fe61CovZrl0Mo5W2B15 cylinders with diameters of 3 and 5 mm (Inoue et al. 1997f). These samples also have smooth surface and metallic luster and no contrast of a crystalline phase is seen over the outer surface. The X-ray diffraction patterns showed a main halo peak with a wave vector Kp(= 47r sin 0/)Q around 29.6 nm -1 and no crystalline peak is observed even for the 5 mmO sample. Besides, the optical micrographs of the cross section of the two samples also revealed a featureless contrast in a sample etched with HE Considering that the bulk cylinder of 7 mm diameter consists of an amorphous phase in an outer surface region with a thickness of about 2 mm and of amorphous and crystalline phases in the inner region, the
Figure 14.20. Outermorphologyof cast amorphous Fe61Co7Zrl0Mo5W2B15 cylinders with diameters of 3 and 5 mm.
406
A. Inoue
tmax for the Fe61Co7Zrl0Mo5W2B15 alloy is determined to be about 6 mm. The tmax is 3 times larger than the largest value (2 mm for FevzA15Ga2P10C6B4Sil) (Inoue et al. 1996b, Inoue et al. 1997b) for Fe-based amorphous alloys reported to date. Figure 14.21 shows the DSC curves of the bulk amorphous Fe61Co7NivZrsNbzB15, Fe56CovNi7ZrgTazB20, Fe60CosZr8NbzMosWzB 15 and Fe60Co8Zrl0MosWzB 15 cylinders with diameters of 1 to 3 mm (Inoue et al. 19970. These amorphous alloys exhibit the sequential transition of glass transition, supecooled liquid and crystallization. The ATx is as large as 55 to 88 K and crystallization occurs by a single exothermic reaction. The crystallites were identified to consist of ot-Fe, Fe2Zr, Fe3B, MoB and W2B phases for the Fe60Co8Zrl0Mo5WzB15 sample heated to a temperature just above the exothermic peak. Thus, the crystallization is due to the simultaneous precipitation of the five crystalline phases. This crystallization mode is in agreement with that (Inoue 1995a, Inoue 1995b, Inoue 1996, Inoue 1997a, Inoue 1997b) for other bulk amorphous alloys. The largest ATx is 88 K for Fe56Co7Ni7Zr8Ta2B20, being larger than the largest values (57 to 67 K) for Fe-(A1, Ga)-(P, C, B, Si) (Inoue and Gook 1995, Inoue et al. 1997b) and nonferrous Pd- and Pt-based amorphous alloys (Chen 1980, Chen 1976, Chen 1978). The Tm was 1420 K for Fe56Co7Ni7Zr10B20 and 1416 K for Fe61Co7Zrl0Mo5W2B15 and the Tg/Tm was evaluated to be 0.60 for the former alloy and 0.63 for the latter alloy. Considering that Tg / Tm is 0.54 for Fe80P12B4Si4 (Inoue and Park 1996) and 0.57 for Fe73A15Ga2P11C4B4Sil (Inoue et al. 1997b), the present Tg / Tm values are believed to be the highest among all Fe-based amorphous alloys. Table 14.4 summarizes the tmax, Tg, A Tx, Tg/Tm, compressive fracture strength (Crc,f), Hv, Bs, He, # e at 1 kHz and )~s for the new amorphous Fe-(Co, Ni)-(Zr, Nb, Ta)-B (Inoue et al. 1997e, Inoue et al. in press) and Fe-Co-Zr-(Mo, W)-B alloys (Inoue et al. 1997f). These amorphous alloys exhibit good soft magnetic properties in an annealed (800 K, 300 s) state, i.e., high Bs of 0.74 to 0.96 T, low He of 1.1 to 3.2 A/m, high #e of 12000 to 25000 and low )~s of (10-14) • 10 -6. The He and #e are superior to those for conventional Fe-Si-B amorphous ribbons (Kikuchi et al. 1975, Zhang et al. 1991), presumably because of the lower )~s. Furthermore, the Fe60CosZrl0MosWzB 15 bulk amorphous alloy has high crc,f and Hv of 3800 MPa and 1360, respectively, which exceed largely those (crf = 1000-2000 MPa, Hv = 300-800) for high-carbon high-alloy tool steels and 25 wt%Ni maraging steel. No bulk alloys with high strength above 3000 MPa for Crc,f and 1000 for Hv have been obtained for any Fe-based alloys including amorphous and crystalline phases. Besides, no weight loss is detected after immersion for 3.6 ks at 298 K in aqua regia. Thus, the Fe-based bulk amorphous alloys possess simultaneously high GFA, high strength, high corrosion resistance and good soft magnetic properties which cannot be obtained for other amorphous and crystalline alloys. 14.10. V I S C O U S FLOW AND M I C R O - F O R M A B I L I T Y OF S U P R C O O L E D LIQUID
In addition to the low Rc and large tmax, bulk amorphous alloys have a wide supercooled liquid region before crystallization (Inoue 1995a, Inoue 1995b, Inoue 1996, Inoue 1997a, Inoue 1997b). The ATx reaches as large as 127 K (Zhang et al. 1991) for the Zr-A1-NiCu alloys and 98 K (Inoue et al. 1996a) for the Pd-Cu-Ni-P alloys. In the supercooled liquid region with such large A Tx values, the viscosity decreases significantly and large
Bulk Amorphous Alloys
9
r
9
i
'
i
....
9.....
407
=
9
i
0 . 6 7 K / s
Fe6oCo~ZrloMo5W2B15 d=3mm~ .
.
.
.
r-
.
FesoCosZraNb2MosW2B15 d--3mm+
-
,
.
.
.
.
.
.
.
.
.
F
.
r-"
._o
v
(--
E FessCo7Ni7ZrsTa2B20 d=lmm t,,.._
(!) r 0 X
Ill
t
FeslCo7Ni7ZreNb2Bls d=lmm~, .
.
.
.
.
Tg
.
I
9
600 Figure 14.21. Fe56C~
1
9
1
9
I
700 800 900 Temperature, T/K
,
1
1000
DSC curves of cast amorphous Fe61Co7Ni7Zr8Nb2B15, Fe60C~ and Fe60Co8Zr10Mo5W2B15 cylinders.
