Molecular Beam Epitaxy: Applications to Key Materials

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Edited by

Robin F. C. Farrow

IBM Almaden Research Center San Jose, California

r::l

NOVES

~

Park

Ridge,

PUBLICATIONS New

Jersey,

U.S.A.

Copyright Q 1995 by Noyes Publications No part of this book may be reproduced or utilized in any form or by any means, electronic or mechanical, including photocopying, recording or by any information

storage and retrieval system,

without permission in writing from the Publisher. Library of Congress Catalog Card Number: 94-31247 ISBN: O-8155-1371-2 Printedin

the UnitedStates

Published in the United States of America by NoyesPublications Mill Road, Park Ridge, New Jersey 07656 10987654321

LibraryofCongressCataloging-in-Publication Molecular

beam epitaxy

: applications

Data

to key materials/edited

by

Robin F. C. Farrow p.

cm.

Includes bibliographical references and index. ISBN O-8155-1371-2 1. Molecular beam epitaxy. 1. Farrow, R. F. C. QC611.6.M64M644 1995 621.3815’2--dc20 94-31247 CIP

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Editors Rointan F. Bunshah, University of California, Los Angeles (Series Editor) Gary E. McGuire, Microelectronics Center of North Carolina (Series Editor) Stephen M. Rossnagel, IBM Thomas J. Watson Research Center (Consulting Editor)

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Contributors

Philip I. Cohen Department of Electrical Engineering University of Minnesota Minneapolis, Minnesota

Vijay P. Kesan IBM Thomas J. Watson Research Center Yorktown Heights, New York

Robin F. C. Farrow IBM Almaden Research San Jose, California

Leslie A. Kolodziejski Department of Electrical Engineering and Computer Science Massachusetts Institute of Technology Cambridge, Massachusetts

Center

Robert L. Gunshor School of Electrical Engineering Purdue University West Lafayette, Indiana

Richard A. Kubiak The Department of Physics University of Warwick Coventry, England

Gerald R. Harp Department of Physics Ohio University Athens, Ohio

Eric C. Larkins Department of Electrical Engineering Stanford University Stanford, California

James S. Harris, Jr. Department of Electrical Engineering Stanford University Stanford, California

Ronald F. Marks IBM Research Division Almaden Research Center San Jose, California

Subramanian S. lyer IBM Thomas J. Watson Research Center Yorktown Heights, New York

Simon M. Newstead The Department of Physics University of Warwick Coventry, England xi

Xii

Contributors

Arto V. Nurmikko Division of Engineering and Department of Physics Brown University Providence, Rhode Island

Darrell G. Schlom Department of Materials Science and Engineering The Pennsylvania State University University Park, Pennsylvania

Nubuo Otsuka Materials Engineering Purdue University West Lafayette, Indiana

Philip Sullivan Fisons Instruments Inc. Danvers, Massachusetts

Morton B. Panish AT&T Bell Laboratories Murray Hill, New Jersey

Henryk Temkin Electrical Engineering Department Colorado State University Fort Collins, Colorado

Stuart P. Parkin IBM Research Division Almaden Research Center San Jose, California

Michael F. Toney IBM Research Division Almaden Research Center San Jose, California

Gale S. Petrich Department of Electrical Engineering and Computer Science Massachusetts Institute of Technology Cambridge, Massachusetts

Dieter Weller IBM Research Division Almaden Research Center San Jose, California

Thomas A. Rabedeau Stanford Synchrotron Radiation Laboratory Stanford, California

Gregory J. Whaley Phillips LMS Colorado Springs, Colorado

NOTICE To the best of our knowledge the information in this publication is accurate; however the Publisher does not assume any responsibility or liability for the accuracy or completeness of, or consequences arising from, such information. This book is intended for informational purposes only. Mention of trade names or commercial products does notconstituteendorsement or recommendation for use by the Publisher. Final determination of the suitability of any information or product for use contemplated by any user, and the manner of that use, is the sole responsibility of the user. We recommend that anyone intending to rely on any recommendation of materials or procedures mentioned in this publication should satisfy himself as to such suitability, and that he can meet all applicable safety and health standards.

Preface

Molecular beam epitaxy was initially developed by J. R. Arthur and A. Y. Chot11t21for growth of GaAs and GaAs/AlxGa,,As structures. It has subsequently been extended to an ever-widening variety of materials while maintaining key advantages over other techniques of epitaxial film growth such as chemical vapor deposition (CVD), liquid phase epitaxy (LPE), metal-organic vapor phase epitaxy (MOVPE), and related techniques. These advantages include the ability to control growth reproducibly to atomic monolayer dimensions and to monitor the growth process in real time. For example, the ultra-high vacuum growth environment of MBE makes it possible to study the dynamics of the growth process itself using modulated molecular beam techniques and RHEED (reflection high energy electron diffraction). In addition, other in-situ techniques such as XPD (X-ray photoelectron diffraction) can be used to examine the formation of interfaces and film growth modes. In this volume, we have set out to describe the use of MBE for a range of key materials fundamental

systems which are of interest for both technological

reasons.

and

Prior books on MBE have provided an introduction

to the basic concepts and techniques characterization

of GaAs-based

of MBE, and emphasise growth and Our aim in this book is structures.

somewhat different; it is to demonstrate the versatility of the technique by showing how it can be utilized to prepare and explore a range of distinct and diverse materials. The impact of MBE in each of these materials systems has been both beneficial and considerable. In Table 1, milestones in the chronological development of MBE are shown. Those in bold type are topics which form the basis of the chapters of this volume. In each

vii

viii

Preface

Table 1. Milestones

in the Development

of MBE

1968

Measurement of sticking coefficients of Ga, As, GaAs growth from molecular beam sources.t’)

1969 - 1970

Growth conditions for epitaxy of GaAs, from beam sources, established using RHEED.t2)

1971 - 1978

MBE established as a powerful, versatile film growth technique for preparation of conventional devices based on Ill-V compound

semiconductor

during

films.t3tt41

1975

First growth of Si,,GeJSi

1978

Observation of electron tion-doped GaAs.t61

1980

Introduction of gas sources compounds.r]

1981

Introduction systems.fel

1982

Discovery of fractional quantized Hall effect in 2-dimensional electron gas (2-DEG) in GaAs.fgl

1983

First observation of RHEED intensity ing growth of GaAs.[lO1[ll]

1984

First growth of pseudomorphic superlatticesJ21

1984

First observation of modulation heterostructures.[131

1987

superlattices.t5) mobility

enhancement

in modula-

for MBE growth

of high-throughput,

of Ill-V

production-style

MBE

oscillations

Si,,Gex/Si doping

dur-

strained layer in Si,,Gex/Si

First use of MBE for growth of high T, superconducting oxide films.[141

1988

Discovery

of giant

Fe/Cr magnetic

magnetoresistance

multilayer

1991

First achievement of high-conductivity ZnSe using nitrogen ion sourceJ61

1991

Achievement MBE-grown,

of 400,000 cm*V-‘s-l high-purity

in MBE-grown

films.[16] P+ doping

electron

GaAs films.[17]

mobility

of in

Preface

ix

of these areas, MBE is making considerable impact in terms of both devices and solid state physics. In several instances, the discovery of new physical phenomena was made possible by MBE synthesis of specific structures. For example, the observation by Dingle et al.t6j in 1978 of enhanced electron mobility in modulation-doped GaAs led to the subsequent discovery by Tsui et al.[a) of the fractional quantized Hall effect in a two-dimensional electron gas. Similarly: the discovetyt14) in 1988 of GMR (giant magnetoresistance) in MBE-grown Fe/Cr multilayers has enlivened the old field of magnetism and magnetic materials, and is leading to GMRbased devices such as rotation sensors and magnetic recording heads. Likewise, MBE techniques permit artificial layering of high-T, oxide superconducting films and provide a promising route to device structures and metastable phases which are difficult to access by more conventional growth techniques. In the field of II-VI semiconductors, the preparationt15) by MBE of high-conductivity, p-type ZnSe films, utilizing a nitrogen plasma source, provided the technological breakthrough to II-VI blue-green lasers, Similarly, the development of heterostructures of elemental semiconductors has been accelerated by the application of MBE to Si,,Ge,/Si heterostructures, We hope that the excitement of these developments and their implications is conveyed by the series of chapters in the present volume. Finally, the editor wishes to thank J. R. Arthur for his help and advice in planning this book, and is most grateful for the efforts and time which the contributors have put into their respective chapters.

1. Arthur, J. R., J. Appl. Phys., 39:4032 (1968) 2. Cho, A. Y., Suti Sci., 17:494 (1969); J. Apple Phys., 4132780 (1970); J. Appl. Phys., 42:2074 (1971)

3. Cho, A. Y,, Arthur, J. R,, Prog. in Solid State Chem., (G. Somorjai and J. M. McCaldin, eds.), 10:157, Pergammon Press, New York (1975) 4. Ploog, K., Crystals, Growth, Properties and Appl, 3:73, (H. C. Freyhardt, ed.), Springer-Verlag, Berlin-Heidelberg (1980) 5. Kasper, E., Herzog, H. J., and Kibbel, H., Appl. Phyq 8:199 (1975) 6. Dingle, R., Stormer, H. L., Gossard, A. C., and Wiegmann, W., Appl. Phys. Left., 333665 (1978)

i X

i

Preface 7. Panish, M. B., J. Electrochem. Sot., 127:2729 (1980) 8. See chapter 1 of this volume. 9. Tsui, D. C,, Stiirmer, H. L., and Gossard, A. C., Phys. Rev. Leti., 48: 1559 (1982) 10. Neave, J. H,, Joyce, B. A., Dobson, P. J., and Norton, N., A@. Phys. Leti., A31 : 1 (1983) 11. Van Hove, J. M,, Lent, C. S., Pukite, P. R., and Cohen, P. I., J. Vat. Sci. Technol., B1:741 (1983) 12. Bean, J. C., Feldman, L. C., Fiory, A. T., Nakahara, S., and Robinson, I. K., J. Vat. Sci. Technol., A2:436 (1984) 13. People, R., Bean, J. C., Lang, D. V., Sergent, A. M., Stdrmer, H. L., 1 Wecht, K. W., Lynch, R. T., and Baldwin, K,, Appl. Phys. Leti., 45 (1985) 14. Webb, C., Weng, S. -L,, Eckstein, J. N,, Missert, N., Char, K,, Schlom, D. G,, Hellman, E. S., Beasley, M. R., Kapitulnik, A., and Harris J. S,, Jr., Appl. Phys. Leti,, 51 :I 191 (1987); Kwo, J., Hsieh, T. C., Fleming, R. M., Hong, M., ,.Liou, S. H., Davidson, B. A., and Feldman, L. C., Phys. Rev,, B36:4036 (1987) 15. Binasch, G., Grunberg, P,, Saurenbach, F,, Zinn, W., Phys. Rev., B39:4828 (1989); Saurenbach, F., Barnas, J., Binasch, G., Vohl, M., Grunberg, P,, and Zinn, W., Thin Solid Films, 175:317 (1989); Van Dau, F. N,, Fert, A., Etienne, P., Baibich, M. N., Broto, J. M., Chazelas, J., Creuzet, G,, Friederich, A., Hurdequint, H., Redoules, J. P,, and Massies, J,, Journale de Physique, 49:C8-1633 (1988); Baibich, M. N., Broto, J. M., Fert, A., Nguyen Van Dau, F., Petroff, F., Etienne, P., Creuzet, G., Friederich, A., and Chazelas, J., Phys. Rev. Leti., 61:2472 (1988) 16. Park, R. M., Troffer, M. B., Rouleau, C. M., De Puydt, J. M., and Haase, MI A., Appl. Phys. Leti., 57:2127 (1990) 17. See Ch. 2 of this volume.

January 1995 San Jose, California

Robin F. C. Farrow

Contents

1

The Technology and Design of Molecular Beam Epitaxy Systems .........................................................

1

Richard A. Kubiak, Simon M. Newstead, and Philip Sullivan 1 .O INTRODUCTION ................................................................. 2.0 MOLECULAR BEAM EPITAXY ...........................................

1 2

3.0

MBE SYSTEM DEVELOPMENT ......................................... 4.0 VACUUM ............................................................................. 4.1 Vacuum Requirements for MBE.. ...............................

5.0

9 9

4.2 4.3

The Ultra-High Vacuum System.. ............................. Pumping ...................................................................

11 16

4.4 4.5

Ctyopanelling ........................................................... System Manufacture ................................................

24 25

MBE COMPONENTS:

SOURCES .................................... K-cells (also known as Thermal Effusion Sources or MBE Furnaces) ....................................................

26

5.1 5.2

Two-Zone

36

Thermal Dissociation

5.3 5.4 5.5

6.0

7

Cells.. ....................

Gas Source MBE (GSMBE) ..................................... Electron Beam Evaporators ..................................... Si-Filament Doping Sources .................................... 5.6 Electrochemical Doping Sources ............................. 5.7 Ion Sources in MBE.. ................................................ MBE COMPONENTS: SHUlTERS AND BEAM INTERRUPTORS ...................................................

xiii

29 38 46 52 53 56 60

xiv

Contents

7.0 MBE COMPONENTS: SUBSTRATE HEATER DESIGNS.. 7.1 Heaters for Ill-V MBE ............................................... 7.2 Substrate Heaters for II-VI MBE.. ............................. 7.3 Substrate Heaters for Si-MBE .................................. 8.0 TEMPERATURE MEASUREMENT AND CONTROL ........ 8.1 Thermocouple Measurements .................................. 8.2 Pyrometer Measurements ........................................ 8.3 Temperature and Process Control ........................... 9.0

10.0

65 67 69 69 71 71 72 74 76

8.4 Control Hardware ..................................................... 76 FLUX MONITORING TECHNIQUES.. ............................... 77 ............................ Ionization Gauge Flux Monitoring 9.1 79 9.2 Quartz Crystal Oscillators ........................................ 9.3 Optical Methods of Flux Measurement. ................... 81 PREPARATION, DlAGNOSTlCS AND ANALYSIS .......... .83 10.1 Vacuum Diagnostics: Gas Analytical Equipment ..... 83 10.2

Reflection Diffraction

High Energy Electron (RHEED) .................................................

10.3

Auger Electron Spectroscopy (AES) and X-Ray Photoelectron Spectroscopy (XPS) .......................... 88 Secondary Ion mass Spectroscopy (SIMS) .............. 92

10.4 11 .O MBE SYSTEM DESIGN: RETROSPECT AND PROSPECT ....................................................................... 11 .l Deposition Uniformity ...............................................

Production MBE: Throughput Considerations for MBE .................................................................... ...................... 12.0 PROCESS AND SYSTEM AUTOMATION REFERENCES ..........................................................................

84

94 94

11.2

2

99 102 103

Molecular Beam Epitaxy of High-Quality GaAs and AlGaAs . . . . . . . . .. . . . . . . . . . . . . . .. . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . . . . . . . . . . . 114 Eric C. Larkins and James S. Harris, Jr. .. .. .. .. . .. . ... .. ... . . .. . . . .. . .. . .. . .. . .. . .. . .. . . .. . .. . .. . .. . 114 1 .O INTRODUCTION 2.0 THE DEVELOPMENT OF HIGH PURITY MBE TECHNOLOGY .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . . .. . .. . .. . .. . 117 2.1 Vacuum Quality .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . . . . .. . .. . .. . .. . .. . .. . .. . 121 2.2 2.3 2.4

Impurities Substrate

Generated by Hot MBE Components .. . . 127 Purity .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . . . . .. . .. . .. . .. . 132 Source Purity .. . .. . .. . .. . ... .. . . . . .. .. . . .. . .. . .. . .. . .. . .. . . .. . . .. .. . .. . 133

Contents

3.0 GROWTH PROCESSES ................................................. ............................ 3.1 Atomic Surface Reconstructions 3.2 Surface Chemisorption ........................................... 3.3

Incorporation

of Chemisorbed

Island Formation 3.4

Surface

3.5

Incorporation

Diffusion..

146

............... 147

.................................................. Species:

152

Surface

Incorporation .......................................................... Gallium Desorption.. ............................................... Thermodynamic Redistribution of the Near-Surface Region.. ............................................ 4.0 SUBSTRATE ORIENTATION.. ........................................ 4.1 Growth on Misoriented (100) Surfaces.. ................. 4.2 Growth on (110) and Misoriented (110) Surfaces ... 3.6 3.7

4.3

137 138

Species:

and Step Propagation..

of Chemisorbed

xv

Growth on (nl l)A and (nl l)B (1 s n s 9) Surfaces

Growth on (221)A, (221)B, (331)A and (331)B Surfaces ..................................................... 5.0 OVAL DEFECTS ............................................................. 6.0 SURFACE MORPHOLOGY AND INTERFACE ROUGHNESS .................................................................. 7.0 SUBSTRATE CLEANING AND MBE GROWTH: ................ IMPURITY AND DEFECT INCORPORATION ...................... 7.1 Substrate Preparation and Cleaning 7.2 Protective Oxide Growth ........................................ 7.3 Wafer Outgassing and Oxide Desorption ............... 7.4 Buffer Layer Design ............................................... 7.5 Choice of Arsenic Species (ASH,, As,, As,). ..........

153 157 158 160 161 162 163

4.4

169 171 172 174 175 176 181

Role of Growth Temperature.. ................................ Role of V/III Ratio ................................................... Role of Growth Rate.. .............................................

183 200

7.9 Role of Growth Interruption .................................... ISOELECTRONIC AND UNINCORPORATED

207

7.6 7.7 7.8 8.0

166 166

DOPANTS 8.1 8.2

.......................................................................

lndium .................................................................... Antimony ................................................................

8.3 Hydrogen ................................................................ 8.4 Lead ....................................................................... 9.0 SURFACE PRESERVATION ...........................................

205

209 209 210 211 212 213

xvi

Contents

10.0

PREPARATION OF AN MBE SYSTEM FOR THE GROWTH OF HIGH PURITY III/V SEMICONDUCTORS.

214

11 .O CHARACTERIZATION TECHNIQUES FOR LAYERS ...................... EPITAXIAL SEMICONDUCTOR 11 .l 11.2 11.3 11.4 11.5

Deep-Level Transient Spectroscopy (DLTS) ......... Hall Effect.. ............................................................. Photoluminescence (PL) ........................................ Optical Absorption Spectroscopy ...........................

224 .225 228 232 235 235

Photoconductivity ................................................... 11.6 Photothermal Ionization Spectroscopy (PTIS) ...... .236 11.7 Secondary-Ion Mass Spectrometry (SIMS) ............ 239 12.0 IMPURITY ENERGY LEVELS IN GaAs AND AlGaAs ..... 240 244 ............................................................ ACKNOWLEDGMENTS 245 REFERENCES ..........................................................................

Gas-Source Molecular Beam Epitaxy: Ga,Jn,.,&_YP,,/lnP MBE with Non-elemental Sources. Heterostructures and Device Properties 275 Morton B. Panish and Hemyk Temkin ................................... 1.0 INTRODUCTION ............................................................. 2.0 CHEMISTRY.. .................................................................. 2.1 Thermodynamic Considerations-Arsenic and Phosphorus ..................................................... 2.2 Group V Dimer Beam Flux Requirements ............. 2.3 Group III Metalorganics .......................................... 3.0

GROUP V GAS SOURCES ............................................. 3.1 High Pressure Gas Source .....................................

275 275 279 279 280 282 284 284

3.2 Low Pressure Gas Sources .................................... 286 4.0 THE MBE AND GAS HANDLING SYSTEMS .................. 288 4.1 MBE System .......................................................... 288 4.2 Gas Handling of ASH, and PH, ........................................ 290 4.3

Gas Handling

of the Group Ill and Group V

Metalorganics ......................................................... 4.4 Dopants and Dopant Sources.. ............................... 5.0 PROCEDURES ................................................................ 5.1 Substrate Mounting and Temperature Measurement .......................................................... 5.2

Substrate

Preparation

............................................

293 295 299 299 300

Contents

6.0 SINGLE 6.1 6.2

BULK LAYERS

7.1

301

GaAs, InP, InGaAs, and InGaAsP by GSMBE 301 (Hydrides and Elements) ........................................ GaAs, InP, GaInAs, and GaInAsP by MOMBE ..... .302

7.0 QUANTUM 7.2 7.3 7.4

..................................................

WELL AND SUPERLATTICE

STUDIES

..... .304

High Resolution X-ray Diffraction by Superlattices .. .305 Optical Properties--Single Quantum Wells ........... 308 Optical Properties of Superlattices ......................... 312 Avalanche Photodetectors and Superlattice

Modulators ............................................................. 7.5 Transport Through The Superlattice ...................... 7.6 Strained Layer Superlattices .................................. 7.7 Heterojunction Bipolar Transistors ......................... ACKNOWLEDGMENTS ............................................................ REFERENCES ..........................................................................

4

xvii

Molecular Beam Epitaxy of Wide Gap II-VI Semiconductor Heterostructures ........................

315 322 324 329 337 338

.344

Leslie A. Kolodziejski, Robert L. Gunshor, Arto V. Nurmikko, and Nobuo Otsuka ........................................... 1 .O GENERAL INTRODUCTION 1 .l Diluted Magnetic Semiconductors .......................... .......................... 2.0 CdTe-BASED HETEROSTRUCTURES 2.1 2.2 2.3 2.4 2.5

344 345 346 346 Introduction ............................................................ Heteroepitaxy of CdTe on (100) GaAs ................... 347 Quantum Well Structures Incorporating (Cd,Mn)Te .358 376 Binary ZnTe/CdTe Superlattices ............................ II-VI Quantum Wells Incorporating MnTe 377 Barrier Layers .........................................................

2.6 InSb Multiple Quantum Wells with CdTe Barriers.. .......................... 3.0 ZnSe-BASED HETEROSTRUCTURES 3.1

Introduction

382 387 387

3.2

............................................................ Homo- and Heteroepitaxy of ZnSe .........................

3.3 3.4 3.5 3.6 3.7

Quantum Well Structures Incorporating (Zn,Mn)Se .405 Epitaxial Growth of the Metastable (Zn,Mn)Se ..... .407 Optical Properties of (Zn,Mn)Se Quantum Wells.. .409 ZnSe/MnSe Magnetic Superlattices ....................... 421 ZnSe/ZnTe Superlattice Structures ........................ 429

3.8

Blue and Blue/Green

387

Laser Diodes and LEDs ...... .433

xviii

Contents

4.0 SUMMARY .. .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . . . . .. . .. . .. . .. . .. . 438 ACKNOWLEDGMENT .. .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . 439 REFERENCES

5

Elemental Semiconductor HeterostructuresGrowth, Properties, and Applications.. ............... .453 Vgay 1 .O 2.0 3.0 4.0 5.0

6

.. .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . ... .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . 439

453 P. Kesan and Subramanian S. lyer ............................... INTRODUCTION ............................................................. 453 GROWTH OF Si,,Ge, ALLOYS ..................................... 453 STABILITY OF Si,,Ge, FILMS ....................................... 463 LONG RANGE ORDER IN THE Si,,Ge, SYSTEM ........ 468 DEVICE APPLlCATlONS OF Si,,Ge, ALLOYS ............. 480 5.1 Heterojunction Bipolar Transistors (HBTs) ............ .484

5.2 Heterostructure FETs ............................................. 5.3 Optoelectronic Devices .......................................... 5.4 Other Quantum Well Structures.. ........................... 6.0 CONCLUSIONS.. ............................................................. ACKNOWLEDGEMENTS .........................................................

485 491 494 497 497

REFERENCES

497

..........................................................................

MBE Growth of High T, Superconductors

. . . . . . . . . . . 505

Darrell G. Schlom and James S. Harris, Jr. 1 .O INTRODUCTION .. . .. . .. . .. . .. . .. . .. ... ... .. . . .. . .. . .. . .. . .. . .. . . .. . .. . .. .. . . 505 1 .l Crystal Structures and Types of Building Layers.... 508 1.2 Chemical Doping .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . . .. . .. . .. . 512 1.3 Phase Diagrams .. .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . 514 1.4

Uncontrolled Intergrowths Inherent in Bulk Methods. 522

1.5

Layer-by-Layer MBE Growth .. . .. . .. . .. . . . .. . . .. . .. . .. . . .. . .. . 527

2.0 OXIDE MBE SYSTEMS 2.1 2.2 2.3

.. . .. . .. . .. . .. . .. . .. . .. . .. . .. . . . . .. . .. . .. . .. . .. . .. . 528

MBE System Configuration .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . 528 In-situ Analysis .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . . . . .. . . .. . .. . .. . 532 Minimum 0, Necessary to Form Structure .. . .. . .. . .. . 537

2.4

Maximum 0, Satisfying MBE Mean Free Path

2.5 2.6 2.7

Constraint .. . .. . .. . ... . .. .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . . . . 540 Alternative Oxidants .. .. . .. . .. . .. . .. . . . . .. . .. . .. . .. . .. . .. . .. . .. . . . . 544 Ozone System .. .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . 550 Composition Control ,.............................................. 552

2.8

Crucibles .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . . .. . .. . .. . .. . 555

Contents

2.9 2.10

Common Substrates.. ............................................. Integration with Semiconductors ............................

3.0 SPECIFIC

HIGH T, MATERIALS

DEMONSTRATED

SYNTHESIS

557 563

AND CAPABILITIES

........... .564 565 30 7-r)...................................................................................

3.1

ReBa&u

3.2 3.3 3.4 3.5

Bi,Sr,Ca,,_,Cu,O,,+, .............................................. 567 571 TI,Ba,Ca,.,_,Cu ” 0 sn+4............................................................. (Ba,K)BiO,. ................................................................................. 572 Superlattices.. ......................................................... 573

3.6 Josephson Junctions .............................................. 3.7 Formation of Metastable Structures ....................... 3.8 Twin-Free Growth.. ................................................. 4.0 FUTURE DIRECTIONS ................................................... 4.1 Hybrid MBE Techniques.. ....................................... 4.2

In-situ Monitoring

Techniques

................................

5.0 CONCLUSIONS.. ............................................................. ACKNOWLEDGMENTS ............................................................ REFERENCES ..........................................................................

7

xix

581 582 584 588 588 590 592 594 594

MBE Growth of Artificially-Layered Magnetic Metal Structures .. ................................ ... ............... . 623 Robin F. C. Farrow, Ronald F. Marks, Gerald R. Harp, Dieter Weller, Thomas A. Rabedeau, Michael F. Toney, Stuart S. P. Parkin 1 .O INTRODUCTION . . . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. ... . .. . .. . . .. . .. . .. . . . . 623 2.0 SEEDED EPITAXY OF MAGNETIC METALS . . . .. . .. . ...a... 626 2.1 2.2

Semiconductor Substrates . . . .. . .. . .. . .. . .. . .. . .. . .. . .. . . .. . .. . 626 Insulating Substrates . . . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . . .. . .. . .. . 632

3.0 STRUCTURAL

AND MAGNETIC

PROPERTIES

OF

ARTIFICIALLY-LAYERED MAGNETIC METAL STRUCTURES . . . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . . . . .. . .. . .. . . .. . .. . .. . .. . 638 3.1 Rare Earth Metal Sandwich Structures . . . .. . .. . .. .. . . .. . 638 3.2 3.3 3.4 3.5 3.6

Fe/Ag Films and Multilayers . . .. . .. ... . .. . .. . .. . .. . .. . .. . .. . .. . 642 Fe/Ag-Seeded Sandwiches of Fe/Q/Fe, Fe/Au/Fe, and Fe/AI/Fe . . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . . . . .. . 642 Seeded Epitaxial Co/Pt Superlattices . . .. . .. . .. . .. . .. . .. . 643 Co-Pt alloy films . ... . .. .. .. .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . .. . 651 Giant Magnetoresistance in MBE-Grown Co/Cu Multilayers

. . . .. . .. . .. . .. . .. . .. .. . . .. . .. . .. . .. . .. . .. . .. . .. . .. . 654

xx

Contents

3.7

Giant Magnetoresistance in 2-Phase Heterogeneous Alloy Films ....................................

660 661 662 662

4.0 CONCLUSIONS.. ............................................................. ACKNOWLEDGMENTS ............................................................ REFERENCES ..........................................................................

8

Reflection High Energy Electron Diffraction Studies ...... .669 of the Dynamics of Molecular Beam Epitaxy 1. Cohen, Gale S. Petrich, and Gregory J. Whaley ..... INTRODUCTION ............................................................. DIFFRACTION GEOMETRY.. ......................................... .................................. DlFFRACTlON FUNDAMENTALS ....................................... Kinematic Approximation 3.1 3.2 Disorder on Low-Index Surfaces ............................ 3.3 Vicinal Surfaces ..................................................... 3.4 Disorder on Vicinal Surfaces .................................. ................................. 4.0 DIFFRACTION MEASUREMENTS 4.1 Low-Index Surfaces ................................................

669 669 670 675 675 679 681 683 685 686

4.2 Vicinal Surfaces ..................................................... 4.3 Strained Layer Growth ........................................... 5.0 SIMPLE GROWTH MODELS .......................................... 5.1 Perfect Layer-Growth ............................................. 5.2 Nondiffusive Growth on a Low-index Surface ........ 5.3 Diffusive Growth on a Low-index Surface ..............

698 719 724 726 727 729

Philip 1 .O 2.0 3.0

5.4 Diffusive Growth on a Vicinal Surface ................... 6.0 CONCLUSION ................................................................. ACKNOWLEDGMENTS ............................................................ APPENDIX: TWO-LEVEL DIFFRACTION ................................ REFERENCES ..........................................................................

Index .............*..................................................................

731 735 735 736 739

745

The Technology and Design of Molecular Beam Epitaxy Systems Richard A. Kubiak, Simon M. Newstead, and Philip Sullivan

1 .O

INTRODUCTION

In essence, MBE is little more than a UHV-based evaporation method. In practice, it is a material deposition technique capable of predictably and reproducibly yielding material with impurity levels below ten parts per billion, with unprecedented control over the precision with which the composition and doping of the structure can be tailored.tl]-flo] Some of these attributes are intrinsic to the MBE process, e.g., slow growth rates and low deposition temperatures. Others, such as material quality and purity, rely on the technology employed. Much of the rapid development of MBE, particularly

in the last five

years, can be ascribed to the willingness of manufacturers of vacuum equipment to come to grips with the complex and often unique instrumentation involved.fllI In particular, the rapid changes in market requirements from custom-made special ultra-high vacuum evaporators to dedicated high-throughput

MBE instruments

have been effectively

challenged,

and

the results have met with considerable success within the semiconductor industry. In this chapter, we chart the progress made in MBE system technology and thereby illustrate current design practice. The development of MBE related components, system geometries, and the MBE process will be covered. Although the physics of growth mechanisms and growth methodology are dealt with elsewhere in this book, we briefly

1

2

Molecular

Beam Epitaxy

consider these aspects where pertinent to the discussion. Where rigorous referencing is difficult, the authors unashamedly draw from their own biased experience.

2.0

MOLECULAR

BEAM EPITAXY

To put in perspective the technology discussed briefly consider the MBE process and those epitaxial

in this chapter, we materials currently

being addressed. MBE involves the generation of fluxes of constituent matrix and doping species (molecular beam) and their reaction at the substrate to form an ordered overlayer (epiraxy). Figure 1 shows a schematic representation of the process, and its components. Elemental or compound constituents are heated (if in the liquid or solid state) or introduced (if gaseous) to cause mass transfer from the flux generators to the substrate, via the vapor phase. To maintain the high purity and integrity of the deposit, stringent vacuum conditions are needed. MBE is essentially a line-of-sight technique from source to substrate, and the fluxes of constituents (and thus the composition of the material perpendicular to the growth direction) can be temporally modulated either by altering the evaporation/ introduction conditions, or by physically interrupting the beam using rapidaction mechanical shutters. A key attribute of MBE is the precision with which the composition and doping of a structure can be tailored, such that atomically abrupt features can be produced. Examples are to be found throughout this book. To achieve this level of control within realistic time spans, deposition rates centered around one atomic layer (a monolayer) per second are used. This places constraints on the operational temperatures of sources, and the speeds with which shutters are required to operate. From the above, the key features to be addressed technology associated with the MBE process are: 1.

Vacuum

2.

MBE components

3.

requirements

l

sources

l

shutters

l

substrate

heating and manipulation

MBE diagnostic

and analytical

facilities

in a discussion

of

Technology

However,

and Design

of MBE Systems

to offer a viable MBE process, other pertinent

4.

Factors governing

5.

Throughput

6.

Automation

7.

Modern manufacturing

areas include:

MBE chamber design

considerations

methods

‘2

I

----\I

SUBSTRATE

1

OPERATOR AN or COMPUTER

Y--

FLUX

3

GENERATORS

,NTRG&s&,ON1 EVAPORATION I

also controlling: VACUUM SUBSTRATE PREPARATION SUBSTRATE HANDLING DIAGNOSTICS

GAS FEED

Figure 1. Schematic representation of the MBE process The evaporation procedure, flux incidence on the substrate, diagnostics are controlled by a supervisory operator or process is supplemented by ex situ substrate preparation procedures.

and control interface. vacuum, and process computer. The MBE and wafer introduction

This chapter deals with the design of components and systems used in MBE. Unfortunately, a generic MBE system configuration does not exist, since this depends on the nature of the deposit and the behavior of the constituent source materials. Nevertheless, Table 1 identifies three methodological areas into which most MBE activities can be categorized, and gives examples of material systems to which each area applies. Also included in Table 1 is a list of the sources encountered in each area (though all are not required for a given material system), and the section in which they are described in this chapter.

4

Molecular Beam Epitaxy

Table 1. Types of source employed in each of the three methodological divisions of MBE (a particular application would not usually require the use of all types of source listed within the appropriate division). Sources

Section

Uses

Hemoepitaxy

Heteroepitaxy

Conventional (K-cell based) K-cell 5.1 r Cracker 5.2 Si-filament 5.5 Electrochemical 5.6 Hg introduction

Ill-v’s II-VI’S IV-VI’S HTcSC

Gas source (GSMBE) Hydride Decomposition 5.3 Metallorganic 5.3 K-cell 5.1 Cracker 5.2

III-V’S ? Si

50/75mm 750°C 75-200mm 900°C

Si & related materials Metals

75-200mm 900°C Small area 900°C

Electrochemical Si-filament

5.6 5.5

High temp. evaporator Electron beam Evaporator 5.4 K-cell 5.1 Ion sources 5.7 B-evaporator 5.1 0, injector 5.3

50/75mm 750°C

Si 50/75 (-2OOmm?)900°C Small areas Small areas Small samples 800°C

{

Si 75-200mm

900°C

Superconductors Small samples 800°C I

Notes: 1. HTcSC = High T, superconductors 2. Temperatures are maximum values

The most active area of MBE is in GaAs/AixGa,,As applications deposited using “conventional” condensed phase (solid or liquid) sources. A schematic of a commercially available III-V:MBE system, described in detail in subsequent sections, is shown in Fig. 2. The equipment used also lends

itself to epitaxy

of other

Ill-V,

II-VI

and IV-VI

semiconductors,

provided that the vapor pressures of the source materials at temperatures

below 1200°C are sufficiently high for evaporation from K-cells (effusion furnaces). The evaporation plane in these systems is near-horizontal so that the sources are angled to prevent stray deposits (e.g., from chamber walls and shutters) from falling vertically into the crucible. The severity of this problem has not, however, been demonstrated. The current generation

Technology

and Design of MBE Systems

5

of III-V:MBE systems is capable of accommodating standard 50 mm and 75 mm wafer sizes processed individually. More recently, 100 mm capability

has become available,

onto three

50 mm wafers

discussed

in Sec. 12.

also offering the possibility

simultaneously.

Larger

scale

of depositing systems

are

GROWTH CHAMBER

CHAMBER

Figure 2. Schematic of a III-V:MBE system. The deposition chamber on the left is fitted with up to 8 thermal effusion (K-cells) and other sources (see Sets. 5.1, 5.2, 5.3, 5.5, and 5.6) configured to achieve optimal uniformity of deposit (see Sec. 12). Shutters (Sec. 6) can interrupt the flux to yield rapid changes in composition or doping. The deposition chamber is connected via an in-line gate valve to the preparation chamber, in which substrate storage and diagnostics can be performed, and to a fast entry chamber (Sets. 2 and 3). Ultra high vacuum conditions are maintained throughout the system to achieve high material quality (see Sec. 3). The MBE system shown handles platens capable of accommodating one lOO-mm, one 75-mm or three 50-mm wafers. (Courfesy VG Semicon.)

The second methodological

area of MBE activities

relates to depo-

sition of Ill-v’s using gaseous source materials, Gas Source MBE. This new technology offers a variety of potential advantages over conventional MBE in terms of control and accuracy over fluxes, indefinite source material lifetime, and suitability to scaling. In general, conventional lllV:MBE systems are used for GSMBE with the K-cells replaced by gas effusion process

and hydride cracker sources, and appropriate pumping for the added. This approach does not necessarily represent the

6


optimum system design for this technology, but no doubt the recent increase in activity in the area can largely be ascribed to the ease with which gas sources can be retrofitted to conventional The remaining

MBE activities

MBE systems.

relate to materials

requiring

source

temperatures in excess of those attainable with K-cells, necessitating electron beam evaporation (MBE of Si and related materials, metals, and superconductors). By their nature, electron beam evaporators dictate the need for a vertical evaporation geometry, with the substrate in the horizontal plane. This permits the substrate to be maintained in the horizontal plane throughout the MBE system, providing simplified larger wafer sizes, particularly in the case of Si:MBE diameter).

A schematic

handling of the (75 to 200 mm

of a Si-MBE system is shown in Fig. 3.

GROWTH

CHAMBER

PREPARATION WAMBER

Figure 3. Schematic of an MBE system used for deposition of Si and related materials, metals, and superlattices. The matrix evaporation sources are electron beam evaporators (Sec. 5.4) although thermal effusion (Sec. 5.1) ion beam (Sec. 5.7) and other species specific (Sec. 5.1) sources can be fitted for matrix and doping flux generation. The deposition geometry is configured to achieve optimal uniformity of deposit (see Sec. 12). Shutters (Sec. 6) can interrupt the flux to yield rapid changes in composition or doping. The deposition chamber is connected via an in-line gate valve to the preparation chamber, in which substrate storage and diagnostics can be performed, and to a fast entry chamber (Sets. 2 and 3). Ultra high vacuum conditions are maintained throughout the system to achieve high material quality (see Sec. 3). The MBE system shown handles wafers up to 150 mm diameter without the need for wafer holders. (Courtesy VG Semicon.)

Technology

3.0

MBE SYSTEM


7

DEVELOPMENT

To aid in an understanding of current MBE methodology and system designs (such as shown in Figs. 2 and 3), it is useful to take a brief look at the historical

development

of the MBE process.

The first MBE systems

incorporated

evaporation

and substrate

heating facilities, and some diagnostic and analytical equipment into a single stand-alone vacuum chamber. In many respects, the MBE sources employed were natural developments of high vacuum precursors, refined to ensure compatibility with the UHV environment. The presence diagnostic and analytical equipment expedited improved understanding

of of

the epitaxial processes involved (just as use of these techniques today offers assistance to the MBE practitioner in establishing optimized or reproducible deposition conditions). Deposition sample areas used in the early systems were small (typically less than several cm2). The small source volumes employed (several cc’s) were adequate when replenished each time the system was vented for the loading of a new substrate. Aside from limiting throughput of samples (to a maximum of one sample a day for the most agile of operators!), the need for air exposure of the system between each deposition run resulted in poor and irreproducible material qualityt12)t13] because system venting contaminated the sources, precluded a thorough de-gassing of sources prior to growth, and exposed the substrate to a poor vacuum during system bake-out. Borrowing from the technology of other vacuum processing, throughput and material quality were improved by use of a “Fast Entry Lock” (FEL) for introduction of substrates into, and removal of processed samples from, the MBE deposition region.n4t A valve between the FEL and deposition chamber ensured high vacuum integrity in the deposition region while the FEL was vented to air.

Transfer

of samples

between

the two

vacuum chambers was performed when the pressure in the FEL was better than 10e6 mbar. Although several methods of sample transfer were reported

(such as insertion

of substrate

manipulators

into the deposition

region,t15] and transfer of entire substrate heater assemblies,[14] the preferred methods proved to be those which required introduction of the substrate and a minimal substrate holder (if any at all), thus minimizing contamination of the deposition environment by outgassing of air-exposed surfaces. With this development, the UHV lifetime of the MBE system became limited only by cell depletion times, typically extending to several tens of microns of material. paramount importance!

Reliability

of MBE components

became

of

8

Molecular

nience

In addition to dramatically enhancing throughput and the conveof system operation, the use of FEL’s significantly improved

material quality. were

held under

Beam Epitaxy

The substrate vacuum

manipulator,

for extended

source cells and materials

time periods

(typically

several

weeks) leading to thorough outgassing. K-cells could be maintained at temperatures 50 to 100°C below their normal operating range when not in use, preventing recontamination by residual gases, and extending crucible lifetime for such materials as Al, which tends to crack boron nitride crucibles if it is allowed to solidify rapidly. Substrates were no longer subjected to extended bake-out periods, but were introduced into the clean deposition area as required. The benefits of load-locked operation are apparent in modern MBE systems, in which the first few samples grown after system bake-out provide for conditioning of the system, with higher quality material being achieved during the second, third or subsequent growth runs.t16)t17j Recent MBE system designstlO] interpose a preparation chamber between the deposition chamber and FEL, as shown in Figs. 2 and 3. Samples are no longer introduced into the FEL individually, but in batches of typically ten. After evacuation of the FEL, the samples are transferred into the preparation chamber “parking stage,” where, once the FEL has been isolated, they are stored under UHV conditions. into the deposition chamber thus does not require

Transfer of a sample exposure of the air-

exposed FEL to the deposition chamber, minimizing contamination of the deposition environment. The preparation chamber also provides a UHV environment in which samples can be pre-processed (e.g., heated to subgrowth temperatures to de-gas the sample) or analyzed. Component-specific ponent maintenance.

load-locks

have also been used for MBE com-

For example, the UHV lifetime of III-V:MBE

extended by reloading the most rapidly depleted P). Retraction of these sources into small volume can be isolated from the deposition chamber with and reloading of the cell without disturbing the Before reinsertion

into its deposition

can be

sources (usually As and load-lock chambers that a valve, permits venting deposition environment.

position, the load-lock station may be

briefly baked, and the source cell de-gassed. Most MBE practitioners would still bake the entire system to ensure vacuum cleanliness after reintroduction of the source cell. The use of component load-locks can expedite simplified and rapid system turn-round, and maintain high material UHV.

quality,

since

all other components

remain

outgassed

under

Technology

Another lesson transferred to increase

throughput

and Design

of MBE Systems

from high vacuum deposition

in MBE systems

is use of substrate

9

techniques movement.

The simplest form of motion, namely rotation, is used in MBE.t’st (More complex motions such as planetaryflg) are avoided in MBE, even for multiwafer systems,

due to the limitations

in bearing lubrication

technology

in

UHV, see Sec. 4.2) Sample rotation significantly improves the uniformity over the sample areas used,t18)f1g) and, even for non-optimized source/ substrate geometries, increases the area over which acceptable deposition uniformity occurs. The optimization of MBE system geometries is discussed

in Sec. 11.

4.0

VACUUM

4.1

Vacuum

Requirements

for MBE

Vacuum provides a unique environment in which materials can be prepared, characterized, and modified, and thus plays an important role in a wide range of technologies.f20)-f30) The quality of the vacuum required (i.e., both the residual gas pressure and its composition) depends on the influence of the residual gases on the process. It is instructive, therefore, to consider the vacuum requirements for MBE. The behavior of gases as a function of pressure is described by the Kinetic Theory of Gases.f31) This yields important relationships between gas pressure, molecular density, the mean-free path of molecules (the distance they travel between collisions with one another), and the impingement

rate of molecules

walls, or a substrate); which corresponds

on a surface

(such as the chamber

see Fig. 4. A useful reference unit is the monolayer, to coverage

of a flat surface

by one atomic

layer.

Given that a flat crystalline surface has between 1014 and 1015 atoms cmm2, Fig. 4 indicates that a pressure of approximately 1Om6mbar corresponds to an impingement rate which would lead to addition of one monolayer in one second,

if all incident

onto the surface.

species

In general,

(assumed

residual

to be atoms)

gas species

adsorbed

are sufficiently

volatile not to adsorb onto a surface. However, in the case of a newly deposited surface, free chemical bonds can enhance the adsorbtion of residual gases, leading either to their incorporation disruption of the growth process.

as an impurity,

or to

10

Molecular

__

I

Beam Epitaxy

PRESSURE (mbar) I

I

I

I

I

I

I

I

I

I

I

I

1

10-1110-1010-910-010-7 10-610-510-~lo-310-210-' 1 10 IO210' I

I

10”

10’

I

I IdO

DENSITY (molecules I

I

IO8

IO9

I

I

I

I

I

I

I

IO"10"10'"10" Id"Id7IO'" Id9

loI0 lo”

MEAN FREE PATH I

cmm3 at 25’C)

(mm at 25°C)

I

I

I

I

I

I

I

I

IO9IO8IO7IO6IO5IO4IO3IO2 IO 1 IO4lo-210-310-4

_ IMPINGEMENTRATE (moles. cm-2sec-1 at 25°C) I

10'"10" lo"1~31d41d51d6 IO"1d81d91020102110221023 _ RATE OF GAS IMPINGEMENT (monolayers I

IiS

I

I

I

I

seed

I

1641631ci2 16' 1 10 IO2IO3IO"IO5IO"IO'IO8

Figure 4. The relationship technology.

between the fundamental

units encountered

in vacuum

Given that the growth rates used in MBE correspond to approximately 1 monolayer set-’ the pressure levels required during MBE can be estimated.

As an example, let us consider the requirements

for the level of

carbon in MBE-grown GaAs. C is a common component of several major residual gases (CO, CO,, CH,, and other hydrocarbons), and is also an effective p-type dopant and deep level in GaAs. For many device applications, C levels in GaAs below 1014 cm3 are required, i.e., an impurity level of 1 atom in lOa. Taking this worst case as the ratio of the impingement monolayer

rates of C-bearing set-‘)

residual

and growth species

gas species

(1 monolayer

(i.e., 10”

set-l),

of a

a maximum

permissible pressure of C-bearing gases in the vacuum of 10”/lOa

= lo-l4

mbar or below is indicated.pj The sticking coefficients of C-bearing species are fortuitously much less than unity, and although total permissible pressures perhaps as high as 10-l’ mbar could be tolerated, even this level represents an exacting requirement of vacuum quality. Similar arguments

apply to incorporation

of all other residual

gas constituents

in

Technology

any MBE-deposited afford to permit Indeed, essential

material.

In practice, then, MBE practitioners

any compromise

it is the availability to many

applications


in achieving

clean

of the UHV environment

deposition

processes,

cannot

UHV conditions.

which

particularly

11

makes MBE

in semiconductor

where the purity of the material is paramount.

Furthermore,

MBE practice involves care in siting and use of hot filaments (e.g., ion gauges) relative to the growth region to minimize the presence of excited gas species surface.t4j 4.2

with

potentially

The Ultra-High

There system:

enhanced

reactivities

with the depositing

Vacuum System

are potentially

four

major

1. High vapor pressure materials,

sources

and materials

of gas in a vacuum with poor thermal

or chemical stability. These can generally be discounted, because materials used in construction of MBE systems are selected to ensure high vacuum quality (i.e., minimized partial pressures of deleterious species). Nevertheless, the MBE process may necessitate use of high vapor pressure source materials (e.g., As and P in III-V:MBE), which may lead to cross contamination of sources, and impart a high load on the pumps. 2.

Gas adsorbed onto surfaces during air exposure. This is the main source of concern in achieving UHV. A surface exposed to air accumulates several monolayers of chemisorbed and physisorbed gases (see Fig. 5),t1g1f251[2gjt30j which slowly de-gas, precluding rapid pump-down of the vacuum system. Several hours, or even days, may be necessary to achieve pressures below 10e7 mbar (depending on exposure history). Therefore, UHV systems are routinely ‘ibaked”t21)t30] after each air exposure to temperatures

of about 25O”C, to accelerate

the desorp-

tion, leaving, after eight hours or more, a conditioned low vapor pressure surface. On cooling to ambient temperatures, pressures of the order of 10-l’ mbar and below can be achieved. An alternative method of stimulating gas desorption from vacuum surfaces (particularly water vapor) is irradiation with UV light.f3*] This method alone cannot achieve the stringent UHV levels required within the MBE deposition chamber, but is useful for

12

Molecular

Beam Epitaxy

speedy de-gassing of entry locks and introduced wafers. Venting of vacuum systems to dry nitrogen rather than air greatly reduces gas adsorption, and reduces the times needed for pumping,

baking, and de-gassing

the components.

MBE com-

ponents which need to be operated at high temperatures

(e.g.,

filaments, source cells) are outgassed at temperatures higher than those of operation (but, in the case of source cells, within the vapor pressure constraints of the source during bake-out, or just prior to use. 3.

materials)

either

Dissolved gases or impurities within the constructional materials which diffuse to the surface and desorb. Such materials can be conditioned by heating in vacua, as #2 above. The generation of hydrogen during evaporation of Si and metals under UHV is an example.

4.

Poor vacuum

integrity

due to leaks.

This

is addressed

by

employing suitable constructional materials and methods.f21j-f30) Early UHV systems were manufactured from glass or quartz.t20) However, the need for easy access into the larger chamber volumes in current use (which imposes large stresses on chamber walls due to the pressure differential) has stimulated the maturation of a UHV technology based on stainless steel,t21]-f30j and, more recently, aluminum alloys.f2e)f301t33j Access into the chambers is facilitated by removableflanges.t25]-[27] Small flanges (usually up to 300 mm port size) use flat copper gaskets as a seal, into which knife-edges on the flanges bite, known as “Conflat” seals. Larger flanges are generally compression seals, using annealed sing, chambers

Au, Cu, or Al wire rings. are often internally

To minimize

electropolished

outgas-

which

re-

duces the vacuum surface area by between two and five times. Argon-arc welding of stainless tubes or plates to form chambers minimizes inclusions and oxidation at welds, and realizes high weld penetration to ensure strength, absence of leaks and low permeability.f27jt30] Grades of stainless steel are selected for high molybdenum (e.g., 316 grade) and low carbon (e.g., 304 and 316 grades) content to minimize chromium diffusion at welds.f27)t30j

Technology

h

-

and Design

of MBE Systems

13

-

5-25 MONOLAYERS OF PHYSISORBED H20

STABLE SURFACE SKINOF CHEMICALLY BONDED OXIDES,CARBIDES, NITRIDES etc..

Figure 5. A representation of the coverage of gases sorbed onto an air-exposed stainless steel surface. The uppermost water-rich layer is weakly physisorbed and is easily removed by baking at temperatures below 120°C or by irradiation with UV light. The remaining chemisorbed gases necessitate higher temperatures (in excess of 200°C) to promote efficient desorption, or reaction with the stainless steel to form a stable low vapor pressure surface. Gases dissolved in the near surface of the stainless steel are also removed by baking.

Chambers and pumps are interconnected by valves which permit isolation of the various sections of the system (e.g., of the pump during chamber venting). Several types of valves have been developed, capable of withstanding many resealing ing on the design constraints).

cycles (in the range lOO-30,000 dependValves with metal (copper, gold or silver)

sealing faces are used where the valves need to be closed during bakeout. For example, between the chambers and permanently plumbed roughing lines (which are outside the bake-out zone), small conductance (up to 100 I/s) right angle valves permit bake-out of the entire UHV portion of the system. To stop oxidation of the sealing material, both sides of the valve must be maintained under vacuum during bake-out. Since transfer of samples between chambers requires linear motion usually through large apertures to accommodate current wafer sizes (typically 100-300 mm diameter), so-called “gate” valves are employed. Although these gate valves are of all-metal construction to ensure UHV comparability, elas-

14

Molecular

Beam Epitaxy

tomer O-rings are most frequently used for valve sealing.f35)f36] The most commonly used O-ring material, Vifon, can be baked to temperatures of 250°C provided not in compression, bake-out.

Being reasonably

i.e., the valves need to be open during

soft, and having a small compression

Viton is tolerant of ingress of evaporant

dust (present in large quantities

MBE).

(100-l

Being of narrow cross-section

50 mm between

set, in

flanges),

gate valves also have high conductances making them ideal for valving pumps. Gate valves with metal sealing rings are also available for applications necessitating valve closure during bake-out, for example to preclude cross-contamination between chambers. Such all-metal valves are prone to leaking due to ingress of evaporant dust, and large bore metal sealing gate valves needed for pumping and sample transfer (150 mm and larger) are costly. Other materials, potentially offering the benefits of both metal and elastomer seals (e.g., Kalrez) are under investigation.t351f36) Other key components (e.g., electrical, water, liquid nitrogen and rotary/linear

motion feedthroughs)

conform

to conventional

UHV prac-

tice.f3q Viewports, which are used for process and sample transfer observation and pyrometry, are generally shuttered with a mechanical flap to prevent “fogging” by evaporant species. Kodial or a similar borosilicate glass is generally used. However, glassware has been shown to generate volatile boron compoundst38t for which the SiO, surface of Si wafers has an affinity. The boron compounds do not desorb during oxide removal, and can lead to p-type doping at the substrate/epilayer interface in Si:MBE.t3s] Quartz viewports (employing direct quartz to metal seals, and not graded through glass) eliminate the B-doping problem,f3s)f40) but large quartz viewports capable of withstanding bake-out temperatures of 250°C are expensive.

The use of other types of glass, or of transparent

coatings

may provide

electrical

insulation

a simpler

solution.

material for conductors

Glass

provides

and components

viewport

a versatile operating

at

temperatures below 300°C (except in Si-MBE for the above reason). The choice of suitable, stable and low vapor pressure metals for use as heater filaments

and heat shields

tungsten,

tantalum,

better than

99.9%,

is restricted

and molybdenum. and in various

to refractory These

forms

elements

are available

including

wire,

such as at purities

rod, and foil.

Tungsten is extensively used for filaments, but is unsuitable for K-cell and other applications containing confined heaters because, being brittle, thermal stressing can lead to failure. Tantalum, which is less affected by thermal cycling, has become more commonly used for these heaters. Tantalum can be easily welded. Of the three metals, molybdenum is the

Technology

and Design

of MBE Systems

15

least difficult to machine, and is thus used for larger support components and threaded components (e.g., nuts and bolts). MO looses its machinability after thermal cycling, cannot be welded, and requires more thorough outgassing

than W or Ta due to formation

of volatile

oxides.

material has to be selected for MBE usage, since commercially MO formed by sintering

MO

available

is laminar and difficult to outgas thoroughly.

“Arc

cast” MO is preferred. Other refra%ry metals, e.g., rhenium, are prohibitively expensive. The key vacuum, mechanical and electrical characteristics of these metals can be found in Refs. 21-30. Insulation of filaments (e.g., inside K-cells and substrate heaters) has been performed using quartz, beryllium oxide, alumina, and hot pressed boron nitride. Unfortunately, these materials have been shown to contribute contaminant fluxes in MBE environmentst41)-t43) due to thermal and chemical (e.g., reduction by the filaments) instability, or the presence of volatile impurities (e.g., BO in BN). Pyrolytic boron nitridet44] (pBN) is now used almost exclusively as an electrical insulator in contact with or near heaters. pBN is considerably less convenient than alumina, because the method of manufacture (pyrolysis in a CVD reactor) permits only simple shapes, and fine dimensional tolerance is difficult to achieve. pBN is also soft, brittle, and delaminates under mechanical stress. Despite these shortcomings, pBN is a near ideal material for MBE by virtue of its high purity, chemical and thermal stability, and non-porous structure. Nevertheless, care is required during extended high temperature treatment to prevent contamination of pBN caused by decomposition of Cbearing species from the residual gases.t45) Some care is required in its use in Si:MBE, since some thermal/chemical decomposition occurs at temperatures above 1300°C which can lead to B-doping.t46) The need for motion

(e.g., substrate

rotation,

sample

movement

around the UHV system) necessitates bearings capable of sustained operation in UHV for both the manipulator and rotary feedthrough. Allstainless-steel ball races are available, compatible with the UHV environment and high temperature lubricants

developed

bake-out.

Unfortunately,

greases and other

for high vacuum (down to 1Om8mbar) are not compat-

ible with the clean, hydrocarbon-free

environment

required for MBE, and

proprietary bearing lubrication processes have been developed by most UHV companies. Also bearings using dissimilar metals are becoming available. The need for high-speed rotation of ever increasing sizes of substrate platens poses a challenge to MBE practitioners and manufacturers alike. These are however being addressed

as discussed

in Sec. 8.

16

Molecular

4.3

Pumping

Beam Epitaxy

Aside from sources of gas in vacuum systems, the other factor influencing the level of vacuum attainable is the efficacy of pumping. The ideal UHV pump would have a high pumping speed for all gases (i.e., would be non-selective) and would itself not contribute to the gas load in the system. For MBE, pumping is achieved by combination of several pumps, aimed at eliminating the most damaging gaseous components and dealing with technology-specific gas loads, for example, H, in GSMBE. Pumps fall into two classes, primary pumps used to achieve and maintain UHV under quiescent conditions, Table 2, and secondary pumps, which provide pumping appropriate to the process, Table 3. In many instances, the same pump provides both functions. Primary pumps (Table 2) are of two types. “Capture” pumps collect gas by gettering/implantation (sputter iont26)t471t4stand titanium sublimationt26)t301 pumps) and freezing of the gas (cryo-pumpst4g]-t53)); the gas remains trapped within the UHV system pump. Throughput (diffusion~~~l~~~l~~~l~~~l-~~~l and turbomoleculart28tt57)-t5gt) pumps compress the vacuum gas to a vented outlet held at 10” mbar or below, where it is removed by a “backing” pump. Detailed discussions of their operation can be found in the expert texts referenced above. UHV pumps are normally connected to UHV chambers via high conductance valvest34) which permit isolation and continued operation of the pumps while the UHV chamber is vented to air. The merits of this arrangement are in maintaining cleanliness of the pump, promoting easy and rapid pumpdown, and, if automated, serving as a safety barrier in case of vacuum or pump failure (e.g., to prevent oil contamination from turbo or diffusion pumps). As is apparent from Tables 2 and 3, the upper pressure at which most UHV pumps will operate is approximately 10” mbar. Table 4 lists the pumps available for evacuation of vacuum systems to this level and for backing throughput pumps, and presents their characteristics. Rotary pumps[231[261[30l[s0lare generally not used for roughing MBE systems due to the potential for oil contamination by backstreaming. Liquid nitrogen cooled sorption pumpst23tt301t61]provide for oil-free evacuation, although care is needed to ensure that sorbate dust does not enter the UHV chamber. The large volume of MBE chambers necessitates two or more sorption pumps to be used sequentially for evacuation to below 10” mbar. Alternatively, the chambers can be pre-evacuated to approximately 100 mbar (removing 85% of the gas) prior to sorption pumping by use of compressed-air Venturi or oil-free rotary vane pumps. The use of preevacuation can greatly expedite the speed, cleanliness, and cost of evacuation. Evacuation (and venting to air) is performed slowly to

Technology

minimize turbulence (at pressures which could cause contamination particles of deposit or dust. Tables 5 and 6 summarize pumps used in the deposition

and Design

of MBE Systems

17

between atmospheric and 50 mbar), of sources due to redistribution of the most common

and appendage

chambers

combinations respectively,

of in

the three methodological areas of MBE. With the exception of gas source MBE (and systems handling phosphorus or mercury), primary pumping by ion-sputter plus titanium sublimation pumps (TSP) suffice to achieve UHV. There is some evidence that operation of TSP filaments during MBE deposition can be detrimental to III-Vf62t and Sifrs3t material quality, although whether this is inherent to TSP operation, or circumstantial, is unclear. The provision of secondary pumping via liquid nitrogen cryopanels is essential for many MBE processes (see Sec. 4.4), and further addition of a cryopump helps to minimize partial pressures of other damaging residual gas species (e.g., CO, CH,). In some cases (e.g., Si:MBE) cryopumps have been successfully used as the primary UHV pump.f51tf52) To date, throughput pumps have generally, though not always,f64) been avoided in MBE due to the potential for oil contamination of the deposition environment. The recent development of oil-free turbomolecular (using magnetic levitation of the rotor) and rotary pumpsf65] removes this objection (albeit presently at high cost). For gas source MBE (and solid source MBE of P-bearing materials), the high gas loads of process gases and hydrogen, and the toxic nature of the gas products necessitates throughput pump use. The safety aspects associated with handling of the toxic and pyrophoric source gases necessitate stringent safety procedures and exhaust scrubbing. The process gases are corrosive and degrade oils and bearings, therefore turbomolecular and rotary pumps are provided with nitrogen gas ballast to dilute the gases, and specially selected oils to maximize service times (typically between 1 and 3 months depending on throughput and type of source gas). To further assist rotary pump operation, phosphorus and metal-organic trapping is performed in the backing line using molecular sieve and activated charcoal, and the backing line is designed to permit Diffusion pumps using easy servicing of the filters and pumps. polyphenylether oils (such as Santavac 5) have been found to be resistant to corrosion from most process gases encountered in GSMBE. The availability of oil-free turbo-molecular and rotary pumps may expedite future development of simpler pumping systems for these processes. MBE of materials containing mercury presents unique problems due to the large volumes of Hg needed for growth. Mercury diffusion pumps are employed.

Table

2(a).

Primary

Pump/Pressure range (mbar)

UHV “capture”

pumps

used in each

methodological

Pumping action and characteristics

division

of MBE.

MBE usage

ION SPUTTER > lo”-1 o*”

Different mechanisms for different species, Gas ionization in ExB field causes implantation, sputtering, burial, gettering, or absorption of gas. Triode pumps normally used. Simple, non-mechanical, reliable and clean. Selective pumping. Poor starting at >10v5mbar. Can have memory effects.

Preferred main primary pump for most MBE systems except where unsuitable, e.g., for GSMBE. “Poisoned” by phosphorous and mercury.

TITANIUM SUBLIMATION 1o-3-1 O-1 1

Titanium sublimed from filaments onto a surface. Ti film reacts with active gases to form low vapour pressure compounds. Higher pumping speeds if Ti film on liquid nitrogen-cooled surface. Simple, high pump speeds easily achievable. Highly selective, does not pump inert or saturated species but compliments ion sputter pump. Pumping action life-time depends on surface area and system pressure.

Used to achieve UHV in most MBE systems. Ineffective in GSMBE systems due to high H, pressures.

CRYOPUMP >l o-3--10“'

Cryocondensation of gas species on series of baffles at -50 and 15 K. Inner 15 K array coated with graphite for cryosorbtion of H,, He, and Ne. High speed pump. Special techniques required to cope with thermal loading during system bake-out,

Concern with deposition of fresh Ti film during MBE growth because of CH, generation by the filament.

Frequently used as primary pump (with TSP) for Si and metals MBE. Secondary pump for all MBE except GSMBE (for safety reasons).

Table 2(b).

Primary UHV “throughput”

pumps used in each methodological

division of MBE.

Pump/Pressure range


MBE usage (mbar)

DIFFUSION

Supersonic jet of high mass oil molecules provides compression of gas towards exhaust. Cheap, simple, reliable. Careful operation needed to preclude oil contamination, but very clean UHV conditions easy to attain and maintain. Pumps all gases, H, and He better than higher mass species,

Has been used, though infrequently, in all areas of MBE. Preferred choice for GSMBE due to high throughput capability, inert nature of oils, and high H, pumping though speed. Ultimate material quality in doubt. Mercury diffusion pumps for mercury-bearing compounds.

Gas compression by impingement of gas molecules onto rapidly moving blades. Fast and effective for system pump-down. Expensive and complex. Pumps all gases. Poor compression for H, and He but compression increases with increasing molecular weight of gas.

Widely used to evacuate FEL’s. Becoming accepted in MBE deposition chambers. Alternative to diffusion pump in GSMBE.

10-3-1

o-11

TURBOMOLECUI.AR 1o-3-_10-l' (but can pump from atmosphere)

Table 3. pumping

Secondary of condensable

Pump/Pressure

UHV pumps,

fitted

to deal

with

process-specific

gas loads

and to provide

highly

effective

vapors.


Liquid nitrogen

Cryocondensation

cryopanelling , (-J5.-10‘11

H,O), and MBE evaporants.

MBE Usage

of certain residual gas specie (notably Surrounding

deposition

region

range (mbar)

Most MBE systems. Al-bearing

Essential

for MBE of

materials to maintain low H,O

provides very high pumping speeds. Used to remove

partial pressure. Used in some Si-MBE

moderate

systems.

thermal

loads in system. Dramatically

enhances

pumping speeds of TSP.

Cryopump

See Table 2(a)

Most MBE applications.

Not suitable for

GSMBE (for safety reasons).

Table 4. Roughing

and Backing

Pumps.

Combinations

of roughing

pumps

pressures at which primary UHV pumps can operate (-1 Os mbar, see text). loads at the outlets of throughput-type primary UHV pumps. Characteristics

are used to evacuate Backing

Pressure

pumps

chambers

MBE uses

Gases pumped

ROTARY PUMP Conventional oil-based

Gas swept by vanes through self sealing chambers. Seals formed by oil. Double stage pumps used. Range of ballast, oil-box bleed, oil and exhaust/mist filters to ensure clean, reliable operation. Cheap, reliable.

atmos-1 0m4mbar

Backing throughput pumps. Need protection by filters and selected oils for GSMBE and P-handling systems. Not used for roughing due to potential for oil contamination. For low gas throughputs, intermittent operation using Ballast volume.

All

ROTARY PUMP Oil free with booster pump

Oil-free pumping. Expensive

atmos-1 0e4mbar

Backing pump for GSMBE System roughing

All

Rotary vane

Simple version of oil-free rotary pump. Cheap. Oil-free

atmos-50mbar

System pre-evacuation

All

Venturi pump

High pressure gas line input pumps via Venturi action. Cheap, reliable (no moving parts). Noisy.

atmos-1 OOmbar

System

All

Sorption

Cryo-sorption by liquid nitrogen-cooled molecular sieve. Cheap, simple. Needs regeneration by heating. Care required to avoid dust transport to vat. system.

atmos-1 0e4mbar

System roughing. Usually preceded by Rotary Vane or Venturi

pump

pumps

Corrosion

resistant

to the

are used to clear the gas

pre-evactiation

All except H,, He, Ne and gases not condensable above 77 K.

Table 5. Common

combinations

of growth-chamber

pumping used in the three

Primary Conventional Solid Source MBE

Gas source MBE

Electron-Beam Evaporator-based

(Most frequent

MBE

methodological

Secondary

III-V’S II-VI IV-VI’S

Ion, TSP. (Diffusion or Turbomolecular evaporated)

CMT

Mercury

(Cd,Hg,,Te)

(TSP)

III-V’S

divisions

of MBE.

Special features

(Cryopump*)

The use of phosphorus necessitates safe pumping via backing line trap. P is pyrophoric therefore safety features.

Extensive LN, cryopanelling

Very high Hg loading needs method of extracting build up of Hg in system.

Diffusion or Turbomolecular. ISP to achieve UHV. (Ion* to maintain UHV during system quiescence).

Extensive LN, cryopanelling

Source materials highly toxic and pyrophoric. Careful safety procedures required. Similar to P above and handling of metal-organics. Output from rotary pumps to be safely vented e.g., to scrubber. Regular maintenance of pumps required.

Si and related materials. Metals.

Ion and/or Cryo. TSP (Diffusion, Turbomolecular)

Water or LN, for removal of heat.

Dominant

High T,superconductor

Diffusion or Turbomolecular

pump listed first. Common

additions

Extensive LN, cryopanelling. if P

Diffusion

(Cryo)

or options in parentheses.

gas load is H,,

For high r, superconductors. high 0, inlet pressures need differential pumping. Care required with pumping of O,-particularly if accumulating on cryopump. Additions

indicated

by *.)

Technology

and Design of MBE Systems 23

24

Molecular

4.4

Ctyopanelling

Beam Epitaxy

Liquid nitrogen condensable

cryopanels

gases, particularly

provide very large pumping

speeds for

H,O and heavier hydrocarbons,

and also,

though less effectively, for C0,.f66) Provision of extensive cryopanelling surrounding the deposition region is an essential secondary pump for achieving high quality with some materials,t’]-t7) e.g., Al-bearing compounds are water sensitive. Cryopanelling is therefore a key design requirement for MBE systems used in the first two methodological areas presented in Table 1. The whole deposition region is surrounded by cryopanelling, with minimal apertures provided for access of components and for substrate transfer. Careful design of the MBE geometry is needed to preclude localized heating of the cryopanel by sources, and of the cryopanel to ensure wide channels for effective liquid nitrogen flow and escape of nitrogen gas. “Hot spots,” caused by trapping of thermally insulating gaseous nitrogen behind the stainless steel (which itself has a poor thermal conductivity), can attain temperatures in excess of 400°C; Mn-doping of GaAs has been observed due to local heating of a cryopanel.t41)f42t This effect can be circumvented by forcing liquid nitrogen flow through the panels. Care is required in construction of the cryopanel, ensuring adequate strength to withstand frequent thermal cycling from 77 K to ambient (and occasional bake-out) temperatures. In applications employing electron beam evaporators (the third area of MBE presented in Table l), the need for liquid-nitrogen cooled cryopanels remains controversial, even if adequately designed to tolerate the significantly higher thermal loads encountered. In Si:MBE, liquid nitrogen cryopanelling is ineffective at collecting the main gas load (H2) and with the most material-damaging gas species (CO and CH4). Furthermore, those residual and outgassed species condensed on the cryopanel can be liberated during process by electron-induced desorbtion stimulated by reflected and secondary electrons from the electron beam evaporators, giving rise to artificially high partial pressures during process.t63)[6fl Another problem with certain materials,

notably Si, is the formation

of loose

powdery deposits on liquid-nitrogen cooled surfaces, which migrate around the vacuum vessel due to poor adhesion. The fine dust thus generated can become charged, again due to electrons in the chamber, and adhere to Si substrates, giving rise to defects.f68) In these applications, water cooled panels are considered preferable by some, though by no means all,

Technology

and Design

of MBE Systems

25

manufacturers and users, since build-up of outgassed species does not occur within the deposition region, and more stable deposits accumulate. The water cooling occurrence adequate

does, however,

need to be very efficient

to preclude

of hot spots on the chamber wall or panel, and thereby ensure gettering

of volatile dopants, as used in MBE of Si and related

materials. The use of closed-loop refrigerators with water/methanol (or glycol) mixture can improve cooling efficiency by proving sub-zero (centigrade) temperatures to the panels. Nevertheless, primary and secondary pumping external to the deposition area must be enhanced to compensate for the loss of the liquid nitrogen surface. 4.5

System

Manufacture

A great deal of care is excercised by MBE practitioners and manufacturers to achieve the clean vacuum conditions necessary for MBE. Chamber materials (stainless steel sheet and tube) are cleaned after each stage of handling and machining, particularly prior to welding stages to preclude inclusions. The completed chambers are thoroughly degreased using solvents and proprietary proceduresf25)t6gt (often based on methods developed originally for the vacuum valve industry) and checked for leaktightness an all welds. Typically, MBE system construction starts with assembly of chambers and pumps onto their support frames, During construction, care is exercised in handling the vacuum components (chambers, cryopanels, pumps, etc.) to avoid contamination from oil, grease, or by human contact. A semi-clean environment is employed, consistent with the need for heavy equipment, e.g., hoists. With the exception of ion gauges and a mass spectrometer,

all ports are fitted with blank flanges,

and after leak check-

ing, the system is baked at 250°C or higher, preferably into an external coupled UHV pump. On achieving satisfactory leak-free vacuum performance, the MBE system pumps are started, and the system is rebaked. The empty vacuum system should achieve specification pressure (the xray limit of commonly used ion gauges of 3 x 10-l’ mbar), and should have a clean residual gas spectrum (dominated by H,, with small CO and H,O peaks, and negligible levels of hydrocarbons). This procedure may need to be repeated if vacuum quality is not sufficiently clean for acceptance of MBE components.

26

Molecular

Beam Epitaxy

The constructional materials for MBE components are treated with even more care than the chambers. All materials and machined components are initially thoroughly

degreased.

Various components

are treated

in a vacuum furnace to further clean and condition their surfaces. bly of the MBE components

Assem-

takes place under clean room conditions.

Prior to loading into the MBE system, the components generally go through vacuum checks and de-gassing. K-cells (without crucibles) are inserted into a dedicated UHV outgassing system, baked to 250°C and operated at their maximum temperature (1400-16OO”C, depending on design) until acceptable vacuum conditions can be achieved. Crucibles are then inserted, and the complete cell re-outgassed. This is often restricted to 1200°C and for a limited time period in the case of pBN crucibles, due to the build up of C on the pBN surface.t45] The substrate manipulators also undergo testing of operation at maximum temperatures and rotational speeds. These components are then introduced into the clean MBE system. The system is once again baked and complete operation of the system is tested. With the exception of the filament assemblies and the source materials, electron beam evaporators operate at near ambient temperatures. Since these sources cannot be tested empty, they are inserted into the MBE system at an early stage of system construction after cleaning to UHV standards, and performing electrical and mechanical checks. Introduction of new components into an operating MBE system can cause disruption of the MBE environment, and degradation of the conditioned state of the system, often leading to degraded material quality. New or repaired components should, if time permits, be thoroughly outgassed either

in a separate,

chamber

dedicated

of the MBE system

UHV chamber, to minimize

or in the preparation

any disruption

of the MBE

for flux generation

in MBE, and

environment.

5.0

MBE COMPONENTS:

SOURCES

A variety of sources are employed

their design depends on the nature of the source materials. A list of sources, used respectively for matrix and dopant flux generation is presented in Tables 7 and 8. These tables describe the main characteristics of each source, their areas of application, and examples of source materials employed.

Table 7. Principal

characteristics

TEMP. OF OPERATION (“C) K-cells

200-1400

MBE

III-V

As K-cell. Cracker -1000

Electron lO% increase in two dimensional densities

electron gas mobilities

and a reduction

from > 700 cme2 to < 70 cm-2.[2701 Failure

arsenic deposits can result in substantially

Table 1. Arsenic

Cracking

for Catalytic Temperature

Re

700

Ta

850 900-950 900-950

W(74%)Re(26%) MO C (graphite)

these

increased epilayer contamination.

Cracking Temperatures Catalyst

in the oval defect to remove

1000-l

100

Cracker Baffles (“C)

132

Molecular

Beam Epitaxy

2.3

Substrate

Purity

In the early 1980’s, GaAs substrates sources of MBE epilayer contaminati0n.f One of the early problems surface,

in which

were identified

as important

6 236][245][374][390][391][452][507l[536][557l

It

was with type conversion

n-type or semi-insulating

of the substrate

GaAs became

p-type at the

surface after annea1ing.t 2s4) This surface-type conversion was first attributed to the indiffusion of carbon, arsenic vacancies, and even gallium vacancies.t2s4)t365) It has since been shown that this type conversion was caused primarily by anomalous manganese outdiffusion from the bulk of the substrate to the crystal sut-face.f 23s)tss7) This diffusion was termed anomalous since the manganese diffused from regions of low (~1 016 cm3) concentration toward the surface, where the manganese concentration increased to well above the bulk value. The principle driving force for this uphill diffusion is probably the indiffusion of vacancies from the surface, which provide energetically favorable, low diffusivity bonding sites for the interstitially-diffusing manganese to occupy. (It can be inferred from a manganese diffusion study that the diffusion coefficient of interstitial manganese at 750°C probably exceeds 1O-lo cm2/sec.)t450) These vacancies increase the solubility of substitutional manganese near the surface and once the manganese impurities occupy substitutional lattice sites, they are relatively immobile. Chromium and iron outdiffusion and surface accumulation was also observed in annealed GaAs substrates~~~~~1~~~1I~~~l~~~~l~~~sl~~~~1 Th ese impurities collected both at the surface of the wafer and at the substrate/epilayer interface, forming thin, highly doped (>>l 017 cm9) layers, often resulting in type conversion of this surface layer. This impurity outdiffusion may also be assisted by electric fields near the surface.f3s2)t450t It resulted in undesirable doping barriers, difficulty in controlling doping profiles, problems controlling planar doped barrier heights, excessive back- and side-gating quantum well luminescence.t2s7)f3so)t3s1)t406) outdiffusion

problem involved annealing

of FETs, and poor, single Early solutions for this

the wafers for 24 hours in purified

hydrogen

at 750°C and etching away the first - 20 pm from the surface prior to epitaxial gr0wth.f 2s7)[390)[392)While this heat/polish cycle reduced the outdiffusion of Mn, it did not prevent chromium outdiffusion.f17*) Today, the manganese outdiffusion problem is much less severe due to the widespread adoption of boule annealing by wafer manufacturers. The problem of chromium outdiffusion has been eliminated through a combination of reducing the background carbon concentration and precise control of the stoichiometry during the growth of bulk GaAs to produce

MBE of High-Guality

GaAs and AlGaAs

133

undoped semi-insulating (1 07-1 0s Q-cm) GaAs. The growth stoichiometry controls the EL2 deep-donor density, which compensates the remaining background

carbon

acceptor

and surface contamination buffer layers

concentration.

However,

substrate

are still significant

problems,

since elaborate

(often involving

high two-dimensional

superlattices)

are still required

electron gas mobilitiesf’*~

purity

to achieve

and low threshold

GaAs/

AlGaAs lasers.f14g)t173) Results as recent as 1990 indicate that the surface quality and cleanliness of as-delivered GaAs substrates are still inadequate, resulting in unacceptably large substrate-induced oval defects.f6s) Lower oval defect densities are obtained on both chemically-cleaned GaAs substrates and on as-delivered silicon substrates. In addition to anomalous impurity outdiffusion from the substrates themselves, it was recognized that the mounting procedures introduced contamination. Early MBE systems utilized indium solder to mount the GaAs wafers onto molybdenum blocks, Problems associated with the trapping of impurities in the indium solder and auto-doping with indium have long been suspected,t174)t205)f3sg] and have been observed on at least one occasion.f1s8) One of the first approaches to reducing the impurities emanating from indium was to use gallium for wafer mounting (gallium was available at higher purity). ~77) The use of gallium solder also eliminated the need to polish the backside of the wafer to remove the solder and associated stress in the crystal. However, gallium quickly erodes the molybdenum sample holders at typical MBE growth temperatures,(s6)f17r) so that it is necessary to use tantalum sample holders. In recent years, there has been a strong tendency to use direct The use of direct radiative heating for substrates.1 ~s~l~*s~1~*ss1~~~~l~~~~lI~s~l radiative substrate tamination

heating eliminated

the problems

from liquid metal solders and eliminated

associated

with con-

the roughness

of the

backside of the wafer. The primary difficulty in using radiative substrate heating has been accurate determination of the substrate temperature, but this can be solved with the use of infrared transmission spectroscopy.[‘~7w7w~i D’erect radiative substrate heating has been used to obtain GaAs with peak electron

mobilities

in the range of 216,000

cm*/

Vsecf265) to 402,000 cm*/ Vsec.fgl] 2.4

Source Purity

The purity of source materials has been an important problem since MBE was first developed. In very early MBE systems, much of the source contamination occurred when the MBE system was vented to load sub-

134

Molecular

&rates,

Beam Epitaxy

but this problem was eliminated

load-locks, as discussed previously. A number of different arsenic been used for molecular

with the introduction

species

and source

beam epitaxy, including

of vacuum

materials

have

arsine (ASH,), tetrameric

arsenic (AQ), dimeric arsenic (As*) obtained from the thermal decomposition of GaAs, and dimeric arsenic obtained by thermal cracking of tetrameric arsenic. The relative purity of these different arsenic sources has changed dramatically in the past decade. In 1981, arsine was available at higher purity than solid arsenic, resulting in the growth of undoped n-type GaAs with a free electron concentration of 2.4 x 10J4 cm”, a 77 K mobility of 110,000 cm2/Vsec and a peak electron mobility of 133,000 cm21 Vsec at 55 KP) At about the same time, solid arsenic resulted in higher purity GaAs than dimeric arsenic produced by the thermal decomposition of GaAsP) It is no longer clear that thermal decomposition results in impure As, sublimation (neglecting the co-sublimation of gallium). Since about 1983, the quality of solid arsenic (As4) sources has improved steadily. In 1983, MBE GaAs reached 77 K electron mobilities of 126,000 cm2/ Vsec[205) and 144,000 cm2/ Vsec[1781 and 146,000 cm2/ Vsec.f367) Not only has the purity improved, but fabrication of solid arsenic slugs, designed to fit the p-BN crucibles, greatly reduces the exposed surface area of the arsenic during source loading and reduces the level of impurities carried into the system. The solid slugs also increased the amount of arsenic loaded, thus allowing the completion of a much larger number of growths before breaking vacuum to reload sources. In 1986, improvements in the arsenic purity contributed to the achievement of MBE GaAs with electron mobilities of 200,000 cm2/ Vsec at 55 Kf353) and 216,000 cm2/Vsec at 46 K (163,000 cm2/ Vsec at 77 ~~~~~~~~~~~~ In 1988, highpurity GaAs was obtained with a free electron concentration of - 3 x 1 013 cm3 and a peak electron mobility of - 300,000 cm2/ Vsec at 60 K.f10~f106)* This material showed no trace of electron traps above a concentration of 5 *The behavior of this material was quite unusual, indicating a shallow donor energy of -1.6 meV and a low temperature mobility which varied as -PO7 instead of -T3~.[1071t108]The unusual low temperature mobility behavior was attributed to differences in screening, since the electron freeze-out occurred at lower temperatures than expected.1lo81It was speculated that the shallow donor was due to carbon,flOflfloe] but the lack of previously observed carbon donors does not support this conclusion. Very early studies occasionally showed low donor activation energies,fs11f5**]but the decreased donor activation energies in those samples was the result of Donor band formation should impurity banding of the excited donor levels. [10*1[*sa1[3861[4esl[s*~l not occur in these more recent samples, since the donor concentrations are much lower. It is suggested that these anomalously low activation energies are impact excitation of the filled shallow donors to the 2s states by the free electrons in the conduction band.f1e51In this scheme, electrons in the excited 2s and 2p states are thermally ionized to the conduction band. While impact-thermal ionization is consistent with the observed data, these apparent ultra-shallow donors require further study.

MBE of High-Quality

GaAs and AlGaAs

135

x lOlo cm3.t107) In 1987, a 77 K electron mobility of 205,600 cm*/ Vsec and a peak mobility of 294,700 cm*/Vsec at 42 K were obtained for silicon doped MBE GaAs, with N, - N, = 3 x 1013 cm-3.p1) (it was estimated that undoped GaAs grown in this system had N, - Nd < 2 x 1013 cm-“). Until quite recently, contamination (particularly sulfur and carbon) of even the best available

solid arsenic

resulted

in increased

unintentional

impurity concentrations when the arsenic beam was thermally cracked into As2.tg11t3531t4731The increased contamination levels probably resulted from cracking of the impurity-containing molecular species emanating from the arsenic source, increasing impurities, Recent improvements

the incorporation efficiency in both the arsenic purity

of these and the

cracker design permit efficient cracking at reduced temperatures, resulting in the growth of GaAs with mobilities as high as 220,000 cm*/Vsec at 77 K. The newest results with cracked arsenic have set new records for GaAs purity with electron mobilities of 402,000 cm*/ Vsec at 28-40 K (Nd - N, = 2.8 x 1013 cm-7.1 473] The principal background donor in this material was silicon. The exact source of the silicon was not identified, but the silicon level decreased as the cracker temperature was decreased, reaching a maximum mobility with a cracker temperature of 650°C. It is believed that the primary reason for the improvement over earlier high purity MBE layers was an increase in the purity of the arsenic source material. In some of the highest purity GaAs layers, there was no trace of the common MBE deep electron traps Ml, M3, and M4, leading to the supposition that these traps are caused by impurities from the arsenic source.f701f71] M2 was observed and it was speculated that it could be caused by a silicon-defect complex. The group III source materials have also created purity-related problems. In particular, the aluminum source purity is critical for obtaining high quality AlGaAs, particularly at lower growth temperatures (< 680%). The use of zone-refined aluminum resulted in the growth of two-dimensional electron gas structures

with low-temperature

mobilities

of 2-3 x 1O6

cm*/ Vsec.f1411t166] A very high-purity aluminum source and decreasing the vacuum pressure to the lo-‘* torr range are both essential for achieving ultra-high two-dimensional electron gas mobilities.t410) The purity of the aluminum and indium sources is critical for the growth of AlGaAs, InGaAs and InAlAs with high luminescence efficiencies.* * Since aluminum reacts with p-BN and can cause crucible delamination, the quality of the aluminum melt may depend as much on the purity and quality of the p-BN crucible as on the purity ofthe aluminum source material. Most p-BN crucibles have an impurity specification of only 4 00 ppm for metallic impurities and specification of the oxygen contamination level is not provided.

136

Molecular

Beam Epitaxy

Early gallium sources had problems with oxygen contamination. In 1981, it was demonstrated that Gas0 impurities originating in the gallium source could be reduced by two orders of magnitude aluminum

into the gallium

forms preferentially

by introducing

s0urce.t 232) This reflects

0.1%

the fact that AI,0

to Ga,O in the molten source and that Ga,O

has a

much higher vapor pressure than AI,O. Since the vapor pressure of aluminum is about an order of magnitude smaller than that of gallium at these temperatures, it would be expected that aluminum would be incorporated at < lo-l9 cme3. Since aluminum is isoelectronic with gallium and since the GaAs/AIAs system is closely lattice matched, this probably does not have a strong effect on the properties of the GaAs. Subsequent improvements in gallium purity and source preparation have reduced the problems associated with oxygen contamination. In particular, gallium is now shipped as high purity ingots and they are not loaded if they have melted at any point. Etching of these ingots in a chilled HCI:methanol solution

and limiting

their

exposure

to the atmosphere

during

source

loading has reduced the surface contamination of the gallium.t265)t388) (Surface contamination is occasionally observed as a dull, lusterless surface. The origin and nature of this contamination has not been determined.) Since oxygen has a low solubility limit in GaAs, contaminated gallium can result in GaAs with reasonable luminescence efficiencies, but the resulting AlGaAs will be terrible. (Sufficiently contaminated gallium can result in semi-insulating AlGaAs, even at moderate doping levels of 1016 cm3. Contaminated gallium may also result in poor InGaAs.) In general, it is essential to minimize the exposure of the source materials to atmosphere during source loading.t’44)t265]t388] The solid arsenic

source

is particularly

since arsenic oxidizes

sensitive

to contamination

rapidly and cannot be outgassed

during

loading

at high tempera-

tures. Arsenic oxides have been observed in the mass spectra during the evaporation of air-exposed arsenic.te6] The introduction of hydrogen during MBE growth has been shown to reduce the total ionized-impurity incorporation, N, + N,, and increase the carrier mobi~ity.[5~1[601~6~1[*~*1[3971[4~~1 The presence of hydrogen during growth also reduced the deep-level

concentrations.t511t3g7)

This impurity

and deep-level reduction is consistent with results showing that shallow acceptors and donors, EL2, and other deep levels, can all be passivated with a hydrogen p1asma.t 405) However, the reduced impurity incorporation is real, since no thermal reactivation of these impurity levels could be induced.f401) A more detailed discussion growth of GaAs and AlGaAs is presented

of the role of hydrogen in Sec. 8.3: Hydrogen.

during the

MBE of High-Quality

3.0

GROWTH

137

PROCESSES

This section describes

the microscopic

growth of GaAs and AlGaAs. semiconductors is performed growth

GaAs and AlGaAs

rates, when

compared

processes

involved

in MBE

Molecular beam epitaxial growth of Ill-V at relatively low temperatures and slow to older epitaxial

techniques

(LPE and

VPE). Low growth temperatures result in growth mechanisms which are strongly dominated by surface kinetics and surface chemistry. The surface kinetics and surface chemistry are also strongly affected by the substrate orientation, the choice of arsenic species (e.g., ASH,, As, or As,), the growth rate and growth interruptions, The kinetics and chemistry of the growth directly affect both impurity incorporation and defect generation. These growth processes also play an important role in the generation of defects and the suppression of surface and interface roughness. The process of MBE growth can be divided into two stages. The first stage relates to the surface processes of molecular chemisorption, surface migration, and subsequent bond formation, The second stage consists of a thermodynamic interaction and reordering to form the bulk crystal, which occurs in the near surface region (the top 3-4 atomic layers). The source molecular beams are usually obtained by evaporation or sublimation of Some approaches to MBE use one or more solid source materials. gaseous sources. In either case, the molecular species are introduced at very low pressures, allowing them to traverse the distance to the substrate without interacting with each other. Some portion of the molecular species reaching the substrate becomes weakly chemisorbed to the surface. Once the molecules are chemisorbed on the surface, they can: (ij migrate to energetically favorable lattice sites and bond there; (ii) gather at surface contamination

to form oval defects;

(ii0 form agglomerations

with similar

species (such as gallium droplet formation); or (iv,) migrate around the surface and eventually desorb. In order to grow acceptable material, growth conditions should be adjusted so that the first process is the most favored one. Following the chemisorption of the group Ill and V molecular species,

epitaxial

growth

occurs via surface

migration

to step edges or

nucleated islands, where chemical bonding occurs.te2] A thermodynamic interaction and redistribution of the near surface crystal layers establishes the final configuration. This last process has recently been observed by Auger electron spectroscopy and x-ray photoemission spectroscopy ing the growth of InGaAs and AlGaAs ternary alloys.t1gej[345j

dur-

138

Molecular

3.1

Atomic

Beam Epitaxy

Surface

A discussion

Reconstructions

of the surface processes

during MBE growth must be

preceded by an introduction to the surface structure of the substrate crystal. In order to minimize the energy of the near-surface region of the crystal, the group III and group V atoms arrange themselves in a regular fashion which exhibits long range order. Each ordered arrangement of the near surface region is known as a surface reconstruction. (An unreconstructed surface is one in which the near surface layers are arranged in exactly the same configuration as the underlying bulk layers, with only slight distortion of the bond lengths.) Surface reconstruction can also be quite different than the bulk crystal and often reflects the stoichiometry of the growth process and influences the growth mechanisms. The surface phase diagram of GaAs, shown in Figure 5, was obtained from RHEED under the following growth conditions: (0.71 pm/hr) for 320°C 5 T 2; 780°C and AsdGa beam equivalent pressure (BEP) ratios of 0.9-l 00.t108a] A similar, though less complete, surface phase diagram for AI,Ga,,As is shown in Fig. 6 under growth conditions of 400°C 5 T s 750°C BEP ratios of 1.5-35, and where the low temperature growth rate and composition (below the Ga desorption temperature) were 0.36 pm/hr and x = 0.35, respective1y.t’ OrJa) Surface reconstructions are commonly observed with reflection high energy electron diffraction @HEED), low energy electron diffraction (LEED) , and scanning tunneling microscopy (STM). The surface reconstructions discussed in the remainder of this section are those observed by scanning tunneling microscopy. All of the STM images were obtained in-situ, after the surface layers were grown by MBE and quenched.t381t3g)t40]t41) There are several static reconstructions observed by STM which are not observed in the dynamic RHEED studies shown in Figs. 5 and 6 (i.e., there is no ~(2x8) or ~(8x2) surface). (A discussion of the notations of surface reconstruction can be found in Ref. 546.) The (100) surface reconstruction changes from ~(4x4) to c(2x8)/ (2x4) to (2x6) to ~(8x2) as the surface stoichiometry changes from arsenicrich to gallium-rich.

STM images of ~(4x4) reconstructed

GaAs surfaces

are shown in Figs. 7(a) and (b).t3e)t391t401 The surface unit cell is superimposed on Fig. 7(a), while Fig. 7(6) is a magnified image showing ordered surface dimers (probably arsenic). Figure 7(c) shows a probable model for the near-surface region. This model is consistent with a charge neutral surface as well as the STM image of this ~(4x4) reconstruction. Other

MBE of High-Quality

GaAs and AlGaAs

139

reconstructions are possible, such as one similar to that shown in Fig. 7(c), but with the four central, underlying arsenic atoms of each unit cell replaced by gallium atoms.t40) The three surface dimers are aligned along the [l lo] direction,

perpendicular

~(2x8) reconstructed structions

surface.

to the arsenic

dimers observed

All of the observed

~(4x4) surface

on the recon-

possess unit cells with three dimers.

facetting

1 1 I 1 I I 0.9 1.0 1.1 1.2 1.3 1.4 1.5 1.G 103/Ts(K-'1 Figure 5. Surface phase diagram of GaAs (001) for MBE growth on (001) 2’ misoriented toward L. DS weritz.)

(1 il)As,

growth rate 0.7 MLs-’ for T, c 630X.*

(Courtesy

of

* Figures 5,6 refer to (001) while we have used (100) throughoutthe chapter. Orientations (001) and (100) are interchangable.

140

Molecular

Beam Epitaxy

T,(OC) 800 700 600 500 II,, 1 I I I

400 I

I

looF GaAs10011 - 14x21 (3x1) lZx&) r I I

roughening

I

I

I

0.9 1.0 1.1 1.2 1.3 1.4 1.5 1.6 d/T,(K-'1 Figure 6. Surface phase diagram of AI,Ga,,As (001) for MBE growth on (001) 2” misoriented toward (11 l)As, growth rate 0.35 MLs-’ for T, < 630°C. The corresponding main phase boundaries for GaAs(OO1) are also shown.* (Courtesy of L. DS weritz.)

l Figures 5,6 refer to (001) while we have used (100) throughout the chapter. Orientations (001) and (100) are interchangable.

MBE of High-Quality

GaAs and AIGaAs

141

Figure 7. (a) The GaAs (100) surface with a c(4x4) reconstruction, imaged by scanning tunneling microscopy; and (b) at higher resolution. (c) The probable atomic configuration of this c(4x4) reconstruction. (Courtesy of D. K. Biegelsen and R. Bringans.)

Figures 8(a) and (b) are STM images of the c(2x8)/(2x4) surface reconstructions of (100) GaAS.[38]-[40]Figure 8 (a) shows coexisting regions with c(2x8) and (2x4) reconstructions (unit cells superimposed) .Figure 8(b) shows the c(2x8)/(2x4) reconstructed surface at higher magnification, indicating the presence of the surface dimers. The energetically favorable c(2x8) reconstruction which corresponds to the STM images of Figs. 8(a & b) is depicted in Fig. 8(c). The (2x4) surface subcells which comprise both

142

Molecular

Beam Epitaxy

the ~(2x8) and (2x4) reconstructions dimers (probably arsenic). is enhanced surface.

by annealing,

possess either two or three adjacent

The presence of unit cells with only two dimers consistent

with the loss of arsenic

from the

The two arsenic dimers in these unit cells are always adjacent to

each other. Figure 9(a) is an STM image of the (2x6) surface reconstruction of (100) GaA~.f~~jf~~) F’rgure 9(b) shows one possible model of the near surface region. This model is consistent with a charge neutral surface as well as the STM image of this (2x6) reconstruction. The surface unit cell appears to have two arsenic dimers with four dimers missing. The incomplete ordering of the (2x6) reconstructed surface suggests that several different surface structures are present. STM studies suggest that (4x6) surfaces observed with LEED are actually due to a superposition of the electron diffraction from coexisting regions of (4x1) and (2x6) symmetry. An STM image of a ~(8x2) reconstructed (100) GaAs surface is shown in Fig. 10(a).f3gjf40) The surface unit cell is superimposed on this image, along with representations of the ordered surface dimers (probably gallium). Figure 10(b) shows the probable model of the near surface region, which is consistent with a charge neutral surface as well as the Each (4x2) subcell has two adjacent STM image of this reconstruction. gallium dimers and two dimer vacancies. Gallium-rich surfaces, such as the ~(2x8) GaAs surface, are difficult to image with the STM because of noise and possibly due to gallium transfer to the tip of the microscope. Scanning tunneling microscopy has also been used to study the --surface reconstructions of (1 1 1) and (110) GaAs. Two different surface --reconstructions of ( 1 1 1) GaAs have been observed with scanning tunneling microscopy.f38)t41j Figure 11 (a) shows an STM image of the arsenicrich (2x2) reconstructed surface. A model for this (2x2) surface is given in Fig. 11(b). This (2x2) reconstruction is characterized by the presence of tightly bonded trimers (probably arsenic).

(The bond length of the arsenic

trimer is shorter than that observed for bulk arsenic.) After annealing, this surface reconstruction changes rapidly to a (Jm x Jm) reconstructed surface. STM images of this (Jm x Jm) surface are shown in Figs. 12 (a) and (b). Figure 12(a) shows the long range order of the surface, which is characterized by terraces with triangular, bilayer-deep (3.3 A) holes. Figure 12(b) shows the (Jm x Jrs) reconstructed surface at higher magnification, with representations of the arsenic atoms superimposed on the image. This reconstruction is characterized by raised hexagonal rings. Figure 12(c) shows a model of the (JKJ

x Jm)

reconstructed

This (Jm x Jl9’) surface is the only known GaAs reconstruction violates the surface autocompensation assumption.

surface. which

MBE of High-Ouality

GaAs and AIGaAs

143

IGaA5(100)c(2xaij

..~ ~ !

~

~. i

!$!l!

~

(Side

.~

...!I!I!

~ !

..

i

..

~ ...

View)

Figure 8. (a) The GaAs (100} surface with a (2x4} or c(2x8} reconstruction, imaged by scanning tunneling microscopy; and (b) at higher resolution. (c) The probable atomic configuration of this surface reconstruction. (Courtesy of D. K. Biegelsen and R. Bringans.)

144

Molecular

Beam Epitaxy

(Side View)

Figure 9. (a) The GaAs (100} suriace with a (2x6} reconstruction, imaged by scanning tunneling microscopy. (b) The probable atomic configuration of this suriace reconstruction. (Courtesy of D. K. Biegelsen and R. Bringans.)

Figure scanning surface

10.

(a) The GaAs

tunneling reconstruction.

(100) surface

microscopy. (Courtesy

with a c(8x2)

(b) The

probable

of D. K. Biegelsen

reconstruction, atomic

imaged

configuration

and R. Bringans.)

by

of this

MBE of High-Ouality

GaAs and AIGaAs

145

Figure 11. (a) The GaAs (111) surface with a (2x2) reconstruction, imaged by scanning tunneling microscopy. (b) A probable atomic configuration corresponding to this surface reconstruction. (Courtesy of D. K. Biegelsen and R. Bringans.)

Figure 12. (a) The GaAs (1TT) surface with a (.J'f9 x .J'f9) reconstruction, imaged by scanning tunneling microscopy; and (b) at higher resolution. (c) A tentative atomic configuration corresponding to this surface reconstruction. (Courtesy of D. K. Biegelsen and R. Bringans.)

146

Molecular

A scanning

Beam Epitaxy

tunneling

microscopy

study of vacuum-cleaved

(110)

GaAs surfaces shows that the (110) surface is unreconstructed.t133j The results of this study suggest that the surface buckles, shifting the arsenic away from the bulk. At the same time, the gallium moves toward the bulk. The bond between the surface gallium and surface arsenic forms a 29”31’ angle with the plane of the surface. The upper and lower limits on this buckling angle are believed to be 23” and 34.8”, respectively. 3.2

Surface Chemisorption

The first step of the epitaxial growth process is chemisorption of the molecular species on the surface of the substrate. Chemisorption is generally a simple process for atomic species. For molecular species, however, chemisorption is often far more complex, involving ing and surface interactions between multiple molecules.

bond breakThe sticking

coefficients of elemental group III sources are assumed to be unity at low temperatures, where desorption is negligible.f221j Fortunately, aluminum desorption is negligible for temperatures up to 1000°C and gallium desorption only becomes significant at temperatures above - 620°C. (On the other hand, indium desorption becomes important at much lower temperatures T > 500°C). The chemisorption of arsenic on the crystal surface depends strongly on the molecular arsenic species employed. Dimeric arsenic adsorbs on gallium arsenide surfaces in a simple fashion, probably as weakly chemisorbed As, precursors, which dissociate into arsenic atoms bound to the group V sub1attice.t *l jf 1431 The sticking coefficient of dimeric arsenic tends toward unity in the presence of a sufficiently large population of chemisorbed gallium. At temperatures below 330°C an associative reaction occurs between these As, precursors, forming A~,.t’~~jt**‘j At temperatures above 180°C tetrameric arsenic adsorbs on gallium arsenide via a dissociative chemisorption process.t142j[221] This process proceeds by the interaction of two As, molecules, resulting in the formation of two chemisorbed As, molecules and the desorption of an As, molecule.

The arsenic surface coverage

than for a comparable

is probably

As, flux, since pairwise

higher for an As, flux

dissociation

of tetrameric

arsenic is believed to require two adjacent gallium atoms exposed

on the

surface, while dimeric arsenic dissociates at a single exposed gallium atom. The sticking coefficient of tetrameric arsenic increases to 0.5 in the presence of a sufficiently large surface population of chemisorbed gallium. The sticking coefficient is independent of temperature from 180°C to 330°C.

Desorption

of chemisorbed

arsenic

is very

rapid for arsenic-

MBE of High-Quality

terminated

surfaces,

allowing

GaAs and AlGaAs

147

the use of a wide range of excess arsenic

fluxes while maintaining good stoichiometry and crystal quality. Thus for conventional growth, the growth rate is controlled by the group III flux. 3.3

Incorporation

of Chemisorbed

Species:

Island Formation

and

Step Propagation The incorporation and growth mechanisms have been studied primarily by reflection high energy electron diffraction (RHEED).* In RHEED measurements, a high energy electron beam is reflected from the crystal surface at a very shallow angle (- 1”). The reflected electrons are diffracted by the atoms at the surface and imaged on a phosphor screen. The shallow angle of incidence makes the electron beam very sensitive to long range order in the reconstruction of the crystal surface. The coherent diffraction of these electrons results in the formation of a series of streaks on the phosphor streaks is directly on the surface. investigated as a RHEED growth.

screen. The number, intensity, and spacing of these related to the structure of the two-dimensional unit cell The structure of the surface reconstruction has been function of growth conditions. Dynamic variations in the

pattern have also been used to explore the kinetics of MBE Since the diffracted electron beam samples a relatively large

surface area (compared to the size of a single surface unit cell), RHEED studies provide information about the configuration of the average unit cell of the surface reconstruction. RHEED also provides a means for observing long-range correlated surface processes, such as surface step propagation and adatom migration lengths, Some of these RHEED studies, particularly the dynamic studies, are discussed in detail in this section. Reflection high energy electron diffraction (RHEED) studies on (100) GaAs surfaces show the existence which

may contribute

to the formation

of staircase

step arrays,t418)f515)

of the characteristic

RHEED

streaks.f41q These step arrays occur even on surfaces misoriented by as little as 6.5 mrad (0.377.f 418It515] Temporal variation of the average terrace width causes oscillations

in the RHEED intensity during growth.f516)

Figure 13 shows a typical RHEED intensity oscillation and recovery measurement. The oscillation results from destructive interference of the diffracted electron beam from different steps, so the oscillations are strongest for electron beams incident midway between Bragg angles. One

* There is an extensive discussion of RHEED for dynamic studies of film growth in Ch. 8 (by P. Cohen) of this book.

148

Molecular

Beam Epitaxy

complete

RHEED

complete

monolayer

oscillations tion through

oscillation

period

corresponds

to the growth

of crystal.[ 370)t516) The existence

and the absence

of spots (from transmission

three-dimensional

growth

nucleation)

tional MBE growth proceeds as a layer-by-layer,

of one

of these RHEED electron

indicate

diffrac-

that conven-

two-dimensional

process.

It has been suggested that the damping of the RHEED oscillation intensity is caused by premature nucleation of a second layer before completion of the first layer.t370) The damping of the oscillation intensity occurs because of the phase mismatch caused by the nucleation of multiple layers. Under some growth conditions, RHEED oscillation damping can also indicate the formation of a steady-state terrace distribution, which encourages a transition from island growth to step propagation.[222t The island formation and step propagation models are shown schematically in Fig. 14. While these damping mechanisms may also be significant, recent work has shown that macroscopic growth-rate variations across the area sampled by the RHEED beam are one of the principal causes of long term decay and beating of RHEED oscillations.fs14] These results suggest that RHEED oscillation damping rates should be interpreted very carefully. It is possible to obtain RHEED oscillations which continue indefinitely, suggesting that there is an optimal set of growth conditions for maintaining the planar nature of the two-dimensional island growth mechanism. The planar growth process proceeds via the formation of twodimensional islands when the group III migration length is shorter than the distance between step edges.f 36g) When the group III migration length is longer than the terrace width, this surface growth process proceeds via a smooth, two-dimensional propagation of atomically high steps or terraces across the crystal surface.

Since the group III diffusion

length increases

with higher substrate temperatures, there exists a critical growth temperature, below which growth proceeds via island formation and above which RHEED oscillations disappear growth proceeds via step propagation. when the growth process changes from island formation to step propagation. The group III diffusion length also increases for reduced group III fluxes and decreases if dimeric arsenic is used in place of tetrameric arsenic

(for equal atomic fluxes).t 36gl For ternary

alloy semiconductors,

this critical temperature is larger than the average of the constituent binary speciest3581 and can create problems of compatible temperature regions for the growth of different alloys.

MBE of High-Quality

149

GaAs and AlGaAs

10

8

I

0 0

60

I

I

I

I

I

I

I,

120

Time Figure 13.

I

I

I

I

180

I

I

I

I

I

240

(set)

A typical RHEED intensity oscillation and recovery measurement. Each oscillation corresponds to the growth of a single atomic layer. The RHEED intensity recovers exponentially when the group Ill fluxes are interrupted. RHEED intensity measurements are used extensively to calibrate growth rates. RHEED intensity measurements have also been used to investigate various aspects of the microscopic growth processes, though agreement on the interpretation of such (Courtesy of J. P. A. van der Wag?, K. L. measurements is far from universal. Bather, and G. S. Solomon.)

150

Molecular

Beam Epitaxy

electron beam

b)

\

Figure 14. Schematic representations of the island formation and lattice steppropagation models often assumed to explain the behavior of RHEED intensity oscillations. (a) In the island formation process, the adatoms form islands on the terraces because they are unable to diffuse to a lattice step edge. RHEED oscillations are attributed to time-dependent electron interference effects caused (b) In the stepby the periodic formation and annhilation of these islands. propagation process, the adatoms have enough energy to diffuse to a lattice step edge, resulting in a smooth drift of the step edges. The lack of time-dependent electron interference effects results in the absence of RHEED intensity oscillations. While these models have proven very useful for understanding MBE growth processes on favorably-oriented substrates, important incorporation and redistribution processes are occasionally overlooked or disregarded in favor of simple explanations based on these models.

The degree of long-range order of the GaAs surface increases dramatically with growth interruption. A RHEED study suggested that the long-range surface order of the (100) GaAs surface in the [l TO] direction increases from 50 nm to 300 nm, after the growth is interrupted.[515al The RHEED intensity recovers through a process which is characterized by a short and a long exponential transient.[ 2791[3701The lengths of the time constants of the two exponential processes depend on where in the growth cycle the growth interruption occurred.[ 2791 The relative importance of the

MBE of High-Quality

two exponential

surface-ordering

tion point.

The time constant

monolayer

coverage

versely,

processes

151

also depends on the interrup-

for the fast process

and shortest

the time constant

GaAs and AlGaAs

is longest for integral

for half monolayer

for the slow process

coverage.

is shortest

monolayer coverage and longest for half monolayer recovery process dominates for integral monolayer

coverage. coverage,

Con-

for integral The fast while the

slow recovery process is most important for half monolayer coverage. It is tempting to speculate on the nature of these two recovery processes as well as their dependence on the monolayer coverage. The thermal activation

energy of the fast transient

is 2.3 & 0.2 eV, indicating

that this process involves the breaking of bonds and not just surface migration of chemisorbed moleculest 370) The fast process may just be the growth of large islands and/or step edges at the expense of smaller islands or at the expense of islands created by second layer nucleation. In this case, the average diffusion distance required for this process would be shortest for half monolayer coverage and the resulting time constant should be smallest for this coverage. The slow process could be due to the thermodynamically-driven interaction between the near-surface layers. The mechanics of such an interaction could be due to vacancy diffusion or to spontaneous step migration, with an effective edge segregation process. Such a spontaneous step migration process could explain the increased order observed by RHEED after growth interruption, the surface enrichment observed during the growth of InGaAs and AlGaAs, and the reduced impurity incorporation observed for growth on (100) surfaces misoriented toward (11 l)A. However, these processes are still speculative and are included to suggest possible models and interesting research directions

to further elucidate the growth mechanisms.

Recent in-situ scanning

tunneling

microscopy

measurements

show

that the step propagation

proceeds via the incorporation of complete (2x4) unit cells at the step edges.1 404) The step edges have less disorder when

the (100) substrate is misoriented toward the (11 l)A surface than when the substrate is misoriented toward the (11 l)B sur-face.t225)[404j Background impurities

also affect the growth kinetics by providing

undesirable

nucle-

ation centers in the middle of the terraces. These impurity-related nucleation centers prevent smooth step propagation, causing surface and interface roughness.vO) Misorientation of the (100) GaAs substrate 3” towards the (11 l)A surface reduces the background impurity incorporation rate by a factor of four. The precise reason for this background contamination reduction

is not immediately

clear, but may result from the lower kink

152

Molecular

Beam Epitaxy

density in the terrace edges of the surfaces misoriented toward the (11 l)A surface.t404] The improved surface order may present fewer suitable bonding sites for impurities, tion.W21PWts211

resulting

in increased

impurity

desorp-

Growth at high temperatures and at a low arsenic-to-gallium ratio results in three-dimensional growth, consistent with the agglomeration of excess gallium on the surface.1 3711 To ensure planar growth, either the amount of gallium applied to the surface needs to be limited (as in migration-enhanced epitaxy) or sufficient arsenic must be supplied to prevent gallium aggregation. (Gallium aggregation also results formation of oval defects, as discussed in Sec. 5: Oval Defects). 3.4

in the

Surface Diffusion A number of gallium surface diffusion

studies have been published,

but only a small amount of surface diffusion work has been done on aluminum and almost none on arsenic. The distance between group III droplets on the surface can be used to determine the migration length of the group III species on both gallium-terminated and arsenic-terminated GaAs surfaces.t213)t547) The gallium and aluminum migration lengths on a group III terminated surface increase with temperature, reaching -0.5-10 pm at 610°C. The gallium migration length on an arsenic terminated surface is as much as three orders of magnitude smaller (- 10 nm) than that on a gallium terminated surface,f 213)t225)consistent with observed transitions between island growth and step propagation.t36g) The diffusion of gallium across the surface of (100) GaAs is anisotropit. This anisotropic surface diffusion is responsible for the shape and orientation of oval defects, as well as for the rippled surface structure observed on AlGaAs for growth conditions resulting in anomalously poor surface morphology. (Cathodoluminescence imaging also recently revealed the presence of ripples in the thickness of GaAs/AIGaAs quantum wells, oriented parallel to the [Oi l] direction.t420a) The anisotropy of the gallium surface mobility has been studied by RHEED oscillation measurements on (100) GaAs substrates misoriented toward the [l lo] and [Oi I] directions.t383) The gallium surface diffusion coefficient is four times as large in the [Oi l] direction as it is in the [l lo] direction. Hence the gallium surface diffusion length is twice as large in the [Oi l] direction as it is in the [l lo] direction. Anisotropic surface diffusion of gallium has also been measured by observing changes in the growth rate at edges near the

MBE of High-Quality

GaAs and AlGaAs

153

intersections of the (100) and the (11 l)A and (11 l)B surfaces.t171) At 560”, the estimated gallium diffusion lengths are - 1 pm in the [l lo] direction and - 8 pm in the [Oi l] direction, The anisotropy of the gallium diffusion from the edge/growth rate measurements agrees qualitatively with the RHEED measurements. However, there is disagreement about the magnitude of the anisotropy, indicating that further work is necessary. On the (100) surface, the activation energy of the gallium surface diffusion coefficient is 2.8 eV for both the [l lo] and [Oi l] directions.t3s3] The maximum growth rate at which good crystal quality can be achieved increases with increasing substrate temperature, due to an increase in the gallium surface diffusion rate. The activation energy obtained from the temperature-dependent measurement of the maximum growth rate is 2.5 eV,t326t which agrees well with the RHEED measurements. The gallium surface diffusion coefficient should depend on the aluminum composition of the epilayer. The activation energy for gallium surface diffusion may also depend on the As/Ga flux ratio and the specific arsenic species used.t3s3) The migration of aluminum on an arsenic-terminated surface is an order of magnitude smaller than that of gallium (- l-3 nm).t213) The diffusion of aluminum on GaAs appears to be comparable to that of gallium on GaAs, but the diffusion length of aluminum on AlAs is much smaller than that of aluminum on G~As.[~~~) About all that is known about the surface diffusion of arsenic is that the activation energy for As, surface diffusion is - 0.24 eV.t143] 3.5

lncorporatlon

of Chemisorbed

Species:

Surface

Incorporation

The microscopic processes occurring during growth are of great interest. The surface processes of most crystal growth techniques are difficult or impossible to observe directly. MBE offers a unique combination of slow growth rates and a very limited number of molecules which can participate in the growth. In addition, since the growth takes place under UHV conditions, it is possible to access the growth sites directly with a variety of analytical tools, including RHEED, scanning electron microscopy (SEM), scanning tunneling microscopy (STM), mass spectroscopy, photoelectron spectroscopy (UPS and XPS), and Auger electron spectroscopy (AES). Thus, it should be possible to accurately determine the microscopic mechanisms by which the epitaxial growth proceeds. Indeed, with the amount of information available for MBE growth of GaAs, it is

154

Molecular

Beam Epitaxy

possible to speculate on possible growth processes. Planar MBE growth processes can be separated into two distinct classes, depending upon how the chemisorbed

species incorporate

of growth processes incorporation

all involve

processes

involve

into the crystal lattice.

edge incorporation the surface

The first class

mechanisms.

migration

Edge

of chemisorbed

species to energetically favorable step edges on the crystal surface, where chemical bonding occurs, The island formation/step propagation model, discussed extensively in the preceding subsections, is an edge incorporation process. The second class of growth mechanisms involves incorporation of the chemisorbed molecular species directly on the surface, without the need for step edges. One particular surface incorporation process has been proposed which satisfies a number of important criteria.f164) In particular, this growth process minimizes the surface energy due to dangling bonds and maintains a constant (2x4) surface reconstruction throughout the growth. This particular process is suitable conditions. Growth processes

only for growth under arsenic-stabilized occurring for gallium-stabilized conditions

and for atomic layer epitaxy are different, since the surface reconstructions are different. The process starts with a unit cell of sixteen atoms, as depicted in Fig. 15(a) (counting the arsenic dimer vacancy).f164] This starting surface is arsenic-rich, with three arsenic dimers located at the surface, as expected for growth under an excess arsenic flux. Since this process relies on the incorporation of gallium and arsenic as pairs or dimers, it is helpful to consider two adjacent surface unit cells. In the first step, two chemisorbed gallium atoms bond to two of the adjacent arsenic dimers to form a gallium dimer, as shown in Fig. 15(b). There are two possible gallium dimer sites per unit cell, probably resulting in a statistical site distribution of gallium dimers across the surface. Since the surface is at an elevated temperature, the gallium dimer probably also hops back and forth from one site to the other. After the gallium dimer is present, no more gallium can be incorporated into an energetically stable configuration until the arsenic dimer is filled. In order to fill the arsenic dimer in an energetically stable manner, a correlated incorporation of the arsenic dimer and a gallium dimer is postulated. In this model, the incorporation of the gallium dimer limits the process, since an excess population of arsenic is available on the surface. The gallium dimer incorporates between two adjacent unit cells at a time when the gallium dimers in the adjacent cells are in phase, as shown in Fig. 15(c). At this point the arsenic dimer also incorporates in the arsenic dimer vacancy. Now, an additional arsenic

MBE of High-Quality

155

GaAs and AlGaAs

dimer can be incorporated across two of the adjacent gallium dimers as shown in Fig. 15(d). (It is also possible to incorporate a gallium dimer, but the presence dimer

intuitively

of excess

arsenic

more attractive.)

makes

the

incorporation

of an arsenic

Next, two gallium atoms bond to form a

dimer across the two adjacent arsenic dimers, as depicted As in the configuration

in Fig. 15(e).

depicted in Fig. 15(b), it is not energetically

stable

to fill the gallium dimer vacancy unless an arsenic dimer is also incorporated as shown in Fig. 15(f). Finally, an arsenic dimer bridges the two adjacent exposed gallium dimers, resulting in the completion of a single growth cycle with the surface unit cell translated Fig. 15(g,I.

fHH~

diagonally

as indicated

in

H-H-H

a)

b)

4

4

9

Figure 15. A proposed surface incorporation growth mechanism for the (100) surface. This process proceeds as follows: (a) initial surface with a (2x4) reconstruction; (b) incorporation of a gallium dimer; (c) correlated incorporation of an arsenic and a gallium dimer; (d) incorporation of an arsenic dimer; (e) incorporation of a gallium dimer; (r) correlated incorporation of an arsenic and a gallium dimer; (g) (not shown) incorporation of an arsenic dimer, resulting in a replication of the initial (2x4) surface translated diagonally.

MBE of High-Quality

3.6

Gallium

GaAs and AlGaAs

157

Desorption

The competing

process of gallium desorption

of GaAs and AlGaAs -620”C.t841t42r)t525)

temperature

The gallium

dependent desorption

makes the growth rate

for temperatures

above

rate is independent

of the

growth rate, but decreases with increasing aluminum fraction, x, for A&Gal &.t s4It421tt525)Gallium desorption is suppressed during the growth of the first 5-6 nm of GaAs on AlAs. 4*1)[470) The change in growth rate caused by gallium desorption as a function of substrate temperature and AIxGa,_.+

composition,

x, is given by the expressiont4*‘):

Bea = -9.545 x 1014 (1 - x) T-‘+O-13452n pm/hr

Eq. (1)

Increasing elevated

the V/III ratio decreases the gallium desorption rate at of temperatures.1 4*1) For example, at a growth temperature

635°C increasing the AsJGa BEP ratio from 16 to 30 and 56 decreased the gallium desorption rates by 29 nm/hr and 43 nm/hr, respectively.[*381[*3Q1[355jTh’ IS increase of 0.04 pm/hr agrees quite well with the predicted GaAs desorption rate of 0.0467 pm/hr predicted for low AsdGa flux ratios and a growth temperature of 63~YC.[~*~t However, when dimeric arsenic is used, there is no noticeable change in the growth rate for substrate temperatures between 570°C and 71 0”C.t1301t13glThe re-evaporation of gallium is probably inhibited by the higher sticking coefficient of dimeric arsenic,t372) along with the subsequent reduction in the quantity of free gallium adsorbed to the surface.t13g] The desorption process can be utilized to etch GaAs in-situ. Thermal etching

rates of GaAs at several

temperatures

pressures are shown in Table 2. The activation

and arsenic

over-

energy of GaAs sublima-

tion or “thermal etching” has been measured as 3.4 f 1 .O eV from thickness measurementst206] and 5.0 eV from RHEED oscillation measurements.t22r)t240) One RHEED oscillation

during GaAs sublimation

cor-

responds to one monolayer of gallium desorption.t*4 The fact that oscillations are observed suggests a layer by layer desorption mechanism. Increasing the arsenic pressure decreases the GaAs sublimation rate. While a P(As,)-” dependence has been reported for the GaAs sublimation rate, other researchers report that an increase in the arsenic pressure from 1.5 x 10m5torr to 5.0 x 10m5torr decreased the thermal etching rate of GaAs by a factor of 3-4.t227)t433)

158


Table 2.

