Properties of Gallium Arsenide

PROPERTIES OF

Gallium Arsenide THIRD EDITION

Edited by M. R. BROZEL Centre for Electronic Materials, UMlST UK

and C. E. STILLMAN University of Illinois, USA

INSPEC

Published by: INSPEC, The Institution of Electrical Engineers, London, United Kingdom © 1996: The Institution of Electrical Engineers

This publication is copyright under the Berne Convention and the Universal Copyright Convention. All rights reserved. Apart from any fair dealing for the purposes of research or private study, or criticism or review, as permitted under the Copyright, Designs and Patents Act, 1988, this publication may be reproduced, stored or transmitted, in any forms or by any means, only with the prior permission in writing of the publishers, or in the case of reprographic reproduction in accordance with the terms of licences issued by the Copyright Licensing Agency. Inquiries concerning reproduction outside those terms should be sent to the publishers at the undermentioned address: The Institution of Electrical Engineers, Michael Faraday House, Six Hills Way, Stevenage, Herts. SG1 2AY, United Kingdom While the author and the publishers believe that the information and guidance given in this work is correct, all parties must rely upon their own skill and judgment when making use of it. Neither the author nor the publishers assume any liability to anyone for any loss or damage caused by any error or omission in the work, whether such error or omission is the result of negligence or any other cause. Any and all such liability is disclaimed. The moral right of the author to be identified as author of this work has been asserted by him/her in accordance with the Copyright, Designs and Patents Act 1988.

British Library Cataloguing in Publication Data A CIP catalogue record for this book is available from the British Library ISBN 0 85296 885 X

Printed in England by Short Run Press Ltd., Exeter

FOREWORD It is approaching 50 years since H. Welker gave gallium arsenide its initial impetus. During that time an immense amount of research has been carried out around the world - to grow bulk GaAs, to study its properties, to develop precise epitaxial growth methods and to realise devices with wide applications. There are now tens of thousands of scientists, engineers and technical support staff involved in work related to GaAs and related substances. For the people concerned and those destined to join them in the next few years this book will be an important reference. The fundamental properties of the materials, as well as the means of characterising them, are covered. Among the latter, deep level transient spectroscopy is present as a precise method of defect characterisation, for example. In addition, the methods best used to process device structures, including etching, ion implantation, making ohmic and Schottky contacts and passivation of surfaces are reviewed. The intense research and development efforts in the areas of the bulk growth of GaAs, and epitaxial growth by VPE, MBE and MOCVD, have been exceptionally successful in yielding sometimes even sophisticated structures to very precise specifications. These results have been the foundation for the success of GaAs devices, now widely used. Over the years the development of various heterojunctions and thin strained (pseudomorphic) materials have been intensively studied. These materials are now bandgap engineered to create sophisticated devices including high performance transistors, lasers and far IR photodetectors. In the second decade of GaAs research simple two-terminal devices were developed. These included Gunn diodes and IMPATT diodes, the former of which are still used in velocity detectors via the Doppler effect in devices such as automobile speed monitors, automatic door openers in public buildings and some alarm systems. Early research on metal-semiconductor field effect transistors and heterojunction lasers followed in the next decade. In the fourth decade heterojunction field effect transistors like the MODFET and various quantum well lasers were studied. In the fifth decade pseudomorphic, compressionally-strained quantum wells have been optimised for use in high performance microwave and millimetre wave transistors which are widely applied in communications areas. Using the same strained quantum wells GaAs lasers are being made and widely applied as optical pumps for Er-doped fibre amplifiers. Superlattice structures have also been made recently and used for far IR cascade lasers and far IR photodetectors. Thus a wide range of communications, via local wireless and satellite systems, as well as fibre optical systems, depend on GaAs and related materials. The total device market is now several billion dollars a year and growing strongly with the sharp rise in the personal communication market. This third edition is intended to be a reference for all those using GaAs. It contains a complete listing of important parameters of GaAs together with state-of-the-art reviews on microwave and digital opto-electronic devices. Each of some 150 Datareviews is individually dated and is as current as the preparation of such a reference book allows. It also guides the user to key references (amounting to over 5000 in all) so that further data can easily be located. This should be a valuable information resource for everyone in the field. Lester F. Eastman Phillips Hall, Cornell University, Ithaca, New York, USA October 1996

INTRODUCTION

This is the third edition of Properties of Gallium Arsenide, the previous volume having been published in 1990. GaAs is the basis of most compound semiconductor devices and whether it is used as a substrate for homoepitaxial or heteroepitaxial growth or for direct ion-implantation, it has a key role in modern high speed and opto-electronic devices. Since the publication of the second edition there have been over 35,000 papers published in this extremely active field. There is clearly a need for an updated edition to reflect the major advances made since 1990. For the first two editions of this book, it was decided to restrict the presentation so that separate Datareviews could be accessed electronically. The discipline involved in doing this resulted in a uniquely valuable work whose content could be used worldwide. However, a drawback of this style of production was the restricted range of letters and mathematical symbols that could be supported by the technology of the time. For the current volume, it was decided to follow the lead given by other more recent volumes in this series in that electronic access was not to be used. As a result we have been able to use virtually all types of presentation in this edition. We have encouraged authors to present tabular, graphical and even photographic material. As the reader will see, this has resulted in a great improvement in the overall quality of presentation of physical phenomena. GaAs has matured from a scientific investigation of potentially useful material to a manufacturing technology which is constantly expanding. Emphasis has turned somewhat into the fields of monolithic microwave integrated circuits (MIMICS) and opto-electronics, although discrete and high speed digital circuits remain important. For these reasons we have included several sections on devices and these can be found towards the end of the book. We hope that these will be of particular interest to those working on the more scientific aspects of GaAs research: the uses and problems of using GaAs in real devices is often imperfectly understood by technologists working outside the device fabrication laboratory. The contents of a review book such as this are only as good as the material that is written by the contributors. We heartily thank all our authors for their efforts. We also wish to thank all our referees (mostly UK) for their time and valuable comments. The assistance of John Sears, Managing Editor of the EMIS Datareviews series, has been invaluable throughout. Lastly, we would like to acknowledge the expert help of Jacqui Gilmore who has, single handedly, prepared the typescript pages for the book. M.R. Brozel (UMIST, Manchester, UK) G.E. Stillman (University of Illinois, USA) October 1996 Note by M.R. Brozel As our authors will know, most of the scientific and linguistic editing was performed at UMIST. In some cases, it was not possible to confirm with the author the scientific accuracy of changes made to the manuscripts in order to improve clarity and overall presentation. Although I believe that all such changes fiilly reflect the intentions of the authors, I must accept full responsibility for any errors, however small, that may have resulted.

CONTRIBUTING AUTHORS CR. Abernathy

University of Florida, MS&E Department, Rhines Hall, Gainesville, FL 32611, USA

16.8

S. Adachi

Gunma University, Department of Electronic Engineering, 1-5-1 Tenjin-cho, Kiryu 376 Gunma, Japan

1.6-1.9, 4.6,14.9

AR. Adams

University of Surrey, Department of Physics, Guildford, Surrey, GU2 5XH, UK

2.7,3.2,3.3, 4.2,4.3,4.7

PJ. Apostolakis

University of Illinois at Urbana-Champagne, Department of Electrical and Computer Engineering, Urbana, Illinois 61801, USA

20.1,20.2

C I H . Ashby

Sandia National Laboratory^ Division 1126^ MS 0603., PO Box 5800, Albuquerque, NM 87185, USA

18A-18.8

MG. Astles

DRA, St Andrews Road, Great Malvern, Worcs, WR14 3PS, UK

16.2

J. Ballingall

Novellus Systems Inc, 3970 N. First Street, San Jose, CA 95134, USA

2.10

S.V. Bandara

Jet Propulsion Laboratory, Mail Stop 302-306, California Institute of Technology, 4800 Oak Grove Drive, Pasadena, CA 91109-8099, USA

21.4

U. Bangert

UMIST, Department of Physics, PO Box 88, Manchester, M60 IQD, UK

10.9

J.M. Baranowski

Warsaw University, Institute of Experimental Physics, Hoza 69, 00-681 Warsaw, Poland

10.2

P.K. Bhattacharya

University of Michigan, Department of Electrical Engineering & Computer Science, 1301 Beal Avenue, Ann Arbor, MI48109-2122, USA

21.1

JC. Bourgoin

Universites Paris 6 et Paris 7, Groupe de Physique des Solides.Tour 23, 2 place Jussieu, 75251 Paris cedex 05, France

16.4

IW. Boyd

University College London, Department of Electronics and Electrical Engineering, Torrington Place, London, WCl 7JE, UK

13.6

MR. Brozel

UMIST, Department of Electrical Engineering and Electronics, PO Box 88, Manchester, M60 IQD, UK

9.10,10.4 10.7

CM. Buttar

University of Sheffield, Department of Physics, Mappin Street, Sheffield, UK

22.2

M. Cardona

Iowa State University, Ames Laboratory, A205 Physics, Ames, IA 50011, USA

4.5

W-H. Chang

University of Illinois at Urbana-Champagne, Department of Electrical and Computer Engineering, Urbana, Illinois 61801, USA

20.1,20.2

Shun Lien Chuang

21.6 University of Illinois at Urbana-Champagne, Department of Electrical and Computer Engineering, 1406 West Green Street, Urbana, Illinois 61801-2991, USA

M. Claassen

Technical University of Munchen, Lehrstuhl fur Allgemeine Elektronik und Angewandte Elektronik, Munchen, Germany

20.8

J P R . David

University of Sheffield, Department of Electronic and Electrical Engineering, Mappin Street, Sheffield, Sl 3JD, UK

4.9

CG. Diskus

Microelectronics Institute, Johannes Kepler University Linz, Attenberger Strasse 69, A-4040 Linz, Austria

20.6

DJ. Dunstan

Queen Mary and Westfield College, 327 Mile End Road, London El 4NS, UK

1.2-1.5

KR. Evans

Wright Patterson Air Force Base, Electron Devices Division, Avionics Directorate, Wright Laboratory, OH 45433-7323, USA

16.7

M. Feng

University of Illinois at Urbana-Champagne, Department of Electrical and Computer Engineering, Urbana, Illinois 61801,USA

20.1,20.2

J. Freyer

Technical University of Munchen, Lehrstuhl fiir Allgemeime Elektronik und Angewandte Elektronik, Munchen, Germany

20.8

K. Fujita

Sumitomo Electric Industries Ltd, Kami Works, 1-1-1 Koya-kita, Itami, Hyogo 664 Japan

15.3

IR. Grant

MCP Wafer Technology, Maryland Road, Tongwell, Milton Keynes, MKl5 8HJ, UK

15.2

C R M . Grovenor

University of Oxford, Department of Metallurgy and Science of Materials, Parks Road, Oxford, OXl 3PH, UK

14.2,14.3

S.D. Gunapala

Jet Propulsion Laboratory, Mail Stop 302-306, California 21.4 Institute of Technology, 4800 Oak Grove Drive, Pasadena, CA 91109-8099, USA

MC. Hanna

National Renewable Energy Laboratory, 1617 Cole Blvd, Golden CO 80401-3393, USA

3.5

E. S. Harmon

MellWood Laboratories, 1291 Cumberland Avenue, Suite E, West Lafayette, IN 47906, USA

2.13,4.8

JJ. Harris

University College London, Department of Electrical and Electronic Engineering, Torrington Place, London, WClE 7JE, UK

2.8,2.9

I. Harrison

11.1-11.4 University of Nottingham, Department of Electrical Engineering, University Park, Nottingham, NG7 2RD, UK

H.L. Hartnagel

Institut fur Hochfrequenztech., Fachbereich 18, 6100 Darmstadt, Merkstrasse 25, Germany

13.2-13.5 13.7

H. Hasegawa

Hokkaido University, Research Center for Interface Quantum Electronics, North 13, West 8, Kita-ku, Sapporo 060, Japan

12.1-12.3

A.M. Hennel

Warsaw University, Institute of Experimental Physics, Hoza 69, 00-681 Warsaw, Poland

7.3,7.4 9.10

MJ. Howes

University of Leeds, Department of Electronic and Electrical Engineering, Leeds, LS2 8JT, UK

20.5

D T J . Hurle

Scotscraig House, Storridge, Malvern, Worcs, WRl 3 5EY, UK

15.1

RJ. Hwu

University of Utah, Department of Electrical Engineering, Salt Lake City, Utah 84112, USA

22.3

R. Kaspi

Wright State University, University Research Center, Dayton, OH 45435, USA

16.1

CJ. Kiely

University of Liverpool, Materials Science Department, Liverpool, L69 3BX, UK

14.1

R.M. Kolbas

North Carolina State University, Department of Electrical 21.3 and Computer Engineering, Raleigh, NC 27695-7911, USA

D. Lancefield

University of Surrey, Department of Physics, Guildford, Surrey, GU2 5XH, UK

2.1-2.5

MR. Leys

COBRA Interuniversity Research Institute, Technical University of Eindhoven, Physics Department, Box 513, 5600 MB Eindhoven, The Netherlands

16.5

D C . Look

Wright State University, University Research Center, Dayton OH 45435, USA

3.7,8.1,15.7 17.2

M l . Lovejoy

Sandia National Laboratories, PO Box 5800, Albuquerque, NM 87185-5800, USA

2.13,3.6

K. Luebke

Johannes Kepler University Linz, Microelectronics Institute, Attenberger Strasse 69, A-4040 Linz, Austria

20.6

M.S. Lundstrom

Purdue University, School of Electrical & Computer Engineering, 1285 Electrical Engineering Bldg, West Lafayette, Indiana 47907-1285, USA

2.13,3.6 3.8,4.8

G B . Lush

University of Texas at El Paso, Department of Electrical and Computer Engineering, El Paso, TX 79969, USA

2.14,3.8

A. Majerfeld

University of Colorado, ECE Department, Boulder, CO 80839, USA

3.5

D K . Maude

CNRS High Magnetic Field Laboratory, 25 ave des Martyrs, BP 166 Grenoble cedex 9, France

7.5

MR. Melloch

Purdue University, School of Electrical & Computer Engineering, 1285 Electrical Engineering Bldg, West Lafayette, Indiana 47907-1285, USA

2.13,3.6 3.8,4.8

CJ. Miner

Bell Northern Research, Ottawa, Ontario, Canada, K l Y 4H7

9.11

U. Mishra

University of California, Department of Electrical Engineering, Santa Barbara, CA 93106, USA

17.3

M. Missous

UMIST, Department of Electrical Engineering and Electronics, PO Box 88, Manchester, M60 IQD, UK

14.5-14.8 17.1

C.JL. Moore

Waterloo Scientific Inc, 419 Philip Street, Unit #9, Waterloo, Ontario, Canada N2L 3XL

9.11

R. Murray

Imperial College, IRC for Semiconductors, Blackett Laboratory, Prince Consort Road, London, SW7 2BZ, UK

R. Murri

Universita di Camerino, INFM, Dipartimento di Matematica e Fisica, V. Madonna Delle Gargeri 62032, Camerino (MC), Italy

7.1

4.4

N.X. Nguyen

University of California, Department of Electrical Engineering, Santa Barbara, CA 93106, USA

17.3

D.D. Nolte

Purdue University, Department of Physics, Lafayette, Indiana 47907, USA

5.1,5.2

O. Oda

Japan Energy Corporation, Materials and Components Laboratories, 3-17-35, Niizo-Minami, Toda, Saitama, Japan

15.5,15.6

E P . O'Reilly

University of Surrey, Department of Physics, Guildford, Surrey, GU2 5XH, UK

4.1

J. Osvald

Institute of Electrical Engineering, Slovak Academy of Sciences, Dubravska cesta 9, 842 39 Bratislava, Slovakia

14.4

JE. Parrott*

University of Wales, Department of Physics and Astronomy, College of Cardiff, PO Box 913, Cardiff, CF2 3YB, UK

21.5

AR. Peaker

UMIST, Centre for Electronic Materials, PO Box 88, Manchester, M60 IQD, UK

10.3,10.5 10.6

SJ. Pearton

University of Florida, Department of Materials Science and Engineering, Rhines Hall, PO Box 116400, Gainsville, Florida 32611-6400, USA

10.8

Universita di Camerino, INFM, Dipartimento di Matematica e Fisica, V. Madonna Delle Gargeri 62032, Camerino (MC), Italy

4.4

R. Riemenschneider

Institut fur Hochfrequenztech., Fachbereich 18, 6100 Darmstadt, Merkstrasse 25, Germany

13.2-13.5 13.7

DJ. Robbins

GEC Plessey Research, Towcester, Northants, NN12 8EQ, UK

6.1-6.3

LP. Sadwick

University of Utah, Department of Electrical Engineering, Salt Lake City, Utah 84112, USA

22.3

LG. Salmon

Brigham Young University, College of Engineering and Technology, Provo, Utah 84602-4101, USA

20.7

H. Samic

University of Sarajevo, Saobracajini Fakultet, Dobrovoljacka 33, 71000 Sarajevo, Bosnia and Herzegovine

N. Pinto

16.4

E.F. Schubert

Boston University, Department of Electrical and Computer Engineering, Photonics Research Center, 44 Cummington Street, Boston, Massachusetts 02215, USA

21.2

BJ. Sealy

University of Surrey, Department of Electrical Engineering, Guildford, Surrey, GU2 5XH, UK

2.6,19.119.6

J.R. Sizelove

Wright State University, University Research Center, Dayton OH 45435, USA

3.7

K. Somogyi

Hungarian Academy of Sciences, Research Institute for Technical Physics, Foti ut 56, H-1047 Budapest, Hungary

16.3

A.L. Springer

Johannes Kepler University Linz, Microelectronics Institute, Attenberger Strasse 69, A-4040 Linz, Austria

20.6

G.E. Stillman

University of Illinois at Urbana-Champagne, Department of Electrical and Computer Engineering, Urbana, Illinois 61801, USA

4.9

S.A. Stockman

Hewlett-Packard Company, Optoelectronics Division, MS 91/ML, 370 W Trimble Road, San Jose, CA 95131, USA

3.4

Y. Sugiyama

The Electrotechnical Laboratory, 1-1-4 Umezono, Tsukuba 305, Japan

22.1

W. Szuszkiewicz

Polish Academy of Sciences, Institute of Physics, Al Lotnikow 32/46, 02-668 Warsaw, Poland

7.2

M. Tatsumi

Sumitomo Electric Industries Ltd, R&D Department, 1-1-3 Shimaya 1-chome, Konohana-ku, Osaka, 554 Japan

15.3,15.4

HW. Thim

Johannes Kepler University Linz, Microelectronics Institute, Attenberger Strasse 69, A-4040 Linz, Austria

20.6

P. Trautman

Warsaw University, Institute of Experimental Physics, Hoza 69, 00-681 Warsaw, Poland

10.2

J.A. Turner*

GEC Marconi Materials Technology Ltd, Caswell, Towcester, Northants NN12 8EQ, UK

20.4

K. Wada

NTT LSI Laboratories, 3-1 Morinosato Wakamiya, Atsugi shi, Kanagawa 243-01, Japan

16.6

R. Wallis

GEC-Marconi Materials Technology, Caswell, Towcester, Northants, NN12 8EQ, UK

20.3

W. Walukiewicz

Lawrence Berkeley National Laboratory 2-200, 1 Cyclotron Rd, Berkeley, CA 94720, USA

2.12

N. Watanabe

NTT LSI Laboratories, 3-1 Morinosato Wakamiya, Atsugi shi, Kanagawa 243-01, Japan

16.6

J. Whitaker

University of Michigan, Center for Ultra Fast Optical 17.4 Science, 2200 Bonisteel Blvd, Ann Arbor, MI48109-2099, USA

V. A. Wilkinson

DRA, 47 Sycamore Grove, Ashvale, Farnborough, Hants, UK

2.7,3.2,4.2 4.3,4.7

R.L. Williams

Institute for Microstructural Sciences, National Research Council, Montreal Road, Ottawa, KlA 0R6, Canada

2.11

E.Yablonovitch

S.Yasuami

Yung Kee Yeo

UCLA, Department of Electrical Engineering, 405 Hilgard Avenue, Los Angeles, California 90024-1594, USA Toshiba Corporation, ULSI Research Laboratories, 1 Komukai Toshiba-cho, Saiwai-ku, Kawasaki 210, Japan

16.9

10.1

Wright Patterson Air Force Base, OH-45433-7765, USA 3.1

J. Zhang

S. Zollner

* See Dedication

Imperial College, IRC for Semiconductors, Blackett Laboratory, Prince Consort Road, London, SW7 2BZ, UK Iowa State University, Ames Laboratory, A205 Physics, Ames, IA 50011,USA

31.1

4.5

DEDICATION

We are aware that between the publication of the 2nd Edition and the present edition of this book, three of our contributors have died. These are J.C. Brice J.E. Parrott J.A. Turner In the course of the preparation of a book such as this, we are in continuous contact with our contributors who become, in many cases, close colleagues with a common purpose. It is therefore particularly saddening when we hear of their passing. For this reason, we wish to dedicate this volume to their memory.

MR. Brozel G.E. Stillman

ABBREVIATIONS The following abbreviations are used in this book: A AAS AC A/D ADP AES AFM AIPC ALCHEMI ALE AM AP APB APD AR AS ASA

acceptor atomic absorption spectrometry alternating current analogue to digital (convertor) Auger depth profiles Auger electron spectroscopy atomic force microscopy ab initio pseudopotential calculation atomic location by chemical microanalysis atomic layer epitaxy air mass atmospheric pressure anti-phase boundary avalanche photodiode anti-reflection admittance spectroscopy atomic sphere approximation

BEP BERT BJT BOA

beam equivalent pressure bit error rate tester bipolar junction transistor Born-Oppenheimer approximation

CAD CADBE CAT CB CBE CBED CCD CCE CCTV CDE CL CMOS CPW CSVT CT CV CVD CW CZ

computer aided design chemically-assisted-ion-beam etching controlled atmosphere conduction band chemical beam epitaxy convergent beam electron diffraction charge coupled device charge collection efficiency closed circuit television chemical dry etching cathodoluminescence complementary metal oxide semiconductor coplanar waveguide close space vapour transport carrier transport capacitance/voltage chemical vapour deposition continuous wave Czochralski (crystal growth)

D D/A DAP

donor digital to analogue (convertor) donor acceptor pair

DBR DBS DBS DC DCFL DCXR DDLTS 2DEG DFB D-FET DFT DH 2DHG DIC DLTS DP DR DRO DT DXRD

distributed Bragg reflector direct broadcast satellite direct broadcasting systems direct current direct coupled FET logic double crystal X-ray diffraction double deep level transient spectroscopy two-dimensional electron gas distributed feedback depletion mode field effect transistor density functional theory double heterostructure two-dimensional hole gas digital integrated circuit deep level transient spectroscopy deformation potential dielectric resonator dielectric resonator oscillator dielectric theory double crystal X-ray diffraction

eB EB EBEP EBIC ECL ECR ECR-RF RIE E/D EDX EELS E-FET EL ELO EPD EPM EPR ER ESD ESR EXAFS

electron-beam-induced current electron beam electron-beam-excited plasma electron beam induced current emitter coupled logic electron cyclotron resonance electron-cyclotron-resonance radio-frequency reactive ion etching enhancement/depletion energy dispersive X-ray analysis electron-energy loss spectroscopy enhancement mode field effect transistor electroluminescence epitaxial lift-off etch pit density empirical pseudopotentials method electron paramagnetic resonance electroreflectance electron-stimulated desorption electron spin resonance extended X-ray absorption fine structure

FB FCC FEC FECTED FET FIB FIR

free to bound face centred cubic folly encapsulated Czochralski field effect controlled transferred electron device field effect transistor focused ion beam far infrared

FM FME FPA FTIR FWHM

frequency modulated flow-rate modulation epitaxy focal plane array Fourier transform infrared (spectroscopy) full width at half maximum

GDMS GF GIXD GPS GR GRIN-SCH GSMBE GW

glow discharge mass spectroscopy gradient freeze graded incidence X-ray diffraction global positioning system growth rate graded refractive index separate confinement heterostructure gas-source molecular beam epitaxy a one particle, Green's function approximation

HB HBT HE HEMT HF HFET HH HHP HOLZ HPLEC HREM HRXD HTET HZM

horizontal Bridgman heterojunction bipolar transistor Hall effect high electron mobility transistor high frequency heterostructure field effect transistor heavy hole Hall effect under hydrostatic pressure higher order Laue zone high pressure liquid encapsulated Czochralski high resolution electron microscopy high resolution X-ray diffraction high temperature electronic technique horizontal zone melting

IBAD IBAE IBE IBO IC ICB ICL IDP IGFET IHET IMA IMPATT IP IR I-S ITC IV IVS

ion beam assisted deposition ion-beam-assisted etching ion-beam etching ion-beam induced oxidation integrated circuit ionised cluster beam interface control layer intervalley deformation potential insulated-gate field effect transistor infrared hot electron transistor ion microprobe analysis impact avalanche transit-time internal photoemission infrared insulator-semiconductor inverted thermal conversion current/voltage intervalley scattering

JFET

junction field effect transistor

LACBED LAN LCAO LD LDA LEC LED LEED LEVB LH LI LLR LMTO LNA LO LP LPE LPEE LS LSI LSS LST LST LT LTG LTMBE LVM LWIR LWP

large angle convergent beam electron diffraction local area network linear combination of atomic orbitals laser diode local density approximation liquid encapsulated Czochralski light emitting diode low energy electron diffraction liquid encapsulated vertical Bridgman light hole light power/current large lattice relaxation linear muffin-tin-orbitals method low noise amplifier longitudinal optical low pressure liquid phase epitaxy liquid phase electro-epitaxy light scattering large scale integration Lindhard-ScharfF-Schiott light scattering tomography Lydane-Sachs-Teller low temperature low temperature grown low temperature molecular beam epitaxy localised vibrational modes/local vibrational mode spectroscopy long wavelength infrared light wave probe

MBE MBT MCD MCDA MCTS MEE MESFET MFP MIC MIE MIS MISFET MITATT ML MLE MMIC MMWIC MO MOCVD MODFET

molecular beam epitaxy miniband transport magnetic circular dichroism magnetic circular dichroism absorption minority carrier transient spectroscopy migration enhanced epitaxy metal semiconductor field effect transistor mean free path microwave integrated circuit magnetron-enhanced reactive ion etching metal-insulator-semiconductor metal-insulator-semiconductor field effect transistor mixed tunnel avalanche transit-time monolayer molecular layer epitaxy monolithic microwave integrated circuit monolithic millimetre wave integrated circuit metalorganic metalorganic chemical vapour deposition, see MOVPE modulation doped field effect transistor

MOMBE MOS MOSFET MOVPE MQW M-S MSI MSM MTTF MWA

metalorganic molecular beam epitaxy metal-oxide semiconductor metal-oxide semiconductor field effect transistor metalorganic vapour phase epitaxy, see MOCVD multiple quantum well metal-semiconductor medium scale integration metal-semiconductor-metal median time to failure multi-step wafer annealing

NEAT NEPM NGC NRZ NTD

noise equivalent differential temperature non-local empirical pseudopotential method normal growth conditions non-return to zero neutron transmutation doping

OA OBIC ODENDOR ODESR ODLTS OEIC OMCVD OMVPE OTCS

optical absorption optical beam induced current (spectroscopy) optically detected electron nuclear double resonance optically detected electron spin resonance optical deep level transient spectroscopy optoelectronic integrated circuit organometallic chemical vapour deposition, see MOCVD organometallic vapour phase epitaxy, see MOVPE optical transient current spectroscopy

PA PAE PC PC PC-LEC PCW PEC PECVD PES PES PIXE PL PLE PPC PR PR PR PRBS PS PSP PZR

positron annihilation power added efficiency photoconductivity photocurrent pressure controlled liquid encapsulated Czochralski photochemical washing photoelectrochemical plasma-enhanced chemical vapour deposition photoelectron spectroscopy photoemission spectroscopy particle induced X-ray emission analysis photoluminescence photoluminescence excitation persistent photoconductivity photoreflectance photoresist photoresponse pseudo-random binary signal photoemission spectroscopy pseudopotentials method piezoreflectance

QCSE

quantum-confined Stark effect

QW QWH QWIP

quantum well quantum well heterostructure quantum well infrared photodetector

RBIBE RBS RCLED RDS RE REM RF RHEED RIBE RIE RRS RT RTA RTP

radical-beam-ion-beam etching Rutherford backscattering spectroscopy resonant cavity light emitting diode reflectance diffraction spectroscopy rare earth rare earth metal radio frequency reflection high energy electron diffraction reactive ion beam etching reactive ion etching resonant Raman scattering room temperature rapid thermal annealing rapid thermal processing

SAG SAGM APD SAMAPD SAW SCH SdH SE SEED SEM SI SIMS SL SLR SO SOS SQW SRH S-S SSMS STEM STM

self aligned gate separate absorption and graded multiplication avalanche photodiode separate absorption and multiplication avalanche photodiode surface acoustic wave separate confinement heterostructure Shubnikov-de Haas selective epitaxy self electro-optic effect device scanning electron microscopy semi-insulating secondary ion mass spectroscopy superlattice small lattice relaxation spin-split-off silicon-on-sapphire single quantum well Shockley-Read-Hall semiconductor-semiconductor spark source mass spectroscopy scanning tunnelling electron microscopy scanning tunnelling microscopy

TA TBA TC TDH TEC TED TEGFET TEM

transverse acoustic tight-binding approximation thermocouple temperature dependent Hall effect thermal expansion coefficient transferred electron device two dimensional electron gas FET transmission electron microscopy

TFL TLV TM TO TOF TRIM TRPL TSC TTTC TUNNETT

trap-filled limited threshold limit value transition metal transverse optical time of flight a statistical simulation program for ion implantation time-resolved photoluminescence thermally stimulated current three-terminal thermoconverter tunnel transit-time

UHV UPS UV

ultra high vacuum ultraviolet photoelectron spectroscopy ultraviolet

VB VB VCO VCSEL VCZ VGF VLE VLSI VME VPE V-S VUV VZM

valence band vertical Bridgman voltage controlled oscillator vertical cavity surface emitting laser vapour controlled Czochralski vertical gradient freeze vapour levitation epitaxy very large scale integration vapour mixing epitaxy vapour phase epitaxy vacuum-semiconductor vacuum ultraviolet vertical zone melting

XPS XPS XRD XRP XTEM

X-ray photoelectron spectroscopy X-ray photoemission spectroscopy X-ray diffraction X-ray photoemission cross-section transmission electron microscopy

ZFTOF ZPL ZTC

zero field time of flight zero-phonon line zero temperature coefficient

Contents

Foreword .................................................................................................................

xi

Introduction ..............................................................................................................

xii

Contributing Authors ................................................................................................

xiii

Dedication ...............................................................................................................

xx

Abbreviations ...........................................................................................................

xxi

1.

2.

Basic Physical Properties of Gallium Arsenide ...........................................

1

1.1

Density of Solid GaAs ..............................................................................................

3

1.2

Lattice Parameter of GaAs .......................................................................................

8

1.3

Bulk Modulus of GaAs .............................................................................................

14

1.4

Stiffness of GaAs .....................................................................................................

16

1.5

Compliance of GaAs ................................................................................................

21

1.6

Thermal Expansion Coefficient of GaAs ..................................................................

23

1.7

Specific Heat and Debye Temperature of GaAs .....................................................

27

1.8

Thermal Conductivity of GaAs .................................................................................

32

1.9

Melting Point of GaAs ..............................................................................................

36

Electron Mobility, Diffusion and Lifetime .....................................................