P 0 to
Table 14.4 Thermal stability, mechanical strength, and magnetic properties of the cast bulk amorphous Fe-(Al, GaHP, C, B, Si), Fe-(Co, Ni)-(Zr, Nb, T a t B , and Fe-CeZr-(Mo, W)-B alloys Thermal stability h a x
Tg
Vickers
Compressive
Hc
Pe
As
(T)
Wm)
(1 kHz)
(lop6)
-
0.96 0.75 0.85 0.74
2.0 1.1 3.2 2.6
19100 25000 12000 12000
10 13 14 14
3800
-
-
-
14
AT, (K)
Tgl
hardness
strength
(K)
Tm
HU
q c . f (MW -
FeyjC07Ni7ZrloB20 FeyjCo7Ni7ZrgNb~B20 Fe61Co7Ni7ZrgNb2B15 Fe56Co7Ni7Zr8TazB20
2 2 2 2
814 828 808 827
73 86 50 88
0.60
-
1370 1370 1340 1360
Fe60CogZr10MogW2B15
6
898
64
0.63
1360
-
Soft magnetic properties
-
B,
? 3 0 fi
m
Bulk Amorphous Alloys
409
viscous flow is obtained (Inoue and Zhang 1997). By utilizing the low viscosity and large viscous flow, the bulk amorphous alloys can be deformed to various complicated shapes while retaining the good mechanical properties (Inoue et al. 1997d, Inoue et al. 1996c, Kawamura et al. 1996, Nishiyama and Inoue to be submitted). This section intends to describe the micro-formability in the supercooled liquid region for the La-A1-Ni, Zr-A1Ni-Cu and Pd-Cu-Ni-P bulk amorphous alloys and to introduce some practical examples of micro-formed bulk amorphous alloys. 14.10.1 Feature of phase transition of bulk amorphous alloys
The phase transition of bulk amorphous alloys heated at a scanning rate of 1 K/s is different from that of conventional amorphous alloys which require high cooling rates above 104 K/s. The bulk amorphous alloys show the sequential transition of amorphous --+ glass transition --+ supercooled liquid region ~ crystallization upon continuous heating, while neither glass transition nor supercooled liquid region is observed for the conventional amorphous alloys. Thus the bulk amorphous alloys have a wide supercooled liquid region before crystallization. The appearance of the wide supercooled liquid region also implies that the supercooled liquid has a high resistance against crystallization. The high thermal stability is the origin for the achievement of the high GFA. The reason for the high stability of the supercooled liquid in the new systems has been described (Inoue 1995a, Inoue 1995b, Inoue 1996, Inoue 1997a, Inoue 1997b) in Section 14.2 on the basis of the formation of a higher dense random packed structure in the multicomponent alloys which satisfy the three empirical rules. 14.10.2 Deformation behavior of supercooled liquid
The Tg and ATx values of the bulk amorphous La55A125Nie0, Zr65A110Ni10Cu15 and Pd40Cu30Nil0P20 alloys measured at a heating rate of 0.67 K/s were respectively 480 and 70 K for the La-based alloy, 650 and 85 K for the Zr-based alloy and 575 and 95 K for the Pd-based alloy. Fig. 14.22 shows the relation between the maximum flow stress (O'max) and the strain rate (~) in the supercooled liquid for the La55A125Ni20 and Zr65A110Ni10Cu15 amorphous alloys (Inoue and Saotome 1993, Inoue et al. 1996c, Kawamura et al. 1996). The O'max increases with increasing ~ and the slope becomes significant in the higher temperature range. The slope corresponds to the strain-rate sensitivity exponent, m. As seen in the figure, the m-value is measured to be approximately 1.0 in the supercooled liquid region, indicating that the supercooled liquid has an ideal Newtonian flow, i.e., ideal superplasticity. Fig. 14.22 also shows that the deformation in the supercooled liquid takes place in a homogeneous mode over the whole strain rate range from 1.5 • 10 -4 to 7 • 10 -1 s -1 (Kawamura et al. 1996). This result also indicates that the high strain-rate superplasticity is obtained in the supercooled liquid. The amorphous alloy exhibits an inhomogeneous deformation mode in the temperature range well below Tg and the flow stress decreases slightly with increasing strain rate. Thus, the deformation behavior differs significantly between the amorphous solid and supercooled liquid. Figure 14.23 shows the change in the viscosity of the supercooled liquid for the Pd40Cu30Nil0Pe0 amorphous alloy during isothermal heating (Hata et al. to be submit-
410
A. Inoue 1000
~ . .I
~aK
100
~
_ f" P" / ~,f~ f " " " .,,eo ,..=..r'" .'~" /'653K o~/" ,/ik'" /
B~s~ w
=E
~
r 473K
fl'" 9
100
/
,
r
//:,/,"//= ,/
4
~ lO rr
~
~,"
/ IV/713K
10
~
10 "4
1 -3
LassAI2oNi25 , 10-1 100
' 10-2
ZresAiloNiloCU15 ,,
1o::4
I
10-3
....
I
I
10-2
10-1
Strain rate, ~ / s'l
Strain rate, ~ I s "1
Figure 14.22. Relation between the maximum true flow stress (O'max) and strain rate (~) in the supercooled liquid for the amorphous La55A125Ni20 and Zr65Al10Ni10Cu15 alloys.
16 14 12 o
>,
4-,
~
o o .(a >
......... qTm = 103 Pa.s q Tm -- 102 Pa. s
Pd40Nil oCu30P20
- ' - - qTm = 101 Pa. s
9TMA result
10
-
Fulcher t y p e e q u a t i o n
8
....... q
6
0 0.75
( Calculated ) _
9.34 x l O 3 e x p
0,,,~0
i 0.80
I 0.85
I 0.90
4135
T - 447
0.95
1.00
Figure 14.23. Temperature dependence of viscosity in the supercooled liquid for the amorphous Pd40Cu30Nil0P20 alloy. The solid circles represent the experimental data and the three lines denote the calculated values on the assumption of different viscosities of the liquid at Tin.
____
100
Bulk Amorphous Alloys
"'4 "'\N ,
~
~
)i!j !!
..1 :,'
"
........
"-';
,
'fj,,
"/;..,',
,.. ,'//
'...i~(:"/ // /'
:7.",~-~'. """ '~'
.,/" "
Figure 14.24. Wire- and gear-shaped amorphous La55A125Ni20 samples prepared by tensile stretching and die forging treatments, respectively, in the supercooled liquid region.
411
412
A. Inoue
. ....
~.~,lb;mm
Figure 14.25. Fine precision amorphous alloy mirrors prepared by die forging an amorphous Zr60All0Nil0CU20 alloy in the supercooled liquid region.
ted). Considering that the flow deformation becomes extremely easy in the viscosity range below 108 Pa.s, the viscous flow deformation can occur during the long time period of 102 to 104 s in the temperature range between 600 and 650 K. The times are long enough to form various complicated shapes because these amorphous alloys have the high strain-rate superplasticity.
14.10.3
Micro-formingof supercooled liquid
By utilizing the ideal superplasticity which can be achieved over the wide strain-rate range in the supercooled liquid region, various micro-forming treatments have been conducted on the La-, Zr- and Pd-based amorphous alloys. Figure 14.24 shows the outer shapes of the La55A125Ni20 amorphous wire obtained by tensile deformation at an initial strain rate of about 1 x 10 -2 s-1 and 500 K (Zhang et al. 1991 b) and the La-based amorphous gear with an outer diameter of 1 mm prepared by die forging into the gear-shape silicon die for 103 s at 500 K (Inoue and Saotome 1993). We have also obtained (Nishiyama and Inoue to be submitted) small amorphous gears by a die extrusion process in the supercooled liquid region for amorphous Zr65Al10Ni10Cu15 alloy with much higher Tg as compared with that for the La-based amorphous alloys. In the relation between extrusion pressure and extrusion velocity at 711 K in the supercooled liquid region for the Zr-A1-Ni-Cu amorphous alloy, there is a linear relation and the m-value is evaluated to be 1.0, indicating that the supercooled liquid also exhibits ideal superplastic behavior even in the die extrusion treatment. Figure 14.25 shows the outer morphology of the precision optical mirrors prepared by die forging the Zr60All0Nil0Cu20 amorphous alloy for 600 s at 673 K. In the as-forged state, the forged sample exhibits a smooth outer surface and good metallic luster. The outer surface keeps good smoothness on a nanometer scale from the data obtained by a surface roughness indicator. The three-stage amorphous gears with outer diameters of 5, 6 and 7 mm were also prepared by press die-forging of the Pd40Cu30Nil0P20 amorphous alloy for 600 s at 630 K (Hata et al. to be submitted). All micro-formed alloys retain the amorphous phase and no distinct difference in the thermal stability and mechanical properties is seen between the as-cast and micro-formed samples (Inoue et al. 1997d, Inoue and Zhang 1997, Kawamura et al. 1996, Nishiyama and Inoue to be submitted). The combination of
B u l k A m o r p h o u s Alloys
413
high GFA, good micro-formability and good mechanical properties has already enabled some practical uses of these bulk amorphous alloys as various micro-forged materials such as fine machinery parts and fine precision optical parts.