Thermal Etching Rates of GaAs (in pm/hr) at Different Temperatures and Arsenic Pressures WS4)

5.0 x 10-s 1.5 x 10-s

700°C

725°C

750°C

co.03

0.048

0.120

0.210

0.180

0.420

0.660

0.072

775°C

The discrepancies in the activation energies and arsenic overpressure dependencies for GaAs sublimation could be due to differences in the temperature measurement or to competition between different gallium desorption pr0cesses.f 554) Modulated molecular beam mass spectroscopy measurements do, in fact, show two gallium desorption processes. Gallium is lost by evaporation from liquid gallium droplets on the wafer surface, with a low activation energy. Direct desorption of gallium atoms from the lattice is also observed with a significantly larger activation energy. At 750°C sublimation of gallium from AlGaAs proceeds until 3-4 monolayers of AlAs are formed at the surface.f**‘) The thickness of this self-limiting AlAs sublimation barrier should depend on the sublimation temperature. Sublimation of AlAs is not observed, even at temperatures as high as 750”C.f22r) 3.7

Thermodynamic

Redistribution

of the Near-Surface

Region

Recent studies suggest that the growth mechanism includes a thermodynamic interaction and relaxation of the near-surface atomic layers. Auger electron spectroscopy (AES) and x-ray photoemission spectroscopy (XPS) measurements reveal* a “surface enrichment leading

* A recent x-ray photoelectron diffraction (XPD) study has shown that a thermodynamicallydriven exchange reaction occurs between Al in and AlAs surface monolayer and gallium in the underlying GaAs subsurface monolayer during the substrate cooldown from 580’7 but not during growth.f44ea) Surface segregation effects arising from Ga-AI competitive incorporation processes were deliberately suppressed in the XPD study, so the XPD observations suggest that relatively little redistribution should occur in binary short-period superlattice (SPSL) pseudoalloys, perhaps explaining the near-surface redistribution differences of GaAs/AIAs SPSL pseudoalloys and ternary AlGaAs. Although competitive incorporation may consitiute the primary surface segregation process observed in the previous AES and XPS studies, these XPD results suggest that post-growth surface exchange may have resulted in an overestimate of the degree of surface segregation occuring during growth.

MBE of High-Quality

GaAs and AlGaAs

159

to a near binary surface” of GaAs during AlGaAs growth at temperatures as low as 600”C.t1ss~t34~t472~ Similar InAs surface enrichment occurs during the growth of 1nGaAs.t 1sslt**1)ts4s) This surface enrichment sistent with RHEED oscillation at 750°C

measurements

showing

is con-

gallium desorption

even when up to three atomic layers of AlAs are present at the

crystal surface.t**~ monolayers

No gallium desorption

of AlAs are present.

is observed when four or more

Mass spectrometry

continue to show gallium desorption gallium desorption rate decreases

measurements

also

when AlAs growth is initiated.t4’O) This by 50% for every two monolayers of

AlAs grown. This suggests that up to four atomic layers can directly participate in molecular beam epitaxial growth, and as many as ten monolayers can participate on a less significant scale. The number of monolayers influencing the epitaxial growth probably depends on the substrate temperature, the growth rate, the V/III ratio and the length of any growth interruption. This thermodynamic interaction and relaxation of the near-surface layers could also explain the long-range order observed in AlGaAs and InGaAs grown on (100) and (110) oriented substrates as well as anomalously poor AlGaAs surface morphologies for certain growth conditions and interface smoothing effects during growth interruption.[165)t218)t253)t472) The degree of ordering is strongly dependent on both the composition of the ternary alloy semiconductor and the growth temperature.t2531t254] It is suggested that the thermodynamically-induced formation of a gallium-rich surface layer during AlGaAs growth at 630-69O”C causes localized gallium agglomeration and increased epilayer roughness.[218)t472] (Thus, differences in the long-range order of this alloy semiconductor is one possible explanation for the growth temperature dependence of the carrier mobilities in AIGaAs.)t s4It284l The degree of interface smoothing caused by growth interruption at “GaAs-on-AIxGa, ,As” interfaces depends on the aluminum

fraction,

x, for x < 0.5.t 4g61 The dependence

of the interface

smoothing mechanism on the AlGaAs composition could be caused by increased surface diffusion of gallium and aluminum or by the redistribution of the group III atoms in the near surface

region.

The self-limiting

gallium desorption studies of AIGaAst 22’1 suggest that the depth of the near-surface redistribution increases as the aluminum fraction decreases. There are several possible thermodynamic driving forces for cation (group Ill) intermixing in the near-surface epitaxial layers, such as lattice strain and differences in bonding energy. For example, high-resolution transmission

electron

microscopy

and RHEED suggest that lattice strain

160

Molecular

relaxation dimensional

Beam Epitaxy

is the dominant

2.5 monolayers.t55] tween

driving force causing

cation mixing and three-

growth in InAs films (on GaAs), with thicknesses RHEED measurements

two-dimensional

strained

growth is abrupt to within which strained

InGaAs

InGaAs

exceeding

show that the transition growth

be-

and three-dimensional

0.2 mono1ayers.t 27g) The critical thickness

begins three dimensional

growth

decreases

at with

increasing indium fraction. Several new characterization techniques are being used to explore the microscopic processes which occur during MBE growth. In addition to RHEED studies, mass spectrometry is once again being used to characterize the species desorbing from the surface.t470)t554) In addition, in-situ scanning reflection electron microscopy (SREM) and micro-RHEED are being emp1oyed.t 2131t547) The scanning tunneling microscope (STM) is being used to characterize as-grown surfaces ex-situ and in-situ.t404] STM offers the ability to look at the actual surface reconstructions on an atomic scale.

Photoemission oscillations have recently been observed during oscillations make it posGaAs and AlGaAs gr0wth.t 1201 Photoemission sible to monitor layer-by-layer growth with the substrate rotating. Photoemission should make it possible to investigate the surface composition and bonding during growth. The surface growth processes are also being investigated by measuring the dependence of the optical reflectance of polarized light on the polarization direction during growth.f163) X-ray photoelectron diffraction (XPD) has been used to observe redistribution processes affecting single monolayers of AlAs in GaAs.t448a) Cathodoluminescence imagingt43)t406a)t420a~and photoluminescence microscopyt523] are being used to investigate thermodynamic redistribution effects, as well as the influence of the substrate (e.g., dislocations, excess arsenic outdiffusion, etc.) on the growth processes and material properties. Thermodynamic and kinetic (numerical) models are also being developed to describe the MBE growth processes.t218)[358)[487)f508) These recent efforts are rapidly elucidating underlying MBE growth.

4.0

SUBSTRATE The orientation

the complex

microscopic

processes

ORIENTATION of the substrate strongly influences

the incorporation

of impurities and deep levels in the epilayer. Epitaxial growth of GaAs and AlGaAs is most commonly performed on (100) oriented GaAs substrates, because of the wide range of growth conditions resulting in smooth

MBE of High-Quality

GaAs and AlGaAs

161

epilayers.

Improved epilayer quality is obtained when (100) substrates are During the past misoriented a few degrees toward the (11 l)A surface.

decade, there has been increasing

interest in the use of other low-index

substrate orientations. Non-planar growth, which requires growth on more than one crystal surface, is being investigated for optoelectronic and quantum wire devices.1 2261 Growth on (11 l)B surfaces is receiving attention

in order

to exploit

the

piezoelectric

field

for

device

applica-

tions.t851t14~t40g) Recently, however, the advantages of epitaxial growth on (11 1)t1241t50gland slightly misoriented (11 0)[151[2671[2681[3g81 substrates have been demonstrated. The growth of high purity epilayers demonstrated for higher index planes.t520)t521] 4.1

Growth on Misoriented For conditions

has also been

(100) Surfaces

in which the growth of GaAs and AlGaAs

directly on

the (100) orientation yield a rippled or “orange peel” morphology, growth on (100) substrates misoriented 2” toward the (110) surface results in smooth epilayers.1 1781 Later, AlGaAs was grown by MBE on convex or lenticular substrates centered around the (100) surface, extending to misorientations of 14”.t248] The best surface morphologies and highest luminescence efficiencies were observed for misorientations toward the (1 ll)A surface. In addition, misorientation toward the (1 ll)A surface significantly reduced or eliminated deep-level luminescence. A misorientation of 3-4” towards the (11 l)A surface also reduced the background impurity incorporation rate by a factor of three to four, but this difference decreases as the material purity increases.PO)tsog] The best material

was obtained

surface.t248]

with

misorientations

The material grown on substrate

of -6”

toward

the (11 l)A

misorientations

toward the

(11 l)B surface was rougher than that grown directly on (100) and had poorer luminescence. Increasing the AsdGa ratio increased the range of misorientations

resulting in smooth epilayers.

Other studies have confirmed

these results and it was further observed that substrate misorientation toward the (11 l)A plane reduced the concentration of two electron traps.t420l (The trap density increased for substrate misorientation toward the (11 l)B face). Photoluminescence (PL), deep level transient spectroscopy (DLTS), and electron transport measurements indicate that fewer impurities and deep level defects are incorporated as the substrate is misoriented toward the (11 l)A surface.1 4201t50g) Misorientation toward the (11 l)A surface results in a terrace structure with gallium atoms exposed at the step edges.

162

Molecular

Beam Epitaxy

Improvements in the surface morphology occur because the terraces suppress the nucleation of islands in favor of step propagation (as discussed above in Sec. 3: Growth gallium

atoms suppress

straight

terraces.

Processes).

the incorporation

Conversely,

terraces

Step edges with exposed

of lattice defects, with

resulting

arsenic-exposed

in

edges

[misorientation toward (11 l)B] are very jagged.f404] The incorporation of impurities at straight terrace edges should be less favorable than for highly disordered terrace edges. 4.2

Growth

on (110) and Misoriented

(110) Surfaces

There has long been an interest in the MBE growth of Ill-V semiconductors on (110) sur-faces.f26jp5)[252)f5201 Originally, (110) epitaxy was pursued for the purpose of zincblende on diamond heteroepitaxy,P5)f2521 since the nonpolar

(110) surface

has no net interface

charge

and the

suppression of antiphase domains was expected.f16g]f2521 GaAs (110) surfaces cleaved under ultra-high vacuum also have greatly reduced surface state densities in the forbidden gap,f463)f4641and (110) surfaces have been studied extensively with surface science techniques such as photoemission spectroscopy and Auger electron spectroscopy. Initial attempts to grow GaAs on GaAs substrates, oriented directly on the (110) surface, resulted in rough surfaces and severe faceting.f261f520) The amphoteric nature of silicon was also enhanced on the (110) surface. Low AsJGa ratios and high substrate temperatures caused silicon to incorporate primarily on the arsenic sublattice as acceptors, Conversely, high AsJGa ratios and low substrate temperatures caused silicon to incorporate primarily on the gallium sublattice as donors.f261f520] High-quality modulation-doped interfaces were formed on the (110) surface using very low growth ratios.f555] substantially

rates and low substrate temperatures, along with high V/III These low temperatures and slow growth rate conditions increased the incorporation

of impurities,

such as carbon.

Interest in growth on (110) surfaces has been revived by the observation that misorientation of the substrate slightly toward the (11 l)A plane (gallium-exposed steps) results in smooth GaAs epilayersf15] and in substantial improvements in the performance of microwave MESFETS.[~~~] (GaAs MESFETs have been fabricated on misoriented (110) GaAs with microwave results comparable to that of the best (100) MESFETs with identical geometry. This suggests a potential for superior performance by (110) MESFETs.) The concentrations of deep levels in GaAs and AlGaAs

MBE of High-Quality

are significantly

lower (and the PL efficiencies

GaAs and AlGaAs

are correspondingly

163

higher)

for epilayers grown on (110) GaAs misoriented 6” toward the (111)A surface than for comparable epilayers grown on (100) GaA~.t*~~t*~~) Growth on substrates

oriented directly on the (110) plane gave very broad

quantum well luminescence, as expected from the rough morphology of the epilayers,t152) but growth on (110) surfaces misoriented 6” toward the (11 l)A surface gave 77 K PL peaks as narrow as 3 meV for a 15 nm quantum well.t2671 Long-range order has been observed in AlGaAs grown on (100) and (110) surfaces.t 253] Perfectly ordered, such an alloy would consist of alternating layers of AlAs and GaAs when viewed along the (110) growth direction or the (100) direction perpendicular to the growth direction. [Identical long range order has been observed for In,,,Gac,As grown on (110) lnP].t254) This long-range order is stronger and occurs at lower growth temperatures for AlGaAs grown on the (110) surface than for alloys grown on the (100) surface.t 253) Long-range order is thought to be the equilibrium state of AlGaAs, suggesting that the (110) surface may be preferable for the growth of high quality AlGaAs. A quasi-planar growth mechanism was recently proposed and demonstrated by the growth of high quality GaAs/AIGaAs heterojunctions and quantum wells on misoriented (110) G~As.[*~~ Finally, surface defects, even oval defects, are substantially reduced for growth on misoriented (110) substrates.[267) The deep-level density reduction and elimination of oval defects shows that impurities incorporate differently or with significantly reduced concentrations because of the (110) quasi-planar growth mechanism. 4.3

Growth on (nll)A

and (nll)B

(1 s n s 9) Surfaces

Initial attempts to grow GaAs/AIGaAs on (11 l)B surfaces resulted in poor epilayer morphology and correspondingly broad quantum well luminescence.t152) Smooth morphology

GaAs and AlGaAs have been achieved

on (11 l)B surfaces, after it was demonstrated that the substrate cleaning, growth initiation, and growth conditions are all critical.[124j Silicon incorporation on (11 l)B surfaces is highly auto-compensated, but good luminescence efficiencies are obtained. Improved silicon-doped GaAs growth has been reported for growth on (11 l)B oriented substrates using migration enhanced epitaxy (MEE) at substrate temperatures of 400-530”C.t14fl (Excess arsenic was detrimental GaAs MEE layers.)

to the quality

of these low temperature

164


MBE growth of GaAs and AlGaAs on higher index surfaces has also been investigated. The (211) surface has been of particular interest for the purposes of suppressing

the formation

of anti-phase

domains

in polar on

nonpolar epitaxy. t251)f4s2) High-mobility modulation-doped GaAs/AIGaAs structures have also been grown on (211)A, (21 l)B, (311)A, (311)B, (51 l)A, (51 l)B, (71 l)A, (71 l)B, (91 l)A, and (91 l)B surfaces.t5*‘) Particular attention has been given to differences in the amphoteric behavior of silicon, Smooth

surface

morphologies

have

been achieved

on both the

(21 l)A and (211)B GaAs surfaces through the use of a GaAs/AIGaAs smoothing super1attice.t 482) However, the RHEED pattern indicates that the (211)A surface is microscopically smooth, while the (211)B surface shows some microscopic roughness. The growth mechanisms resulting in surface roughness on (211)B surfaces are not understood, but may relate to surface contamination, surface disorder, surface migration or thermodynamic redistribution of the near surface region. Silicon is incorporated as an acceptor during MBE growth on the (11 l)A, (21 l)A and (311)A surfaces.1 5*1j On the (21 l)A surface, silicon is incorporated as a partially-compensated acceptor at low As/Ga ratios and as a partially compensated donor at very high As/Ga ratios.f4**j Silicon is incorporated

as adonorforthe

(11 l)B, (21 l)B and (311)B surfaces.f482)f521)

Silicon donor incorporation on (211)B surfaces is essentially uncompensated.f4s2) For orientations with indices higher than (311), it was found that silicon was incorporated as a donor.1 521) Planar p-n junctions and LEDs have been formed by growing MBE layers on substrates etched to simultaneously expose the (100) and (11 l)A surfaces, using the fact that silicon can be incorporated

as an acceptor on the (11 l)A surface and as a donor

on the (100) surface.t 32711 32ej Initial work on these plane-selective

junc-

tions had problems caused by rough morphologies on the (11 l)A surfaces and possibly by defects at the (11 l)A/(lOO) inter-face.f3*s) The quality of plane-selective junctions has recently improved by forming the p-n junction at the intersection of a (31 l)A surface with a (100) surface.t320) The vertical

redistribution

lengths of beryllium

during the growth of

(311)A GaAs grown at 630°C are reduced by a factor of 36 in relation to (100) growth.f344j This reduced beryllium movement is attributed to the large step density, which is believed to suppress the incorporation of highly mobile beryllium

interstitials.

MBE of High-Quality

The exciton luminescence

transitions

GaAs and AlGaAs

165

of GaAs are more intense for

(21 l)A GaAs than for (100) GaAs, while the carbon acceptor transition at 1.49 eV was less intense.t50g] The defect-induced bound exciton (DIBE) transitions

were much stronger for (211)A GaAs than for (21 l)B GaAs.*

AI,,Ga,,As,

on the other hand, the carbon

more intense

for growth

acceptor

on the (211)A surface

luminescence

than for growth

In is

on the

(211)B surface. The increased carbon acceptor formation on (21 l)A AlGaAs (compared to GaAs) may occur because aluminum is more reactive than gallium or because of a reduction in the carbon surfacemigration length on AlGaAs. The photoluminescence linewidths of (21 l)A superlattices are comparable to, or smaller than, those produced by (100) superlattices, but the luminescence linewidths produced by (211)B superlattices are significantly br0ader.f 4s*)fsO~ The photoluminescence efficiency of (211) superlattices is about an order of magnitude higher than that of (100) superlattices. The increased linewidth agrees with the microscopic roughness of (21 l)B surfaces observed with RHEED. Smooth morphologies, resulting in narrow quantum well luminescence transitions, have been achieved with MBE on the (31 l)A and (31 l)B surfaces.f152) However, impurities incorporate more readily at the (31 l)B surface than at the (31 l)A surface.t152) The increased surface roughness of the (21 l)B surface over that of the (21 l)A surface agrees with measurements on misoriented (100) surfaces. (Misorientation of the (100) surface toward the (11 l)B surface increases surface roughness, while misorientation toward the (11 l)A surface reduces the surface roughness). This difference in surface roughness may result from lower diffusion lengths for gallium on arsenic-terminated surfaces (and steps) than for gallium-terminated also possible thermodynamic

that arsenic-terminated redistribution

surfaces

in the epilayers

surfaces

(and steps).

and steps result near the surface.

It is

in less Clearly,

more work is needed to understand the differences between epitaxial growth on arsenic- and gallium-terminated surfaces (and steps).

The DIBE transitions are produced by complexes involving multiple acceptorsf’ 191f3er)(456a) carbon-defect complexes,[55a~[~g~~~~[3~1~[4581 or lattice-defect complexes.f256)f258)[25G) Thus, the (21 l)B surface probably favors the incorporation of carbon as a simple acceptor, while the (21 l)A surface favors the incorporation of carbon-carbon and carbon-oxygen complexes. Careful measurements and numerical modeling of growth on (21 l)A and (21 l)B surfaces could help elucidate the structure of the DIBE complexes and the growth processes leading to their formation.

l

166

Molecular

4.4

Growth

Beam Epitaxy

on (221)A, (221)B, (331)A and (331)B Surfaces

A limited amount of data is available

for growth on (221)A, (221)B,

(331)A and (331)B surfaces of GaAs.f 509tt5101In one experiment, the GaAs luminescence from (100) GaAs was stronger than that of (221)A or (221)B GaAs, while the situation was reversed for AIo,,Gao,,As.f50g] A photoluminescence linewidth as small as - 10 meV was observed for (22l)A Al,.,, Ga,,,,As, but the (221)B AlGaAs linewidth was quite broad. The luminescence from a multiple quantum well structure was more intense for (221)A oriented wells than for (100) wells, but the transition was substantially broader, probably due to interface roughness. Silicon is incorporated as a shallow donor on (331)B GaA~.t~‘~l For growth on (331)A GaAs, however, silicon incorporates as a shallow acceptor if the V/III ratio is low. Increasing the V/III ratio causes silicon to form a shallow donor in (331)A GaAs. The DIBE luminescence transitions were smaller for (331)A GaAs than for (100) GaAs and were even smaller for (331)B GaAs, in good agreement (21 l)B GaAs.

5.0

with the results obtained for (21 l)A and

OVAL DEFECTS

There is a close relationship between the nucleation of surface defects and the presence of impurities and par-ticulates on the substrate.t66)t313] Two oval defect families exist for MBE GaAs/AIGaAs. The first oval defect family is characterized by a spike with a pit at the center of the defect and originates from surface contamination. The second family is characterized by a spike without a pit and originates from liquid gallium agglomeration. Therefore, it is important to discuss recent developments relating to surface defects, particularly

the oval defect family.

Surface contamination of the substrate is an important source of oval defects. Even recent results show that chemical cleaning of GaAs substrates reduces the oval defect density.f681 Further growth comparisons on material suggest that GaAs than that of silicon procedures have

grown on as-delivered GaAs and silicon substrates substrate surface quality and cleanliness is still poorer substrates. Improvements in the substrate preparation reduced the particulate-related oval defect densities.

For example, the pair defect which is linked to sulfur surface contamination from H,SO,-based etches, is eliminated by performing a second etch

MBE of High-Quality

(30% HCI) and thoroughly

rinsing the substrate

Improvements in the substrate the number of wafer cleaning defects due to particulates deionized resistivity

with deionized

water.1641

cleaning environment and a reduction steps have also reduced the number

and surface contamination.

water, used during substrate

oval defect density.

167

GaAs and AlGaAs

The deionized

cleaning,

in of

The quality of the

also strongly

influences

water must have both high electrical

(i.e., low ionized impurity concentrations)

and low organic impu-

rity concentrations (e.g., bacteria). Other sources of impurities internal to the MBE system can cause particulateand contaminant-related oval defects. For example, the process of moving wafers inside the vacuum system increases the oval defect density. Removal of arsenic deposits inside the growth chamber and dust inside of the buffer and entry/exit chambers may also be necessary to obtain oval defect densities below 100 cm-2.t270)t3g6j Impurities emanating from the hot sources and crucibles are another source of oval defect nucleation sites.tg0)t443j Re-using the gallium crucible reduces the density of Ga,O,-related surface defects.t 443j This decreased oval defect density is consistent with the diffusion of impurities out of P-BN[~~] and with the slow decomposition of the p-BN when exposed to molten aluminum.t285)f271) The use of single crystal sapphire crucibles also eliminates the Ga,O,-related oval defects commonly observed with p-BN crucibles.t443] (However, sapphire crucibles have other problems, including contamination of the GaAs with oxygen and deep levels, poor thermal conductivity, and their susceptibility to catastrophic failure induced by thermal shock.) The other major source of oval defects is gallium agglomeration on the crystal surface. If growth is immediately initiated on a surface which has been allowed to go “gallium rich” for long enough to form gallium droplets, the resulting oval defect density can be extremely droplets

on the substrate

are also caused

by gallium

large.

spitting

Gallium from the

gallium s0urce.t 545) (Oval defect formation can be induced in GaAs grown by MOVPE by introducing a “spray of fine globules” of a Ga-In-Al mixture during growth, supporting the gallium spitting defect formation mechanism.)t472j

Gallium

spitting

can occur when the gallium

droplets

which

have condensed on the cooler crucible lip roll back into the molten gallium source. Gallium spitting also occurs when gallium droplets explode.t5gj Explosion of gallium droplets occurs both spontaneously and through collisions between mobile droplets inside of the crucible lip. Gallium droplet explosion is consistent with measurements showing that these droplets are coated with polycrystalline GaAs shells.t4g4j A drop in the oval

166

Molecular

defect density

Beam Epitaxy

is observed

when the gallium

crucible

full.t178tt545t The oval defect density subsequently level drops.

is more than 80%

increases

(Other results suggest that this dependence

as the gallium of the gallium-

related oval defect density on the degree of crucible filling depends on the geometry of the source furnace and its position in the MBE system.)t66)t681 Outgassing the gallium source above its normal operating temperature prior to growth also increases the oval defect density.t68] The oval defect density is directly related to the size and number of gallium droplets at the lip of the gallium crucib1e.t 66I[4g4) The number of droplets inside the crucible lip decreases as the temperature of the crucible lip increases. The droplet density also decreases as the distance between the substrate and the crucible increases.t4g4t The temperature dependence is believed to result from a growth rate reduction of the poly-GaAs shells. The dependence on the source-substrate separation is attributed to a reduction in the arsenic flux coming off the substrate. Large gallium flux variations, flux noise, with a center frequency

of - 0.1 Hz, have been correlated

increased oval defect densitiest 332t Speculation suggests that this noise is related to the gallium spitting phenomenon. Work with hot lip furnaces shows large reductions in both number of gallium droplets on the crucible lip and the gallium-related defect density.t 4431t526tt545t Other work suggests that oval defects can

with flux the oval also

be nucleated by arsenic “spitting. “4881 [ The density of these arsenic-related oval defects can be reduced with the use of a hot lip furnace for the arsenic source. Re-using spent aluminum crucibles (p-BN)* for gallium also eliminates the formation of gallium droplets on the crucible lip and the nucleation of gallium-related oval defects.t66)f68t Increases

in the oval defect density from < 500 crnm2to - 2900 cme2

also occur when the growth rate increases from 0.02 pm/hr to 1.1 pm/ hr.t326) This growth rate dependence occurs because the larger population of adsorbed gallium at higher growth rates increases the probability for gallium agglomeration. (No Ga,O was observed in the residual gas spectrum,t326l supporting the hypothesis that the increased oval defect density is caused by gallium agglomeration and not by Ga,O contamination). Increasing the (100) GaAs growth temperature from 570°C to 680” resulted in an increase in the oval defect density, which is attributed to

* Aluminum reacts with the p-BN to form a layer of boron-doped aluminum nitride, AIN, with a boron-rich interface between the AIN coating and the p-BN crucib1e.L271] This boron-rich AIN alloy allows the gallium metal to wet the crucible surface, thereby preventing the formation of gallium droplets and gallium spitting.

MBE of High-Quality

gallium

agglomeration

surface.t130)

from an increased

GaAs and AlGaAs

population

of gallium

169

on the

The use of dimeric

number of gallium-related

arsenic also dramatically reduces the oval defects.1 1391ts1sjt4s2)The dimeric arsenic

may reduce either the gallium surface populationt13g] or the gallium surface mobility, thus preventing gallium agglomeration and oval defect formation.f313j Buffer layers can be used to suppress the formation of crystallographic defects. The oval defect density is reduced from - 490 cme2 to 70 cm-2 if the initial 50-200 nm thick GaAs buffer layer is grown by migration-enhanced epitaxy (MEE) at low temperatures (300”).t436) (MEE is similar to MBE except that the group III and group V fluxes are supplied alternately, resulting in the growth of one epitaxial layer each cycle.) This MEE buffer layer reduced the number of oval defects caused by microscopic surface contamination. The enforced two-dimensional growth mechanism of MEE probably suppresses three-dimensional growth nucleation at the contaminants and effectively buries these microscopic impurities without forming oval defects. Finally, it has been observed that surface defects, even oval defects, can be eliminated for growth on misoriented (110) substrates.t267) The reduced deep-level density and elimination of oval defects indicates that impurities incorporate differently or with significantly reduced concentration because of the (110) quasi-planar growth mechanism.

6.0

SURFACE Impurities

MORPHOLOGY

and gallium agglomeration

ment of surface and interface the

mobilities

AND INTERFACE

roughness.

of two-dimensional

ROUGHNESS

play key roles in the developInterface

electron

roughness

gases,

the

degrades

luminescence

linewidths of AIGaAslGaAs/AIGaAs quantum wells,t42)t44) the peak current densities of resonant tunnel diodes,t 2721t4°2j and increases the optical scattering losses of quantum well laser diodes. The presence of impurities on the growth surface can reduce the surface diffusion and aluminum,

resulting in increased

length of gallium

interface and surface roughness.t454)

Desorption of surface contaminants may be the major reason that earlier quantum well luminescence studies suggested somewhat higher optimal growth temperatures of 675-710 ’ .t5251 Continued improvement of the source purity (especially aluminum) and vacuum quality has substantially reduced the problems of surface contamination

and interface

roughness.

170

Molecular

Beam Epitaxy

Thus, the system

and source

purity

have a strong

effect on interface

smoothness. If there are too many impurities impinging on the AlGaAs surface, it may not be possible to grow reproducibly smooth interfaces, since the temperature

required to desorb the impurities

which favors three-dimensional

is in the regime

growth.

The morphology of AlGaAs depends strongly on the growth conditions employed (substrate temperature, arsenic species used and the V/III ratio). The growth of AI,Ga,,As (x 2 0.2) in the temperature range of 630°C to 690°C requires an excess arsenic (As,J flux above that required for lower (5 620°C) and higher (2 700°C) growth temperatures.t13]t351)t472] Failure to supply this excess arsenic flux results in very rough surface morphologies. This surface roughness is greatest for aluminum mole fractions of - 0.5, while specular surfaces are obtained for both the GaAs and AlAs binary compounds with the same flux ratios.t13) (Strong damping of AlAs RHEED oscillations at growth temperatures in the 600-700°C range is observed.1 3621 Thus, while the binary AlAs alloy grown in this temperature

regime is probably

better than that of AlGaAs ternary alloys,

the AlAs quality is probably also not optimized.) This surface degradation occurs with As,, but not with As,.[‘~~) Improved surface morphologies and AlGaAs quality are obtained with lower growth rates (c OZpm/hr), and with the use of either high or low V/III flux ratios of - 11 and - 2 .t131t17el This anomalous surface roughness has been variously attributed to an arsenic-deficient surface, surface contamination and reduced aluminum surface mobility. Thermodynamic redistribution of the near-surface layers and the formation of a gallium-rich surface layer could also cause the surface roughness. An arsenic-deficient surface is inconsistent with the observed improvement in morphology at high growth temperatures or low V/III ratios. Surface contamination probably does not cause the surface roughness, since tnis effect appears to be universally observed and since a layer of tin segregating on the surface actually improves the morphology. bution

Low aluminum mobility on the surface and thermodynamic redistriof the near-surface layers remain possible surface degradation

mechanisms. Surface segregation of gallium increases as the growth temperature increases, but should decrease as the V/III ratio increases.t21~t472] Gallium from this gallium-rich layer then diffuses anisotropically across the surface, forming the characteristic surface ripples. The surface segregation effect should be strongest for aluminum mole fractions of - 0.5, in agreement with the observed peak in AlGaAs

MBE of High-Quality

roughness.

At high growth temperatures,

GaAs and AiGaAs

gallium

desorption

important and reduces the surface gallium population. tion is attractive,

more work is necessary

171

becomes

While this explana-

to test its validity.

If this model is

correct, a detailed investigation of AlGaAs roughness as a function of the crystal orientation and epitaxial growth conditions should yield important quantitative information about the growth processes. Other processes can also contribute to interface roughness. For instance, increasing the growth temperature from 580°C to 650°C increases the interface roughness because of a change from two-dimensional growth to three-dimensional growth.t43) On the other hand, if the growth temperature is too low or the growth rate too fast, epitaxial growth will proceed by island nucleation instead of step propagation, resulting in an increase of the inter-facial disorder. There is also evidence to suggest that the use of dimeric arsenic (AS*) suppresses surface roughness. The role of dimeric arsenic is considered in greater detail in the next section. The surface roughness of GaAs and AlGaAs epilayers grown on other substrate orientations also depends strongly on the growth conditions used. Decreasing the substrate temperature from 550°C to 450°C during the growth of (110) GaAs inisoriented 6” toward the (117) plane resulted in a strong reduction in the surface roughness.t26rl Growth on convex or lenticular substrates centered around the (100) surface also showed an increase in the AI,Ga,,As over a wide range of misorientation creased from 620°C to 650”C.f246]

7.0

SUBSTRATE DEFECT

CLEANING

AND MBE GROWTH:

IMPURITY

AND

INCORPORATION

Many parameters MBE growth

(0.3 i; x 5 0.4) surface roughness as the substrate temperature in-

influence

of compound

the incorporation

semiconductors.

These

of impurities

during

parameters

can be

separated into three general categories: (7) Those directly related to the substrate, particularly preparation and buffer layer design. (2) The choice of group V species (As,, As, or ASH,). (3) Those parameters which directly affect the thermodynamics governing growth and impurity incorporation, but are easily changed during growth, and include: growth temperature, ratio of group V flux to group III flux, growth rate, and growth interruption time.

172

Molecular

Beam Epitaxy

7.1

Substrate

Preparation

and Cleaning

Substrate preparation and cleaning is much more important for MBE than for other epitaxial growth techniques, due to the lack of a convenient in-situ etching procedure.

LPE has significant

surface etching due to the

contact of the melt with the substrate and interaction

with hydrogen

and/or

other gases. VPE and MOVPE techniques have hydrogen present and can use HCI and other gaseous surface cleaning techniques. Elaborate attempts to provide in-situ cleaning for MBE using ion sputtering and thermal etching techniquesf 2101f4331 have met with mixed success. Ion sputtering successfully removes carbon contamination, but the resulting damage creates a semi-insulating layer at the substratelepilayer interface.p6] A vacuum of cl Oegtorr is necessary for successful cleaning by ion bombardment. (Residual gas pressures greater than 1 OTgtorr result in recontamination of the surface with CH,+ and COt21). On the higher index (211) GaAs surface, ion-sputtering leaves the surface unstable, resulting in faceted growth.flgg) Thermal etching has also been investigated as a means for in-situ cleaning of GaAs surfaces prior to MBE growth. Thermal etching is reported to reduce the “dip” in the free carrier concentration at the interface.f433) (This reduction of the dip in carrier concentration was originally attributed to a reduction in the concentration of interfacial carbon impurities for temperatures above 750°C. However, this work did not investigate p-type interfaces or the role of deep-acceptor impurities, which are important for regrown interfaces.f 32g] Others results show that “thermal etching” does not substantially reduce the interface contamination, but rather results in the accumulation of dopants from the substrate at the interface.f166)f206) These accumulated dopants reduce the interface depletion by compensating the carbon contaminants. This accumulation is not surprising, since both carbon and silicon have much lower vapor pressures than either arsenic or gallium.) Thermal etching of AI,Ga,_,As (x 2 0.01) still results in carrier depletion at the interface, presumably due to the increased adsorption of residual gases on the thin AlAs layer which forms at the surface due to the preferential aluminum arsenide.f206)

desorption

of gallium arsenide

over

Chemical etching with HCI and HCI + H, gas mixtures has also been used for in-situ cleaning of GaAs substrates prior to MBE growth.f435l Etching with HCI gas reduces the interface carrier depletion from - 1.2 x 1012 cm-2 to - 3.5 x 10” cm-2, independent of the etch depth. Etching with a mixture of HCI and H, reduces the interface depletion to values

MBE of High-Quality

below 1 x lOlo cm-*. magnitude

reduction

pared to thermal

SIMS measurements in the interface

etching.

carbon

The removal

GaAs and AlGaAs

show an additional contamination,

of the carbon

attributed to chemical reactions between the carbon active atomic hydrogen liberated by the HCI. The preparation

of substrates

prior to loading

173

order of

when

com-

contamination contaminants

is and

them into the MBE

system will remain a critically important issue, until such time as a reliable process for in-situ substrate cleaning is developed. The substrates must be prepared in a particulate-free environment to minimize the defect density caused by particulates. interface result in macroscopic

Particulates at the substrate/epilayer defects (up to -100 pm), consisting of

twinned and highly dislocated material. The conventional GaAs substrate preparation consists of a degreasing procedure (trichloroethelene or trichloroethane, acetone, methanol or isopropanol), followed by a two minute 3:l :l ::H2S04:H202:H20 etch to remove - lo-20 pm of polish damage, another etch (HCI or 1 :l :200::NH40H:H202:H20) to remove oxides or any sulfur impurities left by the sulfuric acid etch, followed by several deionized water ~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~ After etching

the substrate

with HCI, it is necessary

to rinse the

substrate for more than ten minutes with deionized water in order to remove the chlorine from the surface.t 3101While chlorine does not seem to have any negative impact on material quality, careful SIMS measurements show that even high purity H3SO4 and HCI can introduce heavy metal (Fe, Mn, Mg, and Cr) contamination at the substrate/epilayer interface.t201) Incomplete rinsing of the substrate following a H2S04:H202:H20 etch results in contamination of the regrown interface with sulfur-containing impurity species and the formation of pair defects.[641 This problem is reduced by following this etch with a concentrated HCI etcht641 or with a sufficiently long deionized water rinse.1 1431[14sl Auger electron spectroscopy shows surface contamination of GaAs with carbon, oxygen, and sulfur after bromine:methanol upon heating to temperatures carbon contamination with bromine:methanol

etching.teg] The oxygen and sulfur desorb of 520 and 540°C respectively. However,

does not desorb at these temperatures. Etching solutions (once widely used during GaAs substrate

preparation) results in persistent carbon contamination of the surface, but etching with sulfuric-acid:hydrogen-peroxide:water solutions results in significantly lower carbon contamination levels.t56] Carbon contamination levels in excess of 0.2 monolayer results in twinned growth and surface facetingJ61

174

Molecular

Beam Epitaxy

when

During the cleaning procedure, the wafer surface must be kept wet transferring from one solution to the next, in order to prevent

impurities and particulates the surface

from adhering to the wafer surface.t205] Rinsing

of GaAs after etching,

without

exposing

the surface

to air,

reduces the amount of carbon on the surface and increases the amount of oxygen.fs6) However, some results suggest that deionized water can also be a source of B, C, Mg, and Fe contamination.t32g) The use of distilled, deionized water results in fewer particulates and lower surface contamination.f151] Thus, the purity of the water used for rinsing is important and there is an optimal rinsing time which depends on the water purity. In most early MBE systems, the substrates were etched and then blown dry with filtered nitrogen. Spinning the wafer dry, rather than blowing it dry with filtered nitrogen gas, reduces the particulate density on the surface of the wafer and is now a routine part of the standard cleaning procedure.f3s6) After drying the wafer, it is mounted onto a molybdenum or tantalum substrate holder. In earlier systems (many of which are still in service), this was done with indium or gallium solder on a hot plate.f84tf17s] In newer systems, which use direct radiative substrate heating, the wafers are mounted on ring-shaped substrate holders with clips or other mechanisms.~~ss1~~~7l~~s~l~~s~l H owever, direct radiative substrate heating eliminates the growth of thermal oxides during In-bonding, resulting in incompletely passivated surfaces. lndium bonding is believed to be a significant source of impurities,[l741[lss1[2051[3891,and causes wafer strain and backside roughness. Thus, it is highly desirable to eliminate indium bonding. Clearly, it is advantageous to introduce a final step to the conventional substrate preparation procedure to grow a clean, non-porous, and reproducible oxide. It is important to avoid contamination of the wafer holders by handling them. (Even “clean” gloves can introduce contamination and cause irreproducible results.) It is beneficial to use a special chuck to hold the wafer holder when the wafer is being mounted and dismounted.

This

same chuck can be used to load the wafer and wafer holder into the MBE system. 7.2

Protective

Oxide

Growth

Since there are no suitable in-situ cleaning techniques

for MBE, the

preparation of a high quality protective surface oxide layer is crucial. This oxide layer protects the surface from contamination and can be thermally

MBE of High-Quality

desorbed in the MBE system electron spectroscopy the H,SO,:H,O,:H,O

GaAs and AlGaAs

prior to growth.

Contrary

175

to early Auger

(AES) studies on indium-bonded GaAs samples,[3561 etch followed by a deionized water rinse does not

produce a passivating surface oxide layer.[ 3101 HCI solutions also produce no significant surface oxide on GaAs, and bromine:methanol solutions resulted in the formation

of a poor quality gallium oxide.[2111 Exposure

to

stagnant deionized water results in a strongly arsenic-deficient oxide.[2111[3101 The oxide produced by the conventional cleaning process is porous and contains arsenic-rich octahedra, which result in surface defects.[1461 Instead, the oxide layer is produced by subsequent exposure to air, particularly during In-bonding to the substrate holder at temperatures in excess of 160”C.[310] Thermal oxides formed by heating the substrates at 250300°C in air for 3-5 minutes result in low surface defect densities, allow the growth of high purity GaAs and high mobility two dimensional electron gases.[ 1451[1461[2371[4341[4731 At 45O”C, the growth rate of a thermal oxide layer is - 2 nm/minute and is linear for times up to ten minutes.[434] Thermal oxides are primarily composed of gallium oxides with nonoxidized arsenic. Thermal oxides reduce the interface depletion from - 1.3 x 1012 cm-2 to 4 x 10” cme2 as the oxide growth time (450°C) increased from O5 min. Oxide growth times greater than five minutes did not reduce the interface depletion further. Passivating oxides on GaAs can also be grown by exposing the substrate to ozone, generated by placing the substrate under an ultraviolet light source in air (T < 60”C).[2111 GaAs oxides produced by ozone exposure are stoichiometric, possessing nearly equal quantities of arsenic and gallium oxide. Arsenic oxides thermally decompose between 140°C and 400X, 500°C and 600”C.[2111[3101 7.3

Wafer Outgassing Outgassing

while gallium

oxides

desorb

between

and Oxide Desorption

the wafer prior to the oxide desorption

is an important

step for reducing the substrate/epilayer interface contamination. Exposing the substrate to an ionizing beam (e.g., RHEED and AES electron beams) prior to heating it to 350-4OO”C in vacuum makes it very difficult to thermally remove carbon impurities from the substrate.[761[841 On the other hand, the presence of a vacuum ion gauge about six inches from the substrate does not noticeably affect the desorption of hydrocarbons.[76] Contaminated surfaces often result in spotty RHEED patterns after the oxides are desorbed. These spots change to the usual streaks obtained for

176


smooth, clean surfaces after - 10 nm of growth. Outgassing the substrates at - 400°C for 1 hour in the preparation chamber, transferring them to the growth chamber and baking them at 400°C until raising the temperature for oxide desorption

and growth is believed to result in lower surface

impurity

than

concentrations

less extensive

baking.f178)[266)

Extensive

outgassing of the substrate and wafer holder also results in higher twodimensional electron gas mobi1ities.f le5) These results are consistent with Auger electron spectroscopy studies showing that most of the surface hydrocarbon contamination can be reduced by heating the substrate above 350-4OO”C, as long as it is not exposed to an electron beam.fr6) Baking a (100) GaAs wafer under an arsenic flux after the oxide is desorbed can result in an increase in the concentration of carbon incorporated in the GaAs epi1ayer.f 497) The principal sources of the carbon contamination are believed to be the arsenic source and furnace assembly. However, there are more recent indications that baking a bare GaAs surface for 1 hour at - 580°C after oxide desorption

under a small (l-l

0x

10s torr) As, flux results in significantly higher two-dimensional electron gas mobi1ities.f la51 This improvement over the earlier studies is a result of the lower arsenic fluxes used, better vacuum quality, higher arsenic source/furnace purity, and a lower oxygen-related background from the oxide desorption. 7.4

Buffer Epitaxial

Layer Design layers grown

between

the substrate

and the active

epi-

taxial device layers are known as buffer layers. Buffer layers serve several purposes, including: (i) moving the device active regions away from the substrate/epitaxial interface to minimize the undesirable effects caused by surface damage and contamination (introduced during wafer sawing and polishing); (ii) gettering impurities outdiffusing from the substrate to the growth surface; (iii) smoothing the growth surface to obtain atomically-flat heterointerfaces; and (iv) reducing the crystallographic defect density (oval defects, dislocations, etc.) in subsequent epitaxial layers. The optimal buffer layer structure depends strongly on the purpose of the buffer layer and the intended device application. Common buffer layers include simple GaAs layers, compositionally graded AlGaAs layers, quantum wells, and superlattice

structures. Simple GaAs Buffer. The most widely used buffer layer is a simple GaAs layer with a thickness of 0.2-2 pm. The principal purpose of a GaAs

MBE of High-Quality

buffer layer is to move critical epitaxial

GaAs and AlGaAs

177

layers away from the initial growth

interface and to reduce the background impurity concentrations in the epitaxial material. Early work on high-mobility two-dimensional electron and hole gases demonstrated the importance of relatively thick GaAs buffer layers,~~~~I~~~~1~~~s1~~~Q1~~Q~1~~~~1~~~~1~~~~1 Large concentrations of chromium deep-levels have been observed at the interface between Cr:GaAs substrates and n-GaAs layers grown by chloride VPE.t61t5511 These interface traps strongly affect the conductivity of the n-GaAs active layer by changing the carrier sheet density and interface depletion width as a function of substrate potential. SIMS and photoluminescence measurements of annealed GaAs substrates reveal a l-3 pm thick initial layer with high concentrations of iron and manganese due to anomalous impurity outdiffusion from the substrate.t 374] SIMS measurements of MBE layers show the accumulation of impurities, such as carbon, sulfur, and silicon at the substrate/epilayer

interface.t 391) Since the resolution

of SIMS is - 1Or4

cme3 (higher for some elements and for thin layers), it is probable that accumulation of undetected species also occurs at the interface. While higher-purity substrates have largely suppressed the outdiffusion of impurities from the substrate, these measurements demonstrate the importance and efficacy of GaAs buffer layers. Low Temperature GaAs and AlGaAs Buffers. Low growth temperatures (150-330°C) result in semi-insulating GaAs for normal growth 1 p m/hr.tQ21t284t4601 L rates of ow-temperature semi-insulating buffer layers are sometimes used to reduce sidegafing and backgating of MESFETs and MODFETs, since the large deep-level densities of these layers pin the Fermi level at the interface and inhibit changes in the channel conductivity caused by potential changes applied through the substrate and buffer layer. Unfortunately, diffusion of traps into the FET channel can subsequently degrade the device performance.t283)[416] Even though the low-frequency performance of MESFETs are improved through the use of low-temperature

buffers, poorer high-frequency

observed than for devices using conventional Resitivities

of

r 5 x lo5

ohm-cm

performance

is

buffer layers.t4781

have

been

measured

for low

temperature GaAs, while substantially larger restivities (2 lo8 ohm-cm) have been measured for low temperature AlGaAs (0.3 s x s 0.5) .tQ21These GaAs and AlGaAs restivities are lo2 and lo3 times larger, respectively, than observed for GaAs and AlGaAs layers grown at 600°C. As the growth temperature is reduced from 620°C to 450°C there is a 20% drop in the electron density of n+-GaAs (4 x 1018 cmm3)and a 30% drop in the electron

178

Molecular

Beam Epitaxy

mobility.f334) As the growth temperature 35O”C, the electron concentrations

is further reduced from 450°C to

in both GaAs and AI,,3Ga,,7As

by a factor of - 106. The GaAs mobilities dropped to 25-30% obtained for growth at 620°C.

(The resistivity

dropped

of the values

of AI,,sGa,,,As

grown with

As, only dropped by - 104.) The photoluminescence efficiency also decreased as the growth temperature was reduced. The electron concentrations and photoluminescence intensities recovered with annealing at 850°C for 5-20 sec. (The electron concentration only recovered to about one third of the values obtained for growth at 620°C). The deep levels responsible for this semi-insulating behavior are believed to be caused by the presence of excess arsenic (up to 1%) in the resulting epilayers, which form AsGa antisite and gallium vacancy defects and complexes, as well as arsenic precipitation.f322)f416) Deep-level transient spectroscopy (DLTS) measurements performed on low temperature GaAs layers revealed deep electron traps with activation energies (capture cross sections) of 0.28 eV (5 x lo-l5 cm2), 0.45 eV (1.5 x lo-l4 cm2) and 0.53 eV (2 x 1 O-l6 cm2), observed at temperatures of 190 K, 260 K and 350 K, respectively. f416j Hole traps with activation energies of 0.38 eV and 0.52 eV were observed at temperatures of 236 K and 320 K. The trap densities increase by a factor of 4-5 after a 24-hour anneal at 400°C. The electrical and optical properties showed no dependence on the choice of arsenic species used during growth. Annealing low temperature (250°C growth) GaAs at 600°C for one hour produces arsenic precipitates with a density of - 1017 cm-3.f323j Average precipitates sizes of 5 nm and 10 nm were observed for films grown with As, and As,, respectively. Annealed, low temperature GaAs/ AlGaAs heterojunctions are characterized by a narrow (20-30 nm) region, which

is free of arsenic

precipitates.f2g6)

No difference

was observed

between normal and inverted heterojunctions. Simple AlGaAs Buffer. Undoped AlGaAs buffer layers are used to suppress the formation of parasitic two-dimensional electron gases in the adjacent GaAs layers. Such parasitic two-dimensional electron gases change the apparent carrier densities and mobilities of doped AlGaAs layers.t1w)f352) Elimination of these parasitic two-dimensional electron gases is particularly important for the electrical characterization of AlGaAs, its dopants and doping calibration. Many measurements performed on AlGaAs prior to - 1982 are suspect due to the effects of parasitic twodimensional electron gases,

MBE of High-Quality

GaAs and AlGaAs

179

Compositionally-Graded AlGaAs Buffer. Compositionally-graded AlGaAs buffer layers are also used to improve the performance of laser diodes grown by MBE. room temperature

Compositionally-graded

PL linewidth

meV.t173) A compositionally-graded old current

of double

buffer layers reduced the

of a 9 nm quantum

heterojunction

well from 13 to 10

buffer layer also reduced the threshlaser diodes

(70 nm active

layer

thickness, 5 pm stripe width, 250 pm cavity length, and x = 0.45 cladding) by a factor of two from 156-240 mA to 76-l 15 mA. The differential quantum efficiency of the DH laser was 50% with the compositionallygraded buffer layer and only 20% without the graded buffer. Finally, the lifetimes of the lasers with and without the compositionally-graded buffer layer were > 1500 hours and c 10 hours, respectively. Multiple-Quantum-Well and Superlattice Buffers. Multiple-quantum-well and superlattice buffer layers effectively getter impurities at the heterointerfaces, thus reducing the background impurity levels in subsequent layers and reducing the microscopic surface roughness. Strained superlattice buffer layers can be used to suppress the propagation of dislocations up into the active layers.1 347j Unstrained GaAs/AIGaAs buffer layers can also prevent the propagation of the dark spot cellular structure of the substrate luminescence (associated with the substrate dislocation networks) into the active epitaxial device layers.t523j Superlattice and multiple-quantum-well buffer layers are most useful when atomically-flat epitaxial layers are required and when electrical currents propagate parallel to the plane of the epitaxial layers. On the other hand, device considerations (e.g., series resistance) occasionally preclude the use of buffer layers with abrupt heterojunctions. Gettering of impurities at GaAs/AIGaAs heterojunctions is now a well established phenomenon. Accumulation of oxygen, AIO, and GaO is observed by SIMS at GaAs on AlGaAs interfaces.t4] No discernible accumulation occurred at the AlGaAs on GaAs interfaces. The oxygen accumulation aluminum. AI,Ga,_yAs

was most severe for epilayers

containing

more than 35%

This accumulation of oxygen also occurs at AI,Ga,& pseudo-alloys interfaces, where x < y. t3] Superlattice

on also

contained 2-3 times less oxygen than comparable random alloy layers.t3j (These accumulation layers are apparent in profiles starting both from the surface and from the substrate, clearly demonstrating that the accumulation is not a SIMS-related artifact. This oxygen accumulation probably causes increased surface roughness, as discussed in Sec. 6: Surface Morphology and Interface Roughness.)

180

Molecular

Beam Epitaxy

Accumulation

of impurities

in the first few

superlattice is also observed with photo1uminescence.t Careful consideration

of the energy and lineshape

quantum

wells

of a

149][174][297l[336][406][482]

of the impurity-related

quantum well luminescence suggests that carbon is responsible for this extrinsic impurity 1uminescence.t 338) Photoluminescence measurements on superlattice buffer layers with different well widths show that most of the silicon and carbon impurities accumulate in the first quantum well.t174] There is a twofold reduction in the carrier concentration of the undoped AlGaAs above and below the superlattice. In addition, the 10 K boundexciton luminescence linewidth of AIo,sGac7As drops from 7 meV below the superlattice to 3.8 meV above the superlattice. Superlattices are very effective for removing the microscopic roughness (typically 3-5 nm) of initial growth surfaces. This smoothing effect is also effective for thick AlGaAs layers, which become rough as the layer thickness increases.t406] Most of the surface smoothing occurs in the first 3-5 quantum wells of a GaAs/AIGaAs superlattice. This smoothing effect is believed to be the result of reduced contamination of the growth surface and a corresponding reduction of step pinning during two-dimensional growth. Superlattices are also effective for smoothing the surface during MBE growth on (21 l)A and (21 l)B GaAs surfaces.t482) Submicron faceting is noticeably reduced for (21 l)B growth and eliminated for (21 l)A growth, when a superlattice buffer layer is used. This surface-smoothing technique is useful for obtaining narrow, uniform quantum well luminescence linewidths, high current densities in double barrier resonant tunnel diodes, and very high mobility two-dimensional electron gases.t115~[127~~~21~[45sl[4721 Superlattice buffer layers have been used to improve the performance of both laser diodes and very high mobility two-dimensional electron gases. Introduction of a five-layer GaAs/AIGaAs (15 rim/l 5 nm) superlattice buffer reduced the threshold current density of a GRIN-SCH laser from - 230250 A/cm2 to - 190-205 A/cm 2.[14g] The superlattice buffer also reduced the 4.2 K PL linewidth of a 5 nm quantum well from - 14 meV to - 7 meV, and eliminated the quantum well acceptor bound-exciton signal. Replacing the ternary AlGaAs barriers of a quantum well with short period GaAs/ AlAs superlattices has resulted in improved quantum well luminescence efficiency and a six-fold increase in the recombination lifetime.t150)t477j These results are consistent with the observed gettering of impurities at the heterojunction interfaces. A ten-period AIAs/GaAs (2.5 nm/2.5 nm) superlattice, and a total buffer layer thickness of 0.1 pm, resulted in two-dimensional electron gas

MBE of High-Quality

GaAs and AlGaAs

181

mobilities as high as 7.1 x 1O5 cm*/ Vs at 5 K, after illumination and with an 18 nm undoped AlGaAs spacer.11461 While this mobility is impressive for such a thin buffer layer, the results were better for a simple 2.0 pm thick GaAs buffer layer (same study).

This is consistent

with other results where

a series of 100-350 nm thick GaAs and AI,,3GacTAs layers followed by a 1 pm GaAs layer (total buffer thickness of 2.65 pm) was used to obtain a two-dimensional electron gas mobility of 5 x lo6 cm*/ Vs at 2 K, under illumination and with a 75 nm undoped AlGaAs spacer.t’*q The mobility of the two-dimensional electron gas in “inverted” GaAs on AlGaAs modulation doped structures

is much more sensitive

to interface

roughness

and

interface contamination than in “normal” AlGaAs on GaAs structures~~~~~1~~*~1~~~~1~~~~1~~~~1~~~*1~~~~1 St rong mobility enhancement in inverted modulation doped field effect transistor (MODFET) structures has been achieved by (i) lowering the growth rate to -0.24pm/hr; (ii) interrupting the growth periodically under an arsenic flux (preferably with a monolayer of GaAs) to allow surface migration and smoothing; (iii) lowering the aluminum composition below that used in normal MODFET structures; and (iv) lowering the growth temperature during the growth of the undoped AlGaAs spacer.t321)t456j The use of low temperatures for the undoped GaAs layer minimizes the surface segregation of silicon from the doped AlGaAs layer immediately below it. Migration-Enhanced Epitaxial Buffer. Buffer layers grown by migration-enhanced epitaxy (MEE) can be used to suppress the formation of crystallographic defects, as discussed in detail in Sec. 5: Oval Defects. 7.5

Choice

of Arsenic

Species

(ASH,, AS*, As,)

The type of arsenic source used affects the quality of MBE films. The most commonly used arsenic source is tetrameric arsenic (As4), sublimed from solid arsenic. However, arsine gas (ASH,) and dimeric arsenic (As*) are now becoming widely used. Dimeric arsenic is obtained either from the sublimation of GaAs or by thermally cracking tetrameric arsenic from a solid source. Since most of the results reviewed in this chapter pertain directly to MBE growth with tetrameric arsenic, only results obtained with arsine and dimeric arsenic will be discussed in this section. Dimeric arsenic is incorporated via a simple first-order dissociative process.t14 Tetrameric arsenic, on the other hand, is incorporated via a second-order process in which two tetrameric arsenic molecules weakly chemisorb to adjacent gallium atoms on the GaAs surface. These molecules

182

Molecular

then

dissociate

evaporate.t142]

Beam Epitaxy

into four dimeric The two remaining

dergo simple dissociation Currently,

arsenic

molecules,

dimeric

arsenic

arsenic

re-

then un-

and incorporation.

most dimeric arsenic is obtained

special two-stage

two of which molecules

cracking

furnace

by cracking arsenic in a

(cracker). Solid arsenic

is

loaded into the low-temperature stage, which is responsible for setting and controlling the magnitude of the arsenic flux. The resulting tetrameric arsenic flux passes into the high-temperature stage, where a combination of thermal and catalytic cracking 0ccurs.t 155~~51 The choice of catalytic materials in the high-temperature stage determines the temperature at which this stage must be operated to obtain dimeric arsenic. A more complete discussion on catalytic cracking of As, into As, is provided in Sec. 2.2: Impurities Generated by Hot MBE Components. Dimeric arsenic increases the tendency of amphoteric dopants to become incorporated on the gallium sublattice as donors.flscl Dimeric arsenic also inhibits the re-evaporation of gallium from the substrate surface. With tetrameric arsenic, the G&s growth rate drops from 1 .OO pm/hr to - 0.92 pm/hr as the substrate temperature increases from 570°C to 680”C.t130t However, when dimeric arsenic is used, no noticeable change in the growth rate occurs for growth temperatures as high as 710”C.t1301t13gl Gallium re-evaporation is probably inhibited by the higher sticking coefficient of dimeric arsenic) 3721 along with the subsequent reduction in the amount of free gallium adsorbed to the surface.t13g] Stronger RHEED oscillations are observed for growth with dimeric arsenic than for growth with tetrameric arsenic.t3721 In addition, the RHEED oscillations persist longer for growth with dimeric arsenic, evidence of increased arsenic

planarity

during

also dramatically

MBE growth reduces

with dimeric

the number

defects.t13g~t313tt432]The surface roughness

arsenic.

Dimeric

of gallium-related

observed

for AlGaAs

oval growth

in the temperature regime from 600-68O”C with As, does not occur when As, is used.t13gl Dimeric arsenic may reduce the gallium surface population and gallium surface mobility, preventing gallium agglomeration and oval defect formation.t13gtt313t

On the other hand,

a reduction

in the

gallium surface population enhances the incorporation of impurities on the gallium sublattice. The use of dimeric arsenic during MBE growth of GaAs/AIGaAs structures reduces the GaAs/AIGaAs interface recombination velocity and increases the radiative efficiency of the Ga~s.t’~~1[~~1 The increased radiative

efficiency

of GaAs

is consistent

with

observations

that the

MBE of High-Quality

GaAs and AlGaAs

183

concentrations of the Ml, M3, and M4 deep electron traps are smaller when As, is used.13671Initial reports suggest that the defect-induced bound excitons (DIBE) can be eliminated through the use of dimeric arsenic.t258)t25g) However, the elimination

of the DIBE could also be due to fewer impurities

in the dimeric

since

arsenic,

the initial

sourcet258jt25g) and the subsequent

study

used GaAs

as the As,

work was performed with cracked solid

arsenic.(114j For example, the GaAs source probably produces less oxygen contamination than the cracked arsenic source, because the solubility of oxygen in GaAs is quite small (= 1016 cme3). This possible difference in the As, purity could change the incorporation of the DIBE. Recently, the use of dimeric arsenic (from cracked As,) has resulted in very high purity GaAs with 77 K electron mobilities as high as 195,000 cm*/ Vsec.tgl) In fact, the highest electron mobilities observed in GaAs (402,000 cm*/ Vsec at 28-40 K for N, - N, - 3 x 1013 cme3) were observed in MBE samples grown with dimeric arsenic.t473j (The principal background donors were silicon and sulfur. The silicon concentration decreased as the cracker temperature decreased, but the precise source of the silicon impurities was not determined.) It is clear from these two separate results that superior material can be obtained with cracked arsenic. 7.6

Role of Growth Temperature

The temperature of the substrate during molecular beam epitaxial growth of GaAs and AlGaAs impacts all aspects of the epitaxial crystal quality. The temperature influences the incorporation and redistribution of impurities, deep levels, and lattice defects. These impurities, deep levels, and lattice defects directly affect the electrical and optical properties

of the

semiconductor. The substrate temperature also influences interface roughness and surface morphology. Thus, the optimal growth temperature depends on the application of the particular epitaxial structure (as well as the cleanliness ambient).

of the molecular

For example,

beam epitaxial

heterojunction

sources

bipolar transistors

and vacuum (HBTs),

metal

semiconductor field effect transistors (MESFETs), and modulation doped field effect transistors (MODFETs) are often grown at 560--62O”C, to minimize dopant diffusion, obtain atomically smooth heterojunctions or maximize carrier mobilities. On the other hand, laser diodes, solar cells and charge-coupled devices are often grown at 680-750°C to reduce nonradiative recombination rates and deep-level densities, particularly in

184

Molecular

Beam Epitaxy

AlGaAs epilayers. At high substrate temperatures (z= 640°C for GaAs based structures and > 480°C for InGaAs based structures), group III desorption also becomes important. Since many devices must be grown under these conditions, are becoming

problems due to substrate temperature

increasingly

uniformity

important.p3tf522)

The purpose of this subsection

is to provide a working knowledge

of

the impact of growth temperature on various aspects of GaAs and AlGaAs material quality. The material characteristics considered in this subsection include: the incorporation of background impurities; the incorporation and redistribution of dopants, including surface segregation; the incorporation of deep levels; luminescence efficiency; clarity of the fine structure for various luminescence features; surface and interface roughness; GaAs decomposition/desorption; and bulk carrier mobilities. The results discussed in this subsection represent the work of numerous researchers, using different reference temperatures and temperature measurement techniques. While most of these techniques are fairly reliable for determining relative temperatures, absolute temperature calibrations can be in error by more than 5O”C.* New reflection and absorption temperature measurement and calibration schemes which tie the temperature measurement to the GaAs band gap energy offer the ability to reliably compare temperatures measured by different researchers in different MBE systems.t187tt2771t278)Photoluminescence has also been used for temperature calibration,t437t but the luminescence spectrum is very sensitive to impurities and the properties of thin epilayers (e.g., InGaAs quantum wells). Incorporation of Background Impurities. The incorporation rate of many background impurities is a strong function of the substrate temperature during growth. As the substrate temperature increases, high vapor pressure

impurities

these species.

On the other hand, higher

tend to desorb,

reducing substrate

the incorporation temperatures

of can

result in cracking or decomposition of complex molecular species, thus increasing the incorporation rate of one or more of the constituent species. The incorporation rates of background calcium, magnesium, carbon, sulfur,

manganese,

and iron impurities

are strongly

temperature

dependent, as is the incorporation rate of Ga,O. SIMS measurements have shown strong decreases in the incorporation rates of calcium, magnesium, iron, and manganese as the substrate temperature increased

We quote the authors’values of temperature, but note that the temperature calibration of various researchers may be quite different. We have not attempted to correct these calibration differ-

l

WlCeS.

MBE of High-Quality

GaAs and AlGaAs

185

from 480°C to 650”C.[g4] In one early experiment, the calcium concentration dropped from 2 x 1014 cm-3 to c 8 x 1012 cmm3. The magnesium concentration

dropped

from 2 x 1016 crne3 to 2 x 1013 crnm3. The iron

concentration

dropped

from 5 x 1015 crnm3to c 1 x 1014 cm-3, and the

manganese concentration dropped from 2 x 10’ 5 cm-3 to < 4 x 10’ 3 cmm3. (These impurities are generally not observed in current MBE materials.) As the substrate temperature increased from 620°C to 650°C a reduction in the free electron concentration of lightly silicon-doped GaAs was observed.t205) This reduction was attributed to an increase in the formation of silicon acceptors, however, in retrospect, it is probably the result of increased sulfur desorption. (Sulfur is the primary n-type background impurity in MBE GaA~.)t~~~)t~~~l The intensities of the carbon acceptor-related photoluminescence transitions show a U-shaped temperature dependence.t256)t4g7) Intensities of the carbon-related luminescence transitions increase as the substrate temperature decreases below 55O”C.t 2s6)ts26) However, as the substrate temperature is increased above 600°C the intensities of the carbonrelated luminescence transitions increase again.t256)t4g7] The minimum carbon acceptor luminescence is obtained for a substrate temperature of 580°C. Similarly, low-temperature PL measurements of AI,Ga,_& (0.20 5 x 5 0.33) show a decrease in the carbon acceptor luminescence as the growth temperature is increased from 800°C to 870”C.t106) Since the carbon acceptor luminescence intensity is proportional to the carbon concentration,t2se)t45a) and since carbon incorporates almost exclusively on the arsenic sublattice,t366] we conclude that the carbon concentration in GaAs is minimized for growth at 580 f 20°C. In AlGaAs, however, the background carbon concentration is not only higher than in GaAs, but is minimized at much higher temperatures, possibly due to an increase in the carbon sticking coefficient. The incorporation rates of impurity-related defect complexes also depend

on the growth temperature.

For example,

the strengths

of the

GaAs defect induced bound exciton (D/BE) photoluminescence transitions* (1.504 eV to 1.511 eV) decrease when the substrate temperature *A series of sharp, low-temperature luminescence transitions in GaAs from 1.504 eV to 1.511 eV were first reported in 1980. These transitions were attributed to excitons bound to lattice defects and given the name defect induced bound excitons (DIBE).[*~~~[*~~] A second series of transitions from 1.471 eV to 1.491 eV was reported in 1982 and were designated defect comp/exes.[55al The energies of the defect complexes are related to the energies of the DIBE transitions by Haynes’ rule, but with a proportionality constant of 0.38--somewhat larger than the

MBE of High-Quality

GaAs and AlGaAs

187

aluminum percentage approaches zero (x = > 0). Recent local vibrational mode infrared absorption studies suggest that this 0.4 eV deep level is GasO, with the oxygen situated interstitially. [171[44s)Thus, it appears that the incorporation

of Ga,O

in both GaAs and AlGaAs

MBE growth temperature increases. Incorporation of Dopant Impurities.

decreases

The substrate

as the

temperature

during MBE growth also affects the incorporation kinetics and redistribution of GaAs dopants. For many device applications (e.g., MODFETs), high purity material must be grown in close proximity to intentionallydoped material. Thus, it is important to have good control over the incorporation and redistribution of the dopant impurities. Dopants which are, or have been, used for MBE growth include beryllium, germanium, silicon, tin, sulfur, magnesium, manganese, chromium, and erbium. The substrate temperature affects the re-evaporation, surface segregation, site incorporation (for amphoteric dopants), and precipitation (at high doping densities). The behavior of dopants with respect to growth temperature restricts the choice of viable MBE doping species to beryllium and carbon as acceptors and silicon and tin as donors. In fact, diffusion and surface segregation of both beryllium and tin make carbon and silicon the most desirable choices for acceptor and donor dopants for MBE GaAs and AlGaAs.* The difficulty in producing a reliable carbon source has limited its use to trimethylgallium in MOMBE. However, carbon filament sources are becoming available for conventional solid source MBE. Almost all of the p-type doping in MBE to date has been with beryllium, because of the lack of a reliable solid carbon doping source. Low-energy accelerated-ion doping can be used for dopants which have low sticking coefficients (e.g., Zn, Mg, S) or which Te) _[s11WW261[‘W

exhibit

strong

surface

segregation

(e.g.,

Se,

Beryllium. A beryllium delta doping study reveals beryllium surface segregation in AI,,,,Ga,,,,, As.[‘~~) (Delta doping, also known as planar doping, is a heavily doped layer deposited during a growth interruption.)

In AI,Ga,_,As, n-type dopants form unusual deep levels called DX centers.f~O4l12161P171Ps~l~~ssllss~l~~sslls~l DX centers have metastable states which are responsible for persistent photoconductivity. The DX center concentration increases with increasing aluminum molefraction and also with increasing donor doping levels. The energy of the DX center depends on the specific donor species. For similar doping levels and AlGaAs compositions, the energies of the DX centers associated with different donors increase in the order Se, Te, Sn, and Si. The relatively deep DX center for silicon can be a significant disadvantage for the use of silicon doping in MODFETs and other devices. f2’s1fs5slA significant improvement in the low temperature performance of MODFETs has been reported using selenium-doped structures. l

188

Molecular

The beryllium

Beam Epitaxy

concentration

decreases

exponentially

toward the surface

with a characteristic length, I,. As the growth temperature increases from 500°C to 600°C I, increases from 2.8 nm/decade to 5.3 nm/decade. Beryllium

surface-segregation

AI,s,Gacs,As.

Beryllium

is significantly accumulation

smaller

occurs

in GaAs than

at GaAs-AIGaAs

in

inter-

faces/69] consistent with the enhancement of beryllium surface-segregation during AlGaAs growth. The diffusion coefficient of beryllium increases dramatically at high concentrations (> 2-5 x 1Olg cm9) ,t12glt206)t316)t4001 making it difficult to obtain high doping concentrations. The increase in the effective beryllium diffusion coefficient is due to a rapid increase in the amount of rapid interstitial beryllium diffusion. (At 680°C the diffusion coefficient of interstitial beryllium is greater than 6 x lo-lo cm2/sec.)t12g) Several points of reference data for beryllium diffusion are given in Table 3. Beryllium interstitials diffuse slower in n-type material than in undoped or p-type material because the electric field of the p-n junction opposes the diffusion of the positively charged interstitial beryllium.[12gl Efficient incorporation of beryllium at high concentrations can be achieved by lowering the growth temperature (2.7 x lO*O cm3 at 480°C) and (1 x lO*O cm-s at 520°C) ,~*s’1~~~*1~~~‘1~~~~1~~~~3 F or some devices (e.g., heterojunction bipolar transistors), reduced growth temperatures and increased V/III ratios are also found to strongly reduce the degradation which is induced by thermal and electrical stress.t47gl The introduction of indium also reduces the beryllium diffusion c0efficient.t 5011 For example, the introduction of only 7% indium into heavily beryllium-doped ([Be] = 7 x 1Olg cm9) Al,,, GacgAs grown at 600°C reduced the beryllium diffusion coefficient by a factor of five. Beryllium redistribution depends on the substrate orientation and is suppressed

by as much as a factor of three during

substrates.t344j This suppressed

beryllium redistribution

growth

on (31 l)A

is attributed to the

large surface step density, which is believed to inhibit the incorporation beryllium interstitials.

Table 3. Effective

[BeI (cm-?

Beryllium

Diffusion

D(600°C)

Coefficient D(800”C)

1.5 x 10’7 l-2

x 10’9

10lg

D(900”C) 5-10 x lo-16

3-5 x 1O-16

3 x 10’9 26x

(in cm*/sec)

5-10 x 10-15 > lo-‘*

5-10 x 10-14

of

MBE of High-Quality

Carbon.

Carbon diffusion

in GaAs is negligible

large as l@O cm3 and growth temperatures diffusion

coefficient

of carbon in undoped

825°C.t105) It increases negligible

carbon

to 2.3 x lo-l6

diffusion

GaAs and AlGaAs

189

for doping levels as

as high as 700”C.t360)

The

GaAs is 1 .O x 1 O-l6 cm*/sec

cm*/sec

occurs in n+ GaAs.

in p+ GaAs at 825°C The dependence

at but

of the

carbon diffusion coefficient on the background doping of the material is caused by the junction electric fields at the diffusion front. Results with resistively heated graphite filaments show well-behaved CAs doping levels up to 102O cmm3, without the introduction of unusual deep levels .I1Q~l~~~~l~~~~1~~~~1~~6oj C ompared with comparably beryllium-doped GaAs, recent experiments show an enhancement in the non-radiative recombination processes of heavily carbon-doped GaAs ([Cl > 1 Olg cm-q .t360)t1g4)t1121 The hole concentration and the hole mobility begin to drop for carbon doping levels above lO*O crnm3, possibly due to C, and/or C, incorporationtssO1or to lattice strain. The maximum delta-doping density obtained with carbon is 3.8 x lo’* cm-2.t360) Currently, the reliability of the graphite filament is poor for doping levels above - 3 x 10lg crnm3, but improvements are being made in the design of these graphite filament sources. Recent results with carbon filaments have shown a small, short term (l-2 growths) memory effect for carbon doping.t361) This memory effect may be the result of hydrogen or other impurities in the graphite. (Such impurities may form or escape as volatile carbon species such as CH,.) It is our belief that this memory effect will be reduced or eliminated. For many important device applications (i.e., laser diodes, HBTs, etc.) a temporary change in the shallow impurity background is less important than the problem of obtaining improved control over the p-type dopant distribution. Germanium.

Germanium

is an amphoteric

dopant,

incorporating

both as a shallow acceptor on the arsenic sublattice and as a shallow donor on the gallium sub1attice.t 4991 Germanium is incorporated primarily as an n-type dopant at low growth temperatures with a gallium-stabilized surface, and primarily as a p-type dopant at higher growth temperatures with an arsenic-stabilized surface.t8s) Germanium also increases the ionized impurity background concentration of epilayers grown in an MBE machine in which germanium has been previously evaporated.t177) As a result of these effects, germanium is no longer commonly used during MBE growth. Silicon. Silicon is an amphoteric dopant, incorporating preferentially on the gallium sublattice (100) oriented

substrates

during MBE growth of GaAs and AlGaAs on

for a wide range of growth

conditions.

The

190


maximum free-electron concentration achieved with silicon is usually - 7 x lo’* cm3, but electron concentrations as large as 1.6 x 10IQ cm” are possible

using low growth temperatures,

tively-heated

silicon

higher V/III ratios, and a resis-

source.t3021t330)t331tt3681

Low-temperature

photoluminescence measurements of the (e,Si,) and (D,Si,) transitions show that the relative incorporation of silicon acceptors increases as the substrate temperature increases,t 3241t4581 even for growth with As,.[‘~~] The diffusion of silicon at typical MBE growth temperatures can usually be neglected.1 1361At 105O”C, the diffusion coefficient of silicon in GaAs is - 2 x 10-l* cm*/sec.t ‘591 In AI,,Ga,,7As at 7OO”C, the diffusion coefficient of silicon is - 3 x lo-l4 cm*/sec for large silicon concentrations ([Si]>6x101*cm 3 ).t 15*1 As the silicon concentration increases above 2 x lo’* cm3, Siti,-SiAs pair formation (and even precipitation) begins to occur, and for [Si] > 6 x lo’* cm-3 the diffusion coefficient increases The diffusion rapidly, due to Si,-Si, p air ~~~~~~~~~~~~1~~31~~5~1~~~Q1~~~~1-~~~~1 coefficient of Si,-SiA, pairs is given by: D = 0.1 lexp(-2.5/kl). At large silicon concentrations, the formation of SiAs-SiGa pairs and silicon precipitates is sensitive to the growth temperature, with higher electron densities achieved as the growth temperature is decreased from 620°C to 520”C.t331) Silicon also surface-segregates, particularly during the growth of AlG~s.~‘~~~~~~~l~~~Q1~*~*l~~*~l~~~~l Th e amount of silicon surface segregation is small compared to tin and escaped detection in early silicon doping studies.t1361 Silicon surface segregation is one of the primary causes of poor two-dimensional electron gas mobilities for structures in which the GaAs channel is grown on top of the doped AlGaAs layer.[176~t212)t321)t4541 Lowering the substrate temperature reduces the amount of silicon surface segregation.[166~[176~~2t2~~3211[40*~[454~ The silicon surface segregation about the same in both GaAs and Delta-doping

rate is

A&&ao,67k[‘67~

is often used to obtain large electron concentrations,

to

create planar-doped barriers and to reduce coulombic scattering of charged carriers. Electron sheet densities of 3 x 1013 cm-* can be obtained with silicon planar doping.t4**) SIMS profiles of silicon planar-doped barriers show asymmetric diffusion tails toward the surface and the bulk for doping below 1.25 x 1013 cm-2.t166)t422)t30)The silicon distribution layers can be described by two exponential concentrations

of delta-doped with character-

istic lengths 1, (leading or surface side) and IT (trailing or substrate side) ,tQQ1trQQ1trQ71 L ength IL is caused primarily by silicon surface segregation and increases with increasing growth temperature with values of 1.4, 2.2, 4.5, and 11.6 nm/decade at 390, 490, 550, and 600°C. Length 1, is

MBE of High-Quality

191

GaAs and AlGaAs

independent of temperature with a value of - 4.0 nm/decade and results from diffusion and/or from the thermodynamic redistribution of the near-surface layers during the growth of the delta-doped diffusion set

coefficient

at 55O”C, and increases

Substantial

region.

for 1013 cm-2 silicon delta-doped for higher

The effective

delta doping

diffusion toward the substrate

silicon

GaAs is - 1O-l6 cm2/ of 5 x 1013 cm-2.

is observed for doping densities

above 1.25 x 1013 cm-2.f422) On the substrate side of planar-doped barriers, with concentrations above 1 x 1014 cm-2, a two-stage concentration profile is observed with a well-defined knee at the point where the concentration reaches 3-4 x 10’s cmm3. Tin. Even though tin is a group IV element, it is incorporated as a shallow donor and displays almost no amphotericity during MBE growth on (100) GaA~.t~~~l Tin displays substantial surface-segregation during MBE growth of GaAs and AlGaAs. 83If412) Tin surface-segregation increases as the substrate temperature increases from 500°C to 615°C with the characteristic length, IL, increasing from - 50 nm/decade to - 500 nm/decade.t83)f16r) The surface-segregation of tin leads to the following expression for the tin incorporation rate (for constant arsenic and tin fluxes): K = K, exp{-1 .35eV/kT).t542) (K,, is a constant which depends on other factors including the substrate orientation, the V/III ratio, and the growth rate.)t165) The diffusion of incorporated tin also increases significantly as the substrate temperature increases from 550°C to 615”C.tE3] The diffusion coefficient of tin is described by D = D, exp(-Q/kT), where D, - 6.0 x 10m4 to 3.8 x 10m2cm2/sec and Cl = 2.5-2.7 eV.t2281 However, tin also diffuses interstitially, with an interstitial diffusion coefficient of - 1O-8-1O-g cm2/sec at 835°C. Tin desorption becomes important for substrate temperatures above 620°C.

These surface-segregation

and desorption

problems

have

precluded the use of tin for uniform doping and limited its use to achieving very high surface doping for ohmic contact layers. Selenium and Tellurium. Selenium and tellurium form shallower DX centers in AlGaAs than si1icon.f 2161 Maximum electron concentrations of - 10lg cm” are obtained with selenium selenide

and gallium

telluride

and tellurium

precipitation

doping, but gallium

occurs at higher

doping

lev-

e1s.t’) The desire to find a column VI source which might be useful for both moderate and high doping levels has led to investigations of tin-telluride (SnTe), lead sulfide (PbS), and lead selenide (PbSe) dopant sources.tggl[541) The incorporation of Sn and Te shallow donors from a SnTe source decreased from - 1 to - 0.1 as the growth temperature increased from 550°C to - 600”C.tgg) Surface segregation of both tin and tellurium was

192

Molecular

Beam Epitaxy

observed for substrate temperatures above 530°C. The diffusion coefficient of selenium is described by Ds, - 3 x lo3 exp(-4.1 6eV/kT).t228t The diffusion coefficient of tellurium is - lo-l3 cm2/sec at 1000°C and - 2 x lo-l2 cm2/sec at 1100°C selenium

suggesting

that tellurium

diffusion

is very similar to

diffusion.

Sulfur.

Re-evaporation

of S, is important

for substrate

tempera-

tures above 590”C,t201 making sulfur a poor dopant choice for nearly all MBE applications. The high vapor-pressure of sulfur is also expected to increase the n-type background doping levels of Ill-V MBE epilayers. The diffusion coefficient

of sulfur is 2 x 1 O-l4 to 2 x 1O-l3 cm2/sec at 81 0”C.t312t

Others report much larger sulfur diffusion coefficients of - 7 x lo-l4 cm2/ set at 600°C and 3-8 x 1O-l3 cm2/sec at 700”C.t228~ Magnesium. Re-evaporation reduces the sticking coefficient of magnesium from 1 0m2to 1O-5 as the MBE growth temperature is increased from 450°C to 580”C.t8*1tg4t Similarly, the incorporation of manganese is dominated by desorption for MBE growth temperatures above 500”C.[542t The low, temperature-dependent sticking coefficients of magnesium and manganese make them poor p-type dopants for MBE. The diffusion coefficient of magnesium is - 3 x lo-l6 cm2/sec at 600”C.t228] Chromium. Chromium is sometimes used to make GaAs semiinsulating. The solubility limit of chromium in GaAs increases as the MBE substrate temperature increases from 500°C to 640X, while the amount of chromium needed to render the MBE GaAs semi-insulating decreases.t3481 Erbium. Erbium doping of GaAs is technologically interesting because it produces very sharp luminescence transitions at - 1.5 pm, a low-loss wavelength for fiber optic communication. The intensity of the 1.54 eV erbium photoluminescence peak is maximized for growth temperatures between 570°C and 590”C.t461] The number of erbium-related luminescence

transitions

decreases

when the MBE growth temperature

dramatically

to eight sharp transitions

is decreased from 600°C to 580”C.t128]

The photoluminescence efficiency of erbium-doped AlGaAs is an order of magnitude stronger than that of erbium-doped GaA~.tl~l This improvement in the erbium luminescence

in AlGaAs

is attributed

to gettering

by

the aluminum, allowing more efficient incorporation of Ers+ centers. Erbium doping concentrations of 1.5 x 10lg cm3 are reported.t132] Rapid damping of RHEED oscillations occurs for [Er] > 4 x 10lg cm”, suggesting that the two-dimensional growth processes are degraded. Erbium diffusion is observed for temperatures as low as 500°C.

MBE of High-Quality

GaAs and AlGaAs

193

Incorporation of Deep-Level Defects. Substrate temperature also plays a key role in the formation and incorporation of deep-level defects. Changes in the substrate temperature levels by changing availability

can affect the incorporation

the sticking coefficient

of lattice defects

(vacancies,

of impurities antisites,

known as Ml, M2, M2’, M3, and M4 are commonly

of deep

and by altering the

etc.).

Electron

traps

found in GaAs grown

by MBE.f40) Occasionally, other electron traps known as MOO, MO, M5, M6, M7, and M8 are also observed.t11’)t262) The general characteristicst40)f51)f111)t262)t260)t306) of the MOO-M8 traps are listed below in Table 4. The concentration of traps Ml and M4 decrease by - 100X (i.e., from - 1015 cm-3 to - 1013 cm-9 as the substrate temperature increases from 520°C to 650”C.f40] The concentration of trap M3 also decreases by - 1 OX (i.e., from - 1014 cm” to - 1013 cme3) as the substrate temperature increases from 550°C to 650°C but the concentration of M3 also decreases for growth temperatures below 550°C. The concentration of M2’ also decreases as the substrate temperature increases from 520°C to 6OO”C, but the concentration of M2 increases for temperatures above 650°C. DLTS depth profiles of the Ml-M4 traps in GaAs grown on silicon substrates suggest that these traps are caused by impurities interacting with lattice defects.f72) The specific impurities responsible for these traps have not been identified, but they often originate in the arsenic source.

Table 4. G&s

Trap

Deep Levels

T eak (p 725°C for

AlGaAs (unless other effects, such as interface roughness or surface segregation/diffusion of dopants limit the maximum acceptable growth temperature).

Table 6 shows the temperature

Table 6. Growth Temperature

Dependence

dependence

of AlGaAs.

of Deep Levels in AI,,,Ga,,eAs*

Trap

660°C

720°C

780°C

ME4 ME5

3 x 10’4 1 x 10’5

5 x 10’2

( 0.2), the longitudinal optical dispersion curves for the CdTe and (Cd,Mn)Te constituents in a superlattice are expected to show only partial overlap, thus allowing for the possibility of optical phonon confinement effects. For the (lOO)-oriented CdTe/Cd,,Mn,Te superlattices, such confinement is readily observed, as illustrated in Fig. 15, for a structure of 57/96 A well/barrier layer thickness ratio and x = 0.24 Mn ion concentration.f4e) In sharp contrast, confinement effects have not been seen in the (ill)-oriented structures; rather, the Raman spectrum

374


contains details which have been identified as interface optical phonon modes.t48) These observations complement those made through photoluminescence entations

and excitation

(see above)

spectroscopy

and lend credence

for the two different

to suggestions

ori-

that interface

geometry for the CdTe/(Cd,Mn)Te superlattices is strongly dependent on orientation in the MBE growth. As an interesting further use of the Raman probes, Ramdas and co-workers have examined magnetic excitations in the (Cd,Mn)Te system.f4r)t48) Among other things they have shown that while for bulk (Cd,Mn)Te it is possible to observe the transition from a paramagnetic to a magnetically-ordered Mn concentration, Raman scattering

phase (magnon) with increasing at least from the (11 l)-oriented

(Cd,Mn)Te superlattices, shows the absence of the ordering. This may be due to a combination of effects due to lower dimensionality and interface structure frustrating ordering tendencies, not unlike the scenario lated for the ultrathin ZnSe/MnSe superlattices discussed below.

I

“.” l36- ~~~

150

z(xx)i

170

‘i/

190

specu-

I

210

RAMAN SHIFT (cni’) Figure 15. Raman spectrum from optical phonons in a (lOO)-oriented CdTe/ Cd,,,Mn,,,Te superlattice for different polarizations. Confined optical phonons are labeled with n = 2, 4, and 6. Inset shows the observed frequencies of the confined LO phonons plotted on the bulk CdTe dispersion curve calculated with a linear chain model.

Wide Gap II-VI Semiconductor

Heterostructures

375

Stimulated Emission in (Cd,Mn)Te Quantum Wells. Further evidence of the optical quality of the material system is supported by the demonstration

of stimulated

Cd,_XMn,Te/Cd,_yMnyTe prepared

emission

in both CdTe/(Cd,Mn)Tet551t56]

and

quantum well structures.t571 Laser samples were

using a selective

chemical

etcht34] to remove the GaAs sub-

strate; the resultant free-standing epilayers were cleaved and mounted on a copper heat sinkt58l for optical pumping experiments. For lasers structured with (111) CdTe as the active material, lasing was obtained at wavelengths of 763 to 766 nm at 25 K with a threshold power density of 1.35 x lo4 W/cm2.t5sl When (111) Cd,,Mn,Te (x = 0.19) was the active quantum well material, lasing occurred in the red spectral region at 665670 nm at 15 K; the threshold for laser action was 2.0 x 1O4W/cm2.t57] The exchange interaction occurring in the DMS material allows a shift in the energy of the quantum well states such that the application of a magnetic field would allow for tuning of the output energy of the DMS laser. For the DMS lasers studied, a magnetic field tuning rate of 3.4 meV/Tesla was obtained at 1.9 K.t5gl This magnetic field-induced shift was approximately one-fifth that obtained from bulk Cd, _,Mn,Te having a comparable x value. Quantum wells which are (lOO)-oriented, containing CdTe as the active well material, have also exhibited stimulated emission up to 119 K.t561The effects of strain in these strained-layer superlattices results in TE-polarized stimulated emission from the sample edge, and has been compared with the edge emission from oppositely-strained (Zn,Mn)Se multiple quantum well structures. (This comparison is described in detail in the section on Polarization-Dependent Luminescence below.) Doping in CdTe/(Cd,Mn)Te Quantum Wells.

A longstanding

problem associated with II-VI compounds is their propensity for defect generation, together with an associated self-compensation, as dopant impurity species are incorporated

into the material.

When indium atoms

are incorporated during MBE growth of CdTe, the photoluminescence is degraded, while little evidence of activation is obtained. Bicknell et al.f60] have employed a technique of photo-assisted MBE to overcome this tendency

for self compensation.

By illuminating

the growing

film with a

low intensity beam from an argon laser (150 mW/cm*), a high degree of dopant activation was obtained. A similar success has been achieved for MBE-grown CdTe doped with Sb.t6’] Comparisons of photoluminescence and transport properties have shown dramatic improvements when laser illumination is employed for both n and p doping. The laser-assisted doping technique has also been used during dopant incorporation in CdTe/

376

Molecular

Beam Epitaxy

(Cd,Mn)Te superlattices where the DMS barrier layer is doped with indium.t3*] In the case of relatively wide CdTe wells, the measured mobility values exceeded

that obtained

the wells became narrower,

in single layer, In-doped

however,

the mobilities

CdTe films.

As

tended to decrease,

suggesting that the interfaces may play a role in transport parallel to the superlattice layers. Having achieved the controlled substitutional doping of MBE-grown CdTe, a variety of devices, such as pn diodes and metalsemiconductor field effect transistors, have been fabricated and studied.t@] Although the mechanism by which the photons interact with the growing surface layer during photo-assisted MBE is not understood, this work has generated a great deal of interest. For the readers convenience, Refs. 63 to 69 involve the molecular beam epitaxy of II-VI semiconductors in the presence of photons which is not detailed here. 2.4

Binary ZnTe/CdTe

Superlattices

The binary superlattice CdTe/ZnTetrOj has also been a subject of optical spectroscopy as a possibly useful pseudomorphic structure. The situation, however, involves a very large lattice mismatch (approximately 6.4%). Initial optical studies had suggested a reasonable agreement between experimentally

derived superlattice

bandgap values and calcula-

tions of the free-standing superlattice limit (assuming also the absence of a finite valence band offset).plj A small valence band offset would follow the situation already established with CdTe/(Cd,Mn)Te and ZnSe/(Zn,Mn)Se, in that the finite lattice mismatch strain (through hydrostatic and uniaxial components) may be the main factor determining the actual valence band offsets, thus making them dependent

on individual

sample

parameters.

Apart from the band offset issue, subsequent resonant Raman scattering experiments have, however, cast doubts on the arguments on attaining the free standing superlattice limit in the CdTe/ZnTe system.p*) One important consequence of a small band offset in a real, highly strained however,

system

subject

to small but finite

been established

structural

irregularities

by using the lowest interband

exciton

has, reso-

nances as an indicator in photoluminescence experiments. In particular, time-resolved and resonantly excited spectroscopies have shown that excitons exhibit unusual localization in the CdTe/ZnTe system.t3’j Qualitatively, variations in the layer thickness on a monolayer scale in a highlystrained structure is sufficient to produce significant fluctuations in the local strain about some mean value. In the absence of strong confinement


Heterostructures

(i.e., small offset), the associated random potential quite efficient in capturing electronic quasiparticles

377

fluctuations may be at low or moderate

lattice temperatures. Low threshold,

optically

pumped

lasers,

emitting

orange portion of the visible spectrum and fabricated

in the yellow-

with Cd,,,,Zn,,,,Te/

ZnTe superlattice structures, have been reported by Glass et al.p3] The binary CdTe/ZnTe superlattices are heavily strained, thus the laser structures incorporated alloys of CdZnTe as the well layer to reduce the strain. In the superlattice structures studied, the lasing wavelength increased from 575 nm at 8K to 602 nm at 31 OK. At low temperatures, the threshold pump intensity was found to be quite low at 7 kW/cm2. At room temperature, the threshold pump intensity increased to - 55 kW/cm2 and represented the first report of room temperature, optically-pumped lasing in a IIVI superlattice. 2.5

II-VI Quantum

Wells Incorporating

MnTe Barrier Layers

MnTe/CdTe Single Quantum Wells. The first epitaxial layers of the zincblende phase of MnTe were recently grown by MBE. Whereas bulk crystals of MnTe have the hexagonal NiAs crystal structure and a bandgap of 1.3 eV, the variation of lattice parameter and excitonic bandgap with Mn concentration for zincblende (Cd,Mn)Te epilayers extrapolated to predicted values of 6.34 8, and 3.18 eV,[74] respectively for zincblende MnTe. Especially important is the dramatic increase in bandgap associated with the formation of the zincblende phase. The MnTe growth studies were undertaken primarily in response to the need for a wide bandgap semiconductor in the lattice parameter range of CdTe, but also as a result of the considerable amount of speculation in recent years concerning the expected physical properties of the “hypothetical” zincblende MnTe.[421~5]-~8] The zincblende MnTe has been incorporated into three types of heterostructures: (i) thick epitaxial layers (up to 0.5 pm thick), (ii) single quantum well (SQW) structures where MnTe forms the barrier when layered with CdTe, and (iii) where MnTe is the barrier when layered with ZnTe.p] Cdle substrates having a CdTe buffer layer were used in fabricating the relatively thick epilayers of MnTe. The configurations incorporating CdTe quantum wells were grown on InSb substrates having buffer layers of InSb and CdTe.LeO] The InSb was grown in an MBE chamber dedicated to Ill-V film growth and transferred under UHV to a

378

Molecular

Beam Epitaxy

second chamber for MnTe and CdTe growth. The MnTe epilayers exhibited a (2 x 1) RHEED pattern during the growth of both thick epilayers and quantum well structures.

Following

phase of MnTe was confirmed

growth, the presence of the zincblende by x-ray

diffraction

diffraction peaks obtained from x-ray measurements,

(6-28 scans).

The

performed on samples

containing relatively thick epilayers of MnTe, could be identified as corresponding exclusively to the zincblende phases of CdTe and MnTe. The microstructure of the MnTe epilayers and MnTe single quantum well structures were examined using cross-sectional TEM.~Q)[80) The TEM study indicated that the MnTe layers have the zincblende structure and have formed epitaxially. Both high resolution images and electron diffraction patterns indicate that the MnTe epilayer has a perfect epitaxial relation with the CdTe crystal. In the diffraction pattern, each diffraction spot of the MnTe crystal forms a pair with a spot of the CdTe crystal, as expected from the zincblende structure of the MnTe crystal. A nearly complete relaxation of the tions, spots. MnTe faces

lattice mismatch at the interface, due to formation of misfit dislocais suggested by the degree of separation of the 220 type diffraction TEM examination of single quantum well structures shows that the layers (having thicknesses of 30 to 40 A) maintain coherent interwith the CdTe layers, forming strained-layer structures. A series of strained single quantum well CdTe/MnTe structures were studied with CdTe well thicknesses ranging from approximately 50 A to 10 A. Raman scattering experiments indicate that the structures are pseudomorphic, i.e., that the lattice mismatch of - 3.2% is coherently accommodated if the MnTe barrier layers remain sufficiently thin (typically - 40 A). Under these conditions photoluminescence has been detected from the quantum wells up to room temperature; at temperatures up to 77K strong PL is often visible to the eye even when the photoexcitation is below the bandgap of the barrier layer.fB1) An example of the effects of the strong quantum confinement in these heterostructures, malized amplitudes), with well thicknesses

Fig. 16 shows photoluminescence

spectra

(nor-

obtained from three CdTe quantum well samples approximately 22 A, 15 A, and 10 A, respectively.

Note that the emissions occur in the red, yellow, and blue for the n = 1 exciton recombination in the structures.1 slal[s*l The blue emission (at 4600 A) shows how confinement effects are responsible for increasing the “effective “ bandgap of CdTe by about 1 eV, perhaps the largest confinement effect in a semiconductor

heterostructure

to date.


$

Heterostructures

379

0.6

A C g 0.4

L, - 24A

-E $

0.2 J\ IIll

,I

“1.8

1.9

2.0

,Illl

2.1

2.2

1111

2.6

2.7

Photon Energy (eV) Figure 16. Photoluminescence from three MnTe/CdTe single quantum well samples at T = 10 K. The quantum well thicknesses are approximately L,,, = 22 A, 15 A, and 10 A,respectively. For reference, the bulk bandgap of CdTe is 1.60 eV.

Optical data at the n = 1 transition (HH and LH) have been used as input to arrive at the band offsets. The uniaxial strain induced HH-LH splitting is calculated and imposes another constraint. The end result, namely a very large conduction band offset, allows some simplifications. On the other hand, it also makes the accurate determination of the valence band offset somewhat less reliable. The effect of uniaxial strain in splitting the HH-LH degeneracy as follows:

by identifying

Eq. (14

~EHH

= 2a[(Cl

Eq.(lb)

~ELH

= WC,,

where the deformation

1 - %)/cl

potential

+

the individual

band energy shifts is

dE%

2C,&,1% and elastic constants

have been defined

above, and Em is the in-plane lattice mismatch strain which we take to be equal to the value 3 x lo-*. Under this strain, the light-hole (mi = + 3/2) moves away by 17 meV. (One assumes that the CdTe quantum well layer is unstrained.) The experimentally-observed HH - LH splitting in a sample with L, = 29 A, for example, electron,

one-hole,

is 44 meV.

square well calculation

In order for the simple oneto agree with this splitting,

as

380

Molecular

Beam Epitaxy

well as with the measured samples, the following respectively,

interband

have been obtained:

GE,+,,., = 0.340 meV.tsl]

photoluminescence

band offsets for the conduction,

energies

for all

LH, and HH bands,

6E, PI 1.280; 6E,,

= 0.160 eV, and

These offsets are substantially

larger than those

described earlier in this chapter for the CdTe/(Cd,Mn)Te MQWs and pose the fundamental question about the dependence of the offsets on the Mnconcentration in this system. As part of the analysis for the CdTe/MnTe SQWs, the following HH exciton binding energies for selected samples were obtained from a variational model: L, = 22 A, E, = 28 meV; L, = 25 A, E, = 27 meV; L, = 49 A, E, = 23 meV. These should be compared with the bulk exciton binding energy E, = 9 meV in CdTe.(81) The temperature dependence of the linewidth of the n = 1 HH exciton has been used to estimate the exciton stability against dissociation by optical phonon scattering in a simple model where the LO-phonon induced dissociation/formation rate yields the following temperature dependent exciton linewidth in absorption: 6E(T) = 6E, + r&[exp(I’&kT) -11, where 6E, is the inhomogeneous broadening (from concentration fluctuations at the heterointer-faces) and rLo is proportional to exciton-LO phonon coupling (Frohlich interaction). The fits to experimental data (solid lines) yielded rLo = 40 meV for the typical CdTe/ MnTe SQW samp1es.t 81It81al This value is significantly larger than that measured for GaAs quantum wells (5-8 mev). Most of the difference is directly due to the larger polaron coupling in CdTe; while there appear to be no direct data on exciton-LO phonon coupling constant, or hole-LO phonon coupling, an approximately five times larger coupling for electrons in bulk CdTe is expected. MnTe/ZnTe Single Quantum Wells. A series of single quantum well structures were grown and evaluatedt7g1[80] where thin layers of zincblende MnTe served as barriers for ZnTe. The ZnTe-based SQW structures

were grown on GaSb substrates

with ZnTe buffer layers.

The

ZnTe epilayers were grown using elemental sources at a substrate temperature of 320% with a growth rate of 1.5 &sec. Cross-sectional TEM has confirmed

that MnTe

layers having

the cubic zincblende

structure

have formed epitaxially on the ZnTe epilayers. One interesting aspect of these structures has been the opportunity to study electron (hole)-LO phonon coupling effects in lower dimensions with the Frohlich interaction dominating hot carrier kinetics. Here single quantum wells of ZnTe/MnTe were investigated in structures where the MnTe barriers are so thin (==20 A) that substantial tunneling of photoexcited


381

Heterostructures

electron-hole pairs from the ZnTe quantum well takes place. This has the consequence that thermalized luminescence is strongly suppressed, making hot luminescence effects with LO-phonon structure particularly visible throughout the recombination spectrum, especially at the quantum well lowest exciton resonance. Figure 17 shows an example of the hot photoluminescence

spectrum while comparing

it to emission from a ZnTe

epitaxiai thin film.ts3) in these experiments, strong resonance effects were seen in multiphonon events at least to 14 orders. Qualitatively, with increasing order, a gradual transition from resonance Raman scattering (dominant at the l-LO transition) to hot luminescence in terms of polarization memory effects and linewidth

ZnTe/MnTe

has been observed broadening.

SQW

Ex=2.602

2.45

2.46

2.54 2.51 PHOTON ENERGY (eV)

eV

2.57

2.6

Figure 17. Comparison of secondary emission spectra with multiple LO-phonon sidebands from a single MnTe/ZnTe quantum well sample and a thin epitaxial film of ZnTe (T = 10 K). Note the large resonant enhancement near the n = 1 quantum well exciton resonance. (The bulk bandgap of ZnTe is at - 2.29 eV where thermalized luminescence disallows the viewing of hot photoluminescence or RRS; hence the use of thin barrier layers in our SQW structures).

382

Molecular

Beam Epitaxy

Such higher order hot luminescence both in external

magnetic

spectroscopy.fss)

spectra have been investigated

fields and through

Moderate

external

fields

Zeeman shifts of the quantum well resonance with the magnetic semiconductor

picosecond

time-resolved

(up to 10 Tesla)

show

how

(from g-factors associated

aspect of the problem) cause a commen-

surate shift in the LO-phonon enhancement energy. Following the saturation of the Zeeman shifts, however, only very small subsequent amplitude or spectral changes are seen in the hot luminescence (studied at 5th or 6th order LO-phonon sideband) up to a field of 23 Tesla. This field corresponds to an electron cyclotron frequency comparable to the LO-phonon energy of ZnTe (26.5 meV); hence these experiments strongly suggest that the photoexcited electron-hole pairs exist in the form of Coulombically bound states, i.e., hot excitons, even at the initial high energy state following photoexcitation (typically more than 100 meV above the bandgap). The primary effects of finite disorder in the quantum well manifest themselves in striking details of the resonance enhancement for higher order (m 2 4) LO-phonon sidebands, which resonate not at the (absorptive) excitonic bandgap but some 1O-l 5 meV below it. In this energy region, dominated by the 2D density of states tail below a mobility edge, deviations from the LO-phonon ladder energy are also seen. Picosecond time-resolved measurements show that bottleneck effects exist, that is, the hot exciton relaxation slows once the quasi-particles 2.6

InSb Multiple

Quantum

The primary motivation

have reached the localized states.

Wells with CdTe Barriers for considering

the CdTe/lnSb

heterostruc-

ture is to obtain InSb quantum well structures; the problem to be solved, by using a II-VI/III-V

heterojunction,

is the absence

of suitable

Ill-V

com-

pounds

as the barrier

layers.

are no available

Ill-V

to serve

compounds

having

lattice constants

There

compatible

with InSb.

predictionsf4*) of band offsets agree fairly well with experimental mentsfs4) and suggest

that these quantum

substantial

and valence

conduction

Theoretical measure-

wells will be of Type I with

band confinement.

Minimal

strain

effects are expected as these two materials are very closely lattice matched (- 0.05%), while a perfect lattice match can be achieved by incorporating a few percent of either Zn or Mn into the CdTe barrier layer. Large quantum shifts in the bandgap energy are predicted for relatively wide quantum wells as a result of the small effective mass of electrons and light holes. For example, a 75 8, quantum well has a ground state


383

Heterostructures

transition energy twice that of bulk lnSb.f8sj Structures involving reasonable well dimensions allow a wavelength range of 2-5.5 I_tm to be accessed. provide

The high carrier mobilities the possibility

and the large de Broglie wavelength

for a wide variety

of interesting

devices.

The

realization of proposed device structures, however, has been hampered by the significant materials problems associated with this II-VI/III-V materials system. Several research groups have reported the MBE growth of InSb on CdTe substrates,fa6j and CdTe on both InSb substrates[84j[8rj-fgoj and InSb epilayers.f8s) These studies have primarily focused on interfaces epitaxial layers and substrates, whereas recent work involves

between epilayer/

epilayer interfaces.fg11fg2) The majority of previous studies employed InSb and CdTe bulk substrates which were ion etched and thermally annealed prior to epitaxy. The resultant epitaxial layers were of very high quality; however, close examination of the inter-facial region revealed a variety of problems. These difficulties must be eliminated as in quantum well structures the interfaces can completely dominate the electronic and optical behavior. Many problems arise due to the widely differing optimum growth temperatures for the two materials; high quality CdTe on InSb is grown at temperatures as low as 160”C,[g3] whereas InSb with superior electrical properties is grown at temperatures at or above 400°C. Interfacial problems occurring at the CdTe/lnSb interface include interdiffusion f84)f861fgo) precipitate formation (metallic indium or Sb segregation) ,fs8jfgol and intermediate layer formation of ln2Te3.f84)fg0j It is unclear which of these problems will prove important for the case of epilayer/epilayer interfaces. In addition to the interface problem, a fundamental difficulty associated with the CdTe/lnSb system is the tendency for autodoping. A recent studyfg4j indicated that when both compounds are grown in the same MBE chamber, Te seriously contaminated the Sb source such that it was difficult to control the carrier concentration of the InSb material. In order to circumvent

some of these problems,

recent workfg1)fg2j

has been: (i) to employ two separate growth chambers, connected by an ultrahigh vacuum transfer module, to eliminate autodoping problems; (ii) to use an antimony

cracker as a source of Sb, in an effort to improve the

low temperature growth of InSb, and (iii) to study the effects on the heterointerface when a large Cd overpressure is used during multilayered growth.fg5) To achieve high quality InSb material at substrate temperatures near 300°C the use of Sb, was employed with anticipation of achieving results similar to those obtained for the low temperature growth

384

Molecular

Beam Epitaxy

of GaAs using As,tgGttg71in place of As,. were also used for the growth configurations;

in addition,

Very low growth rates (0.18 A/s)

of the InSb in various

an epitaxial

heterostructure

buffer layer of CdTe was grown

when a CdTe substrate was used, and an epitaxial layer of InSb was grown on the InSb substrate. (Growths employing the Sb cracker were performed in a single growth chamber.) Golding et al.tg5) have studied the effect of Cd:Te flux ratio and InSb growth rate on the interfacial properties of multilayered structures using Auger electron spectroscopy and depth profiling. It was reported that the tendency for In,Te, formation at the interface during nucleation of CdTe on InSb was associated with a deficiency of Cd at the growth surface. It was found that a considerable reduction in the tendency for inter-facial compound formation resulted when an overpressure of Cd was used during nucleation of the CdTe layer. For the growth of InSb on CdTe at 300X, the evolution of RHEED patterns, from spotty to streaked, indicated three-dimensional nucleation. For growth rates of InSb that were less than 0.15 pm/hr, however, the spotty pattern remained unchanged during the growth period. At the lower growth rates, Auger electron spectroscopy analysis and depth profiling revealed a complete degradation of the 400 A CdTe layer lying below the interface, and severe intermixing throughout. When two growth chamberstgl)tge) were used for the growth of InSb/ CdTe single quantum well structures, the InSb epilayer was transferred from the InSb growth chamber via an ultrahigh vacuum transfer module (4 x 1O-lo torr) to a separate growth chamber for the CdTe epitaxy. Following the growth of CdTe, the structure was returned to the InSb chamber for formation of the quantum well, after which the second CdTe barrier or cap layer was formed. Single quantum wells and multiple quantum wells (20 periods) of InSb/CdTe have also been grown using a single growth chamber; the Sb cracker was employed for the growth of the InSb layers with substrate temperatures of 280°C. Figure 18 shows a dark field TEM micrograph of a 20 period CdTe/ InSb multiple quantum well structure. The dark contrast represents the 163 A InSb well, while the light contrast

is the 168 A CdTe barrier layer.

The presence of an In,Te, layer at CdTe/lnSb interfaces has been reported by Zahn et al.tgO) but is not confirmed in these TEM investigations. In one growth sequence, multiple quantum well structures were grown under similar growth conditions with the exception that differing Ta cracking tube temperatures were used. For one 20 period MQW structure,


the cracking grown

using

zone was kept at 850X, a temperature

investigations

revealed

region of the InSb/CdTe

whereas

of 1040°C

a high density multiple

Heterostructures

a second structure

for the cracking of dislocations

quantum

385

was

zone.

TEM

generated

in the

well for the former structure.

For approximately the same growth rate, the second structure, grown with the higher cracking zone temperature, exhibited an order of magnitude reduction in the number of dislocations occurring in the superlattice region. The generation of dislocations in the multiple quantum well region is unexpected as these two materials are very closely lattice-matched. We speculate that the reduction in the number of dislocations was related to a decrease in the number of Sb precipitates which may form at such low growth temperatures for InSb.

Figure 18. Dark field TEM micrograph of an InSb/CdTe multiple quantum well structure having 20 periods with layer dimensions of 163 A (dark contrast) and 168 A (light contrast), respectively. Structural

characterization

of the InSb/CdTe

MQWs has also been

performed with x-ray rocking curve diffraction. Figure 19 shows a x-ray rocking curve obtained from a 15 period superlattice structure having a periodic spacing of approximately 833 f 10 A. Satellite peaks are present

366

Molecular

Beam Epitaxy

with spacing of 208 arc seconds, and indicate a superlattice periodicity of 870 A, agreeing well with the TEM. The higher angle feature (FWHM = 22 arcsec)

is attributed

to the zero order diffraction

peak of the multilayer

structure, whereas the other high intensity feature (FWHM = 11 arcsec) in the spectrum is the (004) reflection of the InSb buffer layer/substrate. Assignment of the high angle peak as corresponding to the average lattice spacing of the multilayer in the growth direction assumes that each interface of the multiquantum well structure contains ultrathin inter-facial layers of In,Te,. Since the lattice spacing in the growth direction for the InSb buffer/substrate should be smaller than the average lattice plane spacing associated with the periodic structure, the zero order diffraction peak from the superlattice is expected to lie at a lower angle than the peak corresponding to InSb. A simple calculation, involving minimization of the strain energy in the multilayer, will predict the measured angular positions of the diffraction peaks provided several (- 5) monolayers of the assumed interfacial

In,Te,

is present per period of the superlattice.

I 000

800

1llSl)

arc seconds Figure 19. X-ray rocking curve of a 15 period superlattice having a periodic spacing of 833 2 10 A.

structure of InSb/CdTe


Heterostructures

387

infrared photoluminescence has been used to examine the optical properties of the InSb epilayers,tsl] double heterostructures,ts8] and multiple quantum wells.

The quantum efficiency

is lower than that obtained reasonable.

The double

of the double heterostructure

from an epilayer

heterostructure

buffer layer grown on an InSb substrate,

or substrate,

but is still

consisted

of a 0.42 pm InSb

followed

by a 1.63 pm CdTe

buffer layer, the active 160 A InSb layer, and a 2200 A CdTe cap.

(In this case the structure was fabricated by the interrupted growth approach using two separate growth chambers.) The spectrum contained two features of which the higher energy peak was assigned to band-to-band recombination. In comparison with bulk InSb, the spectrum was approximately five times broader, however, measuring approximately 20 meV at T = 10 K. The broadening is tentatively attributed to the presence of inter-facial layers.

3.0

ZnSe-BASED

3.1

Introduction

HETEROSTRUCTURES

Whereas the primary device interest in CdTe results from its lattice constant compatibility with HgTe, ZnSe has a lattice constant and bandgap which are relevant to integrated optoelectronics. The room temperature direct bandgap of ZnSe at 2.7 eV makes this semiconductor especially attractive for device applications operating in the blue portion of the visible spectrum. Potential optical devices include flat panel displays, blue light emitting diodes, and blue semiconductor injection lasers. Of perhaps equal importance is the possibility of passivating the surface of GaAs with ZnSe, as these two compounds

are closely lattice-matched

(0.25%) and

have the potential to provide a close-to-ideal heterointerface. reviews recent advances in the molecular beam epitaxy ZnSe-based superlattice and quantum well structures.

This section of ZnSe and

3.2

Homo-

and Heteroepitaxy

of ZnSe

Nucleation of ZnSe on GaAs. The majority of the work involving theMBE growth of ZnSe has employed GaAs as the substrate material.[99][lo81 GaAs is an attractive substrate material for several reasons: (i) ZnSe has a relatively small lattice constant mismatch with GaAs (0.25%), (ii) GaAs is readily available with high quality at a low cost, and (iii,) surface

388

Molecular

Beam Epitaxy

preparation techniques for molecular beam epitaxy are well understood. The accepted standard GaAs wafer preparation technique involving chemical etching subsequently

results

in the growth

thermally

MBE, oxide desorption

of a passivating

desorbed

in-situ

is generally

oxide

at 582°C.

layer which

is

In the case of II-VI

performed without the presence of an

arsenic flux. (Arsenic is an acceptor impurity in II-VI compounds.) Although a variety of desorption times and temperatures are reported, a typical bulk GaAs surface shows a reconstructed diffraction pattern indicating arsenic deficiency when examined with reflection high energy electron diffraction (RHEED). After oxide desorption, the GaAs substrate temperature is reduced to a value ranging between 250-400°C, and nucleation occurs as the Zn and Se source shutters are opened. (In most cases elemental sources of Zn and Se are used, but a compound ZnSe source has been used by some groups.)[lOs]-[llll Nucleation occurs in a three-dimensional manner as the streaked GaAs RHEED pattern is replaced by a spotty “fishnet” pattern. Under the aforementioned growth conditions, all research groups observe the three-dimensional nucleation; such behavior is also observed in the atomic layer epitaxy (ALE) of ZnSe on GaAs substrates.[“*] The early stages of growth of ZnSe have been studied using RHEED intensity oscillations[ 1°61[1071 where both GaAs substrates and MBE-grown GaAs epilayers were employed. The RHEED intensity oscillations were observed on the specular spot with an incident angle of less than 1” (offBragg conditions). In the study of nucleation on GaAs epilayers, two MBE systems were used. In one system, the GaAs epilayers were grown and the resultant as-grown GaAs surface was maintained via arsenic passivation techniques; another separate system was used for the ZnSe MBE growth.

To deposit an amorphous

cooled toward room temperature

arsenic layer, the GaAs sample was

in an arsenic beam.

After transfer

in air

from the Ill-V MBE, the As layer was desorbed at 290°C in the analytical chamber

of the separate

quent to As desorption, served; depending reconstructed

II-VI MBE.

Prior to As passivation

a (2 x 4) reconstructed

on the time and temperature

pattern was also observed

and subse-

GaAs surface

was ob-

of As desorption,

a (4 x 6)

after desorption.[sl]

The evolu-

tion of the RHEED diffraction pattern[lOq during nucleation of ZnSe on an MBE-grown GaAs epilayer clearly contrasted the nucleation on a GaAs substrate (which is described above). Once the ZnSe was nucleated, the early observation (after 9 seconds) of a strongly streaked RHEED pattern and the early presence of reconstruction lines suggested a more two-


dimensional character of the nucleation.

Heterostructures

The two-dimensional

was confirmed in observations of RHEED intensity oscillations Fig. 20 (top).f106jf10rj Strong intensity by-layer

growth/ “4

were observed

oscillations, for nucleation

389

nucleation as shown in

characteristic

of layer-

on the GaAs epilayer;

RHEED intensity oscillations were not seen when ZnSe was nucleated on a substrate (Fig. 20, bottom). Instead the variation of the specular RHEED intensity on a substrate was similar to observations reported for threedimensional nucleation of InGaAs on GaAs epilayers.f114j The RHEED intensity oscillations just described are unique in that they describe nucleation at a II-VI/III-V interface. two-dimensional nucleation

In a subsequent study by Tamargo et al.,f106j was reported to occur when growth com-

menced on an As-rich GaAs bulk substrate surface, whereas threedimensional nucleation was observed when growth occurred on a Ga-rich MBE-grown GaAs epilayer or substrate. Photoluminescence measurements performed by Tamargo et al. indicated that the ZnSe epilayers grown on the As-stabilized surfaces were of higher quality than layers nucleated on Ga-stabilized surfaces. Based on these experimental results, it has been proposedf 11sj-f117jthat an instability exists at the ZnSe/ GaAs interface due to an electronic imbalance of differing numbers of ZnAs bonds and Ga-Se bonds. The instability subsequently affects the resultant overlayer by the presence of a “disordered” interface. Their model proposed that a (2 x 4) As-surface has half coverage of As such that a mixed layer of As and Se results upon nucleation. The optimum growth would then occur since the number of electron-deficient Zn-As bonds would roughly equal the number of electron-rich Se-Ga bonds. In an effort to provide a closer lattice match to ZnSe than is provided by GaAs, ZnSe has been grown on substrates

of (In,Ga)As,

and epilayers

of AlAs and Ga(AI,As) .fl lel The emphasis in this work was to compare the photoluminescence obtained under the more lattice-matched growth conditions. The excitonic features (dominated by bound excitons) were reported to be considerably narrower (FWHM) than those observed when ZnSe was nucleated on GaAs directly. It is anticipated that x-ray rocking curves (not yet reported) would also become more narrow when using a substrate/buffer layer of closer lattice match. When x-ray rocking curves are obtained for ZnSe nucleated on GaAs substrates or epilayers, narrow FWHM values for relatively thick (> 1 pm) layers are approximately 126 arcsec;tllgj the angular broadening results from the formation of misfit dislocations at the GaAs/ZnSe heterointerfaces and the associated gation of threading dislocations. A detailed study of the crystalline

propaquality

390

Molecular

Beam Epitaxy

ZnSo Nucleation on GaAI Eplleyer

I

(llO]S~wg~ spot ,o

20

30

40

50

The

50

70

50

90

100

110

(seconds)

A

ZnSe Nucleation on GaAs Substrate

0

I 10

20

I 30

40

50

I 60

70

60

I 60

+

Time (seconds)

Figure 20. pop] RHEED intensity oscillations obtained during the nucleation of ZnSe on a MBE-grown GaAs epilayer at a substrate temperature of 320°C and viewed in the [l lo] azimuth. (Some positive peaks were cut off due to recorder (Sottom) Intensity variation of the specular spot RHEED bias limitations.). reflection in the [l lo] observed during nucleation of ZnSe on a GaAs bulk The higher substrate temperature substrate (400°C substrate temperature). would be exected to favor two-dimensional nucleation.