39

2.1

Electron Mobility in GaAs: Overview D. Lance Field ...............................................

41

2.2

Electron Mobility in Bulk GaAs .................................................................................

46

2.3

Electron Mobility in LPE GaAs .................................................................................

48

2.4

Electron Mobility in VPE and MOVPE GaAs ...........................................................

50

2.5

Electron Mobility in MBE GaAs ................................................................................

54

2.6

Electron Mobility in Ion Implanted GaAs ..................................................................

57

2.7

Electron Mobility in GaAs, Pressure Dependence ..................................................

59

2.8

Electron Mobility Enhancement in GaAs Heterostructures .....................................

61

2.9

Ballistic Transport and Velocity Overshoot in GaAs ................................................

66

2.10 Modulation Doping ...................................................................................................

70

2.11 Carrier Concentrations Following δ-Doping .............................................................

73

2.12 Theoretical Electron Mobility/Temperature Dependences on Carrier Concentration and Compensation Ratio .................................................................

78

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v

vi

3.

4.

5.

6.

7.

Contents 2.13 Minority Electron Mobility in Doped GaAs ...............................................................

81

2.14 Electron Lifetimes in p-Type GaAs ..........................................................................

87

Hole Mobility, Diffusion and Lifetime ............................................................

89

3.1

Hole Mobility in Doped and Ion Implanted GaAs ....................................................

91

3.2

Hole Mobility in GaAs, Temperature Dependence ..................................................

98

3.3

Hole Mobility in GaAs, Pressure Dependence ........................................................

100

3.4

Carbon Doping of GaAs ...........................................................................................

101

3.5

Carrier Concentration Dependence of the Hole Mobility in GaAs ...........................

117

3.6

Minority Hole Mobility in GaAs .................................................................................

123

3.7

Theoretical Hole Mobility Curves .............................................................................

131

3.8

Hole Lifetimes in n-Type GaAs ................................................................................

135

Band Structure and Carrier Ionization .......................................................... 143 4.1

Band Structure of GaAs: Overview ..........................................................................

145

4.2

Direct Bandgaps of GaAs, Temperature Dependence ...........................................

151

4.3

Direct Bandgap of GaAs, Pressure Dependence ....................................................

153

4.4

Optical Properties of Amorphous GaAs ...................................................................

155

4.5

Intra- and Intervalley Deformation Potentials for Electrons in GaAs .......................

162

4.6

Intravalley Deformation Potentials for Holes in GaAs .............................................

177

4.7

Electron Effective Mass and Hole Effective Mass in GaAs, Pressure Dependence .............................................................................................................

184

4.8

Effective Bandgap Narrowing in Doped GaAs ........................................................

186

4.9

Carrier Ionization Coefficients of GaAs ...................................................................

190

Optical Functions of Gallium Arsenide ......................................................... 199 5.1

Optical Functions of GaAs .......................................................................................

201

5.2

Table of GaAs Optical Functions at 300 K ..............................................................

207

Electro-Optic Properties of Gallium Arsenide .............................................. 215 6.1

Electro-Optical Coefficients of GaAs .......................................................................

217

6.2

Franz-Keldysh Effect in GaAs ..................................................................................

220

6.3

Piezoelectric Properties of GaAs .............................................................................

223

Infrared Absorption and Energy Levels Due to Impurities ......................... 225 7.1

IR Absorption Bands Due to Localized Vibrational Modes (LVM) of Impurities in Bulk and Epitaxial GaAs ......................................................................

227

7.2

IR Absorption Due to Free Carriers in GaAs ...........................................................

235

7.3

Electronic Absorption Bands of Impurities and Defects in GaAs ............................

244

7.4

Energy Levels Due to Transition Metals in GaAs ....................................................

248

7.5

DX Centres in GaAs .................................................................................................

250

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Contents 8.

9.

vii

Photoconductivity Spectra ............................................................................ 257 8.1

IR Photoconductivity Spectra of SI Bulk GaAs ........................................................

259

8.2

Far-IR Photoconductivity Spectra of Shallow Donors in GaAs Epilayers ...............

262

Photoluminescence Spectra of Gallium Arsenide ....................................... 271 9.1

Photoluminescence Spectra of Undoped SI GaAs .................................................

273

9.2

Photoluminescence Spectra of Undoped LEG p-Type GaAs EMIS Group ..............................................................................................................

276

9.3

Photoluminescence Spectra of LPE GaAs ..............................................................

278

9.4

Photoluminescence Spectra of WE GaAs ...............................................................

282

9.5

Photoluminescence Spectra of MOVPE (MOCVD) GaAs ......................................

286

9.6

Photoluminescence Spectra of MBE GaAs .............................................................

291

9.7

Photoluminescence Spectra of Group II Atoms in GaAs ........................................

298

9.8

Photoluminescence Spectra of Group IV Atoms in GaAs .......................................

302

9.9

Photoluminescence Spectra of Group VI Shallow Donors in GaAs .......................

308

9.10 Photoluminescence Spectra of Transition and Rare Earth Metals in GaAs ...........

310

9.11 Room Temperature Photoluminescence Mapping of GaAs Substrates and Epitaxial Layers ........................................................................................................

320

10. Defects, Deep Levels and Their Detection ................................................... 333 10.1 Structure of Native Defects in Undoped SI LEG GaAs ...........................................

335

10.2 Energy Levels and Fundamental Properties of the EL2 Defect in GaAs ................

341

10.3 Electrical Techniques for the Measurement of Deep State Properties ...................

358

10.4 Defect Densities in Melt-Grown GaAs (a Review) ..................................................

371

10.5 Deep States in As-Grown Epitaxial GaAs ...............................................................

380

10.6 Defect Energy Levels in Electron and Neutron Irradiated Photon Damaged and Ion Implanted GaAs ..........................................................................................

394

10.7 Analysis of LEG GaAs by Near-IR Mapping ............................................................

399

10.8 Passivation of Defects in GaAs by Hydrogenation ..................................................

409

10.9 Transmission Electron Microscopy of GaAs ............................................................

418

11. Diffusion of Impurities .................................................................................... 429 11.1 Diffusion of Impurities in GaAs: Introduction ...........................................................

431

11.2 Diffusion of Shallow Acceptors in GaAs ..................................................................

432

11.3 Diffusion of Shallow Donors in GaAs .......................................................................

438

11.4 Cr Diffusion in GaAs .................................................................................................

442

12. Surface Passivation ........................................................................................ 445 12.1 Technical Aspects of Surface Passivation in GaAs ................................................

447

12.2 Fundamental Aspects of Surface Passivation in GaAs ...........................................

449

12.3 New Approaches for Surface Passivation in GaAs .................................................

451

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viii

Contents

13. Surface Structure and Oxidation ................................................................... 453 13.1 Surface Structure of GaAs .......................................................................................

455

13.2 Oxide Layer Structures of GaAs ..............................................................................

463

13.3 Thermal Oxidation of GaAs .....................................................................................

470

13.4 Wet Oxidation of GaAs ............................................................................................

477

13.5 Plasma Oxidation of GaAs .......................................................................................

482

13.6 Laser-Assisted Oxidation of GaAs ...........................................................................

486

13.7 Miscellaneous New Methods of GaAs Oxidation ....................................................

489

14. Interfaces and Contacts ................................................................................. 497 14.1 The Structure of the Epitaxial Al/GaAs Interface .....................................................

499

14.2 Structure of the Au/GaAs Interface ..........................................................................

505

14.3 Structure of Au-Ge/GaAs Interfaces ........................................................................

510

14.4 Structure of Silicide/GaAs Interfaces .......................................................................

516

14.5 Barrier Height at the GaAs/Al Interface ...................................................................

521

14.6 Barrier Height at the Ag/GaAs Interface ..................................................................

530

14.7 Barrier Height at the GaAs/Au interface ..................................................................

534

14.8 Barrier Height at the Pt/GaAs and W/GaAs Interfaces ...........................................

539

14.9 Conduction and Valence Band Offsets at the AlGaAs/GaAs and (Al)GaInP/GaAs Heterostructure Interfaces ............................................................

545

15. Bulk Growth of GaAs ...................................................................................... 555 15.1 Thermodynamics of the Ga/As System ...................................................................

557

15.2 LEG Growth of GaAs ...............................................................................................

565

15.3 Horizontal and Vertical Bridgman Growth of GaAs .................................................

572

15.4 New Types of LEG Growth and Vapour Controlled LEG GaAs ..............................

579

15.5 Heat Treatments of GaAs Ingots .............................................................................

585

15.6 Heat Treatments of GaAs Wafers ...........................................................................

591

15.7 Carrier Concentrations of Semi-Insulating GaAs ....................................................

596

16. Epitaxial Growth of GaAs ............................................................................... 599 16.1 Introduction to Epitaxy ..............................................................................................

601

16.2 LPE Growth of GaAs ................................................................................................

608

16.3 Halogen Transport VPE Growth of GaAs and Related Compounds ......................

625

16.4 Growth by Close Space Vapour Transport ..............................................................

639

16.5 MOVPE Growth of GaAs .........................................................................................

643

16.6 Unintentional Hydrogen Doping in MOCVD Growth ...............................................

651

16.7 High Temperature MBE Growth of GaAs ................................................................

655

16.8 Metalorganic Molecular Beam Epitaxy and Atomic Layer Epitaxy of GaAs ...........

663

16.9 Epitaxial GaAs Lift-Off ..............................................................................................

672

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Contents

ix

17. Low Temperature MBE GaAs ......................................................................... 677 17.1 Low Temperature GaAs: Growth Dynamics and Effects of As:Ga Flux Ratio .........................................................................................................................

679

17.2 Point Defects in LT GaAs .........................................................................................

684

17.3 Uses of Low Temperature Grown Materials in MESFETs ......................................

689

17.4 Opto-Switches Using Low-Temperature MBE GaAs ..............................................

693

18. Etching ............................................................................................................ 703 18.1 Etching of GaAs: an Overview .................................................................................

705

18.2 Wet Etching of GaAs ................................................................................................

707

18.3 Chemical Dry Etching of GaAs ................................................................................

714

18.4 Plasma Etching of GaAs ..........................................................................................

719

18.5 Ion-Beam Milling and Sputter Etching of GaAs .......................................................

722

18.6 Reactive Ion Etching (RE) and Magnetron-Enhanced Reactive Ion Etching (MIE) of GaAs ...........................................................................................................

725

18.7 Etching of GaAs Using ECR-RF Reactive-Ion Etching ...........................................

741

18.8 Etching of GaAs Using Reactive-Ion-Beam Etching (RIBE), ChemicallyAssisted-Ion-Beam Etching (CAIBE), and Radical-Beam-Ion-Beam Etching (RBIBE) .......................................................................................................

751

19. Ion Implantation and Rapid Thermal Annealing ........................................... 763 19.1 Ion Implantation in GaAs: an Overview ...................................................................

765

19.2 Ion Ranges in GaAs: Discussion .............................................................................

770

19.3 Rapid Thermal Annealing of GaAs: Overview .........................................................

774

19.4 Maximum Concentrations and Activation Efficiencies for Each Ion Species .....................................................................................................................

776

19.5 Controlled Atmosphere Annealing of GaAs .............................................................

779

19.6 Compensation Mechanisms in GaAs at High Dopant Fluences .............................

781

20. Exploitation of GaAs in Microwave and High Speed Digital Circuits ............................................................................................................ 783 20.1 GaAs MESFET: Discrete, Power and MMIC Microwave Devices ..........................

785

20.2 GaAs MESFET: Digital and Opto-ICs ......................................................................

799

20.3 High Electron Mobility Transistors ...........................................................................

811

20.4 GaAs in Bipolar Integrated Circuits ..........................................................................

820

20.5 GaAs MMWICs ........................................................................................................

824

20.6 The Transferred Electron Effect ...............................................................................

830

20.7 Back-Gating in GaAs MESFET Integrated Circuits .................................................

836

20.8 GaAs IMPATT Diodes .............................................................................................

851

21. GaAs in Optoelectronics ................................................................................ 859 21.1 GaAs Based Heterostructures for Optoelectronic Devices ..................................... This page has been reformatted by Knovel to provide easier navigation.

861

x

Contents 21.2 GaAs Light Emitting Diodes .....................................................................................

874

21.3 Quantum Well Heterostructure Lasers ....................................................................

887

21.4 GaAs Mid- and Far- Infrared Detectors ...................................................................

906

21.5 The GaAs Solar Cell ................................................................................................

918

21.6 GaAs Optoelectronic Integrated Circuits and Future Applications ..........................

925

22. GaAs in Other Applications ........................................................................... 931 22.1 Sensors ....................................................................................................................

933

22.2 GaAs Nuclear Particle Detectors .............................................................................

942

22.3 High Temperature Applications for GaAs ................................................................

948

Index ....................................................................................................................... 963

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CHAPTER 1 BASIC PHYSICAL PROPERTIES OF GALLIUM ARSENIDE 1.1 1.2 1.3 1.4 1.5 1.6 1.7 1.8 1.9

Density of solid GaAs Lattice parameter of GaAs Bulk modulus of GaAs Stiffness of GaAs Compliance of GaAs Thermal expansion coefficient of GaAs Specific heat and Debye temperature of GaAs Thermal conductivity of GaAs Melting point of GaAs

1.1

Density of solid GaAs J.C. Brice February 1990 (Revised by M.R. Brozel, May 1996)

A

INTRODUCTION

Anyone who has tried to measure the density of a solid knows that there are many practical problems and that to achieve errors of less than 1 mg/cm3 is very difficult. Thus it is not surprising that there have been few precise measurements of the density of GaAs. The available data [1-4] show that while at any fixed temperature there should be one value for the density of pure perfect stoichiometric GaAs, the densities (D) of real samples lie in the range 5.3150 < D < 5.3180 g/cm3 at 300 K

(Ia)

5.3152 < D < 5.3182 g/cm3 at 25°C

(Ib)

or

B

UNDOPED STOICHIOMETRIC SAMPLES AT 300 K

At 300 K, the X-ray lattice parameter of ideal GaAs5 a0, is 5.6536 A with a last digit uncertainty of about ±6 [5]. From this the X-ray density is D (X-ray) = 5.3168 g/cm3

(2a)

with a last digit uncertainty of ±2. Because the vacancy concentration in GaAs is small [1,2,6], this should be a good estimate of the density of pure perfect GaAs. For comparison the measured values of D from refs [4], [2] and [3] are respectively 5.3170, 5.3164 and 5.3160 g/cm3, which give a mean for undoped stoichiometric GaAs of D (stoichiometric, undoped) = 5.3165 ± 0.0002 g/cm3

(2b)

Because the average sample contains about 0.2 atomic ppm of point defects and up to lOVcm2 dislocations [7-9], EQN (2b) is not directly comparable to EQN (2a). An estimate of the density of ideal GaAs is given later, where it is shown that the imperfections just mentioned could decrease the density of ideal GaAs by about 0.0002 g/cm3 to give an X-ray based value of 5.3166 g/cm3, so that EQN (2b) is the best estimate of the density of real undoped crystals. C

THE EFFECT OF IMPURITIES

Impurities and dopants can change the density of a crystal in two ways: (a)

They can change the mass in the average unit cell, increasing it by addition in the case of an interstitial atom, or in either direction in the case of a substitutional atom.

(b)

They can change the volume of the average unit cell.

Often the effects cancel. Heavy substitutional atoms increase the mass of the cell and, usually, the lattice constant. The first effect increases the density: the second decreases it. Light atoms tend to have reversed effects: interstitial atoms increase both the mass and the lattice parameter. For GaAs we have measured data on the effect of impurities on the lattice parameter [5]. These are given in TABLE 1 where columns 2, 3 and 4 represent the change in ao per unit change in impurity concentration (cm"3), the change in D per unit change in impurity concentration resulting from the observed change in a0, and the change in D per unit change in impurity concentration resulting from the introduction of atoms of different mass, respectively. Column 5 is the total change and is the sum of columns 3 and 4. The calculated effect on density given in the second column is questioned by -(3D/ao)(dao/dC). In order to calculate the effect on mass (8Am. a0"3 where Am is the mass change and the 8 arises from the number of atoms per cell), we need to know where the atoms go. Here it is assumed that Al, B, Be, Cr and In substitute on gallium sites; S5 Se and Te substitute on arsenic sites; while C, Si, Ge and Sn substitute equally on both sites. Actually, we know that Si is primarily a donor so that most Si goes on gallium sites, but the error due to this is small. Clearly these estimates are inaccurate but they do demonstrate that the two effects usually cancel. A more important observation is that the effects are in general small: the units are micrograms per cubic centimetre per atomic part per million. For most of the elements, 25 to 50 atomic ppm (i.e. 1 to 2 x 1018/cm3 atoms ) are needed to produce a significant density change (say 0.0002 g/cm3). Yarmolyuk et al [3] give experimental data to support this. Chromium and oxygen are the exceptions which have large values of dD/dC. D

THE EFFECT OF DISLOCATIONS

Dislocations can affect the mass of unit volume of a crystal, but a moment's reflection on the number of atoms per unit surface area shows that dislocation densities of millions per square centimetre produce changes in the region of parts per billion. So this effect is negligible. However, dislocations increase the lattice parameter a0. IfN is the dislocation density, then it has been suggested that the increase in a0, A a0, is Aa o = 0.015/Nfm[5] and the decrease in density is dD = -4.2* 10" 7 /N so that a typical density (5 x lOVcm2) decreases D by about 0.0001 g/cm3.

(3)

TABLE 1. The effect of impurities on density. Element

dao/dC

dD/dC* due to a0

dD/dC* due to Am

Total

Al

O

0

-3.15

-3.15

B

-0.32±0.04

+9.03

-4.33

+4.70

Be

-0.16

+4.51

-4.46

+0.05

C

<-0.07

+1.97

-4.43

-2.46

Cr

-2.OH/4

+56.4

-1.30

+55.1

Ge

+0.20±0.08

-5.64

-0.20

-5.44

In

+0.50±0.06

-14.1

+3.31

-10.8

S

-K).04±0.04

-1.13

-3.14

-4.27

Se

+0.14

-3.95

40.30

-3.65

Si

-0.16±0.04

+4.51

-3.28

+1.23

Sn

+0.4±0.2

-11.3

+3.41

-7.89

Te

-K).3±0.1

-8.46

+3.87

-4.59

* The units of dD/dC are micrograms per cubic centimetre per atomic part per million.

E

THE EFFECT OF STOICHIOMETRY

The few data available suggest that compared to typical undoped samples (D = 5.3165 g/cm3), arsenic-rich samples have larger densities: values of 5.3169 and 5.3175 g/cm3 have been quoted [2,1]. The second of these samples is arsenic-saturated. For gallium-rich samples, only one value has been measured. That was D = 5.3158 g/cm3. F

THE EFFECT OF TEMPERATURE

Both the density and the thermal expansion coefficient of GaAs vary with stoichiometry and purity [10]. For small changes of temperature near to 300 K, it can be assumed that dD/dT = -0.000096 g/cm3

(4)

Values of the density at other temperatures are given in TABLE 2. These values are for undoped typical samples. TABLE 3 gives the uncertainties at various temperatures. These include variations from all causes.

IABLE 2. Density (D) at various temperatures. T(K)

0-100

150

200

250

300

400

D (g/cm3)

5.332

5.329

5.332

5.322

5.3165

5.307

T(K)

500

600

700

800

900

1000

D (g/cm3)

5.298

5.289

5.279

5.269

5.259

5.249

T(K)

1100

1200

1300

1400

1500

D (g/cm3)

5.239

5.228

5.217

5.207

5.196

TABLE 3. The uncertainties (AD, g/cm3) at T K. T(K)

0

300

600

900

1200

1500

AD

0.004

0.0015

0.005

0.010

0.015

0.020

G

THE EFFECT OF IRRADIATION

Kolin et al [4] show that neutron irradiation causes the density to fall by an amount proportional to the dose. For a dose of 3.5 * 1019/cm2 the density falls to 5.300 g/cm3. Annealing at any temperature in the range 500 to 9000C restores the density to its unirradiated value. H

THE DENSITY OF IDEAL GaAs

Atypical sample ofundoped GaAs has a density of 5.3165 g/cm3. Allowing, as discussed above, for the effects of impurities and dislocations gives a value for ideal GaAs of 5.3167 g/cm3 which is close to the X-ray estimate of 5.3168 g/cm3. I

CONCLUSION

The density of pure, dislocation-free, stoichiometric gallium arsenide should be about 5.3167 g/cm3 and data are given in Sections C and D which allow correction of this value to take account of impurities (and dopants) and dislocations. However, these corrections are small and for many purposes the value of the density of a typical lightly-doped (i.e. < 1018/cm3) sample (5.3165 g/cm3) is a more useful value. The situation with respect to non-stoichiometric samples is less satisfactory. We can be sure that the density increases with increasing arsenic content, but we cannot, with the few data available, quantify the effect.

REFERENCES [1] [2] [3] [4] [5] [6] [7] [8] [9] [ 10]

M.E. Straumanis, CD. Kim [Acta Crystallogr. (Denmark) vol.19 (1965) p.256 ] Y.T. Bublik et al [ Sov. Phys.-Crystallogr. (USA) vol. 18 no.2 (1973) p.218-9 ] N.I.Yarmolyuk et al [ Sov. Phys.-Semicond. (USA) vol. 14 no.7 (1980) p.773-5 ] N.G. Kolin et al [ Phys. Chem. Mater. Treat. (UK) vol.21 (1987) p.223 ] J.C. Brice [ Datareview in this book: 1.2 Lattice parameter of GaAs ] D.T.J. Hurle [ J. Phys. Chem. Solids (UK) vol.40 no.8 (1979) p.613-26 ] J.C. Brice et al [ J. Mater. Sci. (UK) vol.2 (1967) p. 131 ] J.C. Brice [ Nature (UK) vol.209 (1966) p. 1346 ] J.C. Brice [ J. Cryst. Growth (Netherlands) vol.7 (1970) p.9 ] J.C. Brice [ in Properties of Gallium Arsenide, 2nd Edition (DSfSPEC, IEE, London, UK, 1990) ch. 1 p. 18-9]

1.2

Lattice parameter of GaAs J.C. Brice March 1990 (updated October 1995 by DJ. Dunstan)

A

DfTRODUCTION

This Datareview now contains all the data relevant to the lattice parameter of GaAs. Previously this information was spread over two Datareviews [1,2]. The lattice parameter of GaAs has been widely studied [5-23,36,37]. A dozen precise X-ray values for stoichiometric crystals have been reported. Some work has also been done on non-stoichiometric samples. It is also possible to deduce lattice parameters from precise density data which are reviewed in [3]. Here X-ray results are reported on the basis that CuK (al) radiation has a wavelength of 1.537400 kXu = 0.1540592 nm (X-ray wavelengths are uncertain by about 10 ppm [4]). When necessary, values of the lattice parameter have been converted to those appropriate to 25 0 C using a thermal expansion coefficient of 6.02 ppm/K [24]. Values of the lattice parameter (a0) are given in nanometres while deviations are given in femtometres (106 fin = 1 nm). B

HOMOGENEITY

A major problem in determining the lattice parameter of GaAs is that most samples are inhomogeneous. In general, small (< 1 cm) thin (< 1 mm) slices are reasonably homogeneous: for example, variations of ± 3 fin are reported [20]. Large undoped slices are much less homogeneous: typical data [17,25] indicate variations of about ± 20 fin. Large slices of doped material can show variations of up to ± 50 fin [17,25]. If slices are cut into dies a few mm on edge, the variation falls typically by a factor of 10 [20]. C

THE FACET EFFECT

One cause of inhomogeneity (but by no means the only one) is the facet effect in which a (111) facet in the growth face grows by a different mechanism to the other parts of the face. For undoped crystals this adds about 3 fin to a0. For doped crystals, the effect is larger and 7 ± 3 fin should be added to a0. See for example [13,14]. D

SURFACE DAMAGE

References [5] and [8] show that mechanical damage increases the lattice constant near to the surface. Relative to an etched surface the following increases (in fin) were noted: sawn surface surface lapped with 600 grit mechanically polished surface lightly scratched surface

+50 +30 +20 +30

E

LATTICE PARAMETERS OF UNDOPED SAMPLES

The twelve precise X-ray values found for stoichiometric crystals give a mean value of a0 = 0.565361 nm

(1)

and the standard deviation of the data is 7 fin. Adding the parameters calculated from density data [3] does not affect the standard deviation but reduces the mean to a0 = 0.565360 nm

(2)

Note that we do not know much about the purity or perfection of the samples measured. However, bulk material purity and crystallinity has of course been improving in recent years. Two recent values not included in the above analysis are 0.565351 nm for CZ GaAs [36], and 0.565370 nm for modified LEC material with measured (low) impurity concentrations and dislocations below 104 cm"2 [37]. However, we may treat these displayed values as good estimates of the values expected for undoped samples, with a likely error of ±10 fm. Estimation of the lattice parameter of pure perfect GaAs is left until the dependences of a0 have been discussed. F

VARIATION WITH TEMPERATURE

From the expansion coefficient at 250 C (6.02 ppm/K [24]) we can deduce that near to room temperature a0 increases by 3.4 fm per degree increase in temperature. However, the coefficient is a function of temperature which shows a rapid change particularly at low temperatures so that values remote from 250 C must be calculated using the data on the coefficient or the dilation which is given in [24]. G

VARIATION WITH DOPING AND PURITY

Based on data in [5,9,10,12,14,26], the earlier Datareview published in the 1986 EMIS book [2] gave data on the variation caused by Si, Te, Ge, Sn, Pb, O, S and Cr. We can now add data on Si from stoichiometric melts [18], In [27], B [20], C and Be [28] (see also [29]). Data are reported here as the rate of change of a0 (in fin) with concentration C (in atomic ppm). One atomic ppm corresponds to 4.4 x 1016 atoms per cm3. TABLE 1 summarises the data for silicon. TABLE 1. The effect of silicon doping. Crystal Ga

dao/dC (fin/atomic ppm) rich

Stoichiometric As

-0.50 ±0.05 -0.16 ± 0.04

rich

+0.20 ± 0.20

'Rich' is not well defined but implies a crystal grown from a melt differing by more than 2% from the near 50% concentration which produces stoichiometric crystals.

TABLE 2 gives data for other solutes. Because less information is available, we do not know the effects of stoichiometry except for B which behaves like Si (see note (d)). But the large uncertainties in the other data suggest that stoichiometry may have a significant effect. The tabulated values should be near to those for stoichiometric material. It would be reasonable to expect carbon to behave like silicon. TABLE 2. The effects of various solutes. Solute

dao/dC (fin/atomic ppm)

B

-0.32 ± 0.04 (d)

Be

-0.16 (a)

C

<-0.07 (b)

Cr

-2.0 ±1.4

Ge

40.20 ± 0.08

In

-K).50 ± 0.06

O

+1.4±1.0

S

4O.04 ± 0.04

Si

-0.16 ±0.04

Sn

40.4 ± 0.2

Te

40.3 ±0.1 (c)

Notes (a)

Since only one value is given, no estimate of the uncertainty can be made.

(b)

Reference [28] shows that for concentrations over 500 ppm, ao = 0.56533 nm corresponding to a contraction of about 30 fm, so we can only give a limit. The rate of contraction may well exceed 0.07 fin/atomic ppm.

(c)

Reference [26] shows that the anomalously large dilation due to Te at C > 50 ppm can be relieved by annealing. The data are for annealed samples.

(d)

Boron in gallium rich crystals seems to produce a larger change than in stoichiometric crystals. For Ga rich crystals, da^dC = -0.44 fm/atomic ppm. For As rich samples dao/dC could be +0.3 fm/atomic ppm.

Excluding the data for O and Cr, within each group in the periodic table the values of da^dC correlate well with the covalent radii of the solutes [33], On this basis, we can estimate the two obvious omissions and find values of da
H

THE EFFECT OF DISLOCATIONS

The effect of dislocations has not been widely studied but there is little doubt that the presence of dislocations increases the lattice parameter. Typical results [17,22] suggest that a dislocation density of 1 * 104 /cm2 increases the lattice parameter by 1 to 2 fm, which is in accord with simple theory which predicts that Aa0 = 0.015 VN

(3)

where Aa0 is the increase in lattice parameter (in fin) produced by a dislocation density N (in number per cm2). The argument used is the one used to predict misfit dislocations where the density N = 4(Aa0Zb^0)2 where b is the Burger's vector. The square root dependence explains why crystals with large dislocation densities do not have enormous lattice constants: for example if N is one million, Aa0 is 15 fm which is only observable with careful measurements. However, further work in this area is vital: in particular, the squareroot relation needs to be tested rigorously. I

VARIATION WITH STOICHIOMETRY

It is important to recognise that if we grow a batch of crystals and change only the arsenic content of the melt (or equivalently the arsenic partial pressure in the growth system), the crystals are not identical. First, as Brice et al [30] showed in the case of silicon doping, the impurity contents are not all the same: increasing the arsenic partial pressure from 700 to 760 torr caused the silicon content to fall by a factor of 10. Second, as Brice showed for both pulled [31] and Bridgman crystals [32] the dislocation density increases with increasing arsenic partial pressure. Taking into account the likely impurities (Si, C, B) and their likely concentrations, these two effects might well contribute to an increase in lattice parameter in the range 3 to 30 fm between gallium-rich and arsenic-rich crystals. This potentially large indirect change makes it difficult to isolate a direct change caused by a change in stoichiometry, particularly if the change has the opposite sign to that just mentioned. Nevertheless, an earlier version of this Datareview [1] did suggest that the lattice parameter might fall as the arsenic content rises. The difference ao(As-rich) - ao(Ga-rich) was 10 (± 9) fm. Kuwamoto and Holmes [23] also found a negative difference of -15 (± 2) fm. However, all the other reports have given positive differences: +30 (± 14) fin [16], +2 (± 2) fin [20], +7 (± 4) fin [19] and +4 (± 3) fin [21]. Using the reciprocal of the standard deviation of the difference as a weighting factor gives a mean difference of -1.3 (± 3.4) fin, which looks like evidence that there is no difference. However, if we subtract an estimate of the likely indirect effect as discussed above, say +10 (± 5) fin, the mean difference becomes -11.3 (± 6) fin. A one-tail t-test ascribes over 90% confidence to this. Further supporting evidence comes from using lattice parameters calculated from density data [34] which gives a difference of -38 (± 1) fin. Some of this difference could be attributed to the effect of vacancies on the mass in a unit cell but to account for the total effect in this way involves vacancies on more than 0.04% of the gallium sites in gallium-rich crystals. Theory [35] and experiment [15,34] suggest that such a vacancy concentration is unlikely: expected vacancy concentrations are at the ppm level. The most recent results [36] show no effect of melt composition up to ±2% from stoichiometry.

A decrease of lattice constant with an increase in arsenic concentration remains possible and we can quantify this by saying that near to the maximum melting point, a 1 atomic % increase in arsenic in the melt produces a contraction, Aa0, of 0 to 9 fm. Note that the rate of change of lattice parameter with arsenic concentration in the melt must vary: it must go to zero at the melt concentrations corresponding to the extreme solid concentrations. The solid compositions are usually derived experimentally on the basis of lattice parameter and density measurements [15,34]. The calculated values [35] are believed to be reliable. J

THE LATTICE PARAMETER OF IDEAL GaAs

In Section E, it was noted that the X-ray data gave an average lattice parameter for an undoped crystal of 0.565361 nm. From the available data (see for example [30-32,17,22]) we know that an average crystal contains less than 1 atomic ppm of impurities and point defects and has a dislocation density of about 5 x lOVcm2. From the data given in Sections G and H, this crystal should have a lattice parameter about 8 fin greater than an ideal crystal. Taking into account the most recent data [36,37], we deduced that an ideal crystal has a lattice parameter of a0 = 0.565360 nm

(4)

The lattice parameter deduced from the densities of stoichiometric crystals are 0.565344 and 0.565359 nm. These should be insensitive to defects and are therefore good estimates of a0 for ideal crystals. Thus our best estimate of the lattice parameter of ideal GaAs is a0 = 0.565360 nm

(5)

with a standard error of about 10 fin.