14.11. A P P L I C A T I O N S A N D F U T U R E P R O S P E C T S
As described above, the discovery of a number of new alloys with a high stability of the supercooled liquid against crystallization has led to a drastic decrease by 6 to 8 orders in the critical cooling rate for glass formation and enabled the production of bulk amorphous alloys with thicknesses ranging from several to about 80 millimeters by various casting processes. This drastic increase of the stability of the supercooled liquid also implies the advent of new basic science and materials. Table 14.5 summarizes the fundamental characteristics and application fields of the bulk amorphous alloys produced by conventional casting processes. It is confirmed that the bulk amorphous alloys exhibit various characteristics such as high mechanical strength, high elastic energy, high impact fracture energy, high wear resistance, high corrosion resistance, good soft magnetic properties, high frequency permeability, fine and precise viscous deformability, good castability and high consolidation tendency into bulk forms. By utilizing these characteristics inherent to the bulk amorphous alloys, the new alloys are expected to open up new application fields in the near future. Table 14.5 Fundamentalcharacteristics and application fields of bulk amorphous alloys Fundamental characteristic
Application field
High strength High hardness High impact fracture energy High elastic energy High corrosion resistance High wear resistance High viscous flowability Good soft magnetism High magnetostricfion
Machinery structural materials Optical precision material Tool materials Cutting materials Corrosion resistant materials Ornamental materials Composite materials Writing appliance materials Sporting goods materials Soft magnetic materials High magnetostrictive materials
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Author Index
Ananthapadmanabhan, E V., 121 Bhaduri, S. B., 289 Bhaduri, S., 289 Cahn, R. W., 1 Colligon, J. S., 225 Dollet, A., 257 Groza, J. R., 347 Inoue, A., 375 Ishihara, K. N., 5
Jones, H., 23 Kobayashi, K. E, 89 Koch, C. C., 313 Lavernia, E. J., 153 Li, B., 153 Liu, B. X., 197 Suryanarayana, C., 49, 313 Teyssandier, E, 257 Venkatramani, N., 121
417
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Subject Index
ABS s e e arc bond sputtering process (ABS) AC plasma torches, 126 Acetone, as process control agent, 52 Adhesion, ion-induced, 239 AFM s e e atomic force microscopy (AFM) Age-hardening, early studies, 1 Alloy phase formation, 41 ion mixing, 200-3, 217-18 Alloy phases free energy, 7-8 new, 41 nucleation, 15-17 Alloys disordering, 64 eutectics, 29 microstructure selection maps, 97-8 precipitation-hardened, 76 solid solution-strengthened, 76 solidification microstructure, 97 spray-formed, 169-70 supercorroding, 80 s e e a l s o amorphous alloys Alonso's method, 200-1 Alumina coatings, 133 plasma-sprayed, 131 scanning electron micrograph, 143-4 transmission electron micrograph, 143, 145 ultra fine powder, thermal plasma synthesis, 142-7 or-Alumina, synthesis, 306 y-Alumina, formation, 144-7 Aluminum powder plasma oxidation, 142-3 solutes, solid solubility
extensions, 31 aluminum-base alloys, applications, 79-80 aluminum copper alloys, mechanical properties, 185-6 aluminum iron alloys, solidification microstructure selection map, 35 Aluminum-iron-vanadium-silicon alloys, tensile properties, 42 Aluminum-nickel-manganese alloys, HallPetch plots, 41 Aluminum oxide ultrafine, 140 X-ray diffraction, 143 Aluminum-zinc alloys, mechanical properties, 187 AM s e e amorphous materials (AM) Amorphization conditions, 207 kinetics, 98 and milling temperature, 69 solid-state, 67-71, 199, 218 Amorphous alloys, 58, 199, 319 corrosion resistance, 98 formation, 70, 314-16 via ion mixing, 198, 207-8 metal-metal, formation, 197 metal-metalloid, formation, 197 metallic, 332 ternary, 380 s e e a l s o bulk amorphous alloys Amorphous materials (AM), structure, 268 Amorphous phase formation and atomic size, 69-70 conditions, 69
419
420
Subject Index
ion mixing, 200, 205-12 kinetics, 96-7 laser quenching, 98-9 in mechanical alloying, 49, 50, 61, 68, 81 in mechanical milling, 68 predicted composition ranges, 207-8 processes, 67-8 rapid solidification, 376 Amorphous phases enthalpy of formation, 201 free energy of, 6-7 heats of crystallization, 9 processing, 141 Anti-site disorder, 64 APS see atmospheric plasma spraying (APS) Arc bond sputtering process (ABS), 236-7 Arc discharge, 147 Arc melted ingots, 385-7 Arc plating system, 232 Arrhenius parameters, 272, 354 Atmospheric plasma spraying (APS), 134, 135 Atmospheric plasmas, 122 Atomic clusters, nanoscale, 313 Atomic collision cascade, use of term, 199 Atomic force microscopy (AFM), applications, 278 Atomic size and amorphous phase formation, 69-70 and solid solubility extensions, 32 Atomization electro-hydrodynamic, 147 melt, 23, 41 processes, 24-5 see also spray atomization Atomization gas, in spray forming, 160-3 Atomizers circular, 158 in spray forming, 158-63 Attritor mills, 53 Auto-ignition synthesis, of nanocrystalline oxides, 304-6
Automobile emissions, 340 Ball milling, nanocrystalline materials, 67 Ball mills ball-to-powder weight ratios, 57 in mechanical alloying, 49, 53-4, 71 Bearing steels, 41 Borides, reaction temperatures, 294 Brazing foils, 43 Bulk amorphous alloys, 375-413 alloy systems, 376 applications, 413 characteristics, 413 continuous cooling transformation, 3803 density, 392 electrical resistivity, 393 endothermic reactions, 403 glass-forming ability, 376-83 factors, 376-80 glass transition, 387-91 historical background, 375-6 hysteresis, 404-5 mirrors, 412 phase transition, 409 preparation, 383-7 properties, 408 magnetic, 403-6 mechanical, 395-9 physical, 392-5 specific heats, 387-91 structural relaxation, 387-91 supercooled liquids deformation behavior, 409-12 micro-formability, 406-13 synthesis, 403-6 tensile fracture behavior, 396-7 thermal expansion coefficient, 394-5 thermal stability, 384 thermograms, 388-91,393-4 viscoelasticity, 399-403 viscous flow, 406-13 Young's modulus, 399
Subject Index
c-BN s e e cubic boron nitride (c-BN) Calculating the phase diagram s e e CALPHAD CALPHAD, 7, 8 applications, 12 Carbides, reaction temperatures, 293 Carrier concentration, active, 103, 104 Catalysts heterogeneous, 340 nanostructured materials, 340-1 rapidly solidified materials, 41 CCPS s e e counter current plasma synthesis
(ccps) CCT s e e continuous cooling transformation
(CCT) Ceracon processing, 57, 326 Ceramics coatings, 133 nanocrystalline, 330 non-oxide, production, 140-1 processing, thermal plasmas, 140 CFUBMS s e e closed field unbalanced magnetron sputtering (CFUBMS) Charpy test, 397-8 Chemical vapor deposition (CVD), 225, 257-82 applications, 257, 316 approaches solid-gas thermodynamic equilibrium, 267-8 thermodynamics, 263-4 crystal growth, 270-1,278-81 morphology, 279-81 diamond deposition, 141-2 early studies, 2 fluid flows non-reactive, 260-1 reactive, 261-7 gas-phase mechanisms, 264 gas-phase transport, 259-67 growth rates, 263 kinetics, 264, 271-3 theoretical method, 273 layers, structure, 268-70
421