Heterostructures

391

of ZnSe grown on GaAs as measured by x-ray double crystal rocking curves and topography have been performed by Qadri et al.t120) When pseudomorphic

ZnSe films are grown on GaAs, misfit dislocations

do not

form, however the narrowness of x-ray rocking curves for these thin (1000 to 1500 A) layers are limited by the thinness of the films. Alternate Substrates. In addition to the use of Ill-V compounds as a substrate, ZnSe has been homoepitaxially grown on ZnSe substrates, and heteroepitaxially grown on the closely lattice-matched Ge (0.17%) and on lattice-mismatched Si (4.0%) substrates. Because of the difficulty in acquiring high quality, large area ZnSe substrates, only a few initial results of the homoepitaxial growth have been reported. Park et al.n2’l have reported the MBE growth of ZnSe on both (11 l)- and (lOO)-oriented ZnSe substrates which were prepared using ion milling and annealing. Under the conditions reported, although the (100) epilayer was of high quality, the (111) epitaxial layer quality was found to be significantly inferior, showing no excitonic

emission

in photoluminescence.

lsshiki

et al.t122j have

reported the fabrication of very high quality bulk substrates using zonerefined Zn as one of the starting source materials; these substrates may prove to be useful for achieving high quality homoepitaxial ZnSe material. ZnSe has also been grown on Ge substrates,t123] and MBE-grown Ge epilayers (grown on Si substrates). t124l With the presence of a ZnSe/Ge superlattice grown as the buffer layer on a Ge substrate, the quality of the subsequent ZnSe was greatly improved. ZnSe has also been grown on (100) and (111) Si substrates.t125j Monte Carlo Simulations. Monte Carlo simulations of the growth of ZnSe by MBEt 126] have been developed using a kinetic model closely following that developed for GaAs growth by Madhukar[127) and Singh and Bajaj.t126) The basic framework

for the simulation

lattice gas model with Arrhenius-type

rate equations

consisted

of a rigid

to represent

kinetic

processes such as surface migrations and re-evaporations. The kinetic model was modified from that used for GaAs to account for the differences in the growth conditions of the two compounds. Since ZnSe epilayers are typically grown under comparable Zn and Se flux intensities (unlike GaAs), it is expected that both cationic and anionic roles in controlling the quality of the film. surface migration of cationic Zn, the surface considered. The surface kinetic processes

atoms play equally important Therefore, in addition to the migration of Se needs to be considered for the simulation

are: (i) incorporation of monoatomic Zn, (ii) incorporation of Se, (iii) surface migration of Zn and Se, and (iv) re-evaporation of Zn and Se. Although

392

Molecular

Beam Epitaxy

only initial results have been obtained thus far, the insights provided by the Monte

Carlo simulations

conditions

for various

Effects

should

lead to predictions

of optimal

growth

II-VI compounds.

of Source

Many

Purity.

relied on the ability to control

electron

semiconductor

devices

have

and hole concentrations

by the

selective incorporation of donors and acceptors. The wide gap II-VI compounds have exhibited a difficulty in achieving amphoteric doping. In general, it was found that the tellurides were readily doped p-type whereas the selenides were more easily doped n-type. Reasonable carrier concentrations have been achieved in the n-type doping of MBE-grown ZnSe. In these experiments, the dopant specie was incorporated during the growth process itself; annealing at elevated temperatures was not required to activate the substitutional donors. Attempts to produce p-type MBE ZnSe have been reported1 12g)-f134)and photoluminescence has indicated the presence of ionized substitutional acceptors, Studies

of dopant

incorporation

have

been

complicated

by the

presence of unintentional impurities found in both elemental and compound source material used in MBE. In most cases undoped ZnSe grown by MBE, using commercially available source material of six-nines purity, has been n-type with low resistivity (- 1 R-cm).t135) The low resistivity of the ZnSe material implied that the ZnSe was of good stoichiometry, as a relatively small deviation from a unity Zn-to-Se flux ratio toward either Znrich or Se-rich conditions was found to result in high resistivity material.t136] The defects generated during growth under non-stoichiometric conditions appeared to compensate the nonintentionally incorporated impurities. In doping experiments performed in our laboratory, at a given Ga oven temperature,

the resistivity was found to increase by two orders of magnitude when the flux ratio (Se/Zn) went from one to two. (The fluxes of

the elements were measured by a quartz crystal monitor placed approximately at the position of the substrate.) Through enhancement of the purity of the source material by vacuum distillation and/or zone-refining, nominally temperature

undoped

ZnSe,

grown

under similar

and flux ratio, exhibited

enhanced purity photoluminescence

conditions

high resistivity

of substrate

(- lo4 Q-cm).

The

of the resultant ZnSe material was confirmed in measurements, where free exciton features were more

prominent, having intensities similar to, and sometimes greater than, bound exciton-related transitions, The use of purity-enhanced source frc41frc51f1371 Yoneda et al.tlc41 material was reported by three groups. performed a study in which they reported the variation

of carrier concentra-


Heterostructures

393

tion, resistivity, and relative amplitude of free exciton emission for undoped ZnSe as a function of the number of purification cycles wherein the Se source material was vacuum distilled. The carrier concentrations ranged from 1 x 10”

cm” to less than 7 x 1014 cm3 as the number of purification

cycles was varied from 1 to 9, respectively. tently

used vacuum

distilled

source

At Purdue, we have consis-

material

(Zn, Se, Mn, and CdTe)

prepared in-house for the growth of (Cd,Mn)Te,fg] ZnSe,f105] and (Zn,Mn)Se.flcsl In our case, depending on the conditions of the MBE apparatus and particular charge of source material, we have measured both high resistivity (- 1O4 Q-cm) undoped ZnSe and lower resistivity ZnSe (on the order of 3 Q-cm). We have also grown undoped ZnSe using commercially available vacuum distilled source material obtained from Osaka Asahi Mining Company. Again we found nominally undoped ZnSe to have a resistivity greater than lo4 ohm-cm. Our observations agreed with the results of Ohkawa et al.f137j where they obtained high resistivity (lo4 &-cm) undoped ZnSe using the purity-enhanced source material purchased from the above commercial vendor. Material Characterization. A wide variety of characterization techniques are used to study both the quality of the ZnSe material itself as well as the ZnSe/GaAs heterointerface. Studies of photoluminescence, modulated reflectance, x-ray diffraction, and transmission electron microscopy are described in this section. Basic photoluminescence studies involve the examination of optical emissions corresponding to, for example, free and impurity bound exciton features, donor-to-acceptor pair transitions, and deep level emissions usually attributed to impurity/defect complexes. Numerous studies to date have investigated

the effect of Zn-to-Se

flux ratio,f1361f138jsubstrate

tem-

perature,f13g] thermal and lattice mismatch strain, and impurity content of the source material.

Recent developments

have been directed towards: (i)

a better understanding of the role of lattice mismatch and thermal strain and their effect on the energy of near-bandgap optical transitions, and (ii) the effect of source excitonic

purity

on the relative

intensity

of free and bound

features.

One factor

which

contributes

ZnSe/GaAs heterojunction is the match. Zincblende ZnSe has a 5.6676 A whereas GaAs has a 5.65315 A, resulting in a mismatch

to the potential

importance

of the

relatively small lattice constant misroom temperature lattice constant of room temperature lattice constant of of 0.25%. Although the lattice constant

394


mismatch is relatively small for this heterointerface, a finite amount of strain exists in the ZnSe epilayer. In the context of photoluminescence, the strain is evidenced gies, which epilayers

in both free and bound excitonic

differ from those observed

in bulk crystals.

transition

ener-

For very thin

(< 1500 A), the exciton energies lie above those observed in bulk

crystals; as layer thickness increases to greater than a 1 pm, the transition energies exhibit a red shift from the bulk values. The red shift is attributed to strain resulting from the thermal expansion coefficient mismatch between ZnSe and GaA~.t~~~l Measurements of reflectance, as the film thickness is varied, show a particular transition feature shifting in energy with film thickness.t140) The interpretation of the origin of the changing excitonic energies is complicated by the wide range of values for deformation potentials and thermal expansion coefficients reported in the literature. As the ZnSe layer is nucleated on the GaAs substrate, the layer remains pseudomorphic (having the in-plane lattice constant of the substrate) until a critical thickness is reached where misfit dislocations appear. The result of the misfit dislocations is to cause the lattice constant of the film to relax toward that of a bulk crystal. Assuming that a pseudomorphic layer has a known strain, photoluminescence measurements can be used to estimate the deformation potentia1s.t lo 71f1411 Figure 21 shows the low temperature photoluminescence spectrum of a 1000 8, ZnSe epitaxial layer grown on a 1.5 pm MBE-grown GaAs epilayer. The feature at 2.7997 eV is usually associated with a neutral donor bound exciton while the two higher energy features represent free exciton transitions, as identified in modulated reflectancet142] spectra (Fig. 22). The free exciton feature at 2.8064 eV (corresponding

to a heavy-hole

transition)

is shifted, from the bulk transi-

tion energy of 2.802 eV, due to the strain resulting from the 0.25% lattice constant mismatch; the feature at 2.8178 eV is associated

with a light-hole

transition. The magnitude and sense of the strain aredetermined by the pseudomorphism; the amount of valence band splitting, together with the blue-shift of excitonic features, can be used to estimate the ZnSe deformation potentials.

The values of b = -1.05 eV and a = -4.87 eV are found to

compare favorably with the recently reportedfg8] values of -1.2 eV and -5.4 eV, respectively. The subsequent measurement of excitonic energies for pseudomorphic ZnSe layers using modulated reflectancet143a] results in deformation potential values which agree even more closely with the reports in Ref. 143.


Heterostructures

395

2.6064 .V

4400

4420

4440

4460

WAVELENGTH

4460

4500

(A)

Figure 21. Photoluminescence spectrum of a 1000 8, ZnSe epilayer on a 1.5 /.fm GaAs MBE-grown epilayer at 6 K taken with an excitation density of 200 mW/cm”.

90 60

1

30 _ . g

-30

i

-60

x

Q

/

0

-90 -120 -150 -180 2.65

2.70 2.75 2.60 2.85 2.90 2.95 Photon

Energy (EV)

Figure 22. Modulated reflectance spectrum of the 1000 A pseudomorphic ZnSe layer grown on the 1.5pm GaAs MBE-grown epilayer (77 K).[‘~~~] Dominant free exciton features observed in photoluminescence at 77 K occurred at 2.797 and 2.810 eV.

396

Molecular

Beam Epitaxy

X-ray diffraction measurements have been used to observe the variation in lattice constant versus film thickness for ZnSe grown on G~As.[‘~~) In these measurements, at a maximum

for fully strained

toward the bulk lattice constant

the perpendicular pseudomorphic

lattice constant was

layers

as the film thickness

and decreased

approached

1 pm.

For further increases in film thickness, the perpendicular lattice parameter approached a constant value somewhat below the bulk value, as a result of the strain induced by the thermal expansion coefficient mismatch present between the ZnSe and GaAs. The local details of microstructure on an atomic scale is provided by transmission electron microscopy. Samples, ion milled down to dimensions suitable for electron imaging, can be prepared for both crossSample preparation of the II-VI sectional and plan-view examination. compounds is more critical than that for the Ill-V compounds due to the tendency of the II-Vls to sustain significant radiation damage from ion milling. Typically argon ions are used in the milling step, but significant reductions in the degree of radiation damage are observed when xenon or iodine is used in the ion milling process.t144] The microstructure of ZnSe films grown on both GaAs epilayers and GaAs substrates was examined by cross-sectionalt107)f144] and plan-view t14’) transmission electron microscopy. No dislocations or stacking faults were found in observed areas of 1000 8, thick ZnSe films, even in the interface area, thus confirming pseudomorphic growth; the interface between the ZnSe film and the GaAs epilayer appeared as a sharp straight line in cross-sectional dark field imaging. A ZnSe film with a thickness of 1.3 pm grown on a GaAs epilayer, on the other hand, showed an array of misfit dislocations. Figure 23 shows a plan-view (viewed normal to the interface) bright field image of the interface. A well-developed network of misfit dislocations, some of which are dissociated to pairs of partial dislocations, is clearly seen in the image. A comparison of ZnSe epitaxial layers grown on GaAs substrates and MBE-grown GaAs epilayers was provided by TEM. A close examination of the interface

microstructure

was obtained

using cross-sectional

high resolution electron microscope (HREM) images. The film grown on an epilayer (Fig. 24, top) exhibited almost featureless images of the interface, which appeared as an atomically flat boundary over wide areas with a perfectly coherent contact between the two crystals. HREM images

397

Wide Gap II-VI Semiconductor Heterostructures of the interface between typical ZnSe films grown on GaAs substrates, the other hand, revealed presence segments substrate

a wavy step-like

boundary

which

of small pits and steps on the GaAs substrate of stacking interface.

faults were observed

indicated

surface.

Small

along the ZnSe/GaAs

Figure 24 (bottom) shows a high resolution

on the bulk

electron

micrograph of the ZnSe/GaAs epilayer interface as viewed in the [OlO] projection. In this projection, the chemical nature of the interface is probed, unlike the case of the more conventional [Ol l] projection.f145] As can be seen in the figure, the (200) and (002) lattice fringes in each semiconductor layer differ substantially, easily identified.

such that the interface

can be

Figure 23. Plan-view bright field image of ZnSe/GaAs-epilayer interface showing misfit dislocation network. The ZnSe layer thickness is 1.3 pm.

398

Molecular

Beam Epitaxy

Figure 24. High resolution electron microscope image of the (top) [Ol l] projection and (bottom) [OlO] projection of the interface between the pseudomorphic ZnSe-GaAs epilayer viewed in cross-section. The ZnSe layer thickness is 1000 A. Measurements were performed using the 1 MV TEM at the Tokyo Institute of Technology and the 200 kV TEM at Purdue University.


Heterostructures

399

Pseudomorphic ZnSe/n-GaAs MISFET Devices. AI,Ga,_xAs has a direct bandgap of 2.0 eV at x = 0.5 and has been widely used as an “insulator”

for

heterojunction maximum)

GaAs however,

field

effect

transistors.

The

presents a very low interfacial

to carrier flow from GaAs into the (AI,Ga)As.

(Al,Ga)As/GaAs barrier (0.4 eV at ZnSe, having a

direct bandgap of 2.7 eV, is closely lattice-matched (0.25%) to GaAs; in a variety of device applications the ZnSe/GaAs heterointerface could proRecently a prototype device strucvide an alternative to (AI,Ga)As. ture[146)f147j was fabricated wherein pseudomorphic ZnSe formed the “insulator” in a GaAs depletion-mode field effect transistor. The device was fabricated using an interrupted growth technique employing arsenic passivation (described in Sec. 3, Nucleation of ZnSe on G&k). The microstructure of the pseudomorphic ZnSe/GaAs heterointerface (as described above) shows that the degree of coherence and the absence of defects at the interface is similar to observations of cross-sectional TEM samples of (AI,Ga)As/GaAs. The lo-V,, curves for a 45pm gate prototype device is shown in Fig. 25. The FET curves show good depletion-mode characteristics with complete pinch-off and current saturation.[147) The modulation of the channel carrier concentration indicates that the Fermi level positioning at the ZnSe/n-GaAs interface can be varied by at least 0.6 eV. Although the transconductance (g,) appears to be low (3.5 mS/mm), when the effect of series resistance is included, a value of 5.1 mS/mm is obtained and agrees fairly well with a theoretical maximum prediction of 8.5 mS/mm. The possibility offered by the interrupted MBE growth technique is the ability to systematically control the interface and film properties in order to reduce the number of states that contribute to charge trapping. Growth parameters such as substrate temperature, flux ratio, and crystal stoichiometry can be used to alter the interface, leading to improved device performance. ZnSe/GaAs MIS Capacitors. Historically, the passivation of GaAs has met with considerable difficulty due to the inadequate electrical characteristics exhibited by interfaces formed by the deposition of various insulators,

or by the formation

of native oxides onto GaAs surfaces.

The

presence of a high density of interfacial surface states prohibits utilization of the interface in device configurations. As described above, an alternative pseudo-insulating layer for GaAs has been demonstrated, consisting of the wide bandgap II-VI semiconductor ZnSe, and is thus similar to (AI,Ga)As.

The II-VI compound

has a close lattice constant

match to

400

Molecular

Beam Epitaxy

GaAs, and the semi-insulating ZnSe layer can form an epitaxial heterojunction. Recently, the occurrence of both hole accumulation (for ptype GaAs) and inversion

(for n-type GaAs) in post-growth

annealed ZnSe/

GaAs structures has been rep0rted.f 1481 The C-V characteristics of the annealed structures were nearly ideal, exhibiting an integrated (over the GaAs bandgap) interface state density of 2.5 x 10” cm-*, a value which compared favorably with the densities repotted1 14g1f1501 fortypical (AI,Ga)As/ GaAs interfaces. A disadvantage of the post-growth anneal was a tendency for doped ZnSe samples to become compensated, a result which would be undesirable

in certain

device applications.

Figure 25. Room temperature I-V characteristic of the metal/ZnSe/n-GaAs field effect transistor with a gate width and length of 45 pm. The vertical scale is 50 microamps/div and the horizontal scale is1 V/div. The gate bias is decreased by 0.5 volts per trace starting at 0 volts.

Qiu et al.f151) have selection of an appropriate of ZnSe, that resulted in characteristics without the

described a growth technique, involving the GaAs surface stoichiometry prior to nucleation as-grown samples exhibiting nearly ideal C-V necessity for post-growth annealing. Both the


Heterostructures

401

ZnSe and GaAs epilayers were grown in separate growth chambers of a modular MBE system in order to avoid cross-contamination; transfer between growth chambers occurred in an ultrahigh vacuum (UHV) transfer module.

A series of experiments

were performed

in the ZnSe growth

chamber wherein the GaAs epilayer surface stoichiometry was altered prior to the nucleation of ZnSe. Starting from the as-transferred sample which had the arsenic-rich c(4 x 4) reconstructed surface, the GaAs epilayers were heated to different temperatures to reduce the surface As content, resulting in four different GaAs interface reconstructions. When the GaAs was heated to approximately 460°C, a (2 x 4) surface reconstruction was observed. At about 51 O”C, a (4 x 6) surface reconstruction pattern appeared. These latter reconstructed GaAs epilayer surfaces resulted in an essentially two-dimensional (2D) nucleation of ZnSe, where the transition from a streaked GaAs RHEED pattern to a streaked ZnSe pattern occurred within 10 to 15 seconds. The 2D character of the nucleation was further supported by the observation of RHEED intensity oscillations commencing within one second of opening the Zn and Se source shutters. As the GaAs epilayer temperature continued to be raised above the temperature where the (4 x 6) pattern was observed, in the vicinity of 535OC the reconstruction again changed. Although one might expect that a (4 x 2) Ga-stabilized pattern would follow the (4 x 6) as the temperature was increased, the reconstruction pattern recorded was a (4 x 3), changing to (2 x 3). The (4 x 3) has also recently been reported by Kobayashi et al.t152] The latter reconstructions, differing from the conventionally reported GaAs surface reconstruction patterns, may have resulted from the “decoration’q*l) of the heated GaAs surface by high vapor pressure species such as Se or Zn. Nucleation

of ZnSe on the (4 x 3)/(2 x 3) reconstructed

surfaces was of a 3D character, exhibiting a spotty RHEED pattern which became streaked in less than one minute. The observation of 3D nucleation of ZnSe on a Ga-rich

(4 x 2) GaAs surface

has been previously

reported by Tamargo et al.[115) Following the growth of the epitaxial ZnSe/GaAs heterojunctions, a series of C-V measurements were performed. Among the four types of asgrown samples, the capacitors associated with the c(4 x 4) GaAs surface exhibited the most pronounced interface state-induced stretching near the mid-gap, while those capacitors formed on a Ga-rich GaAs surface had no indication of such stretching, indicating a virtual elimination of interface states in this region of the bandgap. Because of the reduction in interface states obtained for the Ga-rich samples, the Fermi level was free to move,

402

Molecular

Beam Epitaxy

and the band bending spanned the entire GaAs bandgap. The C-V characteristics were virtually independent of frequency from 1 kHz to 4 MHz.

At large positive

voltages,

exhibited

deep depletion

instead

existence

of a small conduction

all the ZnSe/p-GaAs

samples

of electron

suggesting

inversion,

band discontinuity,

insufficient

tested the

to confine

electrons, a conclusion which is consistent with previous observations and other published results.[14g1[1501[1521[1531 The interface state density distributions were measured, and the results are shown in Fig. 26. By comparing the interface state densities near the mid-gap for the series of samples, a clear trend showing a reduction in the interface state density as the GaAs epilayer surface became increasingly As deficient was seen. For the samples grown on Ga-rich GaAs, the interface state density, integrated over the lower portion of the GaAs bandgap, was in the low 10” cm-*. Recent TEM studies combining image simulation with the experimental examination of cross-

80t

/’

P

c (4 x 4) (as-grown) \\ I, \

(2 x 4)

,/L

\

‘, I

.,

c (4 x 4) (annealed)

,

I

I

1 1

EF-Ev

tev)

,

I

I

~ a EC

Figure 26. Interface state density distribution calculated using Terman’s method. The two parameters needed for this calculation, GaAs epilayer doping level and the ZnSe layer thickness, are obtained from C-V profiling and TEM, respectively. The dotted line represents the typical interface state distribution for an As-rich GaAs epilayer surface obtained after a post-growth annealing.


Heterostructures

403

sectional samples tend to indicate the presence of 1 or 2 monolayers of a strained interfacial compound at the interfacet1541t155) of samples having low interface densities. Image simulations tend to support a hypothesis that the interfacial The presence

compound

of GazSe,

consists of the zincblende

is further

confirmed

ments of the Se bonding at the interface.

phase of Ga,Se,.

by in-situ XPS measure-

In the structures

grown on As-

rich surfaces, similar inter-facial layers have been observed by TEM, although they appeared to be less distinct compared to those observed in the samples grown on As-deficient surfaces. Dark field images and high resolution electron microscope (HREM) images of cross-sectional samples of each heterostructure were examined in order to find a difference in interface structure due to the change of the As coverage of the GaAs surfaces. The 200 dark field images of (010) cross-sectional samples were found to exhibit the most significant difference among three heterostructures. Figures 27(a), (&I), and (c) are 200 dark field images of (010) cross-sectional samples of the heterostructures grown on the c(4 x 4)) (4 x 6), and (4 x 3) GaAs surfaces, respectively. As seen in the images, the samples grown on the (4 x 6) and (4 x 3) surfaces show a distinct bright line at the interface between the ZnSe and GaAs epilayers. The bright line of the sample grown on the (4 x 3) surface appears to be more continuous than that of the sample grown on the (4 x 6) surface. The 200~type dark field images of the heterostructure grown on the c(4 x 4) surface, on the other hand, do not show such a bright line. The difference among the interface images of the three heterostructures suggests the existence of a transition structure at the ZnSe/GaAs interfaces which have formed on As-deficient GaAs surfaces. In order to analyze the transition structure, HREM images and dark field images of other reflections were examined. All 400~type dark field images show a dark interface line. The crystal structure factors of the 200 reflections for the zincblende structure are given by a difference in scattering factors of atoms occupying two different face centered cubic (fee) sublattices, while the crystal structure factors of the 400 reflections ing factors of these two types of atoms.

are additions

Considering

in scatter-

this relation, one can

propose the following simple model for the transition structure. Between the ZnSe and GaAs crystals, a very thin layer having a zincblende structure exists by maintaining a coherent relation. One of the fee sublattices is occupied by the cation, i.e., Zn or Ga, and the other has the anion, i.e., Se or As. Unlike the GaAs and ZnSe crystals, one of the fee sublattices in the thin layer has a high concentration of vacancies. Be-

404

Molecular

Beam Epitaxy

cause of vacancies in one of the fee sublattices, the crystal structure factors of the 200 reflections of the thin layer become much greater than those of GaAs and ZnSe, which values

of scattering

therefore,

factors

are very small due to the nearly equal

of the constituent

atoms.

The thin layer,

will appear with a brighter contrast in 200 dark field images.

In

400 dark field images, on the other hand, the thin layer will appear as a dark line as a result of having a smaller crystal structure factor than those

Figure 27. The 200 dark field images of the ZnSe/GaAs interfaces in the heterostructures grown on (a) the c(4 x 4) surface, (b) the (4 x 6) surface, and (c) the (4 x 3) surface. In each image, the upper layer is ZnSe and the lower layer is GaAs.


Heterostructures

405

One of the stable phases of (Ga,Se) compounds, Ga,Se,, is known to have a structure identical to that suggested by the present observation.f1561 It has a zincblende structure, and one third of the Ga sites are left as vacancies. As a result of these vacancies, the lattice parameter of Gasses is about 5% smaller than those of GaAs and ZnSe. Based on the model of a thin Gasses layer, which is coherently

inserted into the ZnSe/

GaAs interface, intensities of 200 and 400 dark field images are calculated by utilizing the two-beam dynamical theory and the column approximation. Results of the calculationst154] are in good agreement with observed images despite the use of simple approximations. The profile of the 200 dark field image shows a bright contrast for the Ga,Se, layer, while the 400 dark field image shows a decrease in the intensity at the interface layer. With agreement obtained in the analysis of dark field images, combined with the formation of a transition structure only on As-deficient GaAs surfaces, one may suggest that the ZnSe/GaAs heteroepitaxial interface has a transition structure identical to that of Gasses. The present results therefore, provide further evidence for the formation of III,VI, compounds at II-VI/III-V semiconductor interfaces, 3.3

Quantum

Well Structures

Incorporating

(Zn,Mn)Se

ZnSe has been layered with the II-VI semiconductor compounds ZnTe, ZnS, and Zn(S,Se), the semimagnetic semiconductor Zn,,Mn,Se, and the magnetic semiconductor MnSe to form wide gap II-VI superlattices and multiple quantum well structures. Figure 22 shows a bright field image of a ZnSe/Zn,,,,Mn,~,, Se multiple quantum well obtained by transmission electron microscopy. For all of these materials, the lattice constant varies substantially from that of ZnSe, such that layered structures form strainedlayer superlattices provided the layer thicknesses are below the critical thickness where misfit dislocations form. In this section, some of the more recent results discussed.

involving

the wide

The first reported MBE-grown tures involving

bandgap

superlattice

multiple quantum

structures

are

well (MQW) struc-

ZnSe employed the DMS material (Zn,Mn)Se

as the wider

bandgap barrier layer. As in the case of (Cd,Mn)Te described above, the band structure of the host II-VI semiconductor (ZnSe) is not directly modified by the presence of Mn since the two s electrons of the outer shell replace those of Zn and become part of the band electrons in extended states. The five electrons in the unfilled 3d shell of Mn, however, give rise

406

Molecular

Beam Epitaxy

to localized magnetic moments which are partially The resultant

magnetic

aligned in an external

magnetic

field.

moment interacts

electrons

causing a Zeeman splitting which is orders of magnitude

with the band larger

than for the host II-VI semiconductor at low lattice temperature. The presence of the magnetic ion (with the associated Zeeman shifts) provides a unique and useful feature to the superlattices and multiple quantum well structures in which Mn is incorporated. The magnetic-field-induced changes in optical transition energies in superlattices incorporating the DMS materialt151 provide additional

insight into excitonic

behavior and valence band

offset in strained-layer structures in general. Iron and cobalt, two other magnetic transition metal ions, have also been incorporated into ZnSe by MBE to form thin film DMS alloys. The growth and optical and magnetic characterization of these DMS layers were described by Jonker and coworkerst158)-t186) (The growth by MBE of the first Ill-V based dilute magnetic semrconductor, In,_, MnxAs, has recently been reported by Munekata et al.tr6q)

Figure 28. Bright field TEM image of ZnSe/Zn,,,Mn,,,Se barrier 120 A).

superlattice (well 35 A,


3.4

Epitaxiai

Growth of the Metastable

to investigate

407

(Zn,Mn)Se

The MBE growth of the pseudo-binary has provided an opportunity

Heterostructures

material system ZnSe-MnSe metastabie

over a large range of alloy fractions, which tional bulk equilibrium growth techniques.

zincblende

crystals

are unavailable by convenRelatively thick (l-3 pm)

epiiayers of zincbiende Zn,,Mn,Se have been grown by MBEt168t over the entire (0 < x 5 0.66) composition range, whereas bulk crystals exist with pure zincbiende crystal structure only up to x c 0.10.t16g~~170~Figure 29 shows the variation in lattice parameter versus Mn content for the MBEgrown (Zn,Mn)Se epilayers. The ability to achieve zincblende (Zn,Mn)Se having appreciable Mn content was crucial, as the small variation in bandgap with Mn concentration made it necessary to grow barrier layers with a high Mn fraction, to achieve sufficient band offset for carrier confinement.

5.850 -

6.800

-

2 - lg 5.750

-

_ -----.--

’ ZINOllLEND MIXJ31)PIIASES IIEXAGONAL

6.81 - (4.18)

_ 6.76 (4.13)

0.71 - (4.08)

f3.00 - (4.03)

0

0.1

0.2

0.3

0.4

0.5

0.0

0.7

X Mn Fraction

Figure 29. Lattice parameter (a,J as a function of the Mn mole fraction (x) for zincblende Zn,,Mn,Se epilayers. The broken lines represent data obtained on bulk crystals[1701 which contain mixed phases of zincblende and hexagonal crystal structure.

408

Molecular

Beam Epitaxy

From photoluminescence zincblende (Zn,Mn)Se epilayers, function

of Mn concentration

and reflectance measurements of the the bandgap has been determined as a (see Fig. 30).

reported, the bandgap differences in the range of 100-300 meV. behavior of the MQW structures

For the MQW structures

between barrier and well materials were Detailed studies of the magneto-optical indicated that the magnetic-field-induced

band shifts are primarily due to exchange interactions associated with the penetration of the hole wavefunctions into the (Zn,Mn)Se barrier layers. Concluding that much of the hole envelope wavefunction resides in the (Zn,Mn)Se barrier layer implied that the strained small and is likely to be less than 20 meV.f15q

valence

band offset is

II

Figure 30. Near-bandgap energy (E,) variation versus Mn content (x) for Zn,_ ,Mn,Se epiiayers. Data are shown at 77 K and 6.5 K with the low temperature data extrapolated to yield a bandgap value of 3.4 eV for zincblende MnSe. (Crosshatch is data obtained on bulk crystals by Twardowski et al.)[16g1


3.5

Optical Properties

of (Zn,Mn)Se

Heterostructures

Quantum

409

Wells

The study of optical properties of ZnSe/(Zn,Mn)Se has been focused so far mainly on excitonic transitions near the E, gap in the blue region of the spectrum.

In comparison

tures, two principal differences Zn,,Mn,Se

arise.

with Mn concentration

with CdTe/(Cd,Mn)Te

heterostruc-

First, the increase of the bandgap of

(x) is considerably

weaker so that for x

- 0.20 in the barrier layer, the total bandgap difference is about 150 meV (vs. nearly 400 meV for the Cd,_,Mn,Te structure). Second, the uniaxial component of the lattice mismatch strain is opposite in sign, so that the uppermost valence band has a light-hole character (parallel to superlattice axis at k = 0). At the same time, the exchange effect on hole states near band extrema by the Mn-ion d-electron spins is larger than that in (Cd,Mn)Te. Figure 31 shows the photoluminescence excitation spectrum near the n = 1 exciton ground state from a (lOO)-oriented MQW sample of ZnSe/Zn,,Mn,Se (x = 0.23), with a well width of 67 A.trs71 For such typical parameters, only the n = 1 transition is seen, consisting of the light-hole (LH) and heavy-hole (HH) excitonic resonances. The identification of the resonances was made through magneto-optical studiest157] where both LH and HH transitions are found to exhibit large Zeeman splittings due to finite exciton wavefunction overlap with the (Zn,Mn)Se barrier layers. In contrast with the CdTe/(Cd,Mn)Te quantum well case discussed above, however, details of the band offsets have not yet been extracted from such magneto-optical data. This is in part due to the larger excitonic energies, and due to difficulties in treating the n = 1 light-hole exciton properly. Nonetheless, the qualitative indicators suggest that much of the total bandgap difference is accommodated in the conduction band. Effects of lattice mismatch strain are important; in particular, the uniaxial component deepens the effective well depth for the light-hole state. In the following, we review three examples of recent optical spectroscopy

measurements

which

have elucidated

the nature of electronic

states near the fundamental edge of ZnSe/(Zn,Mn)Se MQWs. Again, excitonic effects play a central role and dominate radiative recombination at low and moderate lattice temperatures. Competition Between Excitons and Mn-ion d-Electron internal Transitions. Apart from increasing the direct bandgap at k = 0, the incorporation of Mn into ZnSe leads to additional features in the optical spectrum due to the d-electron transitions internal to the Mn-ion. The lowest (crystal field split) transition 6A, 4 4T, corresponds to absorption at about 2.1 eV and the zero-phonon line in luminescence is at about 2.0 eV.

410

Molecular

Beam Epitaxy

In heterostructures containing (Zn,Mn)Se, these ‘yellow’ resonances can compete for electronic excitation with the ‘blue’ resonances which are In particular, there are efficient associated with band edge transitions. energy transfer

paths from the band edge exciton

directly into the Mn-ion internal excitation,

something

states in (Zn,Mn)Se which can be graphi-

cally demonstrated by comparing thin alloy films of (Zn,Mn)Se with ZnSe/ luminescence (Zn,Mn)Se quantum wellst 105)t171)through time-resolved soectroscopy.[172)

285

2.80

2.75 Excitation

Energy

( eV 1

Figure 31. Photoluminescence excitation spectrum for a ZnSe/(Zn,Mn)Se MQW structure (see text). The n = 1 light and heavy hole resonances dominate the spectrum.

As an example, Fig. 32 shows luminescence spectra from a ZnSe/ (Zn,Mn)Se MQW sample at T = 2K under CW excitation above the barrier bandgap.[1721 The ZnSe well thickness was 67 A and the Mn-ion concentration in the Zn,,Mn,Se barriers was x = 0.23. The spectrally sharp (blue) exciton recombination dominates the broad yellow Mn-ion recombination from the barriers, the ratio of spectrally

integrated

intensities

being about

17 to 1. In contrast, the blue recombination in a thin film of (Zn,Mn)Se is dwarfed by the now dominant yellow contribution so that the blue/yellow Thus, the quantum well structure is intensity ratio is about 4 x 10”. efficient in collecting electron-hole pairs prior to any significant energy transfer of such bandedge excitation to the d-electron states of the Mn-ion.


ZnSe/

ZI

Heterostructures

411

,

Zn_,,Mn.,,Se

z

L,=

ii?

a, z a, .-> t;; u [I:

67 8

T=2

FWHM

K

N5meVI)

I,,’ 1.9

2.0

2.1 Photon

2.2 Energy

C

I\ 2.795 (eV)

Figure 32. Comparison of blue (- 2.6 ev) and yellow (- 2.1 ev) luminescence emitted from a ZnSe/(Zn,Mn)Se MQW sample. Note the differences in amplitude and energy scale for each.

Once captured in a ZnSe quantum well, the photoenergetic electrons and holes relax by optical phonon emission to the n=l confined particle states. This relaxation step is fast (probably well below one psec in such polar material), The subsequent exciton formation and further energy relaxation (localization by quantum well width fluctuations) is a slower process which can be time-resolved through the use of picosecond pulsed

laser excitation

and a streak camera.[17*]

Figure

33 (left panel)

shows the transient luminescence (dotted line) from the ZnSe/(Zn,Mn)Se MQW sample when excited above the n = 1 exciton resonance. The risetime is approximately

90 psec; this time constant shortens to approxi-

mately 20 psec under resonant

excitation,

thus yielding

a direct measure

of the exciton formation step. Recombination lifetime is, of course, also obtained from the data (- 200 psec); some details of this can be found in Ref. 173. By contrast, the blue exciton decay in a (Zn,Mn)Se thin film is very fast (- 15 psec) as shown in Fig. 33 (right panel); this gives a direct measure of the rate of energy conversion electron excitation of the Mn-ion.

from the exciton state to the d-

412

Molecular

Beam Epitaxy

?b,=2.91

c ::. :: : nw~=291 t .*

eV

.* ::

i: :. :.

T=2K

eV

L : 1.19fim T=2K

: !

Time

0 Time

Wed

100 (psec)

Figure 33. Time-resolved exciton luminescence at the fundamental edge for (/efi) a ZnSe/Zn,,,Mn ,,nsSe MQW and (right) a Zn,,,Mn,,,Se thin film. The dashed

lines indicate the time response of the monochromator/streak camera system.

In order to obtain further information Stark Effect on Excitons. about the excitons and their confinement in the ZnSe/(Zn,Mn)Se quantum well, the influence of an externally applied electric field on the lowest confined particle transitions have been studied in photoluminescence.t174) This is also a subject of some device interest, given the recent success in exploiting the field tuning of excitons in Stark effect-based electro-absorbing quantum wells composed of Ill-V compound heterostructures. Figure 34 graphs spectral shifts induced in photoluminescence (PL) emission at the n = 1 LH exciton transition from a wide single quantum well (SQW) buried within a ZnSe/(Zn,Mn)Se The dimensions La=185Aforth

multilayer

quantum

well sample.

of the SQW section were Lw = 70 A for the ZnSe well and e Z n,,Mn,Se barrier of x = 0.45 Mn alloy concentration.

The experimentally

measured shifts also included the effects of the built-in

Schottky barrier field. (Application of approximately + 9 V of external bias reproduced the PL spectrum from an identical SQW control structure without a Schottky barrier electrode.) The estimated total electric field in the SQW region is shown on the top edge of the figure. Once the Schottky effects are accounted for, an electric field thus produces the expected spectral redshift (and some broadening in the exciton linewidth).


Electric

Field

Heterostructures

( kV/cm )

0.6 0.4 0.2

0

2

4 External

6

413

8

10

Voltage

( V )

12

A .Z 2 g ti

0.0 14

Figure 34. Electric field-induced shifts in the peak position of the recombining n = 1 exciton and changes in its amplitude for the SQW portion of a ZnSe/(Zn,Mn)Se heterostructure (T = 2 K). The total electric field (sum of Schottky and applied fields) is obtained from the applied voltage by comparison with an unmetallized sample. The solid line is a result of calculation; the dashed lines are a guide to the

eye. In the region of biases shown in Fig. 34, the currents measured in the external circuit were negligible (or small) and the PL quantum efficiency was relatively constant as shown in the figure. Beyond approximately 14 V, however, strong forward bias conduction was observed together with a rapid quenching of the SQW PL emission. The conduction corresponds to efficient injection of electrons across the n+ GaAs/ZnSe heterojunction, as also verified by the onset of bright yellow Mn-ion internal luminescence in the structure.

The yellow luminescence,

well known as the basis for thin

film electroluminescent devices, is generally thought to be induced by hot electron impact excitation of the Mn-ion d-electron states. Its presence here shows the high interfacial quality of the nt GaAs/ZnSe heterojunction. A representative treatment of the problem of an exciton in a GaAs/ (Ga,AI)As quantum well in the presence of an electric field, in connection with electroabsorption, has been given by Miller et al.t17q In a II-VI

414

Molecular

Beam Epitaxy

quantum well, where the exciton binding energy may be comparable to one of the band offsets, one must self-consistently include both the externalfield (applied plus Schottky) and the exciton Coulomb field.t176] Such an approach is relevant under the conditions considered here, namely a relatively large exciton binding energy and a small valence band offset. In the calculation we take the valence band offset to be 60 meV, in keeping with AEv .CAEc. While this choice is somewhat arbitrary, it does reflect a finite confinement of the hole part of the exciton as well as a relative stability against field ionization. The calculation also shows the particular feature of the exciton with a large binding energy in wide gap II-VI quantum wells; namely that the internal Coulomb field is quite effective in opposing the tendency of electron-hole Nonetheless, the spectral red shifts separation by the external fields. observed are rather large (several mev) considering the moderate fields employed in these experiments. This net shift is, of course, a result of two opposing effects: (i) a reduction in the exciton binding energy (blueshift) and (ii) changes in the one particle energies (redshift). In this Nonlinear Excitonic Effects in ZnSe Quantum Wells. section, we describe two observations of a nonlinear optical effect associated with excitons in MQW samples of ZnSe/(Zn,Mn)Se. These represent the first attempts to exploit the large exciton binding energy in wide-gap II-VI semiconductor superlattices for inducing phenomena which might be useful in optical switching devices in the blue-green region of the optical spectrum. Nonlinear Excltonic Absorption. Nonlinear optical absorption at excitonic resonances has been extensively studied in Ill-V heterostructurest17rlf17*] with application towards modulators and bistable switches. The main physical phenomena, namely the decrease of oscillator strength at a sharp exciton transition with increasing exciton level, has been theoretically discussed in terms of exciton Coulomb screening and phasespace filling in the interacting electron-hole system. Such effects can also be seen in ZnSe-based quantum wells although their detailed elucidation is still incomplete. Nonlinear excitonic absorption has been measured in ZnSe epilayers and ZnSe/(Zn,Mn)Se MQWs at 77 Kt17gl-f1611and room temperature.[~~91~~~01[~~11A s an example, Fig. 35 shows the change in excitonic absorption in a thin (1.3 pm) ZnSe film and a ZnSe/Zn,,MnxSe MQW sample at T = 77 K.t 17g) The quantum well thickness was 73 8, with For a three-dimensional exciton Bohr x = 0.51 in the barrier layer. diameter of about 60 8, in bulk ZnSe, the exciton (n = 1 LH transition) is, therefore, at least partially influenced by confinement effects in the heterostructure. The nonlinear absorption experimentst17g)t160] were performed with a single tunable dye laser whose intensity-dependent transmission was directly measured. One can obtain a rough quantitative measure for


the saturation homogeneous

Heterostructures

415

of absorption for data in Fig. 35 by assuming the case of a lineshape. (This is not really true here, but still provides a

useful yardstick.) tion intensity

When compared

decreases

with the thin epitaxial

from approximately

film, the satura-

10.7 kW/cm2 to 1.3 kW/cm2

for the MQW sample.f17g]

1.2 z j i d T

r,plO.i

kW/cad

_

1.0 0.6 -

2 C

0.2 -

E

0.4 0.2 0.0 2.7eo

2.780

2.800

2.820

PhotonEnergy (ev)

03

-

I

I330

w/cm’

-

Figure 35. Nonlinear absorption (from transmission measurements) for a ZnSe thin film (upper panel) and a ZnSe/(Zn,Mn)Se MQW structure (Lw = 73 A) (lower panel) at 77 K at the ground state exciton resonance. Note the difference in saturation intensities. (The abscissa is In (l/transmission coefficient).)

416

Molecular

Beam Epitaxy

Using equilibrium

statistics

to estimate

the relative

free electron-

hole pair and exciton densities under these experimental conditions, one obtains some insight into the mechanisms which contribute to the observed absorption

saturation.

However,

(order of 200 psec) makes estimates

the relatively

short-lived

based on thermal equilibrium

exciton some-

what uncertain. Considering the Mott screening of excitons in the Debye limit, one estimates for the bulk ZnSe a critical density n, = 3 x 10” cm” at 77K, while the experimentally generated electron-hole pair density at observed saturation is estimated at 9 x 1016 cm3 . For the MQW sample, estimated screening of excitons occurs at n, = 7 x 1016 cm-s while experimentally measured saturation corresponds to a pair density of 2 x 1016 cm3. Thus Coulomb screening appears to be a dominant mechanism in both cases, although there is additional evidence in thinner ZnSe quantum well samples that the phase space filling becomes an important issue. Excitonic Molecules. The quasi-two-dimensional electron and hole confinement in a semiconductor quantum well has consequences beyond enhancing the binding energy and oscillator strength of an exciton. For example, enhancement in the stability of an excitonic molecule is also expected. This quasi-particle is a molecular-like entity composed of two electron-hole pairs bound together. Following the calculation by Kleinmar$s*] one can estimate that the possible increase in the binding energy of such a molecule (biexciton) in a quantum well may be as much as one order of magnitude. To date, however, there is very scant evidence for biexciton formation in quantum wells,t 16sl in part because the binding in a GaAs or similar Ill-V heterostructure is still weak (less than 1 mev). On the other hand, for bulk ZnSe there are reportsf1s41 of biexcitons with binding energies 6 - 3.5 meV so that in a well-designed heterostructure, the molecular

state should be unambiguously

observable.

biexciton state has been detected in ZnSe/(Zn,Mn)Se part of this work is summarized here. Figure 36 shows an intensity-dependent

Recently,

the

MQW structures;t185)

photoluminescence

spec-

trum at low temperature from a ZnSe/Zn,,Mn,Se MQW sample (with a well thickness of 67 A and barrier thickness of 110 8, with Mn concentration x = 0.23). The photon energy of excitation is well above the bandgap of the barrier layer. The emission X refers to the free exciton (weakly localized by quantum well width fluctuations). The emission XX which increases superlinearly is a first hint for an excitonic molecule. In a thinner quantum well, the XX emission can dominate at relatively modest levels of excitation, growing superlinearly with excitation level until the exciton dissociation

limit (screening].


2.77

2.79

2.79

2.80

2.76

Heterostructures

2.79

2.60

417

2.62

ENERGY (eV)

Figure 36. Photoluminescence from a ZnSe/(Zn,Mn)Se MQW sample (Lw = 67 A and 24 A) showing a superlinear dependence on intensity of excitation of the emission line XX. The line X represents recombination from the n = 1 ground state exciton. The incident intensity is (a) 30 mW/cm2 and (b) 100 W/cm*. The arrows at the high-energy edge in spectra (b) indicate the excitation laser lines.

An excitonic molecular state should also be possible to excite resonantly through direct two-photon absorption with a giant oscillator strength. This is also observed in the ZnSe/Zn,_,Mn,Se MQW structures, as illustrated

in Fig. 37, for a sample with well thickness

of 24 A and barrier

thickness of 160 A with x = 0.28. The two-photon absorption in this case is measured through intensity-dependent photoluminescence excitation spectrum (Fig. 37, a and b) which shows that the two-photon absorption (resonance

at l/2 E,_J is competing

relatively modest levels of excitation

with the one-photon (- 200 W/cm-*).

process at the

The right hand panel

shows how the giant two-photon cross section decreases with excitation by circularly polarized excitation (Fig. 37, c and o); this is expected for the molecular ground state (zincblende crystal) consisting of electrons and holes in singlet states, respectively. Additional supporting evidence for the biexciton in the ZnSe quantum wells has been obtained from temperature and magnetic field dependent studies,1 1861all pointing to the enhanced circumstances which should make the observation of such effects generally possible in wide gap II-VI compound semiconductor

superlattices.

418

Molecular

Beam Epitaxy

II/Z

E,,

(a) (d)

( b)

I

2790

I

2830 EXCITATION

2790 ENERGY

2830

(meV)

Figure 37. Absorption

near the n = 1 exciton ground state for a ZnSe/(Zn,Mn)Se MQW sample (Lw = 24 A), obtained through photoluminescence excitation spectroscopy for two excitation levels: (a) 250 W/cm* and (b) 30 mW/cm*. Note the strong intensity dependence of the low energy peak which is interpreted as corresponding to two-photon resonant excitation of the excitonic molecule. The right hand panel compares the dependence of the excitation spectrum on (c) linear and (d) circularly polarized incident radiation at an intensity of - 200 W/cm*.

Stimulated Emission and Lasing in ZnSe Epilayers Superlattices. Laser oscillations have been observed in ZnSe-based tures under both optical excitation

and electron

beam pumping.

and strucTo our

knowledge the first report of lasing in an MBE-grown ZnSe structure was based on experiments with ZnSe/(Zn,Mn)Se multiquantum wells.t1e7] The experiments were performed on cleaved cavities after removal (by selective etching) of the GaAs substrate. Gain spectra were measured, and thresholds of stimulated emission determined, for various emission wavelengths. Optically pumped lasers (Fig. 38) were fabricated from these


Heterostructures

419

multiple quantum well structures and found to operate in the blue portion of the visible spectrum. Lasing was observed at temperatures up to 80 K. These

first (Zn,Mn)Se

which

exhibited

MQWs

an order

had thresholds

of magnitude

for stimulated

improvement

over

emission previously

reported results[ 1881using single crystal ZnSe grown from a melt. Further improvements

in the thresholds

of the MQWs of this material system are

anticipated based on the addition of cladding layers to provide maximal optical confinement and through optimization of growth parameters.

I

.79 P‘h

)IJ

I

1

d

I

450 451 452 453 454 Wavelength

I

455

45

(nm)

Figure 33. Emission spectra at four different power levels for one of the ZnSe/ Zn,,,,Mn,,,,Se MQW lasers at 5.5K showing the onset of laser oscillation.

420

Molecular

Beam Epitaxy

ZnSe epilayerst l 111[18gland ZnSe/ZnSSe superlatticest” ‘1 have also shown lasing under electron beam pumping. Room temperature lasing of an electron beam pumped ZnSe epilayer was obtained with a threshold 5 A/cm2 whereas

the superlattice

structures

of

had a 12 A/cm2 threshold

current density.t l1 ‘1 Very recently, optical pumping of ZnCdSe/ZnSe quantum well lasers with properly designed waveguide structures has been obtained at room temperature under the usual low duty cycle excitation, and up to 100 K under “quasi-continuous” high repetition rate operation.t190] These lasers exhibit low threshold behavior and suggest that the goal of a room temperature continuous-wave optically pumped laser action may be attainable. In general, electron beam and optically pumped lasers are useful in that heterostructure laser designs can be constructed and tested without the still serious difficulties of achieving both p- and ntype material in wide gap II-VI compounds. Polarization-Dependent Luminescence. The observation that quantum well structures have anisotropic optical properties for light polarized parallel to the layers FE,, and perpendicular to the layers (TM), even though the constituent materials are isotropic in bulk form, is currently the focus of considerable interest. A simple explanation for the anisotropy, neglecting band mixing effects, is as follows. The relative oscillator strength for the heavy-hole conduction-band transition is 3 for the TE polarization

and 0 for the TM polarization,

while for the light-hole

conduc-

tion-band transition it is 1 for the TE polarization and 4 for the TM polarization. In bulk isotropic materials, the heavy-hole and light-hole bands are degenerate at the band edge, so that the absorption and gain depend on the sum for the two bands, a value which is the same for both polarizations. However, in quantum well structures the valence band degeneracy is removed providing for optical anisotropy. In (Ga,AI)As quantum wells, in which the lattice constants of the well and barrier materials are closely matched, the splitting of the heavy and light-hole bands

occurs

primarily

from size quantization.

The size quantization

causes the heavy-hole band to have a smaller blue shift relative to the light-hole band because of the heavy hole’s greater mass. Consequently, when valence-band splitting is principally a result of size quantization, the optical properties near the bandedge are dominated by heavy-hole conduction-band transitions leading to a greater absorption and gain for the TE polarization relative to the TM polarization. Optical gain and absorption spectra for TE and TM modes propagating in the plane of the layers of (1 OO)-oriented (Zn,Mn)Se MQW structures


Heterostructures

421

have been measured.tlgl) In these structures, contrasting the case of (Ga,AI)As MQWs, there is significant strain in both well and barrier layers due to the lattice constant mismatch.

The strain plays an important role in

the splitting

bands in addition

of heavy- and light-hole

to the usual size

quantization effect observed, for example, in (Ga,AI)As MQW structures. In the (Zn,Mn)Se material system having zincblende crystal structure, the lattice constant is found to increase with increasing Mn mole fraction (see Fig. 29). As a result, the ZnSe well regions are subjected to a compressive uniaxial strain in the growth direction. The compressive uniaxial strain acts to lower the energy of the heavy-hole band while raising that of the light-hole band. Thus in (Zn,Mn)Se quantum wells, the strain acts in an opposite sense to the effect of confinement on the shifting of the valence bands. In particular MQWs, the ground state transition energy in the ZnSe well region is actually red-shifted with respect to bulk ZnSe.f105) In measurements of absorption spectra for a (Zn,Mn)Se MQW structure, the relative positions of the TE and TM absorption edges are found to be opposite to that observed in (Ga,AI)As MQW structures; the TM absorption edge is at a lower energy than the TE absorption edge.tlgl) The greater oscillator strength of the TM mode near the bandedge is also reflected in the gain spectra (Fig. 39). In fact only the TM gain spectra were measured; no TE signal could be detected. These wide gap II-VI quantum well structures are the first to exhibit TM polarized stimulated emission originating from a MQW structure.t1g2) In (Cd,Mn)Te MQWs, the TE mode absorption and gain are dominant. (The gain spectra for the (Cd,Mn)Te MQW is seen in Fig. 40.) The opposite behavior of (Cd,Mn)Te and (Zn,Mn)Se MQWs is ascribed to the opposite sense of the uniaxial strain in these systems. 3.6

ZnSe/MnSe

Magnetic

Superlattices

The first growth of the hypothetical

zincblende

magnetic

semicon-

ductor MnSetlg3) resulted from efforts to increase the incorporation

of Mn

in (Zn,Mn)Se. The existence of zincblende MnSe is a consequence of the kinetic (non-equilibrium) nature of the MBE growth method employed; bulk crystals of MnSe have the NaCl crystal structure.

Extrapolations

of

data on lattice constant and bandgap, obtained for zincblende (Zn,Mn)Se epilayers with up to 66% Mn, predict a bandgap and lattice constant for zincblende MnSe of 3.4 eV (6.5 K) and 5.93 A, respectively. A variety of superlattice

structures

composed of ZnSe layered with ultrathin MnSe (3 to

32 A) have been grown by MBE; the layer thickness of the MnSe (with monolayer resolution) was controlled by the use of RHEED intensity oscillations.t~061t~941t1951

422

Molecular

Beam Epitaxy

160140. 120. 100. 60. 60. 40. 20. 01

I

2.700

2.710

2.720

2.730

2.740

WI

Figure 39. Polarization-dependent gain spectra for a (Zn,Mn)Se MQW structure. The spectra were taken at 6 K. For the gain spectra, the excitation density was 4.9 x 1O5 W/cm2. No TE signal could be detected for the gain spectra.

Figure 40. Polarization-dependent gain spectra for the (Cd,Mn)Te MQW structure. The spectra were taken at 25 K. For the gain spectra, the excitation density was 1.57 W/cm2.


Heterostructures

423

RHEED provided the first observation of the zincblende crystal structure for MnSe. During MBE growth, as the Zn shutter was closed and the Mn shutter virtually

was opened

instantaneous

to begin the growth

increase in the intensity

tion streaks was observed.

of MnSe on ZnSe,

of the two-fold

a

reconstruc-

In addition, the RHEED pattern revealed a high

degree of ordering as demonstrated by the presence of Kikuchi lines and, with the exception of somewhat more intense reconstruction lines, was quite similar to the pattern observed from ZnSe. In all cases the zincblende crystal structure was confirmed by TEM examination of cross-sectional specimens. Electron diffraction patterns corresponded phases with no indication of the wurtzite or rock-salt

only to zincblende crystal structures.

The metastabie zincblende MnSe, with thicknesses ranging from 3 to 32 A, was incorporated in strained-layer superlattice structures. When grown as a “thick” epilayer (400 A), the zincblende crystal structure remained in the presence of strain-relieving misfit dislocations. Monolayer Control Using RHEED intensity Oscillations. The zincblende MnSe served as the barrier when combined with ZnSe in By controlling the thickness of, and spacing superlattice structures. between, layers of the magnetic semiconductor, an opportunity was provided to study magnetic ordering as the dimensionality varied from 30 to quasi-2D. it was crucial in a study of this type for the MnSe layer thickness to be controlled to one monolayer. Although RHEED intensity oscillations had not been previously reported for growth of II-VI compounds, they were employed during the superlattice fabrication to provide the requisite one monolayer resolution. RHEED intensity oscillations were observed during the homoepitaxial growth of ZnSe on ZnSe and during the heteroepitaxial growth of ZnSe on MnSe (4.7% lattice mismatch). One period of oscillation was found to be equivalent to one monolayer of growth as confirmed by TEM. It was interesting to compare the homoepitaxy of ZnSe with the growth of ZnSe on MnSe under identical growth conditions. oscillation

periods

were observed

growth; in contrast, the nucleation

For homoepitaxy,

after interruption

six or seven

and re-initiation

of ZnSe on MnSe resulted

of

in up to 30

oscillation periods of enhanced amplitude.f106] Figure 41 shows the effect of alternating the cation species without growth interruption during the fabrication of a binary superlattice composed of MnSe/ZnSe. For this particular binary superlattice, each period consisted of four monolayers of ZnSe separated by three monolayers of MnSe. The layer thicknesses were controlled by counting the number of oscillation periods, and con-

424

Molecular

Beam Epitaxy

firmed by HREM imaging of cross-sectional specimens. Throughout the entire growth of the five period superlattice, intensity oscillations were recorded in both the ZnSe and MnSe layers. intensity

oscillations

layer thicknesses

I

10

The observation

of RHEED

during the growth provided the means to control the

with one monolayer

resolution.

L

20

60

40

60 Time

60

70

60

60

loo

(secwL&)

Figure 41. RHEED intensity oscillations observed during growth of a MnSe/ZnSe binary superlattice structure ([Olo] azimuth and a 400°C substrate temperature). The Se source shutter remained open throughout. The superlattice is grown on a thick (l-l .5 pm) ZnSe buffer layer.

A series of comb-like superlattices consisting of 30 to 100 periods were grown with MnSe layer thicknesses of one, three, and four monolayers; the MnSe layers were separated by approximately 45 A of ZnSe. For comparison the superlattices with one and three monolayer MnSe thicknesses were grown with and without growth interruption techniques at each interface. A fourth related superlattice structure consisted of 30 periods containing ten monolayers of MnSe alternated with 24 A of ZnSe.

Wide Gap II-VI Semiconductor Heterostructures

425

Optical and Magnetic Aspects of ZnSe/MnSe Superlattices. The preparation of semiconductor superlattices, in which layer thicknesses approach molecular monolayer limits, presents a situation where effects of heterointerfaces

can become

a key factor

in determining

the physical

properties of such layered structures. A so far unexploited possibility is to use magnetic phenomena as a complement to conventional electronic probes for obtaining interface-specfic information. Here we review recent measurements on the ‘metastable’ ZnSe/MnSe superlattice.sf1g6) where strikingly large, nearly paramagnetic contributions to the susceptibility have been detected in structures containing ultrathin, highly strained MnSe layers near the 2D magnetic limit. The experimental results show dramatically the importance of real interfaces to magnetic properties, and are interpreted through the presence of inter-facial microstructure which appears to be very effective in frustrating magnetic ordering. The following structures were studied in this work:tlg3) (i) a 102 period MnSe/ZnSe superlattice (SL) where MnSe layers consisting of 3 molecular monolayers (approximately 9 A) were separated by 45 8, spacers of ZnSe (henceforth referred to as the 3 ML SL), (ii) a 30 period MnSe/ ZnSe SL with 10 monolayers of MnSe (30 A) separated by 24 A thick ZnSe layers (a 70 ML SL), and (&I a 400 A thick epitaxial film of zincblende MnSe. In addition, a superlattice with one molecular monolayer was also available and used in optical characterization. To provide a direct characterization of the magnetic properties of these superlattices, magnetization measurements were carried out on a superconducting quantum interference device (SQUID) magnetometer. The magnetic properties arise because of the half-filled d-electron shell of the Mn-ion, which contributes a net spin moment of 5/2 Bohr magnetons of spins

on nearest

neighbor

cation

in the free ion. The interaction

sites in the zincblende

lattice

is

antiferromagnetic (AF) in character; as a consequence, bulk (zincblende) MnSe is expected to undergo a magnetic phase transition into an antiferromagnetically ordered state as the temperature is lowered below approximately 200K. In this limit, only very small net magnetizations will be induced in weak or moderate

externally

applied fields.

A key point is

that the magnetic interactions in question are controlled by very short range interactions and dominated by an exchange process between two Mn-ions through the intervening Se anion. The very small SL or thin film volume made it necessary to employ precautions in order to reduce the mounting materials to a minimum. Figure 42(a)-(c) summarizes the temperature

dependence

of the magnetization

for the three

samples

426

Molecular

Beam Epitaxy

introduced above (measured in a 1 kG field oriented parallel to the (100) layer).[lg6] The striking experimental result, exhibited by both SL samples, is the large positive contribution sample is superposed

at low temperatures

which for the 10 ML

on another distinctly temperature-dependent

(posi-

tive) contribution. a complete

In strong contrast, the thin epitaxial MnSe film exhibits opposite in its temperature dependence, and behaves qualita-

tively as a normal bulk AF insulator, with ordering at a finite temperature (cancellation of net magnetization by opposing sublattice contributions.)

(a)

4

(Cl

_ &p

‘ , .

I

eat/l

.’ . . . J .

0

Figure 42.

loo xl0 Temperature (K 1

3of

Magnetization

monolayer (ML) ZnSe/MnSe (c) a MnSe epitaxial substracted.

film.

0

100

200

3a

Tempera!ure ( K)

(x 10" emu) as a function of temperature for (a) 3 superlattice, (b) 10 ML ZnSe/MnSe superlattice, and Substrate and buffer layer contribution have been

Complementary information about the electronic and magnetic properties was acquired through optical studies near the SL bandgap (197), including SLs in the monolayer limit of MnSe. The efficient photoluminescence emission which originates from such a SL transition is illustrated in Fig. 43 (at T = 2 K in zero and 5 Tesla external magnetic fields in a Faraday geometry). is distinctly blueshifted

As expected, the zero field exciton (at 2.844 ev) from that in bulk ZnSe. The Zeeman effect

immediately suggests that a large magnetization is induced by an external field in the MnSe layers; this can by indirectly inferred by considering the strong exchange interaction of the bandedge states with the Mn-ion d-electrons. In Fig. 43, the strong circularly polarized character of the


Heterostructures

427

emission verifies the transition as connecting the spin-split [l/2,-1/2> conduction and 13/2,-3/2> valence bands. In its simplest description, the Zeeman effect is proportional

to the product of the exchange

d and p-d), the Mn spin system magnetization, overlap with MnSe barriers.

constant

(s-

and the exciton wavefunction

For samples containing

1 and 3 monolayers

of MnSe per SL period, two types of SLs were available:

(i) those of normal

growth and (ii) structures where the MBE growth was interrupted at each MnSe/ZnSe interface. Very little change in the magnitude of the Zeeman shifts resulted from the finite growth interruption. The largest Zeeman shifts occurred in the limit of single monolayer shown in Fig. 44.

MnSe barrier layers,

as

The direct magnetization and magneto-optical measurements crosscheck each other, and display strikingly large low temperature magnetizations which can be induced in the ultrathin layer MnSe SL samples. For example, optical information ensures that these originate from the SL portion of the samples. The key conclusion which emerged from these experiments was that the anomalous, nearly paramagnetic low temperature contributions originated from the heterointerface regions in the MnSe/ ZnSe SLs. That is, the magnetic moments at the interfaces appear to be remarkably free from expected AF couplings.

T=1.8 K

2.80

2.82

Photon

2.84

2.86

Energy ( eV)

Figure 43. Exciton luminescence at the superlattice bandgap for the 3 ML sample at T = 2 K, and circularly polarized emission in a 5 Tesla magnetic field.

428

Molecular

Beam Epitaxy

2.85

I,.

;,--._‘a-““o__ -_ *.a. 0 ‘r,

3 MnSe monolayers/ period 0

-*._

T= 1.13 K

l B I layer 0 B II layer

‘=L~ -..I

*.*

I-.

. .

‘*.A ‘.

monolayer/ period

2.80

.

2.79 0

1 Magnetic

2

4

3 Field

(Tesla

)

Figure 44. Zeeman shifts for 1 ML and 3 ML superlattices at T = 1.6 K. Triangles and dark circles refer to 1 ML structures with and without growth interruption. Field anisotropy is shown for the 3 ML case; here also growth interruption (not plotted) affected only a small change in the slope of the low field Zeeman shift. (Dashed lines are to guide the eye.)


Heterostructures

429

The magnetic “probes” examine microstructure on a scale on the order of chemical bond lengths. This follows from the short range nature of the superexchange paths in insulators such as MnSe (nearest neighbor Mn-ions

coupled

measurements

through

the intervening

an enhanced

sensitivity

Se anion),

to deviations

and give magnetic from ‘perfect’ atomic

arrangements (vs. ideal bulk) within a monolayer or so at the interface region. Some qualitative arguments can be applied concerning the roles of intrinsic and extrinsic microstructure effects at the MnSe/ZnSe heterointerfaces which underlie the frustration of AF interactions seen in our experiments. Diffusion or chemical intermixing effects can provide regions of diluted Mn-ion concentration, but would also cause significant linewidth broadening of the luminescence beyond what is observed. On the other hand, while 2-dimensional growth is characteristic of this SL system,t1Q4)t1Q5)we cannot eliminate the possibility of incomplete layer growth during the heteroepitaxy. Finite size islands at the interfaces, i.e., 2-dimensional clusters, can be effective in frustrating antiferromagnetic ordering. One additional possibility for frustration of magnetic ordering by topological ,reconstruction effects raises the issue of intrinsic reconstruction effects under MBE growth conditions of significant interface strain. High resolution transmission electron microscopy and electron diffraction experiments on the MnSe/ZnSe SLs are at present insufficient to detail the interface on the scale discussed here (< 3 A). At present the answers to these questions remain largely important because of the need to scale of heterointerface formation to generate sophisticated artificial 3.7

ZnSe/ZnTe

Superlattice

a matter of guesswork, yet they are improve understanding on a microscopic in the multitude of contemporary efforts semiconductor layered structures.

Structures

The difficulty in obtaining p-type ZnSe to serve as an injector of holes was the primary motivating factor leading to the growth of ZnSe/ ZnTe

superlattice

Zn(Se,Te)

structures;

ZnTe

is readily

doped

alloy can be either n- or p-type depending

p-type

while

the

on the Te fraction.

There are, however, several potential difficulties associated with ZnSe/ ZnTe heterostructures. The band offsets predicted by the electron affinity rule would suggest a Type II superlattice where holes and electrons are confined in separate layers; the resultant decrease in the oscillator strength of optical transitions could pose a problem for light emitting devices. A more serious consideration is the large lattice constant mismatch between

430


ZnSe and ZnTe (7.4%); however, strained-layer superlattice structures are still possible with layer thicknesses restricted to a few tens of angstroms. For structures

containing

GaAs is still possible, equal

amounts

primarily

whereas

of ZnSe

ZnSe, a reasonable

for structures

and ZnTe,

lattice match to

containing

the average

lattice

approximately constant

ap-

proaches the lattice constant of InP. ZnSe/ZnTe superlattice structures have been grown by a variety of techniques including hot wall epitaxy,t1g8) molecular beam epitaxy,f 1gg)t200)atomic layer epitaxy (ALE),tlggl and by a combination of ALE and MBE.f201) Photoluminescence studiest202) of the MBE-grown superlattice structures show a wide wavelength tunability as a function of the ratio of ZnSe to ZnTe layer thicknesses; photoluminescence emission is observed from the red to the green portion of the visible spectrum. Specially designed ZnSe/ZnTe superlattice structures can also be viewed as a means to circumvent the difficulties encountered in the MBE growth of the Zn(Se,Te) al1oy.t*Or1 Zn(Se,Te) is of particular interest due to the observation that the photoluminescence yield of the alloy can be significantly enhanced over that of bulk ZnSe crystalst203] and epitaxial layers, due to localization of excitons in the random alloy. The growth of the Zn(Se,Te) mixed crystal by molecular beam epitaxy, however, is complicated by a difficulty in controlling the composition. In the work reported by Yao et al.,1*04) over the entire range of Te fraction, a Te-to-Se flux ratio of 3 to 10 was required. In our laboratory we have grown a number of Zn(Se,Te) epilayers with varying fractions of Te; a particular difficulty was encountered when a small fraction of Te was desired, resulting in widely varying compositions under what appeared to be similar growth conditions. To circumvent the problems associated with controlling the alloy concentrations, we have designed ZnSe-based structures consisting of ultrathin layers of ZnTe spaced by appropriate dimensions to approximate

a Zn(Se,Te)

tion. As an illustration,

mixed crystal with low or moderate Te composi-

such a “pseudo-alloy”

quantum

well in a ZnSe/(Zn,Mn)Se

quantum

well structures,

was used to modify the ZnSe

heterostructure.

either one or two ZnTe

In these

multiple

ultrathin

layers were

placed in the center of each ZnSe well; the well thicknesses

ranged from

44 to 130 A. In an effort to optimize the interface abruptness of the ZnTe monolayers, the ALE of ZnTe was performed on a recovered ZnSe surface which made up, for example, the first half of a quantum well, whereas the remainder of the structure was grown by MBE. Although the architecture


of these

structures

optical transitions,

was substantially as viewed

431

Heterostructures

different

from a bulk alloy, their

in photoluminescence,

were dominated

by

features which were quite similar to those found in the bulk Z’n(Se,Te) alloy crystals at low to moderate Te composition. These luminescence features arise from the capture and strong localization of excitons at the isoelectronic Te sites.

Figure 45 dramatically

illustrates

exciton trapping

at the

ZnTe monolayer sheets present in the ZnSe quantum well. For comparison purposes, Fig. 45(a) shows the low temperature photoluminescence spectrum of aZnSe/(Zn,Mn)Se MQW structure in the absence of ZnTe.f15q The luminescence was dominated by the sharp (FWHM < 5 mev), bright, blue exciton recombination at the n = 1 (light hole) quantum well transition. As a striking contrast to Fig. 45(a), the photoluminescence from a ZnSe/ (Zn,Mn)Se MQW, which now incorporates the ZnTe sheets inserted into each quantum well, is shown in Fig. 45(b). The broad luminescence features at lower energy were the result of exciton localization at the ZnTe/ ZnSe heterointerfaces and were similar to those seen from bulk Zn(Se,Te) mixed crysta1s.t 2031Figure 45(c), showing a photoluminescence excitation (PLE) spectrum, indicates that the position of the lowest energy exciton transition has not been significantly shifted by the presence of the ZnTe sheets.t205) In other words, the confined quantum

well valence

particle states of the ZnSe/(Zn,Mn)Se

and conduction

band states are not significantly

perturbed by addition of the ultrathin ZnTe layer in the absorptive process. The luminescence spectrum, however, reflects the energetics of the relaxed exciton, that is, the strong localization of the hole at the ZnTe layer in the middle of the quantum well. Such a self-trapping is triggered by potential energy lowering at the isoelectronic Te sites for valence band states, and followed

by strong local lattice relaxation

effects so that the

final hole Bohr orbit is below 10 A. The two primary features in the PL spectra of Fig. 45(a) and (b) are associated with exciton trapping at single Calculations show that, for one Te and double Te-sites, respectively. monolayer

intermixing

of the anion, the Te distribution

double sites and single sites, followed incorporate

the Te isoelectronic

is dominated

by larger clusters.

trap centers

in a planar

by the

The ability to and spatially

controllable way had provided substantial new insight to the capture process. Considerable microscopic understanding to the exciton trapping process has been obtained from spectroscopy in magnetic fields and measurement of the exciton kinetics through time-resolved spectroscopy in such isoelectronically

delta-doped

ZnSe/ZnMnSe

quantum wells.t205)-t207]

432

Molecular

Beam Epitaxy

2.3

2.4

2.5 Photon

2.6 Energy

2.7

2.6

2.9

(eV)

Figure 45.

Comparison of photoluminescence spectra at T = 2 K of a ZnSel (Zn,Mn)Se MQW sample (a) with that of a similar structure but with the insertion of monolayer sheets of ZnTe in the middle of the quantum well (b). (The amplitude of emission in (a) has been reduced to bring the peak to scale.) The photoluminescence excitation spectrum of sample (b) is shown in the bottom

panel (c).

In an experimental

extension

of this work, the optical properties

of

such unique, isoelectronically doped quantum wells have been studied by increasing the ZnTe layer thickness in monolayer steps up to four. The aim has been to use the ZnSe/ZnTe system to approach the subject of band offset formation from the isoelectronic center point of view. That is, while the intermixed delta-doped ZnTe monolayer case clearly shows the effects of isolated centers, in the opposite limits of a ZnTe/ZnSe quantum well, type II band offset is expected from bulk considerations. Recent spectroscopic


Heterostructures

433

results suggest that the transition from one limit to the other is a continuous one where a structure containing four monolayers of ZnTe in ZnSe show type II-like behavior

in the absorption

spectra,

but is still partially

enced by the exciton trapping processes in the recombination The demonstration

that the recombination

erate lattice temperatures

in these specially

influ-

spectra.t208j

spectra at low and mod-

designed

ZnSe/ZnTe

hetero-

structures is dominated by very pronounced exciton trapping effects at the heterointerfaces suggests that this phenomena should by taken into account in general when considering the possible use of ZnTe/ZnSe-based heterostructrues for light emitting purposes. At the same time, the high quantum efficiency reflects also an advanced degree of epitaxial material quality and shows promise for these artificial microstructures in such applications. 3.8

Blue and Blue/Green

Laser Diodes

and LEDs

Very important developments have occurred recently which have resulted in the long anticipated demonstration of the first blue/green semiconductor diode laser. Achievement of blue semiconductor lasers and light emitting devices (LEDs) has been reported independently by two groups, first at 3M Company and a second group representing a collaboration of researchers from Brown and Purdue Universities. The long period of development of (i) MBE-grown II-VI materials, (ii) advances in the understanding of carrier confinement in II-VI-based quantum wells, and (iii) studies of substitutional doping, all of which are documented in this chapter, have paved the way to achievement of injection lasing in the ZnSe material system. first been reported

Significant

levels of p-type doping in ZnSe have

by Park et a1.f*Ogj and Ohkawa

et al.t210) by using a

nitrogen plasma source developed by Oxford Applied Research. Net acceptor levels approaching 1018 cme3 have been achieved for both ZnSe and the wider gap alloy Zn(S,Se) to enable efficient pn junction devices to be fabricated in these materials for the first time.t211)[212j A prototype

device heterostructure

for a simple laser consists

of a

single or multiple quantum well which provides for efficient electron-hole capture once carriers are injected by the outer pn junction cladding layer. As reviewed in this chapter and in other references,f6g)[213j a large variety of II-VI-based QW structures have been investigated to determine potential layered material systems exhibiting an optimal electronic and optical

434

Molecular

band structure. tions generally

Beam Epitaxy

A significant exhibited

in the QW, whereas addition, materials

number of layered

confinement

the lattice constant was typically

II-VI material

either electron confinement

combina-

or hole confinement

of both carriers was not observed.

mismatch

between

the various

In

constituent

large (in excess of 1%) such that strain-induced

defects precipitously reduced the radiative recombination efficiency at operating device temperatures (room temperature). Two useful heterostructures were finally discovered and were based on (Zn,Cd)Se/ZnSeflQo) and ZnSe/Zn(S,Se)f214j quantum well structures. The MBE growth of the (Zn,Cd)Se/ZnSe QWs was first demonstrated by Jeon et al.f lQoj with optical characterization experimentsf215jf216j identifying this system as a strong new candidate for useful carrier confinement. Direct evidence of the importance of this quantum well structure was obtained from absorption studies which showed that the confinement induced a strong enhancement in the electron-hole Coulomb interaction in the (Zn,Cd)Se-based QW. The resultant 2D exciton state exhibited a very strong optical oscillator strength and a binding energy exceeding the optical phonon energy, typically about 40 meV. The large binding energy provided the ability of the exciton to survive as an entity up to room temperature. As a consequence, the radiative recombination also benefited from the excitonic element. Several different pn junction injection lasers, LED and display structures have been fabricated with, as well as without, sulfur-containing alloys.[*ll1[*‘*1[*‘7]-[**‘1 A schematic illustration of one such variety of heterostructure is shown in Fig. 46. The ZnSe-based structures have ZnSe/(Zn,Cd)Se multiple quantum wells (MQWs) embedded in a ZnSe pn homojunction. Alternate Zn(S,Se)-based structures consisted of Zn(S,Se)/ (Zn,Cd)Se MQWs placed within a pn homojunction formed from Zn(S,Se) layers. The Zn(S,Se)-based structures were essentially the same device configuration as the ZnSe-based configurations, except ZnSe is replaced by Zn(S,Se) with a S mole fraction of about 7%. For lasing, the waveguiding was provided by the index difference between the MQW region and the In other configurations (ZnSe/ adjacent binary or S-containing alloy. Zn(S,Se)-based structures),1 211)f212jthe ZnSe/Zn(Cd,Se) MQW region was positioned within a ZnSe region which is then bounded by Zn(S,Se) cladding layers.


Heterostructures

435

(b) 1.M

7sAl1ooA (W 2.0qm

p-2nSe:N

2.Opm

I---1.m

4.0pm x=4.3%

Figure 46. Three zinc-based epittial heterostructures showing the II-VI active region (top), buffer layer, and substrate: (a) ZnSe-based device; (b) Zn(S,Se)based device; (c) ZnSe/Zn(S,Se)-based device.

The formation of dislocations, due to the lattice constant mismatch with the GaAs substrate, in all laser and LED device structures was minimized by growing the II-VI active region on an appropriate (In,Ga)As buffer layer, and in the case of the structures containing sulfur, by correct choice of the S fraction. Optical devices were also grown on both n-type and p-type GaAs substrates. In the case of the ZnSe-based laser and LED structures, a 4.5 pm thick (In,Ga)As buffer layer with 4.3% In was incorporated for lattice matching to the GaAs substrate. For the structures employing Zn(S,Se) layers where the S fraction was approximately 7%, a 1.5pm GaAs buffer layer was grown prior to the growth of the II-VI active layers. In TEM evaluation of the (In,Ga)As epilayers, imaging revealed dislocation densities in the top region of the (In,Ga)As buffer layers to be in the range of lo5 measured lower

within

crne2 or lower. the ZnSe-based

range of lo6

optimization

The estimated

crnm2,and are expected

of the In fraction.

dislocation

densities

II-VI regions were found to be in the to be reduced

The Zn(S,Se)-based

by further

structures were grown

with a sulfur fraction of 7 to 8%. The full width at half of maximum of (400) x-ray rocking curve peaks obtained from the Zn(S,Se) layers ranged between 16 and 65 arcsec. These values were consistent with the lack of dislocations observed in TEM imaging. It is important to emphasize that the injection-induced

photon emis-

sion from both the laser and LED structures originated from the (Zn,Cd)Se MQW region. The carrier confinement is significant in the MQW region

436

Molecular

Beam Epitaxy

providing excitons that have large binding energies (40 mev), approaching an order of magnitude greater than that exhibited in conventional III-V laser structures. excitons,

There is evidence that, as a result of the robustness

the lasing process

lower temperatures,

is dominated

with excitonic

by excitonic

transitions

of the

recombination

still playing

at

a role at room

temperature. This is in contrast to conventional Ill-V lasers where excitons are fully screened by the mobile carriers, and hence play no role in lasing. In LED device operation, the turn-on voltage for forward conduction, typically 5 volts, but as low as about 3 volts, was found to be coincident with the observation of incoherent light emission emanating from the cleaved facets with a spectrum typical of the photoluminescence observed from the (Zn,Cd)Se quantum wells. We noted that the turn-on voltage for the devices having a p+ ZnSe top layer was somewhat higher (typically 12 V) than an n+ top layer: the difference was likely due to a larger potential barrier at the Au-contact. Although laser operation of a specific structure was similar for both n- and p-type GaAs substrates, the LED light intensity of devices formed on p-type GaAs was found to be somewhat brighter than those grown on n-type GaAs; the difference was attributed to the difficulty of forming an ohmic contact to p-type ZnSe. Laser device configurations consisted of 600 pm to 1 mm long cleaved resonator structures having 20-40pm wide stripe electrodes at the top. lndium was usually evaporated as the contact for those structures having an n+ ZnSe top layer; gold was used to contact the p-type top layers. The structures having an n+ top layer exhibited substantial current spreading, especially at T = 77 K (and below), thus the laser structures were fabricated in a mesa configuration.t212) Figure 47 shows the diode laser output power vs. input current density from 77 K to 273 K for a mesa device with uncoated facets. The threshold

current density at T = 77 Kwas

400 Ncm2, or 160 mA for the current of typical

MQW devices.

Devices

were operated CW with mW average output powers at 77 K, and were operated in a pulsed mode at room temperature by both groups. At room temperature, the threshold current density increased to 1500 A/cm2 (corresponding

to 600 mA actual current).

Threshold

currents

of under 100 A/

cm2 were obtained by the group at 3M Company for single quantum well structures at 77 K. Various laser configurations have provided output powers in excess of 700 mW for pulsed operation at low duty cycles. It is important to emphasize that lasing was obtained from each of the different device configurations described above, and for structures grown on both nand p-type substrates. It should be noted that the 3M group demonstrated pulsed lasing at 300 K. and CW at 77 K from devices with coated facets.


Heterostructures

437

T=200K / T=250K

m

T=273K

nt,

“0

400

Current

800

1200

Density

1600

2000

(A/cm2)

Figure 47. The figure shows diode laser output power versus input current density for a mesa device with uncoated facets. The threshold for this device at 77K corresponded to a current of 160 mA. As room temperature is approached, heating, predominantly at the top contact, creates thermal problems. Except for the heating, as room temperature is approached the device is seen to exhibit a T, of approximately 180 K.

LED devices emitting in the blue (494 nm) at room temperature were prepared by cleaving the Zn(S,Se)-based heterostructures into 2 x 2 mm2 pieces which were contacted by a small indium dot. (In the ZnSe-based structures, the emission wavelengths ranged 5100 8, in the blue-green to 4900 8, at room temperature.) applied across the entire device corresponding to the highest

case of the from about The voltage light output

(P = 120 pW/) was 20 volts; however most of this voltage was needed to overcome the built-in contact barriers in order to achieve adequate initial current flow. Considerable improvement in the overall power

quantum efficiency can thus be anticipated when the contact problems are solved. These devices, when compared with Sic and GaN electoluminescent devices, appeared to be the brightest blue LED sources to date. Lateral transport was found to be quite effective in the heterostructures, such that LED emission devices was uniformly visible.

over the entire front surface

of the

438

Molecular

Beam Epitaxy

In conclusion, laser and LED operation in the blue and green spectral region was obtained from a variety of MBE grown II-VI quantum well device configurations. type (In,Ga)As ture.

The devices, prepared on both p-type and n-

or GaAs buffer layers were operated up to room tempera-

Continuous

wave

operation

was obtained

at 77 K.

Substantial

materials-related problems still remain. A major obstacle is to determine techniques for the formation of low resistance contacts to the widegap II-VI semiconductors. Additional areas for future effort include the optimization of growth techniques to maximize the free hole concentration in p-type layers, and to achieve a reduction in point defects which tend to both degrade the quantum efficiency and increase threshold current at room temperature.

4.0

SUMMARY

As is apparent throughout this chapter, the II-VI semiconductor compound family and associated superlattice systems are viable materials of extreme interest to the optoelectronic community. Incorporation of Mn into the II-Vls results in interesting physics and potential device applications; the bandgap modulation and resultant layered structures, however, were of particular emphasis here. Although essentially of Type I, these strained-layer DMS superlattice structures have most of the bandgap difference appearing as conduction band offset. Wide bandgap II-VI superlattices covering virtually the entire visible spectral region were discussed. The ability of the non-equilibrium MBE growth technique to allow for metastable

structures

is highlighted

magnetic ordering of the MnSe/ZnSe

superlattice

in the work studying structures;

the

the MnSe is

present in the hypothetical zincblende crystal structure. New advances have been made in the controlled substitutional doping of the wide bandgap II-VI compounds, High resistivity, undoped ZnSe

material

specie.

shows

Application

low resistivity

of coherent

upon

incorporation

of the dopant

light during the MBE growth has shown

dramatic effects in the doping of CdTe and (Cd,Mn)Te. The growth of II-VI compounds on Ill-V MBE-grown epilayers (and the inverse interface) results in epitaxial interfaces as opposed to epitaxial/bulk substrate interfaces; these heterostructures can be obtained by either surface passivation techniques or by the use of modular MBE systems.

We expect that the significant

reduction in the density and types


Heterostructures

439

of defects provided by the epitaxial interface will greatly enhance utilization of heterojunctions in device applications. Future directions for the MBE growth of II-VI compound semiconductors and their superlattices may include modifications

to the growth technique

including

photon illumi-

nation of the growth front, and the use of gas and ionized beam sources.

ACKNOWLEDGMENT The work at Purdue has benefited greatly from the collaboration of a number of both faculty and students. We would especially like to thank Professors S. Datta, T. Sakamoto, M. Yamanishi, M. Kobayashi, M. R. Melloch, W. M. Becker, and A. K. Ramdas. Substantial contributions are gratefully acknowledged from a large number of students and collaborators whose names appear as co-authors in cited references. The work at Brown has had key contributions by A. Mysyrowicz, S. -K. Chang, D. Lee, and Q. Fu. The sponsors of our research are Office of Naval Research, Air Force Office of Scientific Research, Defense Advanced Research Projects Agency,

and National Science Foundation.

REFERENCES 1. Turner, A. E., Gunshor, 3155 (1983)

R. L., and Datta, S., Appl. Optics, 22:3152-

2. Gaj, J. A., Ginter, J., and Galazka, B89:655462 (1978) 3. Furdyna,

R. R.,

J. K., J. Appl. Phys., 53:7637-7643

Phys.

Status

Solidi.,

(1982)

4. Diluted Magnetic (Semimagnetic) Semiconductors, (R. L. Aggarwal, J. K. Furdyna, and S. von Molnar, eds.), Pittsburgh: Materials Research Society (1987) 5. Diluted Magnetic Semiconductors, eds.), Vol. 25, in Semiconductors Academic Press (1988) 6. Turner, A. E., Gunshor,

(J. K. Furdyna and J. Kossut, and Semimetals, New York:

R. L., and Datta, S., unpublished.

7. Qadri, S. B., Skelton, E. F., Webb, A. W., Kennedy, Lett., 46:257-259 (1985)

J., Appl. Phys.

8. Dean, B. E., and Johnson, C. J., Proceedings ofthe Third international Conference on //-VI Compounds, Monterey, (July 1987)

440

Molecular

Beam Epitaxy

9. Kolodziejski, L. A., Sakamoto, T., Gunshor, Appl. Phys. Left., 44:799-801 (1984) 10. Cho, A. Y., J. Appl. Phys., 47:2841-2843 11. Mar, H. A., Chee, K. T., and Salansky, 239 (1984) 12. Mar, H. A., Salansky, 900 (1984)

Ft. L., and Datta, S.,

(1976)

N., Appl. Phys. Lett., 44:237-

N., and Chee, K. T., Appl. Phys. Left., 44:898-

13. Cheung, J. T., Khoshnevisan, 43:462-464 (1983)

M., and Magee, T., Appl. Phys. Left.,

14. Nishitani, K., Ohkata, R., and Murotani, T., J. Nect. Mater., 12:619623 (1983) 15. Bicknell, R. N., Yanka, Ft. W., Giles, N. C., Schetzina, J. F., Magee, T. J., Leung, C., and Kawayoshi, H., Appl. Phys. Left., 44:313-315 (1984) 16. Faurie, J. P., Sivananthan, S., Boukerche, Phys. Left., 45: 1307-l 309 (1984)

M., and Reno, J., Appl.

17. Otsuka, N., Kolodziejski, L. A., Gunshor, R. L., Datta, S., Bicknell, R. N., and Schetzina, J. F., Appl. Phys. Lett., 46:860-862 (1985) 18. Otsuka, N., Kolodziejski, L. A., Gunshor, R. L., Datta, S., Bicknell, R. N., and Schetzina, J. F., Mater. Res. Sot. Symp. Proc., 37:449-454 (1985) 19. Faurie, J. P., Shu, C., Sivananthan, Science, 168: 473-482 (1986) 20. Schaffer,

S., and Chu,

X.,

Surface

W. J., private communication.

21. Kolodziejski, L. A., Gunshor, R. L., Otsuka, N., Zhang, X. -C., Chang, S. -K., and Nurmikko, A. V., Appl. Phys. Lett., 47:882-884 (1985) 22. Bull, C. J., and Sealy, 8. J., Philosophical (1978) 23. Kolodziejski, L. A., Gunshor, Sri. Techno/., A4:2150-2151 24. Ponce, F. A., Anderson, 168564-570 (1986)

R. D., and Austin,

26. Srinivasa, 50:1441-l

R., Panish, 443 (1987)

A37:489-500

R. L., Otsuka, N. and Choi, C., J. Vat. (1986)

G. B., and Ballingall,

25. Feldman, (1986)

Magazine,

J. M., Surface Science,

R. F., Appl. Phys. Lett., 49:954-956

M. B., and Temkin,

H., Appl.

Phys. Letf.,


27. Cohen-Solal, 1521 (1986)

441

Heterostructures

G., Bailly, F., and Barbe, M.,App/. Phys. Left., 49:1519-

28. Herman, M. A., Jylha, O., and Pessa, M., J. Crysf. Growth, 66:480483 (1984) 29. Bicknell, R. N., Yanka, Ft., Giles-Taylor, N. C., Blanks, D. K., Buckland, E. L., and Schetzina, J. F., Appl. Phys. Let?., 4592-94 (1984) 30. Kolodziejski, L. A., Bonsett, T. C., Gunshor, R. L., Datta, S., Bylsma, R. B., Becker, W. M., and Otsuka, N., Appl. Phys. Lett., 45440-442 (1984) 31. Zhang, X. C., Chang, S. -K., Nurmikko, A. V., Kolodziejski, L. A., Gunshor, R. L., and Datta, S., Phys. Rev., B31:4056-4059 (1985) 32. Bicknell, R. N., Giles, N. C., and Schetzina, 50:691-693 (1987)

J. F., Appl. Phys. Lett.,

33. Williams, G. M., Cullis, A. G., Whitehouse, C. R., Ashenford, and Lunn, B., Appl. Phys. Lett., 55:1303-1305 (1989)

D. E.,

34. Kolodziejski, L. A., Gunshor, R. L., Datta, S., Bonsett, T. C., Yamanishi, M., Frohne, R., Sakamoto, T., Bylsma, R. B., Becker, W. M., and Otsuka, N., J. Vat. Sci. Techno/., B3:714-717 (1985) 35. Choi, C., Otsuka, N., Kolodziejski, L. A., and Gunshor, Res. Sot. Symp. Proc., 56:235-240 (1986) 36. Warnock, J., Petrou, A., Bicknell, R. N., Giles-Taylor, D. K., and Schetzina, J. F., Phys. Rev., B32:8116-25 37. Hefetz, Y., Lee, D., Nurmikko, A. V., Sivananthan, Faurie, J. P., Phys. Rev., 834:4423-4425 (1986)

R. L., Mat.

N. C., Blanks, (1985)

S., Chu, X., and

38. Chang, S. -K., Nurmikko, A. V., Wu, J. -W., Kolodziejski, Gunshor, R. L., Phys. Rev., B37:1191-1198 (1988) 39. Wu, J. -W. and Nurmikko, 1987)

L. A.,

A. V., Phys. Rev., B., 836:4902-4910

40. Gaj, J. A., Planel, R., and Fishman, 29:435-438 (1979)

G.,

So/id State

41. Pessa, M. and Jylha, O., Appl. Phys. Lett., 45:646-648 42. Tersoff, J., Phys. Rev. Lett., 56:2755-2758

Commun., (1984)

(1986)

43. Chang, S. -K., Nakata, H., Nurmikko, A. V., Gunshor, R. L., and Kolodziejski, L. A., Appl. Phys. Lett., 51:667-669 (1987) 44. Vina, L., Chang, L. L., and Yoshino, Y., Proceedings ofthe Conference on Modulated Semiconductor Structures MSS-Ill, Montpelier, France, (June 1987)

442

Molecular

45. Menendez,

Beam Epitaxy

J. and Cardona, M., Phys. Rev., B31:3696-3704

46. Venugopalan, S., Petrou, A., Galazka, R. R., Ramdas, Rodriquez, S., Phys. Rev., B25:2681-2696 (1982) 47. Venugopalan, S., Kolodziejski, L. A., Gunshor, A. K., Appl. Phys. Lett., 45:974-976 (1984)

(1985)

A. K., and

R. L., and Ramdas,

48. Suh, E. -K., Bartholomew, D. U., Ramdas, A. K., Venugopalan, S., Kolodziejski, L. A., and Gunshor, R. L., Phys. Rev., 836:4316-4331 (1987) 49. Chang, S. -K., Nurmikko, A. V., Kolodziejski, L., Phys. Rev., B33:2589-2593 (1986) 50. Smith, D. L., and Mailhiot, (1987) 51. Jackson, S. and McIntyre, 230 (1987) 52. Goncalves

C.,

Phys.

L. A., and Gunshor,

Rev.

Leti.,

R.

58:1264-1267

C., Mat. Res. Sot. Symp. Proc., 89:225-

da Silva, C., Phys. Rev., B32:6962-6964

(1985)

53. Wu, J. -W., Nurmikko, A. V., and Quinn, J. J., Phys. Rev., B34:10801084 (1986) 54. Zhang, X. -C., Chang, S. -K., Nurmikko, A. V., Kolodziejski, L. A., Gunshor, R. L., and Datta, S., Appl. Phys. Lett., 47:59-61 (1985) 55. Bicknell, R. N., Giles-Taylor, N. C., Schetzina, J. F., Anderson, G., and Laidig, W. D., Appl. Phys. Lett. 46:238-240 (1985) 56. Bonsett, T. C., Yamanishi, M., Gunshor, R. L., Datta, Kolodziejski, L. A., Appl. Phys. Letf., 51:499-501 (1987)

N.

S., and

57. Bicknell, R. N., Giles-Taylor, N. C., Blanks, D. K., Schetzina, J. F., Anderson, N. G., and Laidig, W. D., Appl. Phys. Letf., 46:1122-l 124 (1985) 58. Holonyak, (1971)

N. and Scifres, D. R., Rev. Sci. /n&rum.,

12:1885-1886

59. Isaacs, E. D., Heiman, D., Zayhowski, J. J., Bicknell, Schetzina, J. F., Appl, Phys. Lett., 48:275-277 (1986)

R. N., and

60. Bicknell, R. N., Giles, N. C., and Schetzina, 49:1095-l 097 (1986)

J. F., Appl. Phys. Left.,

61. Bicknell, R. N., Giles, N. C., and Schetzina, 49:1735-l 737 (1986)

J. F., Appl. Phys. Left.,

62. Dreifus, D. L., Kolbas, R. M., Harris, K. A., Bicknell, R. N., Harper, R. L., and Schetzina, J. F., Appl. Phys. Lett., 51:931-933 (1987)


63. Benson, J. D. and Summers, (1988)

Heterostructures

443

C. G., J. Crystal Growth, 86:354-361

64. Farrell, H. H., Nahory, Ft. E., and Harbison, Techno/., B6:779-781 (1988)

J. P.,

J. Vat.

Sci.

65. Benson, J. D., Rajavel, D., Wagner, B. K., Benz, R., II, and Summers, C. J., J. Crystal Growth, 95543-544 (1989) 66. Harper, R. L., Jr., Hwang, S., Giles, N. C., Bicknell, R. N., Schetzina, J. F., Lee, Y. R., and Ramdas, A. K., J. Vat. Sci. Technol., A6:26272630 (1988) 67. Harper, R. L., Jr., Hwang, S., Giles, N. C., Schetzina, J. F., Dreifus, D. L., and Myers, T. H., Appl. Phys. Lett., 54:170-l 72 (1989) 68. Ohishi, M., Saito, H., Okano, H., Ohmori, 95:538-540 (1989)

K., J. Crystal

Growth,

69. Gunshor, R. L., Kolodziejski, L. A., Nurmikko, A. V., and Otsuka, N., Semiconductors and Semimetals, (T. Pearsall, ed.), 33:337-409, Academic Press, San Diego (1990) 70. Monfroy, G., Sivananthan, S., Chu, X., Faurie, J. P., Knox, R. D., and Staudenmann, J. L., Appl. Phys. Lett. 49:152-l 54 (1986) 71. Miles, R. H., Wu, G. Y., Johnson, M. B., McGill, T. C., Faurie, J. P., and Sivananthan, S., Appl. Whys. Leff., 48:1383-1385 (1986) 72. Menendez, J., Pinczuk, A., Valladares, J. P., Feldman, Austin, R. F., Appl. Phys. Leff., 50:1101-l 103 (1987)

R. D., and

73. Glass, A. M., Tai, K., Bylsma, R. B., Feldman, R. D., Olson, D. H., and Austin, R. F., Appl. Phys. Left., 53:834-836 (1988) 74. Lee, Y. and Ramdas, A.K., Whys. Rev., B38:10600-10610 75. Wei, S. H. and Zunger, A., Phys. Rev. Lett., 56:2391-2394 76. Ehrenreich,

H., Science, 235:1029-l

(1988) (1986)

035 (1987)

77. Furdyna, J. K. and Kossut, J., J. Superlaffices 2:89-96 (1986)

and Microstructures,

78. McCaldin, J. 0. and McGill, T. C., J. Vat. Sci. Technol. B6:13601363 (1988) 79. Gunshor, R. L., Kolodziejski, L. A., Kobayashi, M., Nurmikko, A. V., and Otsuka, N., Mater. Res. Sot. Proc., 151:141-149 (1989) 80. Durbin, S. M., Han, J., 0, Sungki, Kobayashi, M., Menke, D. R., Gunshor, R. L., Fu, Q., Pelekanos, N., Nurmikko, A. V., Li, D., Gonsalves, J., and Otsuka, N., Appl. Phys. Lett., 55:2087-2089 (1989)

444


81. Fu, Q., Ding, J., Walecki, Kobayashi, (1990)

W., Nurmikko, A. V., Durbin, S. M., M., and Gunshor, R. L., Phys. Rev., 641:996&9970

81a. Fu, Q., Pelekanos, N., Nurmikko, A.V., Durbin, S., Han, J., Kobayashi, M., and Gunshor, R. L., Surface Science, 229:148-l 51 (1990) 82. Fu, Q., Pelekanos, N., Ding, J., Nurmikko, A. V., Durbin, S., Han J., Kobayashi, M., and Gunshor, R. L., Paper presented at Int. Conf. Quantum Electronics, Anaheim, CA (1990) 83. Pelekanos, N., Ding, J., Fu, Q., Nurmikko, A. V., Kobayashi, Gunshor, Ft. L. Phys. Rev., B43:9354-9358 (1991)

M., and

84. Mackey, K. J., Allen, P. M. G., Herrenden-Harker, W. G., Williams, R. H., Whitehouse, C. R., and Williams, G. M., Appl. Phys. Lefl., 49:354-356 (1986) 85. van Welzenis, R. G. and Ridley, 27:113-l 20 (1984) 86. Sugiyama,

B. K.,

So/id State Electronics,

Koichi, J. Crystal Growth, 60:450-452

87. Farrow, R. F. C., Jones, G. R., Williams, Appl. Phys. Left, 39:954-956 (1981)

(1982)

G. M., and Young,

I. M.,

88. Wood, S., Greggi, J., Jr., Farrow, R. F. C., Takei, W. J., Shirland, A., and Noreika, A. J., J. Appl. Phys., 55:4225-4231 (1984)

F.

89. Williams, G. M., Whitehouse, C. R., Chew, N. G., Blackmore, W., and Cullis, A. G., J. Vat. Sci. Technol., B3:704-708 (1985)

G.

90. Zahn, D. R. T., Mackey, K. J., Williams, R. H., Munder, H., Geurts, J., and Richter, W., Appl. Phys. Let, 50:742-744 (1987) 91. Kolodziejski, L. A., Gunshor, R. L., Otsuka, N., and Nurmikko, A. V., Mat. Res. Sot. Symp., 102:113-123 (1988) 92. Glenn, J. L., 0, Sungki, Kolodziejski, L. A., Gunshor, R. L., Kobayashi, M., Li, D., Otsuka, N., Haggerott, M., Pelekanos, N., and Nurmikko, A. V., J. Vat. Sci. Technol., B7:249-252 (1989) 93. Farrow, R. F. C., Noreika, A. J., Shirland, F. A., Takei, W. J., Wood, S., Greggi, J., Jr., and Francombe, M. H., J. Vat. Sci. Techno/., A2:527-528 (1984) 94. Westwood, D. I., Mackey, K. J., Allen, P. M. G., Herrenden-Harker, W. G., and Williams, R. H., J. Crysfa! Growth, 81:506-507 (1987) 95. Golding, T. D., Martinka, 64:1873-l 877 (1988)

M., and Dinan,

J. H.,

J. Appl.

Phys.,

96. Neave, J. H., Blood, P., and Joyce, B. A., Appl. Phys. Left., 36:31 l313 (1980)


97. Missous,

445

Heterostructures

M. and Singer, K. E., Appl. Phys. Lett., 50:694-696

(1987)

98. Gunshor, R. L., Kolodziejski, L. A., Melloch, M. R., Otsuka, N., and Nurmikko, A. V., Growth and Optical Properties of Wide-Gap II-VI Low Dimensional Structures, (T. C. McGill, C. M. Sotomayor Torres, and W. Gebhardt, eds.), NATO ASI Series B, 200:229-238, Plenum Publishing Corporation, New York (1988) 99. Smith, D. L. and Pickhardt, (1975)

V. Y., J. Appl. Phys., 41:2366-2374

100. Yao, T., Miyoshi, Y., Makita, Y., and Maekawa, Phys., 16:369-370 (1977) 101. Kitagawa, F., Mishima, T., and Takahashi, Soc.,127:937-943 (1980) 102. Yao, T., Ogura, M., Matsuoka, Lett., 43:499-501 (1983) 103. Park, Ft. M. and Salansky, (1984)

S., Jap. J. Appl.

K.,

S., and Morishita,

J. Nectrochem. T., Appl. Phys.

N. M., Appl. Phys. Lett., 44:249-251

104. Yoneda, K., Hishida, Y., Toda, T., Ishii, H., and Niina, T., Appl. Phys. Lett., 45:1300-1302 (1984) 105. Kolodziejski, L. A., Gunshor, Ft. L., Bonsett, T. C., Venkatasubramanian, R., Datta, S., Bylsma, R. B., Becker, W. M., and Otsuka, N., Appl. Phys. Lett., 47:169-171 (1985) 106. Kolodziejski, L. A., Gunshor, R. L., Nurmikko, A. V., and Otsuka, N., Thin Film Growth Techniques for Low-Dimensional Structures, (R. F. C. Farrow, S. S. P. Parkin, P. J. Dobson, J. H. Neave, and A. S. Arrott, eds.), 163:247-260, Plenum Publishing Corporation, New York (1987) 107. Gunshor, R. L., Kolodziejski, L. A., Melloch, M. R., Vaziri, M., Choi, C., and Otsuka, N., Appl. Phys. Lett., 50:200-202 (1987) 108. Tamargo, M. C., de Miguel, J. L., Hwang, D. M., and Farrell, H. H., J. Vat. Sci. Technol., B6:784-787 (1987) 109. Park, R. M., Mar, H. A., Salansky, B3:676-680 (1985)

N. M., J. Vat. Sci. Techno/.,

110. Prinz, G. A., Jonker, B. T., Krebs, J. J., Ferrari, J. M., and Kovanic, F. Appl. Phys. Lett., 48:1756-l 758 (1986) 111. Cammack, D. A., Dalby, R. J., Cornelissen, Appl. Phys., 62:3071-3074 (1987)

H. J., and Khurgin, J. J.

112. Yao, T., and Takeda, T., Appl. Phys. Lett., 48:160-l

62 (1986)

446


113. Joyce, B. A., Neave, J. H., Dobson, P. J., and Larson, P. K., Phys. Rev., B29:814-819

(1984)

114. Lewis, B. F., Lee, T. C., Grunthaner, F. J., Madhukar, A., Fernanadez, IX, and Maserjian, J., J. Vat. Sci. Techno/., B2:419-424 (1984) 115. Tamargo, M. C., de Miguel, J. L., Hwang, D.d M., and Farrell, H. H., J. Vat. Sci. Technol., B6(2): 784--787(1988) 116. Tamargo, M. C., de Miguel, J. L., Turco, F. S., Skromme, B. J., Hwang, D. M., Nahory, R. E., and Farrell, H. H., Growth and Optical Properties of Wide Gap II-VI Low Dimensional Structures, (T. C. McGill, C. M. Sotomayer Torres, and W. Gebhardt, eds.), 200:239 243, Plenum Publishing Corporation, New York (1988) 117. Farrell, H. H., Tamargo, M. C., and de Miguel, Technol., B6:767-768 (1988)

J. L., J. Vat. Sci.

118. Skromme, B. J., Tamargo, M. C., Turco, F. S., Shibli, S. M., Nahory, R. E., and Bonner, W. A., ElectrochemicalSociety Meeting, Chicago, (Oct. 1988) 119. Gunshor, R. L. and Kolodziejski, 24:1744-57 (1988)

L. A., IEEE J. Quantum Electronics,

120. Qadri, S. B., Jonker, B. T., Prinz, G. A., and Krebs, J. J., J. Vat. Sci. Technol., A6: 1526-l 530 (1988) 121. Park, R. M., Mar, H. A., Salansky, 83: 1637-l 640 (1985)

N. M.,

J. Vat. Sci.Technol.,

122. Isshiki, M., Yoshida, T., Igaki, K., Uchida, W. and Suto, S., J. Crystal Growth, 72: 162-l 66 (1986) 123.

Park, R. M. and Mar, H. A., J. Mater. Res., 1:543-546

(1986)

124. Park, R. M., Mar, H. A., and Kleiman, J., J. Crystal Growth, 86:335341 (1988) 125. Mino, N., Kobayashi, M., Konagai, Phys., 58:793-796 (1985)

M., and Takahashi,

K., J. Appl.

126. Venkatasubramanian, R., Otsuka, N., Datta, S., Kolodziejski, and Gunshor, R. L., Growth of Compound Semiconductors, Gunshor, H. Morkoc, ed.), Proc. SPIE, 796:121-129 (1987) 127. Madhukar,

A., Surface Science, 132:344374

(1983)

128. Singh, J. and Bajaj, K. K., J. Vat. Sci. Technol., B2:276-279 129. Yao, T., Crystal Growth, 72:31-40 130. Yao, Takafumi,

(1985)

Jap. J. Appl. Phys., 25:821-827

L. A., (R. L.

(1986)

(1984)


131. Ohkawa, K., Mitsuyu, 86:329-334 (1988)

Heterostructures

T., and Yamazaki,

O.,

132. Park, R. M., Mar, H. A., and Salansky, 58:1047-l 049 (1985)

J. Cry&r/

N. M.,

J. Appl.

447

Growth, Phys.

133. Yao, T., and Taguchi, T., Proceedings of the 73th international Conference on Defects in Semiconductors, Coronado, CA, (Aug. 12-l 7,1984) 134. Cheng, H., DePuydt, J. M., Potts, J. E., and Smith, T. L., Appl. Phys. Leti., 52:147-149 (1988) 135. Yao, T., Makita, Y., and Maekawa, (1979)

S., Appl. Phys. Letf., 35:97-98

136. Yao, T., The Technology and Physics of Molecular Beam Epitaxy, (E. H. C. Parker, ed.), pp. 313-343, Plenum Press, New York (1985) 137. Ohkawa, K., Mitsuyu, T., and Yamazaki, O., Extended Abstracts of the 18th Conference on Solid State Devices and Materials, Tokyo, pp. 635-638 (1986) 138. Cheng, H., DePuydt, J. M., Potts, J. E., Mohapatra, S. K., and Smith, T. L., Growth of Compound Semiconductors, (R. L. Gunshor, H. Morkoc, eds.), 796:91-97, Proc. SPIE (1987) 139. Yao, T., Ogura, M., Matsuoka, Phys., 22:L144-L146 (1983)

S., and Morishita,

T., Jap. J. Appl.

140. Yao, T., Okada, Y., Matsui, S., Ishida, K., and Fujimoto, I., J, Crystal Growth, 81:518-523 (1987) 141. Kolodziejski, L. A., Gunshor, R. L., Melloch, M. R., Vaziri, M., Choi, C., and Otsuka, N., Growth of Compound Semiconductors, (R. L. Gunshor and H. Morkoc, ed.), 796:98-l 03 Proc. SPIE (1987) 142. Lee, Y. R., “Piezo- and photo-modulation studies of semiconductors: I. dilute magnetic semiconductors; II. GaAs/AIGaAs, ZnSe, and CdMnTe heterostructures,” Ph. D Thesis, Purdue University (1987) 143. Blacha, A., Presting, B126:11-36 (1984)

H., and Cardona,

143aLee, Y. R., Ramdas, A. K., Kolodziejski, Phys. Rev., 838:13143-l 3149 (1988)

M.,

Status

Solidi,

L. A., and Gunshor,

R. L.,

144. Cullis, A. G., Chew, N. G., and Hutchison, 17:203-211 (1985)

Phys.

J. L., Ultramicroscopy,

145. Ourmazd, A., Tsang, W. T., Rentshler, J. A., and Taylor, Appl. Phys. Left, 50: 1417-l 419 (1987)

D. W.,

448


146. Studtmann,

G. D., Gunshor, R. L., Kolodziejski, L. A., Melloch, M. Ft., Otsuka, N., Munich, D. P., Cooper, J. A., and Pierret, R. F., Paper presented at the 45th Device Research Conference, Santa Barbara, CA, (June 22-24, 1987)

147. Studtmann, G. D., Gunshor, R. L., Kolodziejski, L. A., Melloch, M. R., Cooper, J. A., Pierret, R. F., and Munich, D. P., Appl. Phys. Lett., 52:1249-1251 (1988) 148. Qian, Q. D., Qiu, J., Glenn, J. L., Sungki, O., Gunshor, R. L., Kolodziejski, L. A., Kobayashi, M., Otsuka, N., Melloch, M. R., and Cooper, J. A., J. Crystal Growth, 95567-571 (1989) 149. Qian, Q. D., Qiu, J., Melloch, M. R., Cooper, J. A., Kolodziejski, A., and Gunshor, R. L., Appl. Phys. Lett., 54:1359-1361 (1989)

L.

150. Qian, Q. D., Qiu, J., Kobayashi, M., Gunshor, R. L., Melloch, M. R., and Cooper, J. A., J. Vat. Sci. Technol., B7:793-789 (1989) 151. Qiu, J., Qian, Q. D., Gunshor, R. L., Kobayashi, M., Menke, D. R., Li, D., and Otsuka, N., Appl. Phys. Lett., 56:1272-1274 (1990) 152. Kobayashi,

N., Jpn. J. Appl. Phys., 27:L1597-1599

153. Kowalczyk, S. P., Kraut, E. A., Waldrop, Vat. Sci. Technol., 21~482-485 (1982)

(1988)

J. R., and Grant, R. W., J.

154. Li, D., Gonsalves, J. M., Otsuka, N., Qiu, J., Kobayashi, Gunshor, R. L., Appl. Phys. Lett., 57(5): l-3 (1990)

M. and

155. Li, D., Otsuka, N., Qiu, J., Glenn, J., Jr., Kobayashi, Gunshor, R. L., Proc. Mat. Res. Sm. Symp., 161 :127-132

M., and (1990)

156. Hahn, H., and Klingler, W., Z. Anorg. Chemie., 259:135-142

(1949)

157. Hefetz, Y., Nakahara, J., Nurmikko, A. V., Kolodziejski, Gunshor, R. L., and Datta, S., Appl. Phys. Lett., 47:989-991

L. A., (1985)

158. Jonker, B. T., Krebs, J. J., Qadri, S. B., and Prinz, G. A., Appl. Phys. Lett., 50:848-850 (1987) 159. Jonker, B. T., Qadri, S. B., Krebs, J. J., and Prinz, G. A., J. Vat. Sci. Technol., A6: 1946-l 949 (1988) 160. Jonker, B. T., Krebs, J. J., Qadri, S. B., Prinz, G. A., Volkening, and Koon, N. C., J. Appl. Phys., 63:3303-3305 (1988) 161. Jonker, B. T., Krebs, J. J., and Prinz, 53:450-452 (1988)

G. A.,

Appl.

F.,

Phys. Lett.,

162. Jonker, B. T., Qadri, S. B., Krebs, J. J., Prinz, G. A., and SalamancaYoung, L., J. Vat. Sci. Techno/., A7:1360-1365 (1989) 163. Krebs, J. J., Jonker, B. T., and Prinz, Magnetics., 24:2548-2550 (1988)

G. A.,

/fEE

Trans.

on


449

Heterostructures

164. Liu, X., Petrou, A., Jonker, B. T., Prinz, G. A., Krebs, J. J., and Warnock, J., J. Vat. Sci. Technol., A6:1508-1510 (1988) 165. Liu, X., Petrou, A., Jonker, B. T., Prinz, G. A., Krebs, J. J., and Warnock, J., Appl. Phys. Left, 53:476-478 (1988) 166. Qadri, S. B., Jonker, B. T., Prinz, G. A., and Krebs, J. J., Thin So/id Films, 164: 11 l-l 14 (1988) 167. Munekata, H., Ohno, H., von Molnar, S., Harwit, A., Segmuller, and Chang, L. L., Phys. Rev. Lett., 63:1849-l 852 (1989)

A.,

168. Kolodziejski, L. A., Gunshor, R. L., Otsuka, N., Datta, S., Becker, W. M., and Nurmikko, A. V., /EEE J. Quantum Electronics, QE-22:16661 676 (1986) 169. Twardowski, 48:845-848

A., Dietl, T., and Demianuk, (1983)

M., So/id State Cornmun.,

170. Yoder-short, D. R., Debska, U., and Furdyna, 58:4056-4060 (1985)

J. K., J. Appl. Phys.,

171. Gunshor, R. L., Kolodziejski, L. A., Otsuka, Surface Sci., 174:522-533 (1986)

N., and Datta,

172. Hefetz, Y., Goltsos, W. C., Nurmikko, A. V., Kolodziejski, Gunshor, R. L., Appl. Phys. Leff., 48:372-374 (1986)

S.,

L. A., and

173. Hefetz, Y., Goltsos, W. C., Lee, D., Nurmikko, A. V., Kolodziejski, L. A., and Gunshor, R. L., Superlattices and Microstructures, 2:455458 (1986) 174. Fu, Q., Nurmikko, A. V., Kolodziejski, Appl. Phys. Lett., 51:578-580 (1987)

L. A., and Gunshor,

R. L.,

175. Miller, D. A. B., Chemla, D. S., Damen, T. C., Gossard, A. C., Wiegmann, W., Wood, T., and Burrus, C. A., Phys. Rev., B32:10431060 (1985) 176. Wu, J. -W., and Nurmikko, (1987)

A. V.,

Phys.

Rev.,

836:4902-4905

177. see e.g., Chemla, D. S., Schmitt-Rink, S., and Miller, D. A. B., Nonlinear Optical Properties of Semiconductors, (H. Haug, ed.), Academic Press, New York (1987) 178. Schmitt-Rink, S., Chemla, 832:6601-6609 (1985)

D. S., Miller,

D. A. B.,

Phys.

Rev.,

179. Andersen, D. R., Kolodziejski, L. A., Gunshor, R. L., Datta, S., Kaplan, A. E., and Nurmikko, A. V., Appl. Phys. Lett., 48:1559-1661 (1986) 180. Andersen, D. R., “Nonlinear optical properties of II-Vi semiconductor compounds grown by molecular beam epitaxy,“Ph.D. Thesis, Purdue University (1986)

450


181. Peyghambarian, N., Park, S. H., Koch, S. W., Jeffery, A., Potts, J. E., and Cheng, H., Appl. Phys. Left., 52:182-l 84 (1988) 182. Kleinman,

D. A., Phys. Rev., B28:871-879

183. Miller, R. C., Kleinman, D. A., Gossard, Phys. Rev., 82565456547 (1982)

(1983) A. C., and Munteanu,

O.,

184. Nozue, Y., Itoh, M., and Cho, K., J. Phys. Sot. Jap., 50:889-894 (1981) 185. Fu, Q., Lee, D., Mysyrowicz, A., Nurmikko, A. V., Gunshor, and Kolodziejski, L. A., Phys. Rev., B 37:8791-8794 (1988)

R. L.,

186. Fu, Q., Lee, D., Mysyrowicz, A., Nurmikko, A. V., Gunshor, and Kolodziejski, L. A., Phys. Rev., B 37:8751-8755 (1988)

R. L.,

187. Bylsma, Ft. B., Becker, W. M., Bonsett, T. C., Kolodziejski, L. A., Gunshor, R. L., Yamanishi, M., and Datta, S., Appl. Phys. Let?., 47:1039-l 041 (1985) 188. Catalano, I. M., Cingolani, A., Ferrara, State Commun., 43:371-374 (1982)

M., and Lugara,

M., Solid

189. Potts, J. E., Smith, T. L., and Cheng, H., Appl. Phys. Leti., 50:7-9 (1987) 190. Jeon, H., Ding, J., Nurmikko, A. V., Luo, H., Samarth, N., Furdyna, J., Bonner, W., and Nahory, R., Appl. Phys. Letf., 57:2413-2415 (1990). 191. Bonsett, T. C., Yamanishi, M., Gunshor, R. L., Datta, Kolodziejski, L. A., Appl. Phys. Lett., 51:499-501 (1987) 192. Bylsma, R. B., Ph.D Thesis, Purdue University,

S., and

May (1986)

193. Kolodziejski, L. A., Gunshor, R. L., Otsuka, N., Gu, B. P., Hefetz, Y., and Nurmikko, A. V., Appl. Phys. Lett., 48:1482-1484 (1986) 194. Gunshor, R. L., Kolodziejski, L. A., Otsuka, N., Gu, B. P., Lee, Hefetz, Y., and Nurmikko, A. V., Superlattices and Microstructures, 35-8 (1987) 195.

Kolodziejski, L. A., Gunshor, R. L., Otsuka, N., Gu, B. P., Hefetz, Y., and Nurmikko, A. V., J. Crystal Growth, 81:491-494 (1987)

196. Chang, S. K, Lee, D., Nakata, H., Nurmikko, A. V., Kolodziejski, A., and Gunshor, R. L., J. Appl. Phys., 62:4835-4838 (1988)

L.

197. Nurmikko, A. V., Lee, D., Hefetz, Y., Kolodziejski, L. A., and Gunshor, R. L, Proceedings of the 18th International Conference on the Physics ofSemiconductors, Stockholm, p. 775, (0. Engstrom, Proc. Ed.), World Scientific (1986)


Heterostructures

451

198. Fujiyasu, H., Mochizuki, K., Yamazaki, Y., Aoki, M., Kuwabara, H., Nakanishi, Y., Shimaoka, G., Surface Science, 174343-547 (1986) 199. Takeda, T., Kurosu, 174348-549 (1986)

T., Lida, M., and Yao, T.,

Surface

M., and Takahashi,

Science,

200.

Kobayashi, M. , Mino, N., Konagai, Science, 174:550-555 (1986)

K., Surface

201.

Kolodziejski, L. A., Gunshor, R. L., Fu, Q., Lee, D., Nurmikko, A. V., Gonsalves, J. M., and Otsuka, N., Appl. Phys. Letf., 52:1080-2 (1988)

202.

Kobayashi, Takahashi,

203.

Lee, D., A. Mysyrowicz, A., Nurmikko, A. V., and Fitzpatrick, Phys. Rev. Lett., 58:1475-l 478 (1987)

M., Mino, N., Katagiri, H., Kimura, R., Konagai, K., J. Appl. Phys., 60:773-778 (1986)

204. Yao, T. , Makita, Y., and Maekawa, 312 (1978)

M., and 8. J.,

S., J. Crystal Growth, 45:309-

205.

Fu, Qiang, Lee, D., Nurmikko, A. V., Kolodziejski, R. L., Phys. Rev., B39:3173-3177 (1989)

L. A., and Gunshor,

206.

Lee, Q. F., Nurmikko, A. V., Gunshor, R. L., and Kolodziejski, Superlattices and Microstructures, 5:345-347 (1989)

L. A.,

207. Trzeciakowski, W., Hawrylak, P., Aers, G., and Nurmikko, So/id State Con-m, 71:653-656 (1989)

A. V.,

208.