K

CONCLUSION

This Datareview shows that the lattice parameter of pure perfect stoichiometric GaAs is about 0.56536(1) nm at 25 0 C. Quantitative estimates of the effects of impurities (and dopants), variation of stoichiometry and dislocation densities are given. These allow estimation of the parameters of characterised real crystals with uncertainties of 7 to 10 fm. Real crystals tend to have larger parameters than ideal ones; for crystals which are not heavily doped it is reasonable to assume a parameter of 0.56536 nm with an uncertainty of 1 or perhaps 2 in the last digit. REFERENCES [1] [2] [3] [4] [5]

J.C. Brice [in Properties of Gallium Arsenide, EMIS Datareview Series No.2, (INSPEC5IEE5 1986) Datareview 1.2] J.C. Brice [in Properties of Gallium Arsenide, EMIS Datareview Series No.25 (INSPEC5IEE5 1986) Datareview 1.3] J.C. Brice [in Properties of Gallium Arsenide, 2nd Edition, (INSPEC5IEE5 London, UK5 1990) ch. 1 p.3-6] A. Bearden, J.S. Thomsen [in AIP Handbook, Ed. D.E. Gray (McGraw-Hill, New York, 1972) p.7-96] C.H.M. Driscoll et al [ Inst. Phys. Conf Ser. (UK) no.24 (1975) p.275 ]

[6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22]

[23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33] [34] [35] [36] [37] [38]

A.S. Cooper [ Acta Ctystallogr. (Denmark) vol.15 (1962) p.578 ] E.D. Pierron et al [ Acta Crystallogr. (Denmark) vol.21 (1966) p.290 ] E.D. Pierron, J.B. McNeely [Adv. X-ray Anal. (USA) vol.12 (1969) p.343 ] J.F.C.Baker et al [ Solid-State Electron. (UK) vol.19 no.4 (1976) p.331-4 ] R. Heckingbottom et al [ Solid-State Electron. (UK) vol.19 no.4 (1976) p.335-9 ] CV. Ozolinsh et al [ Sov. Phys.-Crystallogr. (USA) vol.7 (1963) p.691 ] J.B. Mullin et al [ J. Appl. Phys. (USA) vol.47 no.6 (1976) p.2584-7 ] P.F. Fewster, A.F.W. Willoughby [ EMS Original Data RN=500 (Sept 1980) ] P.F. Fewster, A.F.W. Willoughby [ EMS Original Data RN=501 (Sept 1980) ] M.E. Straumanis, CD. Kim [Acta Crystallogr. (Denmark) vol.19 (1956) p.256 ] Y. Takano et al [ Jpn. J. Appl. Phys. 2 (Japan) vol.24 no.4 (1985) p.L239-40 ] T. Fukumori, K. Futagami [ Jpn. J Appl. Phys. I (Japan) vol.27 no.3 (1988) p.442-3 ] G.P. Watson, D. Ast, A.G. Elliot [Appl. Phys. Lett. (USA) vol.54 no.3 (16 Jan 1989) p.271-3 ] M. Nakajima et al [Appl. Phys. Lett. (USA) vol.49 no.19 (10 Nov 1986) p. 1251-4 ] Y. Okada et al [ Appl. Phys. Lett. (USA) vol.48 no. 15 (14 Apr 1986) p.975-7 ] Y. Okada, F. Orito [Appl. Phys. Lett. (USA) vol.52 no.7 (15 Feb 1988) p.582-3 ] K. Usuda, S. Yasuami, Y. Higashi, H. Kawata, M. Ando [ExtendedAbstracts of 19th Con/. Solid State Devices and Materials, Tokyo, Japan, 25-27 Aug. 1987 (Bus. Center Acad. Soc. Japan, Tokyo, 1987) p. 119-22] H. Kuwamoto, D.E. Holmes [ J. Appl. Phys. (USA) vol.59 no.2 (15 Jan 1986) p.656-7 ] J.C Brice [ in Properties of Gallium Arsenide, 2nd Edition (DSfSPEC, IEE, London, UK, 1990) ch. 1 p.18-9] SJ. Barnett, GJ. Brown, B.K. Tanner [ Inst. Phys. Conf. Ser. (UK) no.87 (1987) p.615-20 ] P.S. Dobson et al [ Inst. Phys. Conf. Ser. (UK) no.45 (1979) p. 163 ] T. Takabe et al [ Inst. Phys. Conf. Ser. (UK) no.79 (1986) p.283 ] K. Saito et al [ J. Appl. Phys. (USA) vol.64 no.8 (15 Oct 1988) p.3975-9 ] J. Bak-Misiuk et al [ Phys. Status Solidi A (Germany) vol. 106 no.2 (1988) p.451 -8 ] J.C. Brice, J.A. Roberts, G. Smith [ J. Mater. Sd. (UK) vol.2 (1967) p. 131 ] J.C. Brice [ J. Cryst. Growth (Netherlands) vol.7 (1970) p.9 ] J.C Brice [ Nature (UK) vol.209 (1966) p. 1346 ] R.D. Harrison [ Book of Data (Longman, London, 1972) p.54 ] V.T. Bublik et al [ Sov. Phys.-Crystallogr. (USA) vol. 18 no.2 (1973) p.218-9 ] D.T.J. Hurle [ J. Phys. Chem. Solids (UK) vol.40 (1979) p.613 ] M. Fatemi [ J. Cryst. Growth (Netherlands) vol.96 (1989) p.316 ] K. Usuda, S. Yasuami, T. Fujii, Y. Higashi, H. Kawata, M. Ando [ Inst. Phys. Conf. Ser. (UK) no.l06(1990)p.l3] O. Madelung (Ed)[ Semiconductors: Group IVElements and III- V Compounds, (Springer-Verlag,

Berlin, 1991) p.151]

1.3

Bulk modulus of GaAs DJ. Dunstan September 1995

A

ISOTHERMAL AND ADIABATIC BULK MODULI

The bulk modulus B is defined by SP C11 + 2c19 B = V = — -

(1)

3

av

and is most accurately obtained from the stiflhess constants, qj3 whose values are given in the next Datareview [I]. Several authors cited there for the Cy data also gave derived bulk moduli; there is no need to review them all since they are derived values. The recommended adiabatic (Bs) and isothermal (BT) values in TABLE 1 are therefore calculated from the recommended cfj values. TABLE 1. The bulk moduli and the pressure derivative. B s ,GPa

B T ,GPa

B 7 = ( —)

UPJ 3QQK

Recommended [1]

75.3 ±0.5

75.Q ±0.5

4.5 ±0.5

77 K

Experimental [2]

7^8

7^5

4.5 ±0.5

0K

Experimental [2]

TLO

76/7

0, 77 K

B

I Dilatation

|

7^3

|

76X)

|

VARIATIONS OF THE BULK MODULUS

Like the elastic stiffness constant, the bulk modulus is expected to have different values for adiabatic and isothermal conditions, and to vary with temperature, pressure and doping. The discussion of these variations is given in the next Datareview and need not be repeated here. All significant values are given in TABLES 1 and 2, and only two comments are necessary. TABLE 2. The rate of change of bulk moduli with temperature above 300 K dB s /3T

5B1VdT

Experimental [2]

-0.007GPaK 1

-0.007GPaK"1

Dilatation

-0.006 GPa K"1

-0.006 GPa K"1

It is noteworthy that B' = 4.5 has been used to convert pressure to lattice constant in highpressure experimentation using an equation of state [3] and doing this results in a linear dependence of bandgap on lattice constant. This suggests this is the right value of B'.

Unfortunately, no authors have reported the sensitivity of this procedure to the exact value of B ' used, so we cannot use it to refine the accuracy of B'. Secondly, one paper distinguishes B / S and B /T , giving a higher experimental value (4.67) for the latter than for the former (4.49) [4]. This is physically improbable as it would result in B T being greater than B s at pressures above 3.8 GPa (corresponding to a negative coefficient of thermal expansion or a negative heat capacity). Consequently we give the same value of 4.5 to both. The value of dB/dT can be used to determine values of B s or B T at temperatures above 300 K, T, by BS'T (T) = BS'T (300 K) + (T - 300 K) (3B/dT) The different experimental and theoretical temperature dependences of C11 and C12 given in [1] result in slightly different predictions for the temperature dependence of B. As in [1] for cij3 we recommend the use of the dilatation values for temperatures above and below 300K. REFERENCES [1] [2] [3]

DJ. Dunstan [ Datareview in this book: 1.4 Stiffness of GaAs ] R.I Cottam, G.A. Saunders [J. Phys. C(UKJ, vol.6, no.13 (1973) p.2105-18 ] V.A. Wilkinson, A.R. Adams [ Datareview in this book: 4.7 Electron effective mass and hole

[4]

effective mass in GaAs, pressure dependence] HJ. McSkimin, A. Jayaraman, P. Andreatch [J. Appl. Phys. (USA) vol.38 (1967) p.2362 ]

1.4

Stiffness of GaAs DJ. Dunstan September 1995

A

EVTRODUCTION

The cubic symmetry of GaAs requires three independent constants to define the elasticity. Most convenient are the elastic stiffness constants, the elements C11, C12 and C44 of the stiffness matrix (reduced from the stiflhess tensor c p by index substitutions). They are most accurately measured by measuring the speed of an ultrasonic wave or by finding the resonant frequency of a crystal excited electrostatically. Then if the density of GaAs is known, the elastic constants can in principle be derived very accurately. B

ADIABATIC ELASTIC STIFFNESS CONSTANTS

The elastic constants c^ have been measured by a large number of authors, mostly by acoustic wave velocities, with varying accuracies. The accuracy claimed by Beilin [1] is, however, typical: 0.06% absolute and 0.03% for relative values. Cottam and Saunders [2] paid great attention to errors and compared their results with previous authors; Blakemore [3] and Brice [4] also gave recommended values. Inspection of some of the reported values (TABLE 1) for values at 300 K implies either large sample-to-sample variations or that the errors are considerably more than reported. The mean values and standard deviations of the first six rows of data are given in TABLEl. TABLE 1. Values of the adiabatic elastic stiffiiess constants at 300 K. C11 GPa

C12 GPa

C44 GPa

Reference

117.60

52.70

59.60

[5]

118.10

53.20

59.40

[6]

118.10

53.72

59.44

[7]

118.85

53.75

59.45

[1]

119.04

53.84

59.52

[8]

119.80

53.40

60.40

[9]

119 ± 0 . 8

53.4 ± 0 . 4

59.6 ± 0 . 4

Mean

118.4 ± 0.4

53.7 ± 1 . 5

59.1 ± 0.2

[2] (recommended)

119(1)

53.8(1)

59.5(1)

[3]

118.8

53.8

59.4

[4]

The Cottam and Saunders values [2] are both close to the averages and have the most carefully considered errors, and are therefore the values recommended here for the adiabatic stiffiiess

constants. Note that the error on the C12 value is much larger than the others; this is because C11 and C44 are measured directly from single waves while C12 is derived from three. C

VARIATIONS OF THE ELASTIC CONSTANTS

The elastic constants are reported to vary in a systematic manner with temperature, with doping and with hydrostatic pressure; also the adiabatic and isothermal values are different. Cl

Isothermal Elastic Constants

The adiabatic moduli are larger than the isothermal values because as a crystal is compressed it tends to heat up; thermal expansion means it compresses less - is stiffer - than in the isothermal case. The difference between the isothermal and adiabatic elastic moduli is much too small to be determined experimentally. It may be deduced from the thermodynamic relationship between the isothermal and adiabatic compliance tensors s, the thermal expansion tensor a and the heat capacity C [10]. In [11] we find that ss - sT = 6.31 x 10"15 m2N"\ This value may be converted for the stiffness constants using the relationships between c and s [10,11], giving 5c u = 0.325 GPa, 5c12 = 0.325 GPa and 5c44 = 0.110 GPa. Making these corrections to the values of the adiabatic constants of Cottam and Saunders [2], we obtain the recommended values for the isothermal stiffness constants given in TABLE 2. TABLE 2. Recommended values for the isothermal stiflhess constants at 300 K.

C11GPa 118 A 0.4

C2

C12GPa I

53.5 ± 1 . 5

C44GPa I

59.0 ± 0.2

Second-order Stiffness Constants and Pressure Dependence

Second order elastic constants have been determined by Drabble and Brammer, applying uniaxial strain to the crystal in the acoustic wave experiments [8]. The values given (TABLE 3) are appropriate for the thermodynamic definition in which cijk is numerically equal to the full tensor element Cy^ [12]. The reliability of these data is hard to assess. The authors give errors of about ±10 GPa, but this appears to be assessed from the uncertainties in the acoustic part of the experiment only. In fact, large errors are likely in uniaxial strain experiments in which the crystal is squeezed between anvils, because the decomposition of the applied force into a hydrostatic and an axial strain depends on the friction or slippage between the sample and the anvils. The error induced by this effect may be as high as a factor of two. However, cijk, the pressure derivatives of C1J and the pressure derivative of the bulk modulus are all related, and values of dc12/3P derived from cijk (using expressions from Birch [13] and Brugger [12]) are given in TABLE 4 for comparison with the measured pressure derivatives. The large discrepancies suggest that the cijk of TABLE 3 may indeed be considerably in error.

TABLE 3. Second order elastic stiffiiess constants (see text for errors).

C 111 GPa

-675

C 112 GPa

I

-400

C123GPa

I

-5

C144GPa

I

-70

C166GPa

|

-320

C456GPa

|

-70

Reference

| [8]

The hydrostatic pressure dependence is shown in TABLE 4 (note that this is a dimensionless quantity). Hydrostatic pressure conditions are relatively accurate compared with uniaxial stress and so these values are taken as reliable. However, they came from very small shifts of acoustic wave velocity with pressure, (of about 1%), and should be taken as only accurate to about ±0.5. TABLE 4. Pressure dependence of elastic moduli.

^c1 t/aP

dc12/dP

dc44/dP

4.63

442

1.10

5.0 C3

I

1.7

1

0.5

Reference [7]

1 From TABLE 3

Dependence on Doping

Some authors reported different values for the stiffhess constants in heavily doped n-type and ptype GaAs [1,5]. We are sceptical whether the effect of doping has really been seen. Bruner and Keyes [14] saw a significant decrease in n-type Ge, of 6%, but only at a doping level of 5><1019 cm"3. Furthermore, as expected from theory [14], the effect was seen only on the shear modulus C44. In contrast, the differences reported in GaAs are limited to doping levels of a few 1018 cm"3, and are about 1% for p-type and 0 to 0.5% for n-type. This is within the scatter of results on different undoped samples (TABLE 1). Furthermore, they are reported to be comparable for C11, C12 and C44 whereas in both p-type and n-type GaAs they should be comparable for the two shear moduli, C44 and C11-C12, but should vanish for the bulk modulus B = (c n + 2c12) / 3. It may be noted that the lattice constant of GaAs with various impurities changes consistently with Vegard's law even for n-type and p-type dopants [15]. It is quite implausible that there should be an effect of the Fermi level on the shape of the interatomic potential which changes the position of the minimum (giving the lattice constant) at all pressures except at ambient (to give a change in B). Consequently we recommend that doping is not assumed to affect Cy.

C4

Dependence on Temperature TABLE 5. Temperature dependence of the elastic stiflhess constants (GPa). Ref [1] Ref[2] C11 (OK)-C 11 (300 K) C11 (77 K)-C 1 1 (300 K)

2.3

ac n /6T(T;>300K)

Ref[6]

Dilatation

2.6

4.5

1.0

2.4

4.0

1.0

-0.010

C12 (0 K) - C12 (300 K)

1.2

C12 (77 K) - C12 (300 K)

0.95

ac n /aT(T^300K)

C44 (77 K) - C44 (300 K)

1.15

ac44/6T(T^300K)

I -.005

-0.014

1.1 -0.005

C44 (0 K) - C44 (300 K)

Refs[5, 13]

-0.0063 3.9

1.0

3.4

1.0

-0.0058

-0.0060

1.3

0.6

0.25

1.2

0.5

0.25

-0.006 |

-0.0070

-0.0015

We come finally to the temperature dependence of Cy. Cottam and Saunders [2] made measurements from 300 K downwards, as did Garland and Parks [6]. Beilin et al [1], Burenkov et al [5] and Jordan [16] considered also data above 300 K. All authors find a linear temperature dependence around 300 K and upwards (TABLE 5). There are, however, some discrepancies between the various results. Theoretically, a major contribution to dc/dT comes from the dilatation due to thermal expansion. Its effect may be estimated using the values given in TABLE 4 for the rate of change of C8 with pressure, and such estimates are also given in TABLE 5. It is remarkable that Cottam and Saunders find exactly the same percentage change (2%) for all three moduli between 300 K and 77 K, since the dilatation terms do not behave in this way. Perhaps there are unrecognised experimental difficulties. Nor, theoretically, is it clear what the contribution of other terms might be. Consequently, we recommend that TABLE 1 and the last column of TABLE 6 should be used to obtain stiffness constants above and below 300 K; however, the reader who prefers to use experimental results is recommended to use those of refs [1] and [2], which are in good agreement with each other. REFERENCES [1] [2] [3] [4] [5] [6] [7] [8] [9]

V.M. Beilin, Yu.Kh. Vekilov, O.M. Krasil'nikov [ Sov. Phys.-Solid State (USA) vol.10 (1969) p.2443] RJ. Cottam, G.A. Saunders [ J. Phys. C (UK) vol.6 (1973) p.2105 ] J.S. Blakemore [ J. Appl. Phys. (USA) vol.53 (1982) p.R123 ] J.C. Brice [ ^Properties of Gallium Arsenide Second Edition (INSPEC, IEE, London 1990) ch. 1.4 p.15] Yu. Burenkov, Yu.M. Burdokov, S. Yu. Davidov, S.P. Nikanorov [ Sov. Phys.-Solid State (USA) vol.15 (1973) p. 1175] CW. Garland, K.C. Parks [ J. Appl. Phys. (USA) vol.33 (1959) p.759 ] HJ. McSkimin, A. Jayaraman, P. Andreatch [ J. Appl. Phys. (USA) vol.38 (1967) p.2362 ] J.R. Drabble, AJ. Brammer [ Solid State Commun. (USA) vol.4 (1966) p.467 ] B.A. Bobylev, A.F. Kravchenko [ Sov. Phys.-Acoust. (USA) vol.13 (1973) p.242 ]

[10] [11] [12] [13] [14] [15] [16]

J.F. Nye [ Physical Properties of Crystals (Clarendon Press, Oxford, 1957) ] DJ. Dunstan [ Datareview in this book: 1.5 Compliance of GaAs ] K. Brugger [ Phys. Rev. A (USA) vol. 133 (1964) p. 1611 ] F. Birch [ Phys. Rev. (USA) vol. 71 (1947) p.809 ] LJ. Bruner, R.W. Keyes [ Phys. Rev. Lett. (USA) vol.7 (1961) p.55 ] J.C. Brice [ Datareview in this book: 1.2 Lattice parameter of GaAs ] A.S. Jordan [ J. Cryst. Growth (Netherlands) vol.49 (1980) p.631 ]

1.5

Compliance of GaAs DJ. Dunstan September 1995

A

INTRODUCTION

The compliance constants, sii9 are an alternative way of representing the same information as the stiffness constants, Cy. There are no independent determinations of them. Accordingly, we take the recommended values of the measured adiabatic Cy from the previous Datareview [1] and convert them. It is simple to calculate the difference between the adiabatic and isothermal values of S1J. The result is then used to obtain the isothermal c^ given in [I].

B

ADIABATIC COMPLIANCE CONSTANTS

These are obtained from the recommended values of adiabatic stiffness constants using the relationships [2] C 11 + C 12

s

11

(C11 -C 1 2 )(C 1 1 + 2 c 1 2 )

S1, =

(1) c

c

c

( n - i2)( n

844 =

+2c

i2)

7" C

44

Taking due account of the errors, the results are presented in TABLE 1. Note that S12 is negative. TABLE 1. Adiabatic compliance constants at 300 K.

C

S11In2N"1

S 12 In 2 N' 1

S 44 In 2 N' 1

11.8 ±0.2 x 10-12

-3.7 ±0.1 x 10-12

16.92 ±0.05 x 10"12

ISOTHERMAL COMPLIANCE CONSTANTS

The difference between the isothermal and adiabatic elastic constants is much too small to be determined experimentally, particularly as the most accurate techniques using acoustic waves determine only the adiabatic values. The difference may be deduced from the thermodynamic relationship [2] between the isothermal and adiabatic compliance tensors s, the thermal expansion tensor a and the heat capacity C:

HP

*i

s

" i

= a

a

ij ki—

(2)

where the superscripts S, T and a refer to constant entropy, constant temperature and constant stress respectively. Since the thermal expansion of a cubic crystal is isotropic, the difference is the same for each component OfSp1. Using the values for 300 K, where a = 6.03 x 10"6K"1, the specific heat is 0.325 J ^1K"1 and the density is 5.3165 g cm"3, giving C° = 1.73 x 106 J m"3. Then s s - sT = 6.31 x 10"15 Hi2N"1. This quantity may be subtracted from the values in TABLE 1 to obtain the isothermal values, but it is much smaller than the uncertainties in S1J. Here it is important to note that S44= 4s2323, so that s44s - s 44 T= 4 (ss-sT). D

VARIATIONS IN THE COMPLIANCE CONSTANTS

The compliance constants for temperatures other than 300 K, and for pressures above ambient, may be obtained if required from the corresponding c^ values from [1] using EQN (1). REFERENCES [1] [2]

DJ. Dunstan [ Datareview in this book: 1.4 Stiffness of GaAs ] J.F. Nye [ Physical Properties of Crystals (Clarendon Press, Oxford, 1957) ]

1.6

Thermal expansion coefficient of GaAs S. Adachi July 1995

A

INTRODUCTION

The thermal expansion coefficient (TEC) is a second-rank symmetric tensor relating temperature T (scalar quantity) and strain [e] (second-rank tensor) by [e] = [a]T

(1)

The thermal expansion tensor [a] in crystals with the zinc-blende structure has only one tensor component, ath = a^ = O57 = Ct22. The thermal expansion coefficient a± is known to be proportional to the specific heat cv (Griineisen's rule) [I]:

^ = J-f *»! = I E £ *

O0UTJP

3V

(2) (>

where S0 is the lattice constant, y is the averaged Gruneisen parameter, C0 is the compressibility and V is the volume of the crystal. The TEC depends markedly on the temperature and is positive for most crystals. Knowledge of the TEC is often essential in variable temperature processes and experiments. B

UNDOPED GaAs

The TEC is usually determined by measuring the temperature dependence of the lattice constant. Experimental evaluation of a± has been carried out by many authors [2]. The data are taken from Sparks and Swenson [3] for the range O-40 K (open circles), from Novikova [4] for the range 28-337 K (solid circles), from Feder and Light [5] for the range 307-609 K (open triangles) and from Brice [2] from 350 K into the molten range above 1500 K (solid triangles). TABLE 1 also lists the a± values for GaAs reported by Brice [2]. It is understood from FIGURE 1 that with decreasing temperature a± decreases from its positive value and then passes through zero (T-50 K) undergoing a reversal in sign.

TABLE 1. Thermal expansion coefficient a A for GaAs. Note that the expansion coefficient has a maximum at 12 K and a minimum at 31 K. T(K)

C^(IO-6K-1)

T(K)

CCd1(IO-6K-1)

J

0.001

J00

6J03

JO

0.018

J50

6J9

J2

0.022

J00

628

J5

0

J50

^36

JO

-0.06

J00

6A4

J5

-014

J00

6.59

JO

-017

J00

674

Jl

-018

J00

6J89

J5

-016

JOO

7X)4

JO

IOH

1000

7.19

JO

I O02

1100

7.34

100

2X)7

1200

749

150

^90

1300

7J54

200

AM

1400

7.80

250

5,45

1500

7.95

273

1 5.89

The experimental data reported by Sparks and Swenson [3] and Novikova [4] showed negative Ct1I1 values at low temperatures. Smith and White [6] measured a^ in the temperature range 4-30 K, and also observed negative a^ above 12 K. Above 10 K the Smith-White data agreed reasonably with those of Sparks and Swenson, but below this temperature there was a considerable difference. It is known that not only zinc-blende materials but also diamond-type semiconductors show such 'unusual' negative thermal expansion for temperatures below T ~ 0.10D(T=O K), where 6 D (T=0 K) is the limiting value of the Debye characteristic temperature as T-O K. Biernacki and Scheffler [7] carried out density-functional-theory calculations of thermodynamic potentials to study the temperature dependence of a^. Their results showed excellent agreement with published experimental data of Si. They concluded that the origin of the negative expansion effect is traced back to the entropy contribution of the Gibbs free energy.

GaAs

TEMPERATURE

( K)

FIGURE 1. Experimental a A (TEC) versus temperature for GaAs. Data are taken from Sparks and Swenson [3] for the range 0-40 K (open circles), from Novikova [4] for the range 28~337 K (solid circles), from Feder and Light [5] for the range 307-609 K (open triangles) and from Brice [2] from 350 K into the molten range above 1500 K (solid triangles).

C

EFFECT OF STOICHIOMETRY

Brice [2] discussed the effects of stoichiometry on <%. The data in FIGURE 1 (TABLE 1) probably represent values for stoichiometric crystals. Data for gallium-rich samples lie below the curve and data for arsenic-rich samples lie above the curve. The effects of deviations from stoichiometry can be roughly quantified by saying that increasing the arsenic content of the melt by 1 atom % increases a± by the order of 3% [2]. D

EFFECT OF DOPING

Temperature dependence of the lattice constant in doped GaAs was studied by Bak-Misiuk et al [8-10] in the range 300 to 800 K. They found that at high temperatures the TEC depends on dopants and their concentrations. For most of the dopants studied [8] the TEC was increased. Typical increases for 100 ppm (4.4 x 1018 cm"3) concentrations were 1.1, 0.04, 0.23 and 0.18 x 10"6 K"1 for Si, Zn, Sn and Te, respectively, at 650 K. On the other hand, a decrease in the TEC was observed in S-doped GaAs [10]. The TEC in this case can be given by a± = (6.683 -6.8OxIO-20N8)XiO-6

K1

(3)

where Ns is the sulphur content in cm"3. The decrease in a± was explained by the strong covalency of the Ga-S bond in comparison with that of GaAs [10].

E

EPITAXIAL FILMS

The OC^ in a strained material system, (GaAs films grown on (10O)Si substrates), was studied by Lucas et al [11] in the temperature range between 20 and 450 0 C by means of X-ray scattering. It was found that the measured a± of GaAs in the direction parallel to the film plane follows the thermal expansion of the Si substrate and is therefore smaller than in bulk GaAs. The a^ values for GaAs films in the directions parallel and perpendicular to the film plane were 3.46 * 10"6 and 8.91 x 10-6K-1, respectively. The thermal expansion of low-temperature (190"2200C) MBE-grown GaAs (LT GaAs) was measured by Leszczynski and Walker [12] using X-ray diffraction methods. It is noted that LT GaAs usually possesses high concentrations of nonstoichiometric arsenic present as arsenic antisites and interstitials (of concentrations even higher than 1020 cm"3). They determined that the as-grown LT GaAs layers have smaller values of TEC than semi-insulating GaAs substrate (by only about 0.05 x 10"6 K"1). They also found that both values become almost identical for an annealed sample with a diminished amount of arsenic excess. From this fact, they concluded that the difference in the TEC originates from the presence of point defects, such as arsenic interstitials, arsenic antisites and possibly gallium vacancies. However, the compensating mechanisms of a variety of defects in LT GaAs cannot be excluded at present. The anisotropy of the TEC in GaAs-based multilayer structures has also been discussed in [13,14]. F

CONCLUSION

The TEC of GaAs depends markedly on the temperature and shows negative values below about 50 K. Although the majority of data have been obtained on bulk samples, some studies of epitaxial materials have been performed and variations from bulk values seen as expected. More detailed studies are needed in this field to understand thoroughly the effects of stoichiometry, doping (free carriers) and point defects on TEC. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] II1] [12] [13] [14]

S.I. Novikova [ Semicond. Semimet. (USA) vol.2 (1966) p.33-48 ] J.C. Brice [ Properties of Gallium Arsenide, 2nd Edition, EMIS Datareviews Series No.2 (INSPEC, IEE, 1990) ch. I p . 18-9] P.W. Sparks, CA. Swenson [ Phys. Rev. (USA) vol. 163 (1967) p.779-90 ] S.I. Novikova [ Sov. Phys .-Solid State (USA) vol.3 (1961) p.129-30 ] R. Feder, T. Light [ J. Appl Phys. (USA) vol.39 (1968) p.4870-1 ] T.F. Smith, G.K. White [J. Phys. C, Solid State Phys. (UK) vol.8 (1975) p.2031-42 ] S. Biernacki, M. Scheffler [ Phys. Rev. Lett. (USA) vol.63 (1989) p.290-3 ] J. Bak-Misiuk, H.G. Bruhl, W. Paszkowicz, U. Pietsch [ Phys. Status Solidi A (Germany) vol. 106 (1988)p.451-7] J. Bak-Misiuk, G. Kiihnel, W. Siegel, U. Pietsch [ Phys. Status Solidi B (Germany) vol. 158 (1990) p.Klll-14] J. Bak-Misiuk, J. Kohler, U. Pietsch [ Phys. Status Solidi A (Germany) vol. 126 (1991) p.K43-7 ] N. Lucas, H. Zabel, H. Morko?, H. UnIu [Appl. Phys. Lett. (USA) vol.52 (1988) p.2117-9 ] M. Leszczynski, J.F. Walker [Appl. Phys. Lett. (USA) vol.62 (1993) p.1484-6 ] G. Clec'h, G. Calvarin, P. Auvray, M. Baudet [J. Appl. Crystallogr. (Denmark) vol.22 (1989) p.372-5] J. Bak-Misiuk, J. Wolf, U. Pietsch [ Phys. Status Solidi A (Germany) vol. 118 (1990) p.209-17 ]

1.7

Specific heat and Debye temperature of GaAs S. Adachi July 1995

A

INTRODUCTION

Investigation of the thermal properties of solids is an old topic which arises in strong connection with the fundamental properties of the solids. We discuss here the specific heat (heat capacity) and the Debye temperature of GaAs. B

SPECIFIC HEAT

The specific heat at constant pressure, cp, can be given by

where AQ is the heat input and AT is the corresponding change in temperature. In order to obtain the heat capacity at constant volume c^ which is the quantity usually resulting from theoretical calculations, one can use the following expression:

cp-cv~-^—

(2)

where ath is the coefficient of linear expansion, V is the volume of the crystal and C0 is the isothermal compressibility. There have been several experimental data on the specific heat of crystalline GaAs over a wide range of temperature (1
Temperature (K) FIGURE 1. Experimental specific heat cp as a function of temperature for GaAs determined from calorimetric studies. The experimental data are taken by Cetas et al [3] for the range 1 ~34 K (open circles), by Piesbergen [1] for the range 12-273 K (solid circles) and by Lichter and Sommelet [4] from 300 K into the molten range above 1513 K (open triangles). The inset shows the calculated results of EQN (3) (dashed line) and EQN (4) and (5) (solid line), together with the experimental data of Lichter and Sommelet [4] (open triangles).