mechanisms, 258-9 modified, 265 nucleation enhancement, 276-7 epitaxial, 275-6 heterogeneous, 273-8 multi-component, 276 solid particle, 265-7 studies, 277-8 theories, 274-5 operating parameters, 261 plasma-assisted, 248 reactors, 257-8,260-1 research, 282 selective deposition, 277 solid phase formation, 267-81 supersaturation, 270-1 surface mechanisms, 271-3 thermodynamics, 263-4 Chemical vapor infiltration (CVI), 257 CHEMKIN/SENKIN package, 264 CIP s e e cold isostatic pressing (CIP) Closed field unbalanced magnetron sputtering (CFUBMS), 236, 237 Co-injection, and spray forming, 163-6 Coatings adhesion, 134 anti-reflective, 243,244 ceramic, 133 corrosion protective, 245 decorative, 225 hard, 246-8 ion-assisted, 249 nanocomposites, 248,249 nanostructured materials, 339 optical, 243-5 plasma spraying, 131,132-4 thermal barrier, 133, 134, 339 wear-resistant, 132, 339 Cobalt/tungsten carbide powder, production, 318 Cold isostatic pressing (CIP), 358 Cold-wall reactors, 257,258 Cold welding, 52, 55, 79
422
Subject Index
Combustion synthesis (CS), 63,289-307 advantages, 290 apparent activation energy, 295 applications, 290 auto ignition, 294 "azide" process, 294, 297-9 conversion efficiency, 297 development, 2, 289-90 enthalpy of formation, 291 field-activated, 295-6 gas phase reactions, 299 green density, 303 ignition chemical oven, 293 explosion mode, 293 kinetics, 294-5 mechanisms, 290 non-equilibrium effects, 307 propagating reactions, 296 reaction temperatures, 291-4, 299, 3012 thermodynamics, 291-4 use of term, 304-5 Combustion wave propagation, 304 Containerless processing, 41 Continuous cooling transformation (CCT) bulk amorphous alloys, 380-3 diagrams, laser processing, 96, 112, 113 Continuous wave lasers, 89, 94-5 cooling rates, 94 Copper-zirconium alloys, mechanical properties, 188 Corrosion protective coatings, physical vapor deposition, 243,245 Corrosive atmospheres, materials for, 77, 79 Counter current plasma synthesis (CCPS), 141 Crucibles, in spray forming, 158 Cryomilling, processes, 51 Crystal faces classification, 278-9 growth rates, 279-81 Crystal growth
chemical vapor deposition, 270-1,27881 laser processing, 96 mechanisms, 278 metastable phases, 12, 18 undercooling, 18 unstable, 280-1 Crystal-to-amorphous transformation, 205 Crystalline materials, synthesis, 268 Crystalline morphology, 279, 280 Crystalline phase formation laser quenching, 99-100 in mechanical alloying, 49 Crystallization explosive, 99 kinetics, 96 Cubic boron nitride (c-BN), coatings, 247 Cutting tools, nanostructured materials, 337 CVD see chemical vapor deposition (CVD) CVI see chemical vapor infiltration (CVI) CW lasers see continuous wave lasers DC plasma arc process, 139 DC plasma torches, 124-6 Dental amalgams, 43 Deposition dual ion beam, 233,234 electron beam vapor, 316 energy-assisted, 231,233,248 ion-assisted, 228,232, 238,249 ion beam, 233-5 ion beam induced chemical, 233,234 ionized cluster-beam, 226-7,233-5 laser-enhanced chemical vapor, 257 modified chemical vapor, 265 organometallic chemical vapor, 257 plasma-activated chemical vapor, 257 plasma-enhanced chemical vapor, 257, 268 rapid thermal processing chemical vapor, 257 thin film, 92, 230
Subject
chemical vapor deposition (CVD); electrodeposition; physical vapor deposition (PVD); pulsed laser deposition (PLD); spray deposition; vapor deposition diamondlike carbon (DLC), coatings, 246 Diamonds free energy, 5 plasma deposition, 141-2 Raman spectra, 142, 143 vapor-phase synthesis, 141 DIB s e e dual ion beam deposition (DIB) Differential calorimetry, 26 Differential scanning calorimetry (DSC), 9, 324 Direct reduced iron (DRI), 138 DISPAL, production, 50 Dispersion-strengthened materials, 162 research, 189 Dispersoids distribution, 71 in mechanical alloying, 50-1, 52 Displacement per atom (dpa), determination, 199-200 Displacement reactions, 49, 50 in mechanical alloying, 71-3 DMC s e e dynamic magnetic compaction (DMC) dpa s e e displacement per atom (dpa) DRAM s e e dynamic random access memory (DRAM) DRI s e e direct reduced iron (DRI) DSC s e e differential scanning calorimetry (DSC) Dual ion beam deposition (DIB), 233,234 Dynamic magnetic compaction (DMC), 368 Dynamic random access memory (DRAM), 100 see also
Index
423
Electromotive force (emf), measurement, 9, 10 Electron beam lithography, 341 Electron beam vapor deposition, applications, 316 Electron gun evaporation, 204 Electroplating, nanocrystalline materials synthesis, 328 Electrosleeve steam generator, 337 Elements, free energy, 5-7 Emf s e e electromotive force (emf) Emulsification, processes, 25-6 Enthalpy of mixing and glass formation, 30 mechanical alloying, 9-11 Miedema's model, 9 Epitaxy, 268 Equal-channel-angular technique, 318 Equilibrium phases, 29 Ethanol, as process control agent, 52 Evaporation, in physical vapor deposition, 228-9 EXCIMER lasers, 98-100, 103, 115 Explosive forming, 57
FACS s e e field-activated combustion synthesis (FACS) FAST s e e field-activated sintering technique (FAST) Fatigue crack growth, resistance, 79 Fe-6.3wt%Si, grain size, 32, 36 Feedstocks, rapid solidification in, 41 Ferromagnetic materials, nanoscale, 333, 334 Ferrous alloys s e e iron alloys Field-activated combustion synthesis (FACS), 294, 295-6 system, 298 Field-activated sintering technique (FAST), 363,364, 366-8 Field ion microscopy (FIM), 272 Electric arc spray process, 167-8 Filamentary structures, 313 Electro-discharge compaction, 57,326 Electrodeposition, nanostructured materi- Filtered arc system, 233 FIM s e e field ion microscopy (FIM) als preparation, 317, 335
424
Subject Index
FINEMET materials, 319, 333,335,340 Finite differences, and laser beam processing, 92, 94 Fluxing, in rapid solidification processes, 26 Frank-Read sources, 331 Frank-Van der Merwe mechanism, 275, 276 Free energy of alloy phases, 7-8 of elements, 5-7 interfacial, 219 of metastable crystalline phases, 216-18 see a l s o Gibbs free energy Free energy-concentration diagrams, 16 Free energy hierarchy, 7, 9, 13 Fuel cells, 43 Functional nanostructures, 313 Furnace racks, 79 Fusion welding, 91 heat-affected zones, 112
maps, 30 nominal, 211-12 Glass-forming composition range (GFR), and ion mixing, 198,205,206-7, 21112 Glass transition temperature, 7 Glassy phases, devitrification, 322 Glow discharge plasmas, 121 Glycine-nitrate process, 304 Golf club heads, 43 Grain growth abnormal, 324 inhibitors, 324 isochronal, 324 isothermal, 323 kinetics, 325-6 kinetics, 323 nanostructured materials, 323-6 Grain refinement, 64 Grain size and enthalpy of mixing, 9-11 and mechanical alloying, 71, 76 and rapid solidification, 32-3 Grinding balls, 53 Grinding media in mechanical alloying, 51, 74 velocity, 53
Gallium free energy hierarchy, 9 metastable phases, 15 Gas turbines combustion chambers, 79 vanes, 77, 78 Gasless combustion see combustion Hall-Petch equation, 183,240, 247, 329 synthesis (CS) Hall-Petch plots, 37, 41 GFR see glass-forming composition range Hall-Petch strengthening, nanocrystalline (GFR) materials, 67 Gibbs free energy, 6 Hard coatings, physical vapor deposition, determination, 201 246-8 Glass formation, 2, 32, 94 Hard facings, 43 bulk, 29 Hazardous wastes, treatment, 129, 141 critical cooling rates, 29 HE Mill, 54 and enthalpy of mixing, 30 Heat exchangers, components, 79 via combustion synthesis, 307 Heats of mixing Glass-forming ability (GFA) excess, 19 bulk amorphous alloys, 377-83 negative, 12, 69 and heats of formation, 206-8 Heptane, as process control agent, 52 intrinsic, 211 Heteroepitaxy, 269 and ion mixing, 198,205,206 Hexane, as process control agent, 52