Ding, J., Fu, Q., Pelekanos, N., Watecki, W., Nurmikko, A. V., Durbin, S. M., Han, J., Kobayashi, M., and Gunshor, R. L., Proc. of 20th Int. Conf. on Physics of Semiconductors, World Publishing, Thessaloniki, Greece (1990)

209.

Park, R. M., Troffer, M. B., Rouleau, C. M., DePuydt, Haase, M. A., Appl. Phys. Left., 57:2127-2129 (1990)

210.

Ohkawa, K., Karasawa, 3O:L152-L155 (1991)

211.

Haase, M. Qiu, J., DePuydt, 59:1272-1274 (1991)

T., and Mitsuyu,

J. M. and

T., Jap. J, ofApp/.

J., and Cheng,

Phys.,

H., Appl. Phys. Lett.,

212. Jeon, H., Ding, J., Patterson, W., Nurmikko, A. V., Xie, W., Grille, D. C., Kobayashi, M., and Gunshor, R. L., Appl. Phys. Leff.,_59:36193621 (1991) 213.

Kolodziejski, L. A., Gunshor, R. L., and Nurmikko, A. V., Compound Semiconductor Strained-Layer Superlattices, (R. M. Biefeld, ed.), pp. 199-230, Trans Tech Publications, Switzerland (1989)

452

Molecular

Beam Epitaxy

214.

Nakanishi, K., Suemune, I., Yoshihisa, F., Kuroda, Y., and Yamanishi, M., Jap. J. ofApp/. Phys., 3O:L1399-L1401 (1991)

215.

Ding, J., Jeon, H., Nurmikko, A.V., Luo, H., Samarth, Furdyna, J.K., Appl. Phys. Left, 57:2756-2758 (1990)

216.

Ding, J., Pelekanos, N., Nurmikko, A. V., Luo, H., Samarth, N., and Furdyna, J. K., Appl. Phys. Left., 57:2885-2887 (1990); Pelekanos, N. T., Ding, J., Nurmikko, A. V., Luo, H., Samarth, N. and Furdyna, J. K., Phys. Rev., B 45:6037-6042 (1992)

N., and

217. Xie, W., Grille, D. C., Gunshor, R. L., Kobayashi, M., Hua, G. C., Otsuka, N., Jeon, H., Ding, J., and Nurmikko, A. V., Appl. Phys. Lett., 60:463-465 (1992) 218.

Jeon, H. Ding, J., Nurmikko, A. V., Xie, W., Grille, D. C., Kobayashi, M., Gunshor, R. L., Hua, G. C., and Otsuka, N., Appl. Phys. Leti., 60(17):2045-2047 (1992)

219. Xie, W., Grillo, D. C., Gunshor, R. L., Kobayashi, M., Jeon, H., Ding, J., Nurmikko, A. V., Hua, G. C., and Otsuka, N., Appl. Phys. Lett., 60(16):1999-2001 (1992) 220.

Hagerott, M. M., Jeon, H., Nurmikko, A. V., Xie, W., Grille, D. C., Kobayashi, M., and Gunshor, R. L., Appl. Phys. Letf., 60:2825 (1992)

221.

Xie, W., Grille, D. C., Gunshor, R. L., Kobayashi, M., Hua, G. C., Otsuka, N., Jeon, H., Ding, J., and Nurmikko, A. V., J. Vat. Sci. and Technol., B10(2):921-923 (1992)

Elemental Semiconductor Heterostructures-Growth, Properties, and Applications Vijay I? Kesan and Subramanian

1 .O

S. /yer

INTRODUCTION

Silicon, the mainstay of VLSI technology, owes its predominance to its processing simplicity. However, its elemental nature precludes the elegant sophistication of band-engineered device design that is common in the compound semiconductor material systems. Nevertheless, several elemental semiconductor-based heterostructures-Si/silicide (CoSi, and NiSiJ structures, Si/(Ca, Ba, Sr)F, structures, SnGe alloys, microcrystalline (PC) hydrogenated silicon (Si:H), SIC, and Si, _xGex, Si, _yCy,and Si, _x,Ge,C, alloys-have been studied in the last few years. Amongst these, we have seen significant developments in the growth of high quality Si-based heterostructures in the Si/Si,,Ge, material system. These alloys have inspired

new research

in Si-based

band-engineering

physics

and device

technology. In this chapter, we review the principal developments in Si,,Ge, growth and applications, and identify key research areas. The reader is also referred to the Proceedings of the International Symposia on Si MBE, the Topical Conferences on Si-based chapters for more details.f’)-fg]

2.0

GROWTH

Silicon compositional

OF Si,,Ge,

heterostructures,

and other review

ALLOYS

and germanium are completely miscible over the entire range and give rise to alloys that retain the diamond crystal

453

454

Molecular

structure.

Beam Epitaxy

The lattice constant

given by Vegard’s

aSi(,_xjGe,=

Eq. (1)

of Si,,Ge,

alloys at room temperature

is

rule

asi + x(aGe - asi)

for low atomic concentrations (x) of Ge. The lattice mismatch between Ge and Si is 4.17% at room temperature and increases only slightly with temperature. Early work on bulk Si,,Ge, alloys by Braunstein and co-workersflO) showed that the bandgap as determined by optical absorption decreased with increasing Ge content. There was strong motivation to grow such alloys epitaxially on Si so that the bandgap change with Ge content could be exploited in heterostructure devices. For pseudomorphic growth of Si,, Ge, on Si, the in-plane lattice constant must be accommodated to that of the Si substrate. The crystal must therefore undergo a tetragonal distortion along the growth direction. The loss of cubic symmetry also affects the band structure of the pseudomorphic Si,,Ge, film by splitting the conduction band and valence band minima (see Sec. 5). From a material stability point of view, the tetragonally distorted crystal is elastically strained, leading to metastability of the grown film. The stability criteria and strain relief of these films via misfit dislocations have been dealt with by van der Merwef11)f12) and by Matthews and Blakesleef13)f14) and is addressed in Sec. 3. In addition to strain, heteroepitaxy of such alloys has to contend with morphological instabilities that arise due to differences in surface energy between the two materials. In the early seventies, there were several attempts to grow epitaxial Si,,Ge, alloys on Si substrates

using MBE by Kasper and co-workersf15) and later using CVD

by Manasavit

and co-workers.f16] These early efforts were plagued by poor

surface

morphology

due to three-dimensional

growth

and limited

Ge

content in the films primarily because of the relatively high temperature employed for growth. In 1984, Bean and co-workers achieved a notable result in the growth of Si,_,Ge, alloys.f171 In these experiments,

using MBE,

they showed that lowering the growth temperature well below what was conventional at that time for Si, preserved the metastability of the film both in terms of defect generation and film morphology. This result sparked renewed interest in the growth of pseudomorphic Si,,Ge, alloys. There have been several refinements to the MBE growth of Si,,Ge, alloys and The extension of MBE to other low Si/Si,,Ge, multilayer structures. temperature growth techniques will be discussed later in this section.

Elemental

The mechanisms are kinetically inhibited

Semiconductor

Heterostructures

455

that lead to relaxation and degraded morphology at lower growth temperatures. Since the driving

force for such mechanisms

is Ge content,

going to higher

Ge content

requires the use of lower growth temperatures. Figure 1 shows a plot of growth temperature as a function of Ge content in the film identifying the regimes which result in planar (2D) and islanded (30) growth. At high Ge content it is necessary to drop the growth temperature even further to attain sharp Si/Si,,Ge, heterointerfaces (see Fig. 1). The growth of Si on top of Ge or Si,,Ge, introduces additional complications. There is a tendency for Ge to segregate to the surface through thin Si overlayers. This has been observed by Raman spectroscopy and Medium Energy Ion Scattering (MEIS) as shown in Figs. 2 and 3.[‘s] In this study a thin fourmonolayer (ML) Ge film is sandwiched between Si layers grown on a Si(lO0) substrate. The Raman and MEIS spectra are shown for different growth temperatures. The Raman spectrum, taken at room temperature, of a thick crystalline Ge film on a Si substrate covered by a crystalline Si

800 lslandedGrowth

600 Planar Growth

200

’

I

0

0.2

I

I

0.4 0.6 Ge Fraction

1

0.8

I

1.0

Figure 1. Growth temperature as a function of Ge content in the Si,_xGe, alloy film showing regimes that result in 2D growth (planar) and 3D growth (islanded). At high Ge content it is necessary to drop the growth temperature even further to attain sharp Si/Ge interfaces.

466

Molecular

layerwould

Beam Epitaxy

show sharp lines near 301 and 520 cm-’ due to the k = 0 optical

modes of Si and Ge. For a pseudomorphic would

film, the Ge phonon

shift to higher values due to the strain introduced

by the lattice mismatch

between

Si and Ge. A more dramatic

the spectrum arises from the existence

frequency

into the Ge layer change

in

of bonds between Ge and Si which

introduces new Raman active optical phonons into the vibrational density of states near 400 cm-‘. The observation of Ge-Ge-like (near 300 cm-‘), Si-Ge-like (near 400 cm-l), and Si-Si-like vibrational modes in the alloy system shows the importance of local bonding in determining the Raman active vibrational

structure.

150

250

350

Raman Shift Figure 2. Raman scattering spectra in (1OO)Si for different growth/anneal Si,,,Ge,,, alloy is shown in curve (d) corresponding to unstrained Ge-Ge,

450

550

(cm\

for a 4 monolayer (- 6 A) Ge film embedded conditions (a)-(e). The spectrum for a thick for comparison. The positions of the shifts Ge-Si, and Si-Si are indicated.

Elemental

(e)250°

Semiconductor

Heterostructures

457

c

zT& 310

320

330

340

ENAGY GeW Figure 3. Random (solid lines) and channeled (dotted) MEIS spectra for Ge embedded in Si(lO0) for the different growth temperatures and post-growth anneals indicated. The interface and surface positions on the spectra are indicated.

458

Molecular

Beam Epitaxy

The relative intensities

of the Ge-Ge and Ge-Si lines in Si-Ge alloys

are a sensitive probe of the alloy composition and the degree of intermixGe ing in a Ge layer imbedded in bulk Si. For the case of a four-monolayer film sandwiched

between silicon substrate

and cap layers, the ratio of Ge-

Ge to Ge-Si bonds is 3, close to what we observe for the sample grown at 250°C in Fig. 2. This ratio will decrease if there is appreciable intermixing between the Ge and Si layers. Such a decrease is progressively observed with increasing growth temperatures and this trend is shown in Fig. 2 for samples grown at 35O’C and 6OO’C. While the quantitative details of the degree of mixing depend on a variety of experimental factors, the Raman spectra in Fig. 2 provide direct information on the ordering and the amount of intermixing in a simple 4 ML Ge film grown at different temperatures. The lower temperature temperature growth.

growth produces

less intermixing

than the higher

Figure 2(a) shows that growth of a Ge layer in Si(lO0) by MBE at 15O’C results in Raman spectra showing no new sharp features but only broad bands characteristic of amorphous Si and Ge. Growth of Ge on Si at 15O’C is not epitaxial but disordered. On the other hand, growth at 25O’C results in the disappearance of over 95% of the amorphous scattering background and the appearance of a sharp line just above 300 cm-‘. This mode is about 7 cm-l above the energy of the Raman active mode in bulk Ge. Such a Raman frequency is consistent with the existence of a thick Ge,,,,Si,,,, alloy grown on Si(100). Such an alloy composition is also consistent with the bond composition of an unalloyed four-monolayer Ge sample if we count the Ge-Ge and Ge-Si bonds. There is a relatively weak Ge-Si mode near 418 cm-l in Fig. 2 (b) which is of comparable width to the line at 300 cm-‘, suggesting that disorder effects on the modes due to the interfaces are no more severe than on the interior “bulk” modes. These spectral features, including the relative intensities, are consistent with epitaxial growth of a well defined Ge layer at 250°C. The Raman spectra in Figs. 2 (c)and (d) show significant changes when the sample is grown at temperatures above 250°C. The ratio of Ge-Si to Ge-Ge scattering intensity

increases with growth temperature.

There is also a shift to lower

energies on the part of Ge-Ge scattering. Both of these changes are consistent with intermixing of Si into the 3-4 monolayer Ge film. The Raman spectrum of the sample grown at 6OO’C suggests that the Ge layer contains more Si than Ge. This can be seen by comparing Fig. 2(d) with the spectrum of the homogeneous Si,,,Ge,,, alloy shown in Fig. 2(f).

Elemental

Semiconductor

In Fig. 3 we show the intensity atoms for growth temperatures 600°C. annealed

Figure

of He+ ions backscattered

of (a) 150°C, (&I) 250°C

3(e) is for a sample

at 600°C for 8 hours.

well as the interface

between

Heterostructures

grown

The position the capping

The dotted (solid) lines are for channeled

at 250%

459

from Ge

(c) 350% and (d) and post-growth

of the Ge surface peak, as

Si layer and Ge, is indicated. (randomly

incident)

He+ ions.

The samples prepared at 150% show near ideal Ge peak widths in the random spectra (see Fig. 3a). However, since this film is not epitaxially grown, the minimum yield is near unity h,r” = 1.0). As the growth temperature is increased to 250°C, the crystallinity improves remarkably (Xmrn= 0.04), but there is no further broadening of the Ge peak, consistent with the fact that we have retained a low temperature abrupt interface without any further intermixing. Additional elevation of the growth temperature further decreases the channeling yield to 0.03 showing an additional improvement in crystallinity. However, a broadening in the Ge peak is observed for growth temperatures of 35O’C and above. This broadening occurs asymmetrically into the capping Si layer, and this trend is amplified at higher temperatures. The results from MEIS thus strongly corroborate the results obtained from the Raman spectra. The ratio of Ge-Ge to Ge-Si Raman intensities correlates with the degree of intermixing observed by MEIS as does the disappearance of the amorphous scattering background. Additionally, the MEIS results show that growth-induced mixing occurs into the silicon capping layer to a much larger extent than into the underlying Si layer. A more sophisticated analysis using resonant Raman scattering shows that interfacial asperities on an atomic scale persist at all temperatures of growth and are probably a direct result of the stepped surface on which growth occurs. This effect has been confirmed for Si,_,Ge, alloys using Secondary co-workers.flQ) abrupt Si/Si,,Ge,

Ion Mass Spectroscopy

(SIMS) by Gravesteijn

and

Thus, in order to suppress this effect and obtain atomically interfaces,

it is necessary to reduce the growth tempera-

ture even further. This is also shown in Fig. 1. Clearly, there is a trade-off between crystal quality, especially respect to point defect concentration and stacking fault density, lowering growth temperature.

However,

with and

since we are dealing with effects

that depend on enhanced sub-surface diffusivity, growth temperatures need be lowered only for the initial portion of the overlayer. More recently, Cope1 and co-workerst20j showed that the presence of an adlayer such as Sb or As during growth (i.e., at the growth front) had a surfactant permitted the growth of non-islanded

effect and

films with abrupt interfaces

even at

460


higher temperatures. temperatures

However, the use of such adlayers,

of interest leads to significant

Strain also plays a significant

especially

at the

dopant incorporation.t*‘)

role in determining

the morphology

of

Si,,Ge,films. Kuan et al.t**] have shown that morphological perturbations may provide strain relief under certain conditions. Figure 4 shows TEM cross-sectional images of a 18 period 5 nm Si,,,Ge,,J20 nm Si superlattice grown at 580°C on (a) (loo), (I) (111) and (c) (110) oriented Si. It is well known that growth on these latter two substrate orientations gives rise to highly defective films.t’l As a result, superlattices grown on these planes partially relax via the generation of twins or dislocations. However, on Si(lO0) substrates, defect generation is minimal and morphological instabilities develop in the Si,,Ge, layers. The surface planarity recovers completely after the thin Si layer is overgrown. The instability then redevelops during the subsequent Si,,Ge, layers. Furthermore, the thickness modulation of the Si,,Ge, layers displays amazing regularity and increasing amplitude indicating that the film possesses inhomogeneous lateral strain (see Fig. 4(a)). There are indications that as the amplitude of the interfacial roughness increases, dislocations are ultimately nucleated. When dislocations are permitted to nucleate, as in the (111) and (110) cases, the Si/Si,,sGe,,s interfaces within the superlattice are indeed quite smooth. Lowering the growth temperature to 400°C on Si(100) surfaces lowers the surface mobility considerably and both the morphological instabilities and strain relaxation seen in Fig. 4(a) are completely inhibited. The success of MBE in growing a variety of Si/Si,,Ge, device structures, as discussed in Sec. 5, has led to the development of several alternative approaches, mainly involving chemical vapor deposition (CVD). Since the key to obtain good quality

Si,,Ge,

films is low temperature

growth, and the fact that CVD rates are necessarily slow at low temperatures, utmost purity in the CVD system and in the gas delivery system is required. Pioneering work in this regard has been done by Meyersont23) using a technique called Ultra High Vacuum CVD (UHV-CVD) and by Sedgwick and co-workerst24) using a method called Atmospheric Pressure CVD (APCVD); this latter approach eliminated the need for pumps on the growth system. Gibbons and co-workerst25] at Stanford University have developed a technique called Limited Reaction Processing (LRP) which is also referred to as Rapid Thermal CVD (RTCVD). in UHV-CVD,

The source gases used

a batch process, and Gas Source MBE (GSMBE),

wafer process which operates under somewhat similar are usually the hydrides of Si, Ge, and the dopants.

a single

conditions,t26] In the case of

Elemental Semiconductor Heterostructures

461

nm)/ Figure 4. Top.’ Bright field (011) cross-sectional image of a Si,sGe,,(5 Si(20 nm) superlattice grown on a Si(lO0) substrate at 580°C and then annealed at 450°C for 30 min. The Si,,sGec.s la y ers, appearing in dark contrast, exhibit substantial surface roughness. Middle: Bright field (110) cross-sectional image of a Sic.sGec.s(5 nm)/Si(20 nm) superlattice grown on a Si(ll1) substrate at 600°C and then annealed at 45O’C for 30 min. Growth along the direction produces irregular Si c,sGe,,, thickness and a high density of dislocations. Bottom: Bright field (111) cross-sectional image of a Si,.sGe,,,(5 nm)/ Si(20 nm) superlattice grown on a Si(ll0) substrate at 580% and then annealed at 45O’C for 30 min. Lattice mismatch strain in this superlattice is partially relaxed by misfit dislocations and also by twinning.

462

Molecular

LRP/RTCVD

Beam Epitaxy

and APCVD the silicon source is usually a halide.

other methods are described summarize

in detail in the references

some of the main features of these techniques

Table 1. A comparison

of major epitaxial techniques

These and

cited, and we only in Table 1.

for growing Si,_,Ge,

alloys.

SUMMARY OF EPITAXIALGROWTH TECHNIQUES

The distinguishing feature of MBE lies in the fact that it is a physical deposition technique. As such, the deposition rate and substrate temperature are independently controllable. Furthermore, at least for the major constituents and for some dopants like B, the film composition accurately reflects the impinging flux composition, and surface chemistry plays only a limited role. On the other hand, hydride-based chemistries at low growth temperatures

using CVD result in growth

faces.~~t2s] As a consequence, surface composition has thereon,

on hydrogen

passivated

sur-

hydrogen desorption, and the effect plays an important role in determining

the growth rate and film composition.

In general,

CVD reactions

are site

specific and deposition proceeds via the exchange of a precursor with a surface terminating species. Thus, the growth process can also be selective.

While to-date, all CVD techniques have been unable to demonof Ge content, while maintaining good film

strate a wide range control

morphology especially for layers with high Ge content heteroepitaxial layers, more elaborate growth chemistries

and ultra thin and conditions

do not preclude CVD techniques from achieving such results in the future. Such control is essential for some device applications employing thin multilayers, such as quantum well structures. On the other hand, CVD is

Elemental

the process of preference throughput surfaces, deposition

are crucial.

in Si technology The hydrogen

and lack of particle systems,

Heterostructures

where high volume

passivation

contamination

have a beneficial

especially when low Ge-content, bipolar transistors

Semiconductor

and high

in most CVD grown

that is common

impact

463

on epitaxial

in physical film quality

thin SiGe layers, such as in heterojunction

(HBTs), are required.

Nevertheless,

the ease, relative

safety, and elegance of MBE makes it the preferred deposition technique, especially in a research environment, and to-date, most new device concepts in the Si/Si,,Ge, system have been first demonstrated by MBE. Clearly, the availability of a variety of growth techniques has been important for the rapid maturity of Si/Si,,Ge, technology, and the optimal growth technique is specific to the application under consideration. 1 summarizes some key features of the various growth techniques deposit Si,,Ge, epitaxial films,

3.0

STABILITY

OF Si,,Ge,

The thickness

Table used to

FILMS

of the Si,_,Ge,

layer is often an important

device

design consideration. The maximum thickness for pseudomorphic growth, often referred to as the “critical thickness,” of Si,_,Ge, alloys is an important parameter in this lattice-mismatched system. There has been considerable debate on how critical thickness should be defined or experimentally determined. Van der Merwef”) introduced the concept of critical thickness based on equilibrium theory. He defined critical thickness as the film thickness below which it was energetically favorable to contain the misfit by elastic energy stored in the distorted crystal, and above which it was favorable heteroepitaxial critical thickness threading account uniform

to store part of the energy in misfit dislocations at the interface. Matthews and Blakeslee (MB)f13)f14) defined in terms of the mechanical

dislocation.

equilibrium

It is now recognized

for kinetic limitations spacing. Treatments

to dislocation which

account

of a pre-existing

that these theories generation for these

do not

and their limitations,

nonfor

example, by considering irregularly spaced or isolated dislocations,f2gj or by accounting for kinetic barriers to dislocation movement, such as the Pierls barrier,f30) generally give larger values for critical thickness. The inclusion of kinetic barriers to the formation of dislocationsf3’j is another promising approach. While growing Si,_,Ge, alloys on high quality Si substrates, the effect of kinetic barriers to dislocation formation assumes

464

Molecular

Beam Epitaxy

added importance and metastable

layers far in excess

of the “critical

thickness” are possible. This is especially true when the observation onset of strain relaxation is limited by the experimental technique Structural

analysis

methods,

detect a lattice mismatch

of the used.

such as triple crystal x-ray diffraction,

as low as 0.01%.

If strain relaxation

can

occurs

through the formation of dislocations, the deleterious impact on device performance may be seen long before strain relaxation is determined by structural characterization techniques. Electron Beam Induced Currents (EBIC)t3*1 measurements have been used to image dislocations because of the increased carrier recombination in the vicinity of the dislocation. Other electrical evaluation methods to determine the structural quality of Si/Si,,Ge, heterojunctions include the measurement of the reverse leakage current density across a p-n heterojunction, the ideality of the forward injection current, and the measurement of band discontinuities. The electrical activity of the different types of dislocations and their various components is not yet a well characterized parameter and cannot be quantitatively correlated to strain relaxation. Nevertheless, these measurements serve as a very sensitive qualitative measure of the heterojunction quality. Electrical measurements of MBE grown Si/Si,,Ge, p-n junctions indicate that excellent junctions are attainable. Experience suggests that extraneous factors, such as proper substrate preparation and particulate contamination during growth, all play a kinetic role in dislocation formation, and may be more important than the energy and mechanical equilibrium theories would suggest, particularly for the growth of low (5 25%) Ge-content films on near perfect Si(100) substrates. The concept of critical thickness does play an engineering role and provides a useful frame of reference for device design. Figure 5 shows the critical thickness also delineates

(theory and experiment) different

useful expression

as a function

regimes used for various

for critical thickness

of Ge content and

device applications.

based on the MB criterion,

A

is given

by

Eq. (2)

effectivestrain

(11)= 0.046 x ln( 10.4hJ lh,

where h, is the critical thickness. The effective strain as obtained from Eq. (2), is averaged over the thickness of the SiGe layer under consideration, i.e., the layers need not be of uniform Ge concentration. In addition, for low Si,,Ge, layers (x 5 0.15) it is safe to assume that about twice the thickness

obtained from this expression

is stable.[33j

Elemental

0

Semiconductor

1

2

Heterostructures

465

Misfit(%)

3

~OOnm

1Onm

lnm 0

0.2

0.6 0.4 Germanium Fraction

0.8

1.0

Figure 5. Critical thickness plotted as a function of Ge content. Shown is the empirical curve of People and Beant2g] and their experimental points obtained from channeling and TEM; EBIC determined points of Kohama et al.t32]; and, points determined from bandgap measurements of HBTs from King et al.f1271 The more conservative mechanical equilibrium theoryt13] and the thermodynamic equilibrium theorytlrl are also shown for comparison. The range of Ge content spanned by the various device applications of Si,,Ge, alloys is also indicated in the figure.

Misfit dislocations phic Si,,Ge, Si,,Gex,

are not the only means of stabilizing

films. In the case of covalent

the onset

of relaxation

pseudomor-

alloy semiconductors

for single

layers

is gradual

such as and film

thickness much greater than the critical thickness may be required before significant relaxation has occurred, as suggested by Fiery et al.t34] and Tuppen et al.f35] An empirical relation between the amount of relaxation and the thickness of film, as normalized to the critical thickness, is shown in Fig. 6.fsc] After the initial partial relaxation, subsequent relaxation occurs more slowly, with a characteristic relaxation thickness, nh,. In practice n is somewhat dependent on film composition and growth condi-

466

Molecular

Beam Epitaxy

tions but is of the order of 10 for low growth temperatures content.

The use of such partially

lattice parameter and Si,,Gex

modification

Sec.

5).

scheme

layers with different

strain respectively, These

relaxed Si,_,Ge,

has been exploited[37] to grow Si

amounts

which allows for even

principles

electron gases (2DEGs)

and moderate Ge

layers as a substrate

of tensile

greater

band

and compressive engineering

have been used to fabricate and resonant

tunneling

(see

two-dimensional

structures.

80 h z

Empirical relation: Relaxation = 1 - exp(-t/nt,) where n - 10 t, = “critical” thickness

6.

s ._ 75 g 40 ; 20 0

I

0

To accelerate relaxation: Increase Ge concentration Raise buffer growth temperatul I

t

*

20 40 60 80 Normalized thickness (t/Q

100

Figure 6. An empirically determined relationship between strain relaxation (fully relaxed corresponds to 100% strain relaxation) and film thickness (normalized to the appropriate critical thickness for any particular Ge content, see Fig. 5).

After growth, these metastable Si,,Ge, films may relax in the course of subsequent device processing. Once again, kinetic factors relating to the creation, propagation, and annihilation of dislocations govern these relaxation mechanisms. The initial crystalline perfection and interface quality of the as-grown structure play a crucial role. We have foundf38)f3Q]that for superlattices with few dislocations to begin with, strain relaxation,

if at all, proceeds via interdiffusion

(i.e., layer mixing).

On the

Elemental

other hand, a preponderance

Semiconductor

Heterostructures

of defects in the original film stimulates

467

rapid

relaxation via dislocation multiplication, while maintaining the integrity the Si/Si,, Gex heterointerfaces. The stability of Si/Si,_,Ge, structures enhanced

of is

by the following:

1. High quality as-grown

interfaces

2. Layers with graded interfaces

and films

and graded Ge content

3. Layers with Si caps For some device applications, such as HBTs that employ Si,,Ge, films with low Ge content and graded Ge profiles, thermal stability during subsequent processing is not a problem. In fact, Patton et al.t40) report excellent p-n junctions even after a 105O’C 30s anneal with no significant difference between the Si,,Ge, heterojunction and homojunction Si control devices. Strain relaxation of thicker Si,,Ge, layers and those with high Ge content is more complex.t4’] As before, we find that strain relief proceeds rapidly at first but slows down considerably thereafter.f36) For example, Fig. 7 shows the thermal relaxation for a 250 nm thick Si, _xGex layer with x = 0.32 grown at 350°C as a function of annealing temperature for isochronal anneals of 30 minutes. The as-grown layers are essentially unrelaxed Strain relaxation beyond 66% Ge is as determined by x-ray diffraction. fairly sluggish. Both with respect to growth and annealing, relaxation appears to proceed sluggishly after about two-thirds of the strain is relieved. Undoubtedly, the driving force for efficient strain relaxation is Houghton and co-workersf4*] have significantly reduced at this point. addressed the problem of quantifying the allowable thermal budget for pseudomorphic Si, +Ge, layers and obtained analytic expressions for plastic strain relaxation under the driving force of the effective stress for an arbitrary thermal cycle in terms of dislocation glide velocities and heterogeneous nucleation sites. Such models based on experimental data provide valuable help in device processing. Experimentally, the reduction of strain upon partial relaxation still leaves a high degree of residual strain in Si,,Ge,

layers as seen in Figs. 6 and 7 which, interestingly,

value at least for high Ge content films (x r 0.25) grown

is similar in

by MBE.

The

above arguments bear on the techniques used to tailor strain in as-grown Si,_xGe, layers. Thus, for example, we need to compensate for incomplete relaxation by augmenting the Ge content in the SiGe buffer layer above what is desired to achieve a relaxed Si,,Ge, film with the appropriate lattice constant/Ge

content.

488

Molecular

100

Beam Epitaxy

I

I

I

I

250 nm Si,_c,Ge,,a2 / Si grown at 35OOC

80 T e s ‘i; 2 s d

I

604020 0 200

I 300

700 400 500 600 Anneal temperature ( “C)

Figure 7. Strain relaxation as a function of annealing Si,,Ge,,, film grown on Si(lO0) at 350°C.

4.0

LONG RANGE ORDER

IN THE Si,,Ge,

temperature

for a 250 nm

SYSTEM

Atomic ordering has been studied extensively several elemental and compound semiconductor Advances in heteroepitaxial growth techniques,

800

in metallic alloys and

material systems.f43]-[451 such as MBE, have per-

mitted the observation of new stable and metastable atomic arrangements which cannot be predicted from bulk phase diagrams at the same temperature range or composition.[44)-f48] Ordering in elemental semiconductor alloys, such as Si,,Ge,,

has been observed

by many workers,f48)-f53)

and the presence of long range order (LRO) had been initially attributed to strain in lattice-mismatched epitaxial Si,,Ge, layers, until it was shown by LeGoues et al.f53] that ordering exists in unstrained, bulk-like Si,_xGe, films. This experimentally observed ordered phase does not correspond to the predicted lowest energy phase of the Si,_,Ge, alloyf46) which consists

Elemental

Semiconductor

Heterostructures

469

of alternating Si-Ge-Ge-Si layers along the direction. Instead, the observed phase was shown to be a microscopically strained structure consisting

of bi-layers

of Si and Ge along

all four

equivalent

cl 1 l>

directions.t53) Further, LeGoues et al.t54] have shown that such ordering is induced

by the quenching

may be explained

of site-specific

segregation.

This segregation

with the help of Fig. 8 which shows a schematic

cross-

section of a 2 x 1 reconstructed Si(lO0) surface. The 2 x 1 reconstruction on the growth front results from a dimerization of the Si atoms at the surface and are present during growth by MBE. We can see from the cross-section of the Si(100) 2 x 1 reconstructed dimer structure (see Fig. 8(a)) that alternating atomic sites in the third and fourth layers of the film are under compressive or tensile stress. These atomic-scale stresses can cause site-specific segregation of Si and Ge atoms in the Si,_,Ge, alloy since sites under compressive stress would rather be occupied by the smaller silicon atom while those under tensile stress would be favored by the larger germanium

atom.

When a double layer is grown on the initial

surface shown in Fig. 8(a), alternating Si-rich and Ge-rich pairs of atoms in the third and fourth layers of the crystal are seen again (see Fig. 8(b)). At low growth temperatures, it is possible to “freeze in” this structure because bulk diffusion coefficients are sufficiently low, and thus elemental distributions seen in the sub-surface layers are sustained throughout the epitaxial SiGe film. This mechanism results in the experimentally observed Ge-Ge ordered phase (see Fig. 8)t53) and not the Si-Ge-Ge-Si

Si-Siphase

predicted by theory at very low temperatures. Hence, epitaxial growth conditions, such as growth temperature and surface reconstructions at the growth front, directly determine the presence or absence of long range order in Si,_xGe, alloy layers, and long range order in Si,_,Ge, films is a direct consequence

of surface growth kinetics.

Molecular beam epitaxy (MBE) and ultra-high vacuum chemical vapor deposition (UHV-CVD) grown Si,_,Ge, (0.1 s x ~0.8) films deposited on Si substrates with different orientations over a wide range of substrate temperature this study,

and growth conditions surface

reconstructions

have been studied by Kesan et al.t55) In occurring

at the growth

front were

modified in situ using an adlayer such as gallium, antimony, or boron and monitored using low energy electron diffraction (LEED). The presence or absence of long-range order was determined using planar and crosssectional transmission electron microscopy (TEM). In addition, the Si,_, Ge, layers were annealed at different temperatures for varying amounts of time to study the stability of the ordered SiGe phase.

470

Molecular

Beam Epitaxy

Figure 8. (a) Cross section of the (100) 2 x 1 surface, projected onto a (110) plane. Surface dimers are at the top. Large so/id circles correspond to sites under compressive stress, favoring Si occupancy. Large open circles denote sites under tensile stress, favoring Ge. (Dimer sites are also shown as large open circles, despite their small stress, because the surface energy favors Ge occupancy for those sites.) Small circles denote sites with little preference for Si or Ge. (b) Proposed growth process: The third and deeper layers in (a) are assumed immobile, while two more layers are added. Thus the circles in the fifth and sixth layers denote Si or Ge occupancy due only to kinetics (past history), rather than to actual stress or any equilibrium preference.

Si,,Ge,,s

films,

5000 A thick,

were

grown

on Si(lO0)

substrates

by

MBE at growth temperatures between 39O”G590°C. Figure 9(a)-(c) shows diffraction patterns from a planar-view sample taken along the (110) zone axis for three different growth temperatures, (a) 390°C, (b) 490°C and (c) 59OOC. The diffraction pattern in Fig. 9(a) clearly shows the presence of additional superlattice reflections present at l/2(1 1 l}, indicating strong order in the sample. The same extra reflections persist in Fig. 9(b), but are considerably

weakened

in intensity.

In Fig. 9(c) we can see

Elemental

Semiconductor

Heterostructures

471

from the overexposed diffraction that these spots have disappeared completely, indicating absence of long range order in the Si,,Ge, films grown at high temperatures Figure

Si,,Ge,

by MBE.

10 shows the effect of annealing

films as a function of Ge composition.

temperature

Ordered Si,,Ge,

on LRO in (0.1 saO.8)

films were annealed for 2 hours at temperatures between 450°C and 800°C. Ordering in Si,,Ge,,s films persists up to annealing temperatures of 650°C. However, Si,,Ge,films with Ge compositions either less than or greater than 50% are less stable to high temperature annealing, and ordering is destroyed at temperatures around 500%. In addition, once ordering is destroyed, long term annealing for several hours over a wide range of annealing temperatures fails to restore LRO.

(a)

(b)

Figure 9. Diffraction patterns from a planar-view sample taken along the (110) zone axis for three different growth temperatures, (a) 390°C (LJ)490°C, and (c) 590°C. The diffraction pattern in (a) clearly shows the presence of additional superlattice reflections present at l/2(1 1 l}, indicating strong order in the sample. The same superlattice reflections persist in (b), but are considerably weakened in intensity. In (c)we can see from the overexposed diffraction that these spots have disappeared completely, indicating no order.

472


The fact that, once ordering annealing

or high temperature

long term annealing alloy is a metastable

in Si,,Ge,

is annihilated

by either

growth, it cannot be restored by subsequent

clearly indicates that the ordered phase of the Si,,Ge, phase which occurs under certain epitaxial growth in Si,,Ge c.s persists at higher annealing tempera-

conditions. Ordering tures than in Si,,Ge, films with unequal percentages of Si and Ge (see Fig. 10). This is consistent with the fact that fewer atomic displacements are necessary to cause destruction of long range order in Si,,Ge, films with unequal amounts of Si and Ge. Since ordering in Si,,Ge, is a strong function of temperature, it is important to establish the degree of long range order that exists in these epitaxial Si,,Ge, films. Kesan et al!551 and Tsang et al.t56] have used grazing angle x-ray diffraction and resonant Raman scattering to determine the extent of ordering in Si,,5Ge,,5 layers grown at 39O’C and 490°C. In other words, the relative proportion of Si and Ge in their respective bilayers was quantified. For a Si,,,Ge,,s film grown at 390°C the ratio of the intensities of peaks at l/2(777) and l/2(888) in the x-ray diffraction spectra was found to be 0.114. Taking into account the structure factor associated with the change in unit cell parameters due to the strain induced by the presence of Ge and the correcting factor related to the experiment geometry, the ordering parameter was found to be 0.64. Hence, the Si (or Ge) bilayers for a Si,,sGe,,, film are expected to contain about 18% Ge (or Si). The extent of ordering was also determined using resonant Raman scattering by comparing the relative strength of the Si-Si, Ge-Ge, and Si-Ge vibrations and found to be around 70-80%.f56] Si,,sGe,.s films grown at 490°C show 50% weaker ordering by x-ray diffraction compared to films grown at 390°C consistent with the electron diffraction spectra seen in Fig. 9. The extent of long range order in the Si, _xGex films and the range of temperature

for which ordering is strong enough to be detected

cantly higher than that predicted

by Kelires and Tersoff.f5’]

is signifi-

It is difficult to

gauge how accurately stress at the growing surface can be described by the empirical classical potential used in their model. Further, since the Kelires and Tersoff model refers to thermodynamic

equilibrium

only, this

discrepancy is not surprising. The nature of the reconstructed Si surface during growth plays a determining role in the ordering of Si,,Ge, films. Surface reconstructions on a Si(100) substrate can be modified in situ during growth using antimony, gallium, and boron, and thus the effect of surface reconstructions on LRO can be examined. Initially, a 5000 A thick Si,,,Ge,,, film was

Elemental

Semiconductor

grown on a normal 2 x 1 reconstructed temperatures

Si(100)

to produce an ordered Si,,,Ge,,,

surface was saturated

by 0.25-0.50

Heterostructures

substrate

at low growth

layer. Subsequently,

monolayer

473

(ML) of antimony

the Si which

causes the reconstruction on the Si(lO0) surface to change from a 2 x 1 to a 1 x 1 structure.t5s) The change in surface reconstruction was confirmed by in situ LEED of the quenched growth surface. Sic.sGee.s growth was initiated again on the 1 x 1 reconstructed surface while continuing to maintain antimony coverage to ensure a 1 x 1 LEED pattern at the growth front. Figure 11 shows the cross section of this sample with a 5000 A thick SiGe bottom layer followed by another 5000 8, of SiGe grown on a 1 x 1 reconstructed surface. Also shown in Fig. 11 are the corresponding electron diffraction patterns, together with the in situ LEED patterns to elucidate the two different growth conditions. The electron diffraction patterns in Fig. 11 clearly show that the bottom Si,,sGe,,s layer is strongly ordered, while the top Si,,,Ge,,s layer is not.

750

I

I

I

ci

6

2

g

I

I

I

0 /-0

0 WEAK ORDER -.

Q 0

\

0 \ ‘\@

I

350 0.1

I I I I I I 0.7 0.3 0.5 GERMANIUM MOLE FRACTION

0.9

Figure 10. Effect of annealing temperature on LRO in Si,_,Ge, films as a function of Ge composition. Ordering in Si,,,Ge,,, films persists up to annealing temperatures of 650°C. However, Si,,Ge, films with Ge compositions either less than or greater than 50% are less stable to high temperature annealing, and ordering is destroyed (absence of both SiGe-GeSi type of order, referred to in the literature as rhombohedral structure 1, RHl,f4e)f5*1 and the SiSi-GeGe type of order referred to in the literature as rhombohedral structure 2, RH2f48)f52)) at temperatures around 5o0°c.

474


Elemental

Semiconductor

Heterostructures

475

Saturation of the growth surface with gallium instead of antimony gives similar, if somewhat less striking, results. Figure 12 shows a crosssectional view of a 1 .Opm Si,.,Ge,, under

conditions

growing

surface.

to sustain

film where the top 5000 A was grown

a 0.25-0.50

monolayer

of gallium

We can see from the electron diffraction

at the

pattern in Fig.

12 that the bottom 5000 A is strongly ordered, and the LEED pattern taken during the growth of this layer shows a distinct 2 x 1 structure at the growth front. The surface was then dosed with gallium which caused the surface reconstruction to change to a weak 2 x 1 or a disordered 1 x 1 structure. Any additional gallium coverage immediately caused gallium precipitation on the SiGe surface, and it is thus difficult to maintain a good 1 x 1 reconstruction during growth. The next 5000 8, of Si,,,Ge,,s was grown by maintaining the resulting 1 x 1 structure at the growth front. The electron diffraction pattern corresponding to this layer (this pattern has been overexposed in order to determine the presence of additional superlattice reflections) in Fig. 12 shows very weak ordering in the Si,,,Ge,,, film. The effect of doping with boron is drastically different since the growing surface retains a 2 x 1 reconstruction (i.e., there is no adlayer) in the presence of a boron flux. Figure 13 shows a cross-sectional view and diffraction pattern of a boron-doped Si,,,Ge,,, film. The structure in Fig. resonant tunneling structure (see Sec. 13 corresponds to a Si/Si,,,Ge,,, 5.4), but in these discussions only the 1 .O pm thick Si,,sGe,,, layer is relevant. The electron diffraction pattern in Fig. 13 clearly shows that the boron-doped Si,,Ge,., film is strongly ordered. It is clear from the adlayer surface modification experiments that simply changing the surface reconstruction from a 2 x 1 to a 1 x 1 in situ during

growth

demonstrates

alters the phase that

ordering

of the SiGe

is entirely

alloy

and unambiguously

due to surface

kinetics.

It is

important to rule out any extensive incorporation from the adlayer species into the SiGe film which might result in a change in the bulk properties of the alloy.

Figure 14 shows a SIMS profile of the sample described

11 and the extent of bulk antimony layer from the antimony

incorporation

in Fig.

into the top Si,,,Ge,s

adlayer at the growth front is around 2 x 1020/cm3.

This amount of antimony incorporation (c 0.5%) into the Si,,sGe,,, is sufficiently low as not to have any effect on the bulk phase stability of this layer. A weak 2 x 1 LEED pattern is observed during the growth of films saturated with a gallium adlayer and this results in a Sio.5Geo.5 weakly ordered Si,,,Ge,,s

film.

The use of gallium

adlayers

results in an

476


Elemental

Semiconductor

Heterostructures

477

Figure 13. A cross-sectional view and diffraction pattern of a boron doped Si,,,Ge,,, film. This structure corresponds to a Si/Si,,Ge,, resonant tunneling structure, but in our discussions only the 1 .O pm thick Si,.5Ge,.5 layer is relevant. The electron diffraction pattern clearly shows that the boron doped Si,.,Ge,., film is strongly ordered. intermediate

set of results between antimony

and boron which show that

dynamic changes in the 2 x 1 reconstruction directly impact the extent of ordering in the SiGe layer. This reinforces the important role played by the atomic arrangement at the growing surface on ordering in the bulk Si,, Ge,.

Figure 15 shows a SIMS profile of gallium in the Si,,Ge,,

sample

described in Fig. 12, and the amount of gallium incorporated in the film is again sufficiently low (< 0.15%) as not to cause any bulk change in the Sio.5Geo.5.

Boron, which

is a p-type dopant in Si and Si,,Gex,

incorporates

readily into the bulk and does not result in a change in the surface reconstruction at the growth front. Hence, a 2 x 1 structure can be maintained at the Si(lO0) surface while exposing the Si,.,Ge,~, surface to a boron flux, and this produces an highly ordered Si,,,Ge,, film.

478

Molecular

Beam Epitaxy

Figure 14. SIMS profile of antimony in the Sio.5Geo,s sample described in Fig. 11. The extent of bulk antimony incorporation into the top Si,sGe,, layer from the antimony adlayer at the growth front is around 2 x 1020/cm3 (< 0.5%).

DEPTH

(pm)

Figure 15. SIMS profile of gallium in the Sio.5Geo.5 sample described in Fig. 12. The amount of gallium incorporated in the film is again sufficiently low (c 0.15%) as not to cause any bulk change in the Sio.5Geo.s.

0.0

0.2 DEPTH

0.4 (/t’m)

0.6

Elemental

Semiconductor

Heterostructures

479

Si,_xGe, films grown at low temperatures (450%) by UHV-CVD are Due to intrinsic not ordered, but films grown at 560°C are ordered. differences temperatures distinctly

in MBE and CVD growth

techniques,

which produce ordered Si,,Ge,

the range

of growth

films in the two cases are

different and actually the inverse of each other.

These Si,,Ge,

films grown by UHV-CVD permits one to examine the effect of surface reconstruction on LRO without the use of adlayers. This is because in the UHV-CVD environment at low growth temperatures, there is significant hydrogen coverage and the Si(lO0) surface is 1 x 1 reconstructed.t27)f2s) At higher growth temperatures, around 550°C, the hydrogen coverage is less complete and a 2 x 1 reconstructed Si surface is seen. Indeed, in keeping with the previous observations in the adlayer experiments, it is found that UHV-CVD SiGe films grown at higher temperatures are ordered but those grown at lower temperatures are not. While this is seemingly contradictory with the MBE grown SiGe results, a closer examination of the conditions which produce LRO in both MBE and UHV-CVD films indicates that both conditions relating to low temperature growth and a 2 x 1 surface reconstruction must be maintained to observe LRO. Si,,Ge, films grown by MBE on Si(ll1) substrates do not show any ordering under the set of growth conditions described earlier. In particular, the same growth temperatures that produce strong ordering in Si,,Ge, films grown on Si(100) substrates do not result in LRO when grown on Si(ll1) substrates. Growth on Si(ll1) substrates allows one to examine Si,,Ge, films grown on a 7 x 7 reconstructed Si surface. These Si,,Ge, films show no ordering under any set of growth or annealing conditions. This once again indicates that a 2 x 1 reconstructed surface is necessary to see LRO. In summary,

ordering

in Si,,Ge,

is an entirely

kinetic phenomenon

governed completely by growth conditions and surface reconstructions and not by bulk thermodynamic equilibrium, as was previously believed. The experimentally-observed ordered phase in Si, _xGex films, which consists of bi-layers of Si and Ge along all four directions, is a metastable phase that does not correspond to the lowest energy phase of the alloy and is irreversibly destroyed by annealing. A set of experiments involving the use of annealing, surface modification through adlayers, UHV-CVD growth, and growth on Si(ll1) substrates all indicate that both low temperature growth and a 2 x 1 reconstructed surface are individually necessary, but not sufficient conditions, to observe LRO in Si,,Ge,. Ordering in Si,,Ge, occurs independent of strain, and indeed Si/Si,,Ge,

480

Molecular

Beam Epitaxy

strained layer superlattices grown on Si substrates

with different

tions show no order.tsQ) We have also discussed suggests that ordering in Si,,Ge, stresses at the growing surface. experimental

5.0

findings

a mechanism

is related to local segregation This mechanism

successfully

orientawhich

induced by explains

all

of ordering in Si,_,Ge,.

DEVICE APPLICATIONS

OF Si,,Ge,

ALLOYS

Si,,Ge, alloys show smaller fundamental bandgaps compared to Si principally because of a larger lattice constant and altered lattice constituents, and in pseudomorphic layers, due to the acquired tetragonal symmetry. Strain in the Si,,Ge, alloy layer results in a further modification of the band structure; these changes include the splitting of degenerate valence and conduction bands minima,f60) the extent of which is determined by the strain in both the Si and Si,,Ge, layers. The bandgap energy as a function of Ge content is shown in Fig. 16. The top solid curve shows the bandgap for a bulk alloy layer due to Braunstein et al.tlo) Note that until the alloy composition reaches about 85 atomic percent Ge, the characteristics retain a silicon-like band structure with the smallest gap occurring near the X Brillouin zone (BZ) boundary. At greater concentrations, the smallest indirect gap acquires a germanium-like character and occurs near the L Brillouin zone boundary. The growth of pseudomorphic Si,,Ge, on Si substrates, causes an even greater shrinkage of the indirect gap as shown in the lower two curves of Fig. 16. The strain-split heavy hole band (labelled upper band) determines the smallest gap and has been calculated determined

by optical absorption

by People and Beant61)f62) and

experiments

by Lang and co-workers.f63]

Here it is important to note that the lowest-lying hole band in a compressively strained Si,,Ge, film is strongly anisotropic and actually has a light effective mass in the plane of the Si,,Ge,fiIm.f64) The bandgap shrinkage is significantly

increased

by the presence of tetragonal

strain in the layers.

In fact, a pseudomorphic Si o,4Geo,6 layer on silicon shows a gap that iS smaller than pure bulk germanium. The band alignment for strained Si,_, Ge, on Si assumes a type I nature depicted in Fig. 17(b).f6’l The entire bandgap of the alloy is contained within the Si gap with the conduction band offset remaining below 0.02 eV and decreasing to almost 0 at higher Ge contents. Hence, most of the bandgap difference with increasing content manifests itself principally in the valence band.

Ge

Elemental

Semiconductor

Heterostructures

481

Band I

0

I

0.2

I

I

I

I

I

0.4 0.6 Germanium Fraction

I

1

0.8

1.0

Figure 16. Bandgaps for unstrained bulk Si,,Ge, alloys[lO] and pseudomorphic Si,,Ge, alloys. Both the strain-split light and heavy hole band@] are shown along with the experimental

optical absorption

data.[63]

The effect of strain on the bandstructure of pseudomorphic structures provides for a wide range of control of the energy gap depending on the substrate chosen. The gap energy of Si,_,Ge, layers as a function of Ge content, x, is summarized in Fig. 17(a) for substrates with varying Ge content and is based on theoretical calculations.[“1[651-[68] Clearly, a greater degree of bandgap adjustment is possible for growth on relaxed Si,,Ge, substrates with different lattice constants. The band offsets themselves are also sensitive to the relative strains in the constituents of the heterostructure,

and strained Si grown on top of unstrained

acquires a type ii alignment

Si,,,Ge,,

as shown in Figs. 17(c) and (d). Note that it is

possible to obtain both large valence band and conduction band offsets by suitably engineering the strain in the Si/Si,,Ge, heterostructure system. The ability to engineer the bandgap

of an elemental

semiconductor

has already led to several new applications; these include heterojunction SiGe MOSFETS,[~~)-~~~ Modulation Doped Field Effect Transistors (MODFETS),~~)-~~) resonant tunneling diodes,f7gj-f83) optoelectronic devices such as pseudomorphic strained layer multiple quantum well p-i-n photodetector.s,f84)-f87) avalanche photodetectors,f88) long wavelength (1.55 pm) photoconductive detectors,f8g) and heterojunction bipolar transistors (HBTs) ~[4cjf9cj+Js)

482


1.2

I

I

I

I

I

1.1 T

-

I

I

I

1

bulk alloy

- - .

Si

-.

Sir&ec,n

.-

Sig,~Cec,50 substrate

substrate substrote

Ce substrote

1.0

.

; 0.9 =r 0.8 k m 0.7 --. . c

0.6-

I I I I I I I I 0.5* I 0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1.0 Ge atomic fraction

0a Figure 17. (a) Bandgap of Si,,Ge, alloys grown pseudomorphically on different substrates as indicated. The band alignments for Si/Si,,sGec,, are also shown for growth on different substrates: (b) on Si, (c) on Si,,,Ge 0,25,r (d) on Sio,,,Geo,,,, (e) on Ge. The substrates have their bulk (i.e., relaxed) lattice constants but the layers are pseudomorphic. (Based on the calculations of Refs. 62 and 65-68.)

Elemental

Figure 17.

(Conr’d)

Semiconductor

Heterostructures

483

484

Molecular

Beam Epitaxy

5.1

Heterojunction

Bipolar Transistors

(HBTs)

Si,,Ge, alloys, by virtue of their narrower bandgap with respect to Si, have found extensive uses in a variety of device applications. One of the most important

applications

from a technological

standpoint

has been

the narrow bandgap base heterojunction bipolar transistor. The proof of concept using Si,,Ge, and both DCfg4j[g5jtg7j and ACtg8j performance advantages were demonstrated using MBE grown Si,,Ge, base layers. Both LRP, UHV-CVD and APCVD have been used to fabricate SiGe base HBTs. An important development has been the process integration of a Si,,Ge, base with a polycrystalline Si emitter process by Patton et al.t40j These and other developments have been extensively reviewed elsewhere.tgoj The high performance HBTs being actively researched today combine both drift transistortggj and HBT action. The drift field is obtained by grading the Ge concentration in the base from a high value at the collectorbase junction to a lower value at the base-emitter junction. Devices with unity gain cut-off frequencies as high as 75-80 GHz at room temperature have been reportedf40jtg8jt100j using a variety of growth techniques such as MBE, UHV-CVD, and APCVD to deposit the Si,_,Ge, base. Recently, HBTs fabricated using MBE grown Si/Si,_,Ge, layers have exhibited cutoff frequencies of 100 GHz.[‘~~] Si,,Ge, base HBTs are one of the most promising device applications for SiGe for the following reasons: impressive performance advantage can be realized using relatively low concentrations of Ge-for example, the high frdevices referred to above have peak Ge concentrations below 8 atomic percent. Furthermore, the Ge content in the base layer is graded in a trapezoidal profile, and the base widths employed in advanced bipolar devices is quite low, below 100 nm. These features, the low Ge content, the fact that it is graded, layer, make it a more processing is done with metallurgical junctions successful incorporation

and the low thickness of the pseudomorphic base stable structure, especially since most thermal a Si cap over the Si,,Ge, layer. Also, the graded impart further stability to the structure. Clearly, of Si,_,Ge,into

advanced digital bipolartechnology

rests on circuit leverage, manufacturing compatibility, and reliability. More recently, there has been considerable interest in using advantages gained by having a narrow bandgap Si,,Ge, base layer analog circuit applicati0ns.t loll The product of the Early voltage and common emitter current gain, 13, is an important figure of merit analog applications. The Early voltage depends exponentially on difference between the bandgap at the collector side of the base and

the for the for the the

Elemental

Semiconductor

Heterostructures

485

By designing the Si,,Gex layer to largest bandgap in the base. maximize the product of the Early voltage, V,, and common emitter current gain, f3, Si,,Ge, HBTs with pV, products of over 100,000 V have been demonstrated.flO1) 5.2

Heterostructure

FETs

The bandgap difference between Si and strained Si,,Gex manifests itself predominantly in the valence band. Hence, p-channel heterostructure Si,,Ge, FETs have been investigated in some detail. Ge has the highest hole mobility amongst all common semiconductors, including Ill-V to MOSFET compound semiconductors.t l”) The application of Si,,Ge, devices, especially for p-channel devices, holds promise because of the predicted high mobility of holes in Si,,Ge, channels.f103)-f10q Performance enhancement of the low mobility p-channel MOSFET in Si-based CMOS circuits can lead to better matching of p-channel and n-channel MOSFET characteristics, with consequent performance and density enhancement in CMOS integrated circuits. The use of Si,,Gex alloys for p-channel

high transconductance

MOSFETs requires a high quality dielectric system. Direct oxidation of Si,,Ge, alloys or even low temperature deposition of SiO, directly on Si,, Ge, results in a very high interface state density. lyer et al.f6Q)have shown that the use of a thin (6-7 nm) Si cap layer grown epitaxially on the Si,_xGe, channel layer followed by either low temperature oxidation of the Si cap or plasma enhanced chemical vapor deposition (PECVD) of silicon dioxide gives a low (below 10”) interface state density. The Si cap layer leads to a sequential turn-on of the Si,,Ge, channel and the Si cap channel and is clearly seen at low temperatures There are several design trade-offs needed to suppress the parasitic Si cap channel, namely the thickness and Ge content in the SiGe channel and the thickness of the Si cap. Briefly, abrupt Si/SiGe interfaces and the thinnest Si cap layers consistent with good oxide/semiconductor interface properties are required. The Si cap thickness can be increased for low temperature operation when carrier confinement in the SiGe channel is readily obtained. Modulation-doped Si/Si, _xGe, two-dimensional hole gases (2DHGs) p2)p6)p7)f106)-f10*)have reported hole mobilities as high as 5000 cm*/ V-s at 4 K. Si,,Gex quantum well long channel p-MOSFETs have shown about 50% higher mobi~ity[6Ql-[711[731-(7~~ over comparable Si p-MOSFET structures. A peak hole mobility of 9000 cm* / V-s at 77 K has been reported in a modulation-doped Ge p-channel device.p*)

486

Molecular

In a Si,,Ge,

Beam Epitaxy

strained

quantum

well, the heavy hole and light hole

band degeneracy is lifted, and due to the combined effects of strain and quantum confinement, the lowest lying sub-band is expected to have a low effective mass for in-plane transport, MOSFET.t641t10gl However, in a filled sub-band

as in a MODFET or due to band-mixing and

large non-parabolicities in the valence band, the experimentally measured in-plane effective mass is not as low as might be expected. Cheng et al. have determined the cyclotron resonance (CR) mass of the 2DHG in a strained 7.5 nm Si,,,,Ge o,37quantum well to be (0.29 + O.O2)m, for a 2D hole density of 2.3 x 1012/cm2 at 3 K. This reduced in-plane effective mass is clearly a consequence of strain-induced valence band splitting and non-parabolicity effects in strained Si,,Ge,. This CR mass of 2DHGs in strained Si,,Ge, is comparable to measurements of the CR mass of 2DHGs in strained InyGa,_yAs with similar 2D hole densities. The successful integration of Si,_,Ge, channels into a CMOS-based process technology has been demonstrated by Kesan et al.r4] Quarter micron Si,,Ge, p-MOSFETs (no modulation doping) with either thermal or PECVD oxides have been fabricated using a CMOS compatible process sequence (see Fig. 18). These devices were fabricated using standard LOCOS oxide isolation compatible with conventional CMOS processing. Identical phosphorus threshold-adjust and deep-well implants and anneal were carried out for both the Si and SiGe p-MOSFET structures. Hence, no modulation doping was required for threshold-voltage control. In these experiments the undoped Si,_,Ge, (05x ~~0.25125-300 A) channel and Si (70 A/105 A) cap were then deposited by selective UHV-CVD at 530%. TEM analysis indicates that the Si/SiGe overgrows the field oxide epitaxially, without faceting, thus avoiding leakage at the point where the gate polysilicon crosses over onto the field oxide. Gate oxides (nominally 70 A) were either thermally grown or deposited using a high-quality PECVD process.

For devices with thermal

oxides, the silicon cap thickness

was

thicker (105 A) to account for silicon consumption during oxidation. SIMS analysis shows that the Ge profile in the channel was not degraded by thermal oxidation at 7OO”C, as indicated by the identical Ge profiles in both the PECVD and thermal oxide structures. In situ boron-doped p+ polysilicon was then deposited. After gate definition, source-drain junctions were formed using antimony pre-amorphization and boron implant, with either a furnace or rapid thermal anneal. The junction depth is estimated to be around 1200-1500 A. Self-aligned titanium silicide was formed after an oxide/nitride sidewall, followed by conventional Ti/AI metallization. The devices fabricated were non-LDD devices compatible with low powersupply (2.5 V) operation.

Elemental

Field Implamts)

Semiconductor

I

Heterostructures

487

Threshold-adjusl Implant

’ \

I

\

I

\ ,‘Si

I

Figure 18. MOSFET.

(a) Schematic

cap

SiGe channel /

\

and

(b)

\ \

TEM cross-sections

of a Si/Si,,Ge,

p-

488

Molecular

Beam Epitaxy

The output characteristics

for Si and Si,,,Ge,,,

p-MOSFETs

indicate

that the current drive capability is enhanced both at 300 K and 82 K by the use of a Sr,,sGe,,s channel. A short channel (L,, = 0.25 pm; L, = 0.4 pm) saturated

transconductance

at 300 K of 167 mS/mm

is obtained

for a

Si,,Ge,,s channel compared to 139 mS/mm for a Si control device (t,, = 78.3 A for the Si devices; t,, = 71.4 8, for the SiGe devices; V,, = Vds = -2.5 V for all devices). The saturated transconductance, gm,sat at 82 K for the Si,,sGe,,, device increases to 201 mS/mm, compared to 160 mS/mm for the Si control. These are the highest transconductances reported for pMOS devices with this channel length and oxide thickness. The external series resistance for both the Si and SiGe p-FETs was less than 1200 ohmpm. The linear region field-effect mobilities [peti = (dlJdVc)(L / W)/VdsC,,J at 300 K and 82 K for a long channel (Leti = 1.85 pm) Si and Si,.,Ge,,, pMOSFET increases from 95 cm*/V-s (Si) to 150 cm*/V-s (Si,,sGe,,2) at 300 K (50% increase), and from 250 cm*/V-s (Si) to 400 cm*/V-s (Si,,,Ge,,2) at 82 K (60% increase). The dependence of the transconductance in saturation for the Si and Si,,,Ge,,2 p- MOSFETS as a function of effective channel length at 300 K and 82 K, normalizing differences in oxide thickness, is shown in Figs. 19(a) and (b). The transconductance improvement decreases as the channel length is reduced due to the onset of carrier velocity saturation. In these experiments, the Ge content in the SiGe channel was kept below 25% to maintain the integrity of the SiGe layer during subsequent device processing. It is clear from Fig. 19 that substantial improvement in p-MOSFET performance can only be obtained with thin high Ge-content (z=50%) channels. The impact of using such high Ge content SiGe layers in VLSI technology will undoubtedly raise serious reliability issues. N-channel heterostructure FETs~~) can be fabricated using the large conduction band offset that exists between Si,,Ge, and strained Si layers,f67] i.e., the Si channel is grown on a relaxed Si,_,Ge, layer (see Fig. 17 (c) and (d)). Very high electron mobilities of around 200,000 cm*/V-s at 4 K have been reported by several groups1 1101-f112)using a MBE-grown modulation doped structure consisting of a strained Si two-dimensional electron gas (2DEG) grown on a SiGe buffer layer. Quarter micron, ntype, UHV-CVD grown, Si/Si,,Ge, modulation-doped FETs have been fabricated with a transconductance of 600 mS/mm at 77 K.f113) The high mobilities obtained in strained Si 2DEGs has spurted tremendous interest in making magneto-transport measurements, similar to the work done in GaAs/AIGaAs 2DEGs, and the quantum Hall effect and the fractional

Elemental

0

I

0.5 1

I

Effective

I

I

Semiconductor

1I.o

1.5 1 Length

Channel

I

8

Heterostructures

I

I

2.0

Il.0

489

c’

(urn)

,

I

I

82K V,,=V,,=2.5V

Effective

Channel

Length

(pm)

Figure 19. Transconductance in saturation (Vds = V,, = -2.5 v) as a function of effective channel length for Si and Si,,,Ge 0,2 p-MOSFETs at (a) 300 K and (b) 82 K. Also shown in the figure is the transconductance improvement plotted as a ratio between the Si,,,Ge,,, and Si devices. The values for transconductance are normalized to account for differences in gate oxide thickness between the Si and Sic,sGec,, devices. Note that the data shown in this figure corresponds to devices with different channel lengths fabricated using a nominal 0.25 pm technology.

490


quantum Hall effect have been observed in Si/SiGe heterostructures.[1081[1 141 The future of Si/Si,,Ge,

n-type

generate high quality Si,,Ge,

heterostructure

FETs rests on the ability to

buffers with low dislocation/defect

densities.

There has been some progress in this regard, and through the use of SiGe superlattices to control relaxation, it is now possible to grow relaxed Si,, Ge, buffer layers that are essentially resolution (i.e., < 105/cm2).

defect-free

(see Fig. 20) within TEM

Figure 20. Cross-sectional TEM image of a relaxed Si,,Ge,, buffer layer grown on a Si/Si,,Ge, superlattice. The Ge content in the Si,,Ge, layers in the superlattice region is progressively increased in several steps from 0 to 60% Ge. A thick, relaxed Si,,Ge,, film is then grown on top on the superlattice, and this layer is essentially dislocation free within TEM resolution.

Elemental

Semiconductor

Heterostructures

491

One key distinguishing feature about the work on SiGe heterostructure FETs, compared to the SiGe HBTs described in the previous section, is that SiGe FET structures lattice-mismatched

use significantly

higher concentrations

material system like Si/Si,,Ge,,

of Ge.

In a

the use of high ger-

manium-content SiGe layers will continue to raise concerns regarding the uniformity, yield, and reliability of devices fabricated in these epitaxial films. 5.3

Optoelectronic

Devices

The indirect nature of the bandgap of Si and Si,,Gex alloys precludes the extensive optoelectronic applications seen in compound semiconductor

heterostructures.

Nevertheless,

there have been several inter-

esting studies of the optical properties of Si,,Ge, quantum wells and strained layer superlattices and their optoelectronic applications. The concept of zone-folding to generate a quasi-direct bandgap in ultra-thin superlattices has been known since the early seventies. Theoretical studies of small period (100) Sir,jGe, superlattices suggest the possibility of a quasi-direct optical transition across the fundamental gap of the superlattice which is enhanced to within two orders of magnitude of that associated with a direct bandgap material.f115] Experimental verification of this conceptf 1161has been inconclusive since dislocation-related luminescence often occurs in the same wavelength regime, and definitive results suggesting the existence of a direct bandgap in Si/Ge superlattices continue to be sought after. Well-resolved band-edge exitonic photoluminescence from Si, _xGe, quantum wells grown by MBE and CVD has been observed by several workers.[1081[1171[1181In addition, near infrared electroluminescence has been observed from p-i-n structures where the i-region consists of Si/Si,_xGe, multiple quantum wel1s.f ‘lQ) These results and reports of efficient photoluminescence from porous silicon, silicon nanocrystals, etc.f120)f121) continue to drive research efforts to obtain an efficient light source in silicon. Si/SiGe waveguides pm) and long wavelength

and photodetectors (1.3 pm) applications

for short wavelength

(0.8

have been investigated.

The use of silicon-germanium heterostructures permits the realization of Si-based optoelectronic detectors in the 1.3 pm long wavelength regime, without the use of Ill-V technology. Silicon-on-Insulator (SOI) structures are useful for Si-based integrated optoelectronics since the buried oxide layer forms a low-index confinement region that permits effective waveguiding in the silicon overlayer. The use of SOI thus permits the integration of active optoelectronic devices with passive waveguide

492

Molecular

Beam Epitaxy

elements. Si/Si,,Ge, strained layer multiple quantum well p-i-n photodetectors,fs4)-fa7) avalanche photodetectors,f88) and long wavelength (1.55 urn) photoconductive

detectorsfaQ) have been demonstrated.

low loss (0.5 dB/cm) SiGe waveguides can then be coupled to photodetectors

In addition,

have been fabricatedf122]f123) which or Mach-Zehnder

interferometers.

Figure 21 shows a cross-section of a multiple quantum well Si/SiGe integrated rib waveguide-photodetector for long wavelength applications.fe5)fe6) Low loss (l-2 dB/cm at 1.3 pm) silicon waveguides on SOI have been used to achieve remote coupling of the optical fiber to the photodetector through the silicon waveguide. This integrated waveguidephotodetector exhibited low reverse leakage currents (1 O-30 PA/pm2 at 15 V reverse bias) with a frequency response bandwidth of l-2 GHz at 1.3 pm (see Fig. 22). The work of Si-based photodetectors clearly suggests that a complete Si-based integrated receiver will be forthcoming in the near future.

MULTiPiE

QUANTUM

WELL ABSORBING

Figure 21. Schematic cross-section of a Si/Si,_,Ge, multiple quantum well p-i-n photodetector monolithically integrated with a Si waveguide on a silicon-oninsulator (SOI) substrate. The Si waveguide evanescently couples into the active region of the waveguide photodetector.

Elemental

Semiconductor

Heterostructures

493

IO

0

FWHM *251.85

PS

6 4

(a)

2 0 0

I

2 TIME

3

4

(ns)

04

I.0 FREQUENCY

(GHz)

Figure 22. (a) Device impulse response when illuminated by 100 ps pulses from a 1.3pm Nd:YAG laser at 10 V reverse bias by coupling evanescently through the Si waveguide. The full-widthhalf-maximum of the impulse response is 250 ps. (b) Response as a function of frequency for the Si/Si,,Gex p-i-n photodetector. The frequency corresponding to a 3 dB response roll-off is 1 X-2.0 GHz.

In addition to photodetectors at near infrared wavelengths, farinfrared (8-12 pm) photodetectors have been fabricated using Si/Si,,Ge, technology. Lin et al. have demonstratedf 1241a Si/SiGe infrared photodetector using the concept of internal photoemission across the Si/SiGe heterointerface. Infrared photodetectors and multiple element detector arrays, based on intersubband transitions in quantum wells both in the valence and conduction bands, have been demonstrated by several workers.t125)f126) The ability to fabricate far-infrared detector arrays on large area Si substrates is a significant low-cost alternative to HgCdTe or Ill-V far-infrared technologies, and this application is perhaps the most promising application of Si/Si,,Ge, technology.

494

Molecular

Beam Epitaxy

5.4

Other Quantum

Well Structures

In addition to quantum well heterostructure

FETs and optoelectronic

multiple quantum well devices, there have been other quantum well structures fabricated using Si,,Ge, alloys. Resonant tunneling diodes (RTDs), which represent the simplest quantum well structure, have been demonstrated by several groups ~‘Ql-ts*)in this material system. The growth of a double barrier hole resonant tunneling diode (RTD) illustrates the flexibility and versatility of the MBE growth technique.ts*) Figure 23 shows a schematic band diagram of two hole RTDs representing (a) a pseudomorphic structure involving tunneling through a strained Si,,Ge, quantum well, and (6) a structure with a unstrained Si,,Ge, quantum well and strained Si barriers grown on a Si,,Ge, buffer layer. For the structure shown in Fig. 23(b), a relaxed buffer layer consisting of a 7000 A thick p+ Si,,Ge,,, layer was first grown on a Si(lO0) substrate at 425’C. Note that a higher Ge concentration was used in the buffer layer to compensate for the fact that Si,_,Ge, layers with high Ge content do not completely relax even at thicknesses many times beyond the critical layer thickness, as described in Sec. 3. An undoped tunneling structure, consisting of a 50 A Si,,sGe,,s quantum well surrounded by 35 A Si barriers and Si,~,Ge,,, spacer layers of varying thickness from 90 A to 360 A, was then grown at 390°C. This structure was capped with a p+ Si,,,Ge,,, contact layer. We can see from Fig. 23(b) that there are two light hole states and three heavy hole states in the Si,,sGe,,, q uantum well. A cross-sectional TEM image of the pseudomorphic RTD structure in Fig. 23(a) is seen in Fig. 24. The graded Ge content in the Si,_,Ge, spacer layers is clearly seen in the TEM image. The low temperature (77 K) I-V characteristics displayed two distinct resonances corresponding to tunneling through the heavy hole and light hole states in the unstrained Si,,sGe,,s quantum well (see Fig. 25). While tunneling through the light hole state yielded a peak-to-valley current ratio of about 2:l at 77 K, no resonances could be seen at room temperatures. Note that the heavy hole resonance is considerably suppressed in Fig. 25 due to the large heavy hole effective mass. The 3.5 nm Si barriers hence act as efficient, effective mass filters in this structure. Similar tunneling characteristics can be seen for the RTD structure shown in Fig. 23(a). Gennser et al.tQ3)have used these Si/Si,,Ge, hole RTD structures to study valence band anisotropy in strained and unstrained Si,_,Ge, through magnetotunneling. Using angle-resolved magnetotunneling spectroscopy,

Elemental Semiconductor Heterostructures

495

it is possible to map out the dispersion relations and constant energy surfaces of the heavy hole and light hole subbands in Si,,Ge.Je31 Hole resonant tunneling cated valence

represents

band structure

an interesting of Si,,Ge,

approach to study the compli-

alloys.

I

I

I

I

I

0

15

20

28

37

I 52

distance (nm)

0

10

18.5

32

distance (nm)

Figure 23. Schematic band diagram for two Si/Si,,Ge, resonant tunneling structures. (a) A pseudomorphic structure with a strained Si,,Ge, quantum well and graded Si,,Ge, spacer layers. (b) A structure with a relaxed Si,.xGe, quantum well grown on a Si,,Ge, buffer layer with strained Si barriers.

496

Molecular

Beam

Epitaxy

Figure 24. Cross-sectional TEM image of the pseudomorphic SI/Si,.,,Ge,.,, hole resonant tunneling structure shown in Rg. 23(a). It can be clearly seen that the Ge content in the Si,.,Ge, spacer layers is graded, and the quantum well heterointerfaces look abrupt within TEM resolution.

5.1(Y2 I

I

I

I

I

I

4.10-' - Device area 3.10m5 cm2 3.1o-2 2.1o-2 z z 2

1.1o-2 ,-*

0.0 -1.10-2

-

-2.1o-2 -

Spacer thickness 15 nm

-3.10-z. -4.1o-2 I -5.1o-2 -0.8 -0.6 -0.4 -0.2 0 0.2 Voltage (V)

I

0.4 0.6 0.8

Figure 25. A represenrarlve I-V cnaracterlstic for a SIISI, SGeo &I CIOUOI~ harrier hole resonant tunneling diode (RTD) at 77 K. The RTD structure IS slmllar to the one shown in Fig. 23(b) and IS grown on a relaxed SI, SGe, 5 buffer layer.

Elemental

6.0

Semiconductor

Heterostructures

497

CONCLUSIONS Significant progress in the growth and application

has taken place over the last several years. resulted

in a better understanding

choice of growth techniques

of proper growth

including

of Si, _xGex alloys

These developments conditions,

have a wider

CVD methods, and a greater appre-

ciation of the thermal stability of Si,,Ge, layers and their application to a variety of device structures. Our treatment here only highlights some of the important developments in this material system. The use of heterostructures in Si-based devices and technology will continue to see increased activity and wider applications electronics in the future.

in less traditional

areas of micro-

ACKNOWLEDGEMENTS We are grateful to our numerous colleagues at the IBM T. J. Watson Research Center for fruitful collaborations and numerous discussions over the years.

REFERENCES 1. lyer, S. S., Silicon Epitaxy, (8. J. Baliga, ed.), pp. 97-175, Academic Press (1985) 2. Silicon Molecular Beam Epitaxy, vols. 1 and 2, (E. Kasper Bean, eds.), CRC Press, Boca Raton, FL (1988)

and J. C.

3. Proc. Is? Int’l. Symp. on Si Molecular Beam Epitaxy, vol. 85-7, (J. C. Bean, ed.), Electrochem. Sot. (1985) 4. Proc. 2nd /nt% Symp. on Si Molecular Beam Epitaxy, vol. 88-8, (J. C. Bean and L. J. Schowalter, eds.), Electrochem. Sot. (1988) 5. “Proc. 5th Int’l Conf. on MBE,” (Y. Shiraki and H. Sakaki, Sapporo (1988); J. Crystal Growth, vol. 95 (1989)

eds.),

6. “Proc. of 3rd Int’l. Conf. on Si MBE,” Thin So/id Films, vol. 183, (E. Kasper and E. Parker, eds.) (1989) 7. “Proc. 4th Int’l. Symp. on Si Molecular Beam Epitaxy,” vol. 220, (J. C. Bean, S. S. lyer, and K. L. Wang, eds.), Mat. Res. Sot. Proc., (1991)

498

Molecular

Beam Epitaxy

8. “Proc. of 1st Topical Symp. on Silicon-based Heterostructures,” (S. S. lyer, D. C. Houghton, and M. L. Green, eds.), 37th Nat. Symp. of the Am. Vat. Sot., Toronto (1990) 9. Herman, M. A. and Sitter, H., Molecular Verlag, Berlin, Heidelberg (1989)

Beam Epitaxy, Springer-

10. Braunstein, R., Moore, A. R., and Herman, F., Phys. Rev,, 109:695710 (Mar 1958) 11. van der Merwe, J. H., J. Appl. Phys., 34(1):117-122

(1963)

12. van der Merwe, J. H., J. Appl. Phys., 34(1):123-127

(1963)

13. Matthews, (1974)

J. M. and Blakeslee, A. E., J. Crystal Growth, 27:118-l

14. Matthews, J. W. and Blakeslee, 273 (1975) 15. Kasper, E., Herzog, (1975)

25

A. E., J. Crystal Growth, 32:265-

H. J., and Kibbel, H., Appl. fhys.,

16. Manasavit, H. M., Gergis, I. S., and Jones, Materials, 12(4):637-651 (1982)

A. B.,

8:199-205 J. Nectronic

17. Bean, J. C., Feldman, L. C., Fiory, A. T., Nakahara, S., and Robinson, I. K., J. Vat. Sci. Techno/. A, 2(2):436-440 (1984) 18. lyer, S. S., Tsang, J. C., Copel, M. W., Pukite, P. R., and Tromp, R. M, Appl. Phys. Letf., 54:219-221 (1989) 19. Gravesteijn, D. J., Zalm, P. C., Van de Walle, G. F. A., Vreizema, C. J., Van Grokum, A. A., and Ijzendoorn, L. J., Thin So/id Films, 183:191-196 (1989) 20. Copel, M., Reuter, M. C., Kaxiras, Leff., 63(6):632-635 (1989)

E., and Tromp, R. M, Phys. Rev.

21. lyer, S. S., Metzger, R. A., and Allen, F. G., J. Appl. Phys., 52:56085612 (1981) 22. Kuan, T. S., lyer, S. S., and Yeo, E. M., Proceedings of47thAnnual Mfg. of Electron Microscopy Society of America, pp. 580-581, San Francisco Press (1989) 23. Meyerson,

B. S., Appl. Phys. Lett., 48(12):797-799

24. Sedgwick, T. O., Berkenblit, 54:269&2692 (1989)

(1986)

M., and Kuan, T. S., Appl. Phys. Left.,

25. Gibbons, J. F., Gronet, C. M., and Williams, 47:721-723 (1985)

K. E., Appl. Phys. Lett.,

Elemental

Semiconductor

Heterostructures

26. Hirayama, H., Tatsumi, T., Ogura, A, and Aizaki, Lett., 51:2213-2215 (1987)

N., Appl. Phys.

27. Liehr, M., Greenlief, C. M., Kasi, S. R., and Offenberg, Phys. Lett., 56:629-631 (1990) 28. lyer, S. S., Arienzo, (1990)

499

M., Appl.

M., and de Fresart, E., Appl. Phys. Lett., 57:893

29. People, R. and Bean, J. C., Appl. Phys. Lett., 47(3):322324 30. Chidambarrao, D., Srinivasan, C. S., Appl. fhys. Lett., 54:l3

G. R., Cunningham, (1990)

(1985)

B., and Murthy,

31. Dodson, B. W. and Tsao, J. Y., Appl. Phys. Lett., 51(17):1325-1327 (1987) 32. Kohama, Y., Fukuda, Y., and Seki, M., Appl. Phys. Lett., 52:38& 382 (1988) 33. Stiffler, S. IX, Comfort, J. H., Stanis, C. L., Harame, Fresart, E., J. Appl. Whys., 70(3):1416-1420 (1991) 34. Fiory, A. T., Bean, J. C., Hull, R., and Nakahara, 31:4063-4065 (1985)

D. L., and de

S., Phys. Rev. B,

35. Tuppen, C. G., Gibbings, C. J., Davey, S. T., Lyons, M. H., Hockly, M., and Halliwell, M. A. G., Proc. 2nd In?‘/. Symp. Si MBE, ECS, 888:36-47 (1988) 36. lyer, S. S. and Pukite, P. R., unpublished 37. Kasper, E., Herzog, H. J., Daembkes, Res. Sot. Proc., 56:347-357 (1986) 38. lyer, S. S. and LeGoues,

(1988) H., and Abstreiter,

F. K., J. Appl. Phys., 65:4693-4698

39. Goorsky, M. S., Kesan, V. P., Ott, J. A., and Angilello, Technol. B, 1 O(2) :927-929 (1992)

G., Mat. (1989)

J., J. Vat. Sci.

40. Patton, G. L., Harame, D. L., Stork, J. M. C., Meyerson, B. S., Scilla, G. J., and Ganin, E., /EEE Nect. Dev. Lett,, 11:171-173 (1990) 41. Hull, R., Gibson, J. M., and Bean, J. C., Appl. Phys. Lett., 46:179181 (1985) 42. Houghton, D. C., Perovic, D. D., Baribeau, J. -M., and Weatherly, C., J. Appl. Phys., 67(4):1850-l 862 (1990) 43. Hansen, M. and Anderko, Hill, New York (1958)

K., Constifution

ofBinaryAl/oys,

G.

McGraw-

44. Kuan, T. S., Encyclopedia of Physical Science and Technology, A. Meyers, ed.), p. 521, Academic Press, New York (1989)

(K.

500

Molecular

Beam Epitaxy

45. Shahid, M. A., Mahajan, S., Laughlin, Rev. Leti., 58(24) :2567-2570 (1987)

D. E., and Cox, H. M., fhys.

46. Martins, J. L. and Zunger, A., Phys. Rev. Lett., 56(13):1400-1403 (1986) 47. Flynn, C. P., Phys. Rev. Lett., 57(5):599-602

(1986)

48. Bernard, J. E. and Zunger, A., Phys. Rev. f3, 44:1663 (1991) 49. Ourmazd,

A., Bean, J. C., Phys. Rev. Leff., 55(7):765-768

(1985)

50. LeGoues, F. K., Tromp, R. M., Kesan, V. P., and Tsang, J., Phys. Rev. 9, Rapid Comm., 47(15):10012-10015 (1993) 51. Miller, E., Nissen, H. -U., Ospelt, M., and von Kanel, H, Phys. Rev. Left, 63(17):1819-l 822 (1989) 52. Miller, E., Nissen, H. -U., Mader, K. A., Ospelt, M., and von Kanel, H., Phi/OS. Mag. Left., 64(4):183-189 (1991) 53. LeGoues, F. K., Kesan, V. P., and lyer, S. S., Phys. Rev. Lett., 64:4&42 (1990) 54. LeGoues, F. K., Kesan, V. P., lyer, S. S., Terspff. J., and Tromp, R., Phys. Rev. Lett., 64:64-66 (1990) 55. Kesan, V. P., LeGoues, 46:1576-l 581 (1992)

F. K., and lyer, S. S.,

56. Tsang, J. C., Kesan, V. P., Freeouf, J. L., LeGoues, Phys. Rev. 9,46:69076914 (1992)

Phys.

Rev. 9,

F. K., lyer, S. S.,

57. Kelires, P. C. and Tersoff, J., Phys. Rev. Lett., 63(11):1164-1167 (1989) 58. Metzger, R. A. and Allen, F. G., J. Appl. Phys., 55:931 (1984) 59. Kuan, T. S. and lyer, S. S., Appl. Phys. Leti., 59:2242-2244 60. Baslev, J., fhys.

Rev., 143:636-647

61. People, R. Phys. Rev, 9, 32:1405-l

(1991)

(1966) 408 (1985)

62. People, R. and Bean, J. C., Appl. Phys. Left, 48:538-540

(1986)

63. Lang, D. V., People, R., Bean, J. C., and Sergent, A. M., Appl. Phys. Leff., 47:1333-l 335 (1985) 64. Cheng, J. -P., Kesan, V. P., Grutzmacher, D. A., Sedgwick, T. O., and Ott, J. A., Appl. Whys. Left, 62(13):1522-1524 (1993); Cheng, J. -P., Kesan, V. P., Grutzmacher, D. A., and Sedgwick, T. O., Appl. Phys. Leti., 64(13):1681-1683 (1994) 65. Van de Walle, C. G. and Martin, 4:1055-l 059 (1986)

R. M., J, Vat.

Sci. Technol.

9,

Elemental

66. Zeller, Ch. and Abstreiter,

Semiconductor

G., ZPhys.

Heterostructures

B, 64:137-143

(1986)

67. Abstreiter, G., Brugger, H., Wolff, T., Jorke, H., and Herzog, Phys. Rev. Lett., 54(22) :2441-2443 (1985) 68. People, R. Phys. Rev. B, 34:2508-2510

501

H. -J.,

(Feb 1986)

69. lyer, S. S., Solomon, P. M., Kesan, V. P., Bright, A. A., Freeouf, J. L., Nguyen, T. N. and Warren, A. C., /EEE Elect. Dev. Leff., 12(5):246248 (1991) 70. Nayak, D. K., Woo, J. C. S., Park, J. S., Wang, K. -L., and MacWilliams, K. P., /EEE Elect. Dev. Lett., 12(4):154-156 (1991) 71. Garone, P. M., Venkataraman, V., and Sturm, J. C., IEEE Elect. Dev. Left., 12(5):23&232 (1991) 72. Murakami, E., Nakagawa, K., Nishida, Elect. Dev. Lett., 12(2):71-73 (1991)

A., and Miyao,

M.,

/EEE

73. Subbanna, S., Kesan, V. P., Tejwani, M., Restle, P. J., Mis, D., and lyer, S. S., Pfoc. Symp. VLSl Technology, pp. 103-I 04 (1991) 74. Kesan, V. P., Subbanna,

S., Restle, P. J., Tejwani, M. J., Aitken, J. M., lyer, S. S., and Ott, J. A., Int’l. Elect. Dev. Mfg. (IEDM) Tech. Dig., pp. 25-28 (1991)

75. Verdonckt-Vandebrock,

S., et al., IEEE Elect. Dev. Letf., 12(8):447-

449(1991) 76. People, R., Bean, J. C., Lang, D. V., Sergent,

Wecht, K. W., Lynch, R. T., and Baldwin, 45(11):1231-1233 (1984) 77. Pearsall, T. P. and Bean, 7(5):308-310 (May 1986)

J. C.,

A. M., Stormer, H. L., K., Appl. Phys. Lett.,

/EEE Elect. Dev, Lett,,

EDL-

78. Dambkes, H., Herzog, H. -J., Jorke, H., Kibbel, H., and Kasper, E., /EEE Trans. Elect. Dev., ED-33(5):633-638 (May 1986) 79. Liu, H. C., Landheer, D., Buchanan, Phys. Lett., 52:1809-l 811 (1988)

M., and Houghton,

D. C., Appl.

80. Rhee, S. S., Park, J. S., Karunasiri, R. P. G., Ye, Cl., and Wang, K. L., Appl. Phys. Lett., 53:204-206 (1988) 81. Kasper, E., Herzog, H. -J., Jorke, H., and Abstreiter, and Microstructures, 3(2) : 141-I 46 (Feb 1987)

G., Superlattices

82. Gennser, U., Kesan, V. P., lyer, S. S., Bucelot, T. J., and Yang, E. S., J. Vat. Sci. Technol. B, 8:210-213 (1990) 83. Gennser, U., Kesan, V. P., Syphers, D. A., Smith, T. P., lyer, S. S., and Yang, E. S., Phys. Rev. Lett., 67:3828-3831 (1991)

502

Molecular

Beam Epitaxy

84. Temkin, H., Pearsall, T. P., Bean, J. C., Logan, R. A., and Luryi, S., Appl. Phys. Lett., 48(15) :963-965 (Apr 1986) 85. Kesan, V. P., May, P. G., Bassous, E., and lyer, S. S., In/‘/ Elect. Dev. Mtg. (IEDM) Tech. Dig., pp. 637-640 (1990) 86. Kesan, V. P., May, P. G., LeGoues, Growth, Ill:936942 (1991)

F. K., and lyer, S. S., J. Crystal

87. Luryi, S., Pearsall, T. D., Temkin, H., and Bean, J. C., /EEE Electron Dev. Letters, EDL-7(2):104-106 (Feb 1986) 88. Pearsall, T. P., Temkin, H., Bean, J. C., and Luryi, S., /EEE Electron Dev. Letters, EDL-7(5):33&332 (May 1986) 89. Temkin, H., Antreasyan, A., Olsson, N. A., Pearsall, T. P., and Bean, J. C., Appl. Phys. Lett., 49(13):809-811 (Sept 1986) 90. lyer, S. S., Patton, G. L., Stork, J. M. C., Meyerson, B. S., and Harame, D. L., /EEE Trans. Electron. Dev., 36:2043-2064 (1989) 91. Patton, G. L., lyer, S. S., Delage, S. L.,Tiwari, S., and Stork, J. M. C., /EEE Electron Dev. Lett,, 9:165-167 (Apr 1988) 92. King, C. A., Hoyt, J. L., Groonet, C. M., Gibbons, J. F., Scott, M. P., and Turner, J., /EEE Electron Dev. Lett., 10:52-54 (Feb 1989) 93. Taft Ft. C., Plummer, J. D., and lyer, S. S., IEEE Electron Dev. Lett., 9:14-l 6 (Jan 1989) 94. Temkin, H., Bean, J. C., Antreasyan, A., and Leibenguth, Phys. Leti., 52:1089-l 091 (Mar 1988) 95. Tatsumi, T., Hirayama, 897 (Mar 1988)

R., Appl.

H., and Aizaki, N., Appl, Phys. Lett., 52:895-

96. Gruhle, A., Kibbel, H., Konig, U., Erben, U., and Kasper, E., /EEE Nectron Dev. Lett., 13(4):206-208 (1992) 97. lyer, S. S., Patton, G. L., Delage, S. L., Tiwari, S., and Stork, J. M. C., Int’l. Elect. Dev. Mtg. (IEDM) Tech. Dig., pp. 874876 (Dee 1987) 98. Harame, D. L.., Stork, J. M. C., Patton, G. L., lyer, S. S., Meyerson, B. S., Scilla, G. L., Crabbe, E. F., and Ganin, E., /nt’/. Elect. Dev. Mtg. (IEDM) Tech. Dig., pp. 889-891 (Dee 1988) 99. Kroemer, H., Proc. IEEE., 70:13-25

(Jan 1982)

100. Gruhle, A., J. kc. Sci. and Technol. S, 11(3):1186-l 189 (1993); Kasper, E., Kibbel, H., Herzog, H.-J., and Gruhle, A., Extended Abstracts of the So/id-State Dev. Mtg. Conf., Chiba, Japan, 419, Jpn. Sot. of Appl. Phys. (1993) 101. Prinz, E. J. and Sturm, J. C., /EEE Elect. Dev. Lett., 12(12):661-663 (1991)

Elemental

Semiconductor

Heterostructures

102. Streetman, B. G., So/id State Necfronic Hall, Englewood Cliffs, NJ (1990) 103. Levitas, A., Phys. Rev., 99:1810-1814

503

Devices, p. 439, Prentice(1955)

104. Manku, T. and Nathan, A., /EEE Elect. Dev. Leff., 12(12):704-706 (1991) 105. Smith, C. and Welbourn, A. D., Proc. /EEE 7987 Bipolar Circuits and Technology Meeting, pp. 57-64 (1987) 106. Wang, P. J., Meyerson, B. S., Fang, F. F., Nocera, J., and Parker, B., Appl. Phys. Leff., 55(22):2333-2335 (1989) 107. Fang, F. F., Wang, P. J., Meyerson, B. S., Nocera, J. J., and Ismail, K. E., Surf. Sci., 263:175-178 (1992) 108. Grutzmacher, D. A., Sedgwick, T. O., Northrop, A., Powell, A. R., and Kesan, V. P., J. kc. 11(3):1083-1087 (1993) 109. Schirber, J. E., Fritz, I. J., and Dawson, 46(2):187-l 89 (1985)

G. A., Zaslavsky, Sci. Technol. f3,

L. R., Appl.

Phys. Lett.,

110. Mii, Y. J., Xie, Y. H., Fitzgerald, E. A., Monroe, D., Thiel, F. A., Weir, B. E., and Feldman, L. C., Appl. Phys. Leff., 59(13):1611-1613 (1991) 111, Schuberth, G., Schaffler, F., Besson, M., Abstreiter, E., Appl. Phys. Left., 59(25):3318-3320 (1991)

G., and Gornik,

112. Schaffler, F., Tobben, D., Herzog, H. J., Abstreiter, G., and Hollander, B., Semicond. Sci. Techno/., 7(2):260-266 (1992) 113. Ismail, K., Meyerson, B. S., Rishton, S., Chu, J., Nelson, Nocera, J., /EEE Necf. Dev. Left., 13(5):229-231 (1992)

S., and

114. Nelson, S. F., Ismail, K., Nocera, J. J., Fang, F. F., Mendez, E. E., Chu, J. O., and Meyerson, B. S., Appl. Phys. Let, 61(1):64-66 (1992) 115. Morrison,

I. and Jaros, M., Phys. Rev. B, 37(2):916-921

116. Presting, H., Kibbel, H., Jaros, Abstreiter, G., and Grimmeiss, 7(9):1127-1148 (1992)

(1988)

M., Turton, R. M., Menczigar, U., H. G., Semicon. Sci. Technol.,

117. Sturm, J. C., Manoharan, H, Lenchyshyn, Rowell, N. L., Noel, J. P., and Houghton, 66(10):1362-1365 (1991)

L. C., Thewalt, M. L. W., D. C., Phys. Rev. Leff.,

118. Robbins, D. J., Canham, L. T., Barnett, S. J., Pitt, A. D., and Calcott, P., J. Appl. Phys., 71(3):1407-1414 (1992)

504

Molecular

Beam Epitaxy

119. Rowell, N. L., Noel, J. P., Houghton, Phys. Left., 58(9):957-958 (1991)

D. C., and Buchanan,

120. Mat. Res. Sot. Syrnp. froc., (S. S. lyer, R. T. Collins, Canham, eds.), vol. 256 (1992)

M., Appl. and L. T.

121. Kesan, V. P., Bassous, E., Munguia, P., Pesarcik, S. F., Freeman, M., lyer, S. S., and Halbout, J. -M., J. kc. Sci. Technol. A, 11(4):1736-l 738 (1993) 122. Splett, A., Schmidtchen, J., Schuppert, B., Petermann, E., and Kibbel, H., Elect. Leff., 27:1486-1488 (1991)

K., Kasper,

123. Pesarcik, S. F., Treyz, G. V., lyer, S. S., and Halbout, J. -M., Elect. Lett., 28:159-l 60 (1992) 124. Lin, T. L., George, T., Jones, E. W., Ksendzov, M. L., Appl. Phys. Lett., 60:380-382 (1992)

A., and Huberman,

125. Tsaur, B. Y., Chen, C. K., and Marino, S. A., /EEE Elect. Dev. Lett., 12:293-295 (1991) 126. Karunasiri, R. P. G., Park, J. S., Wang, 61(20):2434-2436 (1992)

K. I., Appl.

Phys. Lett.,

127. Gibbons, J. F., King, C A., Hoyt, J. L., Noble, D. B., Gronet, C. M., Scott, M. P., Rosner, S. J., Reid, G., Laderman, S., Nauka, K., Turner, J., and Kaminis, T. I., Int’l Elect. Dev. Mtg, (IEDM) Tech. Dig., pp. 566-569 (Dee 1988)

6

MBE Growth of High TC Superconductors Darrell G. Schlom and James S. Harris, Jr.

1 .O INTRODUCTION The discovery by Bednorz and Miillern) of a class of layered crystalline materials which exhibit superconductivity at unprecedentedly high transition temperatures (high T,)* opened new possibilities for future electronic devices. Because of the layered nature of these materials, MBE is a natural method to explore their growth. However, in contrast to the relatively fixed group V stoichiometry of Ill-V compounds grown by MBE, small variations in oxygen or cation composition change these materials from superconductors to semiconductors to insulators, thus requiring a deposition technique with very precise control of stoichiometry. While troublesome for deposition of superconducting films, this dramatic variation in conductivity, together with the existence of chemically and structurally compatible ferroelectric and piezoelectric,f3)-t6) ferrimagneticp] and metallictst-tlo) oxides provide exciting possibilities for new electronic devices, as the spectrum of physical properties exhibited by these oxides is unparalleled by any other materials system. Many new device concepts will likely employ alternately layered structures with dimensions as thin as * The relative term “high Tc”is used here to signify superconductors with Tc significantly greater than those compounds which until 1986 were known as “high Tc” (i.e., the Al 5 compounds with Tc up to 23 Kfor NbsGe, which have since become known as “low Tp). These new compounds, with the exception of the recently discovered fullerenes (e.g., A$,, where A is an alkaline earth), are oxides and are distinguishable from other superconductors by an anomalously high T, for their density of states at the Fermi level.f2) The growth of fullerenes is not discussed in this chapter.

505

666

Molecular

Beam Epitaxy

the superconducting

coherence

length, 5, which

stroms to a few tens of angstroms superconducting

in high

results were achieved

ranges from a few ang-

7, materials.

with polycrystalline

these materials which had very low critical current densities, unsuitable

for electronic

applications.

The earliest bulk forms of making them

Since the current carrying capability

of high 7, superconductors is known to be degraded by grain boundaries,f”) the ideal form of these materials for most electronic applications is likely to be that of epitaxial films, prepared in such a way that composition and structure can be controlled at the level of single atomic layers. The first epitaxial films of a high 7, superconductor were prepared by sputtering.tl*) The discovery of additional high T, superconductors, each in previously unmapped regions of composition (i.e., an absence of relevant phase diagram information), provided a challenging task for those attempting to prepare thin films of these materials. The first epitaxial films with Tc’s exceeding the vaporization temperature of nitrogen, 77 K, were synthesized using electron beam (e-beam) evaporation to codeposit constituent elements.[13)[14] This was followed by a high temperature

the ex-

situ oxygen anneal, during which the high T, phase formed by solid phase epitaxy. Other thin film deposition methods soon followed, including sputtering,f15)f16) molecular beam deposition,f17)[1el and a relatively new technique known as laser ablation or pulsed laser deposition (PLD).f1g1t201 These methods produced epitaxial films with drastically improved critical current densities compared to their polycrystalline bulk counterparts; however, the high temperature (800-900°C) ex-situ annealing step involved in all these methods was not compatible with the controlled synthesis of layered heterostructures. The main obstacle to the in-situ growth of high

T, phases by tradi-

tional thin film methods was the necessity for high oxygen pressure during growth, as discussed in detail in Sec. 2.3. For diode and magnetron sputtering, negative ion bombardment was also a significant problem.f21)-f271 The use of relatively high oxygen pressures in e-beam evaporation (reactive evaporation) ,f*s)t*91 PLD,f201f301 ion beam sputtering,[311 and offaxisf24)t27)t32)and hollow cylinderf*“)

reactive sputtering

geometries,

or with

total pressures of several hundred mtorr to mitigate negative ion bombardment in the conventional planar sputtering geometry,t25) led to the in-situ formation of epitaxial high T, superconductor films with excellent electrical transport properties. These in-situ techniques are now routinely used for the preparation of layered heterostructures, both for device structuresf33] and for investigating the physics of high T, superconductivity.t34)-t36)

High Tc Superconductors

507

The advantages of these other techniques over MBE are primarily cost and, in the case of PLD and off-axis sputtering, the nearly faithful composition

transfer from target to substrate,

which allows a single multi-

component target with the same composition as the desired film to be used,t20)f27) alleviating the need for accurate composition control, a critical component of MBE. MBE, on the other hand, is free from the micron-sized “boulders” common to PLD filmst20j as well as the energetic species present in sputtering, which can lead to interlayer mixing due to ion bombardment effects.f3g1t40j The multi-element deposition control, growth flexibility, and in-situ monitoring advantages of MBE are well suited to the growth of high T, phases which cannot be produced in single phase form by bulk techniques, including the customized growth of new metastable materials, and heterostructures containing these phases. Other deposition techniques, in particular PLD and off-axis sputtering, are, from an economic and process simplicity perspective, generally better suited than MBE to the synthesis of heterostructures made up of phases, each of which can be produced by bulk techniques in single phase form (i.e., the formation energy of each phase is sufficiently favored over other phases that could accommodate its composition). Interestingly, most, if not all, high T, materials are metastable at room temperature and below.t41j Although no firm theoretical connection has been made between metastability and increasing T,, from what is known about the existing high T, phases, it appears that techniques capable of synthesizing metastable structures will be required to achieve superconducting materials with ever greater T,‘s.[~~] Bulk chemical synthesis routes have led to the vast majority of high T, materials discovered to date, but considering the increasing difficulty and complexity of the synthesis

routes needed to form these metastable

method with atomic-scale structure

monitoring,

decomposition

methods.

materials,

a synthesis

control, in-situ process and crystalline

and a low growth temperature

of the deposited

over conventional synthesis

layering

structures

solid state reaction

methods

MBE is an excellent

to kinetically

offers significant

limit the

advantages

as well as other thin film

synthesis

technique

not only for

meeting the materials synthesis challenge of current and future high T, materials, but also for incorporating these high T, layers into device microstructures. Motivated by the potential of creating high T, microstructures with customized atomic layering, together with the ability of MBE to control composition and layering with atomic scale precision and compatibility

888

Molecular

with vacuum

Beam Epitaxy

in-situ characterization

techniques,

researchers

developed

MBE techniques for the controlled in-situ growth of high T, superconductors.t42]j46] In contrast to the relative simplicity of the other materials systems to which MBE has been successfully applied,t4’t the growth of fully oxidized multi-element high T, materials by MBE involves significant challenges. A powerful oxidant is necessary to oxidize the constituent elements sufficiently during growth to form the desired structure at a pressure low enough to preserve the long mean free path necessary for MBE. In addition, composition control is crucial in order to provide the proper mix of constituent species to form the desired multi-component phase and controllably dope it, while avoiding the formation of unwanted impurity phases. Finally, to realize layered metastable structures, a low temperature in-situ growth method is required. After discussing the relevant physical properties of high T, superconductors, this chapter describes the challenges, experimental methods, current status, and outlook of the growth of high Tc superconductors and related oxides by MBE. This new branch of MBE has already demonstrated many promising capabilities which are unmatched by other growth techniques. Custom-layered oxide heterostructures, including high T, Josephson junctions and metastable structures, are currently being made by MBE with unit cell layering precision. In addition, MBE offers an ideal environment for the use of vacuum in-situ characterization methods, allowing the growth process itself to be studied. Such advantages make MBE a promising growth technique for the controlled preparation of high T, superconductors, As the MBE technique is further perfected for the growth of high T, oxide heterostructures, it is expected that this growth technique will be particularly useful for device fabrication. In addition, it will offer an excellent technique to fabricate unique metastable structures or superlattices to test high T, theories, materials, 1.1

Crystal Structures

which may lead to the growth of even higher

and Types

of Building

T,

Layers

Many structurally-related superconducting oxide phases have been identified, all of which have perovskite-related structures.f‘@) The crystal structure of perovskite (CaTiOs) and some of the better known structurallyrelated high T, superconductors are shown in Fig. 1. High T, structures may be viewed as a superlattice comprised of more fundamental layers, where the direction of layering is along the c-axis (the vertical direction in Fig. 1).

Such a perspective

provides

a framework

for the generalized

High Tc Superconductors construction

of existing and potentially

T, structures fundamental

can be assembled layer types.

509

new high T, structures.f50) All high

from a relatively

These 2-dimensional

small number of more

sheets shall be referred to

as building layers in analogy to O-dimensional building blocks. Just as building blocks are assembled in three dimensions, building layers are stacked in one dimension assemble l

=Ba

l

=Bi

l =Ca l

=cu

(the c-axis and the direction

Sr . =Ti

0 =

90K

. =Tl l ,y

0”

l =La l =Nd

of MBE growth) to

high T, structures.

,$llO 40 K

25 K

K

90 K

l=O

ED CaTiO,

1 (Nd (La,Sr)2Cu04

Figure 1. The crystal structure of perovskite (CaTi0.J and some of the better known high Tc superconductors (CaTiO,, (La,Sr),CuO,, (Nd,Ce),CuO,, YBqCu30,~, Bi,Sr$aCu,Os, TI,BqCqCu,O,, and TIBa&qCu,Os). Two equivalent representations of these crystal structures are shown: the atomic positions (above) and the copper coordination polyhedra (below). The oxygen atoms occupy the vertices of the copper coordination polyhedra. The tetragonal subcells of the (La,Sr),CuO, and Bi,Sr,CaCu,O,+, structures are shown for clarity and to illustrate the similarities between these perovskite-related phases. The relative sizes of the atoms reflect their relative ionic radii as given by Ref. 49. The atomic shading given here is used throughout this chapter. The approximate superconducting transition temperatures are also shown.

510


A common feature of the crystal structures of all known coppercontaining high T, superconductors is the presence of CuO, layers. These CuO,

layers, which

are often slightly

puckered,

consist

of an array of

corner-sharing CuO, squares, as shown in Fig. 2. The CuO, layers are separated from one another along the c-axis of the crystal structure by various intervening Depending on the intervening layers, are square-planar

layers, the known types of which are shown in Fig. 3. presence of and positions of oxygen atoms in the the copper coordination polyhedra of the CuO, layers (4-coordinated Cu), pyramidal (5coordinated Cu), or

octahedral (&coordinated Cu), as shown in Fig. 2. The superconducting charge carriers flow in these CuO, layers. High T, superconductivity with electrons as the charge carriers (n-type superconductivity) has been attained only in square-planar coordinated CuO, layers, whereas superconductivity with holes as the charge carriers (p-type superconductivity) has been achieved only in pyramidal and octahedral coordinated CuO, layers.[50]

Square-Planar (4-Coordinated Cu)

Pyramidal (5-Coordinated Cu)

Octahedral (6-Coordinated Cu)

Figure 2. The atomic structure of a CuO, layer, the common structural component of all known superconductors with T, exceeding the boiling point of liquid nitrogen. Depending on the surrounding layers, oxygen atoms may be present at apical positions, changing the copper coordination from 4 to 6. (After Ref. 50.) There is only one known high contain CuO, layers:15’1 (Ba,K)BiO,,

Tc superconductor

which

does not

which has a T, of about 30 K. Like the

structurally-related solid solution[52] Ba(Pb,Bi)O, with T, of approximately 13 K, it has a simple cubic perovskite structure, shown in Fig. 1, with Bi occupying the octahedrally coordinated Ti position and Ba and K occupying the larger Ca site. From a building layer standpoint, (Ba,K)BiO, consists of alternating

BiO, and (Ba,K)O building layers.


(b)

BaO

cue,.* BaO

(4

(e>

BaO TIO BaO

BiO sro

511

(Sr,Pb)

(R&a)

Tao,

Cl

02 (C&W

Figure 3. The building layers of all known layered CuO, structures:

(a) [La,Sr)O(La,Sr)O] or [(Sr,Ca)(Br,CI)-(Sr,Ca)(Br,CI)]*; @) [BaO-Cu),_&-BaO]; (c) [BaO(d) [BaO-TIO-BaO] or [BaO-HgO,-BaO] or [SrO-(Bi,Cu)O,-SrO] CuO-CuO-BaO]; or [SrO-(Bi,Cd)O-SrO] or [SrO-(Pb,Sr)O-SrO] or [SrO-(Pb,Cu)O,-SrO] or [SrO(Pb,Cd)O-SrO] or [SrO-(Ce,Cu)O,-SrO]* or [SrO-(Ce,Cd)O-SrO]; (e) [BaO-TlOTIO-BaO] or [BaO-HgO,-HgO,-BaO] or [SrO-BiO-BiO-SrO] or [(Ba,Sr)O(Pb,Cu)O,-(Pb,Cu)O,-(Ba,Sr)O]; (y9 [SrO-PbO-CuO,-PbO-SrO]; (&j [Lao-SnO,Lao]* or [BaO-Tao,-BaO]* or [SrO-Tao,-SrO] or [SrO-NbO,-SrO]; @) [SrOGaO-SrO] or [SrO-Coo-SrO]* or [SrO-AIO-SrO]*; 0 [Sr-CO,-Sr] or [(Re,Ba)BOs-(Re,Ba)]; (j) [(/?e,Ca)]; (k) [(Ce,Re)-O,-(Ce,Re)]; and (I) [(Sr,Pb)-Cl-(Sr,Pb)]*. The tetragonal subcell of each building layer is outlined. CuO, layers are shown above and below each building layer to illustrate their attachment positions. Partial substitution for the constituents of these basic building layers is frequently possible and may be used for doping purposes (e.g., Bi - Pb, Tl - Pb, Sr - Ba, Sr - La, Ca - Re). Most of these building layers are constituents of the coppercontaining high T, superconductors discovered to date. Those intervening layers known to occur between CuO, layers, but so far not constituents of superconducting structures, are marked with an asterisk (*) . (After Ref. 50.)

Although high T, phases are commonly referred to by formulae implying that they are stoichiometric, e.g., Bi,Sr,CaCu,O,, this is an oversimplification. In reality, significant cation mixing, anion vacancies, or cation vacancies often occur within the building layers, with several percent

mixing

quite common

in some layers,

although

not within

the

512

Molecular

CuO, layers.

Beam Epitaxy

For example,

site-sensitive

crystal of the high T, superconductor overall composition

structural

Bi,Sr,CaCu,O,

was Bi,,,,Sr,.,,Ca,,,,Cu,O,.,

refinement

of a single

determined

that its

with the Ca-site occu-

pied by 75% Ca, 19% Sr, and 6% Bi.f53)t54j Other sites in the crystal structure

also showed

cation mixing,

although

to a lesser extent.

Non-

stoichiometry is an important aspect of high T, superconductors. For some phases, the adjustable non-stoichiometry (solid solution) within a particular building layer is explicitly stated, e.g., (La,Sr),CuO,, (Nd,Ce),CuO,, YBa,Cu,O,,, and (Ba,K)BiO,. As described in the next section, non-stoichiometry in the building via chemical doping. Pure Bi,Sr,CaCu,O, since the average oxidation state of Cu would be no charge carriers,t4’j whereas perconducting and the average oxidation 1.2

Chemical

layers provides charge carriers would not be superconducting, would be exactly +2 and there Bi,,,,Sr,.,,Ca,,,,Cu,O,,, is sustate of Cu is +2.21.

Doping

In addition to structural control, realizing the optimal carrier concentration in the CuO, layers is a crucial part of the synthesis of high T, materials. The building layers surrounding the CuO, layers serve as charge reservoirs to dope the CuO, layers by chemical means.t55] Doping the CuO, layers with holes or electrons is a necessary, although not sufficient, condition for achieving superconductivity in these materials. A convenient method for charge counting to assess chemical doping is through the use of oxidation states. t In order to quantify the doping of the CuO, layers it is sufficient to know the oxidation state of Cu, since that of 0 is defined to be -2. Cu has multiple oxidation states, ranging from +l to +3. The average oxidation state of Cu is determined from the phase composition,

the standard oxidation states of the remaining

ions (e.g., 02-,

Sr2+, La3+, etc.), and the condition of charge neutrality.* For example, the average oxidation state of Cu is +(2+x) in La,_,Sr,CuO, and +(2-x) in Nd,_xCe,CuO, (where Ce4+ is present). In the parent compounds, La,CuO, and Nd,CuO,, where no chemical doping is present (x = 0), the oxidation state of Cu is t2 and the compounds are not superconducting. As x increases and the parent compounds are doped, carriers enter the CuO, layers and the average oxidation state of Cu deviates from t2. t Note that the oxidation state of an ion is not the same as the true charge on the ion.fs6) t In compounds containing elements which, in addition to Cu, may assume a variety of oxidation states (e.g., Bi, Pb, TI, Hg) or compounds with non-equivalent Cu sites (e.g., the CuO,“chains” in Fig. 36), additional information is necessary to ascertain the oxidation state of Cu in the CuO, layers.p)-fsr) Non-integral average oxidation states may also be described as a mixed oxidation state. For example, an average Cu oxidation state of +2.2 corresponds to a mixture of 80% Ct?+ and 20% Cu”.


The oxidation

513

state of Cu is greater than 2 for p-type superconductors

and less than 2 for n-type superconductors. The amount of charge transferred from the surrounding (charge reservoir) building layers to the CuO, layers is given by the deviation of the oxidation state of Cu from +2, or equivalently by the deviation of each CuO, unit from -2: [CuO.J**~, where 6 is positive for hole doped (p-type) and negative for electron doped (n-type) superconductors. Studies of the maximization of Tc with dopingt5q-t61) have revealed that a maximum T, for each structure occurs at about 161 C(0.2, as shown in Fig. 4 for several p-type superconductors.

100

80

60 Y 2 40

0

0.1

0.2

0.3

0.4

Holes per CuO2

Figure 4. T, as a function of hole concentration, nh, per CuO, unit for various high T, superconductors as determined by chemical titration. (From Ref. 62.)

514

1.3

Molecular

Beam Epitaxy

Phase Diagrams The multicomponent

space in which

a given

phase high

multitude of phases, including unresolved. many high

diagrams

T, phase

spanning is found

the compositional invariably

many for which the structural

contain

a

details remain

A precise mapping of the equilibrium phase relationships in T, systems is hampered by the large number of elemental

components, a large number of previously reaction kinetics at the growth temperatures

unknown compounds, slow of interest, and the presence

of impurities (e.g., carbonates or fluxing agents). High Tc systems contain at least four components, and some important ones have as many as six. For example, lead is commonly added to Bi,Sr,Ca,Cu,O,, in order to synthesize this 110 K superconductor in single phase form by bulk techniques, making it a 6-component system. The significant variance of the phase relationships as functions of oxygen pressure and temperature further complicates this mapping. Interestingly, no high T, materials are known that are thermodynamically stable at room temperature or below, however, the kinetics of the decomposition of high T, superconductors are so sluggish at low temperatures that this fact does not deter their utility.fss] While some high T, superconductors are thermodynamically stable at high temperatures and may be quenched to avoid decomposition, others (e.g., YBa,Cu,O,) are not stable at any temperature or pressure.t41)t64] Their synthesis is commonly achieved by first forming a thermodynamically stable parent structure (e.g., YBa,Cu,O,) and subsequently altering its composition after the sample has been cooled to a temperature where the parent structure is kinetically prevented from decomposing. The metastability of possibly all high T, superconductors has led researchers to propose that these materials will continue to become more and more metastable as T, increases.t41] Although all high T, materials allow a finite range of atomic substitution for at least one component, which allows their carrier concentration to be altered by chemical doping, these solid solutions are far narrower than the solid

solutions

present

in common

semiconductor

systems.

The

limited range of compositions over which the high Tc materials are stable poses significant composition control constraints for growing films free of second phases, as can be seen from the phase diagrams. The pseudo-ternary phase diagram of the YO, ,,-BaO-CuO system at 900°C in 1 atmosphere of oxygen is shown in Fig. 5.t6q This diagram represents a section through the full Y-Ba-Cu-0 quaternary phase diagram


tetrahedron.

As the oxygen

pressure

(e.g., T = SSO’C), the phases compound

YBa&u,O,,

Below an oxygen

and tie-lines

change

partial pressure, of the various

at a fixed temperature

in contact

considerably,

with the high

as shown

T,

in Fig. 6.f6‘j)

Po2, of 0.30 torr, the YBa,Cu,OrB

phase is no longer thermodynamically ture dependence

is lowered

515

stable at T = 850°C.

reactions

is indicated

The tempera-

in Fig. 7.f6’) The

minimum PQ at which a desired high T, structure is thermodynamically stable has obvious significance for its MBE synthesis under vacuum conditions. The competing constraints of a high oxygen pressure for phase stability and a low oxygen pressure to maintain the long mean free path necessary for MBE have necessitated the use of activated oxygen species for MBE growth, as discussed in Sec. 2.

BaY,O. Ba,Y,O.

B40)

Figure 5. The subsolidus pseudo-ternary YO,,,-BaO-CuO 900°C in 1 atm. oxygen. The numbers compounds YB~&u,O,~, YB~&I,O,,+,, spectively. (From Ref. 65.)

phase diagram at T = 123, 143, 163, and 211 represent the YB~&u,O,,+~, and Y,BaCuO,, re-

516

Molecular

Beam Epitaxy

BeCuO,

BaCu202

W

(a) 14.2 Torr

cue

(‘I,,&I;

BacuO,

(d) _,&,: -202

Bacu,O2

(h) .,;k2 BaCu02

Bacuo2

-2

.I TOIT

BaCu202

CYO

$

O.j’fj Tom

Cyo

(i)_,&;;; BaCu02

Torr

BaCu20,

Figure 6. The effect of oxygen pressure on the pseudo-ternary Y-Ba-Cu-0 phase diagram at T = 850%. The oxygen partial pressure (PQ,) range for each diagram is indicated (from Ref. 66).

High TC Superconductors

T (“c):

900

IA

I

I

850 Ir

1

800 I

517

750 I

! 124 formation from 123.5 8 CuO ,

10-l

1o-2

po2 W-4

1o-3

Plateau #: 0 I 0 II n Ill, IV

lo4

OV A VI A

ifi1

VIII + IX l

ii:_.___,____.____,

10

-5

0.80

0.85

Figure 7. The temperature dependence tions shown in Fig. 6 (from Ref. 67).

The 5-component

Bi-Sr-Ca-Cu-0

0.90

0.95

of the Y-Ba-CA-0

system

decomposition

1.00 reac-

has an even richer phase

diagram. The four pseudo-ternary faces of the pseudo-quaternary BiO,,s-SrO-CaO-CuO tetrahedron are shown in Fig. 8 for T = 850°C in one atmosphere of air.[681[6g]The minimum Po2 at which the Bi,Sr,CaCu,O, and Bi2Sr2Ca.&u,0,,

phases are stable is shown in Fig. 9.f701

518

Molecular

Beam Epitaxy

Sr:illiyOc;

(4 Figure 8. (a) The pseudo-quaternary BiO,,,- SrO-CaO-CuO phase diagram at T = 850 “C in air (from Ref. 68), along with the four pseudo-ternary phase diagrams that form the faces of the pseudo-quaternary BiO,,,-SrO-CaO-CuO tetrahedron: (b) the pseudo-ternary BiO,,,-SrO-CuO phase diagram at T = 850°C in air (from Ref. 69), (c) the pseudo-ternary SrO-CaO-CuO phase diagram at T = 850°C in air phase diagram at T = (from Ref. 69), (d) the pseudo-ternary BiO ,,,-SrO-CaO 850°C in air (from Ref. 69), (e) the pseudo-ternary BiO,,,-CaO-CuO phase diagram at T = 800°C (from Ref. 69). (Cont’d next page.)


519

(4

Figure 8. (Cont’cl)

520

Molecular

Beam Epitaxy

(d)

CO’3

*/2(8Yz03) coo

A

Figure 8.