It was reported [8] that the heat capacity cp of GaAs at temperatures between 300 and 1500 K can be expressed in the power form: cp(T) = 0.343 + 4 . 7 x i o 5 T-3.2xlO 3 T" 2

J/g K

(3)

The inset of FIGURE 1 represents the calculated result of this equation (dashed line). More recently, Yamaguchi et al [6] measured high-temperature heat capacities of some III-V compounds over the temperature range 650 to 1550 K using a drop-calorimeter. They determined the melting point of GaAs to be 1514±1 K. They also found that the heat capacity could be expressed as cp(T) = 0.327 + 5.52xl0' 5 T - 1.66xlO3 T ~2 J/g K

(4)

for solid GaAs (800
(5)

for liquid GaAs (1514
C

cp(mJ/gK)

T(K)

J

0.043

JO

J

0.156

J

cp(mJ/gK)

T(K)

cp(mJ/gK)

IU

J60

320

70

138

_273

323

0.422

J80

160

J80

325

JO

L00

_90

181

J00

327

J2

2A2

J4

3j)6

JlO

216

J00

343

J6

^62

J20

230

J00

351

J8

10.07

243

J00

359

JO

142

J40

254

_800

367

_22

18J*

J50

264

J)OO

375

J4

23/7

J60

272

1000

383

J6

28j)

J70

279

1100

391

_28

3MkI

J80

285

1200

400

JO

394

J90

292

1300

408

J5

524

J00

298

1400

416

JO

654

JlO

304

1500

424

J5

78J

J20

308

JO

904

240

315

100

130

199

400

335

DEBYE TEMPERATURE

The Debye temperature 0D is a useful parameter in solid-state problems because of its inherent relationship to lattice vibration. The parameter 0 D can be used in characterizing the excitation of phonons and to describe various thermal properties associated with lattice phenomena [2], The Debye model for lattice vibrational energy yields the relation [7]: S ~ cv = C01F(OnZT)

(6)

where F(O1/!) is the Debye function. In EQN (6), Ccl is the classical specific heat (=3 kN, where k is Boltzmann's constant and N is the number of atoms per unit volume of solid). For GaAs, Ccl = 0.345 J/g K. We show in FIGURE 2 the temperature dependence of the effective calorimetric Debye variable 0D which results from inverting EQN (6) and using the values of Cetas et al [3] (open circles) and Piesbergen [1] (solid circles). TABLE 2 also lists the 0D values in the temperature range 0-1514 K taken from refs [1-3,9].

TABLE 2. Debye temperature 0D of GaAs. T(K)

On(K)

J)

345

_20

247

JO

302

100

351

150

360

200

358

250

352

300

363

350

364

1514

1 374

Notes

a

b

a

I

a= 0D has minima at 20 and 250 K. b= 0D has a maximum at ~ 150 K.

GaAs

Temperature (K) FIGURE 2. Effective calorimetric Debye variable 0D for GaAs obtained from inverting EQN (6) and using the cp values of Cetas et al [3] (open circles) and Piesbergen [1] (solid circles).

It is found from FIGURE 2 that 8 D has a nearly constant value from high T to T-IOO K, decreases gradually, showing a minimum at T-20 K, and then increases with further decrease of T. Cetas's data [3] suggest that 0D(T = 0 K) = 347 K (where 6 D (T = 0 K) is the limiting value of the Debye characteristic temperature as T - 0 K). Holste [10] also obtained the value of 0 D (T = 0 K) = 344.6 ± 2.0 K, from a detailed analysis of the data of Cetas et al. This value agrees admirably with expectations from the elastic moduli (344 K; [2]). The X-ray Debye temperature for GaAs has also been reported in [11] (255 ± 5 K (Bragg reflection) and 257 ± 3 K (X-ray transmission)). D

CONCLUSION

We have reviewed the specific heat and Debye temperature of GaAs. These parameters are found to depend strongly on temperature. The results presented here can be used in a general discussion of the thermal properties of this material. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] [II]

U. Piesbergen [ Z. Natforsch. (Germany) vol. 18 (1963) p. 141-7 ] U. Piesbergen [ Semicond. Semimet. (USA) vol.2 (1966) p.49-60 ] T.C. Cetas, CR. Tilford, CA. Swenson [ Phys. Rev. (USA) vol. 174 (1968) p.835-44 ] B.D. Lichter, P. Sommelet [ Trans. Metall. Soc. AME (USA) vol.245 (1969) p. 1021 -7 ] AJ. Dash,A. Finch,P.J. Gardner,M. Cottrell [J Chem. Eng. Data (USA) vol.19 (1974) p.113-14] K. Yamaguchi, K. Itagaki, A. Yazawa [ J Jpn. Inst. Met. (Japan) vol.53 (1989) p.764-70 ] J.S. Blakemore [ J. Appl. Phys. (USA) vol.53 (1982) p.R123-81 ] R. Hillel, J. Bouix [ J. Cryst. Growth (Netherlands) vol.38 (1977) p.67-72 ] J.C. Brice [ in Properties of Gallium Arsenide, 2nd Edition, EMIS Datareviews Series no.2 (INSPEC, IEE, 1990), ch.l p.20 ] J.C Holste [ Phys. Rev. B (USA) vol.6 (1972) p.2495-7 ] RN. Kyutt [ Sov. Phys.-SolidState (USA) vol.20 (1978) p.227-9 ]

1.8

Thermal conductivity of GaAs S. Adachi July 1995

A

INTRODUCTION

Thermal conductivity (K), or thermal resistivity (K"1), results essentially from interactions between phonons and from the scattering of phonons by crystalline imperfections. Knowledge of the thermal conductivity of semiconductors forms an important part in the design of power-dissipating devices, such as diodes, transistors and optoelectronic devices (laser and light-emitting diodes). Numerical thermal-conductivity data are also necessary in calculating the figure of merit for thermoelectric devices (e.g., Peltier devices). A useful description of the theoretical and practical aspects of the thermal conductivity in semiconductors is given by Holland [1], Bhandari and Rowe [2] and Srivastava [3] who review some of the work done on III-V compounds, including GaAs and its related alloys. B

UNDOPED GaAs

An exact calculation of lattice thermal conductivity for the III-V compounds is possible in principle, but lack of knowledge of various parameters (e.g., anharmonic forces and lattice vibration spectra) and the difficulty of obtaining exact solutions of phonon-phonon interactions are formidable barriers to progress. Experimental evaluation of K for undoped (or low-doped) GaAs has been carried out by many authors (see [4]). FIGURE 1 plots the experimental K versus temperature T for GaAs. The data are taken below 300 K from Holland (n ~7 x 1015 cm"3; open circles [5]) and above 300 K from Amith et al (n ~ 5 x 1016 cm'3; solid circles [6]). In most semiconductors, we observe experimentally that the thermal conductivity of a pure single crystal is zero at T = 0 K and rises approximately exponentially to a maximum near 10 K, falls somewhat faster than T"1, and then varies approximately as T"1 to the melting temperature. The calculated curves in FIGURE 1 are obtained from the power law: K(T) = AT n

(1)

The dashed and solid lines are calculated with n = 2.25 and A = 0.41 W/cm K325 and n = -1.30 and A = 745 W/cm K03, respectively. It is seen that EQN (1) successfully explains the experimental K in the range 15 - 700 K. We also know that many III-V compounds obey the power law of EQN (1) with -1.55< n ^ -1.20 for temperatures T > 50 K [7]. These are stronger temperature dependences than is predicted for the three-phonon process, and probably indicate the presence of higher-order processes [I]. The anomalous temperature dependence seen in FIGURE 1 at T ^ 500 K (solid circles) is due to the effect of heat transport by photons [8]. This effect becomes detectable for pure samples when the background of the lattice thermal conductivity is small and when free-carrier absorption is absent [6].

K (W/ cm K)

GaAs

TEMPERATURE ( K ) FIGURE 1. Temperature dependence of the thermal conductivity K for GaAs. The experimental data are taken below 300 K from Holland (n - 7 x 1015 cm"3; open circles [5]) and above 300 K from Amith et al (n - 5 x 1016 cm 3 ; solid circles [6]). The dashed and solid lines are calculated with n = 2.25 and A = 0.41 W/cm K3 25 and n = -1.30 and A = 745 W/cm K0 3, respectively.

The data tabulated in TABLE 1 were obtained by drawing a smooth curve through a plot of the measured curves [5,6,9] as a function of temperature. Note that these data show a maximum at T-10 K (see also FIGURE 1). The thermal conductivity of the solid at the melting point (-1500 K) is about 0.07 W/cm K [9]. C

EFFECT OF DOPING

In doped semiconductors, the total thermal conductivity can be generally given by the sum of the lattice (KL) and electronic (Ke) contributions. In a metal, the electronic thermal conductivity Ke and electrical conductivity a are related by the Wiedemann-Frantz-Lorentz law: K6 = L o T

(2)

where L is the Lorentz number. In a semiconductor, a more complicated relationship exists between Ke and o [I]. It was found [1] that the n-type impurities (Te) do not cause as large a decrease in low-

temperature (T < 100 K) thermal conductivity as comparable amounts of the p-type impurities (Zn3 Cd and Mn). At high temperatures (T > 300 K) an increase in the free electron concentration caused a decrease in thermal conductivity. This was attributed to scattering of phonons by electrons [I]. TABLE 1. Thermal conductivity K and difiusivity D for GaAs. T (K)

D

K (W/cm K)

D (cmVs)

_2

^O

71000

4

9.4

41000

J

18

22000

J

25

11000

JO

28

5200

J5

23

830

_20

16

210

_40

62

18

_60

3/7

6A

_80

^5

2J)

100

L9

L8

150

U

078

200

076

048

300

045

026

400

O30

017

600

O19

O10

800

O14

0.072

1000

0.11

O055

1500

I 0.07

I 0.032

EFFECT OF ALLOYING

It is well known that dilute alloying of GaAs with InAs is a very effective technique for reducing the dislocation density of semi-insulating crystals grown by the liquid-encapsulated Czochralski (LEC) method. Ohmer et al [10] reported the effect of this small addition of InAs on the thermal properties of GaAs. All of the dilute Ga^xInxAs alloys (0.00 < x < 0.013) used in this experiment were grown by the high-pressure LEC method from high-purity GaAs melts contained in pyrolytic boron nitride crucibles. Alloying was achieved by adding In to the melt with an excess of As to maintain a stoichiometric balance between group III and V components. They found that for x = 0.005 the thermal conductivity and difiusivity are reduced to 50% and 60% of the values for GaAs5 respectively. It is also important to point out that when large numbers of foreign atoms are added to the host

lattice, as in alloying, the thermal conductivity may decrease significantly. For example, Abraham et al [11] reported a factor of 8 reduction in thermal conductivity OfInxGa1^As for x~0.5, as compared to GaAs.

E

THERMAL DIFFUSIVITY

The thermal difiusivity can be evaluated from the thermal conductivity by means of the definition n D

K

(3)

where c F p and D are the specific heat at constant pressure, crystal density and thermal difiusivity of GaAs, respectively. Introducing the K values (TABLE 1) and cp values in [12] into EQN (3), we obtain the thermal diffusivity of typical samples. The calculated results are shown in TABLE 1. For 100 < T < 300 K, these values agree to within ± 10% with the measurements of Ohmer et al [10] and Sasaki et al [13] who, however, found a value of 3.4 - 4.7 cm2/s at 80 K. The diffusivity value at the melting point (-1500 K) is estimated to be about 0.03 cm2/s. F

CONCLUSION

It is evident from TABLE 1 that the thermal conductivity and diffusivity of GaAs show very strong temperature dependence. The thermal conductivity shows a maximum at T-10 K, while the thermal diffusivity greatly decreases with increasing temperature. Further studies are needed in this field to understand thoroughly the effects of free carriers on the heat transport properties. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] II1] [12] [13]

M.G. Holland [ Semicond. Semimet. (USA) vol.2 (1966) p.3-31 ] CM. Bhandari, D.M. Rowe [ Thermal Conduction in Semiconductors (Wiley, New York, 1988) ] GP. Srivastava [ The Physics ofPhonons (Adam Hilger, Bristol, 1990) ] J.C. Brice [ Properties of Gallium Arsenide, 2nd Edition, EMIS Datareviews Series no.2 (INSPEC, IEE, 1990) ch.lp.21-3] M.G. Holland [ Phys. Rev. (USA) vol. 134 (1964) p.A471-80 ] A. Amith, I. Kudman, E.F. Steigmeier [ Phys. Rev. (USA) vol. 138 (1965) p.A1270-6 ] S. Adachi [ Physical Properties of IH-V Semiconductor Compounds: InP, InAs, GaAs, GaP, InGaAs, and InGaAsP (Wiley-Interscience, New York, 1992) ] G.D. Cody, B. Abeles, D.S. Beers [ Bull. Am. Phys. Soc. (USA) vol.8 (1963) p.296] A.S. Jordan [ J. Cryst. Growth (Netherlands) vol.49 (1980) p.631-42 ] M.C. Ohmer, W.C. Mitchel, G.A. Graves, D.E. Holmes, H. Kuwamoto, P.W. Yu [ J. Appl Phys. (USA) vol.64 (1988) p.2775-7 ] M.S. Abraham, R. Braunstein, F.D. Rosi [ J. Phys. Chem. Solids (UK) vol. 10 (1959) p.204-10 ] S. Adachi [ Datareview in this book: 1.7 Specific heat and Debye* temperature of GaAs ] M. Sasaki, S. Horisaka, M. Inoue [ Jpn. J. Appl. Phys. (Japan) vol.26 (1987) p. 1704-8 ]

1.9

Melting point of GaAs S. Adachi July 1995

A

INTRODUCTION

The melting point is one of the most important thermophysical parameters. We will examine here the melting point of stoichiometric GaAs. B

EXPERIMENTAL DATA

Data about the normal melting temperature Tm for stoichiometric GaAs are given in [1-7]. TABLE 1 summarizes these results. TABLE 1. Melting point T7n for stoichiometric GaAs. T1n(K) 1511

[1] (1955)

1510 ± 3

[2] (1957)

1510

[3] (1963)

1511

[4] (1963)

1513 zb 3

[5] (1969)

1513

[6] (1972)

1514 ± 1

C

Ref(Year)

1 [7] (1988)

PRESSURE EFFECT

The hydrostatic-pressure derivative of the melting point for GaAs reported by Jayaraman et al [8] was given by (in K/kbar) dT — E = -3.4 (for 0 < P < 45 kbar)

D

(1)

CONCLUSION

The available experimental data on the melting point Tm of stoichiometric GaAs have been presented. The recent data suggest that T m = 1513 ± 2 K is an acceptable value of the normal melting point of stoichiometric GaAs. Data are available for the variation with hydrostatic pressure of Tm.

REFERENCES [1] [2] [3] [4] [5] [6] [7] [8]

W. Koster, B. Thomas [ Z Met. led. (Germany) vol.46 (1955) p.291-3 ] J. van den Boomgaard, K. Schol [ Philips Res. Rep. (Netherlands) vol. 12 (1957) p. 127-40 ] R.N. Hall [ J. Electrochem. Soc. (USA) vol. 110(1963) p.385-8 ] D. Richman [ J. Phys. Chem. Solids (UK) vol.24 (1963) p. 1131-9 ] B.D. Lichter, P. Sommelet [ Trans. Metall. Soc. AME (USA) vol.245 (1969) p. 1021-7 ] G.B. Stringfellow [ J. Phys. Chem. Solids (UK) vol.33 (1972) p.665-7 ] K. Yamaguchi, K. Itagaki, A. Yazawa [ J. Jpn. Inst. Met. (Japan) vol.53 (1989) p.764-70 ] A. Jayaraman, W. Klement Jr., G.C. Kennedy [ Phys. Rev. (USA) vol. 130 (1963) p.540-7 ]

CHAPTER 2 ELECTRON MOBILITY, DIFFUSION AND LIFETIME 2.1 2.2 2.3 2.4 2.5 2.6 2.7 2.8 2.9 2.10 2.11 2.12 2.13 2.14

Electron mobility in GaAs: overview Electron mobility in bulk GaAs Electron mobility in LPE GaAs Electron mobility in VPE and MOVPE GaAs Electron mobility in MBE GaAs Electron mobility in ion implanted GaAs Electron mobility in GaAs, pressure dependence Electron mobility enhancement in GaAs heterostructures Ballistic transport and velocity overshoot in GaAs Modulation doping Carrier concentrations following 8-doping Theoretical electron mobility/temperature dependences on carrier concentration and compensation ratio Minority electron mobility in doped GaAs Electron lifetimes in p-type GaAs

2.1

Electron mobility in GaAs: overview D. Lancefield July 1996

A

INTRODUCTION

General properties of the electron mobility in GaAs as a function of temperature and doping density can be considered independently of the growth technique used. The growth techniques considered in the following Datareviews give information about the electron mobility with respect to developments in the growth procedures, incorporation of background impurities and doping characteristics. The electron mobility, \i, is the velocity per unit electric field and at low fields is described by et

where e is the electronic charge, x is the average scattering time, and m* is the electron effective mass. The background physics and experimental techniques used to measure the electron mobility have been reviewed extensively in the last few years in a number of books[l-3].

FIGURE 1. Electron mobility vs. temperature. Solid curves are least-squares fits to data [I].

In GaAs, the principal scattering mechanisms controlling the scattering time, X9 are polar-optic and acoustic phonon scattering at temperatures above about 100 K. Below these temperatures ionized and neutral impurity scattering dominate. An example of the temperature dependence of the mobility in high purity material is shown in FIGURE 1 [I]. This data can be analysed using the approach of Wolfe et al [4], and also using the numerical techniques of Rode [5]. The impurity concentration, compensation ratio and other material parameters may be found from this analysis. For increasing impurity densities the temperature range where impurity scattering dominates is

extended towards room temperature, resulting in a lowering of the mobility with increased doping density. Measurements of electron mobility at 300 K and, in particular, at 77 K are used as a standard means of GaAs assessment. Look-up tables have been given to determine the compensation ratio, 0 = Na/Nd, from measurements of the carrier density and mobility by Walukiewicz et al [6] and by Benzaquen et al [7], FIGURE 2 shows the variation of electron mobility with carrier density for a large number of samples [8].

Mobility (cmVVs) Carrier concentration (cm 3 )

FIGURE 2. Mobility vs. electron carrier concentration in GaAs at 300 K. Solid and dashed curves give Hall and drift mobilities for the compensation ratios indicated [8].

B

CAUGHEY-THOMAS PARAMETERS

Using the data of Walukiewicz et al [6], Maziar and Lundstrom [9] give the following Caughey-Thomas parameters for calculating the mobility at 300 K r

/

\

r-rmn

I+Ua where Ji1113x = 8200 cm2/ Vs and the other parameters are calculated using the following expressions:

I!,* = 9.63 x io 1 8 0 3 - 4.2 x 1018 0 2 + 3.7 x 1017 0 + 9.85 x 10 1 6 , for 0 < 0.2 nef =

lo [i7.o7(i-0)

OO4

]5

for

0

< O2

VLn^n = -16500 0 3 + 17450 0 2 - 8080 0 + 2750, for 0 < 0.2

^

= 235O[l-0] 1 4 5 ,for0>O.2.

The utility of such look-up tables is dependent on the accurate modelling of the various scattering mechanisms and the inclusion of Hall scattering factor corrections [10]. For example, at intermediate and high doping densities the usual models of impurity scattering result in mobilities that are larger than observed experimentally. This reduction in the mobility has been attributed to impurity band effects [11] and to the correlated distribution of impurities [12]. C

HIGHER ENERGY CONDUCTION BAND STATES

mobility

density

Electron Density (cm"2)

Electron Mobility (cm2/Vs)

At large electric fields or by optical pumping the electrons can be transferred to higher energy conduction band states [2]. The mobility usually decreases with increasing energy in the F band due to the non-parabolicity increasing the electron effective mass. At sufficiently high energies efficient electron transfer to the L satellite valleys may occur resulting in reduced mobilities of the

Temperature (K)

FIGURE 3. Mobility and 2D electron density vs. temperature for a GaAs/ AlGaAs structure with a spacer layer of 75 nm.

order of 500 cm2/ Vs due to both a larger effective mass in the L-valley plus additional intervalley phonon scattering processes [13]. Further details are reviewed in the book by Adachi [14]. D

TWO DIMENSIONAL ELECTRON GAS EFFECTS

In addition high doping levels within a GaAs layer can be introduced by 5-doping. In this case there is negligible mobility enhancement as electron transport takes place in the GaAs close to the highly doped material. However a number of subbands are often occupied and the mobility in each subband has been found to be a function of the carrier density and the distribution of the dopant [18]. E

MINORITY CARRIER MOBILITY IN p-TYPE MATERIAL

In heterojunctions and quantum well structures where GaAs has the smallest bandgap of the materials forming the interface (as is the case, for example, in GaAs/AlGaAs structures), the low temperature electron mobility can be dramatically enhanced. For example the temperature dependence of the electron mobility shown in FIGURE 3 [15] can be compared with that of FIGURE 1. The physics and technological applications of such structures have been reviewed recently in an excellent book by Singh [16]. The mobility enhancement depends critically on the modulation doping of the wider bandgap material, the spacer layer thickness between the doping and the interface, as well as the quality of the interface and the background doping levels. Persistent photogenerated carriers can also have a marked effect on the mobility. For example, Skierbiszewski et al [17] have used the huge mobility enhancement observed when persistent carriers are photogenerated to infer the properties of the Ge donor in GaAs.

Electron Mobility

(cm 2 /V-s )

P-GaAs

FIGURE 4. Minority carrier electron mobility in p-type GaAs [14]. Solid and dashed curves represent Monte-Carlo and phase-shift analysis modelling, respectively.

For devices, the minority carrier mobility of electrons in p-type material is important, although not straightforward to measure. Recent results are reviewed by Adachi [14]. In particular, electron mobilities at 300 K have ranged from 1150 cm2/ Vs for p = 3.6 x 1018 cm"3 [19] to 3500 cm2/ Vs at p = 1017 cm'3 [20]. Recently an increase in mobility above 8 x io 18 cm"3 has been confirmed by Harmon et al [21]. They found an increase in electron mobility from 1370 cm2/ Vs to 3710 cm2/

Vs when the acceptor density was increased from 9 x 1018 cm"3 to 8 x 1019 cm"3. This effect is attributed to reduced plasmon and carrier-carrier scattering at higher hole densities.

REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] [II] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21]

D.C. Look [ Electrical Characterization of GaAs Materials and Devices (Wiley, 1989) ] P. Y. Yu, M. Cardona [ Fundamentals of Semiconductor: Physics and Materials Properties (Springer-Verlag, 1996) ] P.Blood, J.W. Orton [ The Electrical Characterization of Semiconductors: Majority Carriers and Electron States (Academic Press, 1992) ] CM. Wolfe, G.E. Stillman, W.T. Lindley [ J. Appl. Phys. (USA) vol.41 (1970) p.3088-91 ] D.L. Rode [ Semicond. Semimet. (USA) vol.28 Eds RK. Willardson, A.C. Beer (Academic, New York, 1975) ] W. Walukiewicz, L. Lagowski, L. Jastrzebski, M. Lichtensteiger, H.C. Gatos [ J. Appl. Phys. (USA) vol.50 (1979) p.899-908 ] M. Benzaquen, K. Mazuruk, D. Walsh, AJ. Springthorpe, C. Miner [ J. Electron. Mater. (USA) vol.16 (1987) p. 111-7] D. Lancefield, A.R. Adams, M.A. Fisher [ J. Appl. Phys. (USA) vol 62 (1987) p.2342-59 ] CM. Maziar, M.S. Lundstrom [ Electron. Lett. (UK) vol.22 (1986) p.565 ] D.C. Look, P.C Colter [ Phys. Rev. B (USA) vol.28 (1983) p. 1151-3 ] A.W.R. Leitch [ J. Appl. Phys. (USA) vol.65 (1989) p.2357-60 ] LY. Yanchev, S.K. Evtimova [ J. Phys. C (UK) vol. 18 (1985) p.L377-82 ] M.C Nuss, D.H. Auston, F. Cappasso [ Phys. Rev. Lett. (USA) vol.58 (1987) p.2355-8 ] S. Adachi [ GaAs and Related Materials: Bulk, Semiconducting and Superlattice Properties (World Scientific, 1994) ] T. Saku, Y. Horikoshi, Y. Tokura [ Jpn. J. Appl. Phys. (Japan) vol.35 (1996) p.34-8 ] J. Singh [ Physics of Semiconductors and their Heterostructures (McGraw-Hill, 1993) ] C. Skierbiszewski, P.Wisniewski, T. Suski, Z. Wilamowski [ Mater. Sd. Forum (Switzerland) no.I143/147(1994)p.l013] P.U. Koenraad et al [PhysicaB (Netherlands) \ol.2U (1995) p.462-5 ] M.I. Nathan, W.P.Dumke, K. Wrenner, S. Tiwari, S.L. Wright, K.A. Jenkins [ Appl. Phys. Lett. (USA) vol.52 (1988) p.654-6 ] H. Ito, T. Ishibashi [ J. Appl. Phys. (USA) vol.65 (1989) p.5197-9 ] ES. Harmon, M.L. Lovejoy, M.R. Melloch, M.S. Lundstrom, TJ. deLyon, J.M. Woodall [Appl. Phys. Lett. (USA) vol.63 (1993) p.536-8 ]

2.2

Electron mobility in bulk GaAs D. Lancefield July 1996

A

INTRODUCTION

Bulk GaAs for substrate material is required to have high purity, low defect density, good uniformity, good stability against heat treatment and large diameter [I]. High mobility substrates are required for structures formed by direct ion implantation although clearly it is not so important for devices where the conduction is predominantly in the epilayers. The principal growth techniques are liquid encapsulated Czochralski (LEC), vertical and horizontal Bridgman (HB) and vertical gradient freeze (VGF). LEC GaAs can be semi-insulating (SI) without doping but HB is normally doped with Cr to achieve SI characteristics. However, HB grown GaAs normally has better stoichiometry and reduced dislocation densities. SI material is important for FET technology [1] while doped substrates are important for infra-red lasers etc. [2]. Pimental and Look [3] have developed a non-destructive topographic assessment technique for determining the electron mobility and carrier density. Lyons and Vickers [4] have reported an electro-optic probing technique to measure the Hall effect in bulk GaAs, while a number of authors [5-7] have described the determination of electron mobility from measurements where mixed conduction effects are significant. B

ELECTRON MOBILITIES IN HB, LEC AND VGF SAMPLES

Parsey et al [8] have shown that low dislocation electron trap-free GaAs can be grown reproducibly by the Bridgman method. Carrier mobilities of 2400 and 3300 cm2/ Vs were obtained at room temperature for crystals grown from Ga2O3 doped and undoped melts, respectively. Bonnafe et al [9] reported that the Hall mobilities were in the range 1600 to 4200 cm2/ Vs at 350 K and 2400 to 3350 cm2/ Vs at 470 K. For LEC GaAs, mobilities were 890 to 4050 cm2/ Vs at 350 K and 1940 to 3400 cm2/ Vs at 470 K. Mizutani et al [10] reported that mobilities for various HB GaAs: Cr samples after Si implantation and annealing were 3600 - 4000 cm2/ Vs at room temperature and 4000 - 5000 cm2/ Vs at 77 K. LEC GaAs: Cr showed similar results but undoped LEC GaAs showed higher mobilities (4000 - 4400 cm2/ Vs at room temperature and 5500 - 6500 cm2/ Vs at 77 K). Lent et al [11] found that 400 K Hall mobilities for undoped LEC GaAs were in the range 1375 - 4270 cm2/ Vs depending on the crucible material, dopant concentration, growth rate, water content OfB2O3 encapsulant, post-growth annealing schedule and whether the sample was from the top, middle or bottom of the crystal. The melt composition and purity has been found to have a large effect on the electron mobility and the use of a pyrolytic boron nitride crucible significantly reduces Si contamination [12]. Fornari finds that the incorporation of Si is independent of melt composition but its distribution between donors, acceptors and complexes changes with stoichiometry [13]. The highest mobilities are achieved in substrates grown from As rich melts which are also SI. Look et al [14] report that near-stoichiometric wafers which are often not SI can be reversibly converted to a SI state by thermal processing. The effect on the electron mobility is complex but

normally results in an increased mobility. Rumsby et al [15] found that anneals above 650 0 C resulted in improved uniformity of the mobility across the wafer. Siegal et al [16] correlated electron mobilities in undoped LEC and VGF GaAs at various axial and radial positions with variations in the local carrier density. A number of authors have reported the effects of doping on the dislocation density and mobility. Barrett et al [17] find that In-doping of LEC GaAs significantly reduces the dislocation densities in large diameter wafers while maintaining 300 K mobilities of 5000 cm2/ Vs. Initial studies of implanted FETs fabricated on these wafers indicate a much tighter distribution of threshold voltage and source drain currents. Matsumoto et al [18] have grown three inch diameter low dislocation density wafers using a double crucible LEC design. Indium incorporation along the growth axis is more uniform with a mobility of about 4000 cm2/ Vs which increases slowly from seed to tail. Kimura et al [19] have used complementary phosphorus doping to relieve stresses in In-doped GaAs. They report low dislocation densities and mobilities in excess of 7000 cm2/ Vs at room temperature. In nominally undoped SI GaAs, Gray et al [20] have used quite small amounts of iron to control the fraction of EL2 ionized along the length of the boule. Room temperature mobilities ranged from 2900 cm2/ Vs to 5500 cm2/ Vs for carrier densities in the range 6 x io 11 cm"3 to 2 x io 12 cm"3. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] II1] [12] [13] [14] [15]

[16] [17] [18] [19] [20]

R.Nakai et al [ Sumitomo Electr. Tech. Rev. (Japan) no.23 (1984) p. 159-67 ] MR. Brozel, EJ. Foulkes, LR Grant, D.T.J. Hurle [ J Cryst. Growth (Netherlands) vol.80 (1987) p.323-32 ] E. Pimental, D.C. Look [ J Electron. Mater. (USA) vol. 17 (1988) p.63-6 ] V.R. Lyons, AJ. Vickers [ Semicond Sci. Technol (UK) vol.8 (1993) p.2058-61 ] N.C. Haider, D.C. Look [ J Appl. Phys. (USA) vol.66 (1989) p.4858-61 ] R-S. Tang, L. Sargent, J.S. Blakemore [J Appl. Phys. (USA) vol.66 (1989) p.256-61 ] R. Fornari [ Solid-State Electron. (UK) vol.29 (1986) p.589-90 ] J.M. Parsey Jr., Y. Nanishi, J. Lagowski, H.C. Gatos [J Electrochem. Soc. (USA) vol.128 (1981) p.936-8 ] J. Bonnafe et al [Mater. Res. Bull. (USA) vol.16 (1981) p.l 193-212 ] T. Mizutani et al [ Solid-State Electron. (UK) vol.25 (1982) p.885-91 ] B. Lent, M. Bonnett, N. Visentin, J.P.Duchemin [Microelectron. J (UK) vol.13 (1982) p.5-9 ] EM. Swiggard, S.H. Lee, F.W. VonBatchelder [Inst. Phys. Conf. Ser. (UK) vol.336 (1977) p.23 ] R Fornari [ J Cryst. Growth (Netherlands) vol.94 (1989) p.433-40 ] D.C. Look, W.M. Theis, P.W. Yu, J.R Sizelove, W. Ford, G. Mather [ J Electron. Mater. (USA) vol.16 (1987) p.63-8] D. Rumsby, I. Grant, M.R Brozel, EJ. Foulkes, RM. Ware [ Proc. Conf. Semi-Insulating IU-V Materials, Kah-nee-ta, OR, USA, 24-26 Apr 1984, Eds D.C. Look, J.S. Blakemore (Shiva Publishing, Nantwich, UK, 1984) p. 157-9 ] W. Siegal, G. Kuhnel, U. Kretzer [Mat. Sci. Eng. B (Switzerland) vol.28 (1994) p.84-86 ] D.L. Barrett, S. McGuigan, H.M. Hobgood, G.W. Eldridge, RN. Thomas [ J Cryst. Growth (Netherlands) vol.70 (1984) p. 179-84 ] K. Matsumoto, M. Yamashita, R Nakai, S. Yazu, K. Tada, S. Akai [ Proc. Conf. Semi-Insulating III-VMaterials, Malmo, Sweden, 1-3 June 1988 (Adam Hilger, Bristol, UK, 1988) p.447-52 ] H. Kimura, AT. Hunter, E.-H. Cirlin, H.M. Olsen [ J Cryst. Growth (Netherlands) vol.85 (1987) p. 116-23] M.L. Gray, L. Peterson, R-S. Tang, S.B. Saban, J.S. Blakemore [ J Appl. Phys. (USA) vol.73 (1993) p.3319-25]