Subject
Hierarchy of stability, 6 High-velocity oxyfuel thermal spray process, 167-8 HIP s e e hot isostatic pressing (HIP) Homoepitaxy, 268 Hot corrosion, resistance, 76, 79 Hot extrusion, 56-7, 326 Hot isostatic pressing (HIP), 56-7, 58, 290, 326, 363-6 Hot pressing, 290, 363,364 Hot-wall reactors, 257, 258 Hume-Rothery electron compounds, 62 Hydrogen, absorption, 337, 341 Hydrogen embrittlement, 104-5 Hydrogen storage materials nanostructured materials, 340-1 via mechanical alloying, 80 Hydrostatic extrusion, 326 Hysteresis effect, 237 IAD s e e ion-assisted deposition (IAD) IBC s e e ion beam induced chemical deposition (IBC) IBD s e e ion beam deposition (IBD) ICB s e e ionized cluster-beam deposition (ICB) ID s e e isothermal desorption (ID) Ideal solutions, phase diagrams, 16-17 IM s e e ion mixing (IM) Inert gas condensation, applications, 322 Ingot casting, and spray forming compared, 153-4 Integrated circuit pads, 115 Interaction parameter, 11 Intergranular corrosion, 335 Intergranular fractures, 185 Intermetallic matrix composites research, 189 spray-formed, 171, 172 synthesis, 168 Intermetallics disordering, 64-6 superlattice reflections, 65, 66 X-ray studies, 65
Index
425
ordered, 63 synthesis, 62-4, 67 Interstitials, redistribution, 64 Ion-assisted deposition (lAD), 228, 232, 238,249 Ion beam deposition (IBD), methods, 2335 Ion beam induced chemical deposition (IBC), 233,234 Ion bombardment, 228, 231 Ion implantation, 103, 153 applications, 197 early studies, 2 thermal plasmas, 122 Ion irradiation, 205 Ion mixing (IM), 197-220 alloy phase formation, 200-3 amorphous phase formation, 205-12 applications, 197 early studies, 2 experimentation, 203-5 glass-forming ability, 206-12 physics, 199-200 Rutherford backscattering spectrometry, 205 Ion-plating, 225,226, 227, 248 properties, 231 system, 232 Ionized cluster-beam deposition (ICB), 226-7,233,234-5 Iron alloys applications, 79 laser beam welding, 106-8 laser surface-alloying, 100-1 Iron-molybdenum multilayered films, backscattering spectra, 206 Iron oxide, ultrafine, 140 Iron-titanium carbide coatings, 136 Isothermal desorption (ID), 272 Joule effect, 92 Kiln furniture, 79 Kissinger method, 324, 403
426
Subject Index
Knock-out atom, 199 Lamellar structures see layered structures Lance pipes, coke injection, 79 Laser ablation, 92 Laser alloying, 98 Laser annealing, 90, 91, 102-5 applications, 103-4 Laser beam joining, 106-15 Laser beam processing, 94 and finite differences, 92, 94 Laser beam soldering, 90, 91-2 applications, 112 thermal cycles, 114 Laser beam welding, 90, 91 crack-flee, 112, 114 crack frequency, 113 ferrous alloys, 106-8 non-ferrous alloys, 108-12 weld fusion zones, 112 Laser-enhanced chemical vapor deposition (LECVD), 257 Laser glazing, 91 processes, 93-4 Laser-induced thermal desorption (LITD), 272 Laser irradiation applications, 99 for solution annealing, 104 Laser melting heat transfer, 92 processes, 92 and rapid solidification, 92, 95 Laser processing, 89-115, 122 classification, 89-92 continuous-cooling-transformation diagrams, 96, 112, 113 cooling rates, 96 crystal growth rate, 96 high-power, 98 partition coefficient, 97 solidification modes, 97 time-temperature-transformation diagrams, 96
Laser quenching amorphous phase formation, 98-9 crystalline phase formation, 99-100 nanostructured materials, 99 Laser scanning, in surface solutionizing, 105 Laser-solid interactions, 92-3 Laser surface alloying, 90, 91,100-2 Laser surface cladding, 90, 91, 100-2 Laser surface melting, 89-90, 99 Laser surface modification, 153 Laser surface quenching, 89-90 Laser surface solutionizing, 105 Laser surface treatment, 99 applications, 100-1 Lasers absorption phenomena, 92-3 absorption ratio, 92-3 early studies, 2 EXCIMER, 98-100, 103, 115 principles, 89 Q-switched sources, 89 spot size, 89 see a l s o continuous wave lasers; pulsed lasers Lattice damage, 103 Lattice parameter changes, 59 deviation, 12, 19 in disordered intermetallics, 65 of supersaturated phase, 11-12 Layer-by-layer mechanism, 275,277 Layered structures, 313 Lead shot, manufacture, 24 LECVD see laser-enhanced chemical vapor deposition (LECVD) Levitation (containerless) melting, in rapid solidification processes, 26 Ligaments, formation, 154 LIPS see liquid injection plasma synthesis (LIPS) Liquid crystal displays (LCDs), laser annealing, 103
Subject Index
Liquid injection plasma synthesis (LIPS), 128 applications, 141 Liquid melt quenching (LMQ), 197 LITD s e e laser-induced thermal desorption (LITD) LMQ s e e liquid melt quenching (LMQ) Low-friction materials, 132-3 Low-pressure plasmas, 121, 122
427
byproduct removal, 71-3 crystal defects in, 55-6 defect density, 59 deformations, 64 development, 5, 58 diffusion distances, 59 displacement reactions, 71-3 early studies, 2 enthalpy of mixing, 9-11 and grain size, 71, 76 MA s e e mechanical alloying (MA) grinding media, 51, 74 Magnesium-base alloys, applications, 80 interstitial contamination, 63 Magnetic materials mechanisms, 55-6 hard, 43,333-5 metastable phases, 51 permanent, 333-2, 335, 340 mills, 52-5 modeling, 75-6 remanence-enhanced, 73, 334, 335 nanocrystalline phases, 61 soft, 333,340 s e e a l s o ferromagnetic materials nomenclature, 50-1 Magnetic recording materials, 340 phase formation, 51 Magnetocaloric effect, 335 powders Magnetron sputtering, 230, 248-9 consolidation, 56-8 in physical vapor deposition, 235-8 contamination, 74-5 Magnetron systems, unbalanced, 227, 236, handling precautions, 81 particle size, 51 243 Marine corrosion, resistance, 79 processes, 51-5 Martensite phases, deformation-induced, and rapid solidification compared, 58 raw materials, 51-2 63 Martensitic transformation research, 2 early studies, 1 variables, 75 reverse, 215 Mechanical crystallization, 69 Material transfer, in mechanical alloying, Mechanical grinding (MG), 50 Mechanical milling (MM), 64, 66 50 Maximum possible amorphization range amorphous phase formation, 50, 68 (MPAR), 198,205 Medium-pressure plasmas, 122 MBRS s e e molecular beam reactive scatter- Melt atomization, 23, 41 ing (MBRS) Melt extraction, processes, 27 MCVD s e e modified chemical vapor depo- Melt extrusion s e e melt spinning sition (MCVD) Melt overflow, processes, 27 Mechanical alloying (MA), 49-81, 153, Melt spinning, 23, 41, 188, 340 340 cooling rates, 30 and enthalpy of mixing, 30 amorphous phase formation, 49, 50, 61, free-jet chill-block, 26, 27 68, 80 applications, industrial, 76-80 processes, 26-7 attributes, 49-50
428
Subject Index