(Cont’d)

(4 rBoooi

aJtlSSaJd

Ua6hXO

(we) amssaJd ua6Axo

(UJle)


521

522

1.4

Molecular

Beam Epitaxy

Uncontrolled

Intergrowths

By using CuO,

Inherent

in Bulk Methods

layers and the layers in Fig. 3 as alternate

building

layers, not only can the crystal structures of all known layered CuO,containing compounds (including all of the known copper-containing high T, superconductors) be constructed,f50) but a great many additional structures may be imagined. The realization of additional high T, structures by bulk synthesis techniques (e.g., mixing powders with a mortar and pestle followed by solid state reaction) has been an active area of research. However, bulk synthesis of such structures can be quite elusive, especially for structures with unit cells comprised of a large number of repeated building layers. The crystal structures of the Bi,Sr,Ca,_,Cu,O,,+, phases that have been synthesized in pure single phase form by bulk methods (n = 1 to 3) are shown in Fig. 10 along with their approximate superconducting transition temperatures (T,). The synthesized structures are the first three members of a homologous series of phases. The members of such a series are related to each other by the addition or subtraction of a simple structural element, which in this case is comprised of a [CuO,] and a [Cal building layer. The next two members of this series are also shown in Fig. 10. The striking empirical trend of increasing T, with the number of CuO, layers in the unit cell of these structures (a similar trend exists for TlBa,Ca,_, CU,O~~+~, TI,Ba,Ca,_, CU,O~~+~, and HgBa,Ca,_, CU,O~~+~ phases) led many researchers to attempt to prepare higher order members by bulk methods. However, as the number of (CuO,),Ca,_, layers in these structures become greater, pure single phase specimens become progressively more difficult to synthesize in bulk form, presumably because the differences between the free energies of formation of these phases become smaller and smaller.P1] Indeed, TEM studies of bulk samples reveal uncontrolled syntactic intergrowths of Bi,Sr,Ca,_,Cu,O,,+, phases when the bulk synthesis of n > 2 is attempted as shown in Fig. 11 ,p*) and of ~*Ba*Ca"-lCU"02"+4 Phases for n > 3.p3] Analogous examples of uncontrolled intergrowths homologous

abound for the bulk synthesis

series, including

the Y,Ba,Cu,+,O,+,,

of other high T, related phases, whose struc-

tures are shown in Fig. 12, for n > 3,p4) La4n+4Cu2n+808n+,4 for n> 3,p5) and Ba, + , (Pb~W,O,, + I for n > 2.r6] Uncontrolled intergrowths also occur in the

T~l+xBa2Ca"-lCuno2"+3+x

system

(a random mixture of TIBa,Ca,_, hases) for 0 < x < 1 .r71 The disorp

and TI,Ba,Ca,_, Cu,O2"+4 dered nature of intergrowths is not limited numerous examples have been documented

CU,O*n+3

of layered oxides.p81

to high T, systems; for other homologous

rather series


523

90 K

n=5 n=4 n=3 n=2

n=l

Figure 10. The crystal structures of the Bi,Sr,Ca,_,Cu,O,,+, phases for n = 1 to 5. The tetragonal subcells are shown for clarity. The approximate superconducting transition temperatures of the phases that have been prepared by bulk methods are also shown.

524

: .-I ”

Molecular

:--_rrr-,_

. _

‘.

_

*

;

_:

Beam Epitaxy

. .

:.

:

:


525

80 K

n=2

n=3

n=5

Figure 12. The crystal structures of the Y2Ba4Cu,+50,+,3 phases for n = 1 to 5. The approximate superconducting transition temperatures of the phases that have been prepared by bulk methods are also shown. n = 2 and n = 4 phases are shown.

Uncontrolled

intergrowths

For clarity, only unit cells of the

appear to be a general feature and funda-

mental limitation of using bulk synthesis methods to prepare complex layered perovskite structures. Bulk techniques rely on the existence of sufficiently deep reaction free energy minima to transform the starting materials into a single phase at a particular temperature and pressure. If

526

Molecular

the formation

Beam Epitaxy

energies

of other phases comprised

of the same building

layers (and therefore the same in-plane lattice constants) are nearly the same as the desired phase, the sample will contain uncontrolled syntactic intergrowths of these structurally related, but nearly energetically degenerate, phases. The increase in entropy of a syntactically intergrown phase, compared to a single phase sample, provides the free energy driving force for such intergrowths. This driving force increases with temperature and as the numerous examples sited above show, causes significant syntactic mixing at the high synthesis temperatures used in bulk synthesis methods. The overall sample composition is insufficient to determine the microscopic layering order when numerous syntactic members have nearly the same free energies; homologous series members containing fewer of the relevant structural building layers will be balanced out by those containing more in a randomly ordered syntactic mixture of these phases. For example, the Bi,Sr,Ca,_,Cu,O,,+, p hases are comprised of [SrO-BiOBiO-901, [CuO,], and [Ca] building layers. If there is no energetic preference for the formation of the n member of this series over the n-7 or n+l members (i.e., AH, of Bi,Sr,Ca,_&u,_,O,,+,, Bi,Sr,Ca,_,Cu,O,,+,, and Bi,Sr,Ca,Cu,+, Ozn+s are identical), then the enthalpy (AH,) of the formation reaction of the n member, Eq. (1)

FrO-BiO-BiO-SO]

t n[CuO,] + n-7 [Cal ---, Bi,Sr,Ca,_,Cu,O,,+,

is the same as that of the enthalpy

(AHJ of a reaction forming a mixture of

the n-7, n, and r-r+7 members from the same reactants, Eq. (2)

[SrO-BiO-BiO-SrO] + +

e.g.,

+ n [CuO,] t n-7 [Cal

1/3Bi,Sr2Ca,,Cu,_, 02n+2 + 1/3&&Can_, CU,O~~+~ l/3 Bi2Sr2Ca,Cu,+,02n+6

Note that the free energy of the latter reaction (AG2) will be lower than that of the former (AG,) because of the increased entropy of the randomly layered mixture of the n-7, n, and nt-7 members. represents just one example of a random mixture

Of course, composition

energetically

In general, all of the

favored over the pure n member phase.

Eq. (2) that is

members of the homologous series may participate and their fractions are free to vary such that the overall equation is balanced. As long as the enthalpies of formation of participating members are sufficiently equivalent, the free energy of the mixture will always be lower than that of the pure phase.


Calculation of the energy of formation of perovskite

related

phases

indicates

of several homologous

that the differences

energy become smaller and smaller as more building into the parent structure

(i.e., with increasing

The low growth temperature widely

utilized

for the

growth

527

series

in formation

layers are inserted

n) .f7gj,+

and atomic layering of metastable

capability

layered

of MBE,

semiconductor

superlattices, have enabled the controlled customized layering of high T, phases whose phase-pure growth is unattainable by bulk synthesis methods. These results demonstrate the capability to grow customized layered structures and metastable phases within oxide systems encompassing the high T, superconductors. The broad spectrum of electrical and optical properties possessed by oxides suggests that engineering these materials at the unit cell level will yield device enhanced properties. 1.5

Layer-by-Layer

heterostructures

with significantly

MBE Growth

The intriguing connection between high T, superconductivity and metastability is a major motivation for developing an MBE technique with atomic-level layering control for the growth of these oxides. By choosing growth conditions which kinetically limit decomposition of the deposited layers, MBE may be used to directly synthesize metastable structures which lie at local energy minima, provided that the activation energy needed to surmount the barriers separating the desired metastable phase from more energetically favored phases (including the equilibrium phases) are sufficiently high. The energy minimum of a desired metastable phase may frequently

be made deeper and the energy minima of the undesired

non-lattice matched phases raised by utilizing a substrate with a suitable template for the epitaxial growth of the desired phase. The second law of thermodynamics requires that the free energy of the entire thermodynamic system

must be lowered

in order for the products

to form from

the

reactants. However, the relatively high free energies of the constituent elements used as reactants in MBE compared to the free energy of the equilibrium products allows, in principle, significantly metastable phases to be formed which lower the free energy of the entire system. The t For example, in the Sr”+,TinOBn+, homologous series of compounds (another perovskiterelated series of compounds), calculations indicate that the formation enthalpy remains essentially constant for n > 2.frQ] Since there is insufficient enthalpic driving force for forming a low entropy phase-pure compound, disordered intergrowths are to be expected in the synthesis of these compounds by bulk methods. Indeed, TEM images of these Srn+,TinOsn+, phasesfeO) show disordered syntactic intergrowths when n ranges from 2 to 8, as would be expected for the bulk preparation of essentially energetically degenerate phases.

528

Molecular

development

Beam Epitaxy

of a theoretical

understanding

of the mechanism

of high T,

superconductivity or at least a framework for predicting which structures should have higher Tc’s, will allow the customized layering capability of MBE to be most effectively focused on this task. Several factors distinguish MBE for the synthesis

of custom-layered

high T, superconductors. Because of the desired nanoscale layering control, the layers must be deposited in a state which does not require any subsequent high temperature processing which could result in intermixing of the layers or which would decompose metastable layer orderings. In addition, low growth temperatures are desired to minimize undesired bulk diffusion between the layers or between the substrate and film. These considerations require a synthesis method in which the deposited layers are crystallized in-situ during growth. In-situ growth exploits the significantly higher surface diffusion coefficients of the depositing species compared to bulk diffusion coefficients, resulting in lower synthesis temperaOf course, abrupt interfaces and tures and thus less bulk diffusion. nanoscale layering control are also requirements of a suitable technique for the customized growth of high T, superconductors. From the standpoint of MBE growth, the layering along the c-axis of these structures represents a superlattice of the constituent building layers; semiconductor superlattices of comparable dimensions are routinely grown by MBE.fs’)f**) MBE is capable of synthesizing new custom-made structures (even metastable ones), provided the growth can be achieved through the deposition of complete, flat, monolayers of the constituent building layers. An idealized view of the desired MBE growth process is shown in Fig. 13.

2.0

OXIDE MBE SYSTEMS

2.1

MBE System

Configuration

The configuration

of an MBE system

for the growth

of high

T,

superconductors differs in several important ways from today’s more conventional MBE systems designed for the growth of semiconductors. The major differences are the requirements to introduce an activated oxidant species, to provide more accurate composition control, and to have adequate pumping to handle the oxidant gas load. A schematic diagram of the growth chamber of an MBE system which contains most of the features commonly shown in Fig. 14.

used in the growth of high T, superconductors

is


529

Figure 13. A schematic representation of the MBE growth of Bi,Sr,CaCu,Os on SrTiO,. The sprayed beams are individually controlled by shutters which control the sequence and quantity of species reaching the growth surface. Note that this figure is highly schematic; the growth unit and growth mechanism of sequentially deposited oxide structures by MBE is not well understood.

530

Molecular

Beam Epitaxy

Atomic Absorption Light In

FombliV Mechanical Pump

Figure 14. A schematic diagram of an MBE growth chamber for the growth of high T, superconductors. The growth chamber shown contains most of the features commonly used: elemental source beams, shutters, activated oxidant introduction, RHEED, and in-situ composition monitors.

The introduction of a reactive oxidizing agent may involve its metered introduction and delivery from an external vessel (e.g., the ozone distillation and delivery system shown) or the generation of the activated oxygen species within the MBE chamber, e.g., utilizing an electron cyclotron resonant (ECR) plasma source. The materials that come into contact with the reactive carefully

oxidizing

agent before reaching

chosen in order to minimize

decomposition

the substrate

must be

of the oxidant.

The

temperature of potential decomposition surfaces should also be considered, as radiant heat emanating from the substrate heater, effusion cells, ion gauges, etc., may heat the surfaces of oxidant delivery materials to temperatures where significant decomposition of the oxidant occurs, even though at room temperature these materials are quite compatible with the oxidant. The amount of pumping necessary to maintain the long mean free path necessary for MBE growth depends on the oxidant used, the material


grown, the growth temperature

and growth

rate, and how efficiently

oxidant is introduced at the substrate. More than 500 I/s of pumping common; some systems have as much as 4680 l/s.ts3) Getter-pumping oxygen by the rare-earth often increases

531

the is of

and alkaline earth species on the chamber walls

the pumping

rate well beyond the active pump capacity

during growth. Because of the high pumping speeds involved, turbomolecular and cryogenic pumps are most commonly used. When working with potentially explosive oxidants such as ozone, the use of cryogenic pumps is a serious hazard, since the adsorbed ozone may detonate when the pump is warmed during regeneration.tB4] Minor differences between the configuration of an MBE system for the growth of high Tc superconductors and a conventional MBE system for the growth of semiconductors include the selection of materials for high temperature components (heater filaments, crucibles, substrate holders, etc.) which do not react or decompose on exposure to the oxidant, and confining the oxidizing species to the substrate region through the use of differential pumping. Fluxes of the constituent metals are generated by resistively heating the constituent elements in effusion cells or by heating them with e-beam sources. When effusion cells are used, it is preferable to use designs which minimize condensation of the evaporating species at the effusion cell orifice, since this gradually reduces the orifice area and thus the flux. This problem can be quite severe for copper, but occurs more gradually for alkaline-earth sources. Just as dual-filament effusion cells have proven effective for reducing gallium droplets which condense at the crucible orifice in Ill-V growth,t85] they have also been successfully utilized to solve the crucible-lip condensation problem in the growth of high T, superconductor materials.t86) Although the use of molybdenum for parts that reach high temperatures is common in semiconductor MBE machines, the high volatility of Mo03t8’] makes the use of molybdenum in oxide MBE systems imprudent. High levels of molybdenum contamination have been found by many MBE researchers in samples grown on conventional molybdenum MBE substrate blockst88)-tg2] or when molybdenum clips were used to secure the substrates to the substrate block.tg3) The presence of molybdenum in films of high T, materials has been revealed by Rutherford backscattering spectrometry (RBS),tB81 secondary ion mass spectrometry (SIMS),tg41 inductively coupled plasma emission spectroscopy (ICP),tgl) electron probe microanalysis (EPMA),fs1)ts31 and even x-ray diffraction (Mog02s).f8sl In extreme cases, the molybdenum content has amounted to several atomic

532

Molecular

percent.

Beam Epitaxy

Naturally,

the occurrence

concentrations adversely films. The conventional Ti-sublimation molybdenum

pumps

of such high molybdenum

impurity

effects the superconducting properties of high T, titanium-molybdenum alloy filaments utilized in

have been replaced

contamination

with

pure titanium

to avoid

from this source as well.fgO)~t Hot tungsten

has also been reported to be a problem in high TcMBE systems, again due to the volatility of tungsten oxides when tungsten was used as a crucible material (for copper). ig6j It is thus important to choose either materials compatible with an oxidizing ambient throughout the MBE machine, or to confine the oxidizing species to the region of these compatible materials. This is particularly crucial for hot parts that are in the line-of-sight of the substrate, since volatilization from such parts can lead to unwanted contaminants in the deposited films. Stainless steeltg7) (e.g., 304,tge) Haynes”fgg) alloy #214,tg4) or Inconel@‘tgl)flw)) substrate holders are commonly used since these materials, as opposed to molybdenum, are compatible with the strong oxidizing ambient which is inevitably present at the substrate position. Platinum has also been successfully employed as a substrate holder.fg2) 2.2

In-situ

Analysis

The high vacuum growth environment of MBE permits the simultaneous use of vacuum surface analytical tools. Reflection high energy electron diffraction @HEED), low energy electron diffraction (LEED), Auger electron spectroscopy (AES), ultra-violet and x-ray photoemission spectroscopy (XPS), scanning electron microscopy (SEM), scanning tunneling microscopy (STM), and other vacuum characterization techniques can be readily applied to the as-grown

surfaces of films.

Several of these

techniques

can be used during growth to gather crucial information

nucleation

and growth

mechanisms

on the

in real time, rather than relying

on

“pathology” after the growth. RHEED is particularly useful in this context. The sensitivity of grazing angle diffraction to surface structure is ideal for monitoring the evolution of film accumulation from initial nucleation to the

t The low mechanical strength of pure titanium at tempertures where its sublimation rate is sufficientto achieve significant pumping speeds should be considered when substituting pure titanium filaments.fs5) If unsupported, the pumping speed achievable with pure titanium filaments will be drastically reduced.fQ5)Molybdenum is present in the titanium-molybdenum alloy (85% li, 15% MO by weight) of conventional filaments to provide sufficient hot strength.tg5) Its presence allows titanium-molybdenum filaments to be operated (unsupported) at temperatures considerably higher than the melting temperature of pure titanium.tg5]


533

deposition of each subsequent building layer. The formation of intermediate reaction products or impurity phases can be readily monitored and the growth conditions phase transitions

adjusted

during growfh.

in these multicomponent

lytical tools are key to the achievement

Due to the many phases and oxide systems, such in-situ ana-

of atomic layer oxide engineering.

The fact that the building layers of the copper-containing high T, superconductors shown in Fig. 3 are not electrically neutral is likely to influence the degree to which the layering can be customized, i.e., the minimum growth units are likely charge neutral.t50)t101) The growth unit is well established for the growth of copper-containing high T, superconductors and other perovskites when all of the constituent species are supplied simultaneously to the substrate in a continuous manner. For such codeposition growth conditions, the intensity of the RHEED pattern oscillates as shown in Fig. 15. The period of these RHEED oscillations corresponds to a deposited thickness of an electrically neutral formula unit of each compound.t 1021For the growth of high T, superconductor films with their c-axis aligned normal to the plane of the substrate, this amounts to one c-axis unit cell thickness for YBa2Cu,0,B and other high T, materials (e.g., (Ba,K)BiO, and TIBa,Ca,_,Cu,O,,+,) whose unit cells contain one formula unit, and one half the c-axis unit cell thickness for doped La,CuO, and other high T, materials (e.g., B&Sr,Ca,_,Cu,O,,+, and TI,Ba&a,_, CU,O~~+J whose unit cells contain two formula units. This minimum growth unit determined by RHEED oscillations is consistent with the step heights that have been observed by STM observations on the surfaces of and single cryst&.[l051[‘061 high T, superconductor filmst 1031[1041 On the other hand, the growth unit of copper-containing high T, superconductors formed by sequential deposition is more difficult to assess. Existing data indicate that the sequential deposition growth unit is at least as small as that of continuous deposition, but determination of the minimum attainable growth unit from RHEED data is complicated by the sequential growth process itself and its determination remains an active area of study. For example, although the intensity of the RHEED pattern oscillates as the incident species are modulated during sequential deposition,t441t881t1071 the period of these oscillations spond to the growth unit. The modulation

does not necessarily correof the incident fluxes itself

causes corresponding oscillations in the RHEED intensity due to the differing scattering factors of the species being supplied in a discontinuous manner to the surface, as well as modulation that may be induced by the changing

of the surface reconstruction

surface species.

Oscillations

at the

534

Molecular

Beam Epitaxy

Time( arb.unit) Figure 15. (a) BaTiO,,

RHEED oscillations observed during the continuous codeposition (b) La,CuO,, and (c) YBa,Cu,O,,. (From Ref. 102).

of


shuttering

frequency

occur during

migration

enhanced

epitaxial

growth of III-V materials,[ loal a process which is completely the sequential growth,

deposition

oscillations

technique

occur

frequency

(MEE)

analogous

used for high T, materials.

at the shuttering

535

to

In MEE

even when

the

number of species being deposited does not correspond to a complete monolayer.[108] In such cases, a beat frequency caused by periodic layer completion, and indicating that an incomplete monolayer is being deposited in each flux burst, may accompany the first several MEE oscillations, but the presence of such beating has been shown to depend strongly on the growth conditions (e.g., it is absent at low growth temperatures) for lllV materials[108] and has not been studied for oxides. Indirect evidence indicates that the growth unit during

sequential

deposition is at least as small as that of continuous deposition. As described in Sec. 3, shuttered MBE has achieved atomic substitution at particular sites within the smallest electrically neutral formula unit (the minimum growth unit for continuous growth), enabling the growth of metastable phases and site-specific doping. Like the minimum growth unit for continuous growth, the minimum growth unit for shuttered growth is lo11 Whether these electrically also likely to be electrically neutral.[ 5o11 neutral building blocks are an electrically neutral combination of building layers[501 or electrical neutrality is attained by varying the oxygen stoichiometry (e.g., CuO instead of CuO,, BiO,., instead of BiO, etc.) of the building layers as they are deposited, with subsequent oxygen reordering upon deposition of a complete unit cell, is an unknown but fundamentally important aspect since the minimum building block size will limit the degree to which the layering of high T, superconductors may be customized by MBE. The sequential MBE growth process used to customize the layering of high T, structures (illustrated in Fig. 13) uses physical means to influence the layering order of the constituent building layers. The MBE shutters are opened and closed to provide bursts of the depositing

species

in a sequence corresponding to the building layer order of the desired material. However, even with perfect control of the quantity of species in a burst, the depositing

species

may agglomerate

into islands

on the

substrate surface, rather than forming the desired flat layer as shown schematically in Fig. 16. In-situ RHEED can be used to monitor this occurrence, and should islanding occur, the growth conditions may be altered in order to search for conditions

conducive

to layer-by-layer

growth.

536

Molecular

Beam Epitaxy


Although

the utility of RHEED during growth is a distinct advantage

MBE over other thin film growth techniques, RHEED

must be kept in mind when

observation

several important

interpreting

of RHEED streaks does not necessarily

RHEED

537

of

aspects of data.

The

imply an atomically

smooth surface.

One-dimensional disorder in the growth direction will also give rise to RHEED streaks.1 1091X-ray diffraction can be used to determine the extent of order in the growth direction, so that the RHEED streaks may be properly interpreted. Secondly, an ideally flat well-ordered surface will appear spotty, rather than streaky, due to the finite curvature of the Ewald sphere. These spots, which appear on arcs delineating the Laue zones, are readily distinguished from the 2-dimensional array of spots that results from the transmission diffraction pattern of a rough surface. Finally, even a surface which appears to be quite flat by RHEED may in reality contain deep pits; as long as the thickness of the flat regions is greater than the electron penetration depth, no transmission diffraction spots will be seen.tllO] As described

in Sec. 2.7, composition

control is a crucial aspect of

the successful growth of high T, superconductors by MBE, especially when sequential deposition is used. Although the freezing-in of metastable layerings at the low growth temperatures employed can be a distinct advantage of the MBE technique, without adequate composition control (on a burst-by-burst

level for sequential

deposition),

the lack of layering

control results in uncontrolled layering disorder.tBs) Customized cation substitution within a chosen building layer (e.g., chemical doping) is possible by sequential deposition MBE, but again this also requires excellent composition 2.3

Minimum

control. 0, Necessary

to Form Structure

A major obstacle to the growth of high Tc superconductors by MBE is to provide sufficient oxygen to form the desired structure, while at the same time maintaining a long mean free path. In order to select appropriate growth conditions, it is desirable to know the lowest oxygen pressure at which the desired high T, superconductor structure may be formed. The thermodynamic stability limits of several high T, superconductor structures have been measured as a function of oxygen pressure and temperature. As Figs. 6 and 7 show, the minimum oxygen pressure necessary to sustain YBa,Cu,0,8 (8 = 1) as an equilibrium phase depends on temperature and sample composition. The lowest oxygen pressures at which the YB~&u,O,~ phase was found to be thermodynamically stable

538

Molecular

as a function

Beam Epitaxy

of temperature,

samples with a stoichiometric

reaction

dence of the minimum oxygen pressure CUn02n+4

as equilibrium

tions: Bi2Sr2CaCu20,+B

IX in Fig. 7, were

obtained

1:2:3 cation ratio. The temperature necessary

to sustain Bi,Sr,Ca,_,

phases has only been measured and Bi2Sr2Ca2Cu,0,,+B

Fig. 9. It may seem strange to be concerned

for

depen-

at two composi-

(6 CI 0), and is shown

with thermodynamic

in

stability

constraints for a process whose goal is to enable customized layering and the synthesis of metastable phases. The ability of MBE to provide customized layering arises from the low bulk diffusion coefficients at the growth temperatures used, which kinetically limit decomposition of the deposited layers. On the other hand, the growth temperatures must be high enough that the depositing species have sufficient mobility, at the film Since reactions, surface, to crystallize in a highly ordered manner. including oxidation, between the species reaching the substrate from the various molecular beams occur at the film surface, where surface diffusion is high and kinetic barriers are minimal, thermodynamic constraints are important. Thus, while the formation of each surface layer is significantly dependent on thermodynamics, it is the aim of the sequential deposition MBE process to preserve the customized layering order by limiting bulk diffusion. The measurements of the minimum oxygen pressure necessary to stabilize the high 7, superconductors YBa2Cu,0,_b, Bi2Sr2CaCu20,+8, and have been made under conditions where the transforBi,Sr,Ca,Cu,O,,+, mation kinetics of the bulk samples studied were sufficiently rapid that the thermodynamic stability limits of these phases could be determined.[661[671[701[11’1Note, however, formed at pressures

considerably

that these

experiments

were

higher than the MBE regime.

per-

Only a few

measurements have been made at lower pressures, some of which extend into the MBE pressure range,tg6tt1 12)[113)but with significantly decreased measurement precision. Thus, estimating the equilibrium stability limits under MBE conditions

from currently

available

data involves

extrapolating

the higher pressure results to lower pressure or fitting thermodynamic functions to the many phases present in these systems which are consistent with the experimental data obtained at higher pressures, and calculating the phase relationships at reduced pressure, as has been performed by Degterov and Vor0nin.t 64I[1141The minimum oxygen pressures necessary to thermodynamically stabilize YBa,Cu,O,_,, Bi,Sr,CaCu,O,+,, and Bi,Sr,Ca,Cu,0,,+, as a function of temperature, extrapolated into the


MBE regime, are shown in Fig. 17. Although the different crepancies

copper-containing between researchers

indicate that high oxygen stabilize relevant

the structures growth

10-l

c 3 e

high

pressures

there is some scatter among

Tc superconductors,

investigating

a particular

(by MBE standards)

of copper-containing

539

high

as well as disphase, the data are required

to

T, superconductors

at

temperatures.

900 800

700

600

400 “C

500

1o-2 1o-3

5

1o-4

3

1o-5

k

1o-(j

5 ~

10“

p

1o-s

s

LO2& lo3

10” k E 2 E

0 0.8

&

1.0

1.2

1.4

1.6

1000/T(l/K) Figure 17. The thermodynamic stability lines extrapolated into the MBE regime from Figs. 7 and 9 showing the minimum molecular oxygen pressure necessary to sustain the structures of YBa,Cu,O,d, Bi,Sr,CaCu,O,, Bi,Sr,Ca,Cu,O, s, and CuO (as equilibrium phases) as a function of temperature. The mean free path of Ba in molecular oxygen, from Fig. 18, is also shown to indicate the regime of MBEcompatible pressures. The minimum molecular oxygen pressure stability lines extrapolated into the MBE regime from Refs. 111 (dashed) and 114 (dot-dashed) for YBa&u,O,, are also shown.

The thermodynamic oxygen stability limit of the only copper-free high T, superconductor, (Ba,K)BiO,, has not been measured. However, MBE growth studies have shown that molecular oxygen pressures at the limit of

540

Molecular

Beam Epitaxy

the MBE regime (this limit is discussed

in Sec. 2.4) result in an extremely

low bismuth peratures

sticking coefficient, even at the low (345°C) substrate temused.f 1151 Activated oxygen species are a common feature of

the low-pressure growth methods successfully (Ba,K)BiO, films.f116)f117)

employed

for the growth of

From these considerations, it is not surprising that the utilization of more reactive oxidant species is a crucial aspect in the successful growth by MBE of high T, superconductors with respectable transport properties. 2.4

Maximum

0, Satisfying

MBE Mean Free Path Constraint

As the thermodynamic considerations in the previous section show, significant oxygen pressures are necessary to make high T, superconductor structures stable. On the other hand, in order to grow these materials by MBE, low oxygen pressures must be used. A necessary condition for MBE is maintaining a mean free path for each depositing species which is longer than the source to substrate distance. This distance is about 20 centimeters in a typical MBE machine. While it is possible to grow high T, superconductors by physical vapor deposition at oxygen pressures exceeding these limits,f281f2g1f1 12)f1201-f127),f these high pressures can significantly complicate composition control[g2~[g~[1~21~~2sland in-situ RHEED characterization. For example, due to frequent gas phase collisions at high pressures, as the evaporating species make their way toward the substrate, the flux of each constituent species actually reaching the substrate is strongly dependent on the oxygen pressure, which not only varies considerably with position, but also depends on the fluxes of any oxygen gettering species which are being simultaneously evaporated. The differing

cross-sections

cause the fluxes attenuated

and momenta

of the constituent

unequally,

causing

species

of the evaporated reaching

a shift in the relative

species

the substrate mixture

to be

of arriving

species, which results in a strong interrelationship between the oxygen flow rate, the individual fluxes, and the depositing composition. Growth at MBE-compatible

pressures removes the interdependencies

in this compo-

sition control problem.

+ By definition,tlle] such a deposition process is no longer MBE, since well-defined thermal molecular beams are no longer present at the high pressures used. Deposition under such conditions is commonly referred to as “reactive evaporation.” Only synthesis methods resulting in the epitaxial growth of a material from the reaction of well-defined thermal molecular beams with a crystalline surface (MBEt”s] or reactive MBf~‘lg))are described in this chapter.


Figure 18 shows the variation

541

of the mean free path of many of the

constituent species of high T, superconductors, calculated for a random gas mixture of one such species with molecular oxygen. The appropriate formula for this calculationf 12s) has been corrected for the one diatomic species which was used (oxygen) and is given by the equation:’

where L is mean free path, Fis depositing flux, d is diameter, m is mass, kB is Boltzman’s constant, P is pressure, and T is absolute temperature. The subscript i refers to the beam species, and the subscript 0, to molecular oxygen. The parameters used in this calculation are given in Table 1. The shortest mean free paths are those of Rb, Hg, and Ba, while Cu has the longest. The variation between the mean free paths of the various high T, superconductor constituents is about a factor of three in the region where their mean free paths are limited by scattering with 0,. Of the species with the shortest mean free paths, Ba is the most commonly used and for this reason the maximum oxygen pressure defining the MBE regime in Figs. 17 and 19 is for a depositing Ba flux. For the MBE growth of high T, superconductors containing Rb (e.g., (Ba,Rb)BiO,) or Hg (e.g., currently the highest Tc superconductor with T, = 133 HgBa,Ca,Cu,Os+,, K),f13fl the MBE regime extends only to about 1 x 10e4 torr of oxygen. As the oxygen pressure decreases, the mean free paths increase (L-l 0: Po2) until the mean free paths become limited by scattering between the species themselves (this is the reason for the change in slope of the mean free paths in Fig. 18 at low Po2). The minimum molecular oxygen structure of several copper-containing

pressure necessary to sustain the high 7, superconductors and CuO

(the thermodynamic stability lines from Sec. 2.3), is shown in Fig. 17 as a function of temperature, together with the mean free path of Ba in molecular oxygen. In order to have a mean free path B 20 cm (the typical source to substrate pressure

distance

in an MBE machine),

is about 2 x 10e4 torr.

the maximum

0,

operating

In order to be above the CuO line, the

t The beam nature of the depositing species was ignored. A depositing flux of 1 014atoms cmm25-l (which corresponds to a growth rate of 0.15 metal-oxide monolayers per second or about 0.3 A/s for YB~&u~O~.~ or Bi2Sr2CaCu20e), the atomic diameters from Ref. 134, the 4 diameter from Ref. 135, and typical furnace temperatures (listed in Table 1) for each species and room temperature for oxygen were used in the calculation. Only atomic depositing species were considered @ii, which accounts for about40% of the species in the bismuth molecular beam/ss) was disregarded).

542

Molecular

Beam Epitaxy

Table 1. Parameters Used in Mean Free Path Calculation


543

growth temperature must be lower than 570°C. Although from a thermodynamic standpoint, the MBE growth of copper-containing high T, superconductors utilizing 0, is possible at temperatures lower than about 570°C to 605°C

depending

on the particular

such low growth temperatures

compound,t67)t70) kinetic constraints

at

are likely to make this a futile avenue

to

growing high T, films at reasonable growth rates with acceptable transport properties. Studies using MBE-compatible pressures of molecular oxygen have achieved a non-stoichiometric DyBa,Cu,07a-related phase, which may be made superconducting by a subsequent low temperature (400°C) oxygen anneal.tssl However, its poor superconducting properties and seemingly inherent non-stoichiometric nature indicate the impracticality of the use of molecular oxygen for the MBE growth of high T, superconductor films.tg6) Instead, the use of more powerful oxidants and the effective use of differential pumping to enhance the oxidant pressure at the substrate (where it is desired) and decrease it elsewhere in the chamber, are utilized to grow high T, superconductor by MBE.

films with useful transport

characteristics

10 lo5 lo4 lo3 lo* 10’

Metal Flux = 1~10’~ atoms/(cm2 s) 1(p

’ “,“**’ * ““‘_’ ’ ’ 1’.1”’ ’ “..“.I ’ “‘.a”’ * c”l”” ’ “*“**’ a “‘11 lo-* lo-’ 1o-6 1o-5 1o-4 1o-3 1o-2 10-l

Oxygen Pressure (Torr) Figure 18. The mean free path of atomic fluxes of the constituent T, superconductors and the parameters

as a function of oxygen in Table 1.

pressure.

Calculated

species of high using Eq. (3)

544

Molecular

2.5

Alternative

Beam Epitaxy

Oxidants

From thermodynamic

considerations,

have been identified.i881[1381[1~l assume

equilibrium

between

H owever,

all possible

potential

alternative

thermodynamic species

oxidants

calculations

and are not directly

applicable to the deposition conditions used in this reactive MBE approach for the growth of high T, superconductor films, where an enhanced, nonequilibrium concentration of an alternative oxidant is utilized to enable growth under MBE-compatible vacuum conditions. An indication of the inefficiency of 0, as an oxidant for these materials is provided by considering an ideal oxidant: one having high activity and no kinetic barriers to oxidation. Since even the most powerful oxidant must deliver an oxygen flux at least as great as the oxygen content of the crystal structure to be formed, the required minimum flux (and equivalent pressure+) may be calculated assuming complete incorporation of the oxidizing species for the desired growth rate of a given high T, superconductor phase. For example, the minimum flux necessary to form YBa,Cu,O, at a growth rate of 1 &s for an oxidant that provides one oxygen atom per oxidant molecule is 3 x 1014 cm-2s-1, which corresponds to a pressure of about 3 x 10W7torr. This minimum flux is more than three orders of magnitude lower than the minimum 0, flux (or pressure) necessary to stabilize the growth of YBa,Cu,O, at typical growth temperatures which yield good transport characteristics (i.e., T 2 600°C). As Fig. 17 shows, the minimum amount of oxygen required to thermodynamically stabilize the high T, structures rises rapidly with temperature. By 750°C the performance advantage of an ideal oxidant over molecular oxygen is about five orders of magnitude in pressure. The inefficiency of molecular oxygen leaves considerable room for oxidant pressure reduction well into the MBE range by the use of more reactive oxidants. An estimate of the increased oxidizing power of various alternative oxidants for copper-containing gated

by measuring

high T, superconductors

the minimum

oxidant

flux

has been investi-

needed

to form

CuO

t Because of the beam nature of the incident fluxes, the pressure at the plane of the substrate corresponding to the incident flux is given by 4-L case II

nmikBTi 8

where Ois the angle between theincident beam and the substrate normal, and Pi, 5, ml, kg, and T were defined previously (see Eq. 3).[ 1401[1411 This differs by a factor of 4cos.B from the commonly used formula Pi - FJm, random motion.[’ 421

which is appropriate for a gas having completely

High

compared

to the minimum

0, requirement.

ents of copper-containing

high

highest

pressure

molecular

oxygen

Tc Superconductors

545

Of the binary oxide constitu-

T, superconductors,

CuO requires

to be thermodynamically

the

stable,fs8t

whereas rare-earth and alkaline earth constituents are readily oxidized by molecular oxygen at MBE-compatible pressures.[171[143]-[146] Studying the ability of alternative

oxidants to oxidize Cu to CuO and comparing

it to the

ability of 0, to do so has several advantages over studying their ability to sufficiently oxidize specific copper-containing high T, superconductors. First of all, forming CuO is free from the composition control problems which accompany the synthesis of multi-component copper-containing high T, superconductors. Note from Fig. 6 that the minimum Po2 necessary to thermodynamically stabilize a sample with composition Y:Ba:Cu = 1:2:3 (i.e., YBa,Cu,O,) is significantly different from the Po2 requirement for a slightly copper-rich composition, e.g., Y:Ba:Cu = 1:2:3.5. At 850°C the former is thermodynamically stable down to Po2 = 300 mtorr, while the latter will decompose

below

Po2 = 760 mtorr.f66) Secondly,

forming CuO also gives a more definitive measure of the enhanced oxidation ability of the oxidant, since the minimum oxygen requirements of CuO are thermodynamically well established, while there remains considerable disagreement about the precise minimum 0, stability lines for the copper-containing high T, superconductors at low pressures (e.g., YBa,Cu,07~f671f1 l ll[l 141[1471 and Bi,Sr,CaCu * 0 Q+s17Q1f’4Ql)~ In order to compare the ability of alternative oxidants to form cupric compounds, it is necessary to (i) produce a pure beam of a specific oxidant or of the oxidant diluted in O,, since the oxidation behavior of OQ is understood, and (ii) know the flux (or pressure) of the oxidant at the substrate position itself. Although researchers have extensively utilized plasmas to enhance oxidation, relatively few studies have produced oxidant beams of the requisite purity to investigate the activity of specific oxidant species and quantified their flux at the position where oxidation occurs. Such studies, although over a rather limited range of experimental conditions, have been performed for atomic oxygen, O+, ozone, and NO,. In these studies, the ability of atomic oxygen (0), O+, ozone (O,), and NO, to oxidize copper has been measured by depositing copper onto a substrate in the presence of one of these oxidants.f831f12Q)-f1321The formation of CuO (or Cu,O or Cu) has been monitored by in-situ RHEED,f83)f88)f12Ql in-situ optical reflectance,1 lQQl in-situ x-ray photoelectron spectroscopy ()(pS) ,PslUQQlf’s21and ex-situ x-ray diffraction.f83)fQ4)f12Q)f131) When the formation of CuO is tested by ex-situ methods or after the samples have

546

Molecular

Beam Epitaxy

been cooled, the atmosphere

in which the samples

are cooled must be

carefully considered.+ The limited data in the literature that meet these criteria are shown in Fig. 19. neously

In these exposed

experiments, to an oxidant

a heated

MgO substrate

beam and a copper

was simulta-

beam.ts31t12g]-t132]

Growth conditions yielding CuO are shown as solid points and those resulting in Cu,O are shown as hollow points for the various oxidants. In addition to the data for atomic oxygen, O+, ozone, and NO, plotted in Fig. 19, the thermodynamic equilibrium line denoting the thermodynamic limit to CuO formation with molecular oxygen is also shown. All of these oxidants are significantly better than molecular oxygen, and allow the growth of cupric compounds to be achieved at MBE-compatible pressures. The flux needed to form CuO with O+ is over one order of magnitude lower than the thermodynamic 0, requirement at 510”C.t83~ The flux of NO, used to produce CuO is about two orders of magnitude lower than the thermodynamic 0, requirement at 65O”C.t 1311t132]The pressure needed to form CuO with purified ozone is over two orders of magnitude lower at 600°C and over three orders of magnitude lower at 700°C compared to the equilibrium line for O,.t rs0ltls1l The required pressure for atomic oxygen is over four orders of magnitude lower at 650”C.[12g] This limited experimental data is insufficient to conclude which of these oxidants is the most effective. In all cases, the minimum oxidant fluxes found capable of producing CuO in these experiments approached the minimum amount, based on purely kinetic grounds, that an ideal oxidant would require for the copper fluxes used in these experiments. Thus, these powerful oxidants are nearly ideal oxidants. Side-by-side comparisons of NO, with ozone led Ogihara et al.t1311 and Nonaka et al.f132] to conclude that NO, is slightly less reactive than 0,. Ogihara et al.f131] found that the minimum oxidant pressure needed to form CuO was about five times lower for ozone than for N02.t131] Nonaka et al.f132] attributed the decreased

reactivity

of NO, to the known reaction

in which

t The cooling of a sample in tie same oxidant flux in which it was grown (a horizontal cooling path in Fig. 19) results in an increasing driving force toward oxidation. For example, an as-grown Cu,O sample cooled at the same oxidant flux in which it was grown will eventually cross the Cu,O/ CuO stability line. When this occurs, unless the kinetics of the Cu,O + CuO transformation are sluggish, the sample will transform into CuO and the ex-situ determination of the minimum oxidant pressure necessary to form CuO will be in error. In-situ x-ray diffraction analysis has shown thatthistransformation becomes kinetically limited at temperatures below about SCI”C.~‘~~) Thus, when determining the CuO/Cu,O stability line by ex-situ methods, it is preferable to cool the samples in vacuum rather than to cool them in the same oxidant pressure in which they were grown. This provides a lower-bound of the oxidation enhancement provided by the oxidant.


900 800

0.8

700

1.0

400 “C

500

600

547

1.4

1.2

1.6

1000/T (l/K) Figure 19. A comparison of the enhanced ability of atomic oxygen (triangles), 0+ (squares), ozone (diamonds), and NO, (circles) to form cupric oxide compared to molecular oxygen. A copper flux was oxidized to CuO (solid points) and Cu,O (hollow points) by simultaneous exposure to the indicated pressure (flux) of these oxidants. The 0, line indicates the thermodynamic line of coexistence of CuO with Cu,O as a function of molecular oxygen pressure. Above the line, CuO is the equilibrium phase, while below the line, Cu,O is the stable phase. The right hand axis shows the variation of the mean free path of Ba as the oxygen pressure is varied. Experimental data from Refs. 83, 129-132. Thermodynamic line calculated from Ref. 133.

one of the decomposition reduces it:* Eq- (4)

products

of NO,,

NO, reacts with CuO and

NO + 2CuO + ‘hN2 + 0, t Cu,O.

The minimum flux (or pressure) of oxidant depends on the growth rate used since, by conservation of mass, at least one oxygen atom is required per * NO2 begins to decompose into NO and 0 at temperatures above 150°C; this decomposition is complete at temperatures above 600”C.[1 xl

548

Molecular

Beam Epitaxy

deposited copper atom.

The increase in required

oxidant

flux at higher

temperatures may reflect the oxygen loss rate from the film surface. For CuO films, XPS studies have found that the onset for oxygen loss at the sample surface is about 65O”C.t 12QlAt higher temperatures, the oxidant flux must not only provide sufficient oxygen to oxidize the incident copper, but must also supply sufficient oxygen to make up for that lost by vaporization. Practical issues also affect the choice of oxidant. Although several oxidants are suitable for forming high 7, structures under MBE conditions, only in the cases of NO,, O+, and ozone have relatively pure beams of the oxidants been achieved. A disadvantage of plasma sources to produce active oxygen species for the MBE growth of high 7, compounds is the high pressure of molecular oxygen which accompanies the desired oxidant. This unwanted O,, which typically amounts to 90% or more of the gas f~ux,[44[‘~21[‘s01-[‘521increases the gas load and decreases the mean free path. High purity, mass separated, low energy O+ ion beams have been achieved and their extreme activity has led to the formation of new oxide structures.t153) However, the O+ fluxes are relatively low (typically 3x10140+cm2 s-)l ,[153l and the capital investment necessary is relatively high, compared to the use of NO, or purified ozone beams. The purity of the latter is estimated to be between 26% and 70% ozone, the remainder being 02.t154]-f157] An additional advantage of NO, and ozone is that they may be supplied from a vessel outside of the MBE chamber, and delivered in a very pure form to the substrate region using conventional ultra-high vacuum (UHV) materials like stainless steel,t158~t15Q~ whereas atomic oxygen and oxygen plasma species are much more sensitive to recombination. Note that XPS studies do not indicate the incorporation of nitrogen into CuO, NdBa,Cu,O,,, or Bi,Sr,CuO, films when NO, is used as the oxidant at the growth temperatures

employed.t132)f161)~ t

Even when copper is deposited ture, simultaneous

MBE-compatible

onto a substrate

at room tempera-

fluxes of atomic oxygen,

and NO, are capable of oxidizing the depositing CuO ts31tsQlt’29lt’sOIt’s indicating the absence of kinetic oxidation

for these

reactive

perature.

This behavior

oxidants

at temperatures

is distinct from that of molecular

O+, ozone,

copper into barriers to Cu

above

room tem-

oxygen.

Kinetic

barriers to the oxidation of copper by molecular oxygen exist, as was demonstrated by the inability of 0, to oxidize copper at MBE-compatible fluxes at temperatures below 450”C,~Q4~~13Q~~157~~162~~163~ even though these

f Nitrogen incorporation, via the formation of Sr(NO&, was detected by XPS in Bi,Sr,CuOs layers depsited at room temperature, [lscl however, after heating to 300°C the films were free of nitrogen.f611


549

fluxes are, from a thermodynamic standpoint (Fig. 19), sufficient to form CuO. This kinetic barrier to the oxidation of copper using molecular oxygen is absent by 55O”C,t g6It112) but by that point, the pressures of molecular oxygen pounds are quite discussed molecular

required to thermodynamically stabilize cupric comclose to the limit of MBE-compatible pressures, as

above. Kinetic barriers also exist to the oxidation of bismuth by 0xygen.t ler) Since all of the known high T, superconductors

contain copper or bismuth, the use of an activated oxygen source for the MBE growth of these materials is seen to be a necessity. Although the oxidation sources considered and used to date for the MBE growth of high 7, superconductors are gaseous, solid source oxidants (e.g., Sb,O, and As,O,) have been used to grow other oxides at This chemical very low pressures/ 1~ well within the MBE regime. method of oxidation has the advantage of not requiring an expensive oxygen plasma source or the collection of potentially explosive purified ozone. However, the surface chemistry of growth is complicated by the presence of additional species that must be desorbed in order to grow oxide films free from the unwanted components of the particular molecular oxidant used. Early in the development of the MBE technique for the growth of high TC materials, it was recognized that more reactive oxidants, such as ozone, were desirable.t1431 The ability of ozone to grow YBa,Cu,O,, and Bi,Sr,CaCu,O, under entirely MBE conditions (including substrate cooling after growth) was subsequently demonstrated.t1661t16~ For the reasons outlined above, ozone has become the dominant oxidant of choice for the MBE growth of high T, materials; high quality films of YBa2Cu30,,,~g2~[132~~1661 DyBa2Cu30,~,t6g)~go)t16g) SmBa2Cu30,~,t16g1 and are now routinely grown with ozone. Bi2Sr2Ca,_,Cu,02,+4 [1711[1721 NO, has not only been used to grow YBa,Cu,O,,,

NdBa,Cu,O,+

and Bi,Sr,CaCu,O, films at typical substrate temperatures used for the growth of copper-containing high T, superconductors, Tsub 700°C t131)t173)t174) but has also enabled the growth of superconducting DyBa,Cu,O,, films at temperatures as low as 420”C.t156] Epitaxial growth of the Bi,Sr,Ca,_, CU~O~~+~p hases (n = 1 and 2) has been achieved using NO, at temperatures as low as 3OO”C.t 160)t161)t175)These are the lowest growth temperatures at which the successful epitaxial growth of coppercontaining high T, materials has been reported. An important aspect of this low temperature achievement was the sequential supply of the depositing species (the metallic elemental constituents, as well as the oxidant) on a monolayer by monolayer basis by shuttered MBE to the growth surface.

550

2.6

Molecular

Beam Epitaxy

Ozone System As discussed

in Sec. 2.5, more powerful

oxidants

oxygen are necessary to grow high T, superconductors tions, and ozone is a particularly

effective

than molecular

under MBE condi-

and technologically

convenient

choice. In order to achieve growth conditions having the longest possible mean free path, it is desirable to deliver a pure beam of a strong oxidant to the growing surface. Less oxidizing species (i.e., molecular oxygen) result in a decreased mean free path without providing significant additional oxidation. Commercial ozone generators produce only a few percent ozone in the oxygen gas that passes through them. Since the vapor pressure of oxygen is much greater than that of ozone, it is possible to purify ozone by distil1ation.t 1761-t176)A typical ozone collection, distillation, and delivery system is shown in Fig. 20.tg4t

/-\\

./’

‘.

I, ___--___-Vented Explosion Containment Box

Figure 20.

A schematic (after Ref. 94).

diagram

Mechanical Pump

of an ozone

distillation

and delivery system

The heart of the system is the ozone trap which stores the ozone. A cross-sectional diagram of one example of such a trap is shown in Fig. 21 .tW) Ozone

is trapped

in a silica

gel~66~~g0~~112~~121~~1761 in the design


/

Flexible Bellows Ozone Input

I

-1

551

Thern$;yple Glass I Metal

r

Dewar

Figure 21. A schematic diagram of the glassware containing silica gel used to store ozone (from Ref. 94). shown; other designs collect the ozone as a liquid.~84~[g1~~g2~[162~[166] The use of a silica gel to contain ozone has safety advantages over the collection of liquid ozone. Liquid ozone is far more sensitive to changes in temperature,

pressure,

vibration,

etc. than ozone

adsorbed

in a silica

gel.I1761 When using a silica gel, ozone can easily be collected at temperatures higher than the liquefaction temperature of ozone. Indeed, this should always be done in order to pretent the accumulation of liquid ozone.Ig41[121] This can be conveniently accomplished by using dry ice (solid CO,) as a coolant for the ozone trap, since the sublimation point of CO,, 195 K (as opposed above the liquefaction

to the boiling

temperature

point of nitrogen,

of ozone decomposition increases with temperature, at of ozone stored in a stainless steel vessel is more than stainless steel is phosphated, the half-life exceeds indicating that the rate of ozone decomposition is practical

standpoint.

To minimize

77 K), is well

of ozone (161 K) ,[1211Although

decomposition,

the rate

195 K the half-life a week (and if the four months),[17g] negligible from a

the ozone should only

552

Molecular

be allowed

Beam Epitaxy

to come into contact

with compatible

materials,

e.g., silica

ge1,t176tglass, Teflon”to 1801(and the related material Kel-F@t181)), stainless steel, aluminum, and alumina.f15g) All parts should be thoroughly

cleaned in

orderto eliminate organic residues that could later trigger an explosion.t178)t182) Although

at room temperature

the materials

listed above

may be

used to deliver a reasonably pure ozone flux to the substrate, at higher temperatures significant decomposition of the ozone may occur. Radiant heat from the hot filaments inside the MBE can lead to a significant temperature rise of the ozone-ducting within the MBE. For example, the orifice of a stainless steel ozone delivery tube reached Al 130°C during MBE growth when the tube was not cooled with water.t131) It is for this reason that the stainless steel tubing used to introduce ozone is watercooled in the high T, MBE systems that have demonstrated the most effective oxidation at the minimum ozone pressures.tg0)t130]f1311 2.7

Composition

Control

Achieving accurate composition control is a crucial aspect for the controlled growth of high T, superconductors by MBE. Compared to the growth of conventional semiconductor materials by MBE, the growth of four and five element superconducting compounds requires far more stringent incident flux control. For example, stoichiometric GaAs grows relatively easily over a substantial substrate temperature range and As,:Ga flux ratio range where gallium has a near unity sticking coefficient and excess arsenic does not stick. In the growth of high T, superconductors, the three or four metal components typically all have significant sticking coefficients. Thus, deposition uniformity across the substrate and shutterinduced source temperature

transients

growing film as well as the uniformity

affect both the stoichiometry of film thickness.

Experimental

of the data

on the effect of composition deviation on the superconducting properties of thin films t1s*1trss) as well as statistical considerations of the effect of composition’errors on the chemical doping of high T, superconductors (Fig. 4),tg4] indicate that cation composition control at or better than about + 1 atomic percent is necessary to produce quality films of high Tc superconductors. Off-axis sputtering and PLD have a distinct advantage over MBE in that accurate and reproducible composition control is built into these single-target deposition techniques. The inherent flexibility of the MBE technique is only advantageous when it is accompanied by adequate composition control; without it, the user will be unable to control the formation of the desired custom-made structures and controllably dope them. The relatively late arrival of MBE-grown high T, superconductor


553

films with respectable

and reproducible transport properties is in large part due to the inadequacy of the composition control used in the initial work, and the time that it took to develop and implement more precise composi-

tion control systems. When shuttered MBE growth is employed, relative control of the depositing species to f: 1% is no longer sufficient: the total number of atoms incorporated in each flux burst must be an integral number of monolayers, making absolute composition control a necessity. Reproducibly achieving such a level of composition control by MBE is a considerable challenge, which is only beginning to be met. Several forms of flux measurement are applicable. Ion gauges,[171t1541t154] quartz crystal microbalances,~~~1~4~[861[‘~4l~~~51~~5~1~~~~1~~~~1~~s71~~741~~*51-~~s~1 mass spectrometry,fg01tg21t1261~1271t1521t1g01-t1g21 electron impact emission spectroscopy (EIES),[188][1g3]-1105]cold cathode emission spectroscopy (CCES) ,PlWl and atomic absorption spectroscopy (AA)t’ 121t1081-t2021 are widely used techniques for thin film composition control, and all of them have been used for composition control for the in-situ growth of high Tc superconductors by evaporative means. Unfortunately, the techniques utilizing hot filaments, i.e., ion gauges, chopped ion gauges,f1411t203] mass spectrometry, and EIES are difficult to implement during deposition due to the highly oxidizing conditions necessary for the growth of high T, superconductors. For oxygen pressures above about 1 0m5to 10e4torr, composition control using these hot filament techniques becomes unsatisfactory.[02~[112~[1s*~[1g1~[1g4~ Although the time-constant of fluctuation of resistively-heated MBE effusion cells is quite long, an advantage for precise composition control, deviations of the flux by 10% or more are not uncommon over the several hour duration of a typical growth.t861tg4tt2041 Even larger changes in the fluxes (e.g., for alkaline earth sources) can occur after they are exposed to the high oxidant pressures that accompany growth, due to oxidation of the source materials and the consequent control changes in source emissivity.t 2051 Thus, precise composition methods that are able to monitor the fluxes during growth are desirable for the MBE growth of high T, superconductors. Recently, CCES was used to control yttrium, barium, and copper during growth at oxygen pressures up to 10e4 torr.ng7) Atomic fluorescencet 2061 might also prove useful for accurate flux control during growth. An advantage of the use of quartz crystal microbalances is their ability to measure the active oxygen flux (atomic oxygen or ozone) when they are coated with silve61551t207] or copper_[139[2081 Even the most promising methods, e.g., AA and cross-beam quadrupole mass spectrometry, which have achieved better than + 1% composition

554


control in the MBE growth of high T, superconductors,fQ01f152~f201~t202] require an absolute composition measurement to calibrate the AA or mass spectrometer

signal for the growth conditions

plished by growing calibration each cation/cm*)

is measured

used.

This is often accom-

films whose absolute composition by an ex-situ

method,

(atoms of

e.g., RBS, ICP, or

direct coupled plasma emission spectroscopy (DCP). However, the measurement error of these common composition analysis techniques typically is in excess of the + 1% absolute composition control typically needed for the reproducible growth of high T, superconductors with excellent electrical properties. The error of EPMA is about + 5%,t20Q1t2101 and can be significantly greater for films thinner than about 0.3 pm, at which point the substrate becomes part of the excitation volume and the assumption that EPMA is probing a homogeneous composition is no longer satisfied.f *l ‘If1 ‘*I Computational routines for heterogeneous samples which allow the presence of the underlying substrate to be included in EPMA composition analyses do exist,f*l*] but have yet to be widely implemented in commercially-available EPMA systems. In addition to the comparatively high error of EPMA composition analyses, it is insufficient for calibrating the MBE growth of high T, superconductors by sequential deposition since it measures only the relative sample composition and not the absolute quantity of species present. Ideally, absolute composition measurements can be performed insitu, eliminating the need for, and delays involved with, the growth of calibration films and ex-situ analyses. The quartz crystal microbalance is capable of absolute in-situ flux measurements. Although there are several limitations to its use, particularly in an oxidizing environment where the oxidation state and thus the mass of the depositing sensitivity widespread

species is uncertain,

and ability to make absolute flux measurements

contribute

its

to its

use in MBE systems for the growth of high T, superconduc-

tors. Even when another composition control method is used during growth, a quartz crystal microbalance that can be moved into the substrate position (thus avoiding any tooling factors which might change with time) is frequently

used to calibrate

the composition

control

system.

In MBE

systems where no flux monitoring is performed during growth, calibration with a quartz crystal microbalance is often performed prior to and following each growth. The main disadvantages of quartz crystal microbalances are the uncertainty in the oxidation state (and thus uncertainty in the flux) of species deposited in an oxidizing ambient, its sensitivity to temperature changes, the uncertainty in the acoustic impedance of the various layers


deposited

on the crystal, and its low bandwidth

bandwidth

is not a problem for effusion

as a flux monitor.

555

The low

cells, but is a concern for e-beam

sources. Accurate composition control is important not only for attaining optimal carrier concentration in the CuO, layers during coevaporated growth,

but is also crucial for the controlled

substitution

and doping

of

cation sites when shuttered MBE growth is used. For example, the layering control of MBE may not only be exploited to make metastable layering sequences of nominally pure building layers, but also may be used to control the site substitution (cation mixing) and doping with subunit-cell precision. The ability to customize cation mixing within high T, structures would be a powerful synthesis tool for probing the interrelationship between site doping and superconductivity in these materials. Bulk Bi,Sr,CaCu,O, crystals typically have a (Sr+Ca)/Bi ratio of 1.3 rather than the idealized Bi,Sr,CaCu,O, formula value of 1.5 and significant mixing of the cations between the various sites.f53)[54) The best Bi,Sr,CaCu,O, have also been produced by shuttered MBE with a (Sr+Ca)/Bi

films ratio

significantly less than 1.5.t *13] Depositing pure as well as doped building layers by sequential MBE growth requires accurate absolute composition control. Less stringent composition control is readily tolerated in the growth of equilibrium structures where phase separation may be used advantageously to produce a connected path of the equilibrium phase. Composition control, rather than oxidation, is the major remaining challenge to be overcome in order for the controlled growth of high T, superconductors and related phases to become a more controlled and reproducible synthesis tool. 2.8

Crucibles If the oxidizing

environment

necessary

for the MBE growth of high T,

superconductor materials is not adequately confined to the region of the substrate, it may affect the performance of susceptible MBE components. In addition to hot filaments, the performance of pyrolytic boron nitride (pBN) crucibles

can degrade.

This degradation

is particularly

important

when a molten species is contained in a pBN crucible. Oxidizing species react with hot pBN to form N, and B20,.fg4)fg6j The formation of B,O, was established based on studies in which a mass spectrometer was positioned in a line-of-sight position with a copper charge contained in a pBN crucible.fg4)fg6t After oxygen exposure

(e.g., Po2 = 2 x 10m4torr), not only

556


was there a significant

decrease in the copper flux (as much as a factor of

three lower), but the mass spectrometer detected masses consistent with the sporadic evaporation, or “spitting,” of B,O, when it was exposed to the copper beam.f%l Several studies have found that elemental copper is compatible with BN;f 214)-t216)the incompatibility is solely due to the oxidizing environment. The low density of B,O, compared to copper (and other liquid metal sources) results in the B,O, rising to the surface of the liquid source and forming an encapsulant much like that used for the liquid encapsulated Czochralski (LEC) growth of GaAs. The resultant B,O, encapsulating layer not only leads to boron contamination of films, but also greatly diminishes the flux of the desired evaporant. Even though significant boron contamination exists in films, its presence may not only go unnoticed by routine EPMA or RBS composition analysis, because boron is a light element, but may also escape detection by x-ray diffraction via the formation of amorphous boron phases or by boron substitution for copper in the high T, phase being grown. For example, a solid solution between boron and copper in Y(Sr,Ba),Cu,O, exists allowing up to 90% of the copper in the CuO chain layer to be replaced by boron (note the similarity between building layer (b) and (i) in Fig. 3), which results in a concomitant decrease in Tc.f2171 Several superconducting and non-superconducting structural analogs of high T, materials containing borate layers have recently been discovered, e.g., ReBaCuO,BO,, where Re is a rare earth ion.tQ1~f21*] Of course, oxides present in the source materials themselves may also react with pBN. For this reason, even if the oxidant pressure is adequately confined away from the effusion source in order that pBN crucibles may be used, it is advantageous to use vacuum-cast copper rather than air-cast copper and to etch the copper source material before loading to remove the copper oxide on its surface and thus avoid its reaction with the pBN crucibles. Alumina (AI,O,) crucibles are a viable alternative to pBN for containing both copperf*Q1fQ41t*1Qlf**Ql and bismuth.tQ4~f21Q~f220tThe contact angle between copper and alumina in vacuum at 1200°C is 138”.f**‘] This high angle indicates that alumina is not wetted by copper, making alumina a good crucible choice. Sintered alumina is known to contain many volatile impurities which could limit the obtainable purity in an MBE process.f***) However, the doping of superconductors does not appear to be nearly as sensitive to impurities as the doping of semiconductors, where impurity

species often form undesirable deep levels. If purity should become a limit in the future, single crystal sapphire rity concentrations

are available

crucibles and have

with drastically been

used

lower impu-

in MBE.t85)f223)


Although

the thermal

decomposition

of AI,O,

557

itself at the temperatures

and vacuum conditions typically employed in MBE to contain bismuth and copper is not a concern,f 2241a t care should be taken to avoid contacting the alumina with materials typical

operation

components 2.9

(e.g., tantalumi 225)), that react with alumina and, at

temperatures,

can lead to significant

fluxes

of volatile

(e.g., A120[225)) and film contamination.

Common

Substrates

Since large area single crystal substrates of high T, superconductors such as YBa2Cu,0,B or Bi2Sr2CaCu20s+,, necessary for homoepitaxial growth, are not commercially available, other substrate materials must be selected for the heteroepitaxial growth of high T, superconductors. The ideal substrate is one which is both chemically and structurally compatible with the epilayer. Other factors, including commercial availability, cost, the potential for integration sion coefficient, dielectric

with other device technologies, thermal expanconstant, and the dielectric loss tangent at

microwave frequencies (for high frequency operation of superconducting circuits) must also be considered. Table 2 lists some common substrate materials utilized for the growth of high T, superconductors and their properties; including the room temperature lattice constants, the lattice mismatchf264) between these substrates and a (001) oriented YBa,Cu,O,_ 6 film, kub - %m)/afiim~ at room temperature, the epitaxial relationship, and the approximate thermal expansion coefficient of the substrate between room temperature and 7OO”C, (a7cooc - ~5oc)/[a25oc (700 - 25)], where available. Critical current densities at or in excess of 1O6A/cm* at 77 K (in self-field) have been achieved in YBa,Cu,O,_, films grown on all of the substrates[l03~1*451[*s51-[*s71 and substrate/barrier layer combinations1*601[*681-[*731listed. Unlike the heteroepitaxial growth of semiconductors, where it is imperative to choose a well lattice-matched substrate in order to avoid undesired dislocations in the grown structures, dislocations in superconductorst274) and especially in high 7, superconductors,t275)t276] are believed to act as effective

vortex

higher critical current densities.

pinning

sites and lead to significantly

Much higher critical current densities

are

observed in high T, filmsf276) or multilayers[277) than in high T, single crystals. For example, critical current densities as high as 8 x 1O6 A/cm* at t At temperatures less than 1300°C in UHV, the equilibrium vapor pressure of the main decomposition product of AI,O,, Al, is less than 1Oestorr. ~41 Higher oxidant pressures stablize A&Osfrom decomposition, allowing it to be employed at even higher temperatures.

558

Molecular

Beam Epitaxy

77 K (in self-field) have been achieved in YBa,Cu,O,_, fiImsf103)f278)t27g) (1.1 x lo7 A/cm* in YBa2Cu,07+/ (Nd,Ce),CuO, superlatticesf27r$, which is more than an order of magnitude densities

measured

higher concentration

higher than the highest critical current

in YBa2Cu307_B single crystals.t280) The significantly of defects present in films, including

in excess of log

and perhaps as many as 10’ 1 dislocations

per cm*, are believed responThus, the propertiesf 1W1W*W*W

sible for their improved transport relatively large lattice mismatch between high Tc superconductors and the substrates listed in Table 2 is insufficient reason to dismiss many of them from implementation. Thermal expansion mismatches on the other hand, particularly those that lead to the film being in state of tension upon cooling (i.e., substrates with smaller thermal expansion coefficients), can be a significant problem for these brittle oxide materials and lead to film cracking. The substrates listed satisfy varying aspects of the above criteria. SrTiO,, LaAIO,, and NdGaO, have good structural and chemical match to the copper-containing high T, superconductors (this is not surprising, since these substrates and the high T, superconductors all have perovskiterelated structures). The excellent structural and chemical match of SrTiO, to the copper-containing high Tc superconductors immediately made it the substrate of choice for epitaxial growth and DC transport measurements, and it remains the reference substrate to which growth on other substrates is compared. However, the high dielectric constant and high loss tangent of SrTiO, make it useless for high frequency applications.t23g)f241)f244) Of the common perovskite substrate materials, NdGaO, has the best lattice match to YBa,Cu,O,,, but the relatively high loss tangent of NdGaO, (an order of magnitude higher than that of LaAIO,) precludes its use in many microwave applicati0ns.t 244) PrGaO, has an even better lattice match than NdGaO, to YBa2Cu307_6,t281) but crack-free and twin-free PrGaO, crystals have only been grown with small diameters (5 mm or less)t282) and the critical current densities of YBa,Cu,O,, films grown on these substrates as well as the low temperature dielectric properties of this substrate have not been reported. LaAIO, has the lowest loss tangent of these perovskite substrates and is currently available in diameters as large as 75 mm;t283) 100 mm diameter LaAIO, wafers will soon be available.[284) These factors make LaAIO, the substrate of choice for high frequency applications of the copper-containing high T, superconductors, although a compatible substrate with a significantly lower loss tangent than LaAIO, could significantly improve the performance of high T, superconductor films used for high frequency, high Q (highly resonant) applications.f23g)f244)

Table 2(a).

Properties

of Common High T, Superconductors

Table

2(b).

Properties

of Common Substrates

for the Growth of High T, Superconductors

for c-axis Growth

Y2O3-Zra [I lo] y203-=h

(001)

Table 2(c).

Properties

Mp0

of Common Substrate/Barrier-Layer

((Nll$/

Al203

7.8 (u) 8.4 (c)

1 Ii021

M@(IlO)t/

1loTo

Al203

SrT103 (001)~

I

{ ii021

Al203 Y?o3-zrO?

@II)*/

A1203 (li02) ceo2

(001 )t I

YzO3-ZQ (001$ Si (001)

I

I .5x10-~

7.8 (a) 8.4 (c)

1.5x10-8

7.8 ((1) 8.4 (c)

1.5x10-8

3.8

-9.2 % (7 = 85.70)

7.8 (0) 8.4 (c)

7.8 (u) 8.4 (c)

~1~0~( 17021

1.5x10-”

Combinations

-9.2 % (y = 85.7”) -9.2 I (y = 85.7’) -9.2 % (y = 85.7’)

1.5x10-~

-9.2 w (y = 85.7’) -0.3 w

for the Growth of High Tc Superconductors

MgO[IOU,

,,

Al203 I I I?Ol YBazCujOh.6 I loOII MgO~Wl] II

239.251-253

Al203 [ I?101 YBa$&@.& I1001~ I SrTiO3 [ 1001 II

239.251.252.25,

Al203 [02211 YBa2Cu3Opg [loOI& U Y2O3-ZQ [ 1 IO] II 239.251.252.25 Al?03 IO251 I Y Ba2cujo&S [loo] I Cc02 [ 1101 II Al203 [0221] YBa$&@d [ 1001 I Y2O3-ZrO2 [ 1 IO] II Si Ill01

239.25

I .252.25

257-260

562

Molecular

Beam Epitaxy

Yttria-stabilized cubic zirconia (Y,O,-ZrO,) and MgO are more economical substrate materials and are widely utilized. Y,O,-ZrO, reacts with to form a thin epitaxial interfacial barrier layer of YBa,Cu,O,_, BaZr0,,[2481[266] which effectively the YBa2Cu,0,4

overlayer.

limits further reaction and degradation

of

However, the high loss tangent of Y,O,-ZrO,

makes it unsuitable for high frequency applications.[23Q] MgO has reasonable dielectric properties for high frequency applications, but the multiple in-plane epitaxial orientations that often exist in films grown on MgO and Y,O,-ZrO,, likely due to the poor lattice match of these substrates with the copper-containing high T, superconductors, causes undesirable high angle grain boundaries in the films.1 *851[-[*8*1While these grain boundaries are in general undesirable, their controlled introduction has been utilized to prepare Josephson junctions at specific locations for superconducting microelectronics.[254~[2*Q~[2~~ Note that MgO is quite well suited to the growth of the high T, superconductor (Ba,K)Bi0,,[2Q1~[2Q2] in pat-t due to the significantly larger per-containing high The favorable material of choice

lattice constants of (Ba,K)BiO, compared to the copT, superconductors. dielectric properties of sapphire (AI,04 make it the for high frequency applications, especially when a

device structure with a high Q is desired.I23Q1 However, direct growth of high T, superconductors on sapphire is plagued by chemical reactions. Several barrier layers, including SrTi0,,[270] CaTi0,,[2Q3] Mg0,[2531[2681[26Q1 Y203-Zr02,[2551 and Ce02[2581[272]have been successfully implemented to circumvent this reaction. However, the in-plane alignment of high T, superconductorfilms grown on these hybrid substrates is often poor, which is probably a consequence of the poor lattice match between the high T, superconductor, plane orientations orientation

the barrier layer, and sapphire. or a comparatively

Frequently,

several

in-

large mosaic spread in the in-plane

is observed.[2701[2Q3] This

results

in higher

surface

resis-

tances[288] and lower critical currents.[2851 CeO, and CaTiO, barrier layers on sapphire appear to be the most promising, having demonstrated the best in-plane orientation and lowest surface resistance among these barrier

layers.

The

surface

resistance

at microwave

frequencies

of

YBa,Cu,O,, films grown on these latter barrier layers on sapphire is comparable to the best values obtained on any other substrate (i.e., MgO and LaAlO,) .[*561[*Q31


2.10 Integration

with

The integration

563

Semiconductors of high

T, layers with Si (100) is clearly

desirable

considering that silicon is the workhorse of modern semiconductor nology, is available as large area wafers of unequaled perfection, inexpensive.

However,

copper-containing

despite its relatively

high T, superconductors,

techand is

good lattice match with the its successful

use as a sub-

strate remains elusive. The growth of copper-containing high T, superconductors directly on silicon results in an interface reaction and significant diffusion tails into both materials/ *Q41[*Q51 even when moderate (T < 650°C) growth or annealing temperatures are used.1 1*slI*QQl-f*QQlAlthough these materials are chemically incompatible, making direct integration impossible, suitable barrier layers have been found. The large amount of previous work on silicon-on-insulator technoIogy,f2QQ)f300) coupled with the excellent lattice match between the copper-containing high T, superconductors and Si {loo}, simplifies the search for suitable barrier layers. For example, Si (100) has been found to grow epitaxially on AI303 (1 i 02) ,f301) MgAl,O, {100},t302] and Y,O,-ZrO, (1 OO}.f303] Additionally, the technology has been developed to grow these oxides on silicon,f25Q)f304]-f30Q)making it, in principle, possible to move back and forth between silicon and epitaxial high T, superconductors in heterostructures. The use of buffer layers between silicon and YBa2Cu307+ and in particular the use of Y2O3Zr02,f260) has allowed the fabrication of epitaxial YBa,Cu,O,, films with high critical current densities to be fabricated on silicon. Other buffer layers including Mg0,f61 MgAI,O,,t 3101and Ce02,f31’)-f3131 are significant improvements over the direct growth of high T, superconductors on silicon substrates, densities

but only Y,O,-ZrO, in excess

buffer layers have allowed

of lo6 A/cm*

critical current

at 77 K to be achieved

in high

superconductor films. However, the large difference in thermal coefficient between silicon and the copper-containing high 7, ductors remains a significant problem. This thermal expansion not nearly as severe for epitaxial high 7, superconductor films

T,

expansion superconproblem is grown on

silicon-on-sapphire, where crack-free YBa3Cu30,, films more than eight times thicker than the onset of cracking for such films on silicon substrates (m 500 A)f26o) have been grown and have excellent critical current densities (4.6 x 1O6 A/cm* at 77 K for a 1300 8, thick film) .f314) In light of the desirable high frequency properties of sapphire substrates and the progress that has been made in integrating them with high T, superconductors, sapphire substrates may be the best substrate choice for monolithic superconductor-

564

Molecular

Beam Epitaxy

semiconductor

electronics.

demonstration

that functional

high T, superconductor

A significant

step in this direction

silicon-based

electronics

semiconductor

can be fabricated

is the recent circuitry

and

on the same sap-

phire substrate.f315) The integration of high T, superconductor layers with GaAs and Ill-V materials is also highly desirable from a hybrid monolithic microwave integrated circuit (MMIC) standpoint. An MgO barrier layer has been found to work well for GaA~.f*~~l Not only are the properties of the superconductors grown on such barrier layers quite encouraging, but experiments have shown only minor degradation of Ill-V heterostructures (e.g., a two-dimensional electron gas 1200 A below the MgO interface) after the growth of high T, layers on top of these heterostructures.f316) Critical current densities as high as 1.2 x lo6 A/cm* at 77 K have been achieved with YBa,Cu,O,_JMgO/GaAs films.1 273) Because of the higher thermal expansion coefficient of GaAs compared to silicon, crack-free YBa,Cu,O,, films may be prepared up to a thickness than on silicon substrates.

3.0

about five times greater (= 2500 A)f3171

SPECIFIC HIGH T, MATERIALS SYNTHESIS CAPABILITIES

AND DEMONSTRATED

Despite significant improvements in the capabilities of technique for the growth of high T, superconductor films, it must that the films exhibiting the best transport properties have produced by MBE. Rather, the films with the highest transition tures and highest critical currents have been produced by

the MBE be noted not been temperaPLD and

sputtering.f103)f277H27g) Nor has MBE led to the discovery of new higher T, superconductor materials; indeed it is bulk synthesis techniques which continue to provide such advances.

The strength of the MBE technique

is,

in the opinion of the authors, not to duplicate what can be made by bulk techniques nor single target thin film synthesis methods, which are free from the multiple source composition control issues of MBE, but instead to provide a customized

layering

capability

with control at the atomic layer

level for the synthesis of metastable structures and device heterostructures. This is not to say that high T, superconductor phases that may be readily produced by single target techniques, e.g., YBa2Cu307_S, are of poor quality when synthesized by MBE, but rather that for such a simple structure (a thick film of YBa,Cu,O,_,) the capabilities of MBE are not

High

needed and that other (more economical) more appropriate. observed

For example,

in ReBa&u,O,,

Tc Superconductors

synthesis

techniques

565

would be

the highest critical current densities

(JJ

films (where Re is Y or a rare earth) prepared

by MBE are 4 x 1O6 A/cm2 at 77 K,fgo)f3181compared to 8 x 1O6 A/cm2 at 77 Kwhich has been obtained in films grown by PLD and sputtering.t103)t27e)t27g] Just as a thick semiconductor

layer requiring

high mobility

and excellent

optical properties would be better grown by LPE, for example, than by MBE, only certain high T, structures are appropriate for and warrant the use of MBE. MBE is best suited to structures requiring its layering control and in-situ diagnostic capabilities, and structures for which such layering control would yield a significant improvement in device performance or theoretical understanding. In the examples described below, we highlight the progress of the MBE technique toward customized growth. It is hoped that with continued advances, particularly in composition control, that the customized layering capability of MBE may be adequately exploited to allow the physics and device potentials of new layered structures that cannot be produced by bulk methods, including higher T, superconductors, to be explored. 3.1

ReBa,Cu,O,,

superconducting films, whereRe is Y,[451~g21[g71[1321[1661~1681 ReBa.$u,O,B Nd f173)Sm,flrc) and Dy[ goIf16g1with acceptable transport properties have been grown by numerous groups using MBE. Notable among this research is the synthesis of DyBa2Cu307B films with T,‘s as high as 92 K and Jc’s (77 K, self-field) as high as 4 x lo6 A/cm2 with a pressure at the substrate surface of about 4 x 1O-6torr, using ozone.fgO) YBa,Cu,O,, films of similar quality

have been produced

source with a pressure at the substrate using an oxygen plasma source.f1521

by MBE using an atomic oxygen surface

of about 1.5 x 1 0e4 torr

In-situ RHEED analysis and RHEED oscillations have been used extensively during the MBE growth of these materials, allowing the thickness of the growth unit to be measuredfg0~f168~f173)f31g)f320) as well as the inplane orientation relationship,f18)f42)f46)fg0)fg4)f174)t204)f321) lattice constant relaxation process,f45) and the formation of impurity phasesfg4)f204) to be monitored. During the growth of ReBa,Cu,O,, film by codeposition, the periodicity of the RHEED oscillations corresponds to a layer thickness equal to the c-axis dimension of the unit cell for films oriented with their caxis normal to the plane of the substratefg0)[168)f173)[31g) (confirming the

666

Molecular

Beam Epitaxy

result that was first observed

using RHEED

evaporationt102)) and a layer thickness

oscillations

during

reactive

equal to the a-axis dimension

unit cell for films oriented with their a-axis perpendicular

of the

to the plane of the

substrate.t320] The pristine surfaces prepared by MBE have been analyzed in-situ by low-energy

ion scattering

spectroscopy

(LEISS) to establish that

the surfaces of codeposited YBa2Cu,0,_B films are terminated by the CuO chain /ayer.[3211 This, together with the RHEED oscillation results, implies that the minimum growth unit during codeposition of ReBa,Cu,O,, phases is (BaO-CuO,-Re-CuO,-BaO-CuO), with CuO at the film surface, or in terms of building layers, (-BaO]-[CuO,]-[Re]-[Cu02]-[BaO-CuO-), where the [BaO-CuO-BaO] building layer is split after the CuO chain layer. The step heights revealed by STM measurements on surfaces of MBE-grown DyBa2Cu,0,_Bfilms correspond to the periodicity of the RHEED oscillations (the minimum growth unit) and also indicate the presence of a high density (W1Og cm-*) of screw dislocations.tg0)f31g) Such densities have been seen in MBE films grown on SrTiO, (1 00},fgo1t31g]MgO (1 OO},fgo]and. NdGaO, {lOO}.fgo) These defect sites are not specific to MBE-grown films, but have also been reported in ReBa2Cu30,_B films grown by sputtering,f103)t1041t322) PLD,t3221f323)reactive evaporation,f3241 and CVD.f325) The presence of such a high density of screw dislocations is a common feature of the growth of ReBa,Cu,O,, films by gas phase codeposition methods on oriented substrates.t326] The resulting well-known spiral growth mechanism, first proposed by Frank,1 3*7l occurs quite frequently in the growth of of this growth mechanism is layered materia1s.t 3*8lt3*9) A consequence that the film surfaces are not atomically flat; instead, the surfaces consist of mounds, each about 300 nm in diameter and each containing a screw dislocation at its center. The surface roughness depends on growth conditions, but height variations of 50-100 8, from the top of each mound to its intersection with the neighboring mound are fairly typical in MBE films.tgO) Studies by other gas phase codeposition techniques (sputteringf3**] and PLDf330]) have shown that, if the substrate is misoriented sufficiently (typically a few degrees for the growth of YBa,Cu,O,_, films), a high density of screw dislocations is not nucleated, and instead growth occurs by step propagation of the terraces, which accommodate the tilt of the substrate and are energetically favorable incorporation sites. Detailed studies on the surface morphology of films grown by sequential deposition methods are needed to ascertain how such growth conditions impact the minimum growth unit compared to the more widespread use of codeposition.


567

Device structures utilizing MBE-grown ReBa,Cu,O,B superconductors have been made, but such structures have made use of lateral (inplane) boundaries (edge junctions), ~~333~ which rely on differences in epitaxy at a patterned edge, rather than making use of the layering capability of MBE. Similar device structures can be (and have been) made equally

well

by other thin film deposition

techniques,

and will not be

discussed here. Attempts to grow vertical Josephson junctions using the layering capability of MBE have met with limited success, presumably due to pinholes in the thin insulator layer.t113j In general, the layering capability of MBE has not been extensively utilized in the growth of high T, structures containing ReBa,Cu30,a phases. However, one notable example demonstrates the significant difference that shuttered MBE can have on the growth of ReBa&u,O,, superconductor phases, and has allowed them to be synthesized at much lower temperatures than can be achieved by other techniques. Utilizing sequential

deposition,

superconducting

DyBa,Cu30,B

films with their c-

axis normal to the plane of the substrate (c-axis oriented) have been prepared by MBE at substrate temperatures as low as 420”C.f158) Codeposition of the constituent species under the same conditions results in an amorphous film/ 334) whereas shuttering the incident fluxes in the same order as the c-axis order of the building layers produces a c-axis oriented film.t158] It is well established from a variety of gas phase synthesis techniques that in order to attain c-axis oriented ReBa,Cu,O,, films by codeposition, high temperatures (typically exceeding about 650°C) are needed.t31][335jt336) Lower temperatures result in films with their a-axis normal to the plane of the substratet 31I[sssj-tss7j (a-axis oriented), until below about 500°C where crystallization of the ReBa,Cu,O,, phase is no longer 0bserved.t 3381This is just one example of the layering case orientation) control achievable using shuttered MBE. 3.2

(and in this

Bi,Sr,Ca,.,Cu,O,,+, In contrast to the MBE growth of ReB+Cu,O,,

high T, superconduc-

tor phases, where little use has been made of shuttering,

shuttering

has

been extensively used in the growth of Bi,Sr,Ca,_,Cu,O,,+, phases.[ss1[911[1~c1[1541[1601[1s11[1711[1721[1751[~~91-[~4sl As described in Sec. 1 ,4, bulk methods have only been capable of synthesizing

Bi2Sr2Ca,,_,Cu,02,+,

phases in single phase form for n = 1 to 3, presumably due to the nearly degenerate formation energies of higher n members of this homologous

568

Molecular

series.

Beam Epitaxy

Using shuttered

MBE growth, single phase Bi,Sr,Ca,_,Cu,O,,+,

films, for n = 1 to 11, have been grown/ 1~34~) demonstrating the ability of this technique to select between nearly energetically degenerate phases. In Fig. 22,

6-28 x-ray

diffraction

scans

of as-grown,

c-axis

oriented,

epitaxial

Bi,Sr,Ca,_, CU,O~~+~ films are shown for n = 1 to 5.t130) The corresponding Bi,Sr,Ca,_, CLJ,O~~+~structures are shown in Fig. 10. Crosssectional TEM has confirmed the layered nature of these films, shown LP axis superstructure, and revealed the presence of stacking faults in the films.t8s)t33s~ Atomic force microscopy (AFM) scans over 0.5 I_tm by 0.5 pm regions of the film surface have revealed a surface roughness of less than 10 AP4gt Superconducting films with as-grown zero resistance temperatures as high as 90 Kt 172) and Jc’s (4.2 K, self-field) exceeding 2 x 10’ A/cm2 have been grown by shuttered MBE.t 172) The success of the MBE growth layering precision has of Bi,Sr,Ca,_, C~“02~+4 P hases with sub-unit-cell allowed the growth of metastable phases, ordered superlattices, and device heterostructures, which are described in Sets. 3.5, 3.6, and 3.7.

cm n=4

I....I..~.I.~~~1.~~~~~““““~“‘.

0

20

40

60

20 (degrees) Figure 22. O-20 x-ray diffraction scans of as-grown, single phase, c-axis oriented, epitaxial

Bi,Sr,Ca,_,

CIJ,O~~+~ films (after Ref. 130) with n = 1 to 5.


In-situ RHEED provides direct monitoring the constituent

elements

569

of the effect that supplying

in different ways to the substrate during the p hases has on surface structure. A se-

growth of Bi,Sr,Ca,_,Cu,O,,+,

quence of RHEED patterns is shown in Fig. 23 from the shuttered of a Bi,Sr,CaCu,Os

film, where the cation fluxes were supplied

growth individu-

ally in the same order as the building layers making up the structure (i.e., Bi-Sr-Cu-Ca-Cu-Sr-Bi, ...) an d ozone was supplied continuously.f130) Figure 23(a) shows the bare MgO (100) substrate before growth. Figures 23(b) to 23(f) show the RHEED patterns during the growth of a half unit cell of Bi,Sr,CaCu,O, (from the bismuth layer completing the 28th half unit cell to that completing the 2gth half unit cell). The streaked patterns observed after the deposition of the bismuth atoms [Fig. 23(b) and Fig. 23(f)] and the presence of spots in the other photos indicate that the surface is smoothest after the bismuth layers which complete each half unit cell, and that, within the growth of each half unit cell, islands (as schematically indicated in Fig. 16) are present. The pattern of spots in the RHEED pattern is consistent with the oriented presence of islands of the Sr,Bi,O, phase.f3501f351) Since only the Bi,Sr,CaCu,O, phase was observed by x-ray diffraction after growth, and SEM images of the surface did not indicate the presence of a second phase after growth, it was concluded that the SrsBi,O, phase was only temporarily formed during the shuttered growth process; with the completion of each half unit cell, it is transformed into Bi,Sr,CaCu,0,.f130j The temporary formation of the Sr,Bi,O, phase is believed to be the way that the surface adjusts to becoming strontium rich after the deposition of each strontium burst,f130) since Sr,Bi,O, is the most strontium-rich Bi-Sr-0 phase (see Fig. 8) P8)tss1) The streaked pattern observed after the completion of each half unit cell indicates that shuttered MBE is effective in layering

on a half unit cell by half unit cell basis.

The islanding

of other

phases on a temporary basis during the shuttered growth of Bi,Sr,Ca,_, cu n0 2n+4phases, including SrO and Ca0,f172) has also been seen. As one might expect, the interplay between growth conditions and surface structure is significant, and a distinct advantage of MBE is the ability to explore this interplay in a controlled manner, since the growth conditions (substrate temperature,

depositing

fluxes,

and flux sequence)

may be con-

trolled independently. In order to achieve layering control on a finer scale than a cell, it is necessary to find growth conditions which maintain a surface at all times during the deposition of the building layers. easily studied in-situ using RHEED, or even more sensitively,

half unit smooth This is using a

570

Molecular

Beam Epitaxy

recently developed technique called difference R/-/EED,[~~~I in which a digitally stored RHEED image is subtracted from the current RHEED image in order to discern more sensitively the effects that the species deposited

in the time between the two RHEED images have had on the

surface structure.f345] This technique has been applied in real time to the growth of Bi,Sr,Ca,_, CU,O~~+~ p hases in order to monitor their smoothness.f3451

Figure 23. RHEED patterns observed along the MgO cl OO> azimuth during the (n=2) film on MgO (100) at Tsub EI 650°C (after Ref. growth of a Bi,Sr,CaCu,O, 130). The shuttering sequence for each half unit cell was Sr-Cu-Ca-Cu-Sr-Bi, with ozone on continuously. (a) Bare MgO (100) substrate before growth. (b) After the Bi layer completing the 28th half unit cell. (c) After the first Sr layer of the 29th half unit cell. (d) After a Cu layer of the 29th half unit cell. (e) After the second Sr layer of the 29th half unit cell. (r) After the Bi layer completing the 29th half unit cell.


Growth conditions that discourage during the growth of Bi,Sr,Ca,_,Cu,O,,+, method is to shutter the oxidant Decreasing

571

the temporary formation of islands phases have been found. One

in addition

to the cation species.tQ4]t33Qt

the ozone pressure during the strontium

flux burst was found

to result in a streaked RHEED pattern throughout the shuttered growth of Bi,Sr,CuO, films. This effect has been attributed to interplay between the incident oxidant flux and the surface mobility of the adatoms on the substrate surface.tQ41f33Q] When considered individually, the oxidation requirements of the constituent monolayers are significantly different,t8’] so it is not surprising that modulation of the oxidant flux during growth may be beneficial. Shuttering the oxidant during the MBE growth of high T, superconductors is quite analogous to the MEE growth of Ill-V compound semiconductors,t1081 where the group III species are deposited in the absence of the group V species, resulting in a drastic increase in the surface mobility of the group III species.fJcsl Modulation of the oxidant flux, in addition to the cation fluxes, has been explored by several groups In for the growth of Bi,Sr,Ca,_,Cu,O,,+, p ~~~~~~~Q~l~~~~l~~~~l~~~~1~~~Ql~~~~l some cases, the substrate temperature has also been modulated f16’X1s13t175lP461 d emonstrating the extreme flexibility and range of gro$h

conditions

accessible with the MBE technique, and allowing hases to be formed at substrate temperatures as Bi2Sr2Ca,-ICu,02,+4 P low as 300”C.[160~~1s1~~175] Avoiding the temporary formation of islands by modulating the ozone flux has allowed a superconducting metastable Bi,Sr,Cu,O, high T, superconductor phase to be prepared by MBE, which is described in Sec. 3.7. Another method of deterring the formation of islands during shuttered growth is to codeposit several of the constituent monolayers. For example, codepositing the calcium and copper fluxes has been used to eliminate the temporary formation of islands during the shuttered growth of Bi2Sr2Ca,-I CU,O~~+~ Phases, in which the other cation fluxes are individually supplied under a continuous

flux of ozone.f1721f3411However,

as more

of the constituent species are codeposited, the layering control of the MBE process is diminished and eventually reaches the half unit cell level when all of the Bi,Sr,Ca,_,Cu,O,,+, established by RHEED

constituents oscillation

are codeposited, as has been measurements during

codeposition.f352]f353] 3.3

TI,Ba,Ca,_,Cu,O,,+,

Largely because of their toxicity, thin film research on the TIBa&a,_, and T12Ba2Ca,_,Cu,02,+4 p hases has been relatively limited. CU”02n+3

572

Molecular

Beam Epitaxy

These materials have not been grown by MBE and only recently has their in-situ preparation been realized.f354)-f356) The pressures, = 200 mtorr f3ss)fss6jused in their in-situ growth by a combination sputteringf356) of Ba-Ca-Cu-0

with simultaneous

thermal

of PLDf355) or evaporation

of

TI,O are well above the MBE regime, but the minimum oxygen pressure needed to stabilize these phases has not been measured. However, the oxygen pressure at which TI,Ba,Ca,Cu,O,, decomposes into Tl,Ba,CaCu,O, has been measured as a function of temperature and oxygen pressure,f35q providing some hope that it may be possible to synthesize TI,Ba,Ca,Cu,O,, under MBE conditions using an effective oxidant. 3.4

(Ba,K)BiO,

The growth of (Ba,K)BiO, and (Ba,Rb)BiO, have been extensively The simple cubic perovskite strucstudied by MBE.f ‘151-[11~[1891[2911[292l ture, comparatively long coherence length (longest of the high T, superconductors), low number of constituents, and extremely low growth temperatures (typically 300°C) make this an obvious model system for study. In addition, MBE growth conditions in which the K (or Rb) and 0 incorporation are adsorption-controlled have been found.1 ’ 151This greatly simplifies the composition control requirements, since it becomes only necessary to control the barium and bismuth fluxes. Despite its comparatively low T,, the excellent demonstrated properties of (Ba,K)BiO, and (Ba,Rb)BiO, Josephson junctions make this materials system a serious contender for future applications of high T, superconductors in microelectronics, particularly for superconductor-insulator-superconductor base transistors

utilizing

(Ba,Rb)BiO,

(SIS) mixers.

Metal-

as the metallic base layer have also

been made by MBE.f 3581,t Initial results indicate

a common-base

current

gain, a, near unity for these structures, which is a significant improvement over metal-base transistors made from conventional metals and semiconductors.f256) Since it is the only 3-dimensional

high

T, superconductor,

it is of

fundamental interest to study the effect of dimensional confinement on the superconducting properties of (Ba,K)BiO,. One method of accomplishing

t These devices are operated at temperatures above Tc, and thus do not utilize the superconducting properties of these oxides, but rather utilize their low carrier concentrations (typically 5 x 1O*’ cm”for high T, superconductors), which are more than an order of magnitude lower than those of conventional metals.[3581


573

this would be to use the unit cell layering precision of MBE to intersperse barrier layers into (Ba,K)BiO, at controlled intervals. Dimensionality issues have been widely studied for several of the 2-dimensional coppercontaining high Tc superconductors through the use of superlattices, in which non-superconducting layers separate the superconducting ones.f34]f38)f342) However,

such studies

by controlled

thin film methods

have not

been performed on the 3-dimensional (Ba,K)BiO, superconductor. Bulk synthesis methods have been used to investigate the effect of reduced dimensionality

on

these materials through the synthesis of hases.~6)f359)-(361) These are layered structures in (BaN,+l (Pb,Bi),Os,+l P which (Ba,K)O layers occur between every n perovskite layers. However, the scope of these bulk studies has been severely limited both because of the presence of uncontrolled intergrowths for higher n and the restricted range of solid solution between the constituents.~6)f35g)-f361) Only then = 1, n = 2, and n = co members of this homologous series of compounds have been prepared in single phase form;f s~fs61l the n = 3 structure has been observed locally by TEM, intergrown with higher n compounds in an uncontrolled manner.p6) The layering control and site-selective doping capabilities of MBE are probably well suited to studying these materials. 3.5

Superlattices

Both ordered and disordered superlattices have been prepared by shuttered MBE,fg4)f1301although the latter were not deliberately prepared, but resulted from inadequate composition control.fss)fg4) An example of each will be discussed since other thin film techniques have also attempted to make ordered Bi,Sr,Ca,_,Cu,O,,+, superlattices, thus providing an opportunity to compare the custom-layering capability of MBE with other techniques. In order to interpret the x-ray diffraction patterns from Bi,Sr,Ca,_, superlattices, it is useful to compare the observed patterns to CU”02n+4 two limiting case simulations: an ideal superlattice (perfect building layer ordering)

and completely

disordered

layering

consisting

stacking sequence of the constituent building layers. are shown schematically in Fig. 24.

of a random

These limiting cases

The x-ray diffraction theory necessary to perform these simulations is well established. The perfectly ordered superlattice can be simulated by standard kinematic diffraction theory/ 3621where the diffracted intensity due to all of the cations in the ideal superlattice is appropriately summed,

574

Molecular

Beam Epitaxy

Ordered Superlattice

Random Layer Order

0

0

l

‘0

Figure 24. A schematic representation of the layering sequences of two limiting-

case x-ray diffraction simulations. The perfectly ordered superlattice has regular repeat periodicity &,, while the random layering sequence does not.

including the isotropic temperature factors and the Lorentz polarization factor.t”] For simplicity, flat idealized building layers free of vacancjes, cation substitution, and incommensurate superstructure are assumed, and the oxygen atoms in these structures as well as x-ray absorption effects are disregarded, These simplifying assumptions allow qualitative comparisons to the observed x-ray diffraction patterns to be readily made, whereas accurate quantitative modeling of the patterns requires far more knowledge of the structure, strain, and site substitution actually present in these films, as opposed to their bulk counterparts. The completely disordered superlattice can be simulated using the theory of Hendricks and Tellert3s3] which applies to an infinitely thick, one-dimensional layer lattice with a random stacking sequence. For simplicity, scattering from the light atoms in these structures (oxygen, copper, and calcium) is disregarded. Since the atomic scattering factor, 4, increases with atomic number, this is not a bad approximation since most of the scattering is due to the bismuth and strontium atoms in Bi,Sr,Ca,.,Cu,02n+4 phases. The problem then reduces to scattering from only one type of building layer, [SrO-BiO-BiO-


575

SrO] without oxygen, where the distance between these identical layers takes on distinct values, but for which the sequence of the spacings is random.

For example,

Bi,Sr,Ca,Cu,O,, c-axis

length

a randomly layered mixture of Bi,Sr,CaCu,O,

layers contains two distinct interlayer of Bi,Sr,CaCu,O,

and

half

distances:

the

c-axis

and half the

length

of

Bi,Sr,Ca,Cu,0,,.fQ4) Several groups have reported unusual x-ray patterns from Bi-Sr-CaCu-0 films grown by sputtering, fs~sss) shuttered ion beam sputtering,f366] laser ablation,t367)t368) and shuttered MBE.tss] The growth of superlattices was not attempted in these growths, so the strange x-ray patterns observed were at first not understood. Although their c-axis lengths and peak intensity distributions are similar to Bi,Sr2Can_,Cu,0,,+, phases, the diffraction patterns do not correspond to any of the known Bi$%$a,,_, CIJ,O~~+~ phases. A least squares estimate of the c-axis length based on the peak positions results in values in between those of the known Bi,Sr,Ca,_, Further, indexing the peaks to these intermediate CU,O~~+~ phases. lattice constants requires the use of both even and odd OOPpeaks. Due to the glide plane half way up the unit cell of all known Bi,Sr2Can_,Cu,0,,+, structures, destructive interference results in the absence of all odd 001 peaks, making the observed x-ray patterns quite unusual. The repeated layering of a half unit cell of Bi,Sr,CaCu,O,

followed

by

a half unit cell of Bi,Sr,Ca,Cu,O,, results in a new unit cell with c-axis length halfway between that of Bi,Sr,CaCu,O, and Bi,Sr,Ca,Cu,O,,. Since this ordered layering no longer contains a glide plane halfway up the unit cell, both even and odd OOPx-ray diffraction peaks would be allowed. So one possibility is that the unusual x-ray patterns observed are due to an ordered superstructure

of Bi,Sr,Ca,_,Cu,O,,+,

To test this explanation, c-axis

oriented

the x-ray diffraction

Bi,Sr,Ca,Cu,O,,

phase

phases. pattern of a hypothetical film

(Bi,Sr,CaCu,O,

+

Bi,Sr,Ca,Cu,O,,) was calculated. Another possibility is that a random mixture of Bi,Sr,CaCu,O, and Bi,Sr,Ca,Cu,O,, layers is present in such samples and so the x-ray patterns of c-axis oriented Bi2Sr2CaCu,0,

and Bi,Sr,Ca,Cu,O,,

random mixtures

layers in various

proportions

of

were

also simulated. The simulated diffraction patterns are shown in Fig. 25 and Fig. 26. The most intense peaks of both the ordered and disordered layer lattice simulations lie at similar 20 positions. Comparison to the observed x-ray diffraction patterns requires consideration of more subtle aspects of the simulated patterns. The ordered superlattice contains more peaks, while the randomly layered mixture exhibits periodic broadening

576

Molecular

and narrowing

Beam Epitaxy

of the peak width as a function

of diffraction

vector.

The

observed patterns do not contain the extra peaks which should be present for an ordered sequence of the layers. Further, low-angle x-ray diffraction studies

do not reveal the 001 or 003 peaks expected

layer simulation.tg4)

Thus, the x-ray diffraction

from the ordered

data indicate an absence of

layering order in these structures, which is not surprising since no attempt to attain a superlattice was made in these growths. The lack of repeated layering order has also been confirmed by cross-sectional TEM.tss) However, as discussed below, regularly alternating ordered superlattices in+ Bi,Sr,Ca,Cu,O,,) and cluding Bi,Sr,Ca,CusO,s ( Bi 2Sr 2CaCu,Os Bi,Sr,Ca,Cu,O,, (Bi,Sr,CuO, + Bi,Sr,Ca,Cu,O,,), can be grown by shuttered MBE by using a shuttering sequence corresponding to the desired ordered superlattice.tg4]

Bi,Sr4Ca,Cu,0,

8 (n = 2 I n = 3)

Ordered Superlattice z 0 0

Randomly-Layered Mixture 50%n=2,50%11=3

0

20

40

60

80

28 (degrees) Figure 25. Calculated O-20 x-ray diffraction intensity patterns for a perfectly ordered Bi,Sr,Ca,Cu,O,, phase (alternate half-unit-cells of Bi,Sr,CaCu,O, and Bi,Sr,Ca,Cu,O,,) and an infinitely thick, layered, random mixture of 50% Bi,Sr,CaCu,O, (n=2) and 50% Bi,Sr,Ca,Cu,O,, (n=3) layers. A Gaussian instrumental broadening of 0.5” was used in these calculations.


577

30% n=2 70% n=3 .i

.-AL

~.,,,,,,,,,,,,.,,(., m__..mh

0

20

10% n=2 90% n=3 (,,.,,.,., (..~.

~A-A

60

40

^.. SrJ

28 (degrees) Figure 26. Calculated O-20 x-ray diffraction intensity patterns for an infinitely thick, layered, random mixture of Bi,Sr,CaCu,O, (n=2) and Bi,Sr,Ca,Cu,O,, (n=3) layers in various proportions. A Gaussian instrumental broadening of 0.5” was used in these calculations. Although

in-situ

RHEED patterns observed during the growth of Phases, e.g., those shown in Fig. 23, indicate that

Bi&2Ca,-lCu,C2,+4 shuttered MBE is capable of layering on a half unit cell basis, this has been directly tested by growing superlattices consisting of alternate half unit For example, a shuttering secells of Bi,Sr,Ca,_,Cu,O,,+, p hases.t=l quence corresponding followed

to the deposition

of a half-unit-cell

by a half unit cell of Bi,Sr,Ca,Cu,O,,

of Bi,Sr,CuO,

was repeated

61 times.+

t This particular example is chosen because it clearly demonstrates the ability of MBE to control the layering order. Microscopic studies of bulkt36g)and thin filmt365)samples have shown the existence of aslight energetic preference for the alternate arrangement of Bi2Sr2CarrlCun02n+4 and Bi,Sr,Ca,,Cu,+,O,,+, half-unit-cells in samples of intermediate composition. Fully ordered superlattices, i.e., phases, have not been seen in preparation methods not invoking half unit cell layering control, but there does appear to be a slight driving force toward alternating stacking over a few repeat distances. For the superlattice example discussed, the average composition is Bi,Sr&aCu,Oe (see Eq. 5). Since bulk techniques are capable of preparing the BiiSr2CaCu20s in single phase form, we know that energetically this phase is preferred over an alternating halfunit-cell superlattice of BiaSr2Cu06 and Bi,Sr,Ca,Cu,O, c layers, making this a good test for the layering ability of MBE; the superlattice discussed is metastable with respectto the BiiSr,CaCusOs phase.

578

Molecular

Beam Epitaxy

The x-ray diffraction

pattern of the grown

film is shown

in Fig. 27. The 8-

28 scan contains both even and odd OOJ peaks. The measured c-axis length (30.75 + 0.1 A) is about halfway between that of Bi,Sr,CuO, and Rocking off the substrate (see Fig. 27) results in a Bi,Sr,Ca,Cu,O,,. reduction of the x-ray intensity to the noise level, indicating a highly The epitaxial alignment between the oriented c-axis oriented sample. superlattice and the underlying substrate has been confirmed by an x-ray $-scan of an inclined film reflection. The four peaks in Fig. 28 indicate that the superlattice is epitaxially oriented with respect to the SrTiO, substrate with the same in-plane orientation CI_J,O~~+~films.

relationship

as observed for Bi,Sr,Ca,_,

Aligned

to

( 100) SrTiO, Rocked 5” Off

0

20

40

60

80

2 8 (degrees) Figure 27. The O-20 x-ray diffraction scans of an as-grown superlattice of alternating half-unit-cells of Bi,Sr,CuO, (n=l) and Bi,Sr,Ca,Cu,O,, (n+3). Note the log intensity scale. The O-20 scans aligned to the SrTiO, (100) substrate and rocked 5” in omega off alignment to the SrTiO, (100) substrate are shown (from Ref. 130).

High

Tc Superconductors

579

$ (degrees) Figure 28. An x-ray diffraction $-scan of the 115 peaks of an as-grown superlattice of alternating half unit cells of Bi,Sr,CuO, (n=l) and Bi,Sr,Ca$u,O,, (n=3). 41 = 0 was set parallel to SrTiO, . This scan shows that the in-plane alignment between the film and substrate is with the a and b superlattice film axes parallel to SrTiO, ~11 O> (from Ref. 130).

One possible atom rearrangement which could occur during the attempted growth of a superlattice consisting of alternating Bi,Sr,CuOs and Bi,Sr,Ca,Cu,O,, layers, distinct from the formation of a randomly ordered mixture of these two phases, is the formation of single phase Bi,Sr,CaCu,O, Eq. (5)

by the reaction:

Bi,Sr,CuO,

+ Bi,Sr,Ca,Cu,O,,

--, 2 Bi,Sr,CaCu,O,

However, since the presence of a glide plane halfway up the unit cell of Bi,Sr,CaCu,O, (and all Bi,Sr,Ca,_,Cu,O,,+, phases) leads to only even OOLdiffraction peaks, the presence of odd OOPdiffraction peaks in Fig. 27 rules out this possibility. Comparison of the x-ray diffraction patterns of this superlattice to the two limiting case x-ray diffraction simulations confirms its layered nature and demonstrates that shuttered MBE is able to layer Bi,Sr,Ca,_,Cu,O,,+, phases on a half unit cell basis. The simulated

diffraction

patterns are less

580

Molecular

sensitive

Beam Epitaxy

to errors

trapolated

in the assumed

from the structure

20 angles than at high 26 angles. observed

x-ray diffraction

structural

parameters

of bulk Bi,Sr,Ca,_,Cu,O,,+,

(which

In Fig. 29, a comparison

pattern of the superlattice

are ex-

phases) at low between

the

and the two limiting

case x-ray diffraction simulations is shown for low 28 values. The 001 peak (at 20 = 3.17, which would be absent for random layering, is quite evident in the observed diffraction pattern, repeated ordered layering in this MBE-grown

indicating the presence of superlattice. The regularity

of this and other MBE-prepared superlattices,tg4)f341)t342]as revealed by x-ray diffraction and a comparison to these two limiting-case simulations, compares favorably to attempts to synthesize superlattices of Bi,Sr,Ca,_, C1.1,02~+4 phases

by all other

I

techniques.[401[370]-f372]

I

I

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0

reported

I

2

I

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I

I

4

I

I

I

I

6

I

I

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t

10

28 (degrees) Figure 29. The observed low angle O-20 x-ray scan of an as-grown superlattice of alternating half unit cells of Bi,Sr,CuO, (n=l) and Bi,Sr,Ca,Cu,O,, (n=3), calculated O-20 x-ray intensity pattern for a perfectly ordered 61 -period superlattice, and calculated O-20 diffracted x-ray intensity pattern for an infinitely thick layered random mixture of 50% Bi,Sr,CuO, (n=l) and 50% Bi,Sr,Ca,Cu,O,, (n=3) layers. A Gaussian instrumental broadening of 0.5” was used in the x-ray calculations.


3.6

Josephson

581

Junctions

Because of the extremely short coherence lengths, 5, in high T, superconductors, the fabrication of Josephson junctions, where two superconducting conducting

regions need to be separated

barrier layer of a thickness

by a pinhole-free

comparable

non-super-

to s, is a considerable

challenge. The length, 5, ranges from a few angstroms to a few tens of angstroms in these materials, making the precise layering-control of MBE a necessity for the controlled synthesis of sandwich-type Josephson junctions. The MBE growth of sandwich-type Josephson junctions consisting of successful and has BGWa&u,Os,+~ Phases has been particularly yielded both hysteretic and non-hysteretic Josephson junctions,t343)t344j which are important for superconducting electronics. The junctions are formed by the MBE growth of a superconducting Bi,Sr,CaCu,Os film into which a sing/e half unit cell thick barrier layer of Bi,Sr,Ca,_,Cu,O,,+, (with n = 5 to 11) is interspersed.f 34s1[344j Using the site-selective doping capability of shuttered MBE, the central calcium monolayers of the barrier have been doped with bismuth, strontium, B~,S~,C~,_,CU,O,,+~ and dysprosium, causing significant changes in the conductivity of the barrier and allowing the junction critical current, /,_,of these sandwich-type Josephson junctions to be tuned over four orders of magnitude, while maintaining a nearly constant I,/?” product of about 0.5 mV, where R, is the junction normal state resistance.f 343jt344) The quality of the resulting junctions indicates that the half unit cell thick barrier layers (a thickness of 25 a to 44 8) are free of pinholes over the 30 pm x 30 pm areas of the sandwich junctions.f343jf344] Subsequent work using Dy-doped BiSr,Ca,Cu,O,, barriers has yielded I& products in excess of 5 mV.f172] MBE has also been used to prepare (Ba,K)BiO, Josephson junctions.f373)-f376j The comparatively isotropic

cubic symmetry

guish (Ba,K)BiO, technological

of (Ba,K)BiO,

for the fabrication

(and the related compound

high T, materials from which Josephson

length

are characteristics

from the rest of the high

advantages

deed, (Ba,K)BiO,

long coherence

(= 50 A) and which

Tc superconductors of Josephson (Ba,Rb)BiO,)

junctions

distinand are

junctions.

In-

are the only

with nearly ideal BCS-

like tunnel junction behavior have been fabricated.f3’q Sandwich-type Josephson tunnel junctions have been fabricated using the layering capability of MBE to intersperse a non-superconducting BaBi,Oy barrier layer between two superconducting (Ba,K)BiO, layers.f374) KNbO, has been found to work even better as a barrier layer, presumably because of the presence of potassium in all of the layers of the heterostructure; with it nearly ideal sandwich-type SIS Josephson junctions have been fabricated.[375]

582

Molecular

Beam Epitaxy

Excellent SIS Josephson controlled

on bicrystalline 3.7

junctions

grain boundary

Formation

have also been fabricated

(the SIS junction)

by growing

by creating a

(Ba,K)BiO,

films

substrates.t376j of Metastable

Structures

As described in Sec. 3.2, the ability to avoid the temporary formation of islands is a crucial aspect of achieving layering control on a finer scale than the unit cell or half unit cell level which is accessible through codeposition methods, By shuttering the oxidant flux in addition to the cation fluxes, islanding during the shuttered MBE growth of Bi,Sr,Cu06 can be avoided, as evidenced by the RHEED pattern remaining streaked throughout the growth of Bi,Sr,CuO, films.tg4)fssgj Using such growth conditions, the shuttered growth of a Bi,Sr,CaCu,O,-like structure in which the calcium monolayer is completely replaced by strontium has been attempted.[g4jf33gj This hypothetical crystal structure is shown in Fig. 30. Such a composition may be written as Bi,Sr,SrCu,O, or simply Bi,Sr&u,O,.

Bi,Sr,Cu,O,

exists as an equilibrium

phase, but does not

have a Bi,Sr,CaCu,O,-like structure.1 350jt351)In the Bi-Sr-Ca-Cu-0 phase diagram shown in Fig. 8(a), Bi,Sr,Cu,O, is denoted as 2302. It is also present in the Bi-Sr-Cu-0 phase diagram in Fig. 8(b) and denoted as 3:2:2. Attempts to synthesize a fully Sr-for-Ca-substituted Bi,Sr,CaCu,O,-like structure by bulk techniques have failed.f351j Thus, the MBE synthesis of such a phase is a direct test of the ability of shuttered MBE to form metastable materials unit cell level.

and customize

layering

on a finer scale than a half

The x-ray diffraction pattern of this as-grown Bi,Sr,Cu,O, film is shown in Fig. 31.t g4If33g) Other than the intense SrTiO, (100) substrate peaks present, the most intense peaks are due to a c-axis oriented Bi,Sr,CaCu,O,-like structure with c-axis length of 31.3 f 0.3 A. Some of the other peaks present could be due to the equilibrium polymorph of Bi,Sr,Cu,O,, or due to Bi,Sr,CuO, (a Bi,Sr,CuO, buffer layer was grown underneath the Bi,Sr,Cu,O, film). X-ray diffraction scans of inclined film planes indicated that the a- and b-axes of the metastable Bi,Sr,CaCu,O,like structure were aligned parallel to the SrTiO, 4 1 O> directions in the plane of the substrate.fg41

This sample

K.t33gl These experimental

results indicate the ability of shuttered

exhibited

zero resistance

at 15 MBE to

customize layering on a monolayer level. This metastable Bi,Sr&uO, phase with a Bi,Sr,CaCu,O,-like structure has also been synthesized by sputtering.f37q


583

Figure 30. The crystal structure of a hypothetical Bi,Sr,CaCu,O,-related metastable phase containing no calcium. The tetragonal subcell of the Bi,Sr,CaCu,O, phase is shown with complete strontium substitution for the atoms at the calcium site.

0

20

40

60

28 (degrees) Figure 31. The O-20 x-ray diffraction scans of an as-grown film containing a metastable Bi,Sr,CaCu,O,-like Bi,Sr,SrCu,O, phase (from Ref. 339). For clarity, the intensity of the scan aligned to the SrTiO, {loo} substrate in omega is offset from the scan rocked 5” in omega off alignment to the SrTiO, substrate. The 004 peaks due to a Bi,Sr,CaCu,O,-related metastable phase are labeled (II) in addition to those due to other phases (1).

884

Molecular

3.8

Twin-Free

Beam Epitaxy

Growth

So far in this chapter, all the high T, superconductor have been oriented

epitaxial

layers.

Although

tions in the film have been strongly influenced

layers discussed

the crystallographic

direc-

by the crystallographic

form

of the substrate (epitaxy), the epitaxial films of high T, superconductor materials having lower than tetragonal symmetry (e.g., YBa,Cu,O,_, and In Bi,Sr,Ca,_,Cu,O,,+,) d iscussed so far contain many twin boundaries. addition to the reflection twins common to YBa.&u,O,+ 90” rotation twinsf56] abound in these materials. Considering the square surface net of the common substrate materials employed with equivalent a- and b-axes, it is not surprising that both possible epitaxial alignments (related by a 90” rotation twin) of the non-equivalent a- and b-axes of the lower symmetry high T, structures with the substrate axes occur. Just as growth of GaAs or SIC on slightly misoriented silicon substrates has proven effective in eliminating anti-phase boundaries,f3761[37s1 growth of high T, superconductors on vicinal surfaces can significantly reduce the concentration of twin boundaries in YBa,Cu,O,_b films,f360] and virtually eliminate them in Bi,Sr,Ca,_,Cu,O,,+, and (La,Sr),CuO, films.f340)f361)-f36g) Twin-free high T, films were first reported for the Bi,Sr&a,_, CU"O~~+~system for (001) oriented films grown on slightly misoriented SrTiO, substrates.1 3401 Twin-free growth in the Bi,Sr,Ca,_, cu ,,0 2n+4system is easily discerned from RHEED patterns, due to the long (0 26 A) incommensurate superstructure which exists along the b-axis of these structures, but is absent along the a-axis.f53]f3go]-f3g2] Note that the lattice constants of the a- and b-axes of the Bi2Sr2Ca,_,Cu,02n+4 phases themselves are very close in length (e.g., a = 5.414 A and b = 5.418 A for the orthorhombic subcell of Bi2Sr2CaCu206).f3W) The incommensurate superstructure along the b-axis results in closely spaced satellite streaks (or spots) in RHEED patterns, which are easily discerned from the other RHEED streaks because of their very close spacing. The RHEED patterns for the growth of Bi,Sr,Ca,_,Cu,O,,+, (001) films on vicinal and well oriented perovskite substrates are shown in Fig. 32. For growth on a vicinal SrTiO, (001) substrate, which is misoriented about 4” - [l lo], the RHEED patterns observed along the SrTiO, cl lO> azimuths are clearly different from each other. The closely spaced satellite streaks observed along the SrTiO, [li 0] azimuth indicate the lateral presence of the b-axis and its incommensurate superstructure, while the streaks observed along the SrTiO, [l lo] azimuth indicate the lateral presence

of the a-axis.

Note that in films grown on well oriented

High

Tc Superconductors

585

(001) surfaces viewed along the corresponding RHEED azimuths, asuperposition of these two patterns is observed along both perovskite cl 1O> azimuths,

as shown in Fig. 32(c).

The RHEED patterns observed

along

the [loo] perovskite azimuths for Bi,Sr,Ca,_,Cu,O,,+, (001) films grown on vicinal and well oriented perovskite substrates are shown in Fig. 33. The difference

between the twinned

and untwinned

films is also evident

along this azimuth.

Figure 32. RHEED patterns observed along the ~11 O> perovskite azimuths of Bi,Sr,Ca,_, CU,O~~+~films grown on vicinal and well oriented (001) perovskite substrates: (a) A Bi,Sr,Ca,Cu,O,, (n=3) film on a vicinal SrTiO, (001) substrate, SrTiO, [l lo] azimuth, @,JThe same film shown in (a), but along the SrTiO, [li 0] azimuth, and (c) A Bi,Sr,CuO, (n=l) film on a well oriented LaAIO, (091) substrate, LaAIO, [l lo] azimuth. The RHEED pattern along the LaAIO, [ll 0] azimuth was identical to that shown in (c). Images from Refs. 94 and 340.

666

Molecular

Beam Epitaxy

Figure 33. RHEED patterns observed along the perovskite azimuths of Bi,Sr&a,_, CU,O~~+~films grown on vicinal and well oriented (001) perovskite substrates: (a) A Bi,Sr,Ca,Cu,O,, (n=3) film on a vicinal SrTiO, (001) substrate, SrTiO, [loo] azimuth and (b) A Bi,Sr,CuO, (n=l) film on a well oriented LaAIO, (001) substrate, LaAIO, [loo] azimuth. The RHEED patterns along the [OlO] perovskite azimuths were identical to those shown in (a) and (b) for these two growths, respectively. Images from Refs. 94 and 340.

of rotation and reflection twin boundaries in Bi,Sr,Ca,_, films grown on vicinal perovskite substrates is also evident from CU”02n+4 x-ray diffraction measurements. Just as the incommensurate superstrucThe absence

ture along the b-axis of the orthorhombic subcell of these materials causes satellite streaks to occur in RHEED, it also produces satellite reflections in x-ray diffraction. In particular, the 0212 reflection of the Bi,Sr,Ca,Cu,O,, phase has satellite peaks on either side of it along the b*reciprocal space axis, while the 202 reflection does not have satellite peaks adjacent to it along the a* reciprocal space axis. Figure 34 shows the observed x-ray diffraction intensity gathered by aligning the detector to a satellite reflection of the 0212 peak of a Bi,Sr,Ca,Cu,O,, film grown on a vicinal SrTiO, (001) substrate and then rotating the q-axis of the 4-circle diffractometer around the [OOl] zone axis of the film. If the film contained rotation twin boundaries, this scan would have contained a diffraction peak every 90”. The presence of (110) reflection twin boundaries would also lead to x-ray diffraction peaks about every 90” since the a- and b-axes are


587

of nearly identical length in these Bi,Sr,Ca,_,Cu,O,,+, materials. Instead, the observed x-ray diffraction pattern contains only two peaks, indicating that the incommensurate

superlattice

is aligned

along

only one of the

cl lo>-type SrTiO, surface directions. Specifically, the x-ray data indicate that the &As of the film runs up and down the substrate steps and not along the length of the substrate

I

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steps.

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I

360

$I (degrees) Figure 34. An x-ray diffraction $-scan of an incommensurate superlattice satellite peak of the 0212 peak of a Bi,Sr,Ca,Cu,O,, (n=3) film grown on a misoriented SrTiO, (100) substrate. Q, = 0 was set parallel to SrTiO, . This scan indicates that the b-axis of the Bi,Sr,Ca&u,O,c film runs up and down the steps of the misoriented substrate (45” and 225“ peaks) but not laterally along them (no peak at 135’ or 315’7, indicating the absence of twin boundaries in the film (from Ref. 339). Subsequent

work

using

vicinal

surfaces

of SrTiO,

confirmed these results and examined the microstructure TEM,[381)t382)[3861but has demonstrated that this technique to fabricate

virtually

twin-free

films of Bi,Sr,Ca,_,Cu,O,,+,

has not only in detail using may be utilized (1 In) ,tsss1tss4)

(La,Sr),CuO, (103),t3s8) and even the metallic perovskite SrRuO, (11 O).tgl In an analogous manner, vicinal SrTiO, (110) surfaces have been utilized

888

Molecular

Beam Epitaxy

to prepare nearly untwinned YBa,Cu,O,B films,f3Q3t vicinal MgO (110) surfaces with CeO, buffer layers have been used to grow untwinned Bi,Sr,Ca,_, CU,O~.+~ (01 n) films[385) and vicinal LaSrGaO, (110) surfaces have enabled the growth of untwinned

Bi,Sr,CuO,

(011) films.t3srt Cleav-

age steps on MgO

(100) have also been found to reduce the 4-fold orthorhombic substrates, symmetry of the substrate.1 3941 Alternatively, and either anisotropic e.g., NdGa03f3Q5)t3Q6)and (Y,Nd)AIO,,f 1711f34~f38Q) thermal expansion coeff icientsf 3Q7) or the application of mechanical stressf3Q~f3Qs)may be used for this purpose. Note that the untwinned high T, superconductor films grown on vicinal substrates are themselves vicinal. This not only allows the anisotropy in the transport properties to be investigated,f3B1tt3QQ) but with appropriate patterning, the current may be constrained to flow up or down these steps, which for the case of Bi,Sr,Ca,_,Cu,O,,+, films, may lead to device possibilities since each step appears to be a weak link,f3QQ)and the number of weak links in series connection is determined substrate and length of the patterned region.