2.3

Electron mobility in LPE GaAs D. Lancefield July 1996

In recent years the use of LPE has diminished because of its inability to control layer thicknesses, compositions and doping densities to the fine tolerances required for a wide range of device research and increasingly for device production. However, the electron mobilities measured in GaAs grown by this technique are still informative. In 1968 LPE GaAs with carrier concentrations, n, of about 2 x 1014 cm"3 had 300 K mobilities of 8000 cm2/ Vs [1], while at 77 K Solomon [2] had obtained mobilities of 1.1 x 105 cm2/ Vs for a carrier concentration of about 5 x 1012 cm"3, which he ascribed to residual oxygen impurities incorporated during growth. Since that time the carrier concentration dependence of the electron mobility has been measured by a number of authors including Vilms and Garrett [3], Panish [4], and Rosztoczy and Kinoshita [5]. Vilms and Garrett [3] measured mobilities in Sn-doped GaAs at 300 K varying from 9000 cm2/ Vs for n = 5 x 1013 cm"3 to about 1000 cm2/ Vs at n = 1019 cm"3. At 77 K the mobilities decreased from about 13 x io 5 cm2/ Vs at n = 5 x io 13 cm"3 to about 8000 cm 2 /Vs at n = 4 x 1017cm"3. Abrokwah et al [6] have shown that the simultaneous bake out of both melt and substrate at 7750 C results in layers with carrier concentrations of about 2 x IO14 cm"3 and 77 K mobilities in excess of 105 cm2/ Vs. Morkoc and Eastman [7] obtained mobilities at 77 K of 163000 cm2/ Vs with carrier concentrations of 9 x io 13 cm"3 using a 24 hour bake out and a growth temperature of 7000C. Bake out is considered to remove volatile impurities while C contamination from the graphite boat is the main remaining impurity. Hicks and Manley [8] using a Spectrosil dipping boat, rather than graphite, reported on a high purity layer with liquid nitrogen mobilities in excess of 1.7 x IO5 cm2/ Vs. Otsubo and Miki [9] obtained 77 K mobilities of between 1.7 x io 5 cm2/ Vs and 2 x 105 cm2/ Vs reproducibly by LPE under a mixed gas flow of purified H2 and arsenic vapour which was found to reduce contamination due to oxygen, Si and Cu. Using a Spectrosil quartz reaction tube and a glassy carbon boat they obtained mobilities of up to 2.44 x 105 cm2/ Vs for samples with carrier concentrations down to 3 x io 12 cm"3 [10]. Otsubo and Miki [10] obtained 77 K mobilities of between 1.7 x io 5 cm2/ Vs and 2 x io 5 cm2/ Vs reproducibly by LPE under a mixed gas flow of purified H2 and arsenic vapour which was found to reduce contamination due to oxygen, Si and Cu. A comparison of growth from bismuth and gallium melts has been made by Yakusheva et al [11]. Using a Bi melt they obtained 77 K mobilities of 1.5 x io 5 cm2/ Vs with electron concentrations of 2.5 x IO14 cm"3, with Sn as the main donor and Zn, Mg and Be the main acceptors. Bryskiewicz et al [12] have studied high quality liquid phase electroepitaxy where growth is controlled by passing a direct electric current across the substrate-solution interface while the temperature of the melt is maintained constant. Dislocation densities are found to decrease rapidly and for carrier densities of 9 x io 14 cm"3, mobilities of 7200 and 5.4 x io 5 cm2/ Vs were obtained

at 300 K and 77 K respectively. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] [II] [12]

A.R. Goodwin, CD. Dobson, J. Franks [ Inst. Phys. Conf. Ser. (UK) no.7 (1968) p.36 ] R. Solomon [ Inst. Phys. Conf. Ser. (UK) no.7 (1968) p. 11 ] J. Vilms, J.P. Garrett [ Solid-State Electron. (UK) vol. 15 (1972) p.443-55 ] M.B. Panish [ J. Appl. Phys. (USA) vol.44 (1973) p.2667-75 ] F.E. Rosztoczy, J. Kinoshita [ J. Electrochem. Soc. (USA) vol. 121 (1974) p.439-44 ] J.K. Abrokwah, M.L. Mitchell, J.E. Borell, D.R. Schulze [ J. Electron. Mater. (USA) vol. 10 (1981) p.723-46 ] H. Morkoc, L.F. Eastman [ J. Cryst. Growth (Netherlands) vol.36 (1976) p. 109-14 ] H.G.B. Hicks, D.F. Manley [ Solid State Commun. (USA) vol.7 (1969) p.1463-5 ] M. Otsubo, H. Miki [ Jpn. J. Appl. Phys. (Japan) vol.14 (1975) p.621-8 ] H. Miki, M. Otsubo [ Jpn. J. Appl. Phys. (Japan) vol. 10 (1971) p.509 ] N.A. Yakusheva, K. S. Zhuravlev, S.I. Chikichev, O.A. Shegaj [ Cryst. Res. Technol. (East Germany) vol.24 (1989) p.235-46 ] T. Bryskiewicz, M. Bugaski, J. Lagowski, H.C. Gatos [ J. Cryst. Growth (Netherlands) vol.85 (1987)p.l36-41]

2.4

Electron mobility in VPE and MOVPE GaAs D. Lancefield July 1996

A

INTRODUCTION

Vapour phase epitaxy of GaAs (VPE) has been studied using a number of growth techniques. These include chloride, hydride and metalorganic-VPE (MOVPE) depending primarily on the starting materials [I]. We will consider the first two and then consider MOVPE. However, we note that electron mobilities of up to 9000 cm2/ Vs at room temperature and up to 2.1 x io 5 cm2/ Vs at 77 K with a peak mobility of 3.35 x IO5 cm2/ Vs at 38 K have been achieved by MOVPE [2]. This is amongst the highest purity GaAs reported to date. Yao-Wang et al [3] using an AsCl3-Ga-Ar system in which Si contamination is suppressed have obtained 77 K mobilities of 2.05 x io 5 cm2/ Vs with carrier densities of 7.4 x 1012 cm"3 and peak mobilities at 35 K of 3.78 x 10 5 cm 2 /Vs. B

VPE GaAs

Enstrom and Appert [4] showed that VPE layers grown using the hydride method with cooled HCl had superior mobilities to those grown using uncooled HCl. The HCl cooling increased room temperature mobilities by about 20% and 77 K mobilities by up to 75%. Look and Colter [5] obtained high mobility and very low compensation by the hydride method by pre-baking the Ga source at 8200C for 16 hours immediately prior to growth. The mobility of this material also showed a maximum (about 2 x io 5 cm2/ Vs) at 50 K. An additional maximum, thought to be associated with extremely low compensation, occurred near 9 K. Gudz et al [6] have shown that for high purity VPE GaAs deposited onto semi-insulating GaAs substrates, etching away of the surface layer led to a marked increase in the mobility at high temperatures (100 - 300 K). For both unetched and etched samples there was a maximum in the mobility near 50 K. At 50 K the mobility was 5.07 x 10 cm2/ Vs and 5.17 cm2/ Vs for unetched samples and a sample with 5 microns etched off respectively. Poth et al [7] showed that the Hall mobility of VPE S-, Se-, or Si-doped GaAs films was strongly dependent on free carrier concentration but was independent of growth parameters and they produced a universal mobility vs. carrier density curve. Since then several authors have presented data on decreases in mobility in VPE GaAs with increasing doping density. However, Hollan et al [8] reported that the mobility of S-doped VPE GaAs was less than that of Se-doped layers for carrier concentrations above 4 x io17 cm"3. For carrier densities of 1018 cm"3, Hall mobilities were 2900 cm2/ Vs for GaAs:Se and 2600 cm2/ Vs for GaAs:S. For carrier densities of 4 x io 18 cm"3 mobilities of GaAs:Se and GaAs:S were 1900 and 1500 cm2/ Vs, respectively. An interesting modification of the VPE process is vapour levitation epitaxy (VLE) [9]. Mobilities at 77 K of 9.2 x IO4 cm2/ Vs have been reported. Somogyi has recently reported that the lowering of the mobility compared with theory can usually be accounted for by errors in determining the conducting layer thickness [10]. Kourkoutas et al [11] have reported that electron irradiation of doped GaAs results in a decreased carrier density and mobility.

C

MOVPE GaAs

In MOVPE (or MOCVD or OMVPE, etc.) grown semiconductors room temperature vapours of organometallic compounds are used to transport at least one of the primary constituents. Initially this was normally triethylgallium (TEG), which together with arsine, AsH3, was used to form GaAs. However, the technique was of particular interest because of the wide range of materials that can be grown in both III-V and II-VI materials. Chang et al [12] have reported that films with room temperature Hall mobilities above 7000 cm2/ Vs can be grown by MOVPE using TEG as the Ga source under optimum growth conditions. Smith has obtained 77 K mobilities of above 1.2 x 105 cm2/ Vs at pressures of 1.5 ton* with As:Ga ratios in excess of 300 [13]. Fraas [14] found that GaAs layers deposited from TEG and AsH3 using a vacuum MOVPE system at growth pressures of 10"2 ton* exhibited room temperature mobilities of 4990 cm2/ Vs for unintentionally doped layers with a carrier density of 1016 cm'3. Kuech and Potemski [15] using a similar system obtained mobilities of 9.7 x 104 cm2/ Vs at 77 K with an As:Ga ratio of 20, and a growth pressure and temperature of 0.1 atm and 650 0 C, respectively. Shastry et al [16] using TEG and arsine obtained 77 K mobilities of 2.1 x 105 cm2/ Vs with a compensation ratio of 0.5 and Nd+Na approximately 1014 cm"3. Using similar starting materials, Razeghi et al obtained mobilities of 3.35 x 105 cm2/ Vs at 35 K [2]. Using trimethylgallium (TMG) instead of TEG, Hata et al [17] achieved 77 K mobilities of 1.53 x 105 cm2/ Vs. The residual donors were identified as Ge and Si with C as the main acceptor. Because of the toxicity of AsH3, less toxic arsenic sources have been studied. Using dimethylarsine, Chen et al [18] found that undoped GaAs was p-type, but that with Te doping they achieved room temperature mobilities of 5000 cm2/ Vs. Using TMG and diethylarsine Bhat et al [19] obtained 77 K mobilities of 6.5 x 104 cm2/ Vs at a carrier concentration around 3 x 1014 cm"3. Tertiarybutylarsine has also been used [20] to get 77 K mobilities of 8.0 x 104 cm2/ Vs. Tertiarybutylarsine, although highly toxic, has improved safety as it is stored in a bubbler rather than a high pressure vessel. Exposure to a hydrogen plasma has been found to reduce the free carrier density and increase the electron mobility [21]. n decreased from 8.1x io 15 cm"3 to 3.4 x 1015 cm"3 with a corresponding increase in peak mobility from 1.8 x io 5 to 2.3 x io 5 cm2/ Vs. Studies of vapour flow in reactors have been undertaken, 77 K mobilities of 1.4 x io 5 cm2/ Vs for a carrier concentration of IO14 cm"3 being obtained under laminar flow conditions [22]. Kanber et al [23] have studied multiwafer systems and achieved uniform doping-thickness products with mobilities in the range 5500 to 2000 cm2/ Vs for electron concentrations between 5 x io 16 and 5 x 1018 cm"3. The effects of laser irradiation on growth enhancement and its use for selective area growth have been investigated. Kachi et al [24] found that photostimulated MOVPE irradiation with a Co2 laser enhanced growth rates while UV irradiation improved crystal quality resulting in room temperature mobilities of about 6000 cm2/ Vs. Kukimoto et al [25] used an ArF excimer laser excitation for selective area growth of GaAs with mobilities of 6000 cm2/ Vs and carrier densities of about IO16 cm"3 at room temperature. Jones and Lau [26] used native oxides to achieve selective area growth of single crystal material with 77 K mobilities of 1.05 x io 5 cm2/ Vs. Huelsman et al [27] used an RF discharge for the dissociation of arsine to improve mobilities of layers grown at low temperature.

Common dopants in MOVPE are S, Si, Sn and Te. Lee and Chang [28] used tetraethyltin and obtained room temperature mobilities of 5000 to 3000 cm2/ Vs for carrier densities in the range 2 x 1016 to 2 x 1017 cm"3 with an enhanced growth rate and a decreased As:Ga ratio. Parsons and Krajenbrink [29] found that 300 K mobilities of GaAs: Sn with carrier densities in the range 7 x 1014 to 1019 cm"3 had mobilities of 8000 and 1300 cm2/ Vs, respectively. Te doping has been investigated by Houng and Low [30] using diethyltellurium. Carrier densities from 1015 to 1019 cm"3 were obtained with mobilities from 6500 to 1700 cm2/ Vs, respectively. Memory effects could be reduced by heating the dopant line and reducing the surface area of pipework. Doping with silane and disilane and its incorporation has been considered recently by Tang et al [31]. For disilane, mobilities in the range 4000 to 2000 cm2/ Vs for carrier densities of 1017 to 5 x 1018 cm"3, respectively, were obtained. The growth of GaAs on substrate material other than GaAs has been investigated. GaAs on sapphire [32] had mobilities which were strongly dependent on layer thickness with room temperature mobilities of 3000 cm2/ Vs. Shastry and Zemon [33] have grown GaAs on Si with carrier densities less than 1014 cm"3. Intentionally doped layers of 1016 cm"3 had room temperature mobilities of 5800 cm2/ Vs. Mobilities in low dimensional structures have also been extensively studied. For example, Makimoto and Kobayashi [34] report the Si atomic layer doping of AlGaAs/GaAs. They obtained a maximum low temperature mobility of 1.3 * 106 cm2/ Vs for a sheet carrier density of 4.8 x io 11 cm"2. Basco et al [35] using a conventional low-pressure OMVPE system achieved maximum mobilities of 7.7 x 105 cm2/ Vs at 2 K and 1.71 x io 5 cm2/ Vs at 77 K with a sheet carrier density of about 4.9 x io11 cm"2 after the sample was illuminated. They found that preconditioning of the reactor enhanced the mobility measured prior to illumination while the inclusion of an AlGaAsrelated buffer degraded the mobility. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] II1] [12] [13] [14] [15] [16] [17]

M. Yoshida [ J Cryst. Growth (Netherlands) vol.88 (1988) p. 16-22 ] M. Razeghi, F. Omnes, J. Nagle5 M. Defour, O. Acher, P. Bove [Appl. Phys. Lett. (USA) vol.55 (1989) p. 1677-9] Lin Yao-Wang, Zhang Yan-Yun, Li Hsiu-Lan, Liang Jan-Wu, Lin Lan-Ying [ J. Cryst. Growth (Netherlands) vol.70 (1984) p. 108-11 ] R.E. Enstrom, J.R. Appert [ J Electrochem. Soc. (USA) vol. 129 (1982) p.2566-9 ] D.C. Look, P.C. Colter [ Phys. Rev. B (USA) vol.28 (1983) p. 1151-3 ] E.S. Gudz et al [ Inorg. Mater. (USA) vol. 16 (1980) p. 128-32 ] H. Poth, H. Bruch, M. Heyen, P. Balk [ J Appl. Phys. (USA) vol.49 (1978) p.285-8 ] L. Hollan, M. Boulou, J.P. Chane [J Electron. Mater. (USA) vol.10 (1981) p.193-212 ] H.M. Cox, S.G. Hummel, V.G. Keramidas [J Cryst. Growth (Netherlands) vol.79 (1986) p.900-8 ] K. Somogyi [ Acta Phys. Hung. (Hungary) vol.74 (1994) p. 107-20 ] CD. Kourkoutas et al [ Phys. Status Solidi A (Germany) vol. 135 (1993) p.K21 ] CY. Chang, Y.K. Su, M.K. Lee, L.G. Chen, M.P. Houng [ J Cryst. Growth (Netherlands) vol.55 (1981)p.24-9] F.TJ. Smith [J. Cryst. Growth (Netherlands) vol.67 (1984) p.573-8 ] L.M. Fraas [J Appl. Phys. (USA) vol.52 (1981) p.6939-43 ] T.F. Kuech, R. Potemski [Appl. Phys. Lett. (USA) vol.47 (1985) p.821-3 ] S.K. Shastry, S. Zemon, D.G. Kenneson, G. Lambert [Appl. Phys. Lett. (USA) vol.52 (1988) p. 150-2] M. Hata, N. Fukuhara, Y. Zempo, M. Isemura, T. Yako, T. Maeda [ J Cryst. Growth (Netherlands)

[18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33] [34] [35]

vol.93 (1988) p.543-9 ] CH. Chen, E.H. Reihlen, G.B. Stringfellow [ J. Cryst. Growth (Netherlands) vol.96 (1989) p.497-504 ] R. Bhat, M.A. Koza, BJ. Skromme [Appl. Phys. Lett. (USA) vol.50 (1987) p.1194-6 ] G. Haacke, SP. Watkins, H. Burkhard [ Appl. Phys. Lett. (USA) vol.54 (1989) p.2029-31 ] A. Jalil, J. Chevallier, R. Azoulay, A. Mircea [ J. Appl. Phys. (USA) vol.59 (1986) p.3774-7 ] K. Matsumoto, K. Itoh, T. Tabuchi, R. Tsunoda [ J. Cryst. Growth (Netherlands) vol.77 (1986) p.151-6] H. Kanber, T. Zielinski, J.M. Whelan [ J. Electron. Mater. (USA) vol. 14 (1985) p.769-81 ] T. Kachi, H. Ito, S. Terada [ Jpn. J. Appl. Phys. 2 (Japan) vol.27 (1988) p.L1556-8 ] H. Kukimoto, Y. Ban, H. Komatsu, M. Takechi, M. Ishizaki [J. Cryst. Growth (Netherlands) vol.77 (1986)p.223-8] S.H. Jones, K.M. Lau [ J. Electrochem. Soc. (USA) vol. 134 (1987) p.3149-55 ] A.D. Huelsman, R. Reif, CG. Fonstad [ Appl. Phys. Lett. (USA) vol.50 (1987) p.206-8 ] M.K. Lee, CY. Chang [ J. Appl. Phys. (USA) vol.60 (1986) p.2831-4 ] J.D. Parsons, F.G. Krajenbrink [ J. Cryst. Growth (Netherlands) vol.68 (1984) p.60-4 ] Y.-M. Houng, T.S. Low [ J. Cryst. Growth (Netherlands) vol.77 (1986) p.272-80 ] X. Tang, H.G.M. Lochs, P.R. Hageman, M.H.J.M. De Croon, LJ. Giling, AJ. Bons [ J. Cryst. Growth (Netherlands) vol.98 (1989) p.827-37 ] K. Kasai, K. Nakai, M. Ozeki [ J. Appl. Phys. (USA) vol.60 (1986) p. 1-5 ] S.K. Shastry, S. Zemon [ Appl. Phys. Lett. (USA) vol.49 (1986) p.467-9 ] T. Makimoto, N. Kobayashi [ Jpn. J. Appl. Phys. (USA) vol.32 (1993) p.L648-9 ] R. Basco, F. Agahi, K.M. Lau [ Appl. Phys. Lett. (USA) vol.63 (1993) p. 1960-2 ]

2.5

Electron mobility in MBE GaAs D. Lancefield July 1996

Molecular beam epitaxy (MBE) is a popular growth technique for fine control of semiconductor layer thickness, composition and doping density. Together with MOCVD it constitutes the main technique for research and future commercial applications [I]. Numerous variations on the basic technique are now studied. These include metalorganic MBE (MOMBE) [2], chemical beam epitaxy (CBE) [3], and atomic layer epitaxy for 'digital epitaxy' [4,5]. For pure material, mobilities similar to those achieved in VPE grown material have been reported. They depend on HI: V flux ratio [6] and the elimination of flux transients [7]. Improved mobilities are observed at higher growth rates due to the reduced acceptor incorporation [8]. Chin [9] determined the piezoelectric constant to be 1.77 x 107 V cm"1 for the highest purity material. The peak mobility was about 4.5 x 105 cm2/ Vs at 40 K for samples with a carrier density of 2.8 x 1013 cm"3 with a compensation ratio of 0.2. Holland et al [10] measured similar mobilities in GaAs grown using an As2 source. They observed a progressive improvement in the mobility which they attributed to a clean-up of the As2 source. Ballingall et al [11] obtained mobilities of 6.0 x 104 cm2/ Vs at carrier densities of 1.4 x 1015 cm"3. Stanaway et al [12] have grown samples with peak mobilities of 2.55 x io 5 cm2/ Vs. These samples had acceptor densities of 1.7 x io 13 cm"3 and donor densities of 6.3 x io13 cm"3. Residual donors were identified as Si, S and Sn while C was the principal acceptor. Using arsine for the arsenic source, Cunningham measured peak mobilities near 3.0 x io 5 cm2/ Vs [13]. Chand et al [14] found that changing the arsenic from 6N grade to 7N grade reduced residual acceptor densities by two orders of magnitude and allowed peak mobilities of 2.947 x 105 cm2/ Vs to be achieved at doping levels of 3 x io 13 cm"3. The low temperature electron mobility in high purity GaAs studied by cyclotron resonance gives a maximum mobility of 2.2 x IO6 cm2/ Vs at 1.6 K [15]. In modulation doped GaAs/AlGaAs electron mobilities as high as 11.7 x 106cm2/ Vs have been achieved at low temperatures after illumination. The sheet carrier density was 2.4 x io 11 cm"2 [16]. The highest mobility samples often have superlattice buffers to smooth out substrate interface roughness. Etched and regrown AlGaAs/GaAs interfaces have given mobilities of 1.2 x io 5 cm2/ Vs for sheet carrier densities of 4.5 x io11 cm'2 when measured at 10 K [17]. Yang et al [18] show that DX centres in AlGaAs play an important role in determining the low temperature carrier density and mobility. For 5-doped samples Zheng et al [19] report a strong mobility enhancement of up to a factor of 5 for coupled 5-doped layers separated by about 200 A. Si is the main dopant used for electron densities up to 5 x io 18 cm"3. Pan et al [20] have reported passivation of residual Si by exposure of the sample to a hydrogen plasma. The mobility increases by up to 40%. Sn has been used to achieve doping levels up to 1 x IO19 cm"3 at growth temperatures of about 600 0 C with relatively low compensation and mobilities of 1000 cm2/ Vs at room temperature [21]. High Si doping can be achieved at reduced growth temperatures with mobilities of 800 cm2/ Vs at doping levels of 1.3 x IO19 cm"3 [22]. Missous and Singer [23] note that the mobility is significantly higher, i.e. 4000 cm2/ Vs, using an As2 source rather than As4 at low growth temperatures with a reduced growth rate of 0.2 |im/hr, in agreement with Neave et

al [24] who found a reduced trap concentration in GaAs grown with As2. The use of a variety of other dopant species via low energy ion doping has been reported for electron densities up to 3 x 1018 cm'3 [25]. For smooth surfaces, mobilities were in the range 1500 - 2200 cm2/ Vs. In doping has attracted attention as a means of reducing defect concentrations [26]. Lee et al [27] measured mobilities of 4.7 x 104 cm2/ Vs for In doping levels of 2 x 1018 cm"3. For non-uniform 5-doped samples Schubert et al [28] reported layers with a 2DEG density of 2 x 1013cm'2, corresponding to a bulk doping density of 9 * 1019 cm"3. Such high doping densities are attributed to reduced Si autocompensation and have corresponding electron mobilities of 2200 cm2/ Vs. Growth of GaAs on substrates with various orientations as well as differing lattice constants has also been undertaken. Kasahara et al [29] obtained mobilities of 3500 cm2/ Vs for GaAs grown on InP, which has a lattice mismatch of 3.8%. Wang et al [30] found that undoped layers grown on (100) and (311)A orientations were p-type but that layers grown on (311)B orientation were n-type and achieved mobilities of 1.5 x 105 cm2/ Vs. Allen et al [31] growing on (110) orientation obtained mobilities of 5700 cm2/ Vs at room temperature. Modifications to the MBE system design including direct radiative substrate heating [32,33] and automated substrate transport mechanisms have been investigated [34,35] with the aim of high throughput capability. Using such systems 3 inch wafers with mobilities better than 7000 cm2/ Vs at room temperature and 9.0 x 104 cm2/ Vs at 77 K have been obtained. Betko et al have measured the electron mobility of low temperature GaAs. This material is of interest because of its high resistivity and fast non-radiative recombination. Room temperature mobilities in the range 0.1 to 10 cm2/ Vs were measured [36]. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] [II] [12] [13] [14]

L.L. Chang, R Ludeke [ Epitaxial Growth Ed. J.W. Matthews (Academic, New York, 1975) pt. A p.37] H. Luth [ Two-Dimensional Systems: Physics and New Devices; Proc. Int. Winter School, Mauterndorf, Austria, 24-28 Feb 1986 (Springer-Verlag, Berlin, 1986) p.12-23 ] W.T. Tsang [ J. Electron. Mater. (USA) vol. 15 (1986) p.235-45 ] H. Watanabe, S. Usui [ Inst Phys. Conf. Ser. (UK) no.83 (1987) p. 1-8 ] K. Mori, M. Yoshida, A. Usui, H. Terao [Appl Phys. Lett. (USA) vol.52 (1988) p.27-9 ] S.C. Palmateer, P.A. Maki, W. Katz, A.R Calawa, J.C.M. Hwang, L.F. Eastman [ Inst. Phys. Conf. Ser. (UK) no. 74 (1985) p.217-22 ] P.A. Maki, S.C. Palmateer, A.R Calawa, B.R Lee [J. Vac. Sci. Technol. B (USA) vol.4 (1986) p.564-7 ] Y.G. Chai [ Appl. Phys. Lett. (USA) vol.37 (1980) p.379 ] V.W.L. Chin [ Solid-State Electron. (UK) vol.37 (1994) p. 1345-7 ] M.C. Holland, A.H. Kean, J.M. Chamberlain, RT. Grimes, M.B. Stanaway [ J. Cryst. Growth (Netherlands) vol.111 (1991)p.l4-9] J.M. Ballingall, BJ. Morris, DJ. Leopold, D.L. Rode [ J Appl. Phys. (USA) vol.59 (1986) p.3571-3] M.B. Stanaway et al [ Inst. Phys. Conf. Ser. (UK) no.95 (1989) p.295-300 ] J.E. Cunningham et al [ J. Cryst. Growth (Netherlands) vol.95 (1989) p. 185-8 ] N. Chand, RC. Miller, A.M. Sergent, S.K. Sputz, D.V. Lang [Appl. Phys. Lett. (USA) vol.52 (1988)

[15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33] [34] [35] [36]

p.1721-3] M. Kozhevnikov, B.M. Ashkinadze, E. C(ACT5A. Ron [Phys. Rev. B (USA) vol.52 (1995) p.1716571] L. Pfeiffer, K.W. West, H.L. Stormer, K.W. Baldwin [ Appl. Phys. Lett. (USA) vol.55 (1989) p.1888] Y. Kadoya, H. Noge, H. Kano, H. Sakaki [ J. Cryst. Growth (Netherlands) vol. 125 (1993) p.87780] B. Yang et al [ Appl. Phys. Lett. (USA) vol.66 (1995) p. 1406-8 ] X. Zheng, T.K. Cams, K.L. Wang, B. Wu [ Appl. Phys. Lett. (USA) vol.62 (1993) p.504-6 ] N.Pan et al [ Appl. Phys. Lett. (USA) vol.50 (1987) p. 1832-4 ] H. Ito, T. Ishibashi [ Jpn. J. Appl. Phys. 2 (Japan) vol.26 (1987) p.L1760-2 ] M. Ogawa [ Inst. Phys. Conf. Ser. (UK) no. 79 (1986) p. 103-8 ] M. Missous, K.E. Singer [Appl. Phys. Lett. (USA) vol.50 (1987) p.694-5 ] J.H. Neave, P. Blood, B.A. Joyce [ Appl. Phys. Lett. (USA) vol.36 (1980) p.311-2] S. Cavalieri, Ph. Gaucherel, G. Monnom, C. Paparoditis, J.C. Roustan [J. Vac. Sci. Technol. A (USA) vol.5 (1987) p. 1421-4 ] A. Uddin, T.G. Andersson [ J. Appl. Phys. (USA) vol.65 (1989) p.3101-6 ] M.K. Lee, T.H. Chiu, A. Dayem, E. Agyekum [ Appl. Phys. Lett. (USA) vol.53 (1988) p.2653-5 ] E.F. Schubert, J.E. Cunningham, W.T. Tsang [ Solid-State Commun. (USA) vol.63 (1987) p.591-4 ] K. Kasahara, K. Asano, T. Itoh [ Inst. Phys. Conf. Ser. (UK) no.91 (1988) p. 195-8 ] W.I. Wang, RF. Marks, L. Vina [ J. Appl. Phys. (USA) vol.59 (1986) p.937-9 ] L.T.P. Allen, E.R. Weber, J. Washburn, YC. Pao [ Appl. Phys. Lett. (USA) vol.51 (1987) p.670-2 ] L.P. Erickson et al [ J. Vac. Sci. Technol. B (USA) vol.3 (1985) p.536-7 ] E.S. Hellman et al [ J. Vac. Sci. Technol. B (USA) vol.4 (1986) p.574-7 ] J. Sakai, S. Murakami, K. Hirama, T. Ishida, Z. Oda [ J. Vac. Sci. Technol. B (USA) vol.6 (1988) p. 1657-61] T. Sonoda, M. Ito, M. Kobiki, K. Hayashi, S. Takamiya, S. Mitsui [ J. Cryst. Growth (Netherlands) vol.95 (1989) p.317-21] J. Betko, P. Kordos, S. Kuklovsky, A. Forster, D. Gregusova, H. Luth [ Mater. Sci. Eng. B (Switzerland) vol.28 (1994) p. 147-50 ]

2.6

Electron mobility in ion implanted GaAs BJ. Sealy August 1995

A

INTRODUCTION

This Datareview summarises the electron mobilities measured in donor implanted GaAs. B

DISCUSSION

The electrical properties of all of the donor dopants, Si, Ge, Sn, S, Se and Te, have been studied following implantation into GaAs. Most publications have concentrated on the use of silicon implants because of its widespread use for device fabrication. Due to its relatively small mass, silicon does not produce much lattice damage and, hence, it is straightforward to get high levels of electrical activity following room temperature implants. At the same time, the annealing cycle repairs the lattice and produces good quality material with high mobilities. For example, doses of 1-5 x 1012 Si+ cm"2 (and also Se [1,2]) can be close to fiill activity after annealing at 8500C to 900 0 C, with mobilities in the range 4000 to 4500 cm2 / Vs and peak electron concentrations of around 1017cm"3 [3-5]. As the dose increases to about 1015 Si+ cm"2, the net percentage activity reduces and so does the mobility, to a value of about 1500 cm2 / Vs [6-8]. The decrease in mobility is linked more closely to an increase in the sheet carrier concentration than to the dose and the sheet carrier concentration varies with ion dose, ion species, ion energy and thermal cycle. For doses above about 1014 cm"2, the mobility tends to decrease with increasing annealing temperature due to the simultaneous (and often large) improvement in electron concentration (electrical activity) [7-10]. However, for silicon implants, temperatures above about 9500C should be avoided due to its amphoteric behaviour, when both the carrier concentration and the mobility become degraded. Low dose selenium implants behave very similarly to silicon implants in producing high electrical activities and mobilities. However, the other ion species are not so useful because they tend to introduce damage that is difficult to remove by annealing. In the case of germanium, the atoms are amphoteric and self-compensation occurs, producing very low mobilities. Of the heavier mass donor atoms, tin is the ion which activates most easily with good mobilities, but for most implanters (limited to 200 kV) it has a very small range and is therefore not suitable for device fabrication. For example, for a dose of 1 x 1015 cm"2, it was found that tin implants produce higher mobilities than selenium implants following an anneal at 10000C for 10 s [7]. The implication is that the selenium doped material is more highly compensated than the tin doped material. C

CONCLUSION

The magnitude of the electron mobility (4000 to 4500 cm2/ Vs) for GaAs implanted with low doses of Si or Se ions (1-5 x 1012 cm"2) is close to values measured for good quality crystalline material. High doses produce higher carrier concentrations and lower mobilities, the values tending to be lower than expected from crystalline material due to electrical compensation.