Metal hydrides mechanical alloying, 51 applications, 341 nucleation, 13-15 production, 80 occurrence, 5 Metal matrix composites (MMCs), 43,293 physical vapor deposition, 240 synthesis, 163 processing, 141 Metal-metal systems, ion mixing, 209-10 spray forming, 178 Metallic glasses, 1,199 Metastable powders bulk modulus, 34 consolidation, 349-55 creep, 36 pressure, 363-6 ductility, 34 kinetics, 351 elasticity, 34 sintering fatigue limits, 36 mechanisms, 352-3 fracture, 34 scaling laws, 353-4 liquid melt quenching, 211 stability, 350 production, 141, 197,207, 212 thermodynamics, 350-1 properties, 34-7 Methanol, as process control agent, 52 electrical, 36 MG see mechanical grinding (MG) magnetic, 36 Microjoining, in laser beam processing, shear band offsets, 39 112-15 ultra-high strength, 26 Micromachining, and laser processing, 115 zirconium-based, 26 Microstructure formation, via rapid solidiMetalorganic vapor phase epitaxy process fication, 28-34 (MOVPE), 257 Microstructure selection maps (MSMs), alMetastabilities, types of, 348 loys, 97-8 Metastable alloy phases, 205 Miedema's theory, 9, 66, 198, 199, 200-1, Metastable alloys 218,219 formation, via ion mixing, 198, 199 Milling nanoscale composite, 247 ball, 67 Metastable crystalline alloys, formation, fracture rate-welding equilibrium, 55 199,213-18 high-energy, 55 Metastable crystalline phases, 58 intensity, 68, 74, 75 classification, 213-15 interrupted, 63 f.c.c.-I, 213-14 media, 74 f.c.c.-II, 215 reaction, 50 free energy, 216-18 research, 2 h.c.p.-I, 213 see also mechanical milling (MM) h.c.p.-II, 215 Milling atmosphere, 75 in intermetallic synthesis, 62 control, 50 Metastable phase formation and powder contamination, 74 kinetics, 5, 12-18 Milling temperature, and amorphization, 69 thermodynamics, 5-12 Milling/mixing see mechanical alloying Metastable phases (MA) crystal growth rate, 12, 18 Mills enthalpy of transformation, 9 commercial, 53-4 free energy, 9-11
Subject Index
high-energy, 53 low-energy, 53 new designs, 54-5 powder contamination, 74 types of, 52-5 s e e a l s o ball mills Mirrors, bulk amorphous alloys, 412 MM s e e mechanical milling (MM) MMCs s e e metal matrix composites (MMCs) Modified chemical vapor deposition (MCVD), 265 Molecular beam epitaxy (MBE), in physical vapor deposition, 229 Molecular beam reactive scattering (MBRS), 272 Moving point source model, 94 MOVPE s e e metalorganic vapor phase epitaxy process (MOVPE) MPAR s e e maximum possible amorphization range (MPAR) MSMs s e e microstructure selection maps (MSMs) Multilayered films interfacial free energy, 201-3 ion mixing, 197, 198 MX alloys s e e metastable alloys MX phases s e e metastable crystalline phases Nanocomposite refrigerants, 335 Nanocomposites, 338-9 applications, 339 coatings, 248,249 deformation behavior, 339 densification, 361 sintering, 361 synthesis, 306 Nanocrystalline alloys grain boundary mobility, 325 synthesis, 316 Nanocrystalline materials applications, 337-41 catalysts, 340-1
429
magnetic, 340 structural, 337-9 chemical reactivity, 335 corrosion behavior, 335 deformation mechanisms, 332 diffusivity, 328 dissolution rate, 335 early studies, 313 grain boundaries diffusion, 327 mobility, 325 segregation, 324 structure, 321-3 thickness, 320 grain growth, 323-6 grain size, 320-1 grain structure, 321 Hall-Petch strengthening, 67 hydrogen sorption, 337 magnetic, 319 mechanical twinning, 332 M6ssbauer spectra, 321 plastic deformation, 67 properties, 66, 327-37 chemical, 335-7 magnetic, 333-5 mechanical, 328-32, 337 shear banding, 332 structure, X-ray diffraction, 321 superplasticity, 332, 339 synthesis, 66-7, 104, 129, 306 gas condensation chambers, 316, 317 methods, 315, 319 spray conversion processing method, 318 thermal stability, 323, 361-2 three-dimensional, 317 volume fraction of atoms, 320 Nanocrystalline oxides, auto-ignition synthesis, 304-6 Nanocrystalline phases, in mechanical alloying, 61 Nanocrystalline powders, 319 Nanocrystalline spinel, synthesis, 306
430
Subject Index
Nanocrystallites, structure, 341 Nanodevices, 313 Nanofibers, production, 319 Nanoglasses, 314 Nanoparticles, 313,340 Nanoperm alloys, 333,335,340 Nanopowders ceramic, 355,356 compaction cold, 356-8 dry, 357-8 warm, 358-9 densification, 369 metal, 355,356 problems, 350 sintering, 328, 351,359-61 stability, 350-1 Nanoquasicrystals, 314 Nanorods, production, 319 Nanostructured materials, 313-42 applications, 337-41 catalytic, 340-1 electronic, 341 magnetic, 70, 340 structural, 337-9 atomic structure, 321-2 bulk, 313 classification, 313-14 composite, 70 corrosion behavior, 335 crystal lattices, atomic structure, 321 deformation mechanisms, 331-2 plastic, 318 diffusion, 327-8 diffusional creep, 330-1 ductility, 329, 331 early studies, 2 elasticity, 328 enthalpy, 9-12 ferromagnetic, 335 field of study, 313 grain boundaries, 67 atomic structure, 321-2
grain growth, 323-6 inhibition, 324-5 kinetics, 325-6 grain size, 318, 331 hardness, 329 junctions, 322-3 laser quenching, 99 microstructure, 320-1 optical, 341 powder consolidation, 326-7 properties, 327-37 catalytic, 335-7 chemical, 335-7 magnetic, 333-5 mechanical, 328-32 research, 341-2 retention, 57-8 sinterability, 327-8 stability, 323-6 strain rate, 331 strength, 329 structure, 319-23 superplasticity, 331 synthesis, 70, 314-19 via chemical reactions, 318 via devitrification, 319 via electrodeposition, 317 via inert gas condensation, 314-16 via mechanical attrition, 318 via rapid solidification, 316 toughness, 329 see also nanocrystalline materials New crystalline phases, 29 Ni-5wt%A1, grain size, 32, 36 Nickel, temperature-time curve, 29 Nickel aluminide, coatings, 133 Nickel-base alloys applications, 77-8 glass forming ability map, 30 properties, 77 Nickel-molybdenum system, free energy diagram, 217 Nickel-niobium system free energy diagram, 204
Subject Index
glass-forming composition range, 206-7 Nitrides, reaction temperatures, 299 Nitrogen, source of, 298 Non-equilibrium phases crystalline, 30 formation, 146, 147 nanoscale dispersions, 32 rapid solidification, 29-32 synthesis, 58-73 Non-equilibrium plasmas, 121 Non-equilibrium processing applications, 1 development, 5 early studies, 2 mechanical alloying, 49 in solid solubility extensions, 59 and Vegard's law, 11 Non-ferrous alloys laser beam welding, 108-12 laser surface-alloying, 101-2 Nopcowax-22DSP, as process control agent, 52 Nucleation frequencies alloy phases, 15 homogeneous, 13 metastable phases, 13-14 volume, 13 Octane, as process control agent, 52 ODS materials s e e oxide dispersionstrengthened (ODS) materials OMCVD s e e organometallic chemical vapor deposition (OMCVD) Optical coatings, physical vapor deposition, 243-5 Organometallic chemical vapor deposition (OMCVD), 257 Orowan looping mechanism, 184 Ostwald's step rule, 5, 13, 15 Oxalic acid, as process control agent, 52 Oxidation plasma, 142-3 resistance, 76, 78, 79, 101,133,301 surface, 341
431
Oxide dispersion-strengthened (ODS) materials, 49, 51, 59, 160-2 advantages, 77 applications, 76-80 sintering, 57 Oxide powders, production, 140 Oxides, single-phase, 306 Packaging, high-density, 112 PACVD s e e plasma-activated chemical vapor deposition (PACVD) Palladium-copper-nickel-phosphorus bulk amorphous alloys, 380-2 Palladium-copper-silicon alloys, constituent selection diagram, 35 PAS s e e plasma-activated sintering (PAS) Pattern transfer processes, 341 PCAs s e e process control agents (PCAs) PCM s e e polycrystalline materials (PCM) Pechini process, 304 PECVD s e e plasma-enhanced chemical vapor deposition (PECVD) Permanent magnets, materials, 41 Phase diagrams equilibrium, 7 metastable, 8 Phase separation, effects, 7 Physical vapor deposition (PVD), 225-49, 257 adhesion, 239-40 applications, 225,243-8 coatings, 243-8 corrosion protective coatings, 245 corrosion resistance, 243,245 deposition ion-assisted, 231-5 methods, 228-38 development, 225-7 early studies, 2 energy-assisted, 245 evaporation, 228-9 film morphology, 238-9 future trends, 248-9 hard coatings, 246-8