4.0

FUTURE

4.1

Hybrid

by the tilt of the

DIRECTIONS MBE Techniques

As the MBE technique has been applied and adapted to the growth of high T, superconductors, other thin film techniques have also been modified in order to achieve a customized layering capability of these materials. Layering control at the unit cell level has been demonstrated by reactive evaporation,t400) PLD,t401) and sputtering.1 4c21[4cs)The pressures used in ion beam sputtering are sufficiently low that RHEED has been extensively used during shuttered gr0wth.f 366)[404t Differential pumping has allowed the use of RHEED at the relatively high pressures used in reactive evaporationf126)f187] and PLD.f405) Another method enabling the use of RHEED

during

PLD is by supplying

the high oxidant

pressure

to the

substrate as brief pulses in synchronization with the PLD laser pulses. The resulting average background pressure is significantly lowered, making RHEED observation during PLD possible, as well as enabling the effects of altering the oxidation pulse timing to be studied.f‘rcc) The use of synchrotron x-ray radiation has allowed in-situ monitoring by diffraction techniques at the high pressures used in sputteringt40r) and 0MVPE.f408)


This monitoring

has included the observation

of x-ray intensity

589

oscillations

analogous to those present in l7HEED.f 408) Sequential deposition using multiple-target sputtering,f401f3661t404~f40gl-f414) multiple target PLD,t415t reactive evaporation/

roll pulsed organometallic

organic

vapor

chemical

deposition

MBE,f4161 and pulsed metal-

(MOCVD)f417)

have

been

used to

control the order in which the depositing species reach the substrate, in an effort to form customized and metastable structures in an analogous manner to MBE. While some aspects of these techniques are attractive, and the corresponding hybrid approaches (e.g., laser-MBEf41*)-[420) and organometallic MBEt416)) will likely be utilized for the MBE growth of high T, superconductors to a greater extent in the future, the layering control achieved by MBE remains unsurpassed for the controlled growth of high T, superconductors which cannot be prepared by bulk methods, and for heterostructures containing custom-layered high T, materials. Some advantages of the hybrid laser-MBE technique are its ability to deposit materials with extremely low vapor pressures for in-situ patterning, or to deposit multi-component materials from a single target with faithful composition transfer. The latter capability is advantageous when a thick layer of a material that may be readily synthesized by bulk techniques is desired in a heterostructure which contains other layers, the growth of which requires the customized laying capability of MBE. Utilizing gaseous precursors in high T, MBE has similar advantages to metal-organic molecular beam epitaxy (MOMBE) and chemical beam epitaxy (CBE) over conventional MBE. However, the lack of gaseous precursors which may be easily transported to the growth chamber and reacted, particularly for the alkaline earth elements,f421) has been a significant obstacle to the widespread use of gaseous precursors for the MBE growth of high Tc superconductors. The development of suitable gaseous precursors remains an active area of research. As progress in precursor chemistry

and their use in the growth

of high

T, superconductors

by

MOCVD is made, these gaseous reagents may be more actively used in hybrid MBE approaches to the growth of high T, superconductors, including hopefully, the growth of these materials by atomic layer epitaxy (ALE).f4**) The use of surface chemistry

in ALE growth to chemically

limit

the buildup of depositing species to a single complete monolayer has a distinct advantage over MBE, where even the most perfect composition control system results in the correct total number of species being deposited, but does not preclude their agglomeration (see Fig. 16). Clearly a deposition method in which surface absorption chemistry is used to accu-

590

Molecular

Beam Epitaxy

rately fill each monolayer

making up a custom-layered

high Tc supercon-

ductor structure (i.e., ALE), is superior to a deposition purely physical methods (i.e., MBE) are utilized. Although has been made in using absorption-controlled

epitaxial

method where some progress methods

for the

growth of the high T, superconductors (Ba,Rb)BiO, and (Ba,K)Bi0,,t115)t2g11 and MOCVD researchers are working to develop suitable ALE reagents the customized growth and find appropriate growth conditions/ 421)[42s1[424) of high T, superconductors 4.2

In-situ

Monitoring

by ALE remains a challenge. Techniques

Further enhancements to the MBE technique which allow improved reproducibility and process control during the growth of high Tc superconductors will likely involve improving the resolution and accuracy of the currently used in-situ monitoring methods as well as incorporating other promising in-situ analysis techniques. The precision of atomic absorption spectroscopy for composition control has improved significantly over the last several years, and it is currently the best composition control method for use at the relatively high oxidant pressures used in the MBE growth of high T, superconductors. With improvements in the intensity and stability of the light source used for atomic absorption spectroscopy, this technique should achieve much better composition control. The utility of RHEED oscillations as a composition control method would also be greatly enhanced if the effect of the shuttered supply of incident species on RHEED oscillations were understood sufficiently that the temporal behavior of the RHEED oscillation signal during growth could be used for composition control.

RHEED oscillations

are commonly

used as a calibration

in the MBE growth of Ill-V and other semiconductor unlike

the growth

of Ill-V

materials,

the crystal

materials. structures

method However,

and lattice

constants of the component binary oxides of high Tc superconductors are significantly different from each other and from the high T, superconductors, making it impractical to do flux calibration by RHEED oscillations of the various

binary oxide constituents

of the high T, material under study.

Two emerging in-situ composition analysis techniques which make use of the RHEED beam, and are thus quite compatible with the standard MBE geometry, are total-reflection-angle x-ray spectroscopy in RHEED (RHEED-TRAXS)t425) and reflection-electron energy loss spectroscopy (REELS).t426) Both are highly surface sensitive and give chemical information on the region of the sample upon which the RHEED beam is incident.


The former analyzes

the energy and intensity

of the x-rays emitted from

the sample, as is done in EPMA, while the latter analyzes intensity

of the electrons

elemental

detection

scattered

from the sample.

ability, the sensitivities

591

the energy and With monolayer

of these techniques

are excel-

lent, however, the accuracy of the composition determination needs to be improved significantly for these methods to be useful for composition control during the growth of high T, superconductors. RHEED-TRAXS has recently been applied to the MBE growth of high T, superconductors.f4*7 Another emerging technique for real-time composition control is the use of in-situ spectroscopic ellipsometry (SE) during thin film growth.t428~f42Q~ This optical method provides information on the optical properties of the growing film over a range of wavelengths, and its geometry is entirely compatible with MBE growth. Owing to the comparatively long penetration depth of light over the wavelength range used, SE differs from the aforementioned techniques in that it probes well into the film; it is a bulksensitive as opposed to a surface-sensitive

technique.

The variation

of the

optical properties (density and dielectric function) with thickness may be extracted from the SE data,f430) and if the film microstructure is fully dense and single phase, and the variation of the optical properties as a function of composition are known, as in the case of many semiconductor materials grown by MBE, the composition, thickness, and surface microstructure of the various layers making up the thin film sample may be calculated. The SE-determined surface microstructure may be used to distinguish between layer-by-layer and island growth, independently from RHEED. A useful feature of the SE technique is that the parameters extracted, thickness taken

and composition,

at many different

are overdetermined

wavelengths,

dense and single phase microstructure

allowing

by the ellipsometry the assumption

data of fully

to be explicitly checked. Ellipsometry

at a single wavelength has been used for closed-loop real-time composition control for the growth of Ill-V materials by MBE.f431) The use of a range of wavelengths (Le., SE), should allow the in-situ application of this technique to a broader range of materials, perhaps including the MBE growth

of high

T, superconductors.

For composition

sensitivity,

it is

desirable to perform SE in a wavelength range where the optical properties vary significantly with composition; this range lies in the infrared (IR) for many oxides. Spectroscopic ellipsometers that operate in the far-IR have been built and applied to oxides, including high Tc superconductors.(432) A related, but less complicated technique, reflectometry at a single wavelength, monolayer

has been used in the MEE growth of Ill-V materials to monitor completi0n.f 4331An abrupt change in slope of the reflectivity

as

592

Molecular

Beam Epitaxy

a function of deposited flux accompanies the completion of an MEE monolayer. The desire to accurately limit the flux deposited in each burst to an integral number of monolayers during MBE growth by sequential deposition

indicates that such optical methods may also provide a means

to better control the growth of high optical methods

T, superconductors.

Although

such

have not yet been used in the MBE growth of high

T,

superconductors, researchers growing YBa,Cu,O,, by sequential deposition MOCVD techniques[41q have utilized in-situ reflectometry to observe significant variation in the reflected light intensity accompanying the pulsed deposition and oxidation of the constituent fluxes.[424] A distinct advantage of AA, RHEED-TRAXS, REELS, and SE over other composition control methods is that they can be used to measure either the flux incident upon the substrate (AA) or flux incorporated into @HEED-TRAXS, REELS, and SE) the growing film during growth. Other composition control methods, e.g., quartz crystal microbalances, emission spectroscopic methods, mass-spectrometers, and ion gauges, when used during growth, rely on geometric factors to relate the flux at the position of the detector to the composition reaching the growing surface. Since these factors may change with time, gas pressure, and other process variables, it is advantageous to directly measure the flux reaching or, better yet, being incorporated into the film during growth. A schematic diagram of what a future MBE system for the growth of high T, superconductors might look like is shown in Fig. 35. In addition to the features common in today’s MBE systems for the growth of high T, superconductors, this system incorporates the hybrid laser-MBE technique (a combination of MBE and PLD) allowing a wider range of materials and growth conditions to be utilized as well as in-situ SE characterization.

5.0

CONCLUSIONS The use of MBE for the growth of high T, superconductors

infancy

and many

hurdles,

in particular

accurate

composition

is still in its control,

remain to be overcome for this technique to develop greater structural and doping control at the atomic layer level as well as reproducibility. Nonetheless, in the seven years since the discovery of high T, superconductivity, MBE has become established as the premiere synthesis technique for these layered oxides when customized layering control is needed. As the complexity

and metastability

of desired

structures

increases,

hopefully

High T,_-Superconductors

593

aided by increased understanding of high T, superconductivity and thus an enlightened design of new high T, materials, the requirement for a controlled synthesis superconductors important.

environment and device

MBE appears

capable of atomic layer engineering heterostructures

of new

will become

all the more

to be the most likely technique

to meet this

challenge.

“.

Absorption and u Spectroscopic Ellipsometry Light In

AI..:._.,....... ~\eCtW Gun

”

a

Mechanical

Pump

Figure 35. A schematic diagram of what added features a future MBE growth chamber might contain for the growth of high T, superconductors. The growth chamber shown contains features commonly used, plus the ability to perform combined MBE and laser-MBE, and in-situ spectroscopic ellipsometry.

594

Molecular

Beam Epitaxy

ACKNOWLEDGMENTS We gratefully

acknowledge

our many collaborators

with whom,

be-

ginning in early 1987, we have worked toward the MBE growth of high 7, superconductors, especially E. S. Hellman, I. Bozovic, and J. N. Eckstein. In addition, we acknowledge stimulating interactions with the entire Stanford high T, thin film group, and in particular Ft. H. Hammond, T. H. Geballe, and M. R. Beasley. The financial support of the Joint Services Electronics Program through contract DAAG29-84-K-0048, DARPA/ONR under contract N00014-88-C-0760,

the National

Science

Foundation

Material

Re-

search Laboratory Program through the Center of Materials Research at Stanford University, and the support of a Semiconductor Research Corporation Fellowship for DGS for the period of his thesis research are gratefully acknowledged. DGS acknowledges fruitful interactions with the IBM Zurich high T, group, especially J. Mannhart and J. G. Bednorz, and the financial support of ONR through contract preparation of this review chapter.

NO001 4-93-l -0512 during the

REFERENCES 1. Bednorz, J. G. and Miiller, Bednorz, J. G., Takashige, 3:379-386 (1987)

K. A., Z. Phys. B, 64:189-193 (1986); M., and Miller, K. A., Europhys. Leti.,

2. Batlogg, B., Cava, R. J., Schneemeyer, L. F., and Espinosa, IBM J. Res. Develop., 33:208-214 (1989)

G. P.,

3. Ramesh, R., Inam, A., Chan, W. K., Wilkens, B., Myers, K., Remschnig, K., Hart, D. L., and Tarascon, J. M., Science, 252:944946 (1991) 4. Ramesh, R., Inam, A., Char-r, W. K., Tillerot, F., Wilkens, B., Chang, C. C., Sands, T., Tarascon, J. M., and Keramidas, V. G., Appl. Phys. Lett., 59:3542-3544 (1991) 5. Ramesh, R., Chan, W. K., Wilkens, B., Gilchrist, H., Sands, T., Tarascon, J. M., Keramidas, V. G., Fork, D. K., Lee, J., and Safari, A., Appl. Phys. Lett., 61:1537-1539 (1992) 6. Fork, D. K., Ponce, F. A., Tramontana, Appl. Phys. Leti., 58:2294-2296 (1991)

J. C., and Geballe,

T. H.,

7. Ramesh, R., Inam, A., Bonner, W. A., England, P., Wilkens, B. J., Meagher, B. J., Nazar, L., Wu, X. D., Hegde, M. S., Chang, C. C., Venkatesan, T., and Padamsee, H., Thin Film.s, Appl. Phys. Leti., 55:1138-l 140 (1989)

High

8. Lichtenberg, F., Catana, A., Mannhart, Phys. Lett., 60: 1138-l 140 (1992)

Tc Superconductors

J., and Schlom,

595

D. G., Appl.

9. Eom, C. B., Cava, R. J., Fleming, R. M., Phillips, J. M., van Dover, R. B., Marshall, J. H., Hsu, J. W. P., Krajewski, J. J., and Peck, W. F., Jr., Science, 258:1766-l 769 (1992) 10. Satyalakshmi, K. M., Mallya, R. M., Ramanathan, K. V., Wu, X. D., Brainard, B., Gautier, D. C., Vasanthacharya, N. Y., and Hegde, M. S., Appl. Phys. Leti., 62:1233-1235 (1993) 11. Dimos, D., Chaudhari, 4049 (1990) 12. Suzuki, (1987)

P., and Mannhart,

M. and Murakami,

J., Phys. Rev., B41:4038-

T., Jpn. J. Appl. Phys., 26:L524-L525

13. Chaudhari, P., Koch, R. H., Laibowitz, R. B., McGuire, T. R., and Gambino, R. J., Phys. Rev. Letf., 58:2684-2686 (1987); Chaudhari, P., LeGoues, F. K., and Segmuller, A., Science, 238:342-344 (1987) 14. Naito, M., Hammond, R. H., Oh, B., Hahn, M. R., Hsu, J. W. P., Rosenthal, P., Marshall, A. F., Beasley, M. R., Geballe, T. H., and Kapitulnik, A., J. Mater. Res., 2:713-725 (1987); Oh, B., Naito, M., Arnason, S., Rosenthal, P., Barton, R., Beasley, M. R., Geballe, T. H., Hammond, R. H., and Kapitulnik, A., Appl. fhys. Lett., 51:852854 (1987) 15. Enomoto, Y., Murakami, T., Suzuki, M., and Moriwaki, Appl. fhys., 26:L1248-L1250 (1987) 16. Char, K., Kent, A. D., Kapitulnik, A., Beasley, H., Appl. Phys. Left., 51:1370-1372 (1987)

K., Jpn. J.

M. R., and Geballe, T.

17. Webb, C., Weng, S. -L., Eckstein, J. N., Missert, N., Char, K., Schlom, D. G., Hellman, E. S., Beasley, M. R., Kapitulnik, A., and Harris, J. S., Jr., Appl. Phys., Lett., 51 :1191-l 193 (1987) 18. Kwo, J., Hsieh, T. C., Fleming, R. M., Hong, M., Liou, S. H., Davidson, B. A., and Feldman, L. C., Phys. Rev., 836:4039-4042 (1987) 19. Dijkkamp, D., Venkatesan, T., Wu, X. D., Shaheen, S. A., Jisrawi, N., Min-Lee, Y. H., McLean, W. L., and Croft, M., Appl. Phys. Lett., 51:619-621 (1987) 20. See, for example: Cheung, J. and Horwitz, J., MRSBulletin, 17:3036 (Feb. 1992); Chrisey, D. B. and Inam, A., MRS Bulletin, 17:37-43 (Feb. 1992); Venkatesan, T., Wu, X. D., Muenchausen, R., and Pique, A., MRS Bulletin, 17:54-58 (Feb. 1992)

596

Molecular

Beam Epitaxy

21. Cuomo, J. J., Gambino, Ft. J., Harper, J. M. E., and Kuptsis, J. D., IBM J. Res. Develop., 21580583 (1977) 22. Gilbert, L. R., Messier, R., and Roy, R., Thin So/id Films, 54:129136 (1978); Gilbert, L. R., Messier, R., and Krishnaswamy, S. V., J. Vat. Sci. Techno/., 17:389-391 (1980) 23.

Rossnagel, S. M. and Cuomo, J. J., Thin Film Processing and Characterization of High-Temperature Superconductors, (J. M. E. Harper, R. J. Colton, and L. C. Feldman, eds.), pp.106-113, American institute of Physics, New York (1988)

24. Sandstrom, R. L., Gallagher, W. J., Dinger, T. R., Koch, R. H., Laibowitz, R. B., Kleinsasser, A. W., Gambino, R. J., Bumble, B., and Chisholm, M. F., Appl. Phys. Letf., 53:444-446 (1988) 25. Li, H.C., Linker, G., Ratzel, F., Smithey, R., and Geerk, J., Appl. Phys. Lett., 52:1098-l 100 (1988); Meyer, O., Weschenfelder, F., Geerk, J., Li, H. C., and Xiong, G. C., Phys. Rev., B37:9757-9760 (1988) 26. Xi, X. X., Linker, G., Meyer, O., Nold, E., Obst, B., Ratzel, F., Smithey, R., Strehlau, B., Weschenfelder, F., and Geerk, J., Z. Phys,, 874: 13-l 9 (1989) 27. Eom, C. B., Sun, J. Z., Yamamoto, K., Marshall, A. F., Luther, K. E., Geballe, T. H., and Laderman, S. S., Appl. Phys. Left, 55595-597 (1989) 28. Lathrop, D. K., Russek, S. E., and Buhrman, R. A., Appl. Phys. Lett., 51 :1554-l 556 (1987); Tietz, L. A., de Cooman, B. C., Carter, C. B., Lathrop, D. K., Russek, S. E., and Buhrman, R. A., J. E/e&on Microsc. Tech., 8:263-272 (1988) 29. Terashima, T., lijima, K., Yamamoto, K., Bando, Y., and Mazaki, H., Jpn. J. Appl. Phys., 27:L91-L93 (1988) 30. Wu, X. D., Inam, A., Venkatesan, T., Chang, C. C., Chase, E. W., Barboux, P., Tarascon, J. M., and Wilkens, B., Appl, fhys. Lett., 52:754-756 (1988); Hwang, D. M., Venkatesan, T., Chang, C. C., Nazar, L., Wu, X. D., lnam, A., and Hegde, M. S., Appl. Phys. Left, 54:1702-l 704 (1989) 31. Fujita, J., Yoshitake, T., Kamijo, A., Satoh, T., and Igarashi, A/I/I/. Phys., 64:1292-l 295 (1988)

H., J.

32. Koinuma, H., Kawasaki, M., Funabashi, M., Hasegawa, T., Kishio, K., Kitazawa, K., Fueki, K., and Nagata, S., Appl. Phys., 62:15241526 (1987) 33. Lee, L.P., Char, K., Colclough, M.S., and Zaharchuk, Whys. Left., 59:3051-3053 (1991)

G.,

Appl.

High

Tc Superconductors

597

34. Triscone, J. -M., Fischer, 0., Brunner, O., Antognazza, L., Kent, A. D., and Karkut, M. G., Phys. Rev. Left, 64:804-807 (1990) 35. Li, Q., Xi., X. X., Wu, X. D., Inam, A., Vadlamannati, S., McLean, W. L., Venkatesan, T., Ramesh, f?., Hwang, D. M., Martinez, J. A., and Nazar, L., Phys. Rev. Leff., 64:3086-3089 (1990) 36. Lowndes, 65:1160-l

D. H., Norton, D. P., and Budai, J. D., Phys. Rev. Leff., 163 (1990)

37. Tabata, H., Kawai, T., and Kawai, S., Phys. Rev. Leff., 70:26332636 (1993) 38. Terashima, T., Shimura, K., Bando, Y., Matsuda, Y., Fujiyama, and Komiyama, S., Phys. Rev. Left, 67:1362-1365 (1991)

A.,

39. Kitabatake, M. and Greene, J. E., J. Appl. Phys., 73:3183-3194 (1993); Kitabatake, M., Fons, P., and Greene, J. E., J. Vat. Sci. Technol., A8:3726-3735 (1990) 40. Wasa, K., Adachi, H., Hirochi, K., Ichikawa, Y., Matsushima, Setsune, K., J. Mater. Res., 6:1595-l 604 (1991) 41. Sleight, A. W., Physics Today, 44:24-30

T., and

(1991)

42. Schlom, D. G., Eckstein, J. N., Hellman, E.S., Webb, Harris, J. S., Jr., Beasley, M. Ft., and Geballe, T. Abstracts, High-Temperature Superconductors II, (D. W. H. Butler, B. Batlogg, and C. W. Chu, eds.), Materials Research Society, Pittsburgh (1988)

C., Turner, F., H., Extended W. Capone, II, pp. 197-200,

43. Spah, R. J., Hess, H. F., Stormer, H. L., White, A. E., and Short, K. T., Appl. Phys. Left., 53:441-443 (1988) 44. Schlom, D. G., Eckstein, J. N., Hellman, E. S., Streiffer, S. K., Harris, J. S., Jr., Beasley, M. R., Bravman, J. C., Geballe, T. H., Webb, C., von Dessonneck, K. E., and Turner, F., Appl. Phys. Leff., 53:1660-l 662 (1988) 45. Kwo, J., Hong, M., Trevor, D. J., Fleming, R. M., White, A. E., Farrow, R. C., Kortan, A. R., and Short, K. T., Appl. Phys. Leff., 53:2683-2685 (1988) 46. Harris, J. S., Jr., Eckstein, J. N., Hellman, J. Crysf. Growth, 95:607-616 (1989)

E. S., and Schlom, D. G.,

47. See, for example, The Technology and Physics of Molecular Beam Epifaxy, (E. H. C. Parker, ed.), Plenum Press, New York (1985) 48. Miller-Buschbaum, (1989) 49. Shannon,

H., Angew. Chem. Int. Ed. Engl., 28:1472-l

R. D., Acfa Crysf., A32:751-767

(1976)

493

598

Molecular

Beam Epitaxy

50. Tokura, Y. and Arima, T., Jpn. J. Appl. Phys., 29:2388-2402

(1990)

51. Cava, R. J., Batlogg, B., Krajewski, J. J., Farrow, R., Rupp, L. W., Jr., White, A. E., Short, K., Peck, W. F., and Kometani, T., Nature, 332:814-816 (1988) 52. Sleight, A. W., Gillson, J. L., and Bierstedt, Comm., 17:27-28 (1975)

P. E.,

53. Petricek, V., Gao, Y., Lee, P., and Coppens, P.,fhys. 392 (1990)

So/id State

Rev.,B42:387-

54. Coppens, P., Lee, P., Gao, Y., and Sheu, H. -S., J, Phys. Chem. So/ids, 52:1267-1272 (1991) 55. Cava, R. J., Science 247:656-662

(1990)

56. Sleight, A. W., Science, 242:1519-l 57. Tokura, Y., Torrance, Rev., B38:7156-7159

527 (1988)

J. B., Huang, T. C., and Nazzal, A. I., Phys. (1988)

58. Torrance, J. B., Tokura, Y., Nazzal, A. I., Bezinge, A., Huang, T. C., and Parkin, S. S. P., Phys. Rev. Lett., 61:1127-l 130 (1988) 59. Rao, C. N. R., Phil. Trans. R. Sot. London., A336595624 60. Tokura, Y., Takagi, H., and Uchida, S., Nature, 337:345-347 61. Shafer, M. W. and Penney, 27:191-220 (1990) 62. Mannhart,

63. Sleight, A. W., Physica, C162-164:3-7 S. A. and Voronin,

65. Karen, P., Braaten, 52:805-840 (1992)

(1992)

(1989)

G. F., f’hysica,

O., and Kjekshus,

A.,

C178:213-220 Acfa

Chem.

66. Ahn, B. T., Lee, V. Y., Beyers, R., Gtir, T. M., and Huggins, Physica, Cl 67:529-537 (1990) 67. Beyers, (1991)

R. and Ahn, B. T., Annu.

68. Majewsky, 3:488-491

(1989)

T., Eur. J. Solid State Inorg. Chem.,

J., Mod. Phys. Lett., B6:555-571

64. Degterov,

(1991)

P., Hettich, B., Schulze, (1991)

Rev. Mater.

(1991) Stand,, R. A.,

Sci., 21:335-372

K., and Petzow, G., Adv. Mater.,

69. Phase Diagrams for High T, Superconductors, (J. D. Whitler, and R. S. Roth, eds.), pp. 136-145, Amer. Ceramic Sot., Westerville (1991) See in particular, figures S-198(B), S-205, S-21 1 (B), and S215(A). 70. Rubin, L. M., Orlando, T. P., Vander Sande, J. B., Gorman, G., Savoy, R., Swope, R., and Beyers, R., Appl. Phys. Lett., 61:19771 979 (1 992)


599

71. Tarascon, J. M., McKinnon, W. Ft., Barboux, P., Hwang, D. M., Bagley, B. G., Greene, L. H., Hull, G. W., LePage, Y., Stoffel, N., and Giroud, M., Phys. Rev., B38:8885-8892 (1988) 72. See, for example, Eibl, O., Physica, Cl68:249-256 (1990). In this work, TEM analysis of samples produced by bulk synthesis methods yielding a majority of the n=3 Bi,Sr,Ca,_,Cu,O,,+, phase revealed numerous intergrowths of n=l to n=7. 73. See, for example, Raveau, B., Michel, C., and Hervieu, M.,Advances in Superconductivity: Proceedings of the 1st International Symposium on Superconductivity (ES ‘&I), (K. Kitazawa, and T. Ishiguro, eds.), pp. 151-157, Springer-Verlag, Tokyo (1989). In this work, TEM analysis of samples produced by bulk synthesis methods yielding a majority of the n=4 TI,Ba,Ca,_,Cu,O,,+, phase revealed numerous intergrowths up to n=7. 74. Senaris-Rodriguez, and Alario-France,

M. A., Chippindale, A. M., Varez, A., Moran, E., M. A., Physica, Cl 72:477-480 (1991)

75. Cava, R. J., Siegrist, T., Hessen, B., Krajewski, J. J., Peck, W. F., Jr., Batlogg, B., Takagi, H., Waszczak, J. V., Schneemeyer, L. F., and Zandbergen, H. W., J. So/id State Chem., 94:170-l 84 (1991) 76. Fu, W. T., Zandbergen, H. W., Xu, Q., van Ruitenbeek, J. M., de Jongh, L. J., and van Tendeloo, G., So/id State Comm., 70:11171121 (1989) 77. Holstein, W. L., Parisi, L. A., Fincher, C. R., and Gai, P. L., Physica C, 212:110-118 (1993) 78. Hyde, B. G. and Andersson, S., inorganic I nterscience, New York (1989)

Crystal Structures,

Wiley-

79. Udayakumar, K. R. and Cormack, A. N., J. Am. Ceram. Sot., 7l:C469-C-471 (1988); Udayakumar, K. R. and Cormack, A. N., J. Phys. Chem. So/ids, 50:55-60 (1989) 80. Tilley, R. J. D., J. So/id State Chem., 21:293-30

(1977)

81. Gossard, A. C., Petroff, P. M., Weigmann, W., Dingle, Savage, A., Appl. Phys. Left., 29:323-325 (1976)

R., and

82. Sakamoto, T., Funabashi, H., Ohta, K., Nakagawa, T., Kawai, N. J., Kojima, T., and Bando, Y., Superlattices Microstruct., 1:347-352 (1985) 83. Kita, R., Hase, T., Itti, R., Sasaki, M., Morishita, Appl. Phys. Left., 60:2684-2685 (1992) 84. Berkley, D. D., Goldman, A. M., Johnson, Wang, T., Rev. Sci. Instrum., 60:3769-3774

T., and Tanaka,

B. R., Morton, (1989)

S.,

J., and

600

Molecular

Beam Epitaxy

85. Schlom, D. G., Lee, W. S., Ma, T., and Harris, J. S., Jr., J. Vat. Sci. Technol., B7:296-298 (1989) 86. Coleman, E., Siegrist, 19:235-239 (1990) 87.

T., and Yeh,

J. J.,

J. Necfron.

The Oxide Handbook, 2nd ed., (G. V. Samsonov, Plenum, New York (1982)

Mater.,

ed.) p. 153, IFI/

88. Schlom, D. G., Marshall, A. F., Sizemore, J. T., Chen, Z. J., Eckstein, J. N., Bozovic, I., von Dessonneck, K. E., Harris, J. S., Jr., and Bravman, J. C., J. Cryst. Growth, 102:361-375 (1990) 89. Johnson, B. R., Beauchamp, K. M., Wang, T., Liu, J-X., McGreer, K. A., Wan, J-C., Tuominen, M., Zhang, Y-J., Mecartney, M. L., and Goldman, A. M., Appl. Phys. Lett., 56:1911-1913 (1990) 90. Wang, H. S., Eissler, D., Dietsche, W., Fischer, A., and Ploog, K., J. Cryst. Growth, 126:565-577 (1993) 91. Nakayama, Y., Tsukada, 70:4371-4377 (1991)

I., and Uchinokura,

K., J, Appl.

Phys.,

92. Appelboom, H. M., Growth and Properties of Co-Evaporated Superconducting YBaCuO Thin Films, Ph.D. Thesis, Delft University of Technology, Delft (1992) 93. Matijasevic,

V. C., private communication

94. Schlom, D. G., Molecular Beam Epitaxial Growth of Cuprate Superconductors and Related Phases, Ph.D. Thesis, Stanford University, Stanford (1990) 95. Welch, K. M., Capture Pumping Technology: 203-205, Pergamon Press, Oxford (1991)

An Introduction,

pp.

96. Hellman, E. S., Hartford, E. H., and Fitzgerald, 7:795-800 (1992)

E. A., J. Mater, Res.,

97. Humphreys, R. G., Satchell, J. S., Chew, N. G., Edwards, J. A., Goodyear, S. W., Blenkinsop, S. E., Dosser, 0. D., and Cullis, A. G., Sci. Technol., 3:38-52 (1990); Humphreys, R. G., Satchell, J. S., Chew, N. G., Edwards, J. A., Goodyear, S. W., Keene, M. N., and Hedges, S. J., Mater. Sci. Eng., B10:293-303 (1991) 98. Hammond,

R. H., private communication

99. Haynes’“’ is a registered Kokomo, Indiana.

trademark

100. Inconel’“’ is a registered trademark Huntington, West Virginia. 101.

of Haynes

International,

Inc.,

of lnco Alloys International,

Inc.,

lijima, K., Terashima, T., Bando, Y., Kamigaki, J. Appl. Phys., 72:2840-2845 (1992)

K., and Terauchi,

H.,


102. Terashima, T., Bando, Hayashi, K., Kamigaki, 6526842687 (1990)

601

Y., lijima, K., Yamamoto, K., Hirata, K., K., and Terauchi, H., Whys. Rev. Lett.,

103. Gerber, Ch., Anselmetti, D., Bednorz, J. G., Mannhart, Schlom, D. G., Nature, 350:279-280 (1991) 104. Hawley, M., Raistrick, I. D., Beery, Science, 251:1587-l 589 (1991)

J., and

J. G., and Houlton,

R. J.,

105. Heinzelmann, H., Anselmetti, D., Wiesendanger, R., Girntherodt, H. -J., Kaldis, E., and Wisard, A., Appl. Phys. Left., 53:2447-2449 (1988) 106. Lang, H. P., Ramseyer, J. P., Brodbeck, D., Frey, T., Karpinski, J., Kaldis, E., and Wolf, Th., U/?ramicroscopy, 42-44:715-720 (1992) 107. Kanai, M., Kawai, T., and Kawai, S., Appl. Phys. Lett., 58:771-773 (1991) 108. Horikoshi, Y., Kawashima, M., and Yamaguchi, H., Jpn. J. Appl. Phys., 25:L868-L870 (1986); Horikoshi, Y., Kawashima, M., and Yamaguchi, H., Jpn. J, Appl. Phys., 27:169-179 (1988) 109. Bauer, E., Techniques of Metals Research, (R. F. Bunshah, ll(2):501-558, Interscience, New York (1969)

ed.)

110. Lagally, M. G., Savage, D. E., and Tringides, M. C., Reflection HighEnergy Electron Diffraction and Reflection Electron Imaging of Surfaces, (P. K. Larsen and P. J. Dobson, eds.) pp. 139-l 74, Plenum Press, New York (1988) 111. Lindemer, T. B., Washburn, F. A., MacDougall, C. S., Feenstra, and Cavin, 0. B., Physica, Cl78:93-104 (1991)

R.,

112. Matijasevic, V., Rosenthal, P., Shinohara, K., Marshall, A. F., Hammond, R. H., and Beasley, M. R., J. Mater. Res., 6:682-698 (1991) 113. Humphreys, R. G., Chew, N. G., Satchell, J. S., Goodyear, S. W., Edwards, J. A., and Blenkinsop, S. E., /EEE Trans. Magn., 27:13571360 (1991) 114. Degterov,

S. A. and Voronin,

115. Hellman, E. S. and Hartford, 335 (1990) 116. Enomoto, Y., Murakami, 28:Ll355-1357 (1989)

G. F., Physica, C208:403-411 E. H., J. Vat. Sci. Technol.,

T., and Moriwaki,

(1993) B8:332-

K., Jpn. J. Appl. Phys.,

117. Hellman, E. S. and Hartford, E. H., Jr., Physica, C190:3134

(1991)

118. Cho, A. Y., The Technology and Physics of Molecular Beam Epitaxy, (E. H. C. Parker, ed.), pp. l-13, Plenum Press, New York (1985)

602

Molecular

Beam Epitaxy

119. Yoshida, S., Critical ReviewsTM in Solid State and Materials Sciences, (D. E. Schuele and R. W. Hoffman, eds.), 11:287-316, CRC Press, Boca Raton (1984) 120. Terashima, T., Bando, Y., lijima, K, Yamamoto, K., Hirata, K., Hayashi, K., Matsuda, Y., and Komiyama, S., Appl. Phys. Lett., 56:677-679 (1990) 121. Coleman, E., Siegrist, T., Mixon, D. A., Trevor, P. L., and Trevor, D. J., J. Vat. Sci. Techno/., A9:2408-2409 (1991) 122. Yokoyama, S., Ishibashi, T., Yamagami, J. Appl. Phys., 3O:L106-L109 (1990)

M., and Kawabe,

M., Jpn.

123. Satoh, T., Yoshitake, T., Miura, S., Fujita, J., Kubo, Y., and Igarashi, H., Appl. fhys. Lett., 55:702-704 (1989) 124. Silver, R. M., Ogawa, E. T., Pan, S., and de Lozanne, Trans. Magn., 27:1215-1218 (1991) 125. Berberich, P., Tate, J., Dietsche, Lett., 53:925-926 (1988)

W., and Kinder,

A. L., IEEE

H., Appl. Phys.

126. Baudenbacher, F., Karl, H., Berberich, P., and Kinder, Common Met., 164-l 65:269-278 (1990)

H.,J. Less-

127. Schindler, W., van Hasselt, P., Tontsch, P., Markl, J., Burger, J., Bauer, P., and Saemann-lschenko, G., J. Cryst. Growth, 127:10881090 (1993). 128. Herman, M. A. and Sitter, H., Molecular Beam Epitaxy: fundamentals and Current Status, pp. 3-4, 31, Springer-Verlag, Berlin (1989) 129. Do, K. B., Arnason, S. B., Carey, G. P., Ahn, C. H., Beasley, M. R., Geballe, T. H., and Hammond, R. H., “Phase Stability of CuO and Cu,O Films Grown by Reactive Electron Beam Evaporation Using an ECR Atomic Oxygen Source,” submitted to Appl. Phys. Lett., (1993) The O+ partial pressure data from Fig. 3 of this reference was corrected for the beam nature of the incident fluxes (see Sec. 2.5) before plotting this data in Fig. 19. 130. Schlom, D. G., Eckstein, J. N., Bozovic, I., Chen, Z. J., Marshall, A. F., von Dessonneck, K. E., and Harris, J. S., Jr., SPIE, 1285:234247, SPIE, Bellingham (1990); SP/E1287:152-165, SPIE, Bellingham (1990) 131. Ogihara, M., Abe, H., and Yamada, T., Jpn. J. Appl. Phys. 3O:L703L705 (1991). The NO, and 0, partial pressure data from Fig. 2 of this reference was corrected for the beam nature of the incident fluxes (see Sec. 2.5) before plotting this data in Fig. 19. For this correction, it was assumed that the NO, and 0, introduction tubes were perpendicular to the substrate surface (O=O’).


603

132. Nonaka, H., Shimizu, T., Hosokawa, S., Ichimura, S., and Arai, K., S/A, Surf. interface Anal., 19:353-357 (1992). The pressure of NO, at the substrate position (110 mm from the NO, source) used in the copper oxidation experiments described in this reference and plotted in Fig. 19 was estimated from the NO, pressure measurements at a distance of 70 mm from the NO, source. 133. Chase, M. W., Jr., Davies, C. A., Downey, J. R., Jr., Frurip, D. J., McDonald, R. A., and Syverud, A. N., JANAF Thermochemical Tab/es, 3rd ed., American Chemical Society and American Institute of Physics (1985) 134. American institute of Physics Handbook, 3rd ed., (D. E. Gray, ed.), pp. 7-6-7-8, McGraw-Hill, New York (1972) 135. Landolt-Bornstein: Zahlenwene und Funktionen aus Physik, Chemie, Astronomic, Geophysik, und Technik, (A. Eucken, ed.), 1(1):369, Springer-Verlag, Berlin (1950) 136. Prasad, R., Venugopal, Thermodynamics, 9593-601 137. Schilling, A., Cantoni, 36356-58 (1993)

J.

Chem.

M., Guo, J. D., and Ott, H. R.,

Nature,

138. Hashimoto, T., Koinuma, 30:1685-l 686 (1991)

V., and (1977)

Sood,

H., and Kishio,

D.

D.,

K., Jpn. J. Appl. Phys.,

139. Kubinski, D. J., Hoffman, D. W., Soltis, R. E., and Logothetis, J. Appl. Phys., 71 :1860-l 867 (1992) 140. Collins, 141. Schwarz,

E. M.,

D. M., private communication. H., Rev. Sci. Instrum., 32:194-199

(1961)

142. Loeb, L.B. The Kinetic Theory of Gases: Being a Textand Reference Book Whose Purpose Is to Combine the Classical Deductions with Recent Experimental Advances in a Convenient Form for Student and Investigator, 2nd ed., pp. 19, 106, McGraw-Hill, New York (1934) 143. Hellman, E. S., Schlom, D. G., Missert, N., Char, K., Harris, J. S., Jr., Beasley, M. R., Kapitulnik, A., Geballe, T. H., Eckstein, J. N., Weng, S. -L., and Webb, C., J. Vat. Sci. Techno/., B6:799-803 (1988) 144. Yadavalli, S., Yang, M. H., and Flynn, C. P., Phys. Rev., B41:79617963 (1990) 145. McKee, R. A., Walker, F. J., Conner, J. R., Specht, Zelmon, D. E., Appl. Phys. Lett., 59:782-784 (1991)

E. D., and

146. Muthe, K. P., Vyas, J. C., Kothiyal, G. P., Gandhi, D. P., Debnath, A. K., Gupta, S. K., Sabharwal, S. C., and Gupta, M. K., J. Cryst. Growth, 118:213-217 (1992)

604

Molecular

Beam Epitaxy

147. Bormann, R. and Ndlting, J., Appl. Phys. Lett, 54:2148-2150 148. Bormann,

R. and Ntilting, J., Physica, C162-164:81-82

(1989)

(1989)

149. Miller, D. J., Chiarello, R. P., Kim, H. K., Roberts, T., You, H., Kampwirth, R. T., Gray, K. E., Zheng, J. Q., Williams, S., Chang, R. P. H., and Ketterson, J. B., Appl. Phys. Letf., 59:317&3176 (1991) 150. Missert, N., Hammond, R., Mooij, J. E., Matijasevic, V., Rosenthal, P., Geballe, T. H., Kapitulnik, A., Beasley, M. R., Laderman, S. S., Lu, C., Garwin, E., and Barton, R., /EEE Trans. Magn., 252418 2421 (1989) 151. Yamamoto, K., Hammond, R. H., and Harris, J. S., Jr., J. Vat. Sci. Techno/,, A9:2587-2593 (1991) 152. Chew, N. G., Goodyear, S. W., Edwards, J. A., Satchell, J. S., Blenkinsop, S. E., and Humphreys, R. G.,App/. Phys. Leti., 57:2016201 8 (1990) 153. Kita, R., Hase, T., Takahashi, H., Kawaguchi, J. Mater. Res., 8:321-323 (1993)

K., and Morishita,

T.,

154. Eckstein, J. N., Bozovic, I., von Dessonneck, K. E., Schlom, D. G., Harris, J. S., Jr., and Baumann, S. M., Appl. Phys. Lett., 57:931-933 (1990) 155. Kucera, J. T., Perkins, J. D., Uwai, K., Graybeal, T. P., Rev. Sci. Insfrum., 62:1630-1632 (1991)

J. M., and Orlando,

156. Matijasevic, V. C., Appelboom, H. M., Mathu, F., Hadley, P., van der Marel, D., and Mooij, J. E., /EEE Trans. Appl. Supercond., 3:15241527 (1993) 157.

Ichimura, S., Hosokawa, Technol., A9:2369-2373

S., Nonaka, (1991)

H., and Arai, K., J. Vat. Sci.

158. Norimoto, K., Sekine, R., Mori, M., Hanada, T., Kudo, Kawai, M., Appl. Phys. Lett., 61:1971-1973 (1992)

M., and

159. Horvath, M., Bilitzky, L., and Huttner, J., Ozone, pp. 56-57, Elsevier, Amsterdam (1985), and references therein. 160. Watanabe, S., Kawai, M., and Hanada, 29:Lllll-L1113 (1990)

T.,

Jpn. J. Appl.

161. Kawai, M., Mori, M., Hanada, T., Kudo, M., Ishizawa, and Teratani, S., Physica, Cl 90:27-30 (1991) 162. Hosokawa, (1991)

S. and Ichimura,

S., Rev. Sci. lnsfrum.,

163. Luzeau, P., Xu, X. Z., Lagues, M., Hess, N., Contour, M., Queyroux, F., Touzeau, M., and Pagnon, D., Technol., A8:3938-3940

232, Phys.,

N., Goda, T., 62:1614-1619 J. P., Nanot, J, Vat. Sci.

High

Tc Superconductors

605

164. Hill, D. M., Meyer, H. M., III, Weaver, J. H., and Nelson, D. L., Appl. Phys. Leti., 53:1657-l 659 (1988) 165. Stall, Ft. A., J. Vat. Sci. Technol., B1:135-137 166. Berkley, D. Conroy, L. Mecartney, Phys. Lett.,

(1983)

D., Johnson, B. R., Anand, N., Beauchamp, K. M., E., Goldman, A. M., Maps, J., Mauersberger, K., M. L., Morton, J., Tuominen, M., and Zhang, Y-J., Appl. 53:1973-l 975 (1988)

167. Nakayama, Y., Ochimizu, H., Maeda, A., Kawazu, A., Uchinokura, K., and Tanaka, S., Jpn. J. Appl. Phys., 28:L1217-L1219 (1989) 168. Shimizu, T., Nonaka, H., Hosokawa, S., Ichimura, Physica, Cl 85-l 89:2003-2004 (1991)

S., and Arai, K.,

169. Achutharaman, V. S., Beauchamp, K. M., Chandrasekhar, N., Spalding, G. C., Johnson, B. R., and Goldman, A. M., Thin So/id Films, 216:14-20 (1992) 170. Appelboom, H. M., Matijasevic, V. C., Mathu, F., Rietveld, G., Anczykowski, B., Peterse, W. J. A. M., Tuinstra, F., and Mooij, J. E., Physica C, 214:323-334( 1993) 171. Tsukada, I., Kobayashi, S., Koyama, I., Watanabe, H., and Uchinokura, K., Jpn. J. Appl. fhys., 31 :L1683-L1685 (1992) 172. Eckstein, J. N., Bozovic, I., Klausmeier-Brown, M. E., Virshup, G. F., and Rails, K. S., MRS Bulletin, 17:27-33 (Aug. 1992); Bozovic, I., Eckstein, J. N., Virshup, G. F., Chaiken, A., Wall, M., Howell, R., and Fluss, M., J. Supercond., 7:187-l 95 (1994) 173. Shimizu, T., Nonaka, H., and Arai, K., Appl. Phys. Lett., 59:600-602 (1991) 174. Nonaka, H., Shimizu, 2852 (1990)

T., and Arai, K., Appl. Phys. Lett., 57:2850-

175. Kawai, M., Watanabe, S., and Hanada, T., J. Cryst. Growth, 112:745752 (1991) 176. Cook, G. A., Kiffer, A. D., Klumpp, C. V., Malik, A. H., and Spence, L. A., Ozone Chemistry and Technology, 21:44-52, Advances in Chemistry Series, Amer. Chem. Sot., Washington D.C. (1959) 177. Horvath, M., Bilitzky, L., and Hiittner, J., Ozone, pp. 215-230, Elsevier, Amsterdam (1985), and references therein. 178. Jenkins, A. C., Ozone Chemistry and Technology, 21: 13-21, Advances in Chemistry Series, Amer. Chem. Sot., Washington D.C. (1959) 179. Wailer, J. G. and McTurk, G., J. Appl. Chem., 15:363-366

(1965)

606

Molecular

Beam Epitaxy

180. Teflon@ is a registered trademark Inc., Wilmington, Delaware.

of E. I. Du Pont de Nemours & Co.,

181. Kel-F@ is a registered

trademark

of 3M, St. Paul, Minnesota.

182. Horvath, M., Bilitzky, Amsterdam (1985)

L., and Huttner,

J., Ozone, p. 236, Elsevier,

183. Carlson, D. J., Siegal, M. P., Phillips, J. M., Tiefel, Marshall, J. H., J. Mater. Res., 52797-2801 (1990)

T. H., and

184. Giedd, G. R. and Perkins, (1960)

31:773-775

185. Sauerbrey, 186.

M. H., Rev. Sci. Instrum.,

G., Z. Phys., 155:206-222

(1959)

Lu, C. and Lewis, O., J. Appl. Phys., 43:4385-4390 (1972); Lu, C., J. Vat. Sci. Technol., 12:578-583 (1975); Applications of Piezoelectric Quartz Microbalances, Vol. 7, Methods and Phenomena: Their Applications in Science and Technology Series, (C. Lu and A. W. Czanderna, eds.), Elsevier, Amsterdam (1984)

187. Terashima, T., lijima, K., Yamamoto, K., Hirata, K., Bando, Y., and Takada, T., Jpn. J. Appl. Phys., 28:L987-L990 (1989) 188. Gogol, C. A. and Cipro, C., Surf. Coat. Technol., 37:339-348 189. Hellman, E. S., Hartford, E. H., and Fleming, Lett., 55:2120-2122 (1989)

(1989)

R. M., Appl. Phys.

190. Chang, L. L., Esaki, L., Howard, W. E., and Ludeke, R., J. Vat. Sci. Technol., 10: 1 l-l 6 (1973) 191. Fischer, H., Grob, M., Peter, G., and Koprio, Technol., A6:2103-2108 (1988)

J. A.,

J. Vat.

Sci.

192. Appelboom, H. M., Hadley, P., van der Marel, D., and Mooij, J. E., /EEE Trans. Magn., 27: 1467-l 470 (1991) 193. Lu, C., Lightner, M. J., and Gogol, 14: 103-l 07 (1977)

C. A.,

J. Vat.

Sci. Technol.,

194. Kobayashi, M., Kiyonaga, B., Ohsako, N., and Hayashi, Y., Deposition Rate Monitoring of Oxide Superconductor Film by EIES Sensor, in: 5th International Workshop on Future Electron Devices HighTemperature Superconducting Electron Devices (1988) 195. Naito, M., Physica, Cl 85-l 89:1977-l

978 (1991)

196. Sakai, J., Chen, G., Hirama, K., Murakami, Appl. Phys., 27:319-322 (1988)

S., and Ishida, T., Jpn. J.

197. Sakai, J., Kato, K., Hirama, K., Murakami, C/yst. Growth, 95621-624 (1989)

S., and Ishida,

T.,J.

High

Tc Superconductors

607

198. Stirling, A. J. and West-wood, W. D., J. Appl. Phys., 41:742-748 (1970); Kometani, T. Y. and Wiegmann, W., J. Vat. Sci. Technd, 12:933-936 (1975) 199. Lu, C., Missert, N., Mooij, J. E., Rosenthal, P., Matijasevic, V., Beasley, M. IX, and Hammond, R. H., High T, Superconducting Thin Fi/rns, Devices, andApp/ications, (G. Margaritondo, R. Joynt, and M. Onellion, eds.) pp.1 63-l 70, American Institute of Physics (Conf. Proc. #182), New York (1989); Lu, C. and Guan, Y., “An Improved Method of Non-Intrusive Deposition Rate Monitoring by Atomic Absorption Spectroscopy for Physical Vapor Deposition Processes,” to be published in J. Vat. Sci. Technol. A. 200.

Cabanel, R., Schuhl, Physica, Cl 80:38-41

A., Crete, D. G., Tyc, S., and Touzeau, (1991)

201.

Klausmeier-Brown, G. F., Appl. fhys.

202.

Benerofe, S. J., Ahn, C. H., Wang, M. M., Kihlstrom, K. E., Do, K. B., Arnason, S. B., Fejer, M. M., Geballe, T. H., Beasley, M. R., and Hammond, R. H., J. Vat. Sci. Technol. B, 12:1217-1220 (1994)

203.

Hammond,

204.

Hellman, E. S., Schlom, D. G., Marshall, A. F., Streiffer, S. K., Harris, J. S., Jr., Beasley, M. R., Bravman, J. C., Geballe, T. H., Eckstein, J. N., and Webb, C., J. Mater. Res., 4:476-495 (1989)

205.

Hellman, E. S. and Hartford, 12:1178-1180(1994)

M. E., Eckstein, J. N., Bozovic, Lett., 60:657-659 (1992)

R. H., /fEE

Trans. Magn., 11:201-207

M.,

I., and Virshup,

(1975)

E. H., Jr., J. Vat. Sci. Technol.

f3,

206. Winefordner, J. D. and Vickers, T. J., Anal. Chem., 36:161-165 (1964); McClintock, J. A. and Wilson, R. A., J. Cryst. Growth, 81 :177-l 80 (1987) 207.

Matijasevic, V., Garwin, E. L., and Hammond, Instrum., 61:1747-l 749 (1990)

208.

Clampitt, (1988)

209.

Brown, J. D. and Packwook, (1982)

210.

Bastin, G. F., Heijligers, 8:45-67 (1986)

211. Willich, 212.

R. and Hanley,

P. E., Supercond,

R. H.,

Rev. Sci.

Sci. Techno/,,

R. H., X-Ray Spectrom.,

1:5-6

11 :187-l

93

H. J. M., and van Loo, F. J. J., Scanning,

P. and Obertop, D., Philips Tech. Rev., 42:162-171

Brown, J. D. and Packwood, (1986)

R. H., Appl. Surf

(1986)

Sci., 26:294-305

608

Molecular

Beam Epitaxy

213.

Bozovic,

I., private communication.

214.

DeVries, R. C., Cubic Boron Nitride: Handbook General Electric, Schenectady (1972)

215.

Yasinskaya,

216.

Naidich, Yu. V. and Kolesnichenko, G. A., Poroshk. Metal/. Mater, Vses. Konf. 9”‘, (Proc. 9’h All-Union Conf. on Powder Metallurgy) 142-l 45 (1968)

217.

Zhu, W. J., Yue, J. J., Huang, C205:118-122 (1993)

218.

Li, R. K., Kremer, R. K., and Maier, J., Physica, C213:26-32 (1993); Uehara, M., Nakata, H., and Akimitsu, J., Physica C, 216:453-457 (1993)

219.

Holland, L., Vacuum Deposition New York (1961)

220.

Glang, R., Handbook of Thin film Technology, (L. I. Maissel, and R. Glang, eds.) p.l-37, McGraw-Hill, New York (1970)

221.

The Oxide Handbook, 2nd ed., (G. V. Samsonov, Plenum, New York (1982)

222.

Davies, G. J. and Williams, D., The Technology and Physics of Molecular Beam Epitaxy, (E. H. C. Parker, ed.) pp. 15-46, Plenum Press, New York (1985)

223.

Borodin, V. A., Sidorov, V. V., Steriopolo, A., J. Cryst. Growth, 82:89-94 (1987)

224.

Lamoreaux, 16:419-443

225.

Farrow, R. F. C. and Williams, (1978)

226.

Yvon, K. and Francois,

227.

Cava, R. J., Santoro, A., Johnson, Rev., B35:6716-6720 (1987)

228.

Maeta, H., Kato, T., Ono, F., Haruna, K., Tanaka, I., and Kojima, H., Physica, Cl 85-l 89:1383-l 384 (1991)

229.

Fleming, R. M., Batlogg, B., Cava, R. J., and Rietman, Rev., B35:7191-7194 (1987)

230.

Beyers, R. and Shaw, T. M., So/id State Phys., 42:135-212 and references therein.

231.

Specht, E. D., Sparks, C. J., Dhere, A. G., Brynestad, J., Cavin, 0. B., Kroeger, D. M., and Oye, H. A., Phys. Rev., B37:7426-7434 (1988)

G. A., Refractories,

2:70-72

of Properties,

p. 5,

(1965)

Y. Z., and Zhao, Z. X.,

Physica,

of Thin Films, pp. 11 O-l 11, Wiley,

R. H. and Hildenbrand, (1987)

ed.) p. 401, IFI/

T. A., and Tatarchenko,

V.

D. L., J. Phys. Chem Ref. Data,

G. M., Thin So/id Films, 55:303-315

M., Z. Phys., B76:413-444

(1989)

D. W., Jr., Rhodes, W. W., Phys.

E. A., Phys. (1989),

High

232. Arendt, R. Jorgensen, Arendt, R. Jorgensen, (1992)

Tc Superconductors

609

H., Garbauskas, M .F., Meyer, C. A. Rotella, F. J., J. D., and Hitterman, R. L., Physica, C182:73-78 (1991); H., Garbauskas, M. F., Meyer, C. A. Rotella, F. J., J. D., and Hitterman, R. L., Physica, C194:397-402

233. Schneemeyer, L. F., Thomas, J. K., Siegrist, T., Batlogg, B., Rupp, L. W., Opila, R. L., Cava, R. J., and Murphy, D. W., Nature, 335421-423 (1988) 234.

Fleming, R. M., Marsh, P., Cava, R. J., and Krajewski, Rev., B38:7026-7028 (1988)

J. J., Phys.

235.

Landolt-Blirnstein: Numerical Data and Functional Relationships in Science and Technology, New Series, (K.-H. Hellwege, ed.) Group III, 7:134, Springer-Verlag, Berlin (1975)

236. Giess, E. A., Sandstrom, R. L., Gallagher, W. J., Gupta, A., Shinde, S. L., Cook, R. F., Cooper, E. I., O’Sullivan, E. J. M., Roldan, J. M., Segmijller, A. P., and Angilello, J., IBM J. Res. Develop., 34:916926 (1990) 237. Taylor, D., Trans. J. Br. Ceram. Sot., 84:181-l

88 (1985)

238. Touloukian, Y. S., Kirby, R. K., Taylor, R. E., and Lee, T. Y. R., Thermal Expansion: Nonmetallic Solids, 13:570, Thermophysical Properties of Matter, Plenum, New York (1977) 239. Talvacchio, J., Wagner, G. R., and Talisa, S. H., MicrowaveJournal, 34:105-l 14 (July, 1991) 240. Geller, S. and Raccah, P. M., Phys. Rev,. B2:1167-1172

(1970)

241. Simon, R. W., Platt, C. E., Lee, A. E., Lee, G. S., Daly, K. P., Wire, M. S., Luine, J. A., and Urbanik, M., Appl, Phys. Lett., 53:2677-2679 (1988) 242. Geller, S., Curlander, 644 (1974) 243.

P. J., and Ruse, G .F., Mat. Res. Bull, 9:637-

Sasaura, M., Miyazawa, 3644 (1990)

S., and Mukaida, M.,J. Appl. Phys., 68:3643-

244. Young, K. H., Negrete, G. V., Eddy, M. M., Sun, J. Z., James, T. W., Robinson, McD., and Smith, E. J., Thin So/id Films, 206:116-121 (1991) 245.

Koren, G., Gupta, A., Giess, E. A., Segmiiller, B., Appl. Phys. Lett,, 54:1054-1056 (1989)

246.

Ingel, R. P. and Lewis, (1986)

A., and Laibowitz,

D., III, J. Am. Ceram.

Sot.,

R.

69:325-332

610

Molecular

Beam Epitaxy

247.

Nielsen, T. H. and Leipold, M. H., J. Am. Ceram. Sot., 47:155 (1964)

248.

Tietz, L. A., Carter, C. B., Lathrop, D. K., Russek, S. E., Buhrman, A., and Michael, J. R., J. Mater. Res., 4:1072-l 081 (1989)

249.

Landok-BBrnstein: Numerical Data and Functional Relationships in Science and Technology, New Series, Group III, (K. -H. Hellwege, ed.) 7b:26, Springer-Verlag, Berlin (1975)

250.

Touloukian, Y. S., Kirby, R. K., Taylor, R. E., and Lee, T. Y. R., Thermal Expansion: Nonmefallic Solids, 13~288, Thermophysical Properties of Matter, Plenum, New York (1977)

251.

Landolt-Blirnstein: Numerical Data and Functional Relationships in Science and Technology, New Series, Group III, (Hellwege, K. -H., ed.) 7b:46, Springer-Verlag, Berlin (1975)

R.

252. Touloukian, Y. S., Kirby, R. K., Taylor, R. E., and Lee, T. Y. R., Thermal Expansion: Nonmetallic Solids, 13: 176, of Thermophysical Properties of Matter, Plenum, New York (1977) 253. Talvacchio, J., Wagner, 164:659-660 (1989)

G. R., and Pohl, H. C.,

Physica,

C162-

254.

Char, K., Colclough, M. S., Garrison, S. M., Newman, Zaharchuk, G., Appl. Phys. Lett,. 59:733-735 (1991)

N., and

255.

Wu, X. D., Muenchausen, R. E., Nogar, N. S., Pique, A., Edwards, R., Wilkens, B., Ravi, T. S., Hwang, D. M., and Chen, C. Y., Appl. Phys. Lett., 58:304-306 (1991)

256.

Merchant, P. Jacowitz, R. D., Tibbs, K., Taber, R. C., and Laderman, S. S., Appl. Phys. Lett., 60:763-765 (1992)

257.

Landolt-Bernstein: Numerical Data and Functional Relafionships in Science and Technology, New Series, Group III, (0. Madelung, ed.) 22a:18, Springer-Verlag, Berlin (1987)

258.

Touloukian, Y. S., Kirby, R. K., Taylor, R. E., and Lee, T. Y. R., Thermal Expansion: Nonmetallic Solids, 13: 154, Thermophysical Properties of Matter, Plenum, New York (1977).

259.

Fork, D. K., Fenner, D. B., Connell, G. A. N., Phillips, Geballe, T. H., Appl. Phys. Lett., 57:1137-l 139 (1990)

260.

Fork, D. K., Fenner, D. B., Barton, R. W., Phillips, J. M., Connell, G. A. N., Boyce, J. B., and Geballe, T. H., Appl. Phys. Lett., 57:11611163 (1990)

261.

LandolbB6rnsfein: Numerical Data and Functional Relationships in Science and Technology, New Series, Group III, (0. Madelung, ed.) 22a:88, Springer-Verlag, Berlin (1987)

J. M., and

High

Tc Superconductors

611

262. Touloukian, Y. S., Kirby, R. K., Taylor, R. E., and Lee, T. Y. R., Thermal Expansion: Nonmetallic Solids, 131747, Thermophysical Properties ofMatter, Plenum, New York (1977) 263.

Nashimoto, K., Fork, D. K., and Geballe, T. H., Appl. Phys. Lett., 60:1199-1201 (1992); Fork, D. K., Nashimoto, K., and Geballe, T. H., Appl. Phys. Lett., 60:1621-1623 (1992)

264.

Matthews, J. W., EpitaxialGrowth, (J. W. Matthews, 559-609, Academic Press, New York (1975)

265.

Inam, A., Wu, X. D., Nazar, L., Hegde, M. S., Rogers, C. T., Venkatesan, T., Simon, R. W., Daly, K., Padamsee, H., Kirchgessner, J., Moffat, D., Rubin, D., Shu, Q. S., Kalokitis, D., Fathy, A., Pendrick, V., Brown, R., Brycki, B., Belohoubek, E., Drabeck, L., Gruner, G., Hammond, R., Gamble, F., Lairson, B. M., and Bravman, J. C., Appl. Phys. Lett., 56:1178-l 180 (1990)

ed.) Part B, pp.

266. Alarco, J. A., Brorsson, G., Ivanov, Z. G., Nilsson, P. -A., Olsson, E., and Lofgren, M., Appl. Phys. Lett., 61:723-725 (1992) 267.

Moeckly, 8. H., Russek, S. E., Lathrop, D. K., Buhrman, R. A., Li, J., and Mayer, J. W., Appl. Phys. Lett., 57:1687-1689 (1990)

268.

Berezin, A. B., Yuan, C. W., and de Lozanne, Lett., 57:90-92 (1990)

269.

Moeckly, B. H., Lathrop, D. K., Russek, S. E., Buhrman, R. A., Norton, M. G., and Carter, C. B., /EEE Trans. Magn., 27:1017-1020 (1991)

270.

Char, K., Newman, N., Garrison, S. M., Barton, R. W., Taber, R. C., Laderman, S. S., and Jacowitz, R. D., Appl. Phys. Lett., 57:409-411 (1990)

A. L., Appl. Phys.

271. Wang, F., Siegel, M., Smithey, R., Geerk, J., Linker, G., and Meyer, O., High Tc Superconductor Thin Films, (L. Correra, ed.), pp. 61 l615, North-Holland, Amsterdam (1992) 272. Wu, X. D., Dye, R. C., Muenchausen, R. E., Foltyn, S. R., Maley, M., Rollett, A. D., Garcia, A. R., and Nogar, N. S., Appl. Phys. Lett., 58:2165-2167 (1991) 273.

Prusseit, W., Corsepius, S., Baudenbacher, F., Hirata, K., Berberich, P., and Kinder, H., Appl. Phys. Lett., 61:1841-1843 (1992)

274.

Grigoryeva, 101 (1989)

I. V. and Vinnikov,

L. Ya., J. Low Temp. Phys., 74:81-

612

Molecular

Beam Epitaxy

275.

Pan, V. M., Kaminsky, G. G., Kasatkin, A. L., Kuznetsov, M. A., Prokhorov, V. G., Svetchnikov, V. L., Tretiatchenko, C. G., Flis, V. S., Yushchenko, S. K., Matsui, V. I., and Melnikov, V. S., Supercond. Sci. Technol., 5:S48-S54 (1992); Pan, V. M., Kasatkin, A. L., Svetchnikov, V. L., and Zandbergen, H. W., Cryogenics, 33:21-27 (1993); Pan, V.M., J. A//oys Corrtpd., 195387-394 (1993)

276.

Mannhart, J., Anselmetti, D., Bednorz, J. G., Catana, A., Gerber, Ch., Mtiller, K. A., and Schlom, D. G., Z. Phys,. B86:177-181 (1992)

277.

Gross, Ft., Gupta, A., Olsson, E., Segmtiller, Appl. Phys. Left., 57:203-205 (1990)

278.

Gao, J., Hauser, B., and Rogalla, H., J. Appl. Phys., 67:2512-2515 (1990)

279.

Nagaishi, T., Itozaki, H., Tanaka, S., Matsuura, T., Ota, N., Fujimori, N., and Yazu, S., Jpn. J. Appl. Phys., 3O:L718-L721 (1991)

280.

Pan, V. M., Solovjov, V. F., and Taborov, Technol., 5:s 192-S 195 (1992)

281.

Sasaura, M., Mukaida, 57:2728-2729 (1990)

282.

Sasaura, (1992)

283.

Ralston, R. W., Kastner, M. IEEE Spectrum, 29:50-55 Valentino, A. J., and Brandle, (1991); Berkstresser, G. W., 0yst. Growth, 109:467-471

284.

Hsu, Y. T., Applied Technology communication.

285.

Garrison, S. M., Newman, N., Cole, B. F., Char, K., and Barton, R. W., Appl. Phys. Lett., 58:2168-2170 (1991); Garrison, S. M., Newman, N., Cole, B. F., Char, K., and Barton, R. W., Appl. Phys. Lett., 59:3060 (1991)

286.

Fork, D. K., Garrison, S. M., Hawley, Mater. Res., 7:1641-1651 (1992)

287.

Ramesh, R., Hwang, D., Ravi, T. ., Inam, A., Barner, J. B., Nazar, L., Chan, S. W., Chen, C. Y., Dutta, B., Venkatesan, T., and Wu, X. D., Appl. Phys. Leti., 56:2243-2245 (1990); Ravi, T. S., Hwang, D. M., Ramesh, R., Chan, S. W., Nazar, L., Chen, C. Y., Inam, A., and Venkatesan, T., Phys. Rev., B42:10141-10151 (1990)

M., and Miyazawa,

M. and Miyazawa,

S.,

J. Cryst.

A., and Koren,

V. F., Supercond. S., Appl. Growth,

G.,

Sci.

Phys. Left, 123:126-132

A., Gallagher, W. J., and Batlogg, B., (August 1992); Berkstresser, G. W., C. D., J, Cryst. Growth, 109:457-466 Valentino, A. J., and Brandle, C. D., J. (1991) Enterprises,

Columbia,

M., and Geballe,

SC, private

T. H., J.


288.

613

Laderman, S. S., Taber, Ft. C., Jacowitz, R. D., Mall, J. L., Eom, C. B., Hylton, T. L., Marshall, A. F., Geballe, T. H., and Beasley, M. R., Phys. Rev., B43:2922-2933 (1991)

289. Char, K., Colclough, M. S., Lee, L. P., and Zaharchuk, Phys. Left., 59:2177-2179 (1991)

G., Appl.

290. Wu, X. D., Luo, L., Muenchausen, R. E., Springer, K. N., and Foltyn, S., Appl. Phys. Lett., 60:1381-l 383 (1992) 291.

Hellman, E. S., Hartford, E. H., and Gyorgy, E. M., Appl. Phys. Leti., 58:1335-l 337 (1991)

292.

Norton, M. G., Hellman, E. S., Hartford, E. H., Jr., and Carter, C. B., J. Cryst. Growth, 113:716-721 (1991); Norton, M. G., Hellman, E. S., Hartford, E. H., Jr., and Caner, C. B., Physica, C205:347-353 (1993)

293.

Char, K., Newman, N., Garrison, S. M., Barton, R. W., andzaharchuk, G., “Surface Resistance of Epitaxial YBa,Cu,O, Thin Films on Sapphire,” presented at the Applied Superconductivity Conference, Aspen (1990), unpublished.

294.

Mogro-Campero, A., Hunt, B. D., Turner, L. G., Burrell, Balz, W. E., Appl. Phys. Lett., 52:584-586 (1988)

295.

Madakson, P., Cuomo, J. J., Yee, D. S., Roy, R. A., and Scilla, G., J. Appl. Phys., 63:2046-2053 (1988)

M. C., and

296. Venkatesan, T., Chase, E. W., Wu, X. D., Inam, A., Chang, C. C., and Shokoohi, F. K., Appl. Phys. Lett., 53:243-245 (1988) 297.

Fenner, D. B., Viano, A. M., Fork, D. K., Connell, G. A. N., Boyce, J. B., Ponce, F. A., and Tramontana, J. C., J. Appl. Phys., 69:21762182 (1991)

298.

Nordman, C. A., Wang, T., Chandrasekhar, N., Beauchamp, K. M., Achutharaman, V. S., Schulze, R. K., Spalding, G. C., Lin, Z. -H., Evans, J. F., and Goldman, A. M., J. Appl. Phys., 70:5697-5699 (1991)

299.

Filby, J. D. and Nielsen, S., J. Appl. Phys., 18:1357-1382

300.

Manasevit,

301.

Manasevit, H. M., Miller, A., Morritz, F. L., and Nolder, R., Trans. Metal/. Sot., A/ME 233:540-549 (1965)

302.

Manasevit, (1966)

H. M., J. Crystal Growth, 22:125-148

H. M. and Forbes,

D. H., J. Appl.

(1967)

(1974)

Whys., 37:734-739

303. Golecki, I., Manasevit, H. M., Moudy, L. A., Yang, J. J., and Mee, J. E., Appl. Phys. Lett., 42:501-503 (1983); Manasevit, H. M., Golecki, I., Moudy, L. A., Yang, J. J., and Mee, J. E., J. Nectrochem. Sot., 130: 1752-l 758 (1983)

614

Molecular

Beam Epitaxy

304.

Ishida, M., Katakabe, I., Nakamura, T., and Ohtake, N., Appl. Phys. Left., 52:1326-1328 (1988); Sawada, K., Ishida, M., Nakamura, T., and Suzaki, T., J. Crysf. Growth, 95494-499 (1989)

305.

lizuka, H., Yokoo, K., and Ono, S., Appl, Phys. Left., 61:2978-2980 (1992)

306.

Ihara, M., Arimoto, Y., Jifuku, M., Kimura, T., Kodama, S., Yamawaki, H., and Yamaoka, T., J. Elecfrochem. Sot., 129:2569-2573 (1982); Ihara, M., Microelectron. Eng., 1 :161-l 77 (1983)

307.

Morita, M., Fukumoto, H., Imura, T., Osaka, Y., and Ichihara, Appl. Pbys., 58:2407-2409 (1985)

308.

Legagneux, P., Garry, G., Dieumegard, Gautherin, G., and Siejka, J., Appl. (1988)

309.

Fukumoto, H., Imura, T., and Osaka, 27:L1404-Ll405 (1988)

310.

Miura, S., Yoshitake, T., Matsubara, S., Miyasaka, Y., Shohata, and Satoh, T., Appl. Phys. Left., 53:1967-l 969 (1988); Wu, X. Inam, A., Hegde, M. S., Wilkens, B., Chang, C. C., Hwang, D. Nazar, L., Venkatesan, T., Miura, S., Matsubara, S., Miyasaka, and Shohata, N., Appl. Phys. Left., 54:754-756 (1989)

311.

Inoue, T., Yamamoto, Y., Koyama, S., Suzuki, S., and Ueda, Y., Appl, Phys. Left., 56:1332-1333 (1990); Yamamoto, Y., Satoh, M., Sakurai, Y., Nakajima, S., Inoue, T., and Ohsuna, T., Jpn. J. Appl. Phys., 32:L620-L623 (1993)

312.

Yoshimoto, M., Nagata, H., Tsukahara, T., and Koinuma, H., Jpn. J. Appl. Phys., 29:Lll99-Ll202 (1990); Nagata, H., Tsukahara, T., Gonda, S., Yoshimoto, M., and Koinuma, H., Jpn. J. Appl. Phys., 3O:L1136-Ll138 (1991); Nagata, H., Yoshimoto, M., Koinuma, H., Min, E., and Haga, N., J. Crysf. Growfh, 123:1-4 (1992)

313.

Luo, L., Wu, X. D., Dye, R. C., Muenchausen, R. E., Foltyn, S. R., Coulter, Y., Maggiore, C. J., and Inoue, T., Appl. Phys. Left, 59:2043-2045 (1991)

314.

Fork, D. K., Ponce, F. A., Tramontana, J. C., Newman, N., Phillips, J. M., and Geballe, T. H., Appl. Phys. Left., 58:2432-2434 (1991)

315.

Burns, M. J., de la Houssaye, P. R., Russell, S. D., Garcia, G. A., Clayton, S. R., Ruby, W. S., and Lee, L. P., Appl. Phys. Leff., 63:1282-l 284 (1993)

M., J.

D., Schwebel, C., Pellet, C., Phys. Leff., 53:1506-l 508 Y.,

Jpn.

J. Appl.

Phys., N., D., M., Y.,

316. Tseng, M. Z, Nguyen, C., Tarsa, E., Chang, L. D., Hu, E. L., and Kroemer, H., Appl. Phys. Left., 61:601-603 (1992)


317.

615

Prusseit, W., Corsepius, S., Utz, B., Baudenbacher, F., Hirata, K., Berberich, P., Kinder, H., and Eibl, O., /EEE Trans. Appl. Supercond., 3:1070-l 073 (1993)

318. Goodyear, S. W., Satchell, J. S., Humphreys, R. G., Chew, N. G., and Edwards, J. A., Physica, Cl 92:85-94 (1992) 319. Chandrasekhar, N., Agrawal, V., Achutharaman, A. M., Appl. Phys. Lett., 60:2424-2426 (1992)

V. S., and Goldman,

320. Agrawal, V., Chandrasekhar, N., Zhang, Y.J., Achutharaman, V. S., Mecartney, M. L., and Goldman, A. M., J. Vat. Sci. Technol., A10:1531-1536 (1992) 321. Tanaka, S., Nakamura, T., Tokuda, H., and liyama, M., Appl. Phys. Lett., 62:3040-3042 (1993) 322. Schlom, D. G., Anselmetti, D., Bednorz, J. G., Broom, R., Catana, A., Frey, T., Gerber, Ch., Giintherodt, H. -J., Lang, H. P., and Mannhart, J., Z. Phys., B86:163-175 (1992) 323.

Lang, H. P., Frey, T., and Gijntherodt, 15667-670 (1991)

H. -J.,

Europhys.

324.

Baudenbacher, F., Hirata, K., Berberich, P., Kinder, W., and Lang, H. P., Physica, Cl85-189:2177-2178

325.

Luo, L., Hawley, M. E., Maggiore, C. J., Dye, R. C., Muenchausen, R.E., Chen, L., Schmidt, B., and Kaloyeros, A. E., Appl. Phys. Left., 62:485-486 (1993)

H., Assmann, (1991)

326. Schlom, D. G., Anselmetti, D., Bednorz, J. G., Gerber, Mannhart, J., J. Cryst. Growth, 137:259-267 (1994) 327.

Lett.,

Ch., and

Frank, F. C., Disc. Farad. Sot., 5:48-79 (1949); Burton, W. K., Cabrera, N., and Frank, F. C., Philos. Trans. R. Sot. London, A243:299-358 (1951); Frank, F. C., Adv. Phys., 1:91-109 (1952)

328. Verma, A. R., Crystal Growth and Dislocations, (1953) 329. Baronnet,

A., J. Cryst. Growth, 19:193-l

Butterworths,

London

98 (1973)

330. Lowndes, D. H., Zheng, X. -Y., Zhu, S., Budai, J. D., and Warmack, R. J., Appl. Phys. Left., 61:852-854 (1992) 331. Satchell, J. S., Humphreys, R. G., Chew, N. G., Goodyear, S. W., Edwards, J. A., and Keene, M. N., Physica, Cl80:247-250 (1991) 332. Chew, N. G., Goodyear, S. W., Humphreys, R. G., Satchell, J. S., Edwards, J. A., and Keene, M. N., Appl. Phys. Letf., 60:1516-1518 (1992) 333. Edwards, J. A., Satchell, J. S., Chew, N. G., Humphreys, R. G., Keene, M. N., and Dosser, 0. D., Appl. Phys. Lett., 60:2433-2435 (1992)

616

Molecular

Beam Epitaxy

334.

Kawai, M., private communication.

335.

ham, A., Rogers, C. T., Ramesh, R., Remschnig, K., Farrow, L., Hart, D., Venkatesan, T., and Wilkens, B.,App/. Phys. Lett., 57:24842486 (1990)

336.

Linker, G., Xi, X. X., Meyer, O., Li, Q., and Geerk, J., Solid State Commun., 69:249-253 (1989)

337.

Eom, C. B., Marshall, A. F., Laderman, S. S., Jacowitz, Geballe, T. H., Science, 249:1549-l 552 (1990)

338. Terashima, T., Bando, Y., lijima, K., Yamamoto, Appl. Phys. Lett., 53:2232-2234 (1988)

R. D., and

K., and Hirata, K.,

339. Schlom, D. G., Marshall, A. F., Harris, J. S., Jr., Bozovic, I., and Eckstein, J. N., Advances in Superconducfivify 111:Proceedings of the 3rd lnternafional Symposium on Superconductivity (ES ‘90), (K. Kajimura and H. Hayakawa, eds.) pp. 101 l-l 016, Springer-Verlag, Tokyo, (1991) 340.

Eckstein, J. N., Bozovic, I., Schlom, D. G., and Harris, J. S., Jr., Appl. Phys. Lett., 57:1049-l 051 (1990)

341.

Eckstein, J. N., Bozovic, I., Klausmeier-Brown, M., Virshup, Rails, K. S., Thin So/id Films, 216:8-l 3 (1992)

342.

Bozovic, I., Eckstein, J. N., Klausmeier-Brown, G., J, Supercond., 5:19-23 (1992)

G., and

M. E., and Virshup,

343. Virshup, G. F., Klausmeier-Brown, M. E., Bozovic, J. N., Appl. Phys. Lett., 60:2288-2290 (1992)

I., and Eckstein,

344.

Klausmeier-Brown, M. E., Virshup, G. F., Bozovic, J. N., Appl. Phys. Lett., 60:2806-2808 (1992)

I., and Eckstein,

345.

Bodin, P., Sakai, S., and Kasai, Y., Jpn. J. Appl. Phys., 3l:L949L952 (1992)

346.

Suzuki, K., Karaki, T., Iwashima, K., Shibata, M., Okada, H., Onnagawa, H., and Miyashita, K., Jpn. J. Appl. fhys., 31 :Ll339L1341 (1992)

347. Tsukada, I. and Uchinokura, Ll470 (1991)

K., Jpn. J. Appl.

Phys., 3O:Ll468-

348.

Kasai, Y. and Sakai, S., J. Cryst. Growth, 115:758-761

349.

Bozovic, I., Eckstein, J. N., Schlom, D. G., and Harris, J. S., Science and Technology of Thin film Superconductors 2, (R. D. McConnell, and R. Noufi, eds.) pp. 267-272, Plenum Press, New York (1990)

(1991)

350.

Ikeda, Y., Ito, H., Shimomura, S., Oue, Y., Inaba, K., Hiroi, Z., and Takano, M., Physica, Cl 59:93-l 04 (1989)


617

351.

Roth, R. S., Rawn, C. J., Burton, B. P., and Beech, F., J. Res. /Vat/. Inst. Stand. Technol., 95291-335 (1990)

352.

Ishibashi, T., Okada, Y., Yokoyama, Appl. Phys., 30:3900-3903 (1991)

S., and Kawabe,

M., Jpn. J.

353. Sakai, S., Kasai, Y., and Bodin, P., Jpn. J. Appl. Phys., 31:L399L401 (1992) 354. Chou, H., Chen, H. S., Kortan, A. R., Kimerling, and Wu, M. K., Appl. Phys. Lett., 58:2836-2838

L. C., Thiel, F. A., (1991)

355.

Betz, J., Piehler, A., Pechen, E. V., and Renk, K. F., J. Appl. Phys., 71:2478-2479 (1992)

356.

Face, D. W. and Nestlerode, J. P., Appl. Phys. Lett. 61:1838-1840 (1992); Face, D. W. and Nestlerode, J. P., /EEE Trans. Appl. Supercond., 3:1516-1519 (1993)

357. Ahn, B. T., Lee, W. Y., and Beyers, R., Appl. Phys. Lett., 60:21502152 (1992) 358. Abe, H., Toda, F., and Ogiwara, 14:100-l 02 (1993)

M., IEEE Nectron

Device

Lett.,

359. Cava, R. J., Siegrist, T., Peck, W. F., Jr., Krajewski, J. J., Batlogg, B., and Rosamilia, J., Phys. Rev., B44:9746-9748 (1991) 360. Cava, R. J., Takagi, H., Zandbergen, H. W., Hessen, B., Krajewski, J. J., and Peck, W. F., Jr., Phys. Rev., B46:14101-14104 (1992) 361. Xu, Q., Fu, W. T., van Ruitenbeek, Physica, Cl 67:271-277 (1990)

J. M., and de Jongh,

L. J.,

362. Warren, B. E., X-ray Diffraction, Dover, New York (1969, 1990). See in particular Eqs. 1.12, 2.5, 3.2, 3.24, and 4.7. 363.

Hendricks,

S. and Teller, E., J. Chem. Phys., 10:147-167

(1942)

364.

Kuroda, K., Kojima, K., Tanioku, M., Yokoyama, K., Jpn. J. Appl. Phys., 28:1586-1592 (1989)

365.

Kojima, K., Schlom, D.G., Kuroda, K., Tanioku, M., Hamanaka, K., Eckstein, J. N., and Harris, J. S., Jr., Jpn. J. Appl. Phys., 29:L1638L1641 (1990)

366.

Fujita, J., Tatsumi, T., Yoshitake, T., and Igarashi, H., Science and Technology of Thin Film Superconductors, (McConnell, R. D. and Wolf, S. A., eds.) pp. 175-184, Plenum, New York (1989); Fujita, J., Tatsumi, T., Yoshitake, T., and Igarashi, H., Appl. Phys. Lett,, 54:2364-2366 (1989)

K., and Hamanaka,

367. Krebs, H. -U. and Kehlenbeck, M., Physica, C162-164:119-120 (1989); Krebs, H. -U., Kehlenbeck, M., Steins, M., and Kupcik, V., J. Appl. Phys., 69:2405-2409 (1991)

618

Molecular

Beam Epitaxy

368. Zhu, S., Lowndes, D. H., Chakoumakos, B. C., Budai, J. D., Christen, D. K., Zheng, X. -Y., Jones, E., and Warmack, B., Appl, Phys. Leti., 63:409-411 (1993) 369.

Lang, Ch., Hettich, B., Schwarz, Physica, Cl 82:79-88 (1991)

M., Bestgen,

H., and Elschner,

370.

Matsushima, T., Ichikawa, Y., Adachi, H., Setsune, K., Solid State Commun., 76:1201-l 204 (1990)

371.

Horiuchi, K., Kawai, T., Kanai, Phys., 3O:L1381-L1383 (1991)

372.

moue, M., Yoshida, K., Ohbayashi, Physica, C200:403-408 (1992)

373.

Pargellis, A. N., Sharifi, F., Dynes, Ft. C., Miller, B., Hellman, E. S., Rosamilia, J. M., and Hartford, E. H., Jr., Appl. Phys. Lett., 58:95-96 (1991)

374.

Hellman, E. S., Martin, S., Hartford, E. H., Jr., Werder, D. J., Roesler, G .M., Jr., and Tedrow, P. M., Physica, C201:166-170 (1992)

375.

Fink, R. L., Hilbert, C., and Kroger, H., Appl. Phys. Lett., 62:33603362 (1993)

376.

Kussmaul, A., Hellman, E. S., Hartford, E. H., Jr., and Tedrow, M., Appl. Phys. Lett., 63:2824-2826 (1993)

M., and Kawai,

S.,

S.,

K., and Wasa, Jpn. J. Appl.

K., Takai, Y., and Hayakawa,

H.,

377. Adachi, H., Ichikawa, Y., Hirochi, K., Matsushima, T., Setsune, and Wasa, K., Jpn. J. Appl. Phys., 29:L81-L82 (1990)

P. K.,

378.

The first report of GaAs grown directly on vicinal Si (100) was by: Metze, G. M., Choi, H. K., and Tsaur, B-Y., Appl. Phys. Lett., 45:1107-l 109 (1984). Although the vicinal nature of the Si (100) substrates used in a prior study is not stated, from the data presented it is likely that this earlier study of GaAs growth directly on silicon also used vicinal Si (100) surfaces: Wang, W. I., Appl. Phys. Left,, 44:1149-l 151 (1984)

379.

Shibahara, K., Nishino, 50:1888-l 890 (1987)

380.

Budai, J. D., Chisholm, M. F., Feenstra, R., Lowndes, D. H., Norton, D. P., Boatner, L. A., and Christen, D. K., Appl, Phys. Left., 58:21742176 (1991)

381.

Fujita, J., Yoshitake, T., Satoh, T., Miura, S., Tsuge, Igarashi, H., Appl. Phys. Lett., 59:2445-2447 (1991)

H., and

382.

Satoh, T., Fujita, J., Yoshitake, Cl 91:359-362 (1992)

Physica,

S., and Matsunami,

H., Appl. fhys.

T., and Igarashi,

H.,

Leti.,

High

Tc Superconductors

619

383. Kuroda, K., Wada, O., Tanimura, J., Kojima, K., Takami, T., Kataoka, M., Ogama, T., and Hamanaka, K., Jpn. J. Appl. Phys., 3O:L582L584 (1991) 384. Wada, O., Kuroda, K., Tanimura, J., Kataoka, M., Kojima, K., Takami, T., Hamanaka, K., and Ogama, T., Jpn. J. Appl. Phys., 3O:L1881-L1883 (1991) 385. Tanimura, J., Kuroda, K., Kataoka, M., Wada, O., Takami, T., Kojima, K., and Ogama, T., Jpn. J. Appl. Phys., 32:L254-L256 (1993) 386. Ishibashi, T., Song, K., Okada, Y., and Kawabe, Phys., 31: L406-L409 (1992)

M., Jpn. J. Appl.

387.

Ishibashi, T., Fujita, T., Song, K., Okada, Y., and Kawabe, J. Appl. Phys., 32:L257-L259 (1993)

388.

Kwo, J., Fleming, R. M., Kao, H. L., Werder, D. J., and Chen, C. H., Appl. Phys. Lett., 60:1905-l 907 (1992)

389. Tsukada, (1994)

I. and Uchinokura,

K., Physica,

M., Jpn.

B, 194-196:2329-2330

390. Sunshine, S. A., Siegrist, T., Schneemeyer, L. F., Murphy, D. W., Cava, R. J., Batlogg, B., van Dover, R. B., Fleming, R. M., Glarum, S. H., Nakahara, S., Farrow, R., Krajewski, J. J., Zahurak, S. M., Waszczak, J. V., Marshall, J. H., Marsh, P., Rupp, L. W., Jr., and Peck, W. F., Phys. Rev., B38:893-896 (1988) 391.

Bordet, P., Capponi, J. J., Chaillout, C., Chenavas, J., Hewat, A. W., Hewat, E. A., Hodeau, J. L., Marezio, M., Tholence, J. L., and Tranqui, D., Physica, Cl 56:189-l 92 (1988)

392.

Eibl, O., Physica, C168:215-238

(1990)

393. Aarnink, W. A. M., Reuvekamp, E. M. C. M., Verhoeven, M. A. J., Pedyash, M. V., Gerritsma, G. J., van Silfhout, A., Rogalla, H., and Ryan, T. W., Appl. Phys. Lett., 61:607-609 (1992) 394.

Fujita, J., Yoshitake, T., Satoh, T., Ichihashi, IEEE Trans. Magn., 27:1205-1210 (1991)

T., and Igarashi,

395. Scherer, T., Marienhoff, P., Herwig, R., Neuhaus, Physica, Cl 97:79-83 (1992)

H.,

M., and Jutzi, W.,

396. Young, K. H. and Sun, J. Z., Appl. Phys. Left., 59:2448-2450

(1991)

397. Shlichta, P. J., NASA Tech. Briefs, 89-0548 (1989); Shlichta, P.J., NASA Technical Support Package NPO- 17/684 1, NASA, BWI Airport (1989) 398. Takei, W. J., Wu, S. Y., and Francombe, 28:188-l 98 (1975)

M. H., J, Cryst. Growth,

620

399.

Molecular

Beam Epitaxy

Fujita, J., Satoh, T., Yoshitake, T., Tsai, J. S., and Tsuge, H., Extended Abstracts of the 1992 International Conference on Solid State Devices and Materials, pp. 718-719, Business Center for Academic Societies Japan, Tokyo (1992)

400. Terashima, T., Shimura, K., Bando, Y., lijima, K., Yamamoto, K., Hirata, K., Hayashi, K., Kamigaki, K., and Terauchi, H., Advances in Superconductivity III: Proceedings of the 3rd International Symposium on Superconductivity (ES ‘901, (K. Kajimura and H. Hayakawa, eds.) pp. 851-856, Springer-Verlag, Tokyo (1991); Kamigaki, K., Terashima, T., Shimura, K., Bando, Y., and Terauchi, H., Physica, Cl 83:252-256 (1991) 401.

Pennycook, S. J., Chisholm, M. F., Jesson, D. E., Norton, D. P., Lowndes, D. H., Feenstra, R., Kerchner, H. R., and Thomson, J. O., Phys. Rev. Lett., 67:765-768 (1991)

402.

Eibl, O., Hoenig, H. E., Triscone, J. -M., Fischer, 0., Antognazza, and Brunner, O., Physica, Cl 72:365-372 (1990)

403. Jia, C. L., Soltner, H., Jakob, G., Hahn, T., Adrian, K., Physica, C210: l-1 5 (1993) 404.

Fujita, J., Yoshitake, T., Igarashi, Lett., 56:295-297 (1990)

H., and Satoh,

L.,

H., and Urban, T., Appt. Phys.

405.

Karl, H. and Stritzker,

406.

Gupta, A. and Hussey, B. W., Appl. Phys. Lett., 58:1211-1213 (1991); Chern, M. Y., Gupta, A., and Hussey, B. W., Appl. Phys. Lett., 60:3045-3047 (1992); Chern, M. Y., Gupta, A., Hussey, B. W., and Shaw, T. M., J. Vat. Sci. Tecchnol. A, 11:637-641 (1993)

B., Phys. Rev. Lett,, 69:2939-2942

(1992)

407. Williams, S., Zheng, J. Q., Shih, M. C., Wang, X. K., Lee, S. J., Rippert, E. D., Maglic, S., Kajiyama, H., Segel, D., Dutta, P., Chang, Ft. P. H., Ketterson, J. B., Roberts, T., Lin, Y., Kampwirth, R. T., and Gray, K., J. Appl. Phys., 72:4798-4804 (1992) 408.

Fuoss, P. H., Kisker, D. W., Lamelas, F. J., Stephenson, G. B., Imperatori, P., and Brennan, S., Phys. Rev. Lett., 69:2791-2794 (1992)

409. Adachi, H., Kohiki, S., Setsune, K., Mitsuyu, J. Appl. Phys., 27:L1883-1886 (1988) 410.

T., and Wasa, K., Jpn.

Koinuma, H., Nagata, H., Takano, A., Kawasaki, M., and Yoshimoto, M., Science and Technology of Thin Film Superconductors, (R. D. McConnell, and S. A. Wolf, eds.) pp. 205-213, Plenum, New York (1989)


621

411.

Nakamura, K., Sato, J., Kaise, M., and Ogawa, K., Jpn. J. Appl. Phys., 28:L437-L440 (1989); Nakamura, K., Sato, J., and Ogawa, K., Jpn. J. Appl. Phys., 29:L77-L80 (1990)

412.

Kittl, J. A., Nieh, C. W., Lee. D. S., and Johnson, W. L., Mater. Lett. 9:336-338 (1990); Kittl, J. A., Johnson, W. L., and Nieh, C. W., J. Mater. Res., 7:2003-2016 (1992)

413.

Ohbayashi, K., Anma, M., Takai, Y., and Hayakawa, H., Jpn. J. Appl. Phys. 29:L2049-L2052 (1990); Ohbayashi, K., Yoshida, K., Anma, M., Takai, Y., and Hayakawa, H., Jpn. J. Appl. Phys., 31 :L953-L955 (1992)

414. Yang, K. -Y., Dilorio, M. S., Yoshizumi, S., Maung, M. A., Zhang, J., Tsai, P. K., and Maple, M. B., Appl. Phys. Left,, 61:2826-2828 (1992). 415.

Kanai, M., Kawai, T., Kawai, S., and Tabata, 54:1802-l 804 (1989)

H., Appl. Phys. Letf.,

416. Duray, S. J., Buchholz, D. B., Song, S. N., Richeson, D. S., Ketterson, J. B., Marks, T. J., and Chang, R. P. H., Appl. Phys. Lett,, 59:15031505 (1991) 417.

Fujii, K., Zama, H., and Oda, S., Jpn. J. Appl. fhys., (1992)

31 :L787-L789

418. Cheung, J. T. and Cheung, D. T., J. Vat. Sci. Technol., 21:182-186 (1982); Cheung, J. T. and Sankur, H., Critical ReviewsTM in Solid State and Materials Sciences, (J. E. Greene, ed.) 15:63-l 09, CRC Press, Boca Raton (1988) 419. Kanai, M., Horiuchi, K., Kawai, T., and Kawai, S., Appl. Phys. Left., 57:2716-2718 (1990); Kawai, T., Egami, Y., Tabata, H., and Kawai, S., Nature, 349:200 (1991); Tabata, H. and Kawai, T., Thin So/id Films 225:275-279 (1993) 420. Koinuma, H., Nagata, H., Tsukahara, M., Appl. Phys. Left., 58:2027-2029

T., Gonda, S., and Yoshimoto, (1991)

421. Leskela, M., Molsa, H., and Niinistd, 6:627-656 (1993)

L., Supercond.

422. Atomic Layer Epitaxy, Glasgow (1990) 423.

Niinisto,

(T. Suntola

and M. Simpson,

Sci. Techmol., eds.) Blackie,

L. and Leskela, M., Thin So/id Films, 225:130--l 35 (1993)

424. Zama, H., Sakai, K., and Oda, S., Jpn. J. Appl. Phys., 31 :L1243L1245 (1992); Oda, S., Zama, H., Fujii, K., Sakai, K., and Chen, Y. C., Thin So/id Films, 225284-287 (1993)

622

425.

Molecular

Beam Epitaxy

Ino, S., Ichikawa, 1457 (1980)

T., and Okada, S., Jpn. J. Appl. Phys., 19:1451-

426. Atwater, H. A. and Ahn, C. C., Appl. fhys. Letf. 58:269-271 (1991); Nikzad, S., Ahn, C. C., and Atwater, H. A., J. Vat. Sci, Technol., B10:762-765 (1992) 427.

Kamei, M., Aoki, Y., Usui, T., and Morishita, 31 :1326-l 328 (1992)

T., Jpn. J. Appl. Phys.,

428.

Kim, Y. -T., Collins, R. W., and Vedam, K., Surf. Sci., 233:341350 (1990); Collins, R. W. and Kim, Y. -T., Anal. Chef-n., 62:887A-900A (1990); An, I., Nguyen, H. V., Nguyen, N. V., and Collins, R. W.,Phys. Rev. Lett., 65:2274-2277 (1990); Heyd, A. R., An, I., Collins, R. W., Cong, Y., Vedam, K., Bose, S. S., and Miller, D. L., J. Vat. Sci, Techno/., A 9:81 O-81 5 (1991)

429. Aspnes, D. E., Bhat, R., Caneau, C., Colas, E., Florez, L. T., Gregory, S., Harbison, J. P., Kamiya, I., Keramidas, V. G., Koza, M. A., J. Cryst. Growth, 120:71-77 (1992) 430. Vedam, K., McMarr, P. J., and Narayan, 47:339341 (1985); McMarr, P. J., Vedam, Appl. f’hys., 59:694-701 (1986)

J., Appl. Phys. K., and Narayan,

Lett., J., J.

431. Aspnes, D. E., Quinn, W. E., Tamargo, M. C., Pudensi, M. A. A., Schwarz, S. A., Brasil, M. J. S. P., Nahory, R. E., and Gregory, S., Appl. Phys. Lett., 60:1244-l 246 (1992) 432.

Barth, K. -L. and Keilmann, (1993)

F.,

Rev. Sci. Instrum.,

64:870-875

433.

Kobayashi, N. and Horikoshi, Y., Jpn. J. Appl. Phys., 11 :L1880L1882 (1989); Horikoshi, Y., Kawashima, M., and Kobayashi, N., J. Cryst. Growth, 111:200-204 (1991)

MBE Growth of Artificially-Layered Magnetic Metal Structures Robin E C. Farrow, Ronald E Marks, Gerald R. Harp, Dieter Weller, Thomas A. Rabedeau, Michael E Toney, Stuart S. P. Parkint

1 .O

INTRODUCTION

The artificially-layered structures considered here comprise one, several, or many magnetic films sandwiched by non-magnetic spacer films. In recent years, such structures have exhibited a variety of novel phenomena which are interesting from the standpoint of the physics of magnetism as well as for potential device applications. Collections of papers in this field, including reviews of specific topics, can be found in the proceedings of the Spring 1991 and 1993 Symposia of the Materials Research Societyon magnetic surfaces, thin films, and multilayers,f’) and in the Proceedings of the 1992 NATO Advanced Research Workshop on structure properties

and magnetism

in systems of reduced dimension.t*)

of artificially-layered

magnetic

structures

Interesting

include perpendicular

magnetic anisotropy in films of Fe or Co that are a few monolayers thick, sandwiched by non-magnetic metals or simply left uncoated in ultra-high vacuum.

In some cases, as for Co/Pt and Co/Pd multilayers,

this perpen-

dicular anisotropy persists to high temperatures (~200°C) and has technological applicationst3) in magneto-optical (MO) storage of information. Other interesting properties exhibited by these layered structures include the giant magnetoresistance (GMR) effect,t4)-t6] and oscillatory-AF + This work was supported in part by the Office of Naval Research 0339 and NO001 4-92-C-0084).

623

(Contracts

NO001 4-87-C-

624

Molecular

Beam Epitaxy

(antiferromagnetic) to (FM) ferromagnetic-exchange magnetic (NM) spacer films in FM-NM-FM A detailed and growth

understanding

conditions

effects are present.

coupling across non-

sandwiches

and multilayers.m-ts)

of these phenomena

is not yet available

often influence

and even determine

MBE growth, as opposed to sputtering

whether

the

or conventional

evaporation, provides the opportunity for growth under controlled and monitored conditions with in situ analysis of film structure and composition. Figure 1 shows a schematic diagram of the MBE system used in our laboratory. It incorporates a variety of probes which enable the growth process to be studied in detail. For example, film growth can be monitored in real time using RHEED (reflection high energy electron diffraction). This is particularly valuable in determining optimum growth conditions for specific epitaxial structures. In addition, temporary interruption of growth followed by transfer of the substrate, in ultra-high vacuum, to an analysis chamber allows interface formation to be studied by angle-resolved XPS (x-ray photoelectron spectroscopy) and XPD (x-ray photoelectron diffraction) as well as by angle-resolved AES (Auger electron spectroscopy) and in some cases by SPPD (spin-polarized photoelectron diffraction) and SEM (scanning electron microscopy). A key advantage of MBE over more conventional film growth techniques

is that it enables growth of the layered structure

along specific

crystalline directions. This enables one to probe the dependence of magnetic phenomena on orientation, often providing insight into the mechanisms underlying the magnetic phenomena of interest and providing a test of theoretical predictions. Lattice-matching between the prelayer (seed film) and substrate can be achieved by appropriate choice of materials and the growth axis of the magnetic structure selected. The choice of substrate is largely determined by what type of magnetic property is to be studied.

For example, where magnetoresistance

it is necessary

to grow the structure

(MR) is to be measured,

onto semiconducting

or insulating

wafers as opposed to metal single crystals to prevent the shunting current through the substrate. On the other hand, magnetic anisotropy

of or

exchange coupling can be studied for structures grown onto any substrate. Where a highly perfect, exact-orientation (001) substrate is required, and where the magnetic and structural properties can be studied using localized probes, some groups use a single-crystal metal whisker with (001) growth facets, Since we are interested in magnetoresistance and in preparing samples of large area, suitable for x-ray diffraction, as well as in preparing multiple samples simultaneously, our preferred substrates are

Artificially-Layered

semiconductor

or insulating

films with semiconductor applications.[lO1[lll

wafers. substrates

MBE growth chamber

Magnetic

Metal Structures

Indeed, the integration is of interest

625

of magnetic

for potential

device

Reparation chamkr

c .3 effusion sources 3 elccuon gun Y)“rceS RHEED . TcmpraturesonLrolled substrate.manipulator: -Range 4-X03 K -SIZC 0.5-3” diamtcr -Rotadon 1 Hz Growth pressurr:

f

.

.unv sample nansfer track Sample loadlock .LEED . Transferplusure: 2OO”C) temperatures is needed, for example in studying the effects of interdiffusion and chemical ordering on magnetic anisotropy, we have developed new epitaxial systems: Pt/basal-plane sapphiretl*) and Pt/SrTi0,t13) which are chemically stable at higher temperatures than structures based on GaAs substrates.

626

Molecular

Beam Epitaxy

2.0

SEEDED

2.1

Semiconductor

EPITAXY

OF MAGNETIC

METALS

Substrates

MBE growth of magnetic

metal structures

on semiconductors

usu-

ally requires an intermediate film between the metal and the semiconductor. This is because 3d transition metal elements and rare earths often react with the semiconductor forming interfacial compounds which modify or prevent epitaxy. Such reactions may also generate mobile impurity species which can modify the magnetic properties of the structure.t14) Where the epitaxial growth of magnetic metal films on GaAs substrates is required, two contrasting techniques have been developed, one of which is applicable to the growth of specific intermetallic magnetic compounds, the other to the growth of a wide variety of artificially-layered magnetic metal structures containing elemental magnetic metals. The group of Sands, Harbison, and co-workers at Bellcore have pioneeredn5) the use of intermetallic template films, which are thermodynamically stable in contact with the substrate and which form the template for molecular beam epitaxy of magnetic compounds and alloys where the growth is at elevated temperatures (>l OO’C). In a few cases, the magnetic compound of interest is thermodynamically stable in contact with GaAs, has a near lattice match to GaAs, and can be grown directly at elevated temperature. This is the case for the ferromagnetic b-phase of Mn,,Ga, with x = 0.6. Krishnant16) has recently reported the magnetic properties of this phase grown on GaAs(001). However, a more general strategy is necessary for growth of artificially-layered structures, comprising elemental metals, along specific crystalline growth axes. This strategy requires the selection of a suitable seed film which does not react with GaAs or the subsequent overlayer and which leads to the desired growth axis. In the case of rare earth metal epitaxy on GaAs, the author and co-workers have foundt17) that a suitable prelayer is a rare earth fluoride, LaF, or NdFs. These fluorides have hexagonal symmetry and, when grown on the GaAs( 111) 1 x 1 surface at a substrate plane epitaxy with: RF,(OOOl),(lOiO)IIGaAs(iii), The

in-plane

respectively. dislocations reaction

temperature

(T,) -5OO”C, they exhibit basal-

(110) and RF,(112O)IIGaAs(2ii)

misfits to GaAs are 3.6 and 1.5% for LaF, and NdF, These misfits are accommodated by interfacial edge as illustrated in Fig. 2. The interface exhibits no chemical

and the surface

of the fluoride

film is smoother

than the GaAs


Magnetic

Metal Structures

627

surface. Epitaxial overgrowth by the rare earth metals Dy, Ho, Er at T, 300°C resultst17] in basal plane epitaxy with: R(OOOl), (lOiO)IIRF,(OOOl),

(1120)

and R(2iiO)~~RF,(lOiO)

In this setting, the in-plane misfit between the rare earth metal and the fluoride is very large (-13%) and the interface is either semi-coherent or incoherent. Grazing incidence x-ray diffraction showedt171 that the rare earth metal films relaxed to their bulk lattice constant within -25 A from the interface. No strain could be detected in these films at room temperature. The magnetic properties of these films are discussed in Sec. 3.1.

(a)

(b)

Figure 2.

(a) High-resolution, cross-section transmission electron micrograph (HRXTEM) of -75 A thick film of NdF, grown on GaAs (iii) at 500°C. The Note the chemically abrupt but image is viewed along the GaAs[ii o] direction. physically rough GaAs surface and the planarization of the NdF, surface. Both LaF, and NdF, films exhibit basal-plane epitaxy on GaAs (iii) and act as seed films for rare earth metal epitaxy. (b) Higher magnification image of a region of the interface showing a pair of interfacial misfit edge dislocations and a step in the GaAs surface. Micrographs recorded by C. J. Chien.

628


In the case of epitaxy of the 3d magnetic transition metals (Fe, Co) on GaAs, Ag provides a suitable prelayer since it is thermodynamically stable in contact with GaAs and is immiscible interesting

with Fe and Co.

It is an

and useful fact that, by varying either the method of preparation

of the substrate or its orientation, any one of the three major axes of Ag, [OOl], [llO], [ill], can be selected as the growth axis. For example, we have shown previouslytlsj that pre-deposition of a few monolayers of Fe on the GaAs (001) surface, prior to Ag epitaxy, seeds the Ag [OOl] growth and a single

epitaxial

axis

orientation:

Ag (OOl), [loo]

jj Fe (OOl), [llO]

jj GaAs(OOl),

[llO]

Overgrowth of AS/Fe multilayers on such structures maintains this relationship throughout the layered structure. This is illustrated by the highresolution, cross-section transmission electron micrograph (HRXTEM) image in Fig. 3. The azimuth is along the GaAs[l lo] direction. The lattice fringes in the Fe films represent the (110) planes. At each interface the fringe contrast switches from the characteristic stripes of (110) fringes in the Fe to the square lattice of (200) fringes in the Ag. The azimuth in Fig. 3 is along the cube face normal of Ag confirming the rotation of the Ag lattice by 45” with respect to the Fe lattice. 3 -‘W\ ,

Figure 3. HRXTEM image of seeded, epitaxial Fe-AS superlattice grown on a GaAs( 111) substrate surface. The section is viewed along the GaAs [ilO] direction. This direction is also the Fe [il o] and Ag [loo] direction. The structure is seeded by a -9 A Fe prelayer followed by a -60 A Ag film. Micrograph recorded by C. J. Chien.


Magnetic

Metal Structures

629

This Fe prelayer seeding technique also works well with Co and permits[1g1[20]the growth of single orientation Co/Pt superlattices along the Pt [OOl] axis. This is illustrated

by the RHEED patterns in Fig. 4. The -12

A Co prelayer adopts a tetragonally-distorted

bee (ie., bcQ crystal structure

and grows[lgl in a parallel setting with GaAs, consistent with only a 0.4% inplane misfit between two unit cells of Co and the GaAs unit cell:

Co (OOl), [loo] 11GaAs (Ool), [ loo] RHEED Co-Pt (001) / GaAs (001)

274 Figure 4. RHEED patterns recorded during the seeded epitaxy of a 15 period Co/ Pt: (3 A Co-l 6 A Pt) superlattice. The patterns (top to bottom) were recorded of the clean GaAs(001) substrate surface; after growth of a 12 8, prelayer of bet Co; after growth of a -200 A film of Ag; and of the final capping film of 16 A Pt.

630


The Ag overlayer

grows on the Co in the same setting as on the Fe

prelayer. Pt, which has the same fccstructure as Ag, grows parallel to Ag, as seen from the identical symmetry RHEED pattern (Fig. 4). The RHEED pattern from each subsequent structure

parallel

with Pt.

Co layer shows that it grows

Each subsequent

in the

fee

Pt film grows in a parallel

setting to both the Co throughout the structure Co-l 6 A Pt) superlattice. controlled by the Co and To select Ag [l lo]

and Ag films. This relationship is maintained to the final Pt capping film of the 15 period (3 8, Thus the growth axis of the entire superlattice is Ag prelayers. as the growth axis, the Ag was deposited directly

onto the GaAs surface relationship:

at T, 5 100°C.

Ag (llO),

[ilo]

This

jj GaAs(OOl),

resulted

in the epitaxial

[loo]

Finally, the third major axis of Ag, [ill], was selected to be the growth axis by growing the Ag directly onto a clean GaAs( 1 1 1) surface at T, s 100°C. The Ag [l1 l] axis grew parallel to the GaAs[ 1 1 1] axis but the Ag film contained two types of in-plane crystallites, related by a rotation of 180” about the [l1 l] axis. This can be described as rotational twinning with the two epitaxial relations as follows: Ag (11 l), [Oil] Ag (ill), This twinning

[ilO]

jj GaAs (1 1 l), [Oil] jj GaAs (11 l), [Oil]

may arise from the lack of in-plane lattice matching

between

Ag and GaAs across the (111) interface. The lattice misfit is -28%. The preceding methods result in mutually exclusive orientations of the Ag films, ie., there is no mixing of orientations either in the Ag or of the subsequent

overlayers.

TEMt**j experiments. orientation

This was confirmed The use of these epitaxial

Co/Pt multilayers

magnetic properties

is shown

of the multilayers

by x-ray diffractionflgj

and

Ag films to seed selected

schematically are discussed

in Fig. 5 and the in Sec. 3.4.

The University of Michigan (UMI) group has pioneered the use of GaAs(ll0) substrates for preparing [l1 l] oriented CO/CU,[*~) CO/AU,[*~) and Co/Crf25] multilayers for studies of magnetic anisotropy and magnetoresistance. In the case of Co on GaAs(1 lo), Co adopts the metastable bcccrystal structure as shown originally by Prinzt26j and later confirmed by


[lll]Pt,

Ag t

Magnetic

[OOl]Pt,

Metal Structures

Ag

[l lO]Pt,

t

631

Ag

t

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Ag

-co Figure 5. Schematic diagram of selected-orientation epitaxial Co/Pt superlattices. In each case the superlattices were grown on a Ag film but the orientation of this film was selected by different growth procedures.

ldzerda et al.[27] The UMI group showed that bee Co can also be stabilized on Ge(ll0). Since Ge is a near-perfect lattice match to GaAs, it provides an initial smoothing film for heat-cleaned GaAs(ll0) substrates. The technique used by the UMI group is to grow an initial film of Ge, 500 A thick, at a temperature of 550°C on a GaAs(ll0) substrate, heat-cleaned at 600°C for 10 min. The substrate temperature was reduced to 50°C and bee Co grown onto the Ge to a thickness of -25 8, Following t!-., Co, a Cu film of -60 8, was grown to establish the [ll l] growth axis. Direct growth of Cu on Ge or GaAs results in polycrystalline films and is the reason for the intermediate film of bee Co. During growth of the Cu film, there is a transition in structure from bee Cu[l lo] to fee Cu[l 111. This occurs during the growth of the initial -50A[21] of Cu and has been studied using LEED by Marks et a/.I2a Between 38 and 55 A the residual 2-fold symmetry of the Cu LEED pattern disappears and is replaced by 6-fold symmetry, characteristic of a twinned film of Cu(ll1). This 60 8, thick Cu seed film is sufficiently

thin to permit measurements

of GMR in the subsequent

Co/Cu

632

Molecular

multilayers.paj

Beam Epitaxy

In the case of Co/Au multilayer

growth,

the Cu film was

replaced by Au which established the [l1 l] growth axis of the multilayer. Greig et al.t2Qj find that growth of a very thin (-10 A) film of Au on the bee Co

seed film results in a significantly

flatter surface than even a 200 8, Cu

film. The subsequent Co/Cu multilayer exhibits larger giant magnetoresistance (GMR) as a result. This finding has been reproduced by Clarke et al.pO) and an increase in GMR with decreased interface roughness found for a quite different seed film structure by the present authors (see Sec. 3.6). 2.2

Insulating

Substrates

In order to explore the magnetic

properties

of magnetic

multilayers

grown at higher substrate temperatures (>lOO”C) than on GaAs, it is necessary to select a seed-film-substrate combination for which the seed film is morphologically stable and the interface chemically stable at these elevated temperatures. In the case we have just described, Ag on GaAs, there is a tendency (see Sec. 3.3) for surface segregation of Ag during the multilayer growth and this effect becomes more severe as the substrate temperature is increased above 100°C. In addition, the Ag film may roughen as it is heated above this temperature, especially where it can expose facets of a lower surface energy than the initial film plane. Furthermore, where the magnetotransport properties of a magnetic multilayer are of prime interest, very thin (550 A) seed films are necessary to minimize the effect of current shunting through the seed film on the measurement of these properties. For these reasons, we have developed two novel epitaxial systems: Pt(l11) / basal-plane sapphire and Pt(ll0) / SrTiO,(l 10). Considering the first of these, Fig. 6 illustrates the observed epitaxial

relationship: Pt[ll

l] 11AI,O,

[OOOl] and Pt(ll0)

jj AI,O,

(1070)

The hexagonal, basal-plane unit cell of sapphire is indicated by the crosses and the lines on the left hand side of the figure. The (111) closepacked

plane

of Pt is shown

by the circular dots in the right hand Pt atoms form a hexagon which is unit of sapphire. Thus on purely

figure. Note that the second-neighbor only 0.9% larger than the basal-plane

geometric grounds, basal plane epitaxy with Pt(ll0) parallel to sapphire(lOi0)

of Pt on sapphire is expected This epitaxial as shown.

634

Molecular

Beam Epitaxy

LEED

Pt(111)/A!2203(0001)

AS?,O,(OOOl)

246 eV

308

61eV

Pt(ll1)

Figure 1. LEED patterns from sapphire overlayer of Pt(l11) grown at 600°C. overlayer, thickness 30 A, 61 eV.

(0001) substrate and 30 A-thick epitaxial (b) Pt (a) Sapphire substrate, 246 eV.


Magnetic

Metal Structures

RHEED (13keV) (111) Pt Seed Film on Sapphire

642

635

(0001)

645

Figure 8. RHEED patterns from epitaxial overlayers of 30 A-thick Pt(ll1) grown on sapphire (0001) substrate at 600°C. Patterns from two different films are shown along [112] and [l lo] azimuths. The uppermost pattern is from the final Cu film ([112] azimuth) of a [Co 9 A-Cu 9 A],3 multilayer grown onto the 30 A-thick Pt overlayer (seed film), beginning with Co, at a substrate temperature of -0°C. We can speculate on the intetfacial structure of Pt on sapphire by using the recent calculations of Guo et al.[31] which showed that the likely termination

of basal-plane

sapphire, in vacuum, is as illustrated

schemati-

cally in Fig. 9. The Al atoms (black dots) occupy some of the hollow sites in the quasi close-packed network of 0 atoms (open circles). The Al atoms form a hexagonal

lattice with each Al atom bonded to three underlying

adjacent 0 atoms. The observed epitaxial relationship could be established if the first layer of Pt atoms occupies hollow sites in the oxygen network. These sites are indicated by crosses. The origin of rotational twinning could be that steps on the sapphire substrate, if of minimum height (c/6 = 2.2 A), reveali3*] two possible sets of nucleation

sites for Pt,

related by a 180” rotation about the c-axis. Since the sapphire substrates used in the present work are cut to -0.5“ of the c-axis normal setting,

636

Molecular

Beam Epitaxy

surface steps will be present, though their height distribution

is not known.

XPS studies have confirmedfl*]

at a thickness

that the Pt film is continuous

of only 30 8, and that there is no evidence Pt and the substrate. amplitude

X-ray

reflectivity

of a chemical studies[‘*]

reaction between

suggest

a roughness

of -11 8, for Pt films of this thickness.

Figure 9. Schematic diagram of the Al-terminated (0001) surface of sapphire. The open circles represent the surface network of oxygen atoms bonded to the top Al atoms (black dots). The oxygen atoms are drawn at the size corresponding to the usual ionic radius (1.36 A)to reveal the quasi-close packed arrangement of oxygen. The observed epitaxiai relationship of Pt on sapphire (0001) could be established if Pt atoms occupy hollow sites, indicated by crosses.

Following growth of the Pt seed film, the substrate temperature was reduced to the growth temperature desired for the subsequent magnetic film. In the case of Co/Pt multilayers, the growth temperatures studied

were 100, 200 and 300°C. A schematic diagram of a typical structure is shown in Fig. 10. The Pt seed film was also used to prepare [l1 l]-oriented Co/Cu multilayers by growing the first Co film of the multilayer directly onto the Pt seed film. A RHEED pattern, along the [112] azimuth, from the top Cu surface of a [Co 9 A-Cu 9 A] ,s multilayer grown in this way is shown in


Fig. 8. Note that the multilayer the Pt seed film. discussed

Magnetic

Metal Structures

retains the same epitaxial

Such multilayers

exhibited

giant magnetoresistance

as as

in Sec. 3.6.

Pt

[ill]

[38. co -

Figure 10.

orientation

637

t5A

Pfl,

Schematic diagram of a [l 1l]-oriented Co-Pt multilayer grown on a Pt

seed film on a sapphire

The second

(0001) substrate.

new epitaxial

system

developed

to seed the growth

of [l1O]-oriented Co/Pt and Co/Cu multilayers was Pt(l1 O)/SrTiO,(l 10). In this case, the SrTiO, substrate has a cubic perovskite structure with a lattice

constant

(a = 3.9051

that of Pt (a = 3.9240 at 25°C).

8, at 25°C)

only 0.5% smaller

The growth of Pt(ll0)

and Cu(ll0)

than on

Co(ll0) on Pt (1 lO)/SrTiO,(llO) is described in detail elsewhere.t13] Here we summarize the key result that, at a substrate temperature of 600°C Pt grows with a well-defined single epitaxial relationship: Pt[OOl] ]I SrTiO,[OOl]

and Pt(ll0)

]I SrTiO,(llO).

However, the surface of the Pt film was found to be facetted at all thicknesses, up to at least 1000 8, exposing (111) facets. The LEED patterns showed that the surface was terminated with ridges running parallel to SrTiO,[OOl].

838

Molecular

Beam Epitaxy

In order to grow [l1O]-oriented Co/Cu multilayers, a Co film, typically 20 A thick, was grown onto the Pt at a substrate temperature of -0°C. The Co film is partially interdiffused

with the Pt, but it retains an fee crystal

structure to at least 20 8, The deposition a polycrystalline

film, possibly

because

of Cu directly onto Pt resulted in of interdiffusion

and intermetallic

compound formation. However, the growth of Cu onto the Co resulted in a [llO]-oriented film with a surface morphology which depended on the growth temperature. Growth at temperatures below 300°C resulted in faceting of the Cu surface resulting in (100) facets. On the other hand, growth at temperatures above 300°C resulted in smooth, unfaceted surfaces. For the growth of Co/Cu multilayers for magnetoresistance measurements, the substrate temperature was in the range 0-100°C. In this temperature range, the surface of the growing multilayer was essentially unfaceted since the Cu thicknesses in the multilayer (5-100 A) were sufficient for the initial (111) facets to be smoothed out but insufficient for development

of (100) facets.