REFERENCES [1] [2] [3] [4] [5] [6] [7] [8] [9] [10]

P.A. Leigh [ IEEProc. I (UK) vol. 130 (1983) p.265-74 ] NJ. Bairett, J.D. Grange, BJ. Sealy, K.G. Stephens [J Appl. Phys. (USA) vol.56 (1984) p.3503-7 ] CP. Stewart, RT. Blunt, GR. Booker, LR. Saunders [PhysicaB (Netherlands) vol. 116 (1983) p.635-40 ] K.S. Seo, S. Dhar, P.K. Bhattacharya [ Appl Phys. Lett. (USA) vol.47 (1985) p.500-2 ] M.H. Badawi, J. Mun [ Electron. Lett. (UK) vol.20 (1984) p. 125-6 ] W.H. Haydl [IEEEElectron. Device. Lett. (USA) vol.5 (1984) p.78-81 ] K.K. Patel, R. Bensalem, M.A. Shahid, BJ. Sealy [ Nucl. Instrum. Methods Phys. Res. B (Netherlands) vol.7/8 (1985) p.418-22 ] Y.K. Yeo, R. Kwor, Y.S. Park [J Appl. Phys. (USA) vol.53 (1982) p.1812-4 ] DE. Davies,PJ. McNaIIy, TG. Ryan, KJ. Soda, JJ. Comer [Inst. Phys. Conf. Ser. (UK) no.65 (1983)p.619-25] R. Kwor, Y.K. Yeo, Y.S. Park [J. Appl. Phys. (USA) vol.53 (1982) p.4786-92 ]

2.7

Electron mobility in GaAs, pressure dependence A.R. Adams and V.A. Wilkinson July 1995

Lancefield et al [1,2] showed that at room temperature the electron mobility (^6) in ultra-pure GaAs decreases with pressure at a rate of d[ln(jic / |i J ] / dP = -7.3 x \0'2 GPa 1

(1)

Normalised Hall mobility

(where ^60 is the electron mobility at atmospheric pressure) in the pressure range 0-0.8 GPa. This is due predominantly to the increasing electron effective mass with pressure. The value was determined from Hall effect and resistivity measurements on samples grown by LPE on semiinsulating substrates. Lancefield et al [1,2] compared their results with the theoretical pressure variation of the Hall mobility calculated by iterative solution of the Boltzmann equation including piezoelectric, deformation potential, and polar optical scattering. They were able to obtain good agreement between theory and experiment.

Experimental Uncertainty

Hydrostatic presssure (GPa) FIGURE 1. Pressure dependence of the Hall mobility for samples of LPE GaAs normalised to atmospheric pressure. Reproduced from reference [2] where full details of the line fits are given. Reproduced with permission of Journal of Applied Physics.

Next Page

In the LPE samples studied by Lancefield et al [1,2] it was observed that the pressure dependence of the electron mobility increased as the mobility was decreased by ionized impurity scattering as shown in FIGURE 1. In samples with a mobility of 1600 cm2/ Vs the pressure dependence of the mobility is d[ln(m / j i j ] / dP = -1.63 x 10"1 GPa 1

(2)

Correlated impurity scattering as described by Yanchev et al [3] appears to partially explain the results. Recent measurements [4] indicate that the pressure coefficient of the electron mobility is almost independent of temperature down to 20 K for lightly doped samples and for heavily doped samples down to 4 K, but for intermediately doped samples (1017 cm"3), where hopping conduction is believed to be important below 20 K, the magnitude of the pressure coefficient increases to d[ln(|ie / J i J ] / dP = -2.6 * 10"1 GPa 1

(3)

Above about 3.3 GPa at room temperature Pitt and Lees [5] observed a swift decrease in mobility to 375 cm2/ Vs, while Patel et al [6] saw a decrease to about 150 cm2/ Vs. This can be attributed to electron transfer to the X minima where the mobility is lower due to the higher effective mass and due to intervalley scattering. No measurements of the pressure dependence of the electron mobility in the X valleys have yet been made, but by comparison with GaP it can be expected to increase with pressure at approximately 2.5%/GPa. It should be noted that the DX centre associated with donors in GaAs will normally trap out the electrons before the X-minima become lowest and the material will become semi-insulating. REFERENCES [1] [2] [3] [4] [5] [6]

D. Lancefield, A.R. Adams, BJ. Gunney [Appl Phys. Lett. (USA) vol.45, no.10 (1984) p.1121 ] D. Lancefield, A.R. Adams, M.A. Fisher [ J. Appl Phys. (USA) vol.62 no.6 (1987) p.2342 ] I.Y. Yanchev, S.K. Evtimova [ J. Phys. C (UK) vol. 18 no. 14 (1985) p.L377-82 ] CG. Crookes, D. Lancefield, J.M. Baud, A.R. Adams [ High Pressure Res. vol.3 (1990) p.37 ] G.D. Pitt, J. Lees [ Phys. Rev. B (USA) vol.2 (1970) p.4144 ] D. Patel, T.E. Grambaker, J.R. Sites, LL. Spain [ Rev. Sci. Instrum. (USA) vol.57 (1986) p.2795 ]

2.8

Electron mobility enhancement in GaAs heterostructures Previous Page

JJ. Harris May 1996

A

INTRODUCTION

The observation of mobility enhancement in a modulation-doped heterojunction structure was first reported by Dingle et al [1], for a GaAs/AlGaAs multiple quantum well layer in which the n-type dopant was located in the wide bandgap AlGaAs regions, while the free electrons were confined as a two-dimensional electron gas (2DEG) in the GaAs wells. This spatial separation reduced the strength of the Coulomb interaction between electrons and donors, giving a lower ionised impurity scattering rate by these 'remote' impurities, and hence increased mobility. Mimura et al [2] demonstrated that a similar enhancement could be obtained in a single AlGaAs/GaAs modulation-doped heterojunction, and this effect has subsequently been exploited by many researchers to produce 2DEG samples with extremely high electron mobilities, used for studies of, for example, the fractional quantum Hall effect, Wigner crystallisation and ballistic motion, as well as in the development of high electron mobility transistors (HEMTs, also known as MODFETs or TEGFETs). B

LOW TEMPERATURE MOBILITY MEASUREMENTS

The enhancement in carrier mobility is most readily observed at low temperatures, where ionised impurity scattering is generally dominant and the additional effects of optical phonon scattering are reduced (at 77 K) or can be neglected (at 4 K). TABLE 1 shows the gradual improvement in low temperature mobility in MBE-grown AlGaAs/GaAs 2DEGs over recent years. TABLE 1. Low temperature mobility in MBE-grown AlGaAs/GaAs, 2 DEGs. Reference

77 K mobility li77K (IQ4 cm2/ Vs)

Mimura etal (1980) [2]

^25

Hiyamizuetal(1983)[3]

19^

L25

3X)

Heiblumetal(1984)[4]

20

UO

2X)

3A0

2X)

^30

L6

^50

L5

Harris etal (1987) [5] English etal(1987)[6]

27

Foxonetal(1989)[7] Pfeifferetal(1989) [8T

|

4 K mobility ii4K (IQ6 cm2/ Vs)

Sheet carrier density, ^(10 11 Cm" 2 ) 7X)

6JW

|

24

Below - 50 K in these degenerate 2DEGs, the ionised impurity scattering rate can be considered as temperature independent. However, when this rate is low, as in the more recent samples [5-8], further improvements in mobility have been seen on cooling below 4 K (e.g. Pfeiffer et al [8] achieved 1.17 x 107 cm2/ Vs at 0.35 K), showing that acoustic phonon scattering is also playing a part in limiting the low temperature mobility in these layers. C

FACTORS CONTROLLING MOBILITY ENHANCEMENT

The wide range of 2DEG mobility values reported in the literature is the result of a complex interplay of several important factors, principally differences in the design of the 2DEG structures, improvements in growth techniques, and details of the measurement conditions used. These various aspects will be considered below. Cl

Dependence on Design Parameters

2DEG density ( c m - 2 )

M o b i l i t y ( c m 2 V- 1 s ~ 1 )

Theoretical studies of transport in modulation-doped heterojunctions [9, 10] soon established that the major structural factors influencing the low temperature mobility were (a) the dopant density, (b) the thickness W^ of the undoped AlGaAs spacer layer between the doped AlGaAs region and the GaAs layer (this further reduces the strength of the Coulomb scattering), and (c) the carrier density in the 2DEG. The latter determines the velocity, vF, of the carriers at the Fermi energy, and hence their scattering rate (it can be shown that, for a degenerate 2DEG, one expects \i «

Spacer layer thickness (A)

FIGURE 1. Variation of carrier density (",n) and mobility (•,<>) of 2DEG samples at 4 K, as a function of the undoped spacer thickness, W^. The results are for samples which have been illuminated to saturate the persistent photoconduction effect. Two thicknesses of doped AlGaAs are shown, 400 A (closed symbols) and 500 A (open symbols). (From [5].)

vF3 « n?12 [9]). Experimental verifications of these effects are shown in FIGURES 1 and 2, illustrating the dependence of mobility on the thickness of the undoped spacer layer, Wsp (in a set of otherwise identical samples), and the electron density in three gated 2DEGs with different spacers, respectively. In FIGURE 1, the mobility initially rises with increasing Wsp, due to the reduction in remote impurity scattering, while the fall in mobility at high spacer thickness is the result of scattering by background impurities in the GaAs becoming dominant. Although the density of these impurities is independent of W8^ their scattering rate increases as the carrier density falls, due to the reduced Fermi velocity. Also evident from FIGURE 1 is the effect of the thickness of the doped AlGaAs layer: by increasing this to 500 A, the space charge due to ionised donors in the surface depletion region is moved further from the 2DEG, thus reducing their scattering effect [12]. Very high mobility structures have used much thicker, lightly doped AlGaAs to capitalise on this effect [6]. A variation on this idea is now frequently used, in which, instead of uniform doping, a delta-doped plane is used to supply carriers to the 2DEG (it has been claimed that this produces further mobility enhancement over uniformly doped structures [13]), and a second delta- or uniformlydoped region nearer the surface is used to supply the necessary surface charge [8]. The strong dependence of mobility on electron density at fixed doping level is confirmed experimentally in FIGURE 2. By using front- and back-surface gates, variations of the form |u« n j , with Y - 1 - 1.7 (dependent on spacer thickness), have been observed [H]. C2

Growth-related considerations

C2.1

Interface roughness

The surface of epitaxial AlGaAs is almost always rougher than that of GaAs. For this reason, the usual order of deposition is AlGaAs on GaAs, giving a fairly smooth 'normal' interface, and all of the results quoted above have come from structures of this type. Growth of the so-called 'inverted' structure, GaAs on AlGaAs, generally results in a lower mobility, due to interface roughness scattering, together with the possibility of scattering due to dopant which has surfacesegregated into the GaAs region from the underlying AlGaAs. Nevertheless, with care, good results have also been achieved for such structures, e.g. 4.6 x 105 cm2/ Vs at 4.2 K for a sheet density of 2 x 1011 cm"2 [14]. C2.2

Background impurities

The peak in the mobility vs. spacer curve, FIGURE 1, will increase and move to greater spacer thickness if the background scattering rate falls (and vice versa). Hence the greatest mobility enhancements have been achieved in structures with a very low background impurity density (<~1014 cm"3) in the GaAs layer. Such improvements are obtained by paying particular attention to source material purity and growth system cleanliness, and using such features as an additional AlGaAs/GaAs superlattice buffer layer to trap impurities migrating from the substrate [7]. In

/x n - /x (cm 2 /V s) (ND/l0l8cm"3) NORMALIZED MOBILITY

FIGURE 2. Normalised mobility Jin= ^N 0 plotted as a function of carrier density Nsfor three gated Hall bar samples with W^ of 0, 4.5 and 18 nm. Dashed lines are calculated values. (From [11].)

general, MBE appears to be more successful in reducing the level of unintentional doping than MOVPE3 where reported low temperature mobilities of 4.0 * 105 cm2/ Vs at 5 K [15] may be regarded as typical of results obtained using arsine. However, very recent work using tertiarybutylarsine has demonstrated much improved mobilities of 2 x 106 cm2/ Vs at 0.3 K [16]. C3

Effect of Measurement Conditions

C3.1

Temperature

The increasing effect of phonon scattering as the temperature is raised tends to mask the structural and impurity differences between samples, and mobility values close to those occurring in high purity GaAs are obtained. At room temperature, for example, MBE-grown 2DEG samples have given 9200 cm2/ Vs [4] and MOVPE-grown 2DEG layers have given 8100 cm 2I Vs [15]. C3.2

Illumination

Cooling a 2DEG sample to low temperatures in the dark results in a reduction in carrier density, due to freezing out of electrons onto deep DX centres in the AlGaAs. Below a critical temperature of-150 K, subsequent illumination can persistently release these carriers back into the 2DEG, with a resulting increase in mobility due to the increased Fermi velocity, and the highest reported values are generally obtained under these conditions. However, it has been

pointed out, e.g. by Buks et al [17], that if the same carrier density can be achieved in the dark, (in their case, this was achieved by cooling the sample with an applied gate voltage), the mobility can be up to a factor of 6 times greater. This is believed to be due to a correlation of the positively-charged shallow donor states with negatively-charged occupied DX centres, reducing their effectiveness as scatterers. Such a correlation occurs during cooling in the dark, but is destroyed by subsequent illumination. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] [II] [12] [13] [14] [15] [16] [17]

R. Dingle, H.L. Stormer, A.C. Gossard, W. Weigmann [Appl. Phys. Lett. (USA) vol.33 (1978) p.665-7 ] T. Mimura, S. Hiyamizu, T. Fujii, K. Nanbu [ Jpn. J. Appl. Phys. pt.2 (Japan) vol.19 (1980) p.L225-7 ] S. Hiyamizu, J. Saito, K. Nanbu, T. Ishikawa [ Jpn. J. Appl. Phys. pt.2 (Japan) vol.22 (1983) p.L609-ll] M. Heiblum, E.E. Mendez, F. Stern [Appl. Phys. Lett. (USA) vol.44 (1984) p.1064-6 ] JJ. Harris, CT. Foxon, K.WJ. Barnham, D.E. Lacklison, J. Hewett, C. White [ J. Appl. Phys. (USA) vol.50 (1987) p. 1826-31 ] J.H. English, A.C. Gossard, H.L. Stormer, K.W. Baldwin [ Appl. Phys. Lett. (USA) vol.50 (1987) p. 1826-8] CT. Foxon, JJ. Harris, D. Hilton, J. Hewett, C Roberts [ Semicond. Sd. Technol. (UK) vol.4 (1989)p.582-5] L. Pfeiffer, K.W. West, H.L. Stormer, K.W. Baldwin [Appl. Phys. Lett. (USA) vol.55 (1989) p.88890; Mater. Res. Soc. Symp. Proc. (USA) vol.145 (1989) p.3-12 ] K. Hess [ Appl. Phys. Lett. (USA) vol.35 (1979) p.484 ] K. Lee, M.S. Shur, TJ. Drummond, H. Morkoc [ J. Appl. Phys. (USA) vol.54 (1983) p.6432 ] K. Hirakawa, H. Sakaki [ Phys. Rev. B (USA) vol.33 (1986) p.8291-303 ] Y. Takeda, H. Kamei, A. Sasaki [ Electron. Lett. (UK) vol. 18 (1982) p.310 ] E.F. Schubert, L. Pfeiffer, K.W. West, A. Izabelle [ Appl. Phys. Lett. (USA) vol.54 (1989) p. 1350 ] H. Shtrikman, M. Heiblum, K. Seo, D.E. Galbi, L. Osterling [ J. Vac. Sci. Technol. B (USA) vol.6 (1988) p.670-3 ] Y. Mori, F. Nakamura, N. Watanabe [J. Appl. Phys. (USA) vol.60 (1986) p.334-7 ] H.C Chui, B.E. Hammons, N.E.Harff, J.A. Simmons, M.E. Sherwin [ Appl. Phys. Lett. (USA) vol.68 (1996) p.208-10] E. Buks, M. Heiblum, Y. Levinson, H. Shtrikman [ Semicond. Sci. Technol. (UK) vol.9 (1994) p.2031-41]

2.9

Ballistic transport and velocity overshoot in GaAs JJ. Harris May 1996

A

INTRODUCTION

Electrons travelling on a length-scale which is short compared to their mean free path can exhibit collisionless, i.e. ballistic, acceleration. Under such high electric fields they can achieve velocities which are greater than the scattering-limited values, so that velocity overshoot occurs. These effects are transient, operating on a timescale less than or comparable to the energy and momentum relaxation times, < ~ 1 ps. Beyond this time, the equilibrium scattering rates are approached, and the velocity falls back to its scattering-limited value [I]. In addition to the physics interest inherent in these effects, their potential for improved high-speed performance in ultra-small devices has also been recognised [2]. Since the occurrence of these phenomena in semiconductors was first suggested by Ruch [3] on the basis of Monte Carlo simulations, much theoretical and experimental work has been carried out to investigate the effects. From analytical considerations, Hess [4] estimated the limiting velocities in GaAs to be (i) (ii) (iii)

scattering-limited, 3 x 107 cm/s, overshoot, 5 x io 7 cm/s, ballistic, 108 cm/s.

Other Monte Carlo simulations [5-7] have shown that these values will depend on several parameters (doping level, temperature, electric field, etc.), and this is illustrated in FIGURE 1 for GaAs doped at 1017 cm'3 under various applied fields. The experimental techniques which have

FIGURE 1. Monte Carlo simulation of instantaneous velocity vs. distance in GaAs doped at 1017 cm"3 at 300 K as a function of applied electric field. (From [5].)

been applied to the study of these effects fall into three main categories, namely acceleration of cold electrons under high fields (these carriers may either be extrinsic or photo-generated), collection of high energy ('hot') electrons injected across a low-field region, and transport at low temperatures. The measurements have either been made on 'vertical' structures, i.e. across thin epitaxial layers, or, more recently, on lateral structures, many using high mobility GaAs/AlGaAs two-dimensional electron gas (2DEG) samples, where ballistic effects have been observed over a distance of ~ 100 nm [8]. B

VERTICAL TRANSPORT

TABLE 1 lists electron velocities inferred from measurements of majority carrier transport in bipolar structures. TABLE 1. Electron velocities. Reference

Structure

Velocity (107cm/s)

Adachietal(1982)[9]

submicron GaAs ri7p7n+ diodes

>_1

Ishibashietal(1988)[10]

HBT with i/p7n+ collector

3-5

Bertholdetal(1989)[ll]

|

HBT with resonant tunelling collector

|

12

A number of studies have also been made on unipolar structures, principally hot electron transistor samples, where an emitter barrier, either heterojunction [12] or planar-doped [13], is used to inject hot electrons into a thin base region. The energy spectrum of the transmitted carriers is analysed by varying the bias across an adjacent collector barrier. It has been found that a fraction of the injected electrons can cross the base ballistically, without energy loss; for a 300 A base this can be as high as 75% [13]. By studing magneto-transport in such a device, Imamura et al [14] estimated the velocity in the base to be 108 cm/s. Evidence of ballistic motion in wide (-1000 A) resonant tunnelling diodes has been obtained from low-temperature transport measurements with a magnetic field perpendicular to the direction of the tunnelling current [15]. Oscillations in the tunnel current have been interpreted as electron interference arising from a combination of multiple reflections within the quantum well region and deflection of electrons by the field into 'skipping orbits' along the well. Since the electron coherence would be destroyed by the presence of scattering, the mean free path must be several times the thickness of the well. If the angle between the magnetic field and the well is varied, analysis of the resultant electron motion in terms of unstable periodic orbits shows that these structures exhibit classically chaotic behaviour [16]. C

HORIZONTAL TRANSPORT

Direct evidence for the dependence of electron velocity on distance has been obtained by Ryan et al [17], who studied FETs with a range of ultra-short source-to-dr&in separations, and deduced the velocity versus distance data shown in FIGURE 2. Although the electric fields were not specified in this report, a clear similarity between these results and the modelling in FIGURE 1 is apparent.

Transit velocity (10 cm/scc)

Gate Length (nm) FIGURE 2. Gate-length dependence of the transit velocity determined from DC measurements on GaAs MESFETs doped at 2 x 1017 cm"3. (From [17].)

Supporting evidence for this type of behaviour has been obtained from transient measurements on photo-generated carriers injected into undoped GaAs under a high applied field, either in a lateral p-i-n structure [18], or in a 10 |um photoconductive gap in a stripline circuit [19]. Here, too, velocity overshoot effects on a sub-picosecond timescale have been observed, in good agreement with simulations. The advent of high mobility 2DEG structures, where the low scattering rates at low temperature result in extremely long mean free paths, has allowed a number of elegant experiments on ballistic transport in lateral structures to be performed. Of these, only a small selection can be mentioned here. Sivan et al [20] have constructed a lateral version of the hot electron transistor from two split electrostatic gates in series, and shown that, provided the injection energy is less than the longitudinal optical phonon energy (36 meV), ballistic transport can occur over distances of more than 2 mm. Observation of quantised steps in the low-temperature (30 mK) conductivity of a quasi-ID 'quantum point contact' fabricated on a 2DEG also requires collisionless transport [21,22], as do deflection experiments in a magnetic field. The latter have shown that a beam of electrons emitted from a quantum point contact into a 2DEG can be deflected through 180° and focussed into a collector [23], in one case after travelling ballistically for distances of-100 mm [8]. Other manifestations of ballistic motion in 2DEG samples include the observation of electron interference in ring structures (the Aharanov-Bohm effect) [24], and anomalous Hall effect behaviour in intersecting mesoscopic channels, where specular reflection of the electrons from the sidewall potentials can produce a range of geometry-dependent phenomena [25].

REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] [II] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25]

For a detailed discussion, see: K. Hess, GJ. Iafrate [ Proc. IEEE (USA) vol 76 (1988) p.519-31 ] M.S. Shur, L.F. Eastman [ IEEE Trans. Electron Devices (USA) vol.ED-26 (1977) p. 1677-83 ] J.G. Ruch [ IEEE Trans. Electron Devices (USA) vol.ED-19 (1972) p.652-4 ] K. Hess [ IEEE Trans. Electron Devices (USA) vol.ED-28 (1981) p.937-40 ] TJ. Maloney, J. Frey [ J. Appl. Phys. (USA) vol.48 (1977) p.781-7 ] G.M. Wysin, D.L. Smith, A. Redondo [ Phys. Rev. B (USA) vol.38 (1988 p. 12514-24 ] G.M. Dunn, A.B. Walker, J.H. Jefferson, D.C. Herbert, [ Semicond. Sci. Technol. (UK) vol.9 (1994) p.2123-9] H. Spector, H.L. Stormer, K.W. Baldwin, L.N. Pfeiffer, K.W. West [ Surf. Sci. (Netherlands) vol.228 (1990) p.283-5 ] S. Adachi, M. Kawashima, K. Kumabe, K. Yokoyama, M. Tomizawa [ IEEE Electron Device Lett. (USA) vol.EDL-3 (1982) p.409-11 ] T. Ishibashi [ Ext. Abstr. 20th Int. Conf. Solid State Devices and Materials (Bus. Centre for Acad. Soc. Japan, Tokyo, 1988) p 515-8 ] K. Berthold, A.F.J. Levi, J. Walker, RJ. Malik [Appl. Phys. Lett. (USA) vol.54 (1989) p.813-5 ] M. Heiblum, LM. Anderson, CM. Knoedler [Appl. Phys. Lett. (USA) vol.49 (1986) p.207-9 ] AFJ. Levi, J.R.Hayes, P.M. Platzmann, W. Weigmann [Phys. Rev. Lett. (USA) vol.55 (1985) p.2071-3 ] K. Imamura et al [ Surf. Sci. (Netherlands) vol.174 (1986) p.481-6 ] L. Eaves, M.L. Leadbeater, E.S. Alves, F.W. Sheard, G.A. Toombs [in Physics and Technology ofHeterojunction Devices, Ed. D.V. Morgan, RH. Williams, (Peter Peregrinus UK, 1991) p.33-52 ] T.M. Fromhold, L. Eaves, F.W. Sheard, M.L. Leadbetter, TJ. Foster, P. Main [ Phys. Rev. Lett. (USA) vol.72 (1994) p.2608-11 ] J.M. Ryan, J. Han, A.M. Kriman, D.K. Ferry, P. Newman [ Solid State Electr. (UK) vol.32 (1989) p. 1609-13] W. Sha, J.K. Rhee, T.B. Norris, WJ. Schaff [ IEEE J. Quantum. Electron. (USA) vol.28 (1992) p. 1445-55] K. Meyer, M. Pessot, G. Morou, R Grondin, S. Chamoun [Appl. Phys. Lett. (USA) vol.53 (1988) p.2254-6 ] U. Sivan, M. Heiblum, CP. Umbach [ Phys. Rev. Lett. (USA) vol.63 (1989) p.922-5 ] D.A. Wharam et al [J. Phys. C (UK) vol.21 (1988) p.L209-14 ] BJ. van Wees et al [ Phys. Rev. B (USA) vol.38 (1988) p.3625-7 ] H. van Houten et al [ Phys. Rev. B (USA) vol.39 (1989) p.8556-75 ] C. J.B. Ford et al [ J. Phys. C (UK) vol.21 (1988) p.L325 - L331 ] C.J.B. Ford, S. Washburn, M. Buttiker, CM. Knoedler, J.M. Hong [ Surf. Sci. (Netherlands) vol.229 (1990) p.298-302]

2.10 Modulation doping J. Ballingall March 1996

A

INTRODUCTION

The modulation doping concept spatially separates the dopant ions from the ionized electrons or holes, via the potential energy step at a heterojunction, thereby reducing the ionized impurity scattering rate of the free carriers from their host ions. For n-type doping, a high mobility 2-dimensional electron gas (2DEG) is formed at the heterojunction interface, in the material with the lower bandgap. The first reported demonstration was by ATT Bell Laboratories researchers on an AlGaAs/GaAs superlattice grown by molecular beam epitaxy (MBE), with silicon impurity doping of the AlGaAs layers [I]. Measurements revealed a higher electron mobility for the GaAs than would be expected for bulk GaAs doped to a comparable electron concentration. Also, the rise in mobility with decreasing temperature was much greater than in uniformly doped material, in accord with the temperature dependence of phonon scattering. Further evaluations by the Shubnikov-de Haas technique confirmed expectations that a 2DEG was formed on the GaAs side of the heterojunction [2]. Similarly, a 2-dimensional hole gas (2DHG) was shown to be formed in the GaAs when p-type doping was used in the AlGaAs [3]. B

TRANSISTOR APPLICATIONS

The potential advantages for GaAs-based transistors, which up until then had relied on heavily doped active layers, did not go unnoticed by several groups around the world. The first modulation doped transistors were demonstrated in 1980 by Fujitsu researchers, who dubbed their result the high electron mobility transistor (HEMT) [4]. At about the same time, H. Morko? at the University of Illinois reported a similar device [5] which would later be known as the modulation doped FET [6], or simply MODFET. Both names have persisted and continue to be used interchangeably today, as the device has become the microwave low noise transistor of choice for many commercial applications such as Direct Broadcast Satellite (DBS) TV. An extensive review of modulation doped transistors is the content of Datareview 20.3 in this book. C

STRUCTURAL OPTIMIZATION

New scientific discoveries were made and transistor performance improved as engineers and scientists began to optimize the modulation doping profile for various interests. The effects of the AlAs mole fraction were studied [7,8]. As the mole fraction is increased, the conduction band discontinuity also increases, leading to a higher transferred electron concentration. Also, the measured Hall mobility may appear higher due to less parallel conduction in the AlGaAs. Probably the most significant innovation was the addition of an undoped AlGaAs 'spacer', between the GaAs and the doped AlGaAs [9,10]. Up until this development, the doped AlGaAs and the 2DEG were two contiguous regions. The spacer layer spatially separated the two regions by several atomic layers, thereby reducing the coulomb attraction even further, with a 15 nm spacer boosting the electron mobilities at 4.2 K by nearly an order of magnitude while halving the electron density [10].