432
Subject Index
magnetron sputtering, 235-8 metastable phases, 240 microstructure, 240-3 molecular beam epitaxy, 229 novel phases, 243 nucleation, 239-40 optical coatings, 243-5 sputtering, 229-30 Planar diode sputtering, 230, 231 Planar flow-casting, 23 processes, 27 Planetary ball mills, 53 Plasma-activated chemical vapor deposition (PACVD), 257 Plasma-activated sintering (PAS), 57, 326, 366 Plasma cutting, 129 Plasma decomposition, 137-8 Plasma-enhanced chemical vapor deposition (PECVD), 257,268 Plasma furnaces, 137-41 Plasma machining, 129 Plasma melting, 122, 129 Plasma metallurgy, 138-9 Plasma oxidation, aluminum powder, 1423 Plasma-particle interactions, 128-9 heat transfer coefficient, 128 Plasma pressure consolidation (PPC), 366 Plasma processing applications, 316 development, 2 systems, 129 see also thermal plasma processing Plasma reactors, 137-4 1 Plasma spray forming, 167-8 Plasma spraying, 122, 128, 129-36, 147 advantages, 130 applications, 123, 132-4, 136 atmospheric, 134, 135 coatings, 131, 132-4 cooling rates, 132, 134 degree of flattening, 132 deposits, 131-4
low-pressure, 134 particle size, 131 plasma torches, 125 rapid solidification, 129, 134 reactive, 135, 136 spheroidization, 135-6 surface morphology, 131 vacuum, 133, 134 Plasma torches, 123-7, 129 AC, 126 applications, 138 DC, 124-6 electrothermal efficiency, 126 non-transferred arc mode, 125 plasma gases, 127 RE 126-7 transferred arc mode, 125 Plasmagen gases, 127-8 Plasmas DC, 138-9, 141-2 RE 138, 141 thermal conductivity, 128 types of, 121, 122 see also thermal plasmas Plastic deformation, of solids, 318 PLD see pulsed laser deposition (PLD) Polycrystalline films, morphology, 281 Polycrystalline materials (PCM), synthesis, 269 Powder consolidation, 347-69 densification, 347, 349 grain-size control, 361-2 in mechanical alloying, 56-8 metastability, 347, 348-9 compositional, 348-9 morphological, 348-9 topological, 348-9 methods, 356-68 Ceracon processing, 57 electrodischarge compaction, 57 explosive forming, 57 hot extrusion, 56-7 hot isostatic pressing, 56-7 plasma-activated sintering, 57
Subject Index
nanostructured materials, 326-7 and powder contamination, 355 pressure-assisted, 364-6 and rapid solidification, 347 shockwave, 368 sintering, 356-61 non-conventional, 366-8 Powder contamination in mechanical alloying, 74-5 and powder consolidation, 355 Powder formation, via rapid solidification, 24, 27 Powder metallurgy, and spray forming compared, 153-4 Powder-rolling, strained, 326 Powders agglomeration inhibition, 52 compaction, 357-9 handling precautions, 81 nanocrystalline, 319 oversprayed, 156 oxide, 140 particle size, in mechanical alloying, 51 ultrafine, 140, 305 s e e a l s o metastable powders; nanopowders PPC s e e plasma pressure consolidation (PPC) Precipitation-hardening, 49 Primary dendrite spacing, 33 Printed wiring boards, and laser processing, 114 Process control agents (PCAs), 50 applications, 52 and powder contamination, 74 as surface-active agents, 52 Pulsed laser deposition (PLD), 99 applications, 99-100 Pulsed lasers, 93-4, 98 cooling rates, 94 picosecond, 93 Pulsed plasma methods, 227 PVD s e e physical vapor deposition (PVD)
433
Q-HIP s e e quasi-hot isostatic pressing (QHIP) Q-switched sources, lasers, 89 Quantum Rice-Ramsperger-Kassel (QRRK) model, 264 Quasi-hot isostatic pressing (Q-HIP), 363 Quasicrystalline phases, 29, 49, 58 icosahedral, 62 in intermetallic synthesis, 62 non-equilibrium, 62 Quenching rate, determination, 94 Radial distribution function (RDF), 379, 381 Radiation, coherent, 89 Ranz-Marshall equation, 128 Rapid alloying, 115 Rapid solidification, 23-43 amorphous phase formation, 376 applications, 41 constitution changes, 28-33 development, 5 disordering processes, 64 in feedstocks, 43 free energy curves, 7 front velocities, 24, 28 and grain size, 32-3 growth temperature, 32 growth velocity, 23 and laser melting, 92, 95 materials coarsening, 40 corrosion resistance, 41 creep resistance, 41 properties, 34-42 and mechanical alloying compared, 58 methods, 24-8 droplet, 24-6 spinning, 24, 26-7 surface-melting, 24, 27-8 microstructure formation, 28-33 nanostructured materials preparation, 316 non-equilibrium phases, 29-32
434
Subject
in plasma spraying, 129, 134 and powder consolidation, 347 processes, 23 processing, 1 and spray forming, 168, 173-6 time-temperature-transformation diagrams, 29-32 and tungsten inert gas welding, 110 use of term, 23 Rapid thermal processing chemical vapor deposition (RTPCVD), 257 RBS s e e Rutherford backscattering spectrometry (RBS) R D F s e e radial distribution function (RDF) Reaction milling (RM), processes, 50 Reaction temperatures, in combustion synthesis, 291-4, 299, 301-2 Reactive plasma spraying, 135, 136 Reactive spray forming, 160-3, 168 Regular solution model, 11-12 Remanence, enhancement, 73,334, 335 RF plasma torches, 126-7 Rice-Ramsperger-Kassel-Marcus (RRKM) model, 264 R M s e e reaction milling (RM) Rod mills, 54 Rotating water bath, processes, 27 RRKM model s e e Rice-RamspergerKassel-Marcus (RRKM) model RSP s e e rapid solidification RTPCVD s e e rapid thermal processing chemical vapor deposition (RTPCVD) Rutherford backscattering spectrometry (RBS), in ion mixing, 205
Index
Scanning tunneling microscopy (STM), 272 Schaeffler diagram, 107-8 Secondary dendrite arm spacing, 33 Segregation-free solidification (SFS), 32 Self-propagating high-temperature synthesis (SHS) s e e combustion synthesis (CS) Self-supporting reaction waves, initiation, 294 Self-sustained high-temperature processing, 153 SEM s e e scanning electron microscopy (SEM) Severe plastic deformation consolidation (SPDC), 364 SFS s e e segregation-free solidification (SFS) Shaker mills, 53 Shear stress, critical resolved, 184 Shock waves, 290 SHS s e e combustion synthesis (CS) Silanes, reactive sticking coefficient, 265 Silica, ultrafine, 140 Silicides, reaction temperatures, 292-3 Silicon carbide amorphous, 99 powders, 140 Silicon-carbon films, 247 Silicon irons, applications, 40 Silver-copper alloys, solidification microstructure selection diagram, 34 Silver-copper system, 9, 10 lattice parameter deviation, 12 Silver-molybdenum system, free energy diSAP s e e sintered aluminum powder (SAP) agram, 208 SBE s e e single-beam plus evaporation Simultaneous combustion, 296 Single-beam plus evaporation (SBE), 233, (SBE) Scaling laws 234 Herring's, 354 Sinter-forging, 326 in powder consolidation, 353-4 Sintered aluminum powder (SAP), 162 Scanning electron microscopy (SEM), ap- Sintering plications, 278 conventional, 356-62 field-assisted, 366-8
Subject Index