3.0

STRUCTURAL AND MAGNETIC PROPERTIES ARTIFICIALLY-LAYERED MAGNETIC METAL

3.1

Rare Earth Metal Sandwich

OF STRUCTURES

Structures

The magnetic properties of rare earth metal sandwich structures grown on GaAs substrates with a LaF, prelayer have been reported elsewhere;t17) here we summarize the essential results. Sandwich structures of Dy, Ho, and Er have no measurable strain at room temperature and SQUID measurements of such films in the thickness range 25 to 4000 A confirmed that the ferromagnetic

ordering temperature

was close to the

value for bulk crystals though the transition was somewhat broadened, especially for the thinnest films (cl00 A). This broadening may be related to inter-facial disorder

at the lattice-mismatched

interface.

In the case of

Ho, the saturation magnetization measured at 4 K and in high fields corresponded to nearly 100% of the value for bulk single crystals. However, for Dy, the magnetization reached -76% of the bulk value indicating that some of the Dy atoms in the film do not contribute to the magnetization, possibly as a result of inter-facial Dy-La exchange reactions forming paramagnetic

DyF, during growth.


Recently, magnetic

x-ray

diffraction

Magnetic

measurementst33jt341

order of a 2000 A film of Dy sandwiched

were carried out as a function

of temperature

the National Synchrotron Light Source. ferromagnetic hysteretic.

Metal Structures

transition

occurred

of structural

between

639

and

LaF, films

on the beam line X20A[351 at

These studies confirmed

close to the bulk temperature

This contrasts with the dramatic suppression

that the but was

of 7, reported by

Kwo[~~] for Y/DyiY sandwich structures and the complete absence of ferromagnetic ordering in Y/Dy/Y superlattices.t3r)t36j It was conjecturedt36jt37j that these latter effects were a result of coherency strain which was introduced into the Dy during growth and caused a c-axis contraction. However, no measurements of coherency strain or its temperature dependence were providedt36)t37) to support this view. In view of this and the absence of coherency strain in as-grown LaFJDy/LaF, sandwiches, we designed a structure to test the influence of coherency strain on ferromagnetic ordering in Dy. Figure 11 shows a schematic diagram of this structure. The thick Er film is relaxed to its bulk lattice constant while the Dy overlayer has a larger in-plane lattice constant than Er and will be subjected to in-plane compression. Sandwiching the Dy between Er films is an attempt to make the strain symmetric. The LaF, overlayer provides protection from atmospheric oxidation. A priori, it was not clear to what extent the strain in the Dy would be relaxed by misfit dislocations. However, room temperature x-ray diffraction studies confirmed that the Dy film was under compressive in-plane strain and had a c-axis lattice constant expansion of 0.27 + 0.02% compared with the bulk value. The expected c-axis lattice constant expansion is 0.58%, assuming a coherent, lattice-matched interface and bulk value of Poisson’s ratio which is probably reasonable for a film of this thickness. Thus, one may conclude that misfit dislocations

have reduced but have not fully relaxed the misfit-

induced strain, SQUID measurements for this structure showed that the ferromagnetic ordering temperature for the Dy film was significantly higher than for bulk Dy as illustrated in Fig. 12. Magnetization data is shown for four different values of applied field along the a-axis (the easy axis of a bulk Dy crystal).

The transition

(TN) from the paramagnetic

to

helically-ordered antiferromagnetic state is evident from the cusp at about 180 K. The transition temperature from the helically-ordered state to the ferromagnetic state increases with field, as for bulk single crystals of Dy. The arrows indicate the onset of ferromagnetic ordering in bulk Dy. At all fields, the ferromagnetic ordering in the film is completed at higher temperatures than the onset of ferromagnetic ordering in bulk Dy.

640

Molecular

;

-e 6

ii

Beam Epitaxy

4g Ill

B

-4

&l

-4:

\’

T :.: :.:. ~ \ ::: :::: :.:. ..i :>i :j:: :.:: :::: i:;; :::. .._. ..i :j:j :::: :::: :::: ._._ ::; ::::

:::I ::’ :::: :::: :.:. ..i . .. .


* 2.5x10-3

-

I

I

.

Magnetic

1

Metal Structures

I

I

I 0

I

0.5, 2, 5, 10 kOe

A

(Increasing from lower to top curve)

0

641

l.* . ..(....

, 100

Temperature

-

. . . . . (...~***~~)**~~.****~ 200

300

(K)

Figure 12. SQUID magnetometer data for a sandwich structure comprising: GaAs (111)/50 A LaFsI2000 A ErI500 A Dy/lOOO A Erll 00 A LaF,. Note the paramagnetic-to-helical transition at 178 K and the helical-ferromagnetic transition at lower temperatures. The arrows indicate the onset of ferromagnetic ordering in bulk crystals. For each field, ordering of the films occurs at higher temperatures than for a bulk single crystal.

We attribute this to the imposed elastic strain in the film, since, in the bulk crystal, ferromagnetic ordering is accompanied by a discontinuous c-axis expansion and in-plane contraction. By mechanically imposing a c-axis expansion in the film, we have made it energetically favorable for ferromagnetic ordering to occur at higher temperatures. Recently, it has been shown that by sandwiching Dy between Lu films in multilayerst3Q] or single sandwichest40] a larger coherency strain can be generated in the Dy, with the same sign as for Dy/Er, leading to an increase in 7, by -70°C. These results demonstrate the controlling influence of coherency magnetic ordering behavior in rare earth crystalline films.

strain on the

642

Molecular

Beam Epitaxy

3.2

Fe/Ag Films and Multilayers Seeded epitaxy of single crystal films of Ag grown on an Fe prelayer

on GaAs(001) substratest’s) or on an Fe prelayer grown on a ZnSe film on GaAs(001) substratest4’) form the basis for [OOl] oriented AS/Fe single or multilayer

structures

(see Fig. 3) which

exhibit

perpendicular

anisotropy

for Fe films 13 ml thick. This anisotropy persists only at temperatures at or below room temperature. There is general agreement that it arises from a surface anisotropy component which competes with the demagnetizing field. The surface anisotropy increasingly dominates as the film thickness decreases. Brillouin scattering studiest4*) of the surface anisotropy of Fe single crystals, coated with epitaxial Ag and Au, reveal a surface anisotropy of similar magnitude to that measured for the ultrathin Fe films. Moreover, the positive surface anisotropy for Fe/Ag interfaces measured from spin-wave resonance spectra in relatively thick (- 900-2000 A) epitaxial Fe filmst43) is sufficient to overcome the demagnetizing field in ultrathin Fe films. Thus the perpendicular anisotropy is an interface effect due to symmetry breaking at the interfaces rather than to an intrinsic property of ultrathin Fe films. Fe/Ag multilayers have not, to date, exhibited giant magnetoresistance effects. However, oscillatory exchange coupling across Ag spacer films in MBE-grown Fe/Ag/Fe (001) sandwiches, grown on the (001) facets of Fe single crystal whiskers, has been observed by the NIST (National Institute for Standards and Technology) group t44l who earlier[45) had studied oscillatory exchange coupling in Fe/Cr/Fe sandwiches, grown on Fe(OO1) facets, using SEMPA (scanning electron microscopy with polarization analysis) as the local area probe of magnetic coupling. These elegant experiments will not be reviewed here, since we are confining our review to structures grown on semiconductors and insulators. We conclude this section

by pointing

out that the Fe/Ag seed film method

has

been utilized to improve the structural quality of Fe/Cr/Fe multilayers grownt46) on GaAs(001) substrates, and recently Griinberg and coworkerst47)t46) utilized the Fe/Ag seeding approach to prepare sandwiches of Fe/Cr/Fe, Fe/Au/Fe, and Fe/Al/Fe on GaAs(001) substrates. This work is reviewed in the following section, 3.3

FelAg-Seeded Griinberg

Sandwiches

and co-workers

of Fe/O/Fe,

of the KFA-Julich

Fe/Au/Fe,

and Fe/AI/Fe

group have utilized the

Fe/Ag seed film approach to grow [OOl]-oriented sandwich structures of Fe/Cr/Fe,t47) Fe/Au/Fe,t46) and Fe/AI/Fet48) on GaAs substrates for studies


of oscillatory

exchange

A) films were annealed

grown

coupling.

Metal Structures

643

Both the Fe (-10 8, thick) and Ag (-1500

at 100°C.

at 300°C to improve

Magnetic

Following

growth,

the structural

quality

the structure

was

of the Ag film as

judged by SPA-LEED (spot profile analysis LEED) and RHEED. Grirnberg et al., moved the substrate behind a stationary shutter to grade the spacer film (Cr, Au, Al) thickness in the form of a linear wedge across the substrate. The thickness of the spacer varied from -0 at one side of the sample to 30-100 A at the other side. This allowed the coupling dependence on the spacer film thickness to be probed on one sample using the magneto-optical Kerr effect or Brillouin light scattering as a local area probe of magnetization along the wedge. A key feature of this method is that scatter due to variations in the properties of the Fe films is strongly suppressed. In the case of Cr and Au spacers both long and short period oscillations in coupling were observed. The long period was = 18 A for Cr and = 8 8, for Au. These oscillations were from FM to AF coupling. The short period was only= 2 ml in both cases. However, the short period was fully resolved only for a growth temperature of 250°C; it was not fully resolved for growth temperatures s 250°C. In the case of Al only a long (M 18 A) period was found. Our present theoretical understanding of the coupling oscillations is incomplete. There is agreement that the oscillations arise from excitations of the Fermi surface of the spacer film which couple the ferromagnetic films through a spin density wave. The periods correspondt4s)t50) to specific inter- and intrazone wavevectors. In the case of Cr and Au spacers, a short period of ~2 ml is predictedt50) in agreement with experiment. However, the wavevector responsible for the long period in Cr and Au has not yet been identified. In concluding this section, it should be pointed out that Grinberg et al.[4e) clearly demonstrate that the short period interlayer oscillations in interlayer coupling do not simply represent a switch from AF to FM coupling but are more complex and contain a component 3.4

of biquadratic

coupling.

Seeded Epitaxial Co/Pt Superlattices

Co/Pt and Co/Pd multilayers are well knownt51)t521 to exhibit perpendicular magnetic anisotropy when the Co films are very thin (58 8, for Co/ Pd and ~10 A for Co/Pt).

This anisotropy

is observed

for multilayers

prepared by a variety of techniques, including RF sputtering,t51)t521 DC magnetron sputtering,t53)t54) and conventional evaporation.t55)-t57) Depending on the growth conditions,

perpendicular

magnetic

anisotropy

can be

644

Molecular

Beam Epitaxy

obtained in the multilayers and the Kerr rotation (Q, ie., the rotation of the plane of polarization of a reflected light beam by the magnetization of the sample, increases towards the blue spectral region. These are desirable properties in thin-film media for high-density magneto-optical storage of information and as a result there is considerable current interest in the magnetic and magneto-optical properties of Co/Pt and Co/Pd multilayer structures. For such applications, the room temperature coercivity of the thin film should be several kOe and the shape of the hysteresis loop should be square. Although such properties can be obtained by empirically adjusting the growth conditions, in conventional deposition techniques, key questions remain, What is the origin of the magnetic anisotropy in the multilayers? and, What are the structural parameters which control magnetic anisotropy? Unlike the Fe/Ag and Fe/Au couples discussed above, Co/Pd and Co/Pt couples exhibit extensive mutual solid solubility. As a result, it is not clear to what extent the high-temperature perpendicular anisotropy is influenced by interfacial alloy formation, To resolve these issues, the authors and co-workers have studied the structural and magnetic properties of Co/Pt multilayers oriented along the three major axes of Pt using the techniques of seeded epitaxy outlined in Sec. 2. A key initial findingt”] was that, for structures of [Co 3 A-Pt 16 A],,, aligned along the [ill] axis of Pt, perpendicular anisotropy with full remanence and large (~14 kOe) coercivity was found only for the [l1 l] growth orientation. This is illustrated in Fig. 13. The [OOl]-oriented multilayer exhibited an in-plane easy axis and the magnetization could not be saturated in the perpendicular direction. The [l1O]-oriented multilayer exhibited a uniaxial in-plane anisotropy. Each of the multilayers was grown at 100°C and in each case the Ag prelayer was -200 8, thick. Following this finding of the dramatic dependence of magnetic anisotropy on multilayer growth axis, the structural and magnetic properties of the multilayers were explored by a variety of techniques aimed at identifying differences due to changes in the growth axis. A brief account of the findings is given in the following section and a more detailed description can be found in the original papers. Preliminaryx-ray diffraction studiest19)t20) of these multilayers showed a more rapid fall-off of satellite intensities with increasing order for the [ll l]-oriented multilayer when compared with the [OOl]-oriented multilayer. This was true for low-angle specular x-ray reflectivity and highangle (0-26) scans. The tentative initial conclusiont19) that the [ill]oriented multilayer had a greater degree of chemical interdiffusion than the [OOl]-oriented multilayer was based on an oversimplified interpretation of satellite intensities. For example, inter-facial disorder, due to physical


roughness

Magnetic

Metal Structures

and defects, can modify the shape and intensity

the satellites

in addition

to interdiffusion.

these effects requires an extensive

ing data. The initial x-ray experiments dent studies of interface character. photoelectron

diffraction)

during their growth.

Thus

were followed

Hermsmeier

to examine,

of

between

of x-ray scatter-

by several indepen-

et al.t58t used XPD (x-ray

in situ, formation

In XPD, the angular

distribution

distinguishing

data set and modeling

645

distribution

of the interfaces of photoelectrons,

emitted from the near-surface atoms of an epitaxial overlayer, is recorded. Coordination of overlayer atoms is determined from the presence or absence of diffraction peaks due to forward scattering. From the coordination as a function of coverage, the growth mode and presence of interdiffusion can be deduced. XPD from the Pt film of the initial multilayer period and at various stages of Co film growth, as well as from subsequent Pt films, was recorded for the [ll l] and [OOl]-oriented structures shown in Fig. 5. For both orientations, the interfaces for a 3 A Co film, were intermixed on the scale of 2-4 ml (4-6 A).

(3A Go/16A

P’llS

Sample I

No. 1 [ill]

-Sample No. 2 11101

0.1

1 0

ek(deg)

-0.1

I

-16

I

I

-12

I

I

-6

I

I

81

-4

-

o

I

4

-

13

6

” 12

J 16

H (kOe)

Figure 13. Magnetic hysteresis oriented Co-Pt superlattices with of incident laser = 0.633nm. The surface. i.e., along the directions

loops recorded at 2O”C, using Kerr rotation, for (-3 A Co-l 6.6 A Pt) x 15 periods. Wavelength magnetic field was applied normal to the sample indicated.

646

Molecular

Beam Epitaxy

Subsequently,

the interfaces

in complete

multilayers

were exam-

ined by synchrotron high-angle x-ray diffraction and low-angle x-ray reflectometry. Yan et al.t5s) and Mariner0 et al.t60) reported 0-20 scans for [l 1 l] and [OOl]-oriented multilayers record the scattered intensity

using a position-sensitive detector to along directions both perpendicular and

parallel to the surface (Q, and Cl,, respectively). Interestingly, the peak shapes in Q, were different for the two orientations: near-Lorentzian for the [l1 l] orientation but Gaussian for the [OOl] orientation. If the data are analyzed by integrating the peak intensity only in Q,, then the fall-off of integrated intensity with satellite order is determined by both large scale roughness and atomic scale interdiffusion. However, integration over Q, and Q, gives a fall-off which is dominated by atomic scale interdiffusion. Using this method, an estimate of interdiffusion for the same multilayers described in Sec. 2.1 (magnetic data shown in Fig. 13) gave a Co diffusion length of 2.6-2.8 A at each interface for both [ll l] and [OOl] orientations. Thus, the Co atoms were estimated to extend over a total width of 6.4 A) which is = 3.0 ml for [ll l] and = 3.3 ml for [OOl]. This result is consistent with the XPD data which showed that Co was present 2-4 ml from the nominal interface for both orientations. Thus, the atomic scale interdiffusion and the major difference

was quite similar for the two orientations, between the orientations was a much larger

degree of in-plane disorder for the [ll l]-orientation. Other techniques also found that the interfaces

were not atomically

abrupt. For example, grazing-incidence x-ray reflectometry studiest6’) of [OOl]-oriented multilayers showed clear evidence for interdiffusion in that the data could not be modeled without incorporation of Co in the Pt films and vice versa. High-resolution electron microscopy studies by Chien et al.f22)t621and subsequently by Cho et al.t631t64) and Zhang et al.t65)t66) confirmed that the interfaces were not chemically abrupt for either orientation. In the case of [Ill]-oriented multilayers, the work of Cho et al.t63) showed that twinning was a dominant structural defect and not stacking faults as earlier suggested.t22)t62) Twin boundaries are a possible source of the in-plane structural disorder found by Yan et al. Multilayers of the other two orientations

were found to be untwinned.t1g1t221

Based on the finding of interdiffused interfaces and the relation between magnetic anisotropy and partial chemical ordering in N&Fe alloys, faces,

Chien et al.f62) suggested a partially chemically-ordered

that, near the multilayer interCoPt, phase (Ll, phase) formed.


Magnetic

Metal Structures

647

They suggested that if this phase existed as a uniform composition slab with abrupt interfaces to Pt, then a Neel-type anisotropy model could predict the correct directions

of easy axis in the multilayers.

Since this

suggestion was novel, and since the bulk phase diagram shows that this phase is stable at the multilayer growth temperatures, we have undertaken a search for this phase in both [l1 l] and [OOl]-oriented

multilayers.

Figure

14 shows a schematic diagram of this phase and of the random alloy. Since the unit cell of the ordered phase is doubled in real space, superstructure diffraction peaks at half of the reciprocal lattice vectors of the fundamental peaks provide a fingerprint for chemical ordering. For example, the presence of (110) peaks provides an indication of chemical ordering and their width provides information on the coherence length of such ordering. Grazing-incidence synchrotron x-ray diffraction experiments were carried out by Toney and Rabedeau[6fl-t6g] at the National Synchrotron Light Source, Brookhaven to search for chemical ordering in Co/Pt multilayers grown at 100, 200 and 300°C. Both phi-scans and radial-scans were made to search for the (110) diffraction features indicaFigure 15 shows phi-scans for each growth tive of chemical ordering. temperature. The scan traces a circle in reciprocal space through the (110) Bragg peaks. (110) peaks were observed at all temperatures with the expected 60” angular separation for multilayers, twinned on the (111) plane. The amount of ordering increased with the growth temperature. Radial scans also showed the (110) peaks, and from the peak widths the coherence length for chemical ordering was estimated to be only -10 interatomic spacings in the (111) plane for all growth temperatures. X-ray diffraction measurements were also conducted on [OOl]-oriented multilayers grown on GaAs to investigate the formation of a chemically ordered alloy for this growth orientation. We observed x-ray scattering at the location expected for diffraction from an ordered alloy, e.g., (100) diffraction peaks. Unfortunately, this location also corresponds to the expected position of xray scattering features due to the termination

of the GaAs lattice.

Despite

this ambiguity, we believe that some of the observed intensity is due to ordered alloy formation because this feature was quite broad, similar to our observation for ordered alloys in the [ll l]-oriented multilayers. The diffraction data clearly demonstrate the doubling of real space periodicity characteristic of alloy ordering. However, such doubling can, in principle, arise from one of three ordered phases reportedpot for Co-Pt alloys: CoPt,, CoPt, and Co,Pt. We believe that the first of these

648

Molecular

Beam Epitaxy

(illustrated in Fig. 14) is present for the following occurs[581[5Q1~11 over 3-4 monolayers for Co.

Second, we have made parallel studies

films grown,

by co-evaporation,

reasons.

First, mixing

which favors a Pt-rich environment

on sapphire

high as 600°C and find that CoPt, form&‘*]

of Pt-rich epitaxial

(0001) at temperatures

spontaneously

alloy as

during growth at

temperatures at least as low as 300°C. Finally, since during multilayer growth, Co arrives at a pure Pt surface and has a tendencyP3] to prefer Pt neighbors, it is natural that the Co forms the first monolayer of the CoPt, phase.

CoPt3 Disordered

CoPt3 Ordered

FCC

Figure 14. Schematic diagram of random alloy (Co to Pt ratio 1:3) and ordered CoPt, phase. The unit cells and arrangement of atoms in (111) planes are shown. Note the doubling of real space periodicity for the ordered structure.


I

alloy

Magnetic

I

Metal Structures

649

I

alloy

alloy

‘2; *z 0.0075 -e d3 0.0050

0.0025 -100

-50

0 azimuth

50

100

(“)

Figure 15. Phi scans through (110) Bragg peaks for 100°C 200°C and 300°C multilayers. The CoPts ordered alloy peaks and peaks due to tails of diffraction from the multilayer rods (crystal-truncation rods--&) are indicated. Note that the intensity scales for the data have been offset to show the progressive increase in alloy ordering with film growth temperature: 0 - lOO”C, 0 - 200°C and A - 300°C.

Estimates of the amount of CoPt, formed were made from the ratio of the integrated intensities of the (110) superstructure peaks and the {iii} fee multilayer peaks. For the 100 and 200°C multilayers, the fraction of Co in the ordered phase is 20 f 10% and 60 + 30%, respectively (this is the ratio of Co in the ordered alloy to total Co in the multilayer). For the 300°C multilayer, we estimate this fraction may be as high as 90%. Such large amounts of chemical ordering can be expected to influence the magnetic behavior of the alloys. Indeed, magnetic data for these multilayers support this view. Figure 16 shows polar Kerr loops for [Co 3 &Pt 15 &,s multilayers grown at 100, 200 and 300°C onto sapphire substrates as illustrated in Fig. 10. The easy axis of magnetization is perpendicular to

650

Molecular

Beam Epitaxy

the film plane at all three temperatures. increasing temperature as the coercivity surementsp4)

confirm

comes more positive) correlation

between

Loop squareness increases with falls. Magnetic anisotropy mea-

that the magnetic with increasing

chemical

anisotropy

increases

growth temperature

ordering

and magnetic

(krr

be-

suggesting

anisotropy.

a

In fact,

full chemical ordering is known, from bulk alloy experiments, to lead to a reduced Curie temperature since ordering reduces the number Co-Co neighbors. However, the coherence length for chemical ordering within the multilayers remains very small (530 A) at all three growth temperatures so one would not expect the Curie temperature to be strongly modified.

0.1 5 & 0.0 .z -0.1 g -0.2 E g 0.1 5 & 0.0 .g

; Y E s E _; Y

I

2oo”c

3.75kOe

-0.1

I

I

-0.2 0.1

I I

I

I I i

3oo”c

0.0

1.66kOe

-0.1 -0.2 -20

0 10 magnetic field [kOe] -10

20

Flgure 16. Magnetic hysteresis loops, recorded at 20°C using polar Kerr rotation, of three [l 1 l]-oriented Co/Pt multilayers: [Co 3 &Pt 15 A],, grown on 30 A thick Pt seed film on sapphire (0001) at substrate temperatures, 100, 200 and 300°C.


The discovery implications.

Magnetic

of chemical ordering in Co-Pt multilayers

It shows that not only chemical

should be included of these multilayers.

Metal Structures

in any theoretical

The model proposed

has several

mixing but chemical

models for the magnetic

651

ordering

anisotropy

by Chien et al.t6*) is probably

too simplistic since it ignores several factors including lattice strain and the magnetic polarization of Pt by Co nearest neighbors. Also, the profile of Co concentration through the interfaces remains to be determined before a realistic theory can be developed. The discovery also raises the question as to whether chemical ordering in magnetic multilayers is a more general phenomenon since many other materials systems exhibit solid solubility with chemical ordering. These include Fe-Pd, Ni-Pd, Fe-Pt, NiPt, and Co-Ni. It also implies that ultrathin films of the CoPt, alloy might show perpendicular anisotropy if partially ordered. This is indeed the case and is discussed in the following section. 3.5

Co-Pt alloy

films

Following the finding of the Ll, phase within Co/Pt multilayers, we studied the structural and magnetic properties of seeded, epitaxial, [l li]oriented films of CoPt, alloys. These films were grown directly onto sapphire substrates, since it was found that the alloy grew in the same epitaxial relationship with sapphire as Pt (see Sec. 2.2). A key finding was that perpendicular anisotropy for the alloy was present only at elevated growth temperatures, around 300°C. This is illustrated in Fig. 17 which shows polar Kerr loops for -200 8, thick alloy films grown at temperatures from 100-6OO”C. Note that at the lowest and highest temperatures, the loops show no remanence anisotropy.

and nearly zero coercivity,

indicating

in-plane

On the other hand, the 300°C loop shows full perpendicular

remanence and the maximum coercivity. Recent synchrotron x-ray diffraction studiesp*] show that chemical ordering is present in the films grown at 300 and 6OO”C, and that the degree of order is greater at the higher temperature. Structural and magnetic properties of these films are still in progress and will be reported elsewhere.p2)p4) However, we have shown that unlike the multilayers, the magnetic anisotropy is not sensitive Indeed, perpendicular anisotropy is to the growth axis of the alloy. foundp6)p71 for weakly-textured films grown onto amorphous substrates provided that the substrate

is held at -300°C.

Such films can be used as

an alternative to Co-Pt multilayers as MO media and have a larger Kerr rotation and potentially larger carrier-to-noise (CNR) ratio compared with

652

Molecular

Co/Pt multilayers alloy media. signal P

l

Beam Epitaxy

or with the present

This is illustrated

generation

of amorphous

TbFeCo

in Fig. 18 which showsF6) the static MO

(8K’ + EK2) ln of thick (~1000 A) films of a CoPt, alloy, a Co-Pt

multilayer, and TbFeCo amorphous alloy film. R is the reflectivity, 8K the Kerr rotation, and &K the Kerr ellipticity. Note that while the static Kerr signal for the TbFeCo falls towards the blue spectral range, the signal increases for the CoPt, alloy film and the Co-Pt multilayer. At 420 nm the static Kerr signal is greater by 2.9 dB for the alloy than for the multilayer.

t

T,(“c> H,(kOe) ‘k 600

0.00

/

0.17 /

300

2.20

o.Jr

2/_ 100

1.30

0.15

0.20

0.20

Figure 17. Magnetic hysteresis loops, recorded at 20°C using polar Kerr rotation, of 200 A thick films of CoPt, grown at the temperatures indicated directly onto sapphire(0001). The Kerr rotation and coercive fields are indicated. epitaxial with CoPt,(l 11) 11sapphire(OOO1) (see text).

The films are

Artificially-Layered Magnetic Metal Structures

0.35

I

I

I

653

I

0.30 AdB

p 0.25 EA $_ 0.20 ax “,y T 0.15 2

0.10 0.05 0.00

Tb2s.#%roColo)7s.5

.

.

substrate side

350 450 550 650 750 850 wavelength[nm]

Figure 18. Static MO signal Fl+J,* t Pt multilayer

-

and TbFeCo

amorphous

of (>lOOO A)thick CoPt, alloy film, Co/ alloy film as a function of photon wave-

EK2)”

length. See text.

Recently we have reportedp7] dynamic testing of quadrilayer structures comprising: glass substrate/400 8, S&N,/-200 8, CoPts/200 8, S&NJ 500 8, Al on complete, 3.5 inch diameter glass discs. The alloy films were grown in the MBE system at a substrate temperature of -300°C to maximize coercivity. CNR values of >60 dB were obtained at 488 nm This is a highly promising result and has stimulated interest in Co-Pt alloy media films for MO storage.

654

Molecular

Beam Epitaxy

3.6

Giant Magnetoresistance

in MBE-Grown

Another topical area in magnetic playing

a key role is the giant

reported for MBE-grown

Co/Cu Multilayers

metal multilayers

magnetoresistance

in which MBE is

(GMR)

effect,

first

Fe-Cr multilayers.t4)-t61 This effect was observed

in structures with alternating

epitaxial Fe and Cr films grown onto GaAs(001)

substrates. In the case where Fe films of 30 A thick were alternated with thin (g-l 8 A) Cr films in multilayerst5)t6] with 40-80 periods, the structures exhibited a giant negative magnetoresistance, ie., the resistance decreased from the zero field value by as much as 50% when the magnetic field and the current were both along the same [l lo] axis in the (001) plane of the films. Both magnetometry and neutron diffraction studies showed that, in zero field, the Fe films were aligned antiferromagnetically. In an applied magnetic field, antiferromagnetic coupling is overcome and the spins in the Fe films become aligned ferromagnetically. The giant magnetoresistance effect correlates with this spin alignment behavior in that the magnetoresistance saturates at the field where the spins are aligned ferromagnetically. Baibich et al.t6) and Camley and Barnasps) suggested that the giant magnetoresistance effect arises from spin-dependent transmission of conduction electrons through the Cr films. This mechanism requires non-specular scattering of the conduction electrons at the interfaces. If the interfaces were perfectly sharp and smooth then only specular reflection and diffraction of the electrons would occur and this would not contribute to the magnetoresistance. The magnitude of the effect is in fact knownp91 to depend on film growth parameters such as growth temperature which may affect interface roughness and grain size. Following the initial GMR discovery, Parkin et al.p] found that for magnetron-sputtered Fe/Cr and Co/Cr multilayers, the GMR did not simply saturate as the Cr film thickness increased but oscillated. These oscillations were foundp) to correspond to oscillations from ferromagnetic to antiferromagnetic

interlayer

coupling.

This stimulated

considerable

inter-

est and led to the discoverypI-ts) of a large number of different magnetic systems exhibiting oscillations in both interlayer coupling and GMR. Theoretical models of the coupling, based on approximatetsO) and full band structuretEll calculations have predicted specific phases and periods of the coupling which depend on the crystallographic orientation of the nonmagnetic spacer film. One way to test such theories is to compare the oscillations in GMR for specific orientations, selected by seeded epitaxy techniques, with the theoretical predictions. We have chosen to examine Co/Cu multilayers, oriented along Cu[lll] for several specific reasons.


Magnetic

Metal Structures

655

Firstly, Co/Cu multilayers exhibitt8)tg) the largest GMR of any materials system. Moreover, the original studies by Parkin et al.t8ttg) of this system were for polycrystalline

[l1 l]-textured

multilayers

of [OOl] and [l1O]-oriented

crystallites.

and GMRt28)t831in epitaxial

[l1 II-oriented

ity orientations,

Studies

with detectable of interlayer

multilayers,

amounts

couplingts*)

with no other minor-

have shown no evidence of such oscillations.

On the other

hand, oscillations in interlayer coupling have been observed for epitaxial [O01]t84)t85)t88)and [1 10]t87)-oriented multilayers. This has raised the question as to whether oscillations in interlayer coupling are present for purely [ll l]-oriented multilayers. We have prepared [l Ill-oriented multilayers on both GaAs and sapphire substrates using seeding structures described in Sec. 2. Two types of seed structure were used for GaAs substrates. Initially, multilayers were preparedt*s) using GaAs(ll0) substrates with the sequence:

with the multilayer grown at a substrate temperature of 100°C. Structures with Cu spacer thicknesses (tcJ in the range 5-50 8, were examined. These showedt28) little evidence for AF coupling and no oscillations in saturation magnetoresistance, AR/R. Subsequently, multilayers were grown simultaneously on GaAs( 1 1 1) and sapphire (0001) at a growth temperature of 0°C. These structures were: GaAs(l ll)/Ag

40 &Pt

30 &[Co

16 A/Cu tcJ,s/Co

sapphire (OOOl)/Pt 30 A/[Co 16 A/Cu t&s/Co Both of these structures

exhibited

16 &Pt

16 &Pt

30 8,

30 8,

a peak in GMR at a Cu thicknesses

near

9 8. The peak GMR for the sapphire-based structures was much larger than for those grown on GaAs. Figure 19 shows the resistance vs in-plane field curves for a sapphire-based and 295 K. The low-temperature measurements,

multilayer with tcu = 9 A, measured at 3.5 AR/R value reached 40%. In all our MR

we define the saturation

magnetoresistance,

AR/R, as the

maximum change in resistance over the measurement field range divided by the high-field resistance. This definition is chosen because the zerofield resistivity is often sample history-dependent and therefore not a welldefined quantity. The inset in Fig. 19 shows the normalized film magnetization versus in-plane field for this sample at 295 K. There is a large remanent magnetization in zero field indicating that a large fraction of the sample is ferromagnetically coupled, presumably through local ferromagnetic bridges.

656


10 -

I

O-

4

0

-4

Field

(T)

Figure 19. Resistance vs in-plane field curves measured at 3.5 K and 295 K for a UHV evaporated Co/Cu superlattice of the form, sapphire(OOOl)/Pt 30 A/[Co 11 A-CU 9 &,/Co I I A/Pt 30 A. The inset shows the normalized film magnetization versus in-plane field at 295 K. The saturation magnetization per Co volume is, within experimental error (-lo%), that of bulk Co.

Figure sapphire-based significant

20 shows

the

dependence

of AR/R

The

peaks

superlattices.

AR/R

on tc, near

scatter in the data and it is not evident whether

for a series

of

9 8, but there

is

or not there is a

second maximum near tcu = 18 A. It is clear that one needs to reduce scatter in the data to determine the position of the first maximum and to find whether subsidiary maxima are present. The scatter is due to variations in thickness and structural perfection in the Pt seed film as well as in the Co films. Such variations can be reduced by preparing a wedged multilayer on a single substrate comprising uniform thickness Pt and Co films but Cu spacer films which all vary linearly in thickness across the substrate. The magnetoresistance can then be measured for a series of samples cut from the wedged

multilayer.


I

I

Magnetic

I

I

I

Metal Structures

I

I

l1

93 0

I

1

00 3.5-4.2 30 -

657

K

295

K

0 0

g20

.s

0

n

C

a

%

0

l

0

0

10 -

0 0

l 0

01 0

I

I

I

5

I

I

I

layer

thickness

I

15

10

Cu

a

I

20

(A)

Figure 20. Saturation magnetoresistance vs Cu spacer layer thickness & for a series of Co/Cu superlattices of the form: sapphire(OOOl)/Pt 30&(Co 16&Cu tcUlla /Co 16 A/Pt 30 A. Cl (3.5 K) and w (295 K), resistance saturated by applying fields of up to 6 T; 0 (4.2 K) and 0 (295 K), resistance measured in fields of up to 16 kOe and almost completely saturated.

We have preparedtEE such structures using a novel method which takes advantage of the spatial variation of Cu flux across the 3 inch diameter substrate holder in our MBE machine. Using a partially filled effusion cell for the Cu source, the Cu thickness long sapphire strip was -200% and quasi-linear.

variation across a 50 mm This is shown in Fig. 21.

On the other hand, the Pt and Co sources were broad area electron

gun

sources which produced essentially uniform thickness F’t and Co films. Following growth, the samples were cut into 25 strips, each 2 mm wide providing

a set in which only the Cu spacer thickness

varied systemati-

cally. Thus a single wedged multilayer provided 25 samples for MR data, all with the same thickness Pt and Co films grown under identical conditions.

858

Molecular

Beam Epitaxy

Cu Thickness on Wedge 600L,.,,1~,,'1,,,'1,,',"""'

t(A) = 362 + 8.60x + .0356x2

t t....,....l”‘,,‘,““,.,”

-030

20

-20

Wed;:"Position

(%I,

Figure 21. Thickness of a wedged Cu film on glass as measured by ellipsometry. Note that the variation of Cu layer thickness with position on the sample is quasilinear.

Figure 22(a) shows the saturation magnetoresistance values measured for four wedge samples used to span Cu thicknesses from 6 to 75 A. The multilayers were grown at a substrate temperature of 0°C and measurement temperatures were 4.2 K and 290 K. Note that the data overlap well from wedge to wedge indicating good reproducibility from deposition to deposition. The MR shows a single maximum, near 10 A. The 4.2 K measurements show the same maximum near 10 A and then above b - 16A, a monotonically increasing MR with increasing tcU. The increasing

divergence

of MR at the two temperatures

ascribed to the decreasing

effect of current shunting

may be partially

through

the Pt seed

film as the multilayer conductivity far exceeds that of the Pt seed film. The shunting effect is more noticeable at low measurement temperature since the conductivity

of the Pt seed film increases

greatly at low temperature.

Interestingly, when the growth temperature is increased to 150°C the peak MR increases by about 25% compared with the 100°C data. This is shown in Fig. 22(b). However, only a single MR peak is present. A further increase in growth temperature to 200°C produced multilayers with similar magnetic behavior to that shown in Fig. 22(b).

In order to explore possible

changes in multilayer growth mode with substrate temperature, which might be responsible for the observed change in magnetic behavior, we have carried out in situ XPD studies of the growth of Co on Cu(ll1) and vice versa. These are described in detail elsewhere.ts8] Here we summa-


rize the key result that for growth

Magnetic

at 0°C

Metal Structures

Co growth

on Cu results

659

in

physically rough films but with no chemical interdiffusion or segregation of Cu to the Co surface. For growth at 15O”C, segregation of Cu to the surface of Co occurs but there is a smoothing

of the Co/Cu interfaces.

growth of Cu on Co at 150°C significant occurred, possibly via the twin boundaries

For

diffusion of Co into the Cu in the film. This may be the

reason for ferromagnetic bridging of some regions of the multilayer. Thus growth-related defects, rather than an intrinsic absence of interlayer exchange coupling for [l1 II-oriented spacer layers, are the likely reason for the absence of oscillatory GMR in these multilayers. This view has recently been supported by our observation[8gl of oscillatory interlayer exchange coupling in MBE-grown permalloy (Ni,,FeJ/Au(l 11)multilayers.

I

40

growth

I

I

I

1

I

I

0°C

at

(a>-

30 -

T z 0

20-

.

0

V

.

T 4:2

K -

290 v

K _ v_

v I

10 -

G

%%;*LV0 0

I fY 2Q

I , growth

50040

30-

0

I at

I 04 -

0

4.2

B+=-

10 -

%

'qorbo

I 0

, I I I 150°C

= a,.

20 -

0

v

*V

,

I

-

*V*B

20 Cu

K

290 I 40

thickness

K I

, 60

I

(A)

Figure 22. Magnetoresistance at room temperature (open symbols) and 4.2 K (filled symbols) for ColCu samples deposited at 0°C (a) and 150°C (b). Data are displayed from a total of 6 wedge samples, with data from a single wedge plotted using the same symbol. The maximum applied field for each measurement was 60 kOe. In all cases, only a single MR peak is observed, near 10 A. For Cu thicknesses greater than 15 A, the variation of the MR is monotonic with Cu thickness for all samples and displays no further MR maxima.

660

Molecular

3.7

Giant Magnetoresistance Recently,

comprising

Beam Epitaxy

in 2-Phase Heterogeneous

GMR has been foundtgO)fgl] for non-multilayer

single films of a heterogeneous,

ferromagnetic

Alloy Films

and non-magnetic

metal.

phase-separated initially,

systems mixture of a

GMR was reported

for

polycrystalline films prepared by co-sputtering of Co and Cu, followed by sample annealing to drive phase separation. It is expected that in contrast to films prepared by sputtering, slow co-evaporation under UHV conditions and at moderate substrate temperatures would lead to spontaneous phase separation of Co and Cu, or indeed Co and Ag, since for temperatures below 400X, both of these systems are mutually insoluble (solubility < 0.01%). We have confirmedtg2) that crystalline, as-deposited films of these systems, grown epitaxially on sapphire (0001) and rocksalt (001) substrates, exhibit GMR. This is shown in Fig. 23(a) for Co-Cu alloy films and 23(b) for Co-AS alloy films. The saturation MR is plotted as a function of composition for films grown at a substrate temperature of 200°C onto a 30 8, thick Pt seed film (grown at 600°C) on sapphire (0001). The films were capped with a film of Pt 30 A thick. The dependence of MR on composition is fairly weak. The largest values of MR at 4.2 K are -50% and -70% for Co-Cu and Co-AS, respectively. Note that except for the two data points for Co-Ag at Co concentrations of 19 and 26%, the data are from compositional wedge samples prepared in a similar way to the wedged-multilayer samples described in Sec. 3.6. Preliminary results show that for epitaxial [OOl]-oriented Co-Cu films grown on cleaved rocksalt substrates, the saturation MR was significantly smaller than for the [ill] oriented films. However, this may reflect the higher defect density of these films as indicated

by x-ray measurements

of the mosaic

spread (=1.35” for the (002) peak compared with ~30.84“ for the (111) peak of films grown epitaxially

on Pt/sapphire).

Currently, phenomenological modelsfg1)[g3] of the GMR effect in these films suggest that the effect is due to spin-dependent scattering at the interface between the magnetic particles and the matrix.fg4] Recent experimentsfg41 have determined

the particle

size using x-ray scattering

techniques and compared the experimental GMR data with predictions based on the phenomenological models. This is a quite new field in which it may be possible to modify the film growth conditions to produce GMR in attractively low fields which could open up applications for these materials as magnetic field sensors.


60

I _ co-cu

I

Magnetic

a

I

4.2

4o

K n=

I

, .

===9.

Metal Structures

I (0) _

.

l *. n

. . .

20 -

6.

-290

ooooo

qn0

661

~~000~

-

00

K I

q 5

so0

I - Co-Ag

60 -

I

I

I

. 4.2 K

:

(b) -

‘e. . .

l

40 -

-00

.0_

0 20 -

290

K

0

I-

O) 0

40

10 Co (atomic

%)

Figure 23. Dependence of saturation AR/R (measured with field orthogonal to current path) on composition for (111) (a) Co-Cu and (b) Co-AS alloy films. The measurement temperatures are 4.2 K (closed symbols) and 290 K (open symbols).

4.0 -v CONCLUSIONS

We have described a variety of MBE-grown, artificially-layered magnetic metal multilayer structures, which exhibit non-bulk-like magnetic phenomena

that are strongly

dependent

on growth parameters.

Despite

the great excitement and interest in these materials, in many cases the physics underlying these phenomena is not fully understood. A key advantage of MBE over more conventional deposition techniques is the ability to define particular growth orientations of such structures and to probe, in situ, the physical and chemical nature of the interfaces. In addition, epitaxial structures are more easily and thoroughly characterized. To develop a complete picture of the magnetic metal interfaces will require synthesis of information from manyin situprobes including RHEED, LEED, AES, XPD, and even STM. Application of structures such as the

662


Co/Pt multilayers and the CoPt, alloy films as magneto-optical storage media have already been demonstrated and further developments will depend on engineering

the structures

to optimize their magnetic

ties. The current interest in understanding structure

in multilayers

proper-

the relation between GMR and

and alloy films may result in applications

such as

magnetic read heads, when these materials can be engineered to give large (~10%) changes in MR in fields of a few Oersteds. Overall, one may conclude that for an understanding of the new phenomena and their eventual application, a full picture of the magnetic interfaces should be arrived at by continuing

development

of a range of structural and magnetic probes.

ACKNOWLEDGMENTS We acknowledge

with thanks the assistance

of our colleagues:

C. H.

Lee, E. E. Marinero, Q. H. Lam and R. J. Savoy. The cross-section transmission electron micrographs were recorded by C. J. Chien (Ph.D. Thesis, Department of Materials Science and Engineering, Stanford University, 1991) and are reproduced with permission. This work was supported in part by ONR Contracts N00014-87-C-0339 and NOOOl4-92-C0084.

REFERENCES 1. (a) Mat. Res. Sot. Symp. Proc., (S. S. P. Parkin, H. Hopster, J. P. Renard, T. Shinjo, and W. Zinn., eds.), vol. 231, Mat. Res. Sot. (1992); (b) Mat. Res. Sot. Symp. Proc., (B. T. Jonker, S. A. Chambers, R. F. C. Farrow, C. Chappert, R. Clarke, W. J. M. de Jonge, T. Egami, P. Griinberg, K. M. Krishnan, E. E. Marinero, C. Rau, and S. Tsunashima, eds.), vol. 313 (1993) 2. “Magnetism and Structure in Systems of Reduced Dimension,” NATO AS/ Series B 309, (R. F. C. Farrow, B. Dieny, M. Donath, A. Fert, and B. D. Hermsmeier, eds.), Plenum Publishing Corporation, New York (1993) 3. (a) Zeper, W. P., Greidanus, F. J., Garcia, P. F., and Fincher, C. R., J. Appl. Phys., 654971 (1989); (b) Nakamura, K., Tsunashima, S., Iwata, S., and Uchiyama, S., IEEE Trans. Magn., MAG-25~3758 (1989); (c) Hashimoto, S., Matsuda, H., Ochiai, Y., Appl. Phys. Left., 56:1069, (1990)


Magnetic

Metal Structures

663

4. (a) Binasch, G., Grunberg, P., Saurenbach, F., Zinn, W., fhys. Rev., B39:4828 (1989); (b) Saurenbach, F., Barnas, J., Binasch, G., Vohl, M., Grunberg, P., and Zinn, W., Thin Solid Films, 175317 (1989) 5. Van Dau, F. N., Fert, A., Etienne, P., Baibich, M. N., Broto, J. M., Chazelas, J., Creuzet, G., Friederich, A., Hurdequint, H., Redoules, J. P., and Massies, J., Journale de Physique, 49:C8-1633 (1988) 6. Baibich, M. N., Broto, J. M., Fen, A., Nguyen Van Dau, F., Petroff, F., Etienne, P., Creuzet, G., Friederich, A., and Chazelas, J., Phys. Rev. Lett., 61:2472 (1988) 7. Parkin, S. S. P., Fan, S., More, N., and Roche, K. P., Phys. Rev. Left., 64:2304 (1990) 8. Parkin, S. S. P., Bhadra, 66:2152 (1991)

R., and Roche, K. P., Phys. Rev. Left.,

9. Parkin, S. S. P., Li, Z. G., and Smith, 58:2710 (1991)

D. J., Appl.

Phys. Left.,

10. Prinz, G. A., Paper presented Metallurgical Society, Anaheim,

at 119th Annual Meeting Ca. (Feb. 22, 1990)

11. Prinz, G. A., Science, 250:1092

(1990)

of The

12. Farrow, R. F. C., Harp, G. R., Marks, R. F., Rabadeau, T. A., Toney, M. F., Weller, D., and Parkin, S. S. P., J. Crystal Growth., 133:47 (1993) 13. Harp, G. R., Farrow, R. F. C., Marks, R. F., and Vazquez, Cryst. Growth, 127:627 (1993)

J. E., J.

14. Prinz, G. A., NATO AS/ Series, B163:311, Plenum Press, New York (1987) 15. (a) Sands, T., Palmstrom, C. J., Harbison, J. P., Keramidas, V. G., Tabatabaie, N., Cheeks, T. L., Ramesh, R., and Silberberg, Y., Materials Science Reports, 5:99 (1990); (b) Sands, T., Harbison, J. P., Palmstrom, C. J., Ramesh, R., and Keramidas, V. G., Mat. Res. Sot. Symp. Proc., 221:271 (1991) 16. Krishnan,

K. M., Appl. Phys. Left, 61:2365 (1992)

17. Farrow, R. F. C., Parkin, S. S. P., Speriosu, V. S., Bezinge, A., and Segmuller, A., Mat. Res. Sot. Syrnp. Proc., 151:203 (1989) 18. Farrow, R. F. C., Speriosu, V. S., Parkin, S. S. P., Chien, C. J., Bravman, J. C., Marks, R. F., Kirchner, P. D., Prinz, G. A., and Jonker, B. T., Mat. Res. Sot. Syrnp., 130:281 (1989)

664

Molecular

Beam Epitaxy

19. Lee, C. H., Farrow, R. F. C., Lin, C. J., Marinero, E. E., and Chien, C. J., Phys. Rev., B42:11384 (Rapid Comm.) (1990) 20. Lee, C. H., Farrow, R. F. C., Hermsmeier, B. D., Marks, R. F., Bennett, W. R., Lin, C. J., Marinero, E. E., Kirschner, P. D., and Chien, C. J., J. Magnetism Mag. Mat., 93592 (1991) 21. Lee, C. H., Farrow, R. F. C., He, H., Lamelas, F. J., Vavra, W., and Clarke, R., Mat. Res. Sot. Symp. Proc., 221:243 (1991) 22. Chien, C. J., Farrow, R. F. C., Lee, C. H., Lin, C. J., and Marinero, E., J. Magnetism Mag. Mat., 93:47 (1991) 23.

(a) Lee, C. H., He, H., Lamelas, F. J., Vavra, W., Uher, C, and Clarke, R., Phys. Rev., B42:1066 (1990); (b) Schuler, E., Elagoz, S., Vavra, W., Lamelas, F. J., He, H. D., and Clarke, R., Mat. Res. Sot. Symp. Proc., 231:235 (1992)

24. He, H., Lee, C. H., Lamelas, F. J., Vavra, Clarke, R., J. Appl. Phys., 675412 (1990) 25.

E.

W., Barlett,

D., and

(a) Vavra, W., Elagoz, S., Clarke, R., and Uher, C, Mat. Res. Sot. Symp. Proc., 231:385 (1992) (&I] Elagoz, S., Vavra, W., Uher, C, and Clarke, R., Mat. Res. Sot. Symp. Proc., 231:391 (1992)

26. Prinz, G. A., Phys. Rev. Left., 54:1051 (1985) 27. Idzerda, Y. U., Elam, W. T., Jonker, B. T., and Prinz, G. A., Phys. Rev. Left., 62:2480 (1989) 28. Marks, R. F., Farrow, R. F. C., Parkin, S. S. P., Lee, C. H., Hermsmeier, B. D., Chien, C. J., and Hagstrom, S. B., Mat. Res. Sot. Symp., 221 :15 (1991) 29. Greig, D., Hall, M. J., Hammond, C., Hickey, B. J., Ho, H. P., Howson, M. A., Walker, M. J., Wiser, N., and Wright, D. G., J. Magnetism. Mag. Mat., 11 O:L239 (1992) 30.

(a) Clarke, R., data presented at NATO Advanced Research Workshop on “Magnetism and Structure in Systems of Reduced Dimension”, Cargese, France, Unpublished (June 15-20, 1992); (b) Barlett, D., Tsui, F., Lauhon, L., Mandrekar, T., Uher, C., and Clarke, R., Mat. Res. Soc.Symp. Proc., 313:35 (1993)

31. Guo, J., Ellis, D. E., and Lam, D. J., Phys. Rev., B45:13647

(1992)

32. Guo, J., Chang, H. L. M., and Lam, D. J., Appl. Phys. Leti., 61:3116 (1992) 33. Toney, M. F., Wiesler, D. G., Hermsmeier, B. D., and Farrow, R. F., Mat. Res. Sot. Symp. Proc., 231:119 (1992)

Artificially-Layered Magnetic Metal Structures

665

34. Farrow, R. F. C., Toney, M. F., Hermsmeier, B. D., Parkin, S. S. P., and Wiesler, D. G., J. Appl. Phys., 70:4465 (1991) 35. Specht, E. D., Mak, A., Peters, C., Sutton, M., Birgeneau, D’Amico, K. L., Moncton, Phys., B69:347 (1987)

R. J., D. E., Nagler, S. E., and Horn, P. M., Z.

36. Kwo, J. R., Nate AS/ Series, B:163:355-357, York (1987)

Plenum

Press, New

37. Hong, M., Fleming, R. M., Kwo, J. R., Schneemeyer, L. F., Waszczak, J. V., Mannaerts, J. P., Majkrzak, C. F., Gibbs, D., and Bohr, J., J. Appl. Phys., 61:4052 (1987) 38.

(a) Erwin, R. W., Rhyne, J. J., Salamon, M. B., Borchers, J., Sinha, S., Du, R., Cunningham, J. E., and Flynn, C. P., Phys. Rev., B35:6808 (1987); (b) Rhyne, J. J., Erwin, R. W., Borchers, J., Salamon, M. B., Du, R., and Flynn, C. P., Physica, B159:lll (1989)

39. Beach, R. S., Borchers, J. A., Erwin, R. W., Flynn, C. P., Metheny, A., Rhyne, J. J., Salamon, M. B., J. Magnetism Mag. Mat., 104 107:1915 (1992) 40. Beach, R. S., Matheny, A., Salamon, M. B., Flynn, C. P., Borchers, J. A., Erwin, R. W., and Rhyne, J. J., J. Appl. Phys., 73:6901 (1993) 41.

(a) Koon, N. C., Jonker, B. T., Volkening, F. A., Krebs, J. J., and Prinz, G. A., Phys. Rev. Leff., 59:2463 (1987); (b) Krebs, J. J., Jonker, B. T., and Prinz, G. A., J. Appl. Phys., 63:3467 (1988); (c) Volkening, F. A., Jonker, B. T., Krebs, J. J., Koon, N. C., and Prinz, G. A., J. Appl. Phys., 63:3869 (1988)

42. Dutcher, J. R., Cochran, J. F., Heinrich, B., and Arrott, A. S., J. Appl. Phys., 64:6095 (1988) 43. Farrow, R. F. C., J. Appl. Phys., 64:5315 (1988) 44. Unguris, J., Celotta, R. J., and Pierce, D. T., NATO AS/ Series, B309:101, Plenum Publishing Corporation, New York (1993) 45. Unguris, J., Celotta, R. J., and Pierce, D. T., Phys. Rev. Left., 67, 140 (1991) 46. Etienne, P., Lequien, S., Nguyen-Van-Dau, F., Cabanel, R., Creuzet, G., Friederich, A., Massies, J., Fen, A., Barthelemy, A., and Petroff, F., J. Appl. Phys., 67:5400 (1990) 47. Demokritov, S., Wolf, J. A., Grunberg, Sot. Symp. Proc., 231 :133 (1992)

P., and Zinn, W., Mat. Res.

48. Grunberg, P. A., Fuss, A., Leng, Q., Schreiber, R., and Wolf, J. A., NATO Series, B309:87, Plenum Publishing Corporation, New York (1993)

666

49.

Molecular

Beam Epitaxy

(a) Bruno. P. and Chappert, C., Phys. Rev. Leff., 67:1682 (1992); (b) Bruno. P. and Chappert, C., Phys. Rev. Lett., 67E:2592 (1992)

50. Herman, F., Sticht, J., and Van Schilfgaarde, Symp. Proc., 231:195 (1992)

M., Mat. ResSoc.

51. Garcia, P. F., Appl. Phys. Left., 47:178 (1985) 52. Garcia, P. F., J. Appl. Phys., 63:5066 (1988) 53. Sato, N., J.App/. Phys., 64:6424 (1988) 54. Ochiai, Y., Hashimoto, 25:3755 (1989)

S., and Aso, K., IEEE Trans. Mag. Mat.,

55. Den Broeder, F. J. A., Donkersloot, H. C., Draisma, H. J. G., and De Jonge, W. J. M., J. Appl. Phys., 61:4317 (1987) 56. Draisma, H. J. G., Den Broeder, F. J. A., and De Jonge, W. J. M., J. Appl. Phys., 63:3479 (1988) 57. Draisma, H. J. G., De Jonge, W. J. M., and Den Broeder, F. J. A., J. Magn. Magn. Mat., 66:351 (1987) 58. Hermsmeier, B. D., Farrow, R. F. C., Lee, C. H., Marinero, Chien, C. J., J. Appl. Phys., 69:5646 (1991)

E. E., and

59. Yan, X., Egami, T., Marinero, E. E., Farrow, R. F. C., and Lee, C. H., J. Mat. Res., 7:1309 (1992) 60. Marinero, E. E., Farrow, R. F. C., Lee, C. H., Notatys, H., Yan, X., and Egami, T. A., Applied Physics Communications, II:359 (1992) 61. Huang, T. C., Advances

in X-ray Analysis,

35:143 (1992)

62. Chien, C. J., Clemens, B. M., Hagstrom, S. B., Farrow, R. F. C., Lee, C. H., Marinero, E. E., and Lin, C. J., Mat. Res. Sot. Symp. Proc., 231:465 (1992) 63. Cho, N. -H., Krishnan, Appl. Phys., (1992)

K. M., Lucas, C. A., and Farrow, R. F. C., J.

64. Cho, N. -H., Krishnan, K. M., Lee, C. H., and Farrow, R. F. C., Appl. Phys. Left., 60:2371 (1992) 65. Zhang, B., Krishnan, 51:298 (1993)

K. M., and Farrow, R. F. C., Ultramicroscopy,

66. Zhang, B., Krishnan, K. M., Lee, C. H., and Farrow, R. F. C., J. Appl. Phys., 73, 6198 (1992) 67. Toney, M. F., Farrow, R. F. C., Marks, R. F., Harp, G., and Rabedeau, T. A., Mat. Res. Sot. Symp. Proc., 263:237 (1993) 68. Farrow, R. F. C., Lee, C. H., Marks, R. F., Harp, G. R., Toney, M. F., Rabedeau, T. A., Weller, D., and Brindle, H., NATO Series, B309:215, Plenum Publishing Corporation, New York (1993)


Magnetic

Metal Structures

667

69. Rabedeau, T. A., Harp, G. R., Marks, R. F., Farrow, R. F. C., and Toney, M. F., paper presented at Spring MRS Meeting, San Francisco (Apr. 12 -16, 1993), Unpublished 70.

(a) “BinaryA//oy Phase Diagrams”, (T. Massalski, ed.), 2nd Ed., Vol. 2., Mat. Info. Sot., Materials Park, Ohio (1990); (b) Sanchez, J. M., Moran-Lopez, J. L., Leroux, C., T.Cadeville, M. C., J. Phys. C: Solid State Phys., 21 :L1091 (1988)

71. Farrow, R. F. C., Hermsmeier, B. D., Lee, C. H., Marks, R. F., Marinero, E. E., Lin, C. J., Chien, C. J., and Hagstrom, S. B., Mat. Res. Sot. Symp. Pfoc., 229:115 (1991) 72.

(a) Toney, M. F., Rabedeau, T. A., Farrow, R. F. C., Marks, R. F., and Harp, G. R., in preparation, (6) Huang, T. C., Savoy, R., Farrow, R. F. C., Marks, R. F., Appl. Phys. Lett. 62:1353 (1993)

73. Hansen, M., “Constitution New York (1958) 74.

of Binary Alloys”,

2:493,

McGraw

Hill,

(a) Weller, D., et al., to be published: (IJ) Marinero, E. E., Farrow, R. F. C., Harp, G. R., Geiss, R. H., Bain, J. A., and Clemens, B., Mat. Res. Sot. Symp. Proc. 313:677 (1993)

75. Weller, D., Braendle, H., Farrow, R. F. C., Marks, R. F., and Harp, G. R., NATO Series, B309:201, Plenum Publishing Corporation, New York (1993) 76. Weller, D., Braendle, H., Gorman, Appl. Phys. Letf., 61:2726 (1992) 77.

G., Lin, C. -J., and Notarys,

H.,

(a) Farrow, R. F. C., Geiss, R. H., Gorman, G. L., Harp, G. R., Marks, R. F., and Marinero, E. E.,J. MagneticSoc. ofJpn., 17(S1):14O (1993); (b) Farrow, R. F. C., Weller, D., Toney, M. F., Rabedeau, T. A., Hurst, J. E., Harp, G. R., Marks, R. F., Geiss, R. H., Notarys, H., Mat. Res. Sot. Symp. Proc., 343:375 (1994)

78. Camley, R. E. and Barnas, J., Phys. Rev. Leff., 63:664 (1989) 79. Parkin, S. S. P. and York, B. R., Applied. Phys. Left., 62:1842 (1993) 80.

(a) Bruno. P. and Chappert, C., Phys. Rev., B46:261 (1992); (b) Bruno. P. and Chappert, C., NATO Series, B309:389, Plenum Publishing Corp., New York (1993); (c) Bruno. P., J. Magnetism Magn. Mat., 121:248 (1993)

81. Herman, F., Sticht, J., and Van Schilfgaarde, Symp. Proc., 231 :195 (1992)

M., Mat. Res. Sot.

82. Egelhoff Jr., W. F. and Kief, M. T., Phys. Rev., B45:7795

(1992)

888

Molecular

Beam Epitaxy

83. Schreyer, A., Metoki, N., Zeidler, T., Bddeker, P., Abromeit, A., Morawe, Ch., Romahn, U., Sonntag, P., Brohl, K., and Zabel, H., Preprint, (1992) unpublished (note that this group has subsequently reported evidence for oscillatory exchange coupling in Co/Cu(222) superlattices); Schreyer, A., Brohl, K., Ankner, J. F., Majkrzak, C. F., Zeidler, T., Bodeker, P., Metoki, N., and Zabel, H., Phys. Rev., B47:15334 (1993) 84. Cebollada, A., Martinez, J. L., Gallego, J. M., De Miguel, J. J., Miranda, R., Ferrer, S., Batallan, F., Fillion, G., and Rebouillat, J. P., Phys. Rev., B39:9726 (1989) 85. Cebollada, A., Miranda, R., Schneider, C. M., Schuster, Kirschner, J., J. Msg. A&g. Mat., 102:25 (1991)

P., and

86. Johnson, M. T., Purcell, S. T., McGee, N. W. E., Coehoorn, R., aan de Stegge, J., and Hoving, W., Phys. Rev. Leti., 68:2688 (1992) 87. Coehoorn, R., Johnson, M. T., Folkerts, W., Purcell, S. T., McGee, N. W. E., De Veirman, A., and Bloemen, P. J. H., NATO AS/ Series, B309:295, Plenum Publishing Corporation, New York (1993) 88. Harp, G. R., Parkin, S. S. P., Farrow, R. F. C., Marks, R. F., Toney, M. F., Lam, Cl. H., Rabedeau, T. A., and Savoy, R. J., Phys. Rev., B47:8721 (1993) 89. Parkin, S. S. P., Farrow, R. F. C., Marks, R. F., Cebollada, G. R., Savoy, R. J., Phys. Rev. Lett., 72:3718 (1994)

A., Harp,

90. Berkowitz, A. E., Mitchell, J. R., Toney, M. J., Young, A. P., Zhang, S., Spada, F. E., Parker, F. T., Hutten, A., and Thomas, G., Phys. Rev. Lett., 68:3745 (1992) 91. Xiao, J. Q., Jiang, J. S., and Chien, C. L., Phys. Rev. Lett., 68:3749 (1992) 92. Parkin, S. S. P., Farrow, R. F. C., Rabedeau, Harp, G. R., Lam, Q., Chappert, C., Toney, Geiss, R., Europhysics Lett., 22:455 (1993)

T. A., Marks, R. F., M., Savoy, R., and

93.

(a) Zhang, S., Appl. Phys. Lett., 1855 (1992); Levy, P. M., J. Appl. Phys., 73:5315 (1993)

(b) Zhang,

S. and

94.

(a) Rabedeau, T. A., Toney, M. F., Marks, R. F., Pat-kin, S. S. P., Farrow, R. F. C., and Harp, G. R., Phys. Rev., B, 48:16810, (1993); @) Marks, R. F., Farrow, R. F. C., Harp, G. R., Parkin, S. S. P, Rabedeau, T. A., Toney, M. F., Cebollada, A., Thangaraj, N., Krishnan, K. M., Mat. Res. Sot. Symp. Proc., 313:411 (1993)

8 Reflection High Energy Electron Diffraction Studies of the Dynamics of Molecular Beam Epitaxy Philip 1. Cohen, Gale S. Pefrich, and Gregory J. Whaley

1 .O

INTRODUCTION

The success of molecular beam epitaxy is largely due to its compatibility with in-situ surface characterization. Most other crystal growth environments require ambient pressures or fields too hostile to apply the electron probes developed for the study of surfaces. The vacuum requirements and geometry of MBE, however, conveniently permit the use of reflection high-energy electron diffraction (RHEED). Only an electron gun and a phosphor-covered viewport are needed for these measurements. With this simple apparatus, it is easy to obtain important information about the character of the surface and the nature of the growth mode. From the symmetry and separations of the diffracted beams, one can determine the surface

reconstruction

and lattice constant.

From the disappearance

of

reconstructions, one can determine flux ratios. From phase transitions, From the qualitative appearance of the one can calibrate temperature. pattern, cluster growth or layer growth can be distinguished. But far more information can be obtained by careful analysis of the diffraction. A difficulty is that, at this next level of interpretation, the patterns are quite complicated and not completely understood. The purpose here is to describe current progress in the use of quantitative electron diffraction measurements to understand the dynamics of the microscopic processes of epitaxial growth.

669

670

Molecular

RHEED absolute

Beam Epitaxy

is primarily

diffracted

sensitive

intensities

atoms.

The relative

intensities

depend

more

surface

distribution

upon

of atomic steps.

to surface

depend

atomic

on the positions

structure.

and the shape of the diffracted morphology,

The

of the surface

i.e., the surface

beams

order

The first of these is just beginning

and to be

attackedt2j-tsjf2’j with success. It is based on the first dynamical calculations appropriate to RHEED.t6j It is a hard problem to which there is little data that theory can compare. The second is the main focus of this chapter. Diffraction suffers from the disadvantage that it is not as direct as microscopy. Other techniques, such as reflection electron microscopy,~jt6j scanning tunneling microscopytgj and low-energy electron microscopyftOj are able to image the same steps that can cause streaks in a RHEED pattern. Diffraction has the distinct advantage that it measures statistics To understand epitaxial and does not follow an individual step. growth, the latter is required, though both would be exceedingly useful. Unfortunately, heroic measures are required to apply these high resolution imaging techniques to the growth of a complicated film at elevated temperatures. Since the main application of RHEED has been to the growth of semiconductors, this dominates our discussion. Two kinds of surfaces will be distinguished: low-index or singular surfaces, and vicinal surfaces. In the next two sections, the fundamentals two types of surfaces are summarized.

of electron diffraction from these The special diffraction features

particular to RHEED are explained. In the fourth section, measurements during the growth of GaAs and AlGaAs are discussed in terms of these fundamentals. to compare

To avoid artifacts due to multiple scattering, to trends

at several

incident

angles.

it are essential

In the fifth section,

RHEED measurements on a lattice mismatched system are reviewed. Finally these measurements are discussed in terms of simple mathematical models of epitaxy. Our main theme is that at the growth rates important to molecular beam epitaxy, all of these processes demand consideration

2.0

of the role of both 2D cluster formation and step propagation.

DIFFRACTION

GEOMETRY

In a typical MBE, application a 1O-20 keV beam strikes a sample at a glancing angle of less than 5” so as not to block the Knudsen cells or gas

RHEED Studies of the Dynamics of MBE

671

sources used to provide the incident fluxes. At these high energies, stray magnetic fields from the Knudsen cells and sample heater have little effect on the electron trajectories.

Figure 1 shows the scattering

geometry.

incident and final glancing angles, Qi, Fiji and the incident azimuthal q+, are indicated.

Three electron trajectories

The angle,

are also shown, the specular

beam, the part of the incident beam that misses the sample, and the shadow edge that corresponds to the locus of low-energy secondaries that can just leave the surface. We use the convention that all glancing angles, 6, are measured with respect to the low-index plane and not the macroscopic surface plane. This distinction will be important for diffraction from a vicinal surface. For now, it suffices to recognize that, by measuring the position of the primary and specular beams, the angle of incidence determined.

can be

RHEED gun IO keV u

k9

source

kl

Ovens

-E sample motion

phosphor E screen ‘xY stage

data

II Q 0 --

lens slit aperture photomultiplier

acquisition

Figure 1. Schematic of a diffraction apparatus used during MBE. The specular, shadow edge, and primary electron trajectories are shown. Angles are defined with respect to the low-index surface plane.

672


There has not been a systematic effort to determine the relative merits of increasing or reducing the electron energy. It might be that kinematic arguments are more successful at lower energy. Higher energy is better suited to systems with e-beam deposition. Energies as low as 3 keV can more easily be combined with simultaneous Auger spectrosc~py.t~~l At this point, all we can say is that reducing the energy below about 5 keV degrades the operation of our electron gun. Going to higher energy spatially compresses the pattern, as the square root of the energy. The resulting

diffraction

pattern is shown in Fig. 2 for a GaAs(lOO)

surface that has been annealed in an As, flux. This is the characteristic pattern from a well-ordered surface and consists of a set of sharp beams arrayed along a circle of radius k sin eI The primary beam is also evident, though its intensity has been reduced through partial occlusion by a shutter. This is the fourfold pattern of the 2 x 4 GaAs(lO0) reconstruction. Every fourth beam would be present even without the reconstruction and are termed integral order beams. The additional beams due to the super period are termed fractional orderbeams. Because reconstructions do not always cover entire terraces formed by steps, there can be relative displacements between reconstructed domains and hence scattering phase differences between the reconstructions over the surface. direct consequence is that the intensities, rather than amplitudes,

The dif-

fracted from the domains must be added incoherently. Thus the behavior of the beam shape and the time development of the intensities of the fractional order beams and integral order beams can be very different.n2] In this discussion, only the behavior of the integral order beams is described. If growth is initiated by opening the Ga shutter, then the sharp diffraction

pattern of Fig. 2 becomes

sharp beams are now elongated slight changes evident

the pattern shown

perpendicular

in their width parallel to the surface.

if the intensity

is measured.

in Fig. 3.

The

to the surface with only

The intensity

The latter is more at the peak positions

also changes. In fact, the intensity of the specular and other integral order beams often oscillates in time with a period corresponding to the growth of a monolayer of GaA~.f~~jf~~j Th ese changes in shape and intensity are the main subject of this discussion. Other broadening mechanisms are also observed. For example, in the twofold GaAs(l00) pattern, antiphase mistakes elongate the half order streakst15jt16j and, with some multiple scattering,

the integral order streaks.

RHEED Studies of the Dynamics

of MBE

673

Figure 2. Photograph of the four-fold diffraction pattern of a well annealed 2 x 4GaAs(lO0) surface. The sample is at growth temperature (SOO’C) in an As, flux.

Figure 3. After growth is initiated on GaAs(lOO), the sharp beams of the four-fold pattern of the 2 x 4 reconstruction elongate into streaks perpendicular to the surface. The elongation is due to the diffraction from random atomic steps. Though not perfect, the surface is still relatively smooth.

674

Molecular

Beam Epitaxy

There are several choices of detecting schemes. the intensity

of the diffraction

The distribution

of the intensity

pattern is measured versus scattering

angle is determined

scanning with a magnetic deflection system. A schematic is shown

In our laboratory,

with a photomultiplier. by

of the apparatus

in Fig. 1, After striking the sample, the 10 keV beam forms a

diffraction pattern on the IT0 coated phosphor (JEOL P15) covered screen. A photomultiplier lens assembly is positioned with micrometer stage.t14) Then the intensity in the desired region is measured by deflecting the pattern with an external magnetic field. Note that the detector in this mode is fixed. The output is then measured with a fast preamplifier and sampled with a 12 bit analog-to-digital converter. This system has several advantages. First, it is fast with high spatial resolution and wide dynamic range. Second, it can be positioned with a minimal software interface. Most importantly, after positioning the detector over a given phosphor grain, it is insensitive to non-uniformities of the phosphor screen. This is especially necessary in an MBE environment where the electron beam catalytically decomposes the As, background gas onto the phosphor. The disadvantage of this arrangement is that simultaneously following the intensity in different portions of the pattern at once is cumbersome, though two detectors are still quite practical.t17)t1*) For the measurements

discussed

here, the scanning

system

is

mainly used to measure the diffracted intensity along the length of the streak, perpendicular to the sample surface. To measure the intensity across the width of the streak, parallel to the surface, or to measure the separation of two streaks (for lattice constant determination), a different set of scanning coils must be used. Alternative video systems are, of course, practical. Larsen has measured the intensity along the length of the streakst1Ql-t21) and Grunthaner has made intensity measurements.t**] Unfortunately, the spatial resolution of CCD or SIT video systems is limited to between 500 and 1000 pixels and are relatively

slow because of

the enormous amount of data that must be transferred. Image dissectors with custom software should be faster, not spending the time required to measure unwanted

portions of the pattern.

For some measurements,

useful to combine the scanning arrangement with the video system. general, the video system is somewhat more convenient.

it is In


675

3.0

DIFFRACTION FUNDAMENTALS

3.1

Kinematic Approximation

theory

Care must be taken when applying kinematic or single scattering to the analysis of electron diffraction from surfaces. Strictly

speaking, kinematic theory is applicable if the inelastic mean free path is less than the elastic mean free path. It is, then, unlikely that any elastically scattered electron has scattered more than once. Unfortunately, this condition does not hold in RHEED. In the RHEED geometry from surfaces, the path traversed by the electron beam is comparable in length to the inelastic mean free path so that the incident and exiting electrons can undergo multiple elastic scattering. A calculation of the absolute diffracted intensity must include all such processes. It is a dynamic calculation in the sense that it requires consideration of the scattering potential. From this calculation one expects to be able to extract the atomic positions in a reconstructed surface.t21t61 A calculation beams, however, need only consider positions are determined by momentum kinematics of the scattering process. If

of the positions of the diffracted the surface symmetry. These and energy conservation, i.e., the this symmetry is slightly reduced

due to disorder or due to the finite lateral dimensions of the surface, then the diffracted beams will broaden. The essence of our argument is that the broadening of the diffracted beams is determined by the relaxation of momentum conservation and, to first order, is given by kinematic calculation.PW41 This approximation is admittedly limited, as discussed below, and somet27l even argue inappropriate. It should only be appropriate when two-dimensional

islands

on a surface

are large enough

that coherent

scattering

between islands is not important and when island edges do not

represent

a large fraction

of the scattering.

Other approximations

have

been attempted, such as relating the diffraction to the step density on the surface during epitaxy.t28] This approximation must also be incomplete. Though it gives some agreement to Monte Carlo calculations at a fixed incident angle, it is difficult to believe that very small clusters scatter the same as a large cluster or that the angle of incidence dependence is so simple. The discussion of the kinematic treatment below is offered as a reference point to illustrate what a dynamic treatment must ultimately deal with and because it has been found to give agreement with experiment in some cases. However, it is crucial that its predicted angular dependence be checked.

The discussion

is probably

more correct for low energy

676


electron diffraction, where the normal incidence geometry makes step edge scattering a smaller fraction of the total and coherent scattering between islands less important.

The treatment

of disorder is probably the

most important impediment to the application of diffraction to understanding epitaxial growth. It is being pursued by a number of workerst6)t2g)t30) and progress is being made. Our expectation is that some modification of the kinematic result presented here will be practicable. The application of the kinematic approximation is equivalent to the column approximation of electron microscopy. The important points are that (7) several layers near the surface are important to the diffraction, so that shadowing of lower layers by step edges on an upper layer is ignored, (2) the full dynamical scattering under a finite region of a surface is included, and (3) though the subsurface atoms are included, only the toplayer coverages appear in the result. The full application of dynamic theory to calculate the diffraction from a disordered surface is not yet possible, though a start has been made.p)f31) To apply the column approximation of electron microscopy, consider a surface with random steps with edges at Rr = dii + Lr. If each terrace is large enough, the diffracted amplitude is the kinematic sum of the full dynamically scattered amplitude from each block, A+ i.e.,

Eq-(1)

A(S,

@,q~)= 2

A#,

6,

q~)@*~’

i

This neglects multiple scattering between the blocks.

Here @ and cpare the set of appropriate polar and azimuthal scattering angles. As the blocks become small, it is not expected to work well. For example, if one were to

try to calculate the diffraction from a periodically stepped surface with a repeat that was several unit-cells, then this should not be a good approximation. In that case, one needs to do a full dynamic calculationt31) since the surface is more like a reconstruction than a stepped surface. Our second assumptionf23) is that the most rapid variation of Ai with S is due to the finite size of the block. This is not always the case, especially when taking

data near conditions

rapidly with

3,.

of high symmetry

or if the intensity

In the latter case, some correction

varies

can be made by

including the measured S, dependence of the diffracted intensity from a more perfect surface.t32) With this assumption on the variation OfA;, we let f@,q) be the dynamic amplitude from a perfect surface and redefine A;(S) to only contain the kinematic size effects of a block, including the number of scatterers.

Then in this approximation,

the diffracted

intensity

is


A(S, 6, v) = f(S, 9)

Eq. C-3

of MBE

677

x A;(S)e’S’R’ 1

with the recognition

that one should keep the scattering

angles 6 ,cpfixed.

Because of this last point, we will, for the most pat-t, set f = 1. Kinematic shape effects are most easily understood by performing the sum indicated in Eq. (2) for a two-dimensional rectangular net of dimensionsN,a, andI&. This is a special case in which a single block is considered to be a plane of scatterers. (It is important to realize that, more generally, by including a three dimensional block in Eq. (2), one is able to include single and multiple scattering within a given column.) But with only one plane, there is no kinematic dependence on S, and the famous result[33] is that for one block

In the .S, direction, the FWHM of this function is approximately 2x/N,a,. The broadening is similar for theS,, direction. Figure 4 illustrates two cases in diffraction experiments in which the incident wavevector, ki, is fixed and the final wavevector, k/is uncertain due to this broadening. In Fig. 4a, the incident electron is in the; direction, along q~= 0, making a glancing angle D; to the surface plane. For near forward scattering as shown, the range of final angles, 6q1~that correspond to the uncertainty in momentum transfer satisfies k6cpf = 2xlL, where L, = N2a2. It is important to note that this broadening will not be observable if the range of angles in the incident beam is greater than this. Alternatively, if 8~ is the range of angles in the incident beam, then order over a distance less than about 2xlki5rp will broaden the beam measurably.

This distance is the resolvable

transfer

width of the diffractometer

in the direction

incident

beam.

the same argument

Figure 4b illustrates

parallel to the direction of the incident beam.

distance or

perpendicular

to the

for the direction

In this direction,

because of

the small glancing angle, Q the diffraction is much more sensitive to disorder on the surface. Here the range of final glancing angles, 6ej, that correspond to the uncertainty in momentum transfer is given by k 66sin6 = 2x/L,, where L, = Np,. If the range of angles in the incident beam is greater than X$, then this will not be observed. Hence order over distances less than 2z/k sinD@6, where 66 is the range of angles in the incident beam, will broaden the diffracted beam. At an angle of incidence

676

Molecular

Beam Epitaxy

of pi = @I = 2”, disorder over the crystal will broaden the beam in the i direction about 30 times that of the j broadening. This is the fundamental cause of streaks in RHEED. In subsequent sections, the impact of a particular

class of disorder, atomic steps, are emphasized.

, intersectionwith / Ewald sphere

\ \ 3.

\

a)

0

I

]

b)

0

Figure 4. An Ewald construction showing the conservation of energy and momentum for the case in which the sharp reciprocal lattice lines of the 2D surface are broadened into rods by some disorder. (a) A given amount of disorder greatly broadens the beam in fif (b) The same amount of disorder has less effect on the width of a streak, i.e., cp. The asymmetry in the instrumental sensitivity to disorder is purely a result of the low glancing angle of incidence, pi.

The transfer

width

or distance

over which

order is resolvable

is

sometimes referred to as (certainly incorrectly) coherence length. The important point is that the length is asymmetric. For our instrument, distance over about 10,000 a in the direction of the incident beam can be resolved. The main impact is that when studying a misoriented surface, i.e., one in which there is a staircase of steps, it is possible to be in a situation that, when the beam is pointing down the staircase, more than one terrace contributes coherently to the diffraction (add amplitudes). But


of MBE

679

when the beam is parallel to the step edges on the same surface, the steps (if straight and parallel) could be so large that one must add intensities diffracted 3.2

from each.

Disorder

on Low-Index

Surfaces

During perfect layer-by-layer growth, single-layer islands first form on an otherwise featureless plane, and then fill in until the layer is smooth and complete. During this process, the two-dimensional islands that form could have a range of sizes and shapes. In the diffraction pattern, there will be interference because of the difference in the path length traversed by electrons scattering off the different levels of the surface. Depending upon the scattering geometry and the distribution of islands, the diffracted beams will broaden, giving rise to the long streaks typically observed during growth. From the shape of the beams, one hopes to determine the distribution of island sizes. For fixed glancing and azimuthal angles of incidence, fii and vi, the intensity along a streak can be measured versus the final glancing angle ef. For the specular streak, for example, scattering angle can be related to momentum transfer via S, = k

Eo. (4)

COS ~~ -

k

COS pi

With this connection, the intensity along the specular streak is expected to have a simple general form if only a few layers contribute to the diffraction. For example, suppose there are islands on top of a flat surface, with the coverage of the islands 8. Then there will be interference between the top layer with coverage This two-level

8 and the layer below with exposed

case (see the Appendix),

tion of island sizes, the diffracted Eq. (5)

I&S*)

coverage

1 - 0.

shows that, for a Markov distribu-

intensity

is given by

= [@ + (1 - cl)2 + 28( 1 - 0) cos S,d] 2n 8(S,) + 20(1 - 0) (1 - cosS*d) [2h/(P+

S,z)]

This says that, for a two-level system, the diffracted intensity along the streak can be separated into a broad part and a central spike. The FWHM of the broad part is 2/h or twice the reciprocal of the sum of the average hole and island size.f34] This should be compared to the width of about 2nlL

680


given by Eq. (3). One can think of S, as roughly determined from $and .S, determined from f+ - tii, but these must be calculated more exactly using Eq. (4). The two terms of Eq. (5) are: a central spike that results from the long range order, and a broad part that is due to the disorder.

The first term

is written as a delta function, but there is a range of angles in the instrument and so this term is broadened. As shown in Eq. (3), the width determines the size of the coherently scattering region. The second term depends upon the form of the distribution of island sizes. A form more general than the simple exponential distribution given here can be derivedf32) but this will suffice. Further, more than two levels can be considered by including additional terms in Eq. (5). To observe this distribution of intensity, we deposited a submonolayer coverage of Ge onto a Ge(ll1) sur-face.f35) The diffusion of Ge is low so that this system gives an ideal compromise between continuous epitaxy Figure 5 shows the results and the time required to make a measurement. of experiment and a calculation based on Eq. (5). The left panel shows a plot of the intensity versus tif along the specular streak for several incident angles Q. The right panel is a fit using Eq. (5) after converting to angle with Eq. (4). In both there is a central spike and a broad component. The spike dominates at in-phase conditions in which .S,d is an integer times 27~ and the broad part dominates at out-of-phase conditions half way in between. These conditions are crucial to the formation of streaks in RHEED. At in-phase conditions, the extra scattering path length between different terraces is an integral number of electron wavelengths. The diffraction is insensitive to the step disorder and only the central spike, the long range order, is seen. At out-of-phase conditions, the diffraction is maximally sensitive to steps. If the coverage were one half layer, the central spike would vanish. However, and this is an important point, because of the second term in Eq. (5)) the total intensity does not go zero. To our minds, this shape is the dominant feature of the diffraction. During layer-by-layer epitaxy, the diffracted intensity is traded between these two components

as growth proceeds.

How each component

contributes to the

measured peak intensity determines the RHEED intensity oscillations to be discussed in Sec. 4.1. The right panel is a one parameter fit to the data. Note that the same asymmetries in the data are present in the calculation and are primarily due to the cut that the Ewald sphere makes through the reciprocal lattice rod. A better fit could be obtained by using a third level in the calculation. These measurements give us a great deal of confidence in the kinematic analysis of beam shape and the interpretations used in the remainder of this chapter.


of MBE

661

IA 1 II covemge=0.4 -I

a*/3=O.Olh 7 =aoolr

0; =

36mr

A

20

IO

0

E L

A

IO

A@, (mrad)

L /

20

IO

17mr

0

A0,

lo-

20

(mrad)

Figure 5. The left panel shows scans of the intensity of the specular streak of a smooth Ge(l11) surface with 0.4 monolayers of Ge deposited. The intensity is measured versus fif at several different glancing incident angles 8, It is compared to a calculationf35t in the right panel that shows that at in-phase Q, the beam is sharp, while at out-of-phase Bi, the beam exhibits a broad component under a sharp spike.

3.3

Vicinal Surfaces A wafer

that

is cut and

polished

from

a boule

is never perfectly

oriented parallel to a low-index plane. This misorientation results in a series of steps separated by low-index terraces. Typically these steps are one atomic layer in height, though double-layer and higher reported.t361-f3g1 D ue to the exceedingly long transfer width interference between waves scattered from different terraces strongest component to the shape of the diffracted beams. In the diffraction complication

from a regular staircase

of steps is considered.

have been of RHEED, is often the this section, Later, the

of disorder will be added. An important point is that, because

682


a staircase has a specific direction, the resulting modification in the shape of the diffracted beams will depend markedly on azimuthal angle of incidence

cp.

An ordered staircase will cause the specular components.

We calculate

dimensional staircase terrace have length L consists of blocks with arguments in Sec. 3.1,

this angle here.

A(S)

a regular,

one-

of monatomic steps in the ; direction. Let each and the steps have height d. Then the staircase corners at R,,, = di - mdi. Further, following the let the full dynamically scattered intensity from the

block under one terrace be f(S,S,v). the staircase is, following Eq. (2)

Eq-(6)

beam to split into two

Consider

Then the diffracted

amplitude

from

= 1 f,,,iiS’Rn m

where for generality the f’s might depend upon m. If the terraces are all identical, then with the beam in the i direction (along go= 0) this becomes

There is constructive interference when there is both energy conservation and S&. -S& = even n. For example, the specular beam will satisfy this at two final angles, er,, and fiff2 corresponding to two values of S, and two values of S,. Taking differences, the separation Atif = Dff.2 - tff,, is given by

Eq. (8)

ASxL - A&d = 2n

Further, for the specular

beam

Eq. (9)

so that

Eq. (10)

A& = - k sinOf AiZ$ and AS= = k AOf

Combining Eqs. 8 and 10, one has for the beam staircase misoriented by 17~= d/L

4. (11)

pointing

down

the


Thus given d and measuring

of MBE

683

Aej, one can determine the misorientation

of

the staircase, fi,_, with great precision. misorientation

Alternatively, one can measure the with 0 - 26 x-ray goniometry t401f41)and then determine the

step height, d, using RHEED. This can easily be extendedf4*] dence.

for other azimuthal

The main points to be mentioned

angles of inci-

are (7) in order to distinguish

between Kikuchi lines, the incident glancing and azimuthal angular dependence must sometimes be checked,f43) and (2) the separation ABf depends upon the disorder in the staircase as described in Sec. 3.4.t3*] Summarizing, the signature of a staircase in the diffraction pattern is a split diffracted beam. The amount of splitting depends upon scattering angles in a characteristic way. 3.4

Disorder

on Vicinal Surfaces

For a regular staircase, a staircase in which each terrace length is identical, the specular beam will be split into two components. The width of these components will be equal to the range of angles in the incident beam. The separation of the components will be determined by the misorientation as described in the previous section. If, however, there is disorder, then one needs a model for the distribution of terrace lengths to calculate the diffraction. The main result is that the width of the components is broadened and the separation of the components is decreased.t3*) As a specific example, assume that the probability of a terrace of length L >Lo is given by P(L) = a(? -a(L - w

Eq. (12)

and zero for L < Lo. uncorrelated.

The order of the terrace lengths

For this distribution

of terrace

lengths,

is assumed

to be

the mean terrace

length is given by L = Lo t l/a and the rms deviation from the mean is l/a. When the diffracted intensity is measured versus fif, both S, and S, are changed. Hence it is a bit simpler to express the intensity at constant S, when considering the width of the components and the separation of the components. With this in mind, the intensity for the shifted exponential distribution of Eq. (12) is obtained at S,d = x from Eq. (19) of Ref. 32 to be 4/L

Eq. (13)

I(s,,s, = x/d) =

2a*t 2a2 cos21~xt (2~c/L~)*x* - (47ca/Lo)x sin2rwc

684

Molecular

Beam Epitaxy

where x = SJcI2n.

This is illustrated

When a is not too large (l/a for this distribution

in Fig. 6 for several values

is small compared to I),

that the separation

of a.

one can also show

between the peaks AS, is AS, = 2x/

E. Using u = l/a as the rms deviation from the mean terrace length, the width of one component, normalized to the separation of the peaks, 6SJ AS,, can then be calculated

ss,

f% (14)

as:

= 1.5 (;)*

K To convert to angular width, 6tif = &S,/k sinei. As discussed in Sec. 4.2, by following the shape of the diffraction from a vicinal surface, one can observe step bunching

-0.10

versus growth.

-0.06

-0.02

0.02

0.06

0.10

s,(P)

Figure 6. Diffracted intensity exhibiting the split components from a vicinal surface. The intensity is plotted versus S, at S, = x/d for the case in which the distribution of terrace lengths is given by the shifted exponential distribution of Eq. (12). Results for several values of a/L, are plotted. The separation of the components is 2n I L.

RHEED Studies

4.0

DIFFRACTION

of the Dynamics

of MBE

665

MEASUREMENTS

The time evolution

of the diffracted

intensity

during epitaxial

growth

is in a sense a record of the development of surface atomic structure and morphology. Changes in surface reconstruction, increases in the density of point defects, the agglomeration

of islands, step bunching,

lattice plane

bowing near dislocations-all of these contribute to the diffraction. An important example is the periodic intensity oscillations observed during layer-by-layer growth. Our goal is to understand these diverse contributions and to deduce the rates of the important microscopic processes. Ultimately a coherent, quantitative, andyredictive picture of epitaxy should emerge. This section focuses on the diffraction from GaAs and related materials since these have been the object of most of the experimental work. Indeed, using RHEED, the first intensity oscillations were first reported on GaA~(l00).f~~j Recent work on metalsf10jf44)f451and elemental semiconductorsf46)-f48) shows that much of the same phenomena can be observed there. This section is organized as follows: first, two general classes of observations are distinguished, reflecting the qualitative differences observed in the diffraction from low-index surfaces as compared to vicinal surfaces. The main difference in growth on these two surfaces is the change that occurs when the incorporation of adatoms at the built-in steps of a vicinal surface competes with island growth on flat terraces. For the case of the low-index surfaces, the underlying mechanisms giving rise to RHEED intensity oscillations are described. Evidence supporting the kinematic view is presented along with results that as yet are not explained so simply. In addition, related intensity oscillations that give the energetics of the sublimation of GaAs are discussed. Second, the development of the diffracted intensity during growth on vicinal surfaces is presented. A model discussing the damping of intensity oscillations on vicinal surfaces is given. The two main topics are diffusion and step bunching. For both, the roles of the two different types of steps on zincblende crystals are important. Third, the dynamics of lattice mismatched growth is discussed. Strain and dislocation formation strongly influence the evolution of the diffraction pattern. Precise lattice constant measurements are compared to measurements

of intensity

oscillations.

666

Molecular

Beam Epitaxy

The main results GaAs(lO0)

are that upon initiation

of growth

on a smooth

surface, the diffracted beams usually broaden into the streaks

shown in Fig. 3 and the diffracted

intensity

to a monolayer

The separation

completion

time.

can oscillate with period equal of diffracted

beams can

also change. Our main message is that the intensity at the peak of the specular beam and the angular distribution of the intensity contain important complementary information on the growth process. The peak intensity is simpler to measure versus time, but to a large extent ignores the lateral structure of the surface. In this section, the time development of both kinds of measurements sized. 4.1

Low-Index

for the two classes

of surfaces

is empha-

Surfaces

Intensity Oscillations. Intensity oscillations observed at the start of MBE growth on a smooth GaAs(lO0) surface are shown in Fig. 7. These are measurements of the peak intensity of the specular beam versus time and occur simultaneously with periodic changes in the shapes of the diffracted beams. For these data, the period of the oscillations corresponds to the time required for the deposition of a layer of GaAs. In this case in which there is excess As,, it is also the time required to deposit a layer of Ga, since that is the rate-limiting stept50] of the growth. The exact form of the oscillations-their amplitude, damping, phase, and sometimes additional frequency components-depends on the scattering geometry (scattering angles Oi and I&) as well as growth parameters, so that care must be taken before growth information can be extracted. It is generally agreed that these intensity mode characterized pure step propagation oscillations

oscillations

by two-dimensional or step flow.

result from an alternation

well understood.

indicate

a layer-by-layer

island formation

It is also generally

to

agreed that the

in surface roughness.

There is no quantitative

growth

as opposed

Little else is

theory that accounts for all of the

features observed. One can try to consider models in which the step edges scatter out of the diffracted beamst51)t52) but not too much data has been compared to these. The kinematic treatment is presented here. As will be seen, kinematic theory is only partially successful and one goal could be to choose experimental conditions and methods to enhance its chance of success.

But another goal should be to develop theories

can account for the demonstrated scattering geometry.

sensitivities

of this measurement

that to


1

of MBE

667

d.

Time

Figure 7. Examples of the peak intensity of the specular beam measured after the initiation of growth on GaAs(l00). The period of the growth is the time to deposit a monolayer. The envelope, phase, and magnitude depend on scattering geometry as well as growth parameters. (a) Beats can be observed on a small sample with non-uniform illumination by the electron beam. The incident beam is on a symmetry azimuth. Note that the initial intensity increases. (b,c,d) The glancing angle was 8i= 33 mrad, the azimuthal angle +i = 7” from a [Ol i Idirection. (b) T = 580% and As,/Ga is 6, (c) T = 550°C (d) T = 580°C and As,/Ga = 80.

688


In Fig. 8, the sequence of specular intensity oscillations is measured during a growth in which the macroscopic surface plane was parallel to the (100) to within

1 mrad.

For these measurements,

the flux variation

over

the surface is less than one percent so that beats due to a range of periods do not contribute

to the envelope.f53]

Intensity

oscillations

on GaAs(100)

surfaces can be observed under a wide range of growth conditions, and for the data in panels 8(a)-(d) the growth rate is changed from 0.01 to 0.2 layers per second while the substrate temperature was maintained at 580°C. The temperature could be varied between 450” and 700” C with oscillations still observable. By contrast on a GaAs(ll0) surface, intensity oscillations can be observed only under a very narrow range of substrate For this (100) surface, the intensity temperatures and growth fluxes. oscillations shown in Fig. 8 exhibit trends that are commonly observed in other systems. After the initiation of growth, there is a sudden change that, depending upon scattering geometry, can be an increase or decrease. After this initial transient, the higher the growth rate, the larger the initial oscillation amplitude and the lower the baseline. Though the period corresponds to the time to deposit a monolayer, the maxima and minima do not necessarily correspond to the times at which integral numbers of layers are deposited. For these low-index surfaces, the baseline is nearly constant and the maxima damp slowly. Similar intensity oscillations in many epitaxial systems are now routinely used to measure growth rates during MBE.

0

TIME

SO

bed

Figure 8. The peak intensity of the specular beam versus time on surface for a variety of growth parameters. The surface misorientation 1 mrad, the incident azimuthal angle was 7” from the (010) and glancing angle was Bi = 30 mrad. This figure is to be compared to vicinal surface shown in Fig. 17.

a low-index is less than the incident results on a


The variation

basic

mechanismf141 of these

in surface

roughness

illustrated

in Fig. 9 where,

nucleated

causing

during

starting

a dramatic

oscillations

lies in the cyclic

layer-by-layer

with a smooth

decrease

669

growth. surface,

in the diffracted

This

islands

intensity.

is are As

growth proceeds, the surface diffusion and the preference for adatoms to bond at steps causes the islands to become larger until the layer is completed. During this process, the kinematic diffraction reaches a minimum at maximum roughness and then a second maximum as the surface becomes smooth. Depending upon the diffusion of the adatoms, it is likely that islands will nucleate on large terraces before the first layer is completed so that the intensity never recovers to its initial value and, in fact, exhibits an envelope which is damped. Ultimately, depending upon the growth of the particular materials, a steady state surface is reached in which there is no net change in roughness. The constant baseline should correspond to the situation in which the first term in Eq. (5) is small, with Even with more than two layers, this the second term dominating. kinematic result holds.fllOj If the steady state intensity equals the baseline intensity, it suggests that the disorder scattering is dominating at both the minima and the long time result. One steady state could be a case in which new islands are formed at the same rate as old islands are assimilated into larger ones.f54j The original surface can be recovered if growth is interrupted and kept at a temperature sufficient for significant adatom diffusion. For GaAs, an incident arsenic flux must also be maintained in order to replace arsenic that desorbs from the surface. The length of time required for the surface to become smooth depends on the same mechanisms that give rise to the intensity oscillations. The time, which is of the order of tens of minutes here, depends upon sample history, misorientation, and As, flux. For these well-aligned surfaces at 580” C, maximum smoothness could take as long as one hour. An important point is that these

intensity

oscillations

result from the competition

between

cluster

formation and step propagation. If the surface mobility of adatoms were such that all adatoms could migrate and attach at a step, then the built-in steps of the surface would just walk across the surface and there would not be a change in surface roughness-there would not be oscillations in the diffraction. On the other hand, if the adatoms did not diffuse, then the surface would just become randomly rough, giving at best very weak intensity 0scillations.f llW1lj Layer-by-layer growth requires both step propagation

and surface adatom migration.

690

Molecular

Beam Epitaxy

Figure 9. Schematic of island growth and coalescence during layer-by-layer Several processes are illustrated, growth, giving rise to intensity oscillations. including step propagation, nucleation on terraces and islands, coalescence of islands, and island growth. One expects that maxima intensity corresponds approximately to minimum roughness.

The kinematic interpretation of these changes in surface roughness gives a simple correspondence to the measured diffraction. In fact, the interpretation includes multiple scattering for large terraces within the limits of the column approximation, for the same reasons discussed further in Sec. 3.1. Complete multiple scattering calculations have not progressed to the point of realistically including random island growth.t2)t3)t4Q) The intensity oscillations can be understood in terms of the shape of the diffracted beam shown in Fig. 5. There the distribution of intensity along a streak is composed of two parts: a central spike due to long range order, and a broad part due to the step disorder. As seen later in Eq. (43), this spike depends mainly on the coverages of each layer. Starting with a smooth surface,

the intensity

consists

only of the central spike and the

diffraction is maximum. As islands are formed, the central spike is reduced in magnitude and the broad disorder scattering is increased. As the islands grow in size, the intensity

returns to the central spike until a

maximum is once again obtained. If third layers start to form before the second layer is finished growing, the intensity will not regain the initial value. The intensity then continues to oscillate as growth proceeds layerby-layer. Finally in steady state, some number of layers on average are always present and though locally the growth may be similar, the diffracted intensity does not vary. The magnitude of the intensity oscillations in this

RHEED Studies

of the Dynamics

of MBE

691

model depends upon the scattering angle corresponding to the dependence shown in Eq. (5). Simply stated, when the scattering angles are such that the path length difference integral

between

number of electron wavelengths,

the steps; when

the angles

different

the scattering

terraces

are an

is insensitive

are such that the path length

to

differences

between different terraces is a half of a wavelength, the diffraction is maximally sensitive to the steps. For example, when the angles are such that scattering from different steps is in phase, the intensity should not oscillate. How well this interpretation works can be seen qualitatively in Fig. 10, where intensity oscillations during the growth of GaAs on GaAs(100) are shown for several glancing angles@, These particular data were taken with the incident electron beam directed along an azimuth that was 7” from the [OlO] direction. At this azimuth there are a minimum of diffracted beams that are strongly excited[551 and the diffracted intensity of the specular beam is strong. One hopes that multiple scattering in the form of a sharing of intensity with other less strongly excited beams is therefore minimized. The main point is that the oscillations are weak at 43, 66, and 83 mrad, which are near in-phase conditions; and they are strong at angles corresponding to out-of-phase conditions. A more quantitative view is given in Fig. 11 where the ratio of the first minimum to second maximum is plotted versus eP This removes the importance of the initial transient. When this ratio is near unity, the oscillations are weak, and when the ratio is small, the oscillations are strong. Figure 11 shows again that the qualitative explanation works well. Further, if we assume that only two layers are involved in the diffraction (perfect layer growth), then from Eq. (5) at& = 0, the ratio of a minimum (half coverage or 8 = %) to a maximum (full coverage or 8 = 1) is 0~1 - cos(2kdBJ.

This ratio, with the constant of

proportionality fit at one point, is plotted as the solid curve in Fig. 11. The fit is remarkably good for such a simple calculation, though at small angles there is a shift from the data. For comparison, a similar plot for Ge on From our perspective, this agreement Ge(ll0) shows no such shift. indicates that kinematic diffraction.

diffraction

from islands

plays a large role in the

There are also serious discrepancies with kinematic theory, even at azimuthal angles at which there is coupling to only a few diffracted beams. Note that after growth is initiated the intensity will either increase or decrease, depending upon the glancing angle fib In the kinematic analyFor this to be a kinematic effect sis, the central spike cannot increase. within the assumptions made so far, the diffuse component due to the step

692

Molecular

Beam Epitaxy

disorder would have to dominate the diffraction. be a reconstruction

Alternatively,

there may

change once small islands are nucleated.

This is, in

fact, seen in the growth when the fractional after growth is begun.

order beams become weaker

To explain the shift of the ratio measurement

from

calculation, one could require that the scattering phase of a small island be different than that of the substrate, so that the simple path length argument is modified. However, it is difficult to explain the shift in the position of the maxima of the intensity oscillations that can be seen to some extent in Fig. 10 and which also depend on cpP These are described in detail by Joyce and coworkers.prl The main difficulty is that the maxima should correspond to a smooth surface in nearly any reconstruction-independent model that has moderate diffusion. Shadowing or a significant role of step scattering should produce strongly asymmetric oscillations, which are not usually observed. The strong ei dependencies are being investigated, and we note that they are evidently not seen in LEED.f561 Recently, Horio and Ichimiyaf57] have shown that dynamical considerations can at least qualitatively account for the phase shift. They treated a growth front with a potential that was proportional to coverage. The phase shift is seen to come from both the surface roughness and the coverage dependent potential.

Figure 10. Specular intensity oscillations measured during the growth of GaAs on lowindex GaAs(lO0) surface. The in-phase angles correspond to 6i= 43,66, and 83 mrad where the oscillations are weak. The out-of-phase conditions are in between and the oscillations are much stronger. The amplitudes are compared to calculation in Fig. 11.

TIME

(orb.

units)


of MBE

693

8i (mrad) Figure 11. Oscillation strength versus 61 from data like that in Fig. 10. The ratio of the first minimum to the first peak is plotted to show that interference between deposited layers is the cause of the intensity oscillations. At the in-phase angle of 65 mrad, the oscillations are weak; at the out-of-phase angle of 76 mrad, the oscillations are strong. At low angles the maximum and minimum are shifted, though the alternating trend is still clear. The solid curve is a calculation for a twolevel system from Eq. (5).

If the surfaces are sufficiently flat and if the growth does not produce defects, then the oscillations damp very slowly. Three cases are illustrated in Fig. 12, GaAs on GaAs(lOO), Ge on GaAs(llO), Fe on Fe(lOO). The first two are MBE-prepared substrates; the last is a whisker. In all cases, the lateral dimensions of the surfaces were sufficiently small that the effects due to flux variation

over the surface

were minimized.

example, Fig. 7(a) shows a case in which there is a 10% variation

For

over the

surface. The surface is small enough that the illumination of the incident electron beam is uniform over the surface. In this case, beats are observed. If the surface is large, giving both a flux variation and variation of electron intensity, then the oscillations are just damped. This must be removed from the problem if growth information is to be extracted from the damping.

694

Molecular

Beam Epitaxy

l~~~~l~~~~l~.~~l,~*~I 0 100 200 Time (set) III,

I ,,,,I,,,,

J

300

Figure 12. Intensity oscillations on very well-oriented substrates. (a) GaAs on GaAs(lO0) oriented to better than 0.7 mrad. (b) Ge on Ge(l10) lattice-matched to GaAs(l10). (c) Fe on an Fe(lOO) whisker. Note that in (b) the nucleation of Ge clusters is affected by residual arsenic.

Not only can cyclic variations in Layer-by-Layer Sublimation. surface roughness be observed during growth, but intensity oscillations during the sublimation of GaAs are seen.f581t591 From the temperature dependence of these oscillations, one should be able to learn about surface binding energies and the relative importance of equilibrium and kinetic processes.

Figure 13 shows oscillations

seen in the specular beam

intensity during the growth and sublimation of GaAs at moderately high temperatures. By first growing GaAs and then interrupting the Ga flux, both growth and sublimation could be observed without changing the sample temperature. (This procedure also obviates the need to compensate for temperature transients.) In Fig. 13, GaAs is first grown at 1 monolayer per 7.2 set; then growth is interrupted and there is a short time during which the surface anneals and the intensity increases. After this anneal, the intensity oscillates corresponding to the time required to


remove a layer. large enough change

We interpret the annealing

for sublimation

in surface

roughness.

period of the sublimation

of MBE

695

process to produce terraces

from the middle of terraces to cause a If the temperature is increased, then the

is reduced.

1

103x I/T

0

100

300

200

TIME (set) Figure 13. Intensity oscillations during both growth and sublimation of GaAs on GaAs(100). With the sample held at 665”C, GaAs is first grown at 1 monolayer per 7.2 s; then growth is interrupted, and after a short anneal period, sublimation oscillations are observed with a 25 s period. Sublimation

is only observed for certain surface structures.

13 shows the transition reconstructions, conditions

line between

plotted versus substrate temperature,

of no growth.

These

and As, flux under

data were determined

surface by measuring

the intensity

substrate

while the temperature

temperature

Figure

the 1 x 1 and c(2 x 4) surface on a low-index

of the quarter order reflection was raised.

versus

For low-index

surfaces only, the intensity is observed to decrease sharply, over 2”, at the transition temperatures shown. The sublimation oscillations could be induced by crossing into the 1 x 1 region, either by raising the temperature or lowering the As, flux. The latent heat of transition is 4.5 eV. This value has been disputed by Newstead et al.f60) as being too high; their value is 3.9 eV.

696

Molecular

Beam Epitaxy

15 pAs

pGa

2x4

10-

5

3x1

/ -

I

I

I

700

4x1

4x6 (76x2) I

750

I

I

800

850

Substrate Temperature

950

900 (K)

Figure 14. Observed reconstructions on GaAs(lO0) versus As, flux and substrate temperature for the case of zero Ga flux.

A mass action analysis gives agreement with these measurements. Equilibrium

between vapor and solid could be described Ga + 1/2As2 -

Eq. (15)

so that the pressures

GaAs

are related according to

PGa = Kp hs,)-

Eq. (16)

by the reaction

‘/2

By detailed balance, if the equilibrium Ga vapor is removed, Eq. (16) should give the sublimation rate of Ga.f6’) The agreement of the data with this mass action analysis is shown in Fig. 15 where the Ga sublimation rate is plotted

versus

arsenic

flux.

For these

measurements,

the sample

temperature was determined by a thermocouple calibrated at the AlSi eutectic. The arsenic flux was determined from an ion gauge calibrated against intensity oscillations in which an excess of Ga was present on the surface so that the supply of arsenic was rate-limiting.fce) At these substrate temperatures, identical results were obtained with either As, or As,; the slope of the curve in Fig. 15 was 0.5.

Similarly,

one can make an

Arrhenius plot of the sublimation rate of Ga versus l/T at several different As pressures; from this, one determined that the enthalpy of formation in Kp is 4.6 +- 0.2 eV, in good agreement

with the bulk data.


As

INCORPORATION IO

6

2

4

PRESSURE

As,

PERIOD 6) 0.6 0.4

I

( TORI?

02

)

Figure 15. From measurements shown in Fig. 13, the Ga sublimation plotted versus arsenic flux.

A similar analysis

697

can be used to determine

rate is

the mole fraction

of

AIxGa,_$\s at temperatures where sublimation of Ga occurs. Treating GaAs as a constituent with activity 1 -x, the sublimation rate of Ga is now

Eq. (17) As a result, the mole fraction,

x, of Al will depend

upon the sublimation

rate. Using Eq. (17), the normalized growth rate, 1 - TAT&~‘, can be calculated and the results for particular growth rates are shown in Fig. 16 to have excellent agreement with measured rates using RHEED. A main result of these studies is that above 580°C no differences between As, or As, fluxes have been observed. This means that As, is quickly equilibrated at the surface. The success of the mass action analysis also calls into question rate equation

models of the growth

kinetics.

For example,

it is tempting

to

argue that if 6 is the surface coverage of As adatoms and 1 - 0 is the fraction of Ga sites available, that the change in As coverage during growth is

Eq. (18)

-

= 2&Q&l at

- e)2 - 2k,02 - JGa t k&

- 0)

698

Molecular

Beam Epitaxy

Here the probability

of As, dissociating

is proportional

to the availability

of

two adjacent Ga sites, each As, molecule leaves behind two As adatoms, and the kiare temperature

dependent

rate constants.

Further, the desorp-

tion of As is assumed to be a second order process and all Ga that arrives sticks. The problem is that at equilibrium, this equation does not satisfy mass-action. For mass-action to be satisfied, the incorporation term for Ga would need to be replaced by .&&I. Tsao has suggested[661 that the difficulty can be fixed by formulating the kinetic equations in terms of chemical potentials.

S

s

1.0

s L

“,

0.5

L

0

N = i

5

2

0.0 650

750

100

Temporoture

.C

Figure 16. Measured growth rate (reciprocal intensity oscillation period) versus substrate temperature during the growth of Al,Ga,,As divided by the rate at 550°C. As the temperature is increased, the sublimation increases in quantitative agreement with Heckingbottom’sb’] mass-action analysis.

4.2

Vicinal Surfaces

Intensity Oscillations. Commercial GaAs(lO0) surfaces are typically misoriented from the (100) by about l/3” or 5 mrad. The terraces that make up this macroscopic misorientation are on the average approximately 500 8, in length and contain about 150 atomic rows. Since intensity


oscillations are easily observed mobility at this temperature with step propagation. one finds conditions

on such surfaces

of MBE

at 600°C

must be such that cluster formation

At higher temperatures at which step propagation

the adatom competes

or higher misorientations, dominates,

with the result

that no intensity oscillations are observed. When step propagation of built-in staircase contributes to the growth, the qualitative form of This is illustrated in Fig. 17 where intensity oscillations changes. results from a sequence of growths on a 5 mrad surface, using The striking change is that parameters as in Fig. 8, are shown. oscillations are now strongly damped, with minima and falling together. For this surface, the strength of the qualitative form of the oscillations are independent of direction, whether down the staircase or parallel to the

TIME

699

the the the the the

maxima rising and damping and the the incident beam step edges.

(sec.)

Figure 17. Sequence of intensity oscillations measured on a GaAs(l00) surface misoriented by 5 mrad. The growth parameters are close to those used in Fig. 8. There

is only a quantitative

difference

between

growth

on vicinal

and low-index surfaces, but when significant numbers of adatoms can reach the step edges of a vicinal surface, variations in adatom and adcluster density on an individual terrace becomes the focus. As growth proceeds, this terrace moves across the surface. The notion of a layer coverage oscillating between zero and complete coverage does not mean

700

Molecular

Beam Epitaxy

as much since, in a monolayer completion time, a terrace just moves a distance equal to its length. In fact, on a staircase, even half-coverage becomes difficult to reach. The more mobile the adatoms, the less change in roughness there can be. As the temperature

is raised, the layer-by-layer

nature is enhanced

at the expense

by the step propagation

of cluster

formation. The minima which correspond to maximum surface roughness are then not as deep. Step Structures on Zincblende. There are two types of steps on GaAs(l00) and these might be expected to affect the growth differently. The distinction arises for the same crystallographic reason that GaAs(ll1) is either As- or Ga-terminated. If the (100) surface is misoriented in the (011) direction, then (111) steps will be forced. For the same reason, on a low-index surface, sides bounded by Ga-terminated steps and terminated ones. One expects that these

This is illustrated in Fig. 18. each cluster will have two two sides bounded by Asdistinct steps might have

different kink densities and different probabilities of Ga attachment. Other processes, such as the step-catalyzed dissociation of As, postulated by Singh,[641 might

depend

on type of step.

c-L--

Figure 18. Ball and stick model of Ga-terminated and As-terminated GaAs(l00) staircases, assuming bulk termination, formed by miscutting from the perfect (100) in the (011) and (007 ) directions. One expects that the real surface will have reconstructed terraces and reconstructed steps. Real steps will meander depending upon the thermodynamics of kink formation.

RHEED Studies

of the Dynamics

of MBE

701

As described in Sec. 3.3, the signature of a vicinal surface is a specular beam that is split into several components. The angular separation of the components measurement

must follow

Eq. (11).

Figure

of the split specular beam for GaAs(lO0)

26a shows

a

surfaces misoriented

by 2” toward the (011) and (011) directions. The misorientation-induced splitting, as well as the disorder (broadening) among the terrace lengths, is seen. In Fig. 27a and b, the beam is directed parallel to the ledges and the width corresponds to the kink density. These measurements were taken without growth but in an As, flux while the surface was in the c(2 x 4) reconstruction. One can see that the Ga-terminated steps are straight with a large variation in terrace lengths while the As-terminated steps have a high kink density but a narrow distribution of terrace lengths. Upon increasing the temperature or lowering the As flux to go into the 1 x 1 reconstruction, this difference disappears. If growth is initiated, this difference is more difficult to observe. Surface Diffusion. To study surface diffusion, one makes use of the observation that the strength of the intensity oscillations on vicinal surfaces is reduced if the mean terrace length is decreasedt53) or if the surface temperature is increased. For vicinal GaAs(l00) surfaces, Neave and coworker@1 fixed the Ga flux, then raised the substrate temperature until the intensity oscillations were extinguished. They then associated the mean terrace length with a diffusion length and extracted a diffusion coefficient at this critical temperature using the relation that Dr = I,*. Fort they chose the time to deposit a monolayer. The rationale is that if adatoms have sufficient mobility to reach a step edge, then the growth is by step flow. There is no change in the surface roughness during growth. On the contrary,

if the adatoms do not have sufficient

step edge, then nucleation

occurs, the roughness

mobility to reach a

of the surface changes,

and intensity oscillations are observed. After one critical temperature was determined for one Ga flux, the Ga flux was then changed and the measurement repeated. Knowing t and I,, a plot of the diffusion coefficient, D, versus reciprocal substrate temperature, l/T, was obtained. For the MBE growth of GaAs, they found thatD = 5.3 x 1O-l* exp[0.3(eV) lk,T]. This result has now been used in several Monte Carlo calculations.~2)~3) The interpretation of the experiments even within the rationale of the measurements is not clear-cut for several reasons. obvious what to take as the diffusion time t or diffusion

Mainly, it is not length. Note that

there are several characteristic

There is the time

times for the problem.

702

Molecular

Beam Epitaxy

required to deposit a monolayer, mean time between

collisions

which after adsorption, incorporates

the residence of adatoms,

time of a Ga adatom, the

and the time for an adatom

diffuses to, and after some number of attempts,

at a step.

A second difficulty with the approach (actually diffusion on this surface in general) is that the diffusion of Ga, as well as the Ga sublimation rate,t61tt561might depend upon the As flux and on the surface reconstruction.p41 The surface reconstruction, in turn, depends upon the As concentration. When the temperature is raised to extinguish oscillations at the different growth rates, the relative As to Ga fluxes will have changed-the Ga sublimation rate and the As concentration will also have changed. One diffusion coefficient might be insufficient to describe the process. Finally, it is difficult to determine the probability with which an adatom incorporates at a step. This means that both an adatom that has a short diffusion length and one that has a high diffusion length, but low probability of incorporation at steps will likely nucleate a 2D island, giving rise to RHEED intensity oscillations. It will be difficult to separate the two effects. For example, one would like to use this method to study diffusion anisotropy on GaAs. Figure 19 shows the results of changing the surface temperature at a fixed As flux for 2” misoriented (100) surfaces. The left panel shows intensity oscillations for the case in which the steps are Gaterminated. The right panel exhibits results for As-terminated steps. As can be seen, these measurements give very different results. On the Gaterminated surface, the intensity oscillations are well behaved and the one can usually determine a clear temperature at which the oscillations are extinguished. On the As-terminated surface, however, this determination is not always

so straightforward.

This is seen in Fig. 19 which

shows

intensity oscillations for various temperatures near the transition between step flow and 2D cluster formation. The growth rate is fixed. There is a strong difference on the two surfaces but one can not say whether adatoms diffuse more easily in the [Oli ] direction or whether incorporation at Asterminated steps has a high probability. In determining diffusion over a wide range of growth conditions, one must be especially wary of changes in surface reconstruction. At growth rates, away from the 2 x 4 to 1 x 1 phase transition, it may possible to apply Neave’s method, but near the appearance of the 1 reconstruction, some other bond-breaking process appears to be erating.t4s)tr4)

low be x 1 op-


of MBE

703

Figure 19. Intensity oscillations on 2” misoriented GaAs(100) A and B surfaces. The behavior of the oscillations is very different indicating that diffusive and/or step incorporation processes are different on these two surfaces. On the 2”B misorientation, the method of determining the transition between step flow and 2D cius-

ter nucleation is difficult.

b.

a. TIME

TIME

In fact, the concept

of diffusion

during

MBE is problematical

be-

cause of the nonequilibrium nature of the process.~51~6~ Monte Carlo calculations indicate that adatoms adsorbed at the start of a layer migrate very large distances, while those adsorbed near the completion of a layer travel on average only short distances. This strong concentration dependence seems to indicate that it makes more sense to examine the microscopic hopping and cluster probabilities more than macroscopic diffusion. One expects that diffusion in which there are multiple types of sites available for hopping will be equilibrated only if all such sites are visited in the correctly weighted way. Vvedenskyp*l has used a diffusion coefficient

with a hopping

probability

that is weighted

(exponentially)

by

the number of near neighbor bonds of an atom at a particular site in its diffusive motion. The argument described in the Sec. 4.2 corresponds to the limiting case of the absence of adatom-adatom collisions and cluster formation

so that equilibrium

diffusion

notions are more likely appropriate.