The 1980s saw a divergence in the structural optimization, as low temperature physicists utilized spacers up to 100 nm or more to maximize mobility for discoveries such as the fractional quantum Hall effect [11], while device physicists narrowed in on a thickness of about 3 nm as a compromise between achieving higher mobility while maintaining adequate electron density for transistor current requirements [12]. Both interests were served by improvements in MBE techniques, as the 2DEG mobilities increased [13-15] due to a reduction to 2 x 1013 cm'3 in the unintentional 'background' impurities introduced into the AlGaAs and GaAs by the growth process [15]. Electron mobility increased from just under 1 x 105 cm2/ Vs in 1981 [10], to over 1 x 107 cm2/ Vs in 1989 [15], distinguishing modulation-doped GaAs as the highest mobility semiconductor at low temperature. Theoretical analyses of the temperature dependence of electron mobility also progressed from, for example, fairly simple concepts involving only coulomb scattering by the remote impurities and polar optical phonons [16], to include acoustic mode phonons [17], alloy disorder and band-edge discontinuity fluctuations [18], and interface roughness scattering [19]. A topic which received intense attention on the part of both experimentalists and theorists was the so-called 'inverted interface problem'. This was the observation that the electron mobility was always inferior if the modulation doping was supplied by AlGaAs which was grown before the GaAs and hence underlaid the GaAs, as opposed to AlGaAs which was grown after and so overlaid the GaAs. The latter was the more conventional approach for devices, and so became known as the 'normal' interface configuration, leaving the other to be known as the 'inverted' interface [20]. This problem was solved by linking the mobility degradation with the surface roughening and impurity uptake of AlGaAs as it is grown by MBE [20-24]. Surface accumulation of silicon dopant atoms on AlGaAs during growth, and diffusion during growth of silicon dopant ions from the underlying AlGaAs into the GaAs, was also a factor [24,25]. Delta doping, also known as atomic planar doping, was a structural innovation introduced in the mid-1980s, mainly for the benefit of devices [26-28], and, like the undoped AlGaAs spacer, became a common design feature of HEMTs [29]. This replaced the uniform silicon dopant concentration of the AlGaAs supply layer by a two-dimensional sheet of silicon donor atoms, boosting the useful 2DEG density relative to what could be achieved with uniform doping. D

CONCLUSION

The modulation doping concept is one of the great examples in recent history of how a simple idea can spawn a new field of scientific research, and at the same time revolutionize an industrial technology. Fruition of the idea and the intensity of its impact on science and technology was dependent on the development and maturity of MBE technology. REFERENCES [1] [2] [3] [4] [5]

R. Dingle, H. Stormer, A.C. Gossard, W. Wiegmann [ Appl. Phys. Lett. (USA) vol. 33 (1978) p.665-7 ] H.L. Stormer, R. Dingle, A.C. Gossard, W. Wiegmann, M.D. Sturge [ Solid State Commun. (USA) vol.29 (1979) p.705-9] H.L. Stormer, W.T. Tsang [J Vac. Set Technol. (USA) vol.17 (1980) p. 1130-1 ] T. Mimura, S. Hiyamizu. T. Fuiii, K. Nanbu [ Jpn. J Appl. Phys. (Japan) vol. 19 (1980) p.L225-7 ] H. Morkos [IEEEElectron Device Lett. (USA) vol.EDL-2 (1981) p.260 ]

[6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26]

[27] [28] [29]

S.L. Su et al [ Electron. Lett. (UK) ] vol. 18 (1982) p.794-6 ] H.L. Stormer, A.C. Gossard, W. Wiegmann, K. Baldwin [Appl. Phys. Lett. (USA) vol.39 (1981) p.912-4 ] TJ. Drummond, W. Kopp, R. Fischer, H. Morkoc [ J. Appl. Phys. (USA) vol.53 (1982) p. 1028-9 ] L.C. Witkowski, TJ. Drummond, CM. Stanchak, H. Morkoc [ Appl. Phys. Lett. (USA) vol.37 (1980) p. 1033-5] H.L. Stormer, A. Pinczuk, A.C. Gossard, W. Wiegmann [Appl. Phys. Lett. (USA) vol.38 ( 1981) p.691-3 ] D. Tsui, H. Stormer, A. Gossard [ Phys. Rev. Lett. (USA) vol.48 (1982) p. 1559 ] K. Lee et al [ J. Vac. Sci. Technol. B (USA) vol. 1 (1983) p.186-9 ] J.H. English, A.C. Gossard, H.L. Stormer, K.W. Baldwin [ Appl. Phys. Lett. (USA) vol.50 (1987) p.1826-8] M. Shayegan, VJ. Goldman, C. Jiang, T. Sajoto, M. Santos [ Appl. Phys. Lett. (USA) vol.52 (1988) p.1086-8] L. Pfeiffer, K.W. West, H. L. Stormer, K.W. Baldwin [ Appl. Phys. Lett. (USA) vol.55 (1989) p.1888-90] TJ. Drummond, H. Morkoc, K. Hess, A.Y. Cho [ J. Appl. Phys. (USA) vol.52 (1981) p.5231-4 ] E.E. Mendez, PJ. Price, M. Heiblum [Appl. Phys. Lett. (USA) vol.45 (1984) p.294-6 ] N.M. Cho, S.B. Ogale, A. Madhukar [ Appl. Phys. Lett. (USA) vol. 51 (1987) p. 1016-18 ] H. Sakaki et al [ Appl.Phys.Lett. (USA) vol. 51 (1987) p. 1934-6 ] W.T. Masselink et al [ Appl. Phys. Lett. (USA) vol.44 (1984) p.435-7 ] T. Achtnich, G. Burri, M.A. Py, M. Ilegems [ Appl. Phys. Lett. (USA) vol.50 (1987) p. 1730-2 ] U. Meirav, M. Heiblum, F. Stern [ Appl. Phys. Lett. (USA) vol.52 (1988) p.1268-70 ] D. Arnold et al [ Appl. Phys. Lett. (USA) vol.45 (1984) p.902-4 ] M. Heiblum [ J. Vac. Sci. Technol. B (USA) vol.3 (1985) p.820-2 ] K. Inoue, H. Sakaki, J. Yoshino, Y. Yoshioka [ Appl. Phys. Lett. (USA) vol.46 (1985) p.973-5] YK. Chen, D.C. Radulescu, PJ. Tasker, G.W. Wang and L.F. Eastman [ Proc. 13th Int. Symp. on Gallium Arsenide and Related Compounds, Las Vegas, NV, USA, 28 Sept.-l Oct. 1986 (IOP, Bristol, UK, 1987) p.581-6] T. Ishikawa, K. Ogasawara, T. Nakmura, S. Kuroda, K. Kondo [J. Appl. Phys. (USA) vol. 61 (1987) p. 1937-40] E.F. Schubert, J.E. Cunningham, W.T. Tsang, G.L.Timp [Appl. Phys. Lett. (USA) vol.51 (1987) p. 1170-2] J.M. Ballingall et al [ Thin Solid Films (Switzerland) vol.231 (1993) p.95-106 ]

2.11 Carrier concentrations following 5-doping R.L. Williams July 1995

A

INTRODUCTION

The investigation of planar doping structures in GaAs5 where dopant atoms are deposited on a single atomic plane of the host lattice, was begun in 1980 by Wood and co-workers [I] 3 in a series of experiments designed to overcome the problems associated with the thermal inertia of dopant effusion cells in molecular beam epitaxy (MBE). To generate rapidly varying dopant profiles in MBE it is necessary to produce correspondingly rapid variations in the temperature of the dopant effusion cell, a requirement which may not be compatible with the thermal inertia of the effusion cell if the required temperature variations are too rapid. One answer to this problem is to reduce the crystal growth rate to such an extent that the required dopant variation is rapid in the spatial domain but not in the time domain. Taking this concept to its ultimate limit, and using germanium as the n-type dopant, Wood et al were able to grow MBE GaAs with doping profiles of arbitrary complexity, by fabricating the doping profile from a series of appropriately doped and appropriately spaced single atomic planes. The doping of these single atomic planes was performed in the absence of an incident gallium flux and consequently in the absence of any layer by layer growth. Following on from these initial experiments, in which atomic plane or 6-doping was treated purely as an experimental convenience, the electronic and structural properties of narrow doping structures in GaAs have attracted considerable attention, whilst device applications have included the non-alloyed Ohmic contact [2], the atomic plane doped field effect transistor [3] and the doping superlattice laser [4]. Early interest was stimulated by the work of Schubert et al [5], who reported classical Hall experiments on silicon 8-doped GaAs, which seemed to suggest a room temperature mobility enhancement over equivalent, uniformly doped material that increased with increasing carrier density, reaching a factor of approximately four at a free carrier density of N2D= 1.1 x 1013 cm"2. The maximum carrier density reported by these authors was 2 x 1013 cm"2 which was interpreted as being equivalent to a uniform density of 9 x 1019 cm"3 (calculated as N3D = N201'5). To account for this high level of n-type silicon incorporation in 5-doped samples it was argued that the lack of group III flux during doping would suppress the number of unsatisfied arsenic sites relative to conventional doping procedures and so lead to an increase in the level of silicon donor activity. However, this argument pre-supposes that the predominant compensating species in highly silicon-6-doped GaAs is silicon on an arsenic site, as reported by Spitzer and co-workers for uniformly silicon-doped material [6]. More recent work by Maguire et al [7] has suggested that the most important compensating species is a complex of silicon with a native defect, Si-X, where X has been assigned as a gallium vacancy. If the compensation mechanism requires a gallium vacancy for its operation then a high arsenic to gallium flux ratio may not be advantageous. In subsequent work it has been shown that the premise of single plane doping upon which the calculation of effective three-dimensional doping levels is based in [5] is unlikely to be valid at high two-dimensional carrier densities or high growth temperatures. In consequence, the

mobilities quoted in [5] should be compared with three-dimensional doping levels below that predicted on a model of true single plane doping. To date, no conclusive evidence has been presented for an enhancement of the room temperature mobility or for an increased compensation threshold in 5-doped GaAs. At doping levels above the metal-insulator transition (approximately 3 x 1011 cm'2 for Si 5-doping of GaAs, for example), electronic conduction in the 8-doping layer must be considered in terms of the electronic subbands that describe quantised motion along the growth direction, together with a free carrier dispersion in the 5-doped plane [8]. As such classical, low magnetic field Hall effect experiments produce values for the diagonal conductivity, Gx^ and the carrier density, nHall, that are only averages over the individual subband carrier densities, IT1, and mobilities, \i{: a

xx(B)

I^ ~r~ i (1 + n?B2)

n

Hall = ( E ^ i ) 2 / E 1W i

i

To obtain the individual subband carrier densities themselves and thus a true measure of the total free carrier density, a number of authors have studied the high magnetic field Shubnikov-de Haas effect [9,10]. Early measurements have shown subband occupancies in silicon-5-doped samples that vary approximately linearly with total doping concentration, until saturation of the free electron concentration at a level that varies with the sample growth temperature [H]. This observation of free carrier saturation has stimulated a large interest and has led to a greater understanding of the possible saturation mechanisms in silicon-8-doped samples and to a study of 5-doping using other dopants. B

CARRIERDENSITIES

Bl

Silicon

The most extensively studied dopant for 6-doping applications has been silicon. Using silicon as the n-type dopant, near ideal 8-doping layers with widths less than 20 A have been realised [10,12,13]. To achieve these narrow doping widths, a number of techniques including secondary ion mass spectroscopy (SIMS) [14], capacitance voltage (CV) - profiling [15-17], high resolution X-ray diffraction (HRXD) [13] and Shubnikov-de Haas (SdH) [10,12] have indicated the need for low (<550°C) growth temperatures. At higher temperatures a number of mechanisms including surface segregation and concentration dependent diffusion [14] operate to broaden the 5-doping profile and so reduce the effective three-dimensional carrier density. At silicon doping levels in excess of approximately Ix 1013 cm"2, the free carrier density has been observed to saturate [11] or even diminish [18] with increasing doping. Two distinct mechanisms have been proposed for this saturation and, depending upon the growth conditions for the sample, either or both have been seen to operate:

(i)

Electronic Saturation

For samples grown at low temperatures and doped to levels in excess of approximately Ix 1013 cm"2 (the level required for an ideal 5-doping layer has been estimated as 5 x 1012 cm"2), an electronic saturation mechanism has been proposed [9]. An electronic mechanism implies a pinning of the Fermi energy by a resonant defect level associated with a higher order conduction band minimum or by occupation of the higher order minimum itself. Shubnikov-de Haas measurements under hydrostatic pressure [9] have been used to investigate this saturation mechanism, by producing a reduction in the energetic separation of the F- and Lconduction band minima. 8-doped samples grown at sufficiently low temperatures and doped to a level below that at which saturation is observed for this growth temperature show a dramatic loss of free carriers above a threshold value of the applied hydrostatic pressure. Accompanying this loss of free carriers is a large preferential mobility increase for carriers occupying the lowest energy, i = 0 electronic subband. These results are interpreted as reflecting the trapping of free carriers onto a resonant defect level as the hydrostatic pressure is increased. The number of carriers trapped in this manner (« 50% in some cases) and the location of the defect level at the position of the doping plane argue strongly for a silicon related defect. The preferential mobility increase observed for the i = 0 carriers in these experiments has been explained as resulting from preferential neutralisation of the silicon donors at the 5-doped well centre, at which position the localised defect level first penetrates the Fermi sea. At the 8-doped well centre the i = 0 subband wavefunction has a maximum intensity, implying a strong dependence for the i = 0 mobility on the ionised impurity concentration at this position. This observed free carrier loss and associated mobility increase is completely analogous to the pressure induced mobility increase observed by Maude and co-workers in uniformly silicon-doped GaAs [19]. (ii)

Structural Saturation

Samples grown at higher growth temperatures, which are not expected to have the high Fermi energies required for electronic saturation, have also exhibited saturation behaviour. Saturation in these samples has been studied extensively using SIMS [14,17], CV-profiling [15-17], local mode vibrational spectroscopy (LVM) [17,18,20] and HRXD [13]. Beall and co-workers [17] have studied a sample consisting of three 8-doped planes, of doping levels 0.4, 1 and 4 x 1013 cm"2, which has been grown at a low temperature of 520 0 C and then been subjected to post growth, high temperature annealing. The as grown 8-layers showed fairly narrow doping profiles (< 70 A), with some evidence for diffusion of the 8-layer furthest from the sample surface and therefore held at the growth temperature for the greatest length of time. After post growth annealing at 6480 C for 3.5 hours, the spreading of the 8-layers was found to depend very strongly on the initial sheet silicon concentration. A SIMS profile of the lowest doped plane was found to spread into a Gaussian peak with a full width at half maximum (FWHM) of approximately 190 A. The two more highly doped planes were found to have significantly non-Gaussian SIMS profiles, consisting of a central, virtually unbroadened peak at the original doping plane, surrounded by a lower concentration flat region extending for distances of the order of 1000 A, on either side of the original doping plane, depending upon the concentration anneal time and temperature. CV profiling of this same sample showed the flat regions of the SIMS profile to arise from silicon donors, whilst the central peak of the SIMS profile was absent in the CV profile and replaced by a shallow dip, which the authors interpret as indicating a region of

silicon acceptors or electrically inactive silicon. Also observed is an increase in the silicon donor activity as a function of increasing anneal time. To explain these results and the observation of free carrier saturation, the authors suggest that during the deposition of high concentration silicon 6-doping layers, the surface diffusion of silicon atoms is sufficient to allow aggregation of silicon clusters. The silicon in these clusters would be electrically inactive, but once covered with GaAs would act as a reservoir for further electrically active silicon on annealing. The lattice locations of silicon in high concentration silicon 8-doping planes have been studied using LVM [17,18,20] and HRXD [13]. Initial LVM measurements [17,18] showed mainly SiGa lines (i.e. lines from silicon atoms on gallium sites), whose integrated intensity increased after post growth annealing, in agreement with the proposed model of electrically inactive silicon aggregates at the doping plane which would supply electrically active silicon, SiGa, on diffusion away from the original site. Subsequent work in this area [20] has extended the silicon coverages investigated and has shown the presence of Si^, Si03-Si^ pairs, Si-X and covalently bonded Si-Si pairs, in good agreement with the model for silicon aggregation at high surface coverages. B2

Tin

Tin has been used to a lesser degree than silicon to generate spatially localised n-type doping layers in GaAs, primarily because of the problems associated with the dopant spreading away from the original doping plane. In the temperature range 500°C-650°C Harris and co-workers [21] have found significant broadening of deposited tin planes due to surface segregation. Two types of surface segregation behaviour are identified, corresponding to an equilibrium situation above 6000C, in which most of the deposited tin is carried along at the growth surface, and a kinetically limited regime below 6000C, where significant amounts of the dopant are incorporated in the bulk if the dopant is not allowed sufficient time to migrate to the growth surface. At doping concentrations above approximately Ix 1013 cm'2 tin has been shown to form clusters in much the same way as silicon. These tin clusters can again act as sources on further GaAs growth to produce wide, flat doping profiles [22]. B3

Beryllium

For p-type 6-doping, beryllium has been used extensively. Schubert and co-workers [23] have used CV-profiling to show the confinement of the dopant to less than 20 A for a growth temperature of 5000C and a doping concentration of 4 x 1012 cm "2. At higher growth temperatures the beryllium 6-doping planes have been found to broaden significantly [22,23] as a result of both surface segregation and concentration dependent diffusion. In the high doping regime free carrier saturation is not observed [23], with hole concentrations as high as 6.6 x 1014 cm"2 being achieved. In this regime Schubert and co-workers have reported a high degree of dopant migration: approximately 1000 A for a layer doped at 4 x 1014 cm"2 and grown at 500 0 C. Considering the lack of electrical compensation and the good structural quality of the samples these authors have proposed Coulombic repulsion between ionised dopant atoms as the mechanism for the large degree of spreading observed. To test this hypothesis, Harris and co-workers [22] have investigated a series of beryllium- and silicon-5-doped layers in which the two dopants have been deposited simultaneously. Broadening of the beryllium SIMS profiles in these co-deposited samples, after post growth annealing at 6500C for 3.5 hours, is virtually eliminated, adding strong support to the model proposed by Schubert and co-workers [23], since

the ionised silicon dopants would act to screen out the mutual repulsion between ionised beryllium atoms. C

CONCLUSION

Both silicon and beryllium have been shown to produce near ideal 8-doping layers in GaAs for sufficiently low growth temperatures and for sufficiently low sheet doping levels. At elevated growth temperatures or for high doping levels significant movement of the dopant away from the original doping plane can occur as a result of surface segregation, diffusion or Coulombic repulsion effects. For silicon 5-doping layers, electron densities are limited to approximately Ix 1013 cm"2 before electronic or structural saturation effects occur, whilst for beryllium, tightly confined 5-layers (FWHM < 50 A) can be realised at growth temperatures of 5000 C for hole densities up to approximately 2 * 1013 cm'2, at which point significant dopant spreading begins. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] II1] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23]

CEC. Wood, G. Metze, J. Berry, LF. Eastman [J. Appl. Phys. (USA) vol.51 no.l (1980) p.383 ] E.F. Schubert, J.E. Cunningham, W.T. Tsang, T.H. Chiu [Appl. Phys. Lett. (USA) vol.49 no.5 (1986)p.292] E.F. Schubert, K. Ploog [ Jpn. J. Appl. Phys. 2 (Japan) vol.24 no.8 (1985) p.L608 ] E.F. Schubert, A. Fischer, Y. Horikoshi, K. Ploog [Appl. Phys. Lett. (USA) vol.47 ( 1985) p.219 ] E.F. Schubert, J.E. Cunningham, W.T. Tsang [ Solid State Commun. (USA) vol.63 no. 7 (1987) p.591 ] W.G. Spitzer, A. Kahan, L. Bouthille [ J. Appl. Phys. (USA) vol.40 (1969) p.3398 ] J. Maguire, R. Murray, R C Newman, R.B. Beall, JJ. Harris [ Appl. Phys. Lett. (USA) vol.50 no.9 (1987)p.516] A. Zrenner, M. Reisinger, F. Koch, K. Ploog [ Proc. 17th Int. Conf. on the Physics ofSemicond., San Francisco, CA, USA, 6-10 Aug 1984 (Springer Verlag, 1985) p.325 ] A.Zrenner, F. Koch, RL Williams, RA. Stradling, K. Ploog, G. Weimann [ Semicond. Sd. Technol. (UK) vol.3 (1988) p. 1203] M. Santos, T. Sajoto, A. Zrenner, M. Shayegan [ Appl. Phys. Lett. (USA) vol.53 no. 25 (1988) p.2504] A. Zrenner, F. Koch, K. Ploog [ Inst. Phys. Conf. Ser. (UK) no. 91 (1987) p. 171 ] P.M. Koenraad et al [ Semicond. Sci. Technol. (UK) vol.5 (1990) p.861 ] L. Hart, MJ. Ashwin, P.F. Fewster, X. Zhang, M.R. Fahy, R.C. Newman [ Semicond. Sci. Technol. (UK) vol.10 (1995) p.32] R.B. Beall, J.B. Clegg, JJ. Harris [ Semicond. Sci. Technol. (UK) vol.3 no. 6 ( 1988) p.612 ] E.F. Schubert, J.B. Stark, T.H. Chiu, B. Tell [ Appl. Phys. Lett. (USA) vol.53 no. 4 (1988) p.293 ] E.F. Schubert, K. Ploog [ Jpn. J. Appl. Phys. 1 (Japan) vol.25 no. 7 (1986) p.966 ] RB. Beall, J.B. Clegg, J. Castagne, JJ. Harris, R Murray, R.C. Newman [ Semicond. Sci. Technol. (UK) vol.4 (1989) p. 1171] MJ. Ashwin, M. Fahy, JJ. Harris, R.C. Newman, DA. Sansom, R. Addinall [ J. Appl. Phys. (USA) vol.73 ( 1993) p.633 ] D.K. Maude et al [ Phys. Rev. Lett. (USA) vol.59 (1987) p.815 ] RC. Newman et al [ MRS Spring Meeting, San Francisco (March 1995) ] JJ. Harris, D.E. Ashenford, CT. Foxon, PJ. Dobson, BA. Joyce [Appl. Phys. A (USA) vol. 33 (1983) p.87] JJ. Harris, J.B. Clegg, RB. Beall, J. Castagne, R. Murray, R.C Newman [ Proc. Electrochem. Soc. (USA) vol.90 (1990) p.423] E.F. Schubert, J.M. Kuo, R.F. Kopf, H.S. Luftman, L.C. Hopkins, NJ. Sauer [ J. Appl. Phys. (USA) vol.67 no. 4 (1989) p. 1969]

2.12 Theoretical electron mobility/temperature dependences on carrier concentration and compensation ratio W.Walukiewicz March 1996

A

INTRODUCTION

Electron mobility is one of the most commonly measured semiconductor characteristics. The mobility is determined by the momentum randomizing scattering processes. These dominant processes can be divided into two groups: intrinsic effects which include acoustic and optical phonon scattering and a number of extrinsic effects including ionized impurities, neutral impurities, dislocations and other extended defects. The maximum mobility that can be achieved at a given temperature is called the phonon mobility limit and is determined by the intrinsic scattering processes. Since the material parameters are known, the temperature dependence of the phonon mobility limit is now quite well established in GaAs [1- 4]. A deviation of the actual mobility from the phonon mobility limit can thus be considered a measure of the strength of extrinsic scattering processes. In doped material, the most important extrinsic scattering mechanism is ionized impurity scattering. Since the scattering rate by an isolated charged impurity centre can be easily calculated, a measurement of the electron mobility can be used to evaluate the total concentration of ionized impurities in the sample [2-8]. In most instances, for practical purposes, the theoretical electron mobility is presented as a function of the electron concentration and the compensation ratio 0 which is defined as the ratio of the concentration of compensating acceptors to the concentration of shallow donors, NA/ND. The functions have been calculated for both room and liquid nitrogen temperature and are routinely presented in the form of graphs [2-6] or tables [7,8]. B

THEORETICALELECTRON MOBILITIES

In this Datareview the theoretical electron drift mobilities as functions of the ratio of the total ionized impurity concentration, Njn^, to the electron concentration, n, are presented in the form of graphs and, in a limited electron concentration range, are also approximated by analytic expressions. They provide an easy method for the determination of the total concentration of charged impurities from the measured values of the mobility \i and the electron concentration n. FIGURES l(a) and l(b) show calculated drift electron mobilities at 300 K [6] and 77 K [7] using a variational method [9]. The total ionized impurity concentration is related to the electron concentration n and the compensation ratio 0 through the equation Nimp = n(l+6)/(l-0)

(1)

The calculated values were obtained assuming that there is no spatial correlation between impurities in the samples [10,11]. Also multiple scattering [10,11] and effects of macroscopic inhomogeneities [12] were ignored in these calculations

Electron mobility ( m 2 / V s )

(b)

GaAs 77 K

Concentration ratio, N.

/n

imp

Electron mobility (m 2 /Vs)

(a) GaAs 300 K

Concentration ratio N.

/n

imp

FIGURE 1. Calculated room (a) and liquid nitrogen (b) temperature electron mobilities as functions of the concentration ratio N^p/n. Electron concentration n is a parameter for the different curves.

The use of such graphs for the determination of the total impurity concentration is not always easy as, in general, it requires interpolation between the curves corresponding to different electron concentrations. In a limited concentration range it is possible to obtain an analytic expression approximating the theoretically calculated functions relating the total ionized impurity concentration with electron mobility and electron concentration. The expression has the following general form: N tep = n f ( n ) ^

(2)

where n and N ^ are the free electron and the ionized impurity concentrations in m"3, and \i is the mobility in m2/ Vs, respectively. For room temperature mobility, the functions f(n) and g(n) can be approximated by f(n)

= - 0.41 + 4.1 x Kr2 log(n)

(3)

g(n)

=-4.808 x 10 4 n 02546

(4)

EQNs (2) to (4) provide a good approximation only for electron concentrations n > 1023 m"3 i.e. in the range where the room temperature mobility is significantly affected by the ionized impurity scattering. Determination of the impurity concentration in lightly doped GaAs requires measurements at low temperatures where the mobility is more sensitive to the ionized impurity scattering. For very lightly doped GaAs with electron concentrations in the range 1020 m"3 < n < 3 x 1021 m"3 the ionized impurity concentration can be obtained from the mobility and electron concentration measurements at 77 K. Again N ^ can be obtained from EQN (2) with the functions f and g given by f(n)

=2.46 x 1028nL2817

(5)

g(n)

=-3855.9 n 01609

(6)

The above phenomenological formulae are approximate and can be used only for rough estimates of the ionized impurity concentrations. It should also be noted that, in general, the use of experimental mobilities to determine the ionized impurity concentration has to be treated with great caution and can be applied only to a homogeneous material with a low density of extended structural defects. C

CONCLUSION

Theoretical room and liquid nitrogen temperature electron mobilities are presented as functions of free electron and ionized impurity concentration. The curves and the formulae provide convenient methods to assess the total concentration of ionized impurities from the experimentally measured carrier concentration and the electron mobility. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] [II] [12]

H. Ehrenreich [Phys. Rev. (USA) vol.120 (1960) p.1951-1963 ] CM. Wolfe, G.E Stillman, J.O. Dimmock, [ J. Appl. Phys. (USA) vol.41 (1971) p.504-7 ] DX. Rode, S. Knight [ Phys. Rev. B (USA) vol.3 (1971) p.2534-41 ] D.L. Rode [ Semicond Semimet. (Eds. R.K. Willardson, A,C. Beer (Academic Press, New York, 1975) vol.10 p. 1-89] G.E Stillman, CM. Wolfe [ Thin Solid Films (Switzerland) vol.31 (1976) p. 69 ] W. Walukiewicz, J. Lagowski, L. Jastrzebski, M. Lichtensteiger, H. C Gatos [ J. Appl. Phys. (USA) vol.50 (1979) p.899-908] W. Walukiewicz, J. Lagowski, H.C. Gatos [ J. Appl. Phys. (USA) vol.53, (1982) p.769-70 ] M. Benzaquen, K. Mazuruk, D. Walsh, AJ. Springthorpe, C Miner [J. Electron. Mater. (USA) vol.16 (1987) p.111-17] D. Howarth, E.H. Sondheimer [Proc. Phys. Soc. (UK) vol.A219 (1953) p. 53 -74 ] EJ. Moore [ Phys. Rev. (USA) vol. 160 (1967) p.607-26 ] J.R Meyer, FJ. Bartoli [J Phys C Solid State Phys. (UK) vol.15 (1982) p.1987-99 ] CM. Wolfe, G.E. Stillman [ Semicond Semimet. (USA) (Eds RK. Willardson, A.C Beer (Academic Press, New York, 1975) vol. 10 p. 175-220 ]

2.13 Minority electron mobility in doped GaAs E.S. Harmon, M.L. Lovejoy, M.S. Lundstrom and M.R. Melloch September 1995

A

INTRODUCTION

GaAs bipolar devices such as heterojunction bipolar transistors (HBTs), solar cells and lasers often contain regions where minority electron transport through heavily doped p-GaAs regions is critical to device performance. Because of the inherent difficulties in measuring the minority electron mobility in heavily doped semiconductors, many researchers have been forced to assume that the minority electron mobility in p-GaAs is the same as the majority electron mobility in similarly doped n-GaAs. This assumption can be off by a factor of two or more due to the different scattering mechanisms of minority and majority carriers [1-3]. In this Datareview, we examine the doping and temperature dependences of the minority electron mobility in P+-GaAs. The measurement of minority electron transport in semiconductors often relies on the measurement of the minority electron diffusivity (Dn) rather than a direct measurement of the minority electron mobility. Because of electron-hole scattering, the Einstein relation applies only at sufficiently low field [2,3], but we can define a 'minority carrier diffusion mobility' by: \im = —

(D

where, in general, the minority electron diffusion mobility (E = 0) is not equivalent to the minority electron drift mobility (E * 0) [2,3]. B

MEASUREMENT TECHNIQUES

There are three basic techniques available for the measurement of minority electron mobilities in P+-GaAs. For a review of techniques including those suitable for lower doped material, see reference [4]. The three basic techniques rely on measurements of very fast signals to determine the transit time of electrons through a region of p-GaAs. High frequency measurement techniques and packaging must be used in order to minimize the effects of carrier recombination and parasitics. Note that lifetimes in heavily doped GaAs are on the order of one nanosecond or less, requiring the packaging and measurement setup to be capable of measuring frequencies in the multi-gigahertz range. The three measurement techniques will be referred to as the conventional time-of-flight technique, the unity gain transistor cut-off frequency technique, and the zero-field time-of-flight technique. Bl

The Conventional Time Of Flight (TOF) Technique

This is the standard technique for measuring mobility (Haynes-Shockley experiment [5]). In the TOF technique carriers are injected at one end of a uniformly doped semiconductor region, and

the average transit time across a well defined distance is measured [5,6]. This is typically done with high applied electric fields so that most of the electrons traverse the distance before recombining. From the measured average transit time and the distance travelled, the average velocity is measured, and the mobility is calculated by dividing the average velocity by the field strength. As noted earlier, the 'minority electron drift mobility' is not equivalent to the minority electron diffusion mobility [2,3]. This is primarily because the applied electric field causes the majority holes to flow in the opposite direction to the minority electrons. Collisions between the majority holes and minority electrons result in a net momentum transfer to the electrons, reducing the average electron velocity [7] ('hole drag' effect). There have even been measurements of negative minority electron mobilities at low temperatures in two-dimensional hole gas experiments [8]. B2

The Unity Gain Transistor Cut-off Frequency Technique

This overcomes the majority carrier drag effect of the TOF technique by measuring the electron transit time through the zero-field region in the base of a bipolar transistor. The cut off frequency (fT) gives the total transit time of electrons from the emitter to the collector [9]. By measuring fT at various collector current densities and extrapolating to infinite collector current, the emitter and collector charging times may be eliminated. The resulting fT (at infinite collector current) is a function of only the base and collector transit times. To separate the two transit times, HBT structures can be designed where the base transit time dominates the total transit time, or devices with varied base widths may be utilized to determine uniquely the base transit time. The base transit time is then directly related to the diffiisivity through the diffusion equation (xB = W2/2D). The unity gain transistor cut-off frequency technique requires HBTs with a current gain greater than unity, which restricts the usable base thicknesses to relatively thin layers. For heavily doped p-GaAs, this restriction results in base thicknesses on the order of 2000 A, which results in base transit time in the ~1 to 25 ps range. Thus, very careful experimental techniques must be used to minimize and de-embed parasitics [10]. B3

The Zero Field Time Of Flight (ZFTOF) Technique

This third method is also used to obtain the minority electron diffusion mobility [11,12]. The ZFTOF technique measures the collection time for electrons generated by pulsed laser illumination near the top of a vertical p+n photodiode. The photogenerated electrons must diffuse through the P+ region of the photodiode before being collected by the junction. The current response of the photodiode may then be modelled with a numerical simulation or an infinite series solution of the diffusion equation to determine the diffusion coefficient. Note that the current response must be dominated by the diffusion response for accurate diffusion coefficients to be obtained. If significant recombination occurs, either in the bulk p+ region or at the top surface, then the diffusion coefficient cannot be accurately determined unless the recombination rates are also known. Fortunately, surface recombination may be minimized by using an AlGaAs passivation layer on top of the P+-GaAs active layer and the effects of bulk recombination may be minimized by ensuring that the electron transit time through the p+ region is shorter than the recombination time. The effects of bulk recombination can be measured by determining the electron collection efficiency for steady state illumination (solar cell quantum efficiency) [H].

C

EXPERIMENTAL RESULTS

The minority electron mobilities in P+-GaAs are presented in FIGURE 1. Included in the figure are the data from TOF [6,13], fT [14-16], and ZFTOF [17-20] measurements. The measured minority electron mobilities are significantly lower than the majority electron mobilities [21] for doping densities less than 1019 cm'3. For higher doping densities a dramatic increase in the measured minority electron mobility is observed as the doping density is increased from 0.9 x 1019 cm"3 to 8.0 x 1019 cm"3. The measured mobility of 3710 cm2 / Vs at 8.0 x 1019 cm"3 is about three times higher than that measured for 0.9 x io 19 cm"3. The substantial rise in mobility as doping density is increased beyond 1 x 1019 cm"3 is in stark contrast to the mobility saturation observed for majority carrier mobilities. Minority electron mobilities in p-GaAs doped with both beryllium and carbon have been investigated. Although doping of GaAs occurs by very different mechanisms for the two dopants, the mobility of minority electrons in P+-GaAs with hole concentrations in the range 0.8-0.9 x 1019cm"3 is independent of dopant [18].

majority electrons 21 theoretical minorttyl ZFTOF (Be) Harmon17 Lovejoy18'19 ZFTOF (C) Harmon17Lovejoy18 ZFTOF Colomb 20 fTKim 1 6 fTLeeiS fT Beyzavi 14 TOF Furuta 13

Diffusivity (Cm2S"1)

mobility (Cm 2 V 1 S" 1 )

It is important to compare the results of the different measurement techniques. We see that all three measurement techniques agree quite well for doping densities less than 1019 cm"3. However, for higher doping densities, the conventional TOF results are significantly lower than those observed for the zero-field measurement techniques. This difference is most likely due to the majority hole drag effect in conventional TOF measurements described above (the minority electron drift mobility is less than the minority electron diffusion mobility). The ZFTOF [17-20] and fT [14-16] results shown in FIGURE 1 are in excellent agreement.

Doping Density (cm"3) FIGURE 1. Minority electron diffiision mobility in P+-GaAs as determined by the ZFTOF technique [17-20] and the fT technique [14-16]. Also shown are the theoretical calculations of Lowney and Bennett [1], the majority electron mobility in n+-GaAs [21], and the minority electron drift mobility [13] in P+-GaAs.

Also shown in FIGURE 1 are the theoretical mobility calculations from Lowney and Bennett [1] which show general agreement with the measured data. Their calculations included all the important scattering mechanisms for GaAs. Lowney and Bennett showed that treating minoritymajority carrier scattering by ionized impurity scattering is not adequate for heavily doped GaAs. Below a hole concentration of 1019 cm"3, plasmon scattering is significant, resulting in significantly smaller overall mobilities than those obtained when only single minority-majority carrier scattering is considered. As the doping density is increased beyond 1019 cm"3, plasmon scattering and carrier-carrier scattering rates are suppressed, resulting in an increase in the overall minority electron mobility. Plasmon scattering is reduced because the hole plasmon frequency becomes too high for interaction with minority electrons to occur, and carrier-carrier scattering is reduced because the number of states available to receive scattered holes is reduced due to degeneracy. D

MINORITY ELECTRON MOBILITY TEMPERATURE DEPENDENCE

mobility (cm2V" V1)

The temperature dependence of minority electron mobility differs greatly from the temperature dependence of majority electron mobility. Shown in FIGURE 2 is a comparison of the minority electron mobility and majority carrier mobilities in comparably doped n- and p-GaAs.

Minority electrons; Beyzavi14 1/Tfit Majority electrons; Lovejoy 22 Majority holes; Lovejoy22

Temperature (K) FIGURE 2. Mobility vs. temperature for minority electrons in 4 x 1018 cm'3 P+-GaAs [14] and both majority holes and electrons in comparably doped (4 x 1018 cm"3) GaAs [22]. The majority Hall mobilities are found to be constant while a 1/T temperature dependence is exhibited by the minority electron mobility.

The minority electron mobility in p-GaAs doped to 4.2 x 1018 cm"3 was measured by Beyzavi et al with HBT unity gain cut-off frequency (fT) measurements for temperatures (T) from ~ 77 K to 300 K [14]. Also shown in FIGURE 2 is the majority electron mobility [22] for n-GaAs doped to 4.2 x 1018 cm"3 and majority hole mobility [22] in p-GaAs doped to 4.2 x 1018 cm"3 which is comparably doped to the sample in which Beyzavi et al measured the minority electron mobility. It is noted that Beyzavi's room-temperature measurement of minority electron mobility is lower

than the minority electron mobility measurement with the ZFTOF technique [19]. In the temperature range investigated, the majority carrier electron mobility and majority hole mobility are relatively constant at -2000 and -120 cm2/ Vs respectively while the minority electron mobility exhibits a 1/T dependence. Walukiewicz et al calculated the temperature dependence of minority electron mobility in p-GaAs doped to 1 x 1018 cm"3 [23]. In their treatment, minority-majority carrier scattering was treated by the Brooks-Herring impurity scattering rate formula which has a T3/2 dependence and they included a screening length expression appropriate for degenerate majority carrier systems. Their result is that minority electron mobility was predicted to have a weaker than 1/T dependence. However, Szmyd et al derived a scattering rate expression for degenerately doped material that is analogous to the Brooks-Herring expression which is for non-degenerate material [24]. This treatment yielded a mobility that is essentially temperature independent. Combining this result with the fact that the Walukiewicz et al theory did not accurately predict the mobility at a hole concentration above 1 x 1019 cm"3 in GaAs suggests that the scattering physics are more complicated than considered by the first theory. A similar temperature dependence has been measured for minority holes in GaAs [22]. It has been shown that carrier freeze-out, which reduces the ionized impurity scattering rate, is not responsible for the 1/T dependence and a degeneracy effect due to decreased temperature has been speculated to be the origin [22]. E

CONCLUSION

Minority electron mobilities in heavily doped p-GaAs are shown to have significantly different doping- and temperature-dependences from those of the majority electron mobilities in similarly doped n-GaAs. For doping densities less than 1019 cm"3, the minority electron mobility is only 50% to 75% of the corresponding majority electron mobility. As doping densities increase from 0.9 x 1019 cm"3 to 8 x 1019 cm"3, a dramatic 300% increase in the minority carrier mobility is observed. This increase in mobility is in agreement with theories that account for plasmon and carrier-carrier scattering between the minority electrons and majority holes. The very high mobilities observed for the highest doping densities indicate that it may be desirable to utilize such heavily doped materials in the critical delay paths of high speed devices. A very different temperature dependence of minority electron mobility as compared to majority electron mobility in comparably doped GaAs is found as well. Minority mobility exhibits a 1/T dependence while the majority mobility is essentially independent of temperature. The different dependences confirm that different scattering mechanisms dominate mobilities for minority and majority carriers. Finally, we stress the need to differentiate between the minority carrier diffusion mobility and the minority carrier drift mobility. In regions of sufficiently high fields, the drift mobility is significantly lower than the diffusion mobility in heavily doped materials because of hole drag effects. ACKNOWLEDGEMENTS This work was supported by the Department of Energy under contract #DE-AC04-94AL85000. REFERENCES [1] [2]

J.R. Lowney, H.S. Bennett [J. Appl. Phys. (USA) vol.69 (1991) p.7102 ] D.E. Kane, R.M. Swanson [ IEEE Trans. Electron Devices (USA) vol.40 (1993) p. 1496 ]

[3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24]

TT. Mnatsakano, B.N. Gresserov, L.I. Pomortseva [ Solid-State Electron. (UK) vol.38 (1995) p.225] M.L. Lovejoy [ AIP Press of the American Insitute of Physics (1995) ] J.R. Haynes, W. Shockley [ Phys. Rev. (USA) vol.81 (1951) p.83 ] T. Furuta, H. Taniyama, M. Tomizawa [ J. Appl. Phys. (USA) vol.67 (1990) p.293 ] W.P. Dumke [ Solid-State Electron. (UK) vol.28 (1985) p. 183 ] R.A. Hopfel, J. Shah, P.A. Wolff, A.C. Gossard [ Phys Rev. Lett. (USA) vol.56 (1986) p.2736 ] M.I. Nathan, W.P. Durnke, K. Wrenner, S. Tiwari, S.L. Wright, K.A. Jenkins [Appl. Phys. Lett. (USA) vol.52 (1988) p.654 ] S. Lee, A. Gopinath [ IEEE Trans. Microw. Theory Tech. (USA) vol.40 (1992) p.574 ] ML. Lovejoy, M.R. Melloch, R.K. Ahrenkiel, M.S. Lundstrom [ Solid-State Electron. (USA) vol.35 (1992)p.251] R.K. Ahrenkiel, DJ. Dunlavy, D. Greenberg, J. Schlupmann, H.C. Hamaker, H.F. MacMillian [Appl. Phys. Lett. (USA) vol.51 (1987) p.766 ] T. Furuta, M. Tomizawa [Appl. Phys. Lett. (USA) vol.56 (1990) p.824 ] K. Beyzavi, K. Lee, DM. Kim,M. I. Nathan,K. Wrenner, S. L. Wright [Appl. Phys. Lett. (USA) vol.58 (1991) p. 1268] S. Lee, A. Gopinath, SJ. Pachuta [ Electron. Lett. (UK) vol.27 (1991) p. 1551 ] D.M. Kim et al [ Appl. Phys. Lett. (USA) vol.62 (1993) p.861 ] E.S. Harmon, M.L. Lovejoy, M.R. Melloch, M.S. Lundstrom, TJ. de Lyon, J.M. Woodall [ Appl. Phys. Lett. (USA) vol.63 (1993) p.536 ] M.L. Lovejoy et al [ Appl. Phys. Lett. (USA) vol.61 (1992) p.822 ] M.L. Lovejoy [ Ph. D. thesis, Purdue University (USA, 1992) ] CM. Colomb, S.A. Stockman, S. Varadarajan, G.E. Stillman [ Appl. Phys. Lett. (USA) vol.60 (1992) p.65] D.A. Anderson, N. Apsley, P. Davies, P.L. Giles [ J. Appl. Phys. (USA) vol.58 (1985) p.3059 ] M.L.Lovejoy, MR. Melloch, M.S. Lundstrom [Appl. Phys. Lett. (USA) vol.60 (1992)p.65 ] W. Walukiewicz, J. Lagowski, L. Jastrzebski, H.C. Gates [J. Appl. Phys. (USA) vol.50 (1979) p.5040 ] DM. Szmyd, M.C. Hanna, A. Majerfeld [ J. Appl. Phys. (USA) vol.68 (1990)p.2376 ]

2.14 Electron lifetimes in p-type GaAs G.B. Lush October 1995

A

INTRODUCTION

Electron lifetimes in p-type GaAs have long been well understood, and recent data have only confirmed the basic behaviour of the lifetime as a function of doping concentration. Remarkably consistent data have been reported over a span of nearly 20 years, nearly all of which indicate that the electron lifetime in p-type GaAs can be modelled as inversely proportional to doping - the signature of radiative recombination (see Datareview 3.8 in this book on lifetimes in n-type GaAs for a discussion of recombination mechanisms, photon recycling, and other factors influencing the overall lifetime of a semiconductor versus doping concentration). The data reported in this Datareview were observed on samples using three different growth techniques, three different dopants, and two measurement techniques on at least four different types of structure. B

REPORTED p-TYPE LIFETIME DATA

Electron Lifetime (ns)

Electron lifetimes reported in p-type GaAs are much better behaved and much more uniform across growth techniques than the data obtained from n-type GaAs. FIGURE 1 displays reported data spanning nearly 20 years, and spanning three growth techniques: liquid phase epitaxy (LPE) [1,2], vapour phase epitaxy (MOVPE) [3,4], and molecular beam epitaxy (MBE) [5].

Nelson and Sobers Acket t Hooft Melloch Radiative Limit

Hole Concentration (cm 3 ) FIGURE 1. Reported data for electron lifetimes in p-type GaAs.

These data represent the observed lifetime of a double heterostructure, and do not include analysis to account for interface recombination or photon recycling (except for Acket [2] who accounted for interface recombination). These data do not represent the radiative lifetime. C

HOW TO USE THESE DATA

Recombination in p-type GaAs seems to be dominated by radiative recombination and photon recycling, but regardless of the mechanisms, the observed electron lifetimes follow the equation

where x is the lifetime, Na is the acceptor concentration and B varies by a small amount with hole concentration. This is true regardless of the film growth technique. FIGURE 1 shows for comparison the expected radiative limit assuming B = 1.0 x 10"10 cm3/s [1,6]. The values plotted in FIGURE 1 are the sum total of all mechanisms present; it is unlikely that newly grown material will deviate significantly from these data. It is unnecessary to be concerned about which recombination mechanisms are involved unless very detailed modelling of devices is necessary. If so, the references must be studied carefully to see their analyses: see also Datareview 3.8 Hole lifetime in n-type GaAs by Melloch and Lundstrom, in this book. The data are remarkably consistent considering the diversity of sources. The l/BNa trend shown and the analysis reported in the references indicate that radiative recombination and photon recycling are the dominant mechanisms determining the electron lifetimes in p-type GaAs. For Na > 1017 cm"3, the reported electron lifetimes follow the line for B = 1.0 * 10"10cm3/s. For Na < 1017 cm'3, the reported electron lifetimes would follow a line for B = 2.0 x 10'10 cnxVs. (Such a line would roughly be drawn through the data of t'Hooft [3,4].) A rarely noted fact is that Acket et al [2] observed longer lifetimes in highly compensated material than those plotted here. The authors attributed the longer lifetimes to increased band tailing in the more highly compensated samples. D

CONCLUSION

Recombination in p-type GaAs is well understood in terms of what mechanisms are responsible. Much analysis has also been done to compare the laboratory data to theoretical predictions of the radiative lifetimes, and theory agrees very well with data [I]. REFERENCES [1] [2] [3] [4] [5] [6]

RJ. Nelson, RG. Sobers [ J. Appl Phys. (USA) vol.49 no. 12 (1978) p.6103 ] G.A. Acket, W. Nijman, H.'t Lam [ J. Appl Phys. (USA) vol.45 no.7 (1974) p.3033 ] G.W. t'Hooft, C. Van Opdrop, H. Veenvliet, A.T. Vink [ J. Cryst. Growth (Netherlands) vol.55 (1981)p.l73-82] G.W. t'Hooft [Appl Phys. Lett. (USA) vol.39 no.5 (1981) p.389 ] [ Data from double heterostructures grown by MBE, measured by photoluminescence decay (unpublished) ] H.C. Casey Jr., F. Stern [ J. Appl Phys. (USA) vol.47 no.2 (1976) p.631 ]

CHAPTER 3 HOLE MOBILITY, DIFFUSION AND LIFETIME 3.1 3.2 3.3 3.4 3.5 3.6 3.7 3.8

Hole mobility in doped and ion implanted GaAs Hole mobility in GaAs, temperature dependence Hole mobility in GaAs, pressure dependence Carbon doping of GaAs Carrier concentration dependence of the hole mobility in GaAs Minority hole mobility in GaAs Theoretical hole mobility curves Hole lifetimes in n-type GaAs

3.1

Hole mobility in doped and ion implanted GaAs Yung Kee Yeo August 1995

A

ACCEPTOR DOPING STUDIES

In addition to the scattering of holes due to lattice vibrations [1], there is also scattering due to imperfections, i.e., impurities or defects. In most cases, ionized impurity scattering is much more important than neutral impurity scattering, and ionized-impurity scattering is reasonably well described by the Brooks-Herring treatment [2]. However, as discussed in [1], a detailed calculation of hole mobility is much more complicated than for electron mobility [3], so that only semi-quantitative agreement between theory and experiment is possible. Furthermore, many samples are compensated, causing a reduction in mobility, but the compensation ratio is not known in most cases. In spite of that, a rough fit can be made to the data by assuming that the hole concentration, p, is approximately the same as the ionized-impurity concentration, and then applying Matthiessen's rule, l/\x ~ V\iL + 1/In1. Wiley [4] has reviewed hole data until 1975, and has plotted hole mobility vs. hole concentration at 300 K, data from many different workers [5-11] spanning hole concentration from 2 x 1014 to 9 x 1019 cm"3. The plotted line was calculated using the Brooks-Herring ionized impurity scattering model together with an assumed lattice-limited mobility of 400 cm2/ Vs. The line that Wiley [4] has drawn through the 300 K data can be roughly fitted by the following equation given by Look [12] HH= 1 / { 2.5 x 10"3 + 1.1 x 10"21 [In(Hb) - b/(l+b)] p } (1) where b » 7.5 * 1019/ p; p has units of cm"3, \iH has units of cm2/ Vs, and the subscript H designates a Hall-effect measurement without correction for the Hall factor. EQN (1) should be reasonably valid for p below ~1018 cm"3. The estimated |iiH from EQN (1) is shown in TABLE 1. TABLEl Carrier concentration (cm"3)

1015

1016

1017

1018

Hall mobility ( c m 2 / V s )

398

387

321

162

Recently, Zhang et al [13] have also plotted the Hall mobility at 300 K versus hole concentration of MBE grown Be-doped GaAs ranging from 1015 to 2 x 1020 cm"3 including other works [14-16]. Their plotted curve is based on the empirical Hilsum formula [17] originally presented for electron mobility,

n H =m/[i+(p/ P o r]

(2)

where the parameters JLi0, P0, and m have empirical values of 430 cm2/ Vs, 6 x 1017 cm'3, and 0.45, respectively. The estimated \iu is shown in TABLE 2.

TABLE 2 Carrier concentration (cm 3 )

1015

1016

1017

1018

1019

1020

Hall mobility (cm2 / Vs)

407

371

297

190

95

39

In addition to the above cited works, there have been numerous other reports of hole mobility measurements in GaAs, and currently, sufficiently reliable experimental hole mobility data are available for hole concentrations up to ~1020 cm'3. Stockman et al [18] have reported hole mobility measurements at 300 K as a function of hole concentration ranging from 1017 to 4 x 1020 cm"3 for C-, Be-, and Zn-doped GaAs grown by MOCVD, MOMBE, and MBE including other studies [19-22]. They observed that for p > 2 x 1018 cm"3, the hole mobilities for Be-doped GaAs grown by MBE are always at least 20% lower than those attainable using C. Similarly, they observed that while Zn-doped GaAs grown by MOCVD exhibits hole mobilities that are nearly equal to those using Be and C for p < 2 x 1018 cm"3, they are also significantly lower than for Cdoped GaAs when p > 2 x io18 cm"3. This may be related to higher compensation when GaAs is doped heavily with Be or Zn. Nowak et al [23] also reported Hall mobilities at 300 K for Zn diffused GaAs together with those for heavily Zn-doped melt-grown bulk GaAs with concentrations ranging from 1018 to 1020 cm"3. Yamada et al [24] have doped GaAs with Be to attain hole concentrations ranging from -10 17 to 1.3 x 1019 cm"3 at 300 K using a conventional MBE system, and also doped GaAs with C from ~1019 to the 1021 cm"3 range using an MOMBE method, and measured the hole mobility vs. hole concentration. At p = 1018 cm"3, their data agree reasonably well with the value of juH = 162 cm2/Vs calculated from EQN (1). For p > 1018 cm"3, the following empirical formula seems to fit their data with sufficient precision [12]: HH =162 (10 1 V p) 0 2 7 5

(3)

Similarly, Nakwaski [25] reported the variation of hole mobility with hole concentration for Cdoped GaAs grown by different techniques for concentrations ranging from 1018 to over 1020 cm"3 at 300 K. They concluded that for C-doped GaAs, all the hole mobility values at 300 K are within ±20 cm2/ Vs from the following expression: ^H = 180-60 (log p - 1 8 )

(4)

However, EQN (4) gives consistently higher \iu values than those obtained from EQN (3). It is worthwhile to mention that Stockman et al [18] also characterized MBE-grown Be-doped GaAs with 1 x 1O19< p < 6 x 1019 cm"3, and reported that the \iH was -20 - 30% lower than that of Cdoped layers with the same doping concentrations over the entire sample temperature range studied. More recently, Kim et al [26] measured the Hall mobilities of C-doped GaAs grown by atmospheric pressure MOCVD as a function of hole concentration ranging from 3.5 x 1017 to 1.5 x 1020 cm"3, including data for C-doped GaAs grown by low pressure MOCVD [27,28]. They showed that the mobilities of C-doped GaAs were much higher than those of Zn-doped GaAs, probably due to a lower compensation ratio for the C-doped GaAs. For example, the mobility for the C-doped GaAs is -80 cm2/ Vs, whereas that of the Zn-doped GaAs is -50 cm2/ Vs for p

* 4 x io 19 cm"3. Kim et al [27] obtained a hole mobility of 42 cm 2I Vs at 300 K for C-doped GaAs grown by low pressure MOCVD for p = 1.4 x 1020 cm"3. Sulllivan et al [29] also measured Hall mobility at 300 K as a function of hole concentration ranging from 1.6 x 1018 to 1020 cm"3 for C-doped GaAs grown by MBE. They obtained, for example, a hole mobility of 139 cm2/ Vs for a hole concentration of 2.6 x io18 cm"3. They showed that mobilities of C-doped GaAs are comparable to high quality Be-doped GaAs grown in the same MBE system. Yang et al [30] presented a table for hole concentration (1 x io 19 - 1.2 x 1020cm"3) versus mobility at 300 K for heavily C-doped GaAs grown by low-pressure OMVPE. The hole mobility decreases from 106 to 65 cm2/ Vs as the hole concentration increases from IO19 to 1.2 x io 20 cm"3 (TABLE 3). TABLE 3 Carrier concentration (x 1019cm-3)

1.0

2.1

3.4

3.8

4.6

5.6

6.0

7.3

7.7

12

Hall mobility (cm2 / Vs)

106

92

82

80

80

74

74

74

69

65

The temperature dependence of hole mobility over the range 77 - 400 K for a given hole concentration can be roughly represented by [12] \iH = 1 / { 5.0 x 10"9 T 23 + 5.5 x 10"18 [In(Kb) - b/(l+b)] p / T L5 }

(5)

where b « 8.3 x io 14 T2/p. EQN (5) is derived from the lattice-limited formula in [1] and from the expected temperature dependence of JLi1. Stockman et al [18] reported the temperature dependent Hall mobility of MOCVD-grown C-doped GaAs. The mobility for p = 2.0 x io 17 cm"3 increases from -250 to -900 cm2/ Vs as the temperature decreases from 300 to -80 K due to a decrease in phonon scattering. However, as the temperature is lowered below 80 K3 the mobility decreases due to increased ionized impurity scattering. In general, at temperatures below -100 K, the mobility is dominated by ionized impurity scattering, and at high temperatures, the total mobility may be roughly estimated by combining the effects of ionized impurity and phonon scatterings according to Matthiesen's rule (1/ji = 1/^+ 1/In1) with HL(T) - Ii0 (297/T)* Stockman et al [18] have determined that ^0 = 435 cm2/ Vs and a = 2.37 from high-purity p-type MBE-grown GaAs. The 77 K mobility in MOMBE-grown C-doped samples for p > IO19 cm"3 was typically 50 to 60% higher than at 300 K. The expression for total mobility described above showed good agreement with the measured data when using a = 1.5 and \i0 = 250 to 150 cm2/Vs for p = 1 x io 19 to 1.4 x 1020cm"3, respectively. In comparison, Rosi et al [9] estimated ^L(T) « 418 (300/T)23 and the temperature exponent of 2.3 gives excellent agreement with experimental data [4]. Zhang et al [13] also plotted the 77 K Hall mobilities as a function of carrier concentration for Be-doped GaAs grown by MBE, including previously reported data for LPE [31] and MBE layers [16,3234]. Their curve fitting was calculated from the empirical Hilsum formula EQN (2) using the parameters (I0= 10 000 cm2/ Vs, p0 = 5 x io 15 cm"3, and m = 0.7. Their data follow the line well for hole concentrations up to 1018 cm"3, but the Hall mobility data decrease much more slowly than the empirical curve as the hole concentrations increase to more than 1018 cm"3. However,

they showed that their mobility data are consistent with those reported by Ploog et al [16]. It must be emphasized that the expressions given by EQNs (1) through (5) have no real theoretical significance, but they are useful for providing mobility vs. concentration curves which follow the general trend of the data. Furthermore, variations in the compensation ratio can cause considerable deviation from the values predicted by EQNs (1) through (5). B

ION IMPLANTATION OF ACCEPTORS

In the past decades, there have been extensive studies of the ion implantation of p-type dopants into GaAs [35-39]. The most commonly used p-type ion species are Be, Mg, Zn, and Cd [40]. The substrates used were usually either Cr-doped or undoped semi-insulating GaAs. Contrary to the behaviour of n-type dopants, acceptor species of Be, Mg, Cd, and Zn show high activation efficiency close to 100% for doses up to 1014 cm'2 [36,39,41]. For C implants in GaAs, a hole concentration close to 1020 cm"3 was achieved [42]. Ion implantation in GaAs is done usually at room temperature. To improve the activation efficiency, dual implantation with a stoichiometry restoring species or hot temperature implantation is performed. In order to remove the damage created during ion implantation of GaAs, and to obtain the best electrical properties, it is necessary to optimise the annealing conditions. Conventional furnace annealing was usually performed at temperatures between 600 and 900 0 C for 5-30 min. A major problem with acceptor impurities in GaAs is in- and outdiffusion during annealing. Currently, halogen lamp rapid thermal annealing (RTA) at 80010000C for 5-20 s is widely used to minimize sample decomposition and dopant redistribution. To protect the sample surface from thermal decomposition during annealing, a dielectric (SiO2 or Si3N4, etc.) cap and/or a bare GaAs wafer proximity cap is used. Bl

Zn and Mg Implants

A great deal of the early work in Zn and Mg implantation and annealing was done with furnace annealing [36,43,44], which showed a significant diffusion of the dopants after high temperature annealing. Nevertheless, Yeo et al [45] found in GaAs:Mg that \iH » 100 cm2/ Vs at 300 K at a peak concentration of p « 5 x 1018 cm"3, which agrees well with the predicted |i H of -10 4 cm W s from EQN (3). For Zn-doped material, Emel'yanenko et al. [46] found that \iH « 1.0 x 10"3 cm2/Vs at 100 K for p « 1.2 x 1016cm'3, while the prediction from EQN (5) would have been -1.8 xlO 3 cm2/ Vs. This difference could easily be explained by compensation but other factors such as residual ion-implantation induced damage may also enter. For heavily Zn-implanted GaAs [47] with p * 8 x 1019cm"3, |i H is -45 cm2/ Vs at 300 K, whereas the prediction from EQN (3) would have been ~ 48 cm2/ Vs. Later work in Zn and Mg with RTA has shown that good activation could be achieved with minimal in-diffixsion, except at very high implant doses where diffusion was still observed, especially near the peak [37,48-51]. Recently, Sherwin et al [35] have shown that well controlled dopant profiles can be obtained with good activation for both Zn and Mg implantation [37,39,48,49,52] after RTA. In order to curtail and to improve activation of Mg and Zn implants, As or P co-implantation [48,49,53] has also been used. Shen et al [54] measured hole concentration and mobility depth profiles in GaAs implanted with Mg alone at 125 keV with a dose of 5 x io 14 cm"2, and co-implanted with P at 160 keV with the same dose, and annealed at 10500C for 5 s with the RTA method. The single Mg implantation produced a peak

concentration of ~2 x 1018 cm'3 with |i H * 100 cm2/ Vs compared to the value of 134 cm 2 / Vs predicted from EQN (3), whereas the Mg and P dual implantation produced a peak concentration of ~2 x io19 cm"3 with ^H « -75 cm2/ Vs, which agrees well with the predicted mobility of 71 cm2/ Vs from EQN (3). B2

Be Implants

Be implantation has also been extensively studied [55,56] due to its low atomic mass, and hence its deep distribution at low implantation energies. However, it presents an anomalous rapid diffusion at high fluences in GaAs [57]. Rao et al [58] studied high energy (1-3 MeV) Be implantation into LEC-grown GaAs in the dose range of 4 x 1012 - 1014 cm'2. The hole mobility was between 264 and 298 cm2/ Vs in the dose range 4 x 1012 to 3 x io 13 cm"2, but decreased considerably to between 163 and 187 cm2/ Vs for the IO14 cm"2 Be implant. B3

C Implants

Recently, C has been used as an acceptor impurity in GaAs [59] due to its low redistribution during annealing, but the activation of C in GaAs is poor compared to other acceptors. Jiang et al [60] studied C-implanted GaAs, and obtained an activation efficiency as high as 75% for low dose implants, while for highest dose implants, it dropped to as low as 0.5%. Hole mobilities showed little dependence on anneal temperature but decreased with increasing implant dose, ranging from ~ 95 cm2/ Vs for low dose (5 x io 13 cm"2) to -65 cm2/ Vs for high dose (5 x 1015 cm"2) samples. This is most likely a consequence of scattering of holes by ionized impurities and implantation-induced defects. They reported that these mobility values are the same or higher than those for Be-, Zn-, or Cd-implanted GaAs annealed under similar conditions [40,61]. Jiang et al [60] also measured hole concentration and mobility profiles of C-implanted GaAs with a dose of 5 x io 14 cm"2. The mobilities ranged between 100 and -80 cm2/ Vs for hole concentrations between ~1018 and ~ IO19 cm"3, which are about the same as or slightly higher than those reported for Zn-, Be-, and Cd-implanted GaAs [40,61,62], but these mobility values are lower than the values predicted from EQNs (2) - (4). Moll et al [63] showed that co-implantation of C into LEC-grown GaAs with Ga dramatically improved the electrical activity of C acceptors. They obtained their highest sheet hole concentration of 3.2 x io 14 cm"2, which corresponds to the volume concentration of about 3 x 1019cm"3. They reported that the mobilities attained for the C+Ga implanted GaAs are similar to those reported for MOCVD and MOMBE GaAs epitaxial layers with similar C doping levels [24]. C

CONCLUSION

Since ion-implanted GaAs usually produces a non-uniform carrier depth profile, and since the volume carrier concentrations are dependent on the implanted ion species, ion energies, doses, and annealing conditions, it is difficult to compare directly the hole mobility values obtained from one implanted sample with another, and especially with uniformly doped epitaxial GaAs layers. However, in general, low dose implants (below IO13 ions/cm2) produce the highest mobilities with resultant carrier concentrations of -2 - 3 x IO17 cm"3, and Hall mobility decreases as the ion dose increases. For example, as the Mg-ion dose increases from IO13 to 1015 cm"2, the hole mobility decreases from -250 to -90 cm2/ Vs [64]. However, the decrease in mobility is more closely related to an increase in the carrier concentration than to the ion dose.

REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [ 10] [II] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33] [34] [35] [36] [37] [38] [39] [40]

D.C. Look [ in Properties of Gallium Arsenide, 2nd Edition, EMIS Datareviews Series No.2 (INSPEC, IEE, 1990), ch.6 p.99-100 ] H. Brooks [ Adv. Electron. Electron Phys