metastable powders, 352-3 microwave, 366 nanopowders, 328, 351,359-61 non-conventional, 366-8 powder consolidation, 356-61 see also field-activated sintering technique (FAST); plasma-activated sintering (PAS) Skull melting, 158 Sliding wear, applications, 319 SMSD see solidification microstructure selection diagrams (SMSD) Sodium azide, as nitrogen source, 298 SOFCs see solid oxide fuel cells (SOFCs) Solid oxide fuel cells (SOFCs), 134 Solid phase formation, chemical vapor deposition, 267-81 Solid phases, interfacial energy, 9 Solid solubility extensions, 32, 59-62 aluminum solutes, 31 in spray forming, 176, 178-9, 185, 188 via non-equilibrium processing, 59 Solid solubility limits, 327-8 Solid solutions, 213,306 composites, 305 enthalpy of formation, 201 hardening, 36-7 in intermetallic synthesis, 64-6 in liquid immiscible systems, 61 phase diagrams, 17 To curves, 7 Vegard's law, 11-12 see also supersaturated solid solutions Solidification cooling rates, 24 kinetics, 96 models, 97 partitionless, 7 see also rapid solidification Solidification front, 23, 28 Solidification microstructure selection diagrams (SMSD), 32, 33-5 Solidification microstructures, alloy phase formation, 41
435
Solute partition coefficients, 33 Solute trapping, 32 Spark erosion, applications, 319 SPDC see severe plastic deformation consolidation (SPDC) Specific heat, measurement, 9 SPEX mills, 53 Spheroidization, 154-5 Splat formation, 25 Splat quenching, 23, 132 cooling rates, 30 early studies, 1 Spot welding, 27 Spray atomization, 154, 155, 166-7, 173, 176, 180, 188 cooling effects, 174-6 gas flow field, 155-6 gas jet nozzles, 160 use of term, 166 Spray conversion processing method, in nanocrystalline materials synthesis, 318 Spray deposition, 23, 25, 43, 154, 155,176, 180 processes, 156 Spray-formed materials grain size, 181 properties, 181-3 Spray forming, 41, 153-89 advantages, 168-72 applications, 168-72 atomizers, 158-63 circular vs. linear, 158-9, 161 close-coupled vs. free-fall, 159, 162 and co-injection, 163-6 coherent particles, 184-5 concept of, 154 cooling rates, 175 crucibles, 158 deposition, 176-8 development, 2 effects direct, 183-5 indirect, 185-8
436
Subject Index
environmental chamber, 156 grain size, 180-1, 183-5 and ingot casting compared, 153-4 limitations, 168 macrosegregation, 168, 172, 179, 183, 185, 187, 188 metastable phases, 178 microsegregation, 168, 172, 179, 183, 185 near net shape, 163 nomenclature, 166-8 non-equilibrium phenomena, 172-88 plasma, 167-8 and powder metallurgy compared, 153-4 principles, 154-6 and rapid solidification, 168, 173-6 reactive, 160-3, 168 research, 189 solid solubility extensions, 178-9, 185, 188 solidification rates, 175 surface tension, 155 undercooling, 173, 174 use of term, 157, 166 variations, 157-66 Spray pyrolysis process, 305 Sputtering, 225-6 applications, 316 ion-assisted, 225 in physical vapor deposition, 229-30 planar diode, 230, 231 thermal plasmas, 122 s e e a l s o closed field unbalanced magnetron sputtering (CFUBMS); magnetron sputtering Sputtering coefficients, 229-30 SRAM s e e static random access memory (SRAM) Stage turbine vanes, 78 Stainless steels, type 316, grain size, 32 Static random access memory (SRAM), 100 Stearic acid, as process control agent, 52
Steels bearing, 41 corrosion behavior, 139 creep strength, 139 laser beam welding, 106-8 stainless, 32 thermal plasma processing, 138-9 tool, 40 Steered arc system, 232 STM s e e scanning tunneling microscopy (STM) Stranski-Krastanov mechanism, 276 Stress corrosion cracking, resistance, 79 Structural difference rule, 198 Structural materials, welding of, 106-12 Subatmospheric-pressure plasmas, 122 Substrate-adsorbate interactions, 275 Sulfidation, resistance, 78 Super Misuni NEV-MA-8 (mill), 54 Superconducting alloys, 1 Superconducting materials, 100 ion implantation, 197 Supercooled liquids, 29 bulk amorphous alloys, 406-13 Superlattice reflections, in ordered intermetallics, 65, 66 Superparamagnetic materials, 333 Superplastic materials, 339 Supersaturated phases lattice parameter, 11-12 Ostwald's step rule, 15 Supersaturated solid solutions, 1, 2, 49, 58, 62,213 formation, 67, 109, 212, 307 Supersaturation, in chemical vapor deposition, 270-1 Surface-active agents, process control agents, 52 Surface alloying, 89 Surface melting, 41 non-equilibrium effects, 27 and rapid solidification, 27-8 s e e a l s o laser surface melting
Subject Index
437
rapid solidification, 29-32 Tip stability model, 33 Ti3SiCe, synthesis, combustion synthesis reactions, 299-301 To curves, metastable phase diagrams, 7-8 Titanium alloys, formation, 101-2 TAB s e e tape automated bonding (TAB) Tape automated bonding (TAB), and laser Titanium aluminides, nanocrystalline, 328 Titanium-aluminum-nitrogen films, added processing, 115 energy, 242 Target poisoning, 237 TEM s e e transmission electron microscopy Titanium-boron binary system, controlled (TEM) reactions, 301-4 Titanium dioxide, production, 140 Temperature-programmed desorption Titanium nitride, production, 140-1 (TPD), 272 Toluene, as process control agent, 52 TFTs s e e thin film transistors (TFTs) Thermal barrier coatings, 133, 134 Tool steels, 41 Thermal expansion mismatch, 133 TPD s e e temperature-programmed desorption (TPD) Thermal fatigue, resistance, 78 Thermal plasma processing, 121-48 Transformers core materials, 104 advantages, 121-3 laminations, 1 applications, 139 furnaces, 137-41 magnetic materials, 333 silicon irons, 41 homogeneous nucleation, 122, 147 Transgranular dimple rupture, 185 power sources, 129 Transmission electron microscopy (TEM), quench rates, 121, 122, 146-7 321 reactors, 130, 137-41 supercooling, 147 Traverse welding, 27 Tribological behavior, 99 systems, 129 Triple defects, 64 Thermal plasmas, 123-9 generation, 123 TTT diagrams s e e time-temperaturetransformation (TTT) diagrams power density, 123 Tungsten inert gas (TIG) welding, 108-11 Thermal shock, resistance, 133 cooling rates, 110 Thermal spike, concept, 200 hardness profiles, 108, 110, 111 Thermocouples, 79 Thin film deposition, 92, 230 weld fusion zones, 109-10, 111 Thin film transistors (TFTs), laser anneal- Turbine blades, 77, 78 ing, 103 Thin-strip casting, 41 Undercooled liquids Ti-6wt%A1-4wt%V, grain size, 32 frozen state, 6 homogeneous nucleation frequencies, 13 TiCN, chemical vapor deposition, 269 TIG welding s e e tungsten inert gas (TIG) Undercooling welding I and crystal growth rate, 18 Time-temperature-transformation (TTT) Ostwald's step rule, 15 and rapid solidification, 23, 24, 32 diagrams bulk amorphous alloys, 382 and spray forming, 173, 174 laser processing, 96 Uni-Ball-Mill, 54 Surface-mount technology, and laser processing, 112, 114
438
Subject Index
Vacuum furnaces, fixtures, 79 Vacuum plasma spraying (VPS), 133, 134 Vapor condensation, 147 Vapor deposition classification, 257 development, 5 see a l s o chemical vapor deposition (CVD); physical vapor deposition (PVD) Vegard's law, 19, 306 solid solutions, 11-12 Very large scale integration (VLSI) technology, ion mixing in, 197 Vibrating frame mills, 54 Vitrification, solid-state, 218 VLSI technology see very large scale integration (VLSI) technology Volmer-Weber mechanism, 276 VPS see vacuum plasma spraying (VPS) Wastes, hazardous, treatment, 129, 141 Welding
cold, 52, 55, 79 full penetration, 108 spot, 27 of structural materials, 106-12 traverse, 27 see a l s o fusion welding; laser beam welding; tungsten inert gas (TIG) welding Wiedemann-Franz law, 92 Yttrium-molybdenum system amorphous phase formation, 208 free energy diagram, 211 Yttrium-niobium system, free energy diagram, 204 Zirconia coatings, 133 production, 137 Zirconium oxide, ultrafine, 140 Zone I structure, 238