Continuum Analysis. To put this measurement of diffusion on a firmer foundation, Nishinaga and coworkersr8] have associated the appearance and disappearance of intensity oscillations with the transition between step flow and nucleated growth in a BCF formulation. The main feature of their analysis is the solution of the BCF equation, carefully considering the importance of desorption of adatoms versus incorporation

704

Molecular

Beam Epitaxy

at step edges. The advantage of their method is that by focusing on the step flow side of the transition, they are able to avoid the intricacies of the coarsening

of the 2D islands

and the complicated

nature

of diffusion

between island edge sites, adatom sites, and steps. By treating the limit of no nucleation,

adatom-adatom

collisions

can be ignored and they can deal

only with incorporation at steps. The method is to solve the diffusion equation, as described below, and then calculate (7) the condition at which islands form and (2) the In terms of understanding surface growth rate on a vicinal surface. diffusion, the second is important since it supplies an independent check on the relevant lengths of the problem. First, we treat the determination of the condition for the disappearance of intensity oscillations in the growth of GaAs. We describe how Nishinaga’s model accounts for differences in incorporation at step edges. Simplifying Nishinaga’s

analysis

slightly, write the surface diffusion

equation for growth on a staircase of steps, each of equal length. Examining growth at typically low MBE temperatures, we assume that there is no sublimation. Then the adatom density is just determined by a balance between diffusion towards step edges and adsorption. Assuming that the density is too low for nucleation, the time rate of change of the mobile adatom population n(x, I) is given in one dimension by

Eq. (19)

Jn(xJ) at

= D

a*n(%t)

+J

lx*

where 0 is the surface diffusion

coefficient

and J is the incident flux of Ga

atoms. For the coordinate system, choose the origin as the center of a step of length L. Initially solve Eq. (19) in steady state subject to the boundary condition that n(+ L/2 ) = 0. This assumption implies that the steps are perfect sinks, always incorporating incident adatoms. This assumption will be removed shortly, but for now it helps to emphasize the main features of the method. Solving Eq. (19) in steady state, the mobile adatom density on the surface is given by

Eq.(20)

n(x) =

& (f4)


705

In this symmetric problem, n(x) has its maximum at x = 0 and is given by n = JL2/8D. The main assumption of Nishinagave) is that nucleation will occur when this is as large as some critical value

Eq. (21)

ncrit =

anon

where nSOis the equilibrium adatom population of the surface held at a temperature T and the factor a (also slightly temperature dependent) is given by classical nucleation theory.~8j Then nucleation will occur when

Eq. (22)

&L2

= 8Dan,

To make this slightly

more general, the boundary condition at the step This can be seen by first considering should be n( + L / 2) = n,,. equilibrium, where there is no growth. The surface is bathed in a flux of Ga There will be some mobile adatom and As that just match desorption.

population on the surface and since there is no growth, the gradient (the diffusive flux to the steps) must be zero. This means that there will be a uniform adatom density n,, over the entire terrace. If the Ga flux is increased so that there is growth, this adatom population will increase, except that, at the steps, there can be sufficient adsorption sites to be able to maintain the equilibrium. Any nonequilibrium distribution at the steps is suppressed. With this more reasonable boundary condition, the condition for 2D cluster formation is slightly modified, becoming J&,2

= 8D(a

- l)n,,

This condition of perfect equilibration will be relaxed again shortly. In Eq. (23) the flux and surface concentration can now be taken as a unit width of surface, or more simply, as two-dimensional fluxes and concentrations. To make it dimensionless, divide each side by n, = l/a2, the density of atomic sites in a monolayer

of a square lattice with cell side a. Then since

J = n,/t, this becomes Eq. (24)

L2 = 8Dt(a

- l)O,,

where 8,, is the equilibrium coverage of mobile adatoms and t is the time to deposit a monolayer of Ga or GaAs. For a givenL and coverage 0,,, the density of mobile adatoms will be sufficiently high that nucleation is the

706

Molecular

Beam Epitaxy

preferred growth mode. At lower steady state coverages, nuclei will not be stable, allowing step flow to dominate. This equation should be compared to the model of Neaver’l

in which they took L* = Dt.

The complication in determining the diffusion coefficient, both D and 8, are exponentially dependent on temperature. possible to separately

measure these quantities.

D, is that It is not

The best one can do is to

factor out the temperature dependence of the equilibrium flux as follows. Under equilibrium conditions for GaAs(lOO), the desorption rate of Ga is n,,/t, where t, is the surface lifetime of Ga. Assuming unity sticking coefficients, this equals the incident equilibrium to convert pressure to flux

Eq-(25)

n

-

flux.

Using kinetic theory

&a% d?zGim

so-

or using the law of mass action,

Eq. (26)

nso=

Kpts P’h AS d2xmkT

Combining this last equation with Eq. (24), one obtains the condition the disappearance of intensity oscillations to be

for

Kp Eq. (27)

L* = 8t(u-

l)ot, PA/: dz!zz7

Nishinaga takes Kp = 4.3 x 1015 exp(-4.6/kT). The measurement procedure is once again to, first, fix the growth rate and As flux and choose the substrate misorientation; then vary the substrate temperature until the intensity oscillations disappear. Assuming that the exponential temperature dependence in KP given here is correct and that mass action applies, one can determine the exponential dependence of the length h, = (DQ)‘~. At higher temperatures this is easily modified to include sublimation in Eq. (23). How well this entire approach works will depend on how clearly the disappearance of the intensity oscillations can be determined and how well the condition for 2D cluster formation is followed. For GaAs(lOO), Nishinagafr8] finds h, = 4.0 x lo8 exp (Ed/kg cm, where Ed = 0.3 eV.


The reasonableness exam-ining the component Yexp(E@)

wherev

= lo’*

of MBE

707

of this measurement can be evaluated by Using TV = quantities in the calculation. and Ed = 1.7 eV, one obtains at an As pressure

of 4 x 10m6torr and, at a substrate temperature

of 55O”C, a surface lifetime of 2.5 x 10m3s and a surface coverage of O,, = 9.2 x 10e6. The latter seems

very low, though the ratio gives a reasonable

sublimation

rate.

On a 1”

surface, this equilibrium surface coverage corresponds to one adatom in a terrace of length 160 A separated by other atoms by about 1O6A. Especially if the steps meander, this might not be appropriate. Nishinaga calculates a to be of the order of 10. At the point of 2D cluster formation, the atoms would, on the average, be separated by 105A. These values are smaller than one might intuitively expect for critical cluster formation. To include partial incorporation at the different types of steps that are created when GaAs(lO0) is misoriented to either the [Ol l] or [Oli] directions, the boundary condition in the preceding analysis must be changed. We assume that the critical nucleation condition is unchanged. Once again following Nishinagafrgl and Burton, Cabrera, and Frank,fr7) let the equilibration time at a step betk Without growth and in equilibrium (an incident flux J,), the surface adatom concentration on a terrace, n,,, must again be uniform. Then, at an edge, the flux to the edge must balance the flux leaving the edge, i.e.

0 -

anso _

J,S’eP

5 Here the first term on the right represents the equilibrium

flux of n,, atoms

moving at a velocity of a/t, toward the step. The possibility of reflections is included in the value of tk. The second term represents the flux desorbing

from the step and is assumed to hold even away from equilib-

rium. Away from equilibrium, then given by

Eq. (29)

J,=- anstql tk

the density of adatoms next to the step is

- - a”,, Tk

where J, is the growth flux. In the case of minimal sublimation the growth flux is just the density of atoms in a layer, n, = l/a*, times the velocity of a step, L/t. The growth flux must be divided by two since we are only considering the flux on one side of a step. The boundary condition that

708

Molecular

Beam Epitaxy

must be applied to solving Eq. (19) is then that the adatom density at x = *L/2

is L

Eq. (39)

Repeating

%tep = %o + -

2a3

the arithmetic

tk 7

followed

L2 (1 + 2

before this gives

) = 8Dt(a

- l)e,,

Thus, adding to the complication that the exponential temperature dependence of 8, must be determined, one needs to correct the measured diffusion coefficient for differences in tk. Further, if tk = zS, then sublimation must be incorporated into the analysis.~8) This analysis is able to handle (7) the transition between step flow and cluster formation, (2) sublimation rates comparable to growth rates, (3) an average step incorporation time, and (4) a variable As, flux. The last is especially important since one expects that, at low As, to Ga flux ratios (or As,), the Ga mobility over the surface should be exceptionally high. This last can be seen in the widths of the RHEED beams and in the success of the technique of migration enhanced epitaxy (MEE).f@‘) In Nishinaga’s analysis, the diffusion coefficient is independent of the As flux. Instead, the As flux enters by controlling the equilibrium surface concentration, n,. This has two effects: the first and main one is that at reduced As flux, the critical concentration for 2D cluster formation is increased because increased.

the equilibrium

surface

concentration,

This more difficult requirement

n,,, and hence

on 20 cluster formation

anSO is means

that Ga adatoms are more likely to reach a step edge and incorporate. The second effect of a reduced As flux and increased nso is an increased sublimation rate. These are somewhat indirect. There is not yet quantitative consideration

in changes

in step incorporationfsll

or As dependent

changes in the defect densities. Nishinaga

has also examined

the validity of the diffusion times and

diffusivities in a beautiful consideration of the variation in mole fraction InGaAs growth on vicinal surfaces, that doesn’t rely on a calculation

of of

densities for critical nucleation. He compared, measured, and calculated mole fractions versus misorientation angle. The main effect is obtained by noting that for mobile adatoms that have large sublimation rates, the chances of their incorporating at steps increases if the mean distance to a

RHEED Studies

step edge is reduced. solving the diffusion

of the Dynamics

of MBE

709

The method is similar to the preceding analysis, also equation to find the incorporation rate of mobile

adatoms in the step flow regime. L is much less than the diffusion

For the case in which the terrace length length h,, the incorporation

rate is found

to be approximately JGa - n$ /ts, i.e., the excess above the equilibrium sublimation rate. For the case in which the terrace length is several times the diffusion

length, the incorporation

rate is

Eq. (32)

If the equilibrium sublimation terms are included when calculating the expected mole fraction X, then the measured mole fraction of In, xmeaS, is given by X

Eq. (33)

X meaS

=

Llh, + x(1 -L&j

This gives results close to the exact formulation of Nishinaga at high values of L/h,. Note that when considering the sublimation, the As flux affects both the Ga and In equilibrium adatom densities and that stoichiometric In,Ga,_+ is required. How well this formulation works suggests that the step flow model of growth and diffusion lengths are quite reasonable. Knowing the onset of intensity oscillations is then important only for ensuring that the growth is in the step flow regime. Collision Time Analysis. Kuroda[s2] has suggested that the collision time between adatoms is a more appropriate choice to determine the condition at which intensity oscillations disappear. In this model, the assumption is that a dimer constitutes a smallest stable nucleus. Substrate temperatures are assumed to be sufficiently low that sublimation does not occur. The rate at which dimers form is proportional to the square of the adatom density so that the time dependence mer) density is Eq. (34)

WI -

at

of the adatom (mono-

=J- Dn2

where n(t) is the adatom population solution, assuming that n(0) = 0, of

and J is the incident flux.

This has

710

Molecular

n(t) = n(m) tanh(f/tJ

Eq. (35) where

Beam Epitaxy

t, = l/mand

adatoms.

n(m) = m

The time tC is interpreted

are deposited.

is the steady

state population

of

as the mean collision time as adatoms

The density of dimers is easily seen to be n(m)(x - tanh x),

where x = t/t=. The argument is that at a time t = t, the monomer density is roughly comparable to the dimer density of the order of a At this point the density of nuclei will not increase since collisions with nuclei that are already established is likely. When this critical density of nuclei have formed, the mean area per nucleus is the reciprocal of the density or m This is the capture area of each nucleus and subsequent deposition is incorporated into these nuclei without the formation of new ones. The diameter of this area is approximately Eq. (36)

h, = (D/Y)”

and Kuroda argues that RHEED intensity oscillations will damp strongly when half a terrace length equals h,. The condition that should be compared to that of Neave is

Eq. (37)

DC=--

L4 16s

where L is the mean terrace length of the vicinal surface, t is the monolayer time, and a is the side of the unit square mesh (J= l/ta*). This third estimate will give the same activation energy for diffusion as Neave but will give a larger pre-exponential factor. Comparison. The determination of diffusion from the point of rapid damping of RHEED intensity oscillations

is thus seen to be largely depen-

dent on the choice of the size of the stable nucleus. The surprise is that the dependence on the terrace length size is very different. In the Neave model and the Nishinaga

model, the point at which

intensity

oscillations

are extinguished depends on the square versus Kuroda’s fourth power of the terrace length. Especially for the latter, at a given substrate temperature, the strength

of intensity

oscillations

will quickly

decrease

as the

terrace length is reduced. Step Bunching. The growth of AI,Ga,_xAs shows a very clear case of step train disordering. This was observed by Tsui and coworkersts3) who found that AlGaAs grown on GaAs(l00) exhibited a textured surface

RHEED Studies

of the Dynamics

of MBE

711

morphology depending upon the direction of the substrate misorientation. This macroscopic disordering can be followed by RHEED during growth. As shown

in Sec. 3.4, the characteristic

surface consists separation

of split peaks.

of the peaks is determined

of the components

is determined

diffraction

pattern

When the misorientation by the misorientation

of a vicinal

is large,t32] the and the shape

by the disorder.

If the terrace lengths in the staircase become more disordered, the components will broaden. Our method is to measure the shape of the split beams versus growth on the two types of zincblende step terminations.

time during

As a function of the Al mole fraction and substrate temperature, the step train order as described by RHEED shows a strong variation with time during growth. Both the Schwoebel effectts4) or step pinning by impuritiests5]-t87) could be involved. Figure 20 shows the RHEED profile during growth of AI,Ga,_xAs with x = 0.25 on a GaAs(l00) substrate that was misoriented by 2” toward the (11 l)A. For this growth, the ratio of column V to column III flux was about 2, the growth rate was 1 pm/hr, and the substrate was held at 675°C. This diffraction profile was measured with the incident beam pointing down the staircase of steps on the surface. First, one should note that though the two components forming the split specular streak were of nearly equal intensity and width at 580°C at this relatively high temperature where one can grow smooth AlGaAs, the relative intensity of the two peaks are very different. This is a reversible change with from a 2 x 4 upon growth there might GaAs(l00)

temperaturef86] and could be due to a reconstruction change to the 1 x 1 that exists at these high temperatures. Second, there is not much of a change and, for this case, after an hour even be a slight ordering. By contrast, for growth on a

surface

same conditions,

misoriented

the staircase

toward the (lll)B, disordered.

at very close to the

This is seen in Fig. 21 where

after starting with split components that were sharp, growth for 40 min left a surface with only a weakly resolved, characteristic step diffraction pattern. This is reversible in the sense that if GaAs is grown on top of this surface, the original ordered step train is developed. The resultst85)t86) for a variety of temperatures

are shown in Fig. 22

for both A and B step terminations. Here the width of the components of the split peak corresponds to the variance in the terrace length distribution as described in Eq. (14). For the data, one sample was used; after each growth a buffer of GaAs was grown and the surface annealed in an As, flux to obtain a sharp diffraction pattern. Differences in the initial peak width reflect this procedure.

The main point is that the amount of disordering

is

712

Molecular

Beam Epitaxy

reduced if the substrate temperature used.

Further,

terminated immediately,

there

steps. while

is raised or if the A step termination

is some disordering

Finally,

even on surfaces

at some temperatures

at others

there

the disordering

is a slow

initial

delay

is

with Gabegins

before

the

disordering begins in earnest. For both surfaces and at all temperatures, the disordering is more rapid as the Al mole fraction is increased.

2” toward the (1ll)A. For this growth the ratio of column V to column Ill flux was about 2, the growth rate was 1 pm/hr, and the substrate was held at 675°C.

r-

2.5 mrad

Figure 21. RHEED profiles of the split specular beam during the growth of AI,Ga,,As on a GaAs( 100) substrate that was misoriented by 2” toward the (1ll)B. For this experiment, the sample was mounted next to the sample used in Fig. 20, so that the growth conditions are essentially identical for the two experiments.

AI,,,Ga,,AS

,

*

1)I/hr

FWHM

As steps ei =72

mrad

_c

-

39 _.I 50

c

(mradl

-

’ 60

a

n

70

ef (mrod)

5.9 80

90


B

A$,,,Ga,,As,

surface

of MBE

v III

I y/hr

713

= 2.5

635’ A

1 0

IO

surface

I

I

20

30

time

40

(rnin)

Figure 22. The peak halfwidths of one component of the split specular beam for both A and B

versus time during growth for several substrate temperatures misorientations (2’). From Eq. (14) the halfwidth corresponds the distribution of terrace lengths.

to the variance

of

Gilmerie8] has developed a simple rate equation model to describe the step bunching process in terms of one parameter. This parameter reflects the asymmetry in the ease at which an adatom can cross a step from different directions. Later the model was reworked by Tokura et al. to obtain a relation for the time dependence of the variance in the terrace length from the mean. The basic mechanism is illustrated in Fig. 23. Here a staircase of steps is shown with terrace length T,, for the nth step. The growth rate is one monolayer in t seconds and one assumes that there is only step flow. The parameterrl describes the asymmetry in attachment to the neighboring

steps.

If adatoms striking the nth terrace could attach at

the (n-1)th terrace as easily as the (ntl)th terrace then q = 0. If, however, there is an asymmetric barrier for crossing a step from the right over the left, then, as shown, there might be (1 t r1)/2 enlarging the (n-l)th terrace and (l- ~)/2 enlarging the nth terrace. There are four such terms which, when summed, yield a rate equation describing change in the length of the nth terrace, T,,:

714


dT, = dt

Eq. (38)

- TrJ +

'3

U', - Tn-11

To solve this,tsg] expand in a finite series and assume periodic conditions.

Set

N -1

q = O,l, . . . . -

to obtain the variance

Eq. (39)

boundary

N

2x

of the step terrace length distribution

AT,;= ):IT,(O)12 exp[-4?l(ti)sin2q/2] rl*O

Depending upon the initial distribution of terrace lengths and on the asymmetry parameter, one can determine the subsequent behavior. If this asymmetry parameter n is positive, then the staircase orders; but if it is negative, the fluctuations diverge. For ordering, one requires that there is a feedback mechanism by which small steps grow faster than larger steps. One should note that the initial distribution is important since if there is no initial disorder, then there is no reason for the growth to prefer one step over another so that no change will take place. If however there is an initial difference between steps, then an asymmetry can take over. To illustrate this, assume that initially the T,, = T + d, with a random distribution.

Then the Fourier components

are identical

and a calculation

of Eq. (39) for several values of rr is shown in Fig. 24. Only slight values of asymmetry

are needed to achieve

modest agreement

with the data.

For

some experimental curves, there is an initial delay that cannot be fit with this assumed random initial terrace length distribution. Finally we should note that impurities could also be important. They could adsorb randomly, pinning some steps at the expensefag] of others, producing disorder.

terrace length


of MBE

715

l/~ monolayers / set

Figure 23. Definition of the asymmetry ordering or step bunching.

parameter

n that gives

either

step

TERRACE WIDTH ROUGHENING/ORDERING

tY

b

7]=0.0001

10

20

30

’ 0

TIME (minutes) Figure 24.

A calculation of the terrace length roughening or ordering from Eq. (39). The small values of ~1give modest agreement with experiment. The long delay before disorder initiation in some of the data of Fig. 22 cannot be fitted with the assumed random initial distribution of terrace lengths.

716

Molecular

Beam Epitaxy

Step Meandering. Though we do not know the detailed structure of the Ga-terminated or As-terminated steps, it does appear from simple analysis of the diffraction data that Ga-terminated

steps are much straighter

than the As-terminated ones. The picture of the GaAs(lO0) surface that is deduced from the data is schematically illustrated in Fig. 25. This result is similar to the difference in kink density of the two types of steps that can form on Si(lO0) .P)fgl The diffraction patterns from surfaces with these two types of step termination are strikingly different. If two samples are placed side by side on the holder, and the beam switched from one to the other the degrees of order. Note that the A and B misorientations can be prepared by taking a wafer that is polished on both sides, cleaving it into two, and then mounting both but with one turned upside down. Examination of the 2 x 4 reconstruction and determination of the staircase direction from RHEEDfgo) or from x-ray diffraction quickly gives the step termination. even the eye can distinguish

Flgure 25. Schematic diagram showing the order (disorder) in the two types of stepped surfaces that can be created on GaAs(lO0). The data to be presented in the next figures show that Ga-terminated steps are straight while there is severe terrace length disorder. For As-terminated steps, the step edges meander though the variation in terrace lengths is less. This holds for GaAs(lO0) at 600°C with the 2 x 4 reconstruction and not the 1 x 1 reconstruction.

RHEED Studies

of the Dynamics

of MBE

717

To determine the terrace length order, the incident beam is directed down the staircase of steps and the diffracted beam intensity is measured along the length of the streaks.

This is described

for the two types of step termination

in Sec. 3.3. The results

are shown in Fig. 26. For these data

two GaAs wafers were mounted side by side on the sample holder to minimize differences in sample history, incident flux, and temperature. A GaAs buffer was grown by MBE and the measurements were made under an As, flux and with the surface in a 2 x 4 reconstruction. Note that in the top curve of the figure, the cut off peak is an artifact. The main result is that the components of the peaks for the Ga-terminated surface are much broader than those for the surface with As-terminated steps. This broadening is removed if the temperature is raised to cause the higher temperature 1 x 1 reconstruction and can be recovered reversibly. Based on the analysis of Sec. 3.4, we estimate that the rms deviation in terrace lengths twice as high for the Ga-terminated step structure.

‘iii C

5 2

e

&

I

iTI b.

AS steps

I

1.6 mrad

-5

-c-

60

70

80

90

8f (mrad) Figure 26. Intensity profiles along the specular beam for GaAs(lO0) misoriented by 2”. The flat top is an artifact. The Ga-terminated stepped surface gives broad components corresponding to terrace length disorder. The As-terminated stepped surface exhibits much less disorder in terrace lengths.

718

Molecular

Beam Epitaxy

To examine

the meandering,

the electron beam is directed parallel

to the steps. Once again the difference is striking.

between the two step terminations

The data are shown in Fig. 27 and are from the same samples

as in Fig. 26. The top curve corresponds

to the As-terminated

surface (the

beam is in the [Ol l] direction parallel to the steps) and the bottom curve corresponds to the Ga-terminated step surface. For these data, a slit detector was used to integrate over one of the split components. At this point, these measurements have not been evaluated over the range of scattering geometries, terrace lengths, and substrate temperatures to make quantitative statements about observed trends. Nonetheless, the results are reproducible for a few misorientations and many sample preparations, indicating that the Ga-terminated steps meander less than those that are As-terminated. Recently, Pashley[gl] has confirmed these results with scanning tunnelling microscope images of A and B GaAs(lO0) surfaces. In addition, he has given an electron counting argument describing the relative stability of the various structures that are observed and not observed.

3i=30 mrad

a.

Ga steps

4

-

-

3.6 mrad Figure 27. Intensity profile of one of the split components with the incident beam directed parallel to the step edges. A slit detector is used. The Ga-terminated steps meander less.

As steps b.

I IO

20

30

40

8f (mrad)

50

60


4.3

Strained

719

Layer Growth

Recent interest has focused

of MBE

in device applications

investigations

for strained

on the influence

of coherent

epitaxial

layers

strain

in the

kinetics of epitaxial growth.tg2]-tg4) B esides devices such as MODFETs and quantum well lasers that utilize coherently

strained epitaxial films in active

regions of the device, one can use strain to tailor the band structure of materials or to combine optical materials with Si technology. In Ill-V MBE, an important system is the growth of In,Ga,_ &s on GaAs(100). The lowest energy state of an epitaxial film with a lattice parameter that is slightly different than the substrate is coherently strained to accommodate the mismatch. In this case, known as pseudomorphic growth, the in-plane strain in the film is equal to the lattice mismatch. Since the lattice parameter of InAs is approximately 7% larger than that of GaAs, strains of between 0 and 7% can be induced by varying the mole fraction of In contained in the alloy. A major consequence of the very large strains used in these films is the generation of dislocations at the film/substrate interface that can relieve the misfit strain and allow the film to relax toward its bulk value. The thickness at which dislocations are created is termed the critical thickness and is an inverse function of strain and film thickness. At this point, the effect of strain on nucleation and growth of clusters, surface reconstruction, step propagation, or on surface adatom mobility is not well understood. We cannot predict which lattice-mismatched systems will give pseudomorphic growth or, even, epitaxy. Intensity Oscillations. The main difference between the growth of In,Ga, _,+s and GaAs is that the former contains two components with very different surface mobilities. This could partially be due to strain but is also due to the weaker bonding of In. This can be seen to some extent in the conditions under which intensity oscillations are observed during homoepitaxy on InAs(lOO) substrates. Like the case for GaAs(lOO), on InAs(lOO) Intensity

the surface oscillations

rich c(4 x 4) surface

mobility

behaves

are observed

as if it is very As-dependent.

under growth conditions

reconstruction

at around

giving the As-

430” C and cannot

be

observed in the In-rich C(8 x 2) on a surface with a misorientation as small as 0.5”.t1sj In comparison with Ga on GaAs(1 OO), the In adatoms are quite mobile. InxGa, _& grown on GaAs(lO0) has a lattice mismatch from 0 to 7% depending on the In mole fraction in the film. The measured specular intensity during growth of In,Ga, _,As at 51O”C, as shown in Fig. 28 exhibits

720

Molecular

Beam Epitaxy

RHEED intensity oscillations over the entire range of x, thus yielding the film growth rate and the mole fraction, assuming the GaAs growth rate is known.

Knowledge

strain via Vegard’s

of the mole fraction

allows calculation

of the misfit

law, which states that the alloy lattice parameter

is a

linear function of the respective mole fractions. For the data in Fig. 28, the growth rate was 0.3 layers per second, x = 0.33, the (100) substrate was misoriented by less than 1 mrad, and a GaAs buffer was grown at 580°C The intensity oscillations in Fig. 28 are prior to the InGaAs growth. different from those in Fig. 27 in several respects. First, and most striking, the envelope of the diffracted intensity oscillations from the pseudomorphic film decreases more rapidly than in the case of homoepitaxy. The fast damping is followed by relatively weak but sustained oscillations. The change in the envelope appears to be different for the case of compressive versus tensile strair$jgl which we speculate has to do with a changing surface Debye-Waller factor. Bergerf67) and coworkers have suggested that the rapid decay is due to enhanced adatom nucleation, causing increased surface roughness. There is also an increase in the diffuse background between the diffracted streaks, indicating a rising point defect density. Finally, at a strain-dependent thickness, there are new diffraction features appearing as inverted V’s or chevrons, at points of the bulk diffraction beams. These are interpreted to correspond to transmission diffraction and refraction through 3D clusters on the surface.fls) For InGaAs these clusters have 114 facets.

20

30

Time in Seconds Figure 28. Intensity oscillations

during the growth of InGaAs on GaAs(lO0). In mole fraction was 0.33 and the substrate temperature was 510°C.

The


critical

of MBE

721

A simple picture that might explain this behavior is that after some thickness, the strain energy in the film is so great that misfit

dislocations

form at the interface.

The strain field around these disloca-

tions is now such that at the surface, in the vicinity of the dislocations, there is partial relaxation. Adatoms on the surface, seeking to find the least costly adsorption

site in terms of having to distort its bonding, will nucleate

in these regions, producing three-dimensional clusters. This is consistent with results on the (loo), but recent work on GaAs(l1 1),tg5t indicates that dislocation formation does not necessarily preclude strong layer-by-layer growth, as indicated by persistent InxGa,_& growth on that surface.

RHEED

intensity

oscillations

during

Though diffraction techniques are not very sensitive to the onset of dislocation formation, it is useful to plot the film thickness at which chevrons are observed versus misfit strain. This is shown in Fig. 29 for In,Ga,& on GaAs(100).[181~g6~[g8~~wl For comparison, the MatthewsBlakeslee (MB) critical thickness is plotted as the solid curve.t6s] These measurements delineate two regimes of 3D cluster formation. For those films with strain greater than 2%, the measured “critical” thickness is less than the MB prediction; for those films with strain less than 2% the critical thickness is larger than the MB value. The latter does not contradict the MB model since sluggish relaxation due to a variety of kinetic limitations can affect the measured critical thickness.tg3] For the former, Pricetg7] has included a surface energy term to account for the discrepancy. Alternatively, a different mechanism could be responsible for the relaxation and

_.__

0

100 50 Film Thickness in A

150

Figure 29. The measured thickness at which 3D features are observed in the diffraction pattern versus mole fraction of in. The solid line is a calculation of the Matthews-Blakeslee critical thickness for single kink relaxation.

722

Molecular

Beam Epitaxy

decay of the RHEED intensity oscillations, OrrflOO] and Srolovitztloll have suggested that at strains larger than about 2%, a surface instability is responsible thickness

for roughening is dependant

the surface

at some critical thickness.

on surface free energies

This

as well as on surface

kinetics. Lattice Relaxation. A more direct measurement of the generation of dislocations during lattice mismatched epitaxy is to follow the surface inplane lattice parameter during growth and when growth is interrupted. When dislocations form to relieve strain in the pseudomorphic film, the lattice constant should revert toward its unstrained, bulk value. Since RHEED senses only the last few atomic layers, in-plane lattice constants can be measured without averaging over the entire film. This was done first for the case of InGaAs growth on GaAs(l00) by Whaleyt17] and then confirmed by Berger.pOl The method relies on the most basic diffraction measurement: in any diffraction technique the separation between two beams is inversely proportional to the lattice parameter. To avoid having to know the distance from the sample to the screen and to know the energy of the beam accurately, measurements are performed relative to the GaAs substrate. Figure 30 shows the apparatus for measuring the angular separation between two diffracted beams. The positions of two beams must be measured to eliminate errors due to drift. (In principle, the position of one beam relative to the fixed specular beam could be measured.) In this method, two diffracted beams are focused on the entrance slits of two separate detectors. The slits are oriented parallel to the diffraction streaks. Then the beams are magnetically deflected across the stationary slits to obtain the two intensities versus applied field. As each beam is swept across the detectors,

a signal is measured that near each

maximum approximates

These data are then fit to a parabolic

a parabola.

function and the center determined. Knowing the deflection produced by the field, one can measure relative changes in the separation of the beams. The biggest source of error in the measurement is due to changes in shape of the diffracted beams during growth, i.e., the parabolas become broader.

Likely

because

of a combination

of multiple

scattering

and

disorder,tlo21 the beams are broadened asymmetrically, mimicking changes in lattice parameter. Though precise measurements are difficult, if sufficient care is taken it is possible to use RHEED to measure the lattice parameter to about 0.003 A.


Phosphor Screen

of MBE

723

DeteCl0rS

\

I

10 keV RHEED Gun Applied B Field sweeps both diffracted beams simultaneously.

Figure 30. Schematic of the apparatus used to measure the separation between two diffracted beams. The beams are magnetically deflected across slit apertures. The asymmetry of the beams is the limit of the measurement.

31 shows the results of measurements of the surface lattice parameter for the growth of In,Ga,_.#s with x = 0.33 for various substrate Figure

temperatures. At 51 O”C, smooth layer-by-layer growth is observed until, after about 13 layers, the lattice parameter abruptly begins to increase. At this point, one also observes chevrons in the diffraction pattern. Hence we associate 3D cluster formation with the measurable change in lattice parameter. For the measurements, care was taken to only measure the separation between diffraction streaks, away from the positions of the diffuse, bulk-like 3D chevrons. At lower temperatures the lattice relaxation is suppressed; intensity oscillations are observed to continue as wellboth of these indicate that kinetic factors limit the formation of dislocations. A surprise was that if growth were interrupted while maintaining the As, flux, then even the higher temperature

data ceased to relax.

This

further suggests that surface properties strongly affect the kinetics of the dislocation generation and motion. Current work has attempted to use adsorbates to modify the relaxation. For example, Sn which rides the surface during growth might be expected to act as a surfactant and change the critical thickness, relaxation rate, and growth mode. But no effect was observed.tq By contrast, in the growth of SiGe films, Sb has been observed to modify the transition to 3D cluster formation.t103]

724

Molecular

Beam Epitaxy

0

10

30 40 50 20 Film Thidmess in Monolayers

60

70

Figure 31. Surface lattice parameter versus measured film thickness growth of In,Ga,,As with x = 0.33 at several substrate temperatures. relaxation is suppressed at only slightly reduced growth temperatures.

5.0

SIMPLE

GROWTH

during Strain

MODELS

Our goal is to apply electron diffraction to understand the time evolution of surface structure during MBE growth. We would like to understand changes in the long range order of step trains on vicinal surfaces, the agglomeration of islands, the preferential evaporation of one of the components of an alloy, and the diffusion or migration of adatoms. The purpose of this section is to develop an intuition for the requirements of oscillatory behavior in layer-by-layer growth. For this it is useful to have some simple limits in mind. This will give two important capabilities: first, if the model is not too crude and contains some elements of the data, then one can calculate the diffraction. In our case, this will illustrate some of the difficulties of comparing calculation to measurements from real instruments. Second, one can appreciate the difficult job that many-parameter Monte Carlo calculations face in distinguishing among growth modes. In this section, simple growth models with varying amounts sion will be developed. formation and dissolution,

of surface diffu-

Unfortunately an essential ingredient, cluster has not yet been incorporated into the models.

We consider growth on a low-index surface first. As shown in Fig. 32, imagine a (100) surface with atoms or scatterers distributed in such a way as to produce clusters and clusters on top of clusters. Let the vector

RHEED Studies

of the Dynamics

of MBE

725

Figure 32. Model of growth on a low-index surface showing the layer coverages In addition, growth and interdiffusion processes are indicated.

0,.

from an origin to an atom at r to be given by r = x + ndi where d is the interlayer spacing and n is an integer. Two useful quantities are the layer coverage of the nth level, Q,,, and the exposed coverage, c,. The layer coverage is the total number of atoms on the nth level divided by the total number of sites. The exposed coverage is the fraction of sites on which there is an atom that is a topmost surface atom. A simple relation exists between them: Eq. (40)

c,, = Q,,-

Then the singly scattered only is just Eq. (41)

%+l

diffracted

amplitude

from top layer scatterers

A(S) = 1 ~(E,~)c’~~~ X.,1

where f is the atomic scattering

factor, the momentum

ki where k = 211/h is the electron wavevector,

transfer is S = kf -

E is the electron energy, and

6 is a scattering angle. Here the sum only includes those x and n that correspond to scatterers occupying top layer sites. As in Sec. 3.1, takef = 1 so that Eq. (42)

If there are N sites per level, then at SII = 0 this is further simplified[104j to yield

726

Molecular

Eq. (43)

Beam Epitaxy

A(SII = 0,SJ = N i

CJ@‘”

= N i (0, - cI,~+,) es&

it=0

/kO

The diffracted intensity I(&) (more precisely, the interference function)t105) is calculated by taking the square modulus. For a perfect instrument, this would be the peak intensity of, for example, the specular beam. It is essential to note that by performing the sum at St,= 0, one completely neglects the intensity due to the lateral distribution of adatoms. It is difficult to extract the peak intensity from a real measurement. The main modification will be that the total intensity never drops to zero. In what follows we will calculate Q,,(i) for several models of growth. The diffracted intensity given by the square modulus of Eq. (43), apart from a factor of N, will be calculated and compared to the rms roughness. With a growth rate of l/t monolayers per second, the roughness is given by:tlo4)

Eq. (44)

5.1

Perfect

A2 = ,io(n - rN2 (% - %+I)

Layer-Growth

First consider the extreme limit that every atom deposited

goes into

the topmost unfilled layer until that layer is completed. For example, for a growth rate of l/t monolayers per second, one has during the interval 0 5 t 5 t that Eq. (45)

0, = 1

0.=On>l This gives a diffracted

intensity

at S,d = J-Cof [l - (2t modt)12 and an rms

roughness of A2 = (tmodt)[l(tmodt)]. The coverages, diffracted intensity, and rms roughness are plotted in Fig. 33, with each repeated every monolayer. In the figure, the straight line segments are the coverages, g,, increasing with slope l/t. The calculated peak diffracted intensity decreases from unity to zero at half coverage and then increases back to unity.

This cusp-like

The rms roughness zero.

behavior

has been observed

is also periodic,

by Van Hove et al.n4]

being roughest when the intensity

is

Studies of the Dynamics of MBE

RHEED

.A

0

4

2

6

8

727

10

TIME (t/T) Figure 33. The expected perfect layer growth.

5.2

Nondiffusive

coverages,

diffracted

Growth on a Low-index

intensity, and rms roughness

for

Surface

A more realistic limit would be to assume that once an atom impinges onto an exposed portion of a layer, the adatom is confined to that layer.

It is not allowed to cross a boundary

model might approximately

defined by a step edge.

hold for growth at low temperature

This

or at high

rates. After a cluster is formed, a smaller cluster could grow on top, and then on top of that, until the surface becomes very rough. With these severe restrictions, a simple recursion can be derived. If the net growth rate is once again l/z monolayers per second, then the fundamental equation

describing

this growth mode is

728

Molecular

Beam Epitaxy

2= (l/t)(O,,, - O,,) e,(t) = 1

e,,(o)= 0 Here 6,,_, - 0, is the fraction of the area of the nth layer that is unfilled. solid-on-solid model[lOrl[llol

Like models, no overhangs are allowed. This is a birth-death since the growth on an unfilled layer is rapid, while the

growth of a nearly completed layer is slow. These equations solve beginning first with n = 1. The solutions are Eq. (47)

are easy to

e,(t) = 1 - e-~~

Eq. (43) or in general,

Eq. (4%

WI = 1 -

emuxji (t/#/j!

The first few solutions, for a growth rate of one monolayer per second are shown in Fig. 34. As expected, the coverages are qualitatively similar to the perfect layer-growth model, but before one layer is complete, another begins. The important change here is that neither the rms roughness nor diffracted intensity is cyclic. Instead, the intensity at SII = 0 and S, = rcld, shown as the solid line, decreases rapidly to zero. Substituting Eq. (49) into Eq. (43) gives

This falls off very quickly with the number of layers deposited, model, where

there

becomes progressively

is no inducement

t/z. In this

for layers to fill in, the surface

rougher with A2 = f/~ Notice that in all models the

intensity will decrease as in Eq. (50) during the initial nucleation, since only the first layer is being filled and there is no transfer between layers. Recently, Evans has shown that if site exclusion is included, then very weak oscillations can be obtained.


of MBE

729

Figure 34. Solution given by Eq. (49) of coverages with growth without diffusion.

5.3

Diffusive

Growth on a Low-index

Surface

Adding the possibility that an atom can jump to a lower level gives a model intermediate between perfect layer-growth and non-diffusive growth. Schematically

Eq. ( 51)

the differential

>=

(l/t)

equations

would be

(O,r_,- O,,) t (jumps from ntl

to n)

- (jumps from n to n-l) A variety of schemes

could be used for the last two terms.tlOr)tlloj

One

could also allow jumping, with different probabilities, from lower to upper levels. In keeping with the birth-death approach, we assume a jump rate in going from, for example, ntl to n that is proportional to the product of the available space on level n and the uncovered area on level ntl. The reasoning edge.

is that once an atom is covered,

Then Eq. (51) becomes

it is not able to diffuse to an

730

Molecular

Eq-(52)

Beam Epitaxy

2 =(l/4(%-1- %I + WA,+1

- %+2)FLl

- WJ,

The coverages,

diffracted

- %+1)&L2

intensity,

- %I

- %-,)

and rms roughness

of this particular

nonlinear model are remarkably easy to evaluate numerically, subject to the conditions 0,(t) = 1 and O,,(O)= 0. The results are shown in Fig. 35 with k = 100. Several points are worth noting. First, the calculated peak intensity drops rapidly as soon as growth begins. This is similar to what is observed experimentally, though there, one expects reconstruction changes and perhaps multiple scattering to be more important. Second, there is a phase shift so that, though the period of the intensity is t, the peak positions do not correspond to integral average film thicknesses. Third, both the rms roughness and diffracted intensity continue undamped, more similar to Ge or Si(100) than GaAs on GaAs(lO0). Finally, as expected, the coverage plots are intermediate between non-diffusive growth and perfect layer-growth.

g

0.6

d y

0.4

E

0.0

012

012

3

3

4

4

5

5

6

6

7

7

6

6

910

910

TIME (t/T)

Figure 35. Solution of Eq. (52) with k = 100. The curves are intermediate between perfect layer growth and non-diffusive growth. The slopes of the coverages are just the growth rates. intensty oscillations.

How the layers are completed

determines

the form of the


5.4

Diffusive

of MBE

731

Growth on a Vicinal Surface

It is difficult to apply the birth-death because there is no distinction

approach to a staircase of steps

between

the statistics

of different

Instead we examine the solution to a one-dimensional

diffusion

levels.

equation

for the simple case in which there are no clusters. Parts of the analysis should carry forward to that more complicated case. The main result here is to see how moving steps modify the classic Burton-Cabrera-Frank solution. The solution to this differential equation will show that step velocity oscillations are associated with RHEED oscillations and that the envelope shown in Fig. 17 arises in a natural way. As a model appropriate to vicinal surfaces, consider a one-dimensional staircase with equal terrace lengths. Describe the adatom concentration on a terrace by the probability +) that a site x is covered by a mobile adatom. Here, x is measured with some origin that is fixed with respect to the bulk. We assume no clustering and no evaporation. Further we assume that the upper and lower edges of terraces are perfect sinks for mobile atoms and we assume periodic boundary conditions. In the fixed reference frame, the diffusion equation is an&, t) -

Eq. (53)

= D a’n(xJ) + 3Xx’

at

1 (1%

where D is the diffusion coefficient of inter-row separation, and t is the time This equation says that adatoms arrive and diffuse to the top and bottom step

non-interacting adatoms, a is the required to deposit a monolayer. at the surface at a prescribed rate boundaries. A gradient of mobile

adatoms is set up, thus driving the diffusion.

The steps are moving at the

same rate, by periodicity, so that the terrace, which has lengthl, is moving across the surface with a velocity v(f). In order to apply the boundary conditions that the mobile adatom concentration

be zero at the steps edges,

change to a reference frame that is moving at the step velocity, v(t), via

II =x - s ;(9)d9

Eq. (54)

so that Eq. (53) becomes Eq. (55)

-

an(U)

at

= D a2n(M

-t

an2

a2n(u,t)

v(t) -

au

t-

1

T-n

732


In the moving reference frame this can be solved subject to the boundary condition that n(u,t) = 0 at u = 0~5. For the numerical solution, one uses - an(L,t)/au]. The resulting solution is a distribution, W) = aD[Ch(O,t)/au n(u,t), of mobile adatoms that is skewed by the moving step. At t = 0, n(u,O) starts at zero and then increases symmetrically across the step, keeping the boundaries fixed at zero. Since n(u,t) = 0 at the step edges, there is a gradient that feeds the steps, causing them to move. As the steps move, there is a pile up of mobile adatoms near u = 0 and a low adatom density near u = L, where new surface is being created. This continues until a steady state is reached, which can be solved exactly.[106] The steady state distribution, for moderate diffusion, approaches l/u near u = 0 and falls linearly to zero at u = L. At this point, convective motion of the steps is balanced by surface diffusion of the mobile adatoms.[108] The Fourier transform of this adatom distribution n(u,c-~) at the appropriate diffraction geometry gives the steady state intensity. To arrive at this value, the distribution n(u,Q undergoes a transient behavior that gives intensity oscillations. Figure 36 illustrates the solution to Eq. (55) forl*/Dt = 20. Initially n(u,t) is relatively flat. At a later time, it is skewed toward u = 0, oscillating about its final steady state value.

Figure 36. (a) Time variation of the position dependence of the density of mobile adatoms on a terrace of a vicinal sur-

0 I

nnsitinn

nn

terrace

0 0

time (s)

20

face. (b) The velocity of a step calculated in the modified BCF The velocity exhibits model. oscillations since a step edge will move very quickly in regions of high step density (or islands); then slow after the adatoms are used up; finally, speed up when the step density recovers. Here L2 lDc = 20.

RHEED Studies

of the Dynamics

of MBE

733

To determine the diffracted intensity, write Eq. (7) with the scattering amplitude of a terrace written explicitly as a Fourier transform. We use a continuum which

approximation

corresponds

and evaluate the result at S, = x/d and S, = &.,

to the peak of one of the components

of the split

specular beam. We add the scattering between two layers separated by step height d by using n(u,i) as the probability there is a scatterer at u on the top level and 1 - n(u,f) as the probability there is a scatterer at u on the lower level. At this scattering geometry, the diffracted intensity is

Eq. (56)

I = I_/$(u,t)&“du

+l6r”l -n(u,t)]BS~” - iSzddU I2

which using the values of S, and S, becomes

I = 1-s;; -2n(u,t)]ei”~Ldu

Eq. (57)

I*

The resulting intensity forL*/Dc = 45 is shown in Fig. 37. As can be seen in this figure, like experiment, the diffracted intensity starts at an initial value and decreases rapidly depending upon the dimensionless parameter L*/Dc. In contrast to the low-index model discussed earlier, here the maxima fall and the minima rise until a nonzero steady state value is reached. Like our experimental results, the calculated oscillations are independent of whether the incident beam is directed parallel to the step edges or down the staircase. However, as discussed in Sec. 2, because of the asymmetric find a situation direction

instrumental response of RHEED, it might be possible to in which amplitudes from each terrace are added in one

and intensities

in the other.

In this case, one expects

different

results. One result that has yet to be confirmed is the prediction that, at low values of L*/Dc, the period of the intensity oscillations is slightly longer than the expected

monolayer

completion

time.

This occurs because the

step edges rob the mobile adatoms on the terraces of some fraction of the incident flux. Proper inclusion of clustering might eliminate this shift, and as of yet we have not been able to separate this small effect from nonuniformities of the incident flux.

734


Figure 37. Calculated diffracted intensity on a vicinal surface using the modified BCF solution with L*lL?c = 45. Though no cluster growth is incorporated, this calculation mimics many features observed in experiment.

In growth

on the vicinal

surface,

the intensity

oscillations

are seen to

be associated with periodic oscillations of the step velocity. Similar conclusions are obtained from Monte Carlo calculation&‘*1 Just before an intensity maxima, the adatom concentration on a terrace is such that the step velocity begins to increase dramatically. The steps move rapidly across the surface, creating something approaching the starting staircase and giving maximum diffracted intensity. After the essentially fresh surface is created, the step velocity slows until the next period. The oscillations reach steady state when the profile is able to feed the steps without these transient constant

velocity

readjustments.

of L/t.

If clustering

At this point, the steps move with a is included,

then one expects

similar enhancement

of the cluster density

near the trailing

reduction

of cluster

density

created

clustering

is significant,

on the newly

then the low-index

description of the growth process. Like the low-index model,

the growth

part of terrace.

methods term

a

edge and a If

may be a better

in Eq. (55) can be

removed and the recovery of the intensity calculated. From this model there is no evidence of the two-process form postulated by Joyce.tlog] Unlike the low-index model, the intensity always recovers to the starting value because calculation.

the steps that act as perfect sinks are included

in the

RHEED Studies

6.0

of the Dynamics

of MBE

735

CONCLUSION Reflection

the microscopic

high-energy

electron

diffraction

is sensitive

processes that occur during crystal growth.

beams broaden and the peak exhibits

intensity variations

to many of

The diffracted due to changes

in the defect and morphological structure of the surface during epitaxy. Contained in these changes are the microscopic processes of surface diffusion, island growth, dislocation formation, step bunching, and step disorder. These changes have been analyzed in terms of kinematic diffraction theories that neglect of multiple scattering between steps, and which assume that the scattering from isolated islands is essentially identical to the underlying substrate. Simple models can easily describe the main features of the data. In this incomplete discussion we presented rate equation methods. The main point is that since it doesn’t take too much to obtain qualitative agreement, the mere existence of intensity oscillations does not need to indicate the reliability of a model. More realistic models are, of course, needed, though from our perspective

the analysis of the diffraction

data should improve in

parallel. Otherwise the many more parameters and interactions that the models could include would be lost in the comparison to the data. As described in the discussion of the broadening of the diffracted beams for both low-index and vicinal surfaces, a major advantage of diffraction over direct imaging is that the statistics of the island structure are measured. But the complexity of non-equilibrium crystal growth processes, as well as uncertainties of the role of two dimensions and multiple scattering, limit the quantitative impact that has so far been felt. Despite progress in understanding the diffraction patterns and its evolution during crystal growth, the subject needs the combined attack of more than one in-situ technique. Ways to apply scanning tunnelling microscopy, lowenergy electron

microscopy,

or reflection

electron

microscopy

are sorely

needed. Nonetheless, there is still much one can say. With advances in analysis and direct imaging to calibrate the diffraction pattern, more definitive answers to fundamental

questions of epitaxy should be forthcoming.

ACKNOWLEDGMENTS This work was partially supported by grants from the National Science Foundation (DMR-93-07852) and from the Air Force Office of Scientific

Research

(AFOSR F49620-93-l-0080).

736

Molecular

APPENDIX:

Beam Epitaxy

TWO-LEVEL

DIFFRACTION

To calculate the diffraction from a surface with a random distribution of islands or two-dimensional

clusters,

the simple sums of the previous

sections no longer suffice. A correlation function approach must be used. This approach has been treated exhaustively elsewhere.t34) In this discussion we will give the simplest, “back of the envelope” treatment of a very special case and then just state how it is generalized. Hopefully, the generality of the method won’t go unnoticed. The two main goals at this point are (7) to obtain the kinematic angular dependence of the intensity oscillations and (2) to describe how to extract the calculated peak intensity of Sec. 3.2. We treat the case of scatterers distributed among two levels in onedimension.

The kinematic

I(S) =

intensity

1: n(r) Pj

from this surface is

2

where n(x,z) is equal to 1 if there is a surface scatterer at coordinates (x,.z) and 0 otherwise, and S = Sjt S>. It is important to note that though this sum is over top layer scatterers only, the same column approximation arguments of Sec. 3.1 apply, so that the dynamically scattered intensity from a cluster and all of the atoms underneath it are included. Since the clusters are, apart from lateral size, the same, we just need to include the path length difference of the origins. If there are No total lattice sites in a layer, then this can be rewritten as

Eq-(5%

I(S) = 2

P’C n(r)

r

z emis"'n(r') =No r’

where we define the correlation

Eq. (60)

1 G+“C(u) ”

function

C(u) = (l/N,) T”(r) n(r t u)

The correlation

function

C(u) is the probability

of finding two scatterers

on

the surface separated by a vector U. Our procedure is to evaluate the correlation function on the surface and then take its Fourier transform to obtain the diffracted intensity.


of MBE

737

To make this simpler, we use the result that it is sufficientf34) to transform Eq. (59) to an integral overx, convolving the result with the onedimensional reciprocal lattice. In the discrete case the correlation function C(u) = C(na, Id) where 1 = -1, 0, 1. Instead we now write na = x but retain the discrete two levels. Our method is to break C(.~,ld) into partial correlatron functrons, C4,4+,(x), which are straightforward to calculate. Following

Eq.(61)

the definition

in Eq. (60),

%z =Id) = 24 Cq,q+lGy)

x 1 is whereCq.4+r(

the probability

that there is a scatterer

at the origin on

level g and also on level q + 1 at a distance x away. These continuum partial correlation functions will now be dealt with for a very simple case. Following the notation of Sec. 5, let the exposed coverage of the top level be c so that the exposed coverage of the second level is 1 - c. Next, let f(x) be the probability that there is a top layer scatterer at X, hence the probability that there is a second layer scatterer at x is 1 - f(x). We then make a Markov assumption to give the probability that there is a step in going from x to x + a in terms of the jump probabilities pd and pu as

Eq. (62)

f&+4 = (1 - PdfQ

Since we only want the continuum

Eq. (63)

+ PlJD - f(41 distribution,

expand about x to obtain

@d+Pkf(X)-Pu=o

a$+

To find C,,(x), we calculate the probability that there is an atom anywhere on level one, i.e., c, times the probability that, given an atom at the origin, there is onex away. This second factor is f(x) given by the equation above subject to the boundary Eq. (64)

conditions

f(w) = c

The solution is, applying some symmetry, f(.~) = (1 - c) exp(-hjxj) + c, where h = (pd + pJ/a and c = p,/(pd + pJ, Thus the first partial correlation function is

738

Molecular

Eq.(65) Similarly

Beam Epitaxy

= c(1

C,,(x)

- c) e Wi + c2

C,,(x) is found by finding the probability

that given no atom at the

origin on the first level that there is an atom at x on the first level. result is (1 - c)f(x) boundary

The

where now f(x) is a solution to Eq. (63) subject to the

conditions

that

f(m) = c

Eq. (66)

f(O) = 0 This gives C,,(x) = (1 - c) c(1 - e -+ Then because there is no distinction between left and right on a two-level surface, C,,(x) = C,,(-X) = C,,(x). Finally, C,,(x) is found similarly to C,, to be

Eq. (68)

C**(x)

= C(1 - C) emAlxl+ (1

- C)*

Putting these partial correlation functions together in Eq. (61), one obtains the actual correlation function to be C(x,O) = C,, + C,, and C(x,l) = C,, , so that

4. (69)

C(x,O) = c, +(l - c)* + 2c(l

- c) e-hlxl

C(x, f 1) = c(1 - c) (1 - emhbI) These correlation

functions

could be compared to Eq. (26) of Ref. 34.

At last we can calculate tion function

Eq. (71)

the diffraction

profile.

Putting the correla-

into Eq. (59) we have

f(S) = Ldx emisxx C(x,O) t 2 cos SJJIcfx

e-is,XC(x,l)

RHEED Studies

Eq. (72)

r(s.&)

of the Dynamics

of MBE

739

= [c? + (1 - cy + 2c(l - cj cosSzdj2n6(Sx) •t 2c(l

- c)(l - coss,d) [2h/(h2 t S,2)]

This last equation says that the diffracted beam versus S, can be separated into a broad part and a central spike. Note that the delta function is broadened by the transfer width of a real instrument. Further since, for example, the specular beam, S, = 2,40, the relative size of these two components will depend upon the angle of incidence, fib When S,d = 2x, an in-phase angle, the second term will vanish, indicating that the diffraction is insensitive to steps. Similarly, at half coverage and at S,d = n, the first term vanishes. However, and in contrast to our calculation in Sec. 5, if there is step disorder, the total intensity

does not drop to zero!

REFERENCES 1. Jamison, K. D., Zhou, D. N. Cohen, P. I. Zhao, T. C., and Tong, S. Y. J. Vat. Sci. Technol., A6:611 (1988) 2. Tong, S. Y., Proc. of the NATO Conference Phys. Left., A128:447 (1988)

on REM and RHEED,

3. Peng, L. -M. and Cowley, J. M., Surf. Sci., 201:559 (1988) 4. Marten, H. and Meyer-Ehmsen,

G., Surf. Sci., 151:570 (1985)

5. Ichimiya,

A., Jpn. J. Appl. Phys., 22:176 (1983)

6. Maksym,

P. A. and Beeby, J. L., Surf. Sci., 1 lo:423

7. Tanishiro, (1983)

Y., Takayanagi,

(1981)

K., and Yagi, K., Ultramicroscopy,

11:95

8. Hsu, T., lijima, S., and Cowley, J. M., Surf. Sci., 137:551 (1984) 9. Pashley, M. D., Haberern, Technol., B6: 1468 (1988) 10

K. W., and Woodall,

J. M., J. Vat. Sci.

Telieps, W., and Baier. E., Surf. Sci., 200:512 (1988)

11. DeMiguel, J. J., Cebollada, A., Gallego, J. M., Ferrer, S., Miranda, R., Schneider, C. M., Bressler, P., Garbe, J., Bethke, K., Kirschner, J., Surf. Sci., 211:732 (1989) 12. Wang, G. C. and Lu, T. -M., Surf Sci., 122:L365

(1982)

13. Harris, J. J. and Joyce, B. A., Surf Sci. Left., 108:L90 (1981)

740

Molecular

Beam Epitaxy

14. Van Hove, J. M., Lent, C. S., Pukite, P. R., and Cohen, P. I., J. Vat. Sci. Technol., B1:741 (1983) 15. Van Hove, J. M., Cohen, P. I., and Lent, C. S., J. kc. Al 546 (1983) 16. Dobson, P. J., Neave, J. H., and Joyce, 119:L339 (1982)

B. A.,

Sci. Technol., Surf. Sci. Leti.,

17. Whaley, G. J. and Cohen, P. I., J. Vat. Sci. Technol., B6:625 (1988) 18. Whaley, G. J. and Cohen, P. I., Proc. Materials 160:35 (1990)

Research

19. Bolger, B. and Larsen, P. K., Rev. Sci. Instrumen.,

Society,

57:1363 (1986)

20. Chalmers, S. A., Gossard, A. C., Petroff, P. M., Gaines, Kroemer, H., J. Vat. Sci. Technol., 7:1357 (1989)

J. M.,

21. Resh, J. S., Jamison, K. D., Strozier, J., Ignatiev, A., Rev. Sci. Instr., 61:771 (1990) 22. Lewis, B. F., Grunthaner, F. J., Madhukar, A., Lee, T. C., and Fernandez, Ft., J. Vat. Sci. Technol., B3:1317 (1985) 23. Henzler, M., Advances

in So/id State Physics, 19:193 (1979)

24. Lagally, M. G., Methods of Experimental Park and M. G. Lagally, eds.), Academic

Physics: Surfaces, (FL L. Press, Orlando (1985)

25. Zhang, J., Neave, J. H., Joyce, B. A., Dobson, P. J., and Fawcett, P. N., Surf. Sci., 231:379 (1990) 26. Larsen, P. K., Dobson, P. J., Neave, J. H., Joyce, B. A., Bolger, B., and Zhang , J., Surf. Sci., 169:176 (1986) 27. Joyce, B. A., Neave, J. H., Zhang, J., Dobson, P. J., Dawson, P., Moore, K. J., and Foxon, C. T., NATO Adv. Sci. Inst., B163:19 (1986) 28. Shitara, T., Neave, 14:1658 (1993)

J. H., and Joyce,

B. A.,

Appl.

Phys.

Leti.,

29. Peng, L. -M. and Cowley, J. M., Surf. Sci., 199:609 (1988) 30. Korte, U. and Meyer-Ehmsen, 31. Kawamura,

T. and Maksysm,

G., Surf. Sci., (1993), in press. P. A., Surf. Sci., 161:12 (1985)

32. Pukite, P. R., Lent, C. S., and Cohen, P. I., Surf. Sci., 161:39 (1985) 33. Kittel, C., Introd. to So/id State Phys., 3:49, John Wiley (1968) 34. Lent, C. S. and Cohen, P. I., Surf. Sci., 139:121 (1984) 35. Cohen, P. I. and Pukite, P. R., Ultramicroscopy,

26:143 (1988)


of MBE

741

36. Fuchs, J., Van Hove, J. M., Wowchak, A. M., Whaley, G. J., and Cohen, P. I., Layered Structures, Epifaxy, and Interfaces, (J. M. Gibson and L. R. Dawson, eds.), 37:431, Materials Research Society (1985) 37. Horn, M., Gotter, U., and Henzler, M., J. Vat. Sci. Technol., B6:727 (1988) 38. Pukite, P. R. and Cohen, P. I., J. Cryst. Growth, 81:214 (1987) 39. Kaplan, Ft., Surf. Sci., 93:145 (1980) 40. Wood, E. A., Crystal Orientation Manual, p. 24, Columbia University, New York (1963) 41. Report from the American Designation F26-76

Society

for Testing

42. Pukite, P. R., Van Hove, J. M., and Cohen, Tech&., B2:243 (1984)

and

Materials,

P. I., J. Vat.

Sci.

43. Pukite, P. R. and Cohen, P. I., Appl. Phys. Left., 50:1739 (1987) 44. Purcell, S. T., Arrott, A. S., and Heinrich, B6:794 (1988)

B., J. Vat. Sci. Technol.,

45. Egelhoff, W. F. and Jacob, I., Phys. Rev. Lett., 62:921 (1988) 46. Aarts, J. and Larsen, P. K., Surf. Sci., 188:391 (1987) 47. Sakamoto, T., Kawamura, T., Nago, S., Hashigushi, G., Sakamoto, K., and Kunujoshi, K., J. Cryst. Growth, 81:59 (1987) 48. Pukite, P. R., Batra, S., and Cohen, P. I., Difiusion at Interfaces: Microscopic Concepts, (M. Grunze, H. J. Kreuzer, and J. J. Weimer, eds.), p 1936, Springer-Verlag, New York (1988) 49. Maksym, P. A., Presented at the International Conference on Crystal Growth, Sendai, Japan (1989); Maksym, P. A., Semicond. Science and Techno/., 3: 594 (1988) 50. Arthur, J. R., Surf. Sci., 43:449 (1974) 51. Lehmpfuyl, (1991) 52. Stoyanov,

G., Ichimiya, A., and Nakahara, S. and Michailov,

53. Van Hove, J. M., Pukite, Technol., B3:563 (1985)

H., Surf. Sci., 245:Ll59

M., Surf. Sci., 202:109 P. R., and Cohen,

(1988)

P. I., J. Vat.

Sci.

54. Johnson, M. D., Sudinono, J., Hunt, A. W., Snyder, C. W., and Orr, B. G., Surf. Sci., (1993), in press. 55. Horio, Y. and Ichimiya, 56. Henzler,

A., Surf. Sci., 133:393 (1983)

M., private communication,

(1987)

742

Molecular

Beam Epitaxy

57. Horio, Y. and Ichimiya, A., Surf Sci., (1993), in press. 58. Van Hove, J. M. and Cohen, P. I., Appl. Phys. Lett., 47:726 (1985) 59. Kojima, T., Kawai, N. J., Nakagawa, T., Ohta, K., Sakamoto, T., and Kawashima, M., Appl. Phys. Left., 47:286 (1985) 60. Newstead, S. M., Kubiak, 81:49 (1987) 61. Heckingbottom,

R. A., Parker, E. H., J. 0yst.

Growth,

Ft., J. Vat. Sci. Technol., B3:572 (1985)

62. Neave, J. H., Joyce, B. A., and Dobson, P. J., Appl. Phys., A34:179 (1984) 63. Lewis, B. F., Fernandez, R., Madhukar, A., and Grunthaner, Vat. Sci. Technol., 4:560 (1986)

F. J., J.

64. Singh, J. and Bajaj, K. K., J. Vat. Sci. Technol., 2: 276 (1984) 65. Kaspi, R. and Barnett, S. A., J. Vat. Sci. Technol., in press (1990) 66. Tsao, J. Y., private communication,

(1991)

67. Berger, P. R. and Singh, J., J. Vat. Sci. Technol., B5:1162 (1987) 68. Matthews, (1977)

J. W. and Blakeslee,

A. E., J. Vat. Sci. Technol., 14:989

69. Lievin, J. L. and Fonstad, C. G. Appl. fhys. 70. Berger, P. R., Chang, K., Bhattacharya, Sci. Technol., B5:1162 (1987)

Left., 51 :1173 (1987)

P. K., and Singh, J., J. Vat.

71. Neave, J. H., Dobson, P. J., and Zhang, J., Appl. Phys. Lett., 47:lOO (1985) 72. Clarke, S. and Vvedensky,

D. D., Appl. Phys. Lett., 51:340 (1987)

73. Singh, J. and Bajaj, K. K., J, Vat. Sci. Technol., A6:2022 (1988) 74. Van Hove, J. M. and Cohen, P. I., J. Cryst. Growth, 81:13 (1987) 75. Madhukar, A. and Ghaisas, S. V., CRC Crit. Rev. Sol. State Sciences, 14:l (1988) 76. Tringides,

M. C., J. Chem. Phys., 92:2077 (1990)

77. Burton, W. K., Cabrera, N., and Frank, F. C., Phil. Trans. Royal Sot., 243A:299, London (1950) 78. Nishinaga, T., and Cho, K. -I., Jpn. J. Appl. Phys., 27:L12 (1988); Nishinaga, T., Shitara, T., Mochizuki, K., and Cho, K. -I., J. Cry&. Growth, 99:482 (1990) 79. Suzuki, T. and Nishinaga,

T., J. Cryst. Growth, 111:173 (1991)

RHEED Studies

80. Horikoshi, Y., Kawashima, Phys., 27: 169 (1988)

of the Dynamics

M., and Yamaguchi,

of MBE

743

H., Jpn. J. Appl.

81. Shitara, T. and Joyce, 8. A., J. Appl. Phys., 71:4299 (1992) 82. Irisawa, (1990)

T., Arima, Y., and Kuroda,

T., J. Cryst. Growth,

99:491

83. Tsui, Ft. K., Curless, J. A., Kramer, G. D., Peffley, M. S., and Rode, D. L., J. Appl. Phys., 58:2570 (1985) 84. Schwoebel,

R. L. and Shipsey, E. J., J. Appl. Phys., 37:3682 (1967)

85. Saluja, D., Pukite, P. R., Batra, S., and Cohen, Technol., 85710 (1987)

P. I., J, Vat. Sci,

86. Cohen, P. I., Petrich, G. S., Dabiran, A. M., and Pukite, P. R., Kinetics of Ordering and Growth at Surfaces, (M. G. Lagally, ed.), p. 225, Plenum, New York (1990) 87. Radulescu, D. C., Wicks, G. W., Schaff, W. J., Calawa, Eastman, L. F., J. Appl. fhys., 63:5115 (1988)

A. R., and

88. Bennema, P. and Gilmer, G. H., Crystal Growth: An Introduction, Hartman, ed.), Ch. 10, North Holland, (1973)

(P.

89. Tokura, Y., Saito, H., and Fukui, T., J. Cryst. Growth, 94:46 (1989) 90. Pukite, P. R., Van Hove, J. M., and Cohen, P. I., Appl. Phys. Leti., 44:251 (1986) 91. Pashley, M. D., Haberern, Left., 58:406 (1991)

K. W., and Gaines,

J. M., Appl. Phys.

92. Bean, J. C., Feldman, L. C., Fiory, A. T., Nakahara, S., and Robinson, I. K., J. Vat. Sci. Technol., A2:436 (1984) 93. Dodson, B. W. and Tsao, J. Y., Appl. Phys. Leti., 51:1325 (1987) 94. Berger, P. R., Chang, K., Bhattacharya, Phys. Lett., 53:684 (1988) 95. Dabiran, (1991)

A. M. and Cohen,

P., and Singh,

J., Appl.

P. I., J. Vat. Sci. Technol.,

B9:2150

96. Whaley, G. J., and Cohen, P. I., Bull. Am. Phys. Sot., 31:52 (1986) 97. Price, G. L. and Usher, B. F., Appl. Phys. Let,

55:1984 (1989)

98. Nakao, H. and Yao, T., Jpn. J. Appl. Phys., 28:L352

(1989)

99. Miki, K., Sakamoto, K., and Sakamoto, T., Proc. Mater. Res. Sot., “Chemistry and Defects in Semiconductors,” p. 323 (1989) 100. Snyder, C. and Orr, B. G., froc. 101. Srolovitz,

Mater. Res Sot., p. 489 (1991)

D. J. and Hirth, J. P., Surf. Sci., 255111

(1991)

744


102. Whaley, G. J., Ph.D. dissertation,

University

of Minnesota

(1989),

unpublished. 103. Copel, M. and Tromp, R. M., Phys. Rev. Leff., 63:632 (1989) 104. Horn, M., Gotter, U., and Henzler, M., Proc. NATO Workshop RHEED and REM, Eindhoven (1987)

on

105. Webb, M. B. and Lagally, M. G., So/id State Physics, 28:301 (1973) 106. Pukite, P. R., Ph.D. dissertation, (1988) unpublished 107. Arrott, A. S., proceedings

University of Minnesota,

of the Denver magnetism

Minneapolis

meeting, (1987)

108. Ghez, R. and lyer, S. S., IBM J. Res. Dev., 32:804 (1988) 109. Joyce, B. A., Zhang, J., Neave, Phys., A-45:255 (1988)

J. H., and Dobson,

110. Cohen, P. I., Petrich, G. S., Pukite, P. R., Whaley, S., Surf. Sci., 216:222 (1989) 111. Evans, J. W., Phys. Rev., B39:5655

(1989)

P. J., Appl.

G. J., Arrott, A.

Index

A

AF couplings

Absorption edge 370 in fiberoptics loss

429

Ag is a suitable prelayer AS/Fe [OOl] oriented 642 Aharonov-Bohm effect AlAs alloy 170 aluminum on 153 ALE 590

277

321

saturation 416 Abundance ratio 55 Acceptor levels 433 Accessory equipment 92 Accommodation coefficient for As, and P, 302 for dimers 280 Activation efficiency 57 Activation energy 178 of gallium surface diffusion Adatom density 708 diffusion 689 migration length 147 population 704 Adlayers and LRO 469 such as Sb or As 459 surface modification 475 Adsorbed gas contamination of UHV 11 AES 348

628

116

4X2&%74As

deep levels 194 AlGaAs 135, 163 buffer 179 growth rate 208 impurity and defect states 241 morphology 159, 170 optical properties 197, 204 roughness 171 temperature dependence 197 ternary alloys 170 traps 194 AI,Ga, ,As affinity for impurities 123 GSMBE 39 AIGaAslGaAs 275 AIGaAs/GaAs/AIGaAs quantum wells 206

153

745

746

index

Alkaline earth species 531 Alloys Si , _xGe, 454 SnGe 453 Alumina crucibles 556 Aluminum crucibles 130 desorption 146 migration of 153 mobility 170 mole fraction 194 Amphoteric doping 392 Analysis in-situ 532, 590, 624 techniques 532, 670 Analytical equipment 83 Angle of incidence 671, 679, Angular dependence 683 Anisotropic optical properties Anisotropy for Fe/Ag interfaces 642 magnetic 644 Anomalous outdiffusion 132 Antiferromagnetic 425 ordering 429 Antimony 209, 210 adlayer 475 cracker 383 passivation 214 saturation of the Si surface Antiphase mistakes 672 APCVD 462 APD 318 Applications analog circuit 484 of artificially-layered structures 623 Ill-V devices 115 high speed digital 114 microwave 114 optical waveguides 115 optoelectronic 114, 491 SiBi, .xGe, 492 of Si,+,Gex 481 Argon peak 219 Arrhenius-type rate equations

682 420

473

391

Arsenic 134, 279 chemisorption 146 cracking temperature 131 desorption 205 oxides 127, 175 passivation of GaAs 213 precipitation 178 species 179, 181 Arsenic-deficient surface 170 Arsenic-exposed edges terraces with 162 Arsenic-stabilized surface 201 Arsenic-terminated 146, 152 Arsine 134 Artificially-layered structures 623 As,/Ga ratio 199, 201 AsJGa 205 ratio 162 AsGa antisite 178 ASH, gas handling systems 290 and PH, 284 As, to P, ratio 302 Atomic epitaxy 359 fluorescence 553 growth 115 layering 507, 527 ordering 307, 468 oxygen 546 positions in a reconstructed surface 675 scattering factor 574 substitution 514 surface 138, 670 Atomic absorption spectroscopy 553, 590 Atomically abrupt 2 Si/Si, _xGe, interfaces 459 Atomically flat interfaces 117 Auger Electron Spectroscopy (AES) 88, 158 Auger spectroscopy 672 Autodoping of CdTe/lnSb 383 Automated process control 102 Avalanche photodetectors 316

Index

6 Backgating 178 Background carbon concentration 185 impurities 184, 201 sulfur donors 206 (Ba,K)BiO, 510, 572 Josephson junctions 581 thermodynamic oxygen stability limit 539 Bakeout 121, 218 Baking growth chamber 125 Ballistic transport 116 Band energy shifts 379 Band engineering 466 Band exchange coefficients 366 Band offset 366, 379 determination 362 for the CdTe 362 in semiconductor heterojunctions 368 ratio 366 Band splitting 394 Bandgap change with Ge content 454 energy 382 energy and Ge content 480 excitonic 363 of Si,,Ge, alloys 480 485 Si and Si,,Ge, Bandgap-engineered devices 116 Bandstructure effect of strain on 481 Ba(Pb,Bi)O, 510 Barrier 378 bandgap 370 formation of dislocations 463 heights 362 layers 562 Basal plane epitaxy of Pt on sapphire 632 Bayard-Alpert gauge 77, 83 BCF formulation 703 Beam Equivalent Pressure (BEP) 200, 301

747

Beam flux of As, and P, 280 distribution 31 Beam interruptors 60 Beam shape kinematic analysis of 680 Bearings in UHV 62 Beats 535, 688, 693 Beryllium 187 acceptor concentration 202 diffusion coefficient 188 dopants 187 doping 298 interstitials 164 Bi-Sr-Ca-Cu-0 517 Bis.rsSrr .s&ao.&u2%1 512 Bi2Sr2Ca2Cu,0,, 514 Bi,Sr,CaCu,O, 512 Bi,Sr,Ca,_,Cu,O,,+, 538, 567, 571 synthesis of 580 Bi,Sr&u,O, 582 Biatomic layer 307 Biexciton 416 formation 344 Binary superlattice CdTelZnTe 376 Bipolar transistors 330, 336 Biquadratic coupling 643 Bismuth oxidation of 549 Bistable switches 414 Blue emission 378 Blue exciton decay 411 recombination 431 Blue light emitting diodes 387 Blue resonances 410 Blue semiconductor lasers 433 Blue shift 420 W’3

spitting of 556 Bohr diameter 414 exciton volume 372 magnetons 425 orbit 431

748

Index

Bond formation 137 length 307 Bonding epitaxial 358 tetrahedral 350 Boron contamination 556 difficult to evaporate 56 Boron doping causes highly ordered 477 of growing surface 475 Boron evaporation crucibles for 35 Bound excitons 389 Bragg angles 147 Bragg scattering condition 86 Bridgman 69, 346 Brillouin light scattering 643 Brillouin zone 327, 373, 480 Broad part 690 Broadening of diffracted beams 675 Bromine:methanol etch 173 Buffer layers 169, 177, 387, 563 CdTe 359 grown by MEE 181 Buffers compositionally-graded 179 low temperature 177 Building layers 509 Bulk carrier mobilities 184 Bulk chemical synthesis 507 Bulk crystal 137 Bulk diffusion coefficients 528 Bulk synthesis 522, 564, 573 of high T, 522 Burgers vector 356

C c-Axis length 578 Calcium concentration 185 Calibration eutectic 348 films 554

of flux 77 of thermocouple or pyrometer 73 Cap layers 213 Capacitance manometer (CM) 41 Capacitance-voltage (CV) profiles 322 Capillary needle ion source 60 Capture pumps for UHV 16 Carbon 189 acceptor 125, 185, 206 amphoteric 283 contamination 129, 173 doping 187, 283 filaments 189 impurities 175 incorporation 201 Carrier concentration 393, 514 concentration in CuO, 512 confinement 407, 435 mobility 228, 383 transport properties 334 CAs acceptor 198 Catalytic baffles 131 decomposition 43 Cathodoluminescence 160, 233 CaTiO, 508 Cation mixing 160, 511, 555 (Cd,Mn)Te 359 barrier layers 369 interface 371 MQW 370 quantum wells 361 Cdo.55Mno.4sTe 346 Cd,,Mn,,,Te 360 Cd,,Mn,Te 358 Cd:Te flux ratio 384 CdTe 346 doped with Sb 375 epitaxial 354 heteroepitaxy of 347 on InSb 383 two orientations of 347 superlattice cell 352 CdTe-GaAs interface 352, 356

Index

CdTe/lnSb MQW 384 CdTe/MnTe SQW 380 CdTe/ZnTe binary superlattice 376 Cdo.s$‘o.,oTe 349 Central cell corrections 237 Central spike 690, 691 Chain layer 566 Chalcogens group VI 53 Chamber bake 215 Channelling of He+ ions 459 Characteristic lengths 190 Characterization techniques 118, 224, 393 vacuum in-situ 508 Charge counting 512 reservoirs 512 Charge-coupled devices 117, 183 Charge sheet density 228 Chemical beam epitaxy (CBE) 589 Chemical bonding model 350 Chemical etching for in-situ cleaning of GaAs 172 Chemical ordering in Co/Pt multilayers 647 in magnetic multilayers 651 Chemical potentials 698 Chemical vapor deposition (CVD) of Si/Si,,Ge, 460 Chemisorption 146 Chemistry of growth 137 Chevron 720, 723 Chlorine 173 Chlorofluorocarbons 220 Chromium 192 outdiffusion 132 Circularly polarized excitation 417 optical transitions 362 Cladding layer 433 Cleaning in-situ 172 Cleavage steps 588 Clustering 733

749

Clusters 724 formation 689, 699, 708 CMOS-based process technology 486 CMOS integrated circuits enhancement 485 co on Cu 659 on GaAs(ll0) 630 on Ge(ll0) 631 prelayer 629 Co-AS alloy films 660 Co/Au 632 Co-Cu alloy films 660 Co/Cu multilayers 636, 638, 655 Codeposit constituent species 571 Coherence length 678 longest 572 within multilayers 650 Coherency strain in as-grown LaFJDylLaF, 639 Coherent electron wavefunction 116 Coherent strain influence of 719 Cold cathode emission spectroscopy (CCES) 553 Collision Time Analysis 709 Colloidal graphite as a substrate-mounting agent 69 Common-emitter characteristics 334 Composition analysis techniques 554 control 507, 537, 552, 555, 572, 590, 591, 592 in-situ measurements 554 Compositionally-graded AlGaAs buffer 179 Compound semiconductor 275 II-VI 368 Computational routines 554 Condensed matter physics 116 Conduction band offset 366 dependence on strain 481 Confinement effects 373 Constant capacitance DLTS 225 Constructional materials for MBE systems 12

750

Index

Constructive interference 682 Contamination boron 556 by furnace parts and crucibles 127 metallic 71 of flux 31 of the source 51 Continuum analysis 703 approximation 733 distribution 737 model 373 Control loop adjustment of 75 Control systems for MBE 102 Coolants alcohol-water 124 Cooling liquid nitrogen 124 Co/Pd 644 Copper-containing high T, superconductors 510, 533 Co-Pt alloy films 651 Co/Pt 644 CoPt, alloys 651 forms spontaneously 648 Corner-sharing CuO, squares 510 Correlation function 736, 738 Coulomb field 414 Coulomb screening 416 Coulombic scattering of charged carriers 190 Coupling electromagnetic 63 oscillations 643 Covalent bonds tetrahedral 352 Cracker cells 36 hydride thermal 280 low pressure 287 temperature 135 zone 37

Cracking of ASH, and PH, 278 inhibited 284 tetrameric arsenic 131, 181 zone 385 Critical current densities 557 Critical density 416 Critical temperature 701 Critical thickness 463, 719 as a function of Ge content 464 and relaxation 465 Crucible 555 absence of 52 end-caps 33 lip condensation 531 material 128, 532 mechanical strength of 35 outgassing 222 pBN 34 preparation 130, 221 single crystal sapphire 128, 556 solid 48 thermal contact 71 Cryogenic pumps 531 Cryopanels for MBE 24 Cryopump 218 Cryopumping 288 Cryopumps 16 Cryoshrouding 123 Cryoshrouds baking of 125 Crystal field split 409 Crystal structure factor 404 high Tc superconductors 508 Crystallinity and growth temperatures 459 cu contamination 48 oxidation state 513 Cubic symmetry loss 454 cue formation of 545 oxidant flux needed 544 cue, layers 510

Index

PQhCan.l layers 522 Cuprate superconductors 510 Curie temperature 650 Current densities 557 Custom-layered oxide heterostructures 508 Customized layering 527 capability 564 CVD growth techniques 479 techniques of Si,,Ge, 462

D d-Electron excitation 411 states 410 Damping 693 of RHEED intensity oscillations 710 de Broglie wavelength 383 De-gassing of MBE systems 26 Debye limit 416 Debye-Waller factor 720 Decomposition of arsine and phosphine 279 kinetics of 514 of ozone 552 Deep-acceptor impurities 172 Deep electron traps 135 Deep level 184 Deep-level defects 193, 202 Deep-level transient spectroscopy (DLTS) 120, 161, 322 Deep-level transient spectroscopy 225 Defect concentration 459 point 459 sites 566 spike with a pit 166 Defect-free surfaces 300 Defect-induced bound exciton (DIBE) 183, 185, 198 transitions 165 Deformation potential 379, 394

751

Degreasing 173 Delta-doping 187, 190 density with carbon 189 Deposition continuous 533 laser-assisted 349 sequential 533, 566 Deposition rates of typical MBE 49 one atomic layer per second 2 Design of MBE systems 3 Desorption of As 698 of Ga 706 oxide 349 Devices 117 DHBT 333 devices 335 2DHGs modulation-doped 485 DIBE 202 Dielectric constant 558 Difference RHEED 570 Differential control of temperature 74 heating and cooling 125 pumping 531, 588 Diffracted beam split 683 Diffracted intensity 676, 679, 685, 733 calculation of 675 Diffraction analysis of 669 detecting schemes 674 pattern 672 two-level 736 Diffusion distance 151 length 701 of silicon 190 surface 701 Diffusion coefficient 701, 702, 706 of SiGa-SiAs pairs 190 Diffusion equation 704, 731 Diffusion length of aluminum 153 group III 148

752

Index

Diffusion pumps 288 for UHV 16 Diffusive growth 729, 731 Diluted magnetic semiconductors (DMS) 344, 362 Dimensional confinement 572 Dimensionality issues 573 Dimeric arsenic 134, 181 Dimers for Ill-V MBE 36 Dipole effects at heterointerfaces 368 Dipole selection rule 319 Direct radiative substrate heating 133, 174 Dislocations 354 formation of 463 nucleated 460 in superconductors 557 Disorder vicinal surfaces 683 Disordered growth 458 interface 389 superlattice 574 Disordering is reduced 711 Dispersion 277 Distortion rhombohedral 361 tetragonal 454 DLTS 225 measurements 228 DMS 345 lasers 375 material 405 quantum well 366 Dopants 202 incorporation 184, 187 isoelectronic 209 precipitation 55 in Ill-V semiconductors 53 in Si-MBE 56 low sticking coefficients 187 Doping 295 of cation sites 555 in CdTe/(Cd,Mn)Te 375 of GaAs 283

laser-assisted 375 for Ill-V MBE 52 of II-VI 345 of MBE-grown ZnSe of superconductors Droop 51 Dry ice as coolant 551 DX centers 191 DyBa,Cu,O,_, 549

392 556

E E-beam deposition 672 Early voltage 484 EBIC imaging of dislocations 464 ECR 530 Edge dislocations 356 incorporation 154 Effective strain equation for 464 Effusion cells 295, 299, 531, 553 furnaces 127 radiative heating source 35 Elastic strain 360 Electrically neutral formula unit 535 Electrochemical sources 54, 5‘5 Electromagnetic couplings disadvantages of 63 Electron trajectories 671 Electron beam evaporators 24, 46, 51 pumping 420 Electron concentration as a function of temperature 231 Electron diffraction analysis 675 measurements 669 patterns 423 Electron energy 672 Electron gas mobilities 181 Electron impact emission spectroscopy (EIES) 553

Index

Electron mobilities of GaAs 123, 125, 126, 134 Electron sheet densities 190 Electron Spectroscopy for Chemical Analysis (ESCA) 88 Electron transport measurements 161 Electron traps 178, 193, 203 Electron-Induced Emission Spectroscopy EIES 81 Electronic states quantum well 362 Elemental sources 295 group III 278 in MBE 277 Ellipsometry 591 EMF controls the sulfur flux 54 Emitter injection efficiency 334 Encapsulating 556 End-caps 31, 33 Energetically favored phases 527 Energy of formation 527 gaps 277 ion manipulation 58 Enthalpy of formation 526, 696 Epilayer contamination 132 Epilayer/epilayer interfaces 383 Epitaxial films superconductors 506 Epitaxial growth 137, 458, 685 of GaAs and AlGaAs 160 kinetics of 719 of magnetic metal films 626 mechanisms 153 Epitaxial layers 176 Epitaxy 115 layer-by-layer 680 EPMA error of 554 Equilibrium adatom population 705 Equilibrium partial pressures 280 as a function of total pressure Equilibrium phase 514, 537 Equilibrium surface coverage 707 Equilibrium theory and critical thickness 463

Erbium 192 in GaAs 198 Error in measurement 722 ESMBE 275 Ethylene 282 Eutectic AlSi 696 phasechanges 348 Evaporation from a virtual source 49 zone dynamics 49 Evaporator electromagnetically-focused 46 Ewald sphere 537, 680 of incident electrons 84 ExB filter 58 Excess arsenic flux 170 Exchange coupling 624 Exchange interaction in DMS 375 Exciton absorption and refraction 116 binding energy 366 diameter of 310 lifetime 372 localization 376 luminescence of GaAs 165 potential fluctuations 372 trapping 431 wavefunction 372, 409 Excitonic AlGaAs linewidth 199 effects 362 molecules 416 Expert systems for process control 102 Explosive oxidants 531 Exponential surface-ordering 151

F Face centered cubic (fee) Faceting 162 submicron 180 Faraday configuration 363 cup detector 60 geometry 368, 427

403

753

754

Index

Fast Entry Lock (FEL) 7 Fe prior to Ag epitaxy 628 prelayer seeding 629 single crystal whiskers 642 FelAg-seeded sandwiches 642 Fe-Cr multilayers GMR in 654 Feedthroughs hollow magnetically-coupled 67 Fermi level 178, 229, 399, 401 Ferrimagnetic 505 Ferroelectric 505 Ferromagnetic ordering temperature of rare earth metals 638 FETs n-channel heterostructure 488 p-channel heterostructure 485 Fiberoptic communication 277 Field effect transistor depletion-mode 399 Filament powered by 51 Si 52 Filter ExB 58 Fixed leak 295 Flat panel displays 387 Flux Ass/P, ratio 293 control 77, 286 detection 81 distribution curves 95 equation 30, 55 intensity variation 94 measurement 553 monitors 51 noise 168 ratio 200 Flux-sensor feedback 49 Forward bias 413 Fourfold pattern 672 Fourier transform 226 Fractional order beams 672, 692 Free carrier 116 concentration 172, 228 Free electron concentration 134 with silicon 190

Free energies 527 Free-standing superlattice limit Frenkel-Poole effect 236 Frohlich interaction 380 Furnace preparation 221 Future MBE system 592

376

G Ga doping of Si 57 sublimation rate 696 (100) GaAs 141, 147 GaAs 142 aluminum on 153 buffer 177, 181, 717 decomposition/desorption 184 deep levels 193 dopants 187 doping with Zn 57 on GaAs 162 growth 302, 694 high quality by MBE 39 impurity and defect states 240 optical properties 197, 204 RHEED diffraction from 685 sublimation 157 substrate 173, 347 surface 166 GaAs(lO0) misorientation 698 GaAs-on-AI,Ga,.,As 159 GaAs/AIGaAs 117, 163, 179, 182 heterojunctions 163 high purity 118 interface 119 modulation-doped 119 GaAs/AI,Ga, ,As MBE 4 GalnAs 304 ballistic transport 337 growth of 302 on InP 300 GalnAs(P)/lnP 275 GaInAsP 304 quaternary 322

Index

Gallium agglomeration 159, 167 aggregation 152 crucibles 128, 130, 167 desorption 146, 157, 205 diffusion 153 dimers 154 ingot etching 136 leaching 222 oxides 175 re-evaporation 182 saturation 475 solder 133, 174 source 136, 224 spitting 167 Gallium-exposed steps 162 Gallium-terminated GaAs 152 Gallium vacancy concentration 199 Galvanic cell PtlAglAgIlAg,S/Pt 54 Ga,O contamination 168 deep level 187 GasO,-related defects 167 O-Ga,O, 353 Gas adsorbed onto surface 11 analytical equipment 83 inlet devices 42 regulation systems 39 source 284 Gas handling system 288, 290 metalorganic 294 GaSb substrates 380 Ga,Se, 403, 405 Gaseous precursors 589 Gases behavior as a function of pressure 9 electron and hole 116 Gaskets 2 15 Ga-Te bonding 352 Gate oxides 486 Ge diffusion of 680

Ge content and growth temperatures 455 and stability 467 control of 462 effect on p-MOSFET performance 488 Germanium 189 Getter-pumping 531 Gettering impurities 177, 180 Giant magnetoresistance (GMR) 623, 654 Glide plane 575, 579 GMR in non-multilayer alloys 660 “g” Peak 209 Grain boundaries high angle 562 Graphite cracker 37 crucibles 129 filament sources 189 high purity 128 plates as substrate heaters 70 Grazing-incidence x-ray diffraction 647 Ground-state resonance 370 Group III cation intermixing 159 desorption 184 metalorganics 282 migration length 148 sources 135, 146 Group V metalorganics 294 Group V and group III ratio 200 Group VI chalcogens 53 Groups Ill and V elements as dopants 56 Growth 2D and 3D 455 controlled independently 569 disordered 458 epitaxial 458 in-situ 528 interrupted 689, 723 interruption 150, 159, 207

755

756

Index

kinetics 116 layer-by-layer 679 mechanism 115, 151 microscopic 160 mode 669 models 724 on (110) surfaces 162 on vicinal surfaces 699 process 686 rate 10 of high T, 506 Growth axis controlled by Co and Ag prelayers 630 Growth rate 688 and background impurities 205 of InSb 384 of MBE 115 Growth temperature 178, 198 and crystallinity 459 and deep levels 206 and Ge content 455 influence 183 Growth unit 533 minimum 566 GSMBE 5, 275, 295, 460 advantages of 38 described 38 vacuum pumps for 17

H H2

a carrier gas 293 H,S gas source doping from 54 Half order streaks 672 Half unit cell layering 577 Hall effect 228, 488 Hall mobility 118 of GaAs 229 Hamiltonian 366 HBT 183, 333 He+ ions channelling of 459 Heat treatment of the substrate 300 as

Heating by radiation 68 of substrates 70 Heavy metal contamination 173 Heavy- and light-hole states 327 Heavy-hole 362 Helium leak diagnostic 219 Heteroepitaxial structures atomic scale 115 Heteroepitaxy of alloys 454 Heterointerface II-VI/III-V 358 Heterojunction 114 II-VI/III-V 382 quality 464 Heterojunction bipolar transistors (HBTs) 463, 481, 484 Heterostructure bipolar transist 333 growth 300 semiconductor-based 453 silicon-germanium 491 synthesis of 507 Hgo.oaCdo.&“no.,,Te 346 High frequency applications 562 High Pressure Gas Source (HPGS) 284 High purity GaAs/AIGaAs 118 High Cl 558 High resolution x-ray diffraction High T, superconductors 538 copper-containing 510 Hole band 480 Hole density of 2DHG 486 Hole mobility Ge has highest 485 of modulation-doped devices Hole trapping 319, 322 Homologous series 522 Hopping probability 703 Hot filaments 553 Hot holes 319 Hot lip furnaces 168 Hot MBE components 127

305

485

Index

Hot photoluminescence spectrum 381 Hot wall epitaxy 430 HPGS 284 HREM study of interfaces 352 HSMBE 275, 278 Hydride-based chemistries 462 Hydride gas cell 42 safety 46 Hydrides complete decomposition of 45 Hydrocarbons desorption of 175 in the residual gas spectra 220 Hydrogen 211 during MBE 136 passivation 463 Hydrogenated silicon 453

I ldeality factor 335 II-VI/III-V interface nucleation at a 389 III-V:MBE systems 5 Image dissectors 674 Impurities background 116 incorporation of 171 at the substrate/epilayer 177 in vacuum chamber 127 In-phase conditions 680 In-Sb liquidus curve 299 In-situ cleaning of GaAs 172 InAs surface enrichment 159 three-dimensional growth 160 Incorporation of background impurities 184 of chemisorbed molecular species 154 rate 709 lndium 209, 375 in beryllium-doped 188 desorption 146

757

glued with 299 mounting 68 solder 133, 174 lndium free mountings 69 Information storage magneto-optical 623 Infrared ellipsometers 591 imaging devices 346 photodetectors 493 photoluminescence 387 pyrometry 72 radiation 68 Initial disorder 714 Initial nucleation 728 Initial peak width 711 Initial transient 691 Injection lasing 433 InP growth of 302 substrates 300 InSb 383 on CdTe 383 ion-milled 359 MQW 382 optical properties 387 InSb/CdTe MQW 385 Integral control of temperature 75 Integral order beams 672 bulk streaks 350 streaks 672 Integration with Si 563 Intensity distribution 690 oscillations 685, 686, 688, 690, 719 Interdiffusion estimate of 646 Interface CdTe/lnSb 383 coherence 360 HREM study 352 microstructure 396 roughness 169, 171, 181, 208 smoothing 207 traps 177

758

Index

Interface state density distributions 402 Interfacial layer 3.52 Interfacial problems 383 Interference function 726 Intergrowths 522 Intermixing of Ge in bulk Si 458 Interrupted growth technique 399 Ion beam sputtering 588 Ion bombardment negative 506 Ion flux measurement of 60 monitoring 77 Ion gauge 553, 696 Ion Gauge Flux Monitor (IGFM) 78 Ion implantation energies 57 Ion milling 91, 396 Ion source 56, 58 Ion sputtering 172 Ionization efficiencies 79 Ionization gauges 83 IR susceptor plate 68 Iron concentration 185 outdiffusion 132 Island 137, 208, 535, 685 formation 148 oriented presence of 569 temporary formation of 571 two-dimensional 148, 675, 679, 686 Isoelectronic dopants 209 J Josephson junctions vertical 567 Jump rate 729 Junction depth is estimated pn devices 433

K K-cell crucible 97 described 29

562, 572, 581

486

designs 34 for MBE 29 temperature regulation 74 Kerr effect 643 Kerr rotation 644 Kikuchi lines 683 Kinematic approximation 676 Kinematic shape effects 677 Kinematic theory 675, 686 Kinetic barriers to dislocations 463 to oxidation 548 Kinetic model 391 Kinetic theory to convert pressure to flux 706 Kinetics of MBE growth 137, 147 Kink density 151, 700, 701 KNbO, as a barrier layer 581 L LaAIO, 558 Laser ablation 506 blue and green 438 blue semiconductor injection 387 device configurations 436 diodes 183 oscillations in ZnSe 418 quantum well 278 Laser-assisted doping 375 Lasing wavelength 375, 377 Lattice defects 162 fringes 352, 354, 397 mismatch 354, 361, 378, 393, 454, 558, 719 mismatch strain 366 model 350 plane bowing 685 relaxation 159, 722 Lattice constant 454, 557, 722 determination 674 mismatch 347 of SrTiO, 637

index

Laue scattering condition 84 zone 350, 537 Layer coverage 725 Layer-by-layer growth 686, 689, 723 Layered structures 505 growth of 624 metastable 508 Layering on a half unit cell basis 577 sub unit cell precision 568 at the unit cell level 588 Lead 212, 514 Leaks checking 218 in UHV systems 12 LED 433 blue 437 blue and green 438 LEED 138, 633, 692 Light emitting devices 429 Light hole 362 excitons 317 Line-broadening 310, 370 Liquid metal solders 133 Liquid nitrogen cryopanels 24 Liquid phase epitaxy (LPE) 114 Load locks 8, 120 Loading the sources 224 Localization energy 372 Long range order 138, 163, 468, 680, 724 as a function of temperature 471 Longitudinal acoustic modes 372 Longitudinal optical (LO) phonons 368 Lorentz polarization factor 574 Loss tangent 558 Low Pressure Gas Source (LPGS) 284 Low-index model 734 plane 681 or singular surfaces 670 surface 670, 695, 724 LPE 115 has surface etching 172

LPGS 286 LRO effect of annealing temperature on 471 temperature dependence 479 LRP 460 Lubricants 15 Luminescence efficiency 135, 184, 197 of single quantum well 311 M Magnesium 185, 192 Magnetic anisotropy 624, 644, 650 excitations in (Cd,Mn)Te 374 impurities 229 moments 429 ordering 425 phenomena 624 polaron effect 372 semiconductor 405 zincblende semiconductor 421 Magnetically tunable optical sources 347 Magneto-optical behavior 408 measurements 427 properties of Co/Pt and Co/Pd multilayers 644 Magneto-photoluminescence 233 Magnetoresistance (MR) saturation 655 study of 624 Magnetotransport properties of multilayers 632 Magnon 374 Manganese acceptors 202 concentration 185 contamination 127 outdiffusion 132 Manifold pressure-controlled 292 Markov distribution of island sizes 679 Masks to shadow 64

759

760

Index

Mass action analysis 696, 697 law of 706 Mass flow control 39 vs pressure control 42 Mass spectrometry 553 Materials purity 114, 117 Matthews-Blakeslee (MB) critical thickness 721 Maxima shift in position 692 MB criterion of critical thickness 464 MBE 115, 124, 669 categorized 3 deposition of Si/Si,,Ge, 462 described 2 development of 7 electron beam evaporation 6 furnaces 29 future 592 growth 147, 160, 214 growth of eprtaxral Sr,,Ge, 454 growth of magnetic metal structures 624 growth process 528 growth techniques 479 history of 117 hybrid laser 589 intrinsic attributes of 1 Josephson junctions 581 new branch of 508 of Si 46 oxide superconductors 505 process 2 sources 29 strength of 564 substrate movement 9 success Of 568 system 120, 528 high T, superconductors 505 technique 588 MBE system construction 25 design 94 hardware automation 102

III-V 4 throughput 99 ME4 trap 186 Mean collision time 710 Mean free path 116 calculation 541 constraint 540 for MBE 508 Meander steps 707, 716 substrates heaters 70 Measurement of ion flux 60 Mechanical equilibrium and critical thickness 463 MEE 199 buffer layers 181 growth of Ill-V materials 591 MEIS spectra for different growth temperatures 455 Melt-back 48 Memory effects dopant 123 MESFET 178, 183 microwave 162 Metal-base transistors 572 Metal Organic Vapor Phase Epitaxy (MOVPE) 115 Metal-organic cells 44 gases 39, 40, 41 substitutes for hydrides 46 Metalalkyls group V 278 Metalorganics 302 group III 282 handling of 293 Metastability 507 Metastable layerings 537 materials 582 phases 527 zincblende 407 MgO 562 Micro-RHEED 160 Microcrystalline silicon 453 Microtwins 354 Microwave devices 336

Index

Migration enhanced epitaxial (MEE) growth 535 Migration enhanced epitaxy (MEE) 163, 169, 708 Migration length 148, 152 Minority carrier devices 333 MIS capacitors 399 MISFET devices 399 Misfit dislocations 356, 394 in-plane 626 strain 719 Misorientation 151, 681, 719 Misoriented surfaces 678 (100) 161 (110) 162 Mixing 511 MMIC 564 Mn-ion concentration 410 Mn, ,Ga, ferromagnetic o-phase of 626 MnSe 423 MnTe 378 MnTe/CdTe SQW 377 MnTe/ZnTe SQW 380 MO sample blocks 299 Mobile adatom density 704 Mobilities field-effect 488 MOCVD 115, 589 Model growth 724 MODFET 181, 183, 278, 481 Modulation depth 320 Modulation doping 116 Modulation-doped 2DHGs 485 Modulators 414 Mole fraction variation 709 Molecular beam deposition 506 Molecular Beam Epitaxy (MBE) 115, 275, 669 DMS 346 Molecular beam methods 276 Molecular chemisorption 137 Molecular oxygen MBE-compatible pressures 543

761

Molybdenum 531 sample holders 133 in UHV systems 14 MOMBE 276, 278 Momentum transfer 679 Monatomic steps 682 Monolayer control 423 defined 9 Monte Carlo calculations 675, 701, 703 simulations of ZnSe 391 MOO, volatility of 531 Morphological instabilities 460 MOSFETs 481 performance of 485 Motion in UHV system 15 Mot-t screening 416 MOVPE 115, 124, 172 Multilayers Co/Au 632 Co/Cr 654 Co/Cu 632, 636, 638, 655 Co/Pd 644 Co/Pt 644 Fe/Cr 654 permalloy/Au 658 Multiple elastic scattering 675 Multiple quantum well (MQW) 359 buffer layers 179 Multiple-target sputtering 589

N Neel-type anisotropy model Nitrogen boiling point 551 NO, 548, 549 Non-planar growth 161 Non-stoichiometry 512 Nondiffusive growth 727 Nonlinear optical absorption Nucleated islands 137 Nucleation centers 151 of CdTe on GaAs 358

647

414

762

Index

of surface defects 166 three-dimensional 350, 357, 384 two-dimensional 357 of ZnSe 401 of ZnSe on GaAs 387

0 0+ 546 O-ring for UHV systems 14 gaskets 216 seal 126 0, 545, 550 Off-Bragg conditions 388 Ohmic contact freeze-out 229 OMCVD 115 OMVPE 115, 588 Optical devices 387 disc storage 346 dispersion curves 373 DLTS 225 flux detection 81 properties of superlattices 362 spectroscopy measurements 409 transition 315, 393 Optical absorption nonlinear 414 spectroscopy 236 Optical phonon emission 411 scattering 380 two-mode behavior 369 Optical properties of GaAs and AlGaAs 197, 204 of superlattices 312 of the growing film 591 Single quantum wells 308 Optically pumped lasers 377, 419 Optoelectronic applications 491 devices 161, 345, 481 Ordering and the proportion of Si and Ge 472 as a function of temperature 472

due to surface kinetics 475 in elemental semiconductor alloys 468 Organometallic sources 299 Orientation effect of surface stoichiometry 358 of Ag films 630 of covalent bonds 350 of the substrate 160 Oscillation asymmetric 692 damping 148 in RHEED intensity 533 strongly damped 699 sublimation 695 Oscillatory exchange coupling 624 in Fe/Ag/Fe (001) sandwiches 642 studies of 643 Out-of-phase conditions 680 Outdiffusion of impurities 132 problems 132 Outgassing difficult 31 high-temperature 129 wafer 175 Oval defect 166, 167 density 133, 168, 207 gallium-related 182 Oxidant alternative 544 choice of 548 flux 571 shuttering of 571 Oxidation behavior 545 of Cu 512, 513 state 512 Oxide desorption 175, 348 layer 349 of high T, 527 protective 174 residual 300 superconductors 505 Oxidized multi-element high T, 508 Oxidizing agent 530

Index

Oxygen contamination 183 Contamination of gallium 136 flow rate 540 loss rate 548 partial pressure 515 peak 219 pressure 506, 537, 538 source 56 Ozone 546, 550 trap 550

P Pair defects 166, 173 Parabolic function 722 Passivating oxide layer 354 surface oxide 175 Passivation 213 deep-level 212 impact on epitaxial film quality 463 Path length difference 679, 691 PBN crucibles 128, 130, 167, 222, 555 decomposition of 31 diffuser zone 44 insulators 34 is a good IR absorber 69 reaction between aluminum 128 PbS 191 PbS, PbSe, and PbTe doping from fluxes of 54 PbSe 191 Peak intensity 680 separation 684 splitting 207 Penumbra effects 97 formed by the crucible 33 Perfect layer growth 691 Perovskite cubic structure 637 related structures 508 substrate materials 558

763

Perpendicular magnetic anisotropy 623, 642, 651 PH3

gas handling systems 290 Phase coherence 116 differences 672 shift 692 Phase diagrams 514 BiO,,,-SrO-CuO 518 effect of PO, 516, 521 pseudo-ternary 5 14 YO, ,,-BaO-CuO 515 Phase-space filling 414 Phonon confinement effects 373 Ge 456 modes 368, 372 Phosphorus 279 cracking 36 red 277 safe handling of 45 white 277 Photoconductivity 236 Photodetectors Si/SiGe 491 Photoemission oscillations 160 Photoionized holes motion of 324 Photoluminescence 120, 161, 232 efficiency of erbium-doped 192 exitonic 491 experiments 308 microscopy 160 Photoluminescence excitation (PLE) 235, 314, 431 Photon emission injection-induced 435 Photothermal ionization spectroscopy 120, 237 PID control systems 74 Pierls barrier 463 Piezoelectric 505 device applications 161 fields 371 p-i-n Diodes 304 Pinning potentials 229 P,/ln ratio 281

764

Index

PL quantum efficiency 413 Planar doping 187 Planar growth 148 Planar MBE 154 Planar p-n junctions 164 Platinum 532 PLD 588 produced by 564 p-MOSFET Ge content 488 hole mobility 485 output characteristics 488 Point defects density of 685 Polarization 319, 327, 420 Polycrystalline CdTe films 354 silicon 52 Post-growth anneal 400 Potential barriers 362 Power requirements for electron beam evaporation rl Power supplies solid state switched 51 Prebaking 221 Precipitation 190 Precursors gaseous 589 Preheat temperature 354 Prelayers Ag 628 Preparation chamber 8 Preservation of as-grown surfaces 213 Pressure vacuum chamber 348 PrGaO, 558 Primary pumps 16 Proportional band response of the power supply 50 Proportional control of temperature 74 Protective oxide 174 Pseudomorphic growth of Si,,Ge, on Si 454 Pt on sapphire 635 seed film 636

Pt(l1 l)/basal-plane sapphire 632 Pt/basal-plane sapphire epitaxial systems 625 Pt/SrTiO, 637 epitaxial systems 625, 632 PTIS 237

51

p-Type doping 187 to n-type shifts 283 Pulsed laser deposition (PLD) 506 Pumping rate 531 Pumping speed 290 Pumps 288 for gas source MBE 17 turbomolecular and cryogenic 531 for UHV 16 Purity of material 117 Pyrolytic boron nitride 127 has poor conversion efficiency 38 as a cracking medium 43 as electrical insulator 15 Pyrometry infrared 72 Q Quadrupole mass analyzer @MA) 84 Quadrupole mass spectrometers for flus monitoring 77 Quality improved by FELs 8 Quantum dots 116 Quantum Hall effect 116, 488 Quantum phenomena 116 Quantum size effect energy shift 310 Quantum well 116, 152, 204, 304 barrier heights 362 calculations 363 (Cd,Mn)Te 361 first 180 InSb/CdTe 345 isoelectronically doped 432 lasers 719 optoelectronic 491 p-i-n photodetectors 481

Index

Si, _xGex 494 strained 486 QUantUIn wire 116, 304, 311 devices 161 Quantum-confined Stark effect 304, 319, 327 Quartz crucibles 128 Quartz crystal microbalance 553, 554 monitor 80, 392 resonators 77

R Radiance ratio pyrometer 73 Radiant heating of the substrate 68 Radiation damage from ion milling 359 Radiative heating 35, 174 recombination 409, 434 Raman spectra for different growth temperatures 455 Rare earth 556 epitaxy on GaAs 626 sandwich structures 638 species 531 Reactive evaporation 588 Reactive oxidants 549 Reactive oxidizing agent 530 ReBasCu,Ord 565 superconductor phases 567 ReBaCuO,BO, 556 Recombination centers 311 efficiency 434 lifetime 41 1 non-radiative 189 velocity 182 Recombining (energy-relaxed) excitons 372 Reconstructed GaAs 348 surface 142

765

Reconstruction 2 x 4 GaAs(lo0) 672 patterns 401 surface 685 Red shift 394, 414 REELS 590 Reference temperatures 73 Reflection twins 584 Reflectometry 591 Refractory metal thermocouples 71 Regrowth processes 213 Relaxation 722 kinetic factors 466 thickness 465 Reordering 137 Residual gas peaks 218 Residual gas spectrum (RGS) 84 Residual gases in MBE 121, 127 Residual oxide 349, 357 Resistance heating 69 Resitivities for low temperature GaAs 178 Resonant Raman scattering (RRS) 368 Resonant tunneling 116 diode 481, 494 structure 466, 475 RHEED 84, 138, 147, 348, 588, 624, 630, 633, 635, 669, 675 applications of 87 as a calibration method 590 fishnet pattern 388 in-situ analysis 565 intensity oscillations 702, 710, 720 lead-induced changes 212 oscillations 148, 533 spotty pattern 349, 350 streaked pattern 401 streaks 537 transfer width of 681 RHEED-TRAXS 590 Rippled surface structure 152 Rocking curve 305 diffraction 385 ZnSe 389

766

Index

Room temperature thermodynamically stable at Rotation to maintain compositional uniformity 299 of substrate 65 twins 584 Rotational symmetry fourfold 350 Roughness 169, 184, 726 of GaAs-on-AIGaAs 207 microscopic 180 surface 689 RTCVD 460

514

S S doping in GaAs 54 Safety 284, 292 of ASH, and PH, 278 GSMBE systems 46 of toxic hydrides 45 SAM-APD response structures 321 Sample preparation of II-VI compounds 396 Sapphire high frequency 562 Sapphire-based superlattices 656 Saturation magnetoresistance 655 for varying Cu thickness 658 Sb cracker 384 doping of Si 57 Scattering angle 676, 679, 691 effect 361 factors 403 geometry 671, 686 Schottky barrier field 412 Schwoebel effect 711 Screening 417 Screw dislocations 566 SDHBT 333 Se source characteristics of 55

Secondaries low-energy 671 Secondary ton Mass Spectrometry (SIMS) 239 Secondary Ion Mass Spectroscopy (SIMS) 88 Seed film 626 Seeded epitaxy 626 Segregation of Co 659 Selective orientation of CdTe 348 Selenides 392 Selenium 191 Self-anneals during deposition 57 Self-compensation 375 Self-field 558 Self-trapping 431 Semi-insulating 199 AlGaAs 136 GaAs 133 Semiconductor 563, 670 hole mobility of 485 II-VI 345, 405 Ill-V 277 materials 117 Semimagnetic semiconductor 344, 405 Separate avalanche-multiplication (SAM) 316 Sequential deposition 589 Shadow masks 64 Shadowing 692 absence of 97 of the evaporant 33 Shutter blades 64 source 60 substrate 64 Shuttered MBE 553, 568, 577 Si cap 485 doping with Sb and Ga 57 filament source 52 MBE of 46 SiAs-SiGa pair formation 202

Index

SiGaSiAs pair formation 190 Si, _xGe, 453 epitaxial growth 463 film stabilizing 465 HBTs stability 484 thickness 463 Si-MBE filaments 52 system 6, 51 Si/SiGe waveguides and photodetectors 491 Si/silicide 453 Sidegating 178 Sidewall recombination 311 Silica gel to contain ozone 551 Silicon 189, 453 amphoteric behavior of 164 diffusion coefficient 191 as a donor/acceptor 166 dopants 187 incorporation rate 213 superconductors 563 Silicon-on-Insulator (SOI) structures 491 SIMS 92, 239 calibration 239 Single quantum well (SW’/) 308, 311, 377 Single scattering theory 675 Single-layer islands 679 Singular surfaces 670 SIS Josephson junctions 582 Site-specific segregation of Si and Ge atoms 469 Slow response component 319 SLS 327 Smoothing effect 180 SnTe 191, 202 Solar cells 183 Solder indium or gallium 174 to mount GaAs wafers 133

767

Source electrochemical 54 ion capillary needle 60 ion plasma 58 purity 133, 392 Sources for flux generation 26 ion 56 MBE 3 not point sources 97 Source-drain junctions 486 Source materials for MBE 134 high vapor pressure 11 loading 223 purity of 50, 121 Spectroscopic ellipsometry (SE) 591 Source-substrate geometry 95 Specular beam 671, 682 intensity 719 streak 679, 680 Spike 679 Spin split 362 Spin-dependent transmission of conduction electrons 654 Spiral growth mechanism 566 Split specular beam 733 Splitting misorientation-induced 701 Sputtering 588 epitaxial films 506 produced by 564 Sr,Bi,Os phase 569 SrTiO, 558 Stability electron beam position 51 of Si/Si,,Ge, structures 467 Stabilizing of pseudomorphic Si, +Ge, films 465 Stacking fault 397 density 459 Stainless steel 532 a source of impurities 221 in MBE chambers 127 Staircase 678, 681, 699, 700, 731 signature 683 step arrays 147

766

Index

Stark effect 321, 412 quantum-confined 304 Steady state surface 689 Step bunching 684, 685, 710, 713 disorder 680 edges 137, 162 grading 335 heights 566 incorporation time 708 meandering 716 on sapphire (0001) 635 propagation 148, 151, 566, 686, 689, 699 scattering 692 spontaneous migration 151 termination 716 terrace length distribution 714 train disordering 710 trains 724 Step-catalyzed dissociation 700 Sticking coefficient 124, 540, 552 of C-bearing species 10 of dimeric sulfur 202 of group III 146 of magnesium 192 Stimulated emission 375, 419 polarized 421 STM 138, 160 Stoichiometric 511 Stokes shift 313, 372 Strain effect on bandstructure 481 coherency 639 coherent 719 compressive uniaxial 421 control of 466 energy 721 hydrostatic component of 3.59 determines morphology 460 relaxation 467 Strain-free limit 367 Strain-induced effects 362 uniaxial symmetry 364 Strain-layer superlattices 360 Strain-split heavy hole band 480

Strained layer 307 growth 719 quantum well 370 superlattice 430 Strained layer superlattices (SLS) 304, 325 Strained SQW CdTe/MnTe 378 Streak in RHEED 678 is composed of two parts 690 Strontium 569 Structural analysis methods of dislocations 464 Structures high T, superconductors 508 Sublimation 694, 708 Substrate cleaning 167 heaters 65 holder 127, 174 large area single crystal 557 motion 65 preparation 166, 172, 173, 299 reference 558 rotation 94, 127 sapphire 563 silicon 563 preheat temperature 354 temperature 72, 98, 350 Sulfur 192 background impurity 185 contamination 129 electrochemical 55 incorporation 201 Sulfur-containing alloys 434 Superconducting oxide phases 508 Superconductivity mechanism of 528 Superconductor-insulatorsuperconductor (SIS) 572 Superconductors 505 on silicon 563 Superlattice 573 buffer layers 179, 181 comb-like 425 incommensurate 587 modulators 316 optical properties 312

Index

sapphire-based 656 small period 491 smoothing 180 strained-layer 359 structures 304 transport 322 vertical (tilted) 116 Surface absorption chemistry 589 analytical equipment 91 atomic structure 685 chemistry 137 composition 462 contamination 166, 170 diffusion 528, 701 enrichment 159 free energies 722 incorporation 153, 156 kinetics 137, 391 lattice parameter 723 migration 137, 391 (111) misorientations 161 mobility 689, 719 morphology 670 passivation 2 13 phase diagram 138 resistance at microwave frequencies 562 roughness 169, 171, 566, 686, 689, 690, 692 segregation 170, 184, 191 step propagation 147 Surface diffusion anisotropic 152 of gallium 152 Surface growth kinetics and LRO 469 Surface recombination InP 331 Surface reconstruction 138, 212, 401, 702 (100) 138 and LRO 472 applications of 86 (100) GaAs 141 modified using an adlayer 469 Switches bistable 414

Synchrotron x-ray diffraction 647 Syntactic intergrowths uncontrolled 522 Synthesis technique 507

T Ta cracker 37 cracking tube temperature 384 Ta-catalyzed low pressure source 287 Target multi-component 507 TbFeCo amorphous alloy film 652 Te monolayer 357 Te-to-Se flux ratio 430 TEG 303 MOMBE with 282 TEI 302 Tellurides 392 Tellurium 191 Temperature cracker zone 37 cracking zone 385 dependence 185 influences during growth 1183 measurement 71, 184 of effusion sources 34 of MBE growth 115 offset 72 substrate 349 Temperature control by PID 74 hardware 76 Temperature dependence for growth of Co/Cu multilayers 638 of magnetization 426 Temperature-dependent trap emission 226 Template films intermetallic 626 Tensional strain 370 Ternary alloy semiconductors 148 Terrace edges 162 width 147

769

770

Index

Terrace length 683, 701 dependence 710 order 717 rate equation 713 Terraces 566 size 698 Tertiarybutylarsine (TBA) 278, 294 Tertiarybutylphosphine (TBP) 278, 294 Tetraethyltin 295 Tetramer molecules 36 Tetrameric arsenic 134, 181 sticking coefficient of 146 Thermal 393 Thermal dissociation two-zone 36 Thermal Effusion Sources for MBE 29 Thermal etching 157 for in-situ cleaning 172 Thermal expansion coefficient 557 mismatches 558 Thermal oxides 175 Thermocouples 71, 299 become brittle 130 for K-cells 34 Thermodynamic of growth 116 redistribution 158, 160 stability 537 Thermodynamic equilibrium arsenic and phosphorus 280 Thickness of Si,,Ge, 463 Threading dislocations 354 Three term Control systems 64 PID 74 Three-dimensional growth 152 Threshold current 124, 436 Threshold pump intensity 377 Threshold-voltage control 486 Throughput of MBE systems 99 Ti-sublimation pumps 532 Tin 191 dopants 187 doping 295

Titanium sublimation pump (TSP) 17, 124 Titanium-molybdenum alloy filaments 532 Tl,Ba,Ca,_, CU,O~~+., 571 TMG for the growth of GaAs 283 Tooling factor 80 Transconductance saturated 488 Transfer width 678 Transition metal elements 626 structure 403 temperatures 564 Transport properties best 564 Trap characteristics 226 density 206, 225 energy 225 Trapping 431 Trialkyl group III compounds 278 Triethyl compounds 302 Trigger Penning ionization gauge 83 Trimethylgallium 187 Trimethylindium 293 Triode valve to reduce droop 51 TSP 125 Tungsten 532 Turbomolecular 288 pump 16, 531 Twin boundaries 584, 586 Twin-free films 584, 587 Twinning 358 a dominant structural defect 646 origin of rotational 635 Two level diffraction 736 Two-level system 679 Two-dimensional electron gases (2DEGs) 466 Type conversion of the substrate surface 132 Type I superlattice 329

Index

U UHV environment for MBE 11 UHV-CVD 460 Ultra-high vacuum (UHV) system 127 Ultra-violet light 127 Ultrathin layers 430 Uniformity deposition 94 Unincorporated molecular species 209 Unit cell 533 UV irradiation for gas desorption 11

v V/ill flux ratio 200, 203 V/Ill ratio 166, 200, 202 Vacancies 511 Vacuum diagnostics 83 grease 216 in-situ characterization 508 interlocks 120 leaks 123 load lock 121 pressures 126 quality 125, 127 requirements for MBE 9 three-chamber systems 121 Vacuum chamber baking 125 construction of 25 Vacuum system sources of gas in 11 Valence band degeneracy 327 dependence on strain 481 splitting 370 Valence states 362 Valved-cracker source 38 Valves all-metal 126 for UHV ysstems 13

771

Van der Pauw 228 Vegard’s law 720 Vegard’s rule 454 Vibrational modes in the alloy system 456 Vicinal substrates 588 Vicinal surface 681, 731, 734 disorder 683 superconductors on 584 Video systems 674 Viewports for UHV systems 14 Virtual leak 218 Virtual source evaporation from 49 Voigt geometries 368 Vortex pinning sites 557 VPE 172 W Wafer preparation GaAs 348 Wannier-Stark effect 116 Water cooling around the K-cell 34 Water vapor 219 in MBE 121 Water-cooled panels 24 Waveguide modulators 322 Si/SiGe 491 Wavelength 383 lasing 375 Wavevector 677, 725 “Weekend effect” 124 Wide bandgap superlattice Writing epitaxial 64

405

X X-ray diffraction 305, 396, 646, 647 pattern 575 theory 573 X-ray goniometry 683 X-ray photoelectron diffraction (XPD) 160, 624, 645

772

Index

X-Ray Photoelectron Spectroscopy (XPS) 88 X-ray Photoemission Spectroscopy (XPS) 158 X-ray reflectivity 636, 646 X-ray rocking curve diffraction 385 X-ray scattering 660 XPS basic equation of 91

Y YBa,Cu,O, 514, 544 YBa,Cu,O, 514 YBa,Cu,O,_, 537 YO,.,-BaO-CuO 514 Yttria-stabilized cubic zirconia

562

Z Zeeman effect 427 photoluminescence 233 shifts 382 split 363 Splitting 345, 364, 366, 406, 409 Zincblende 344, 407, 700 crystal structure 421 MnSe 421 on diamond 162 phase of MnTe 377 reflections 403 step terminations 711 ZnSe 393 Zirconia cubic 562 Zn doping of GaAs 57 Zn-to-Se flux ratio 392, 393 (Zn,Mn)Se 405 i’nSe 388, 405 miCrOStrUCtUre 396 on GaAs 387 optical properties 409 pseudomorphic 345, 399 CjUantUm well 418 substrates 391 well thickness 410 ZnSe-based laser 435

ZnSelGaAs interface 403 post-growth annealed ZnSelMnSe superlattices ZnSe/ZnTe superlattice structures 429 ZnTe-based SQW 380 Zone folding 491 effects 372

400 425

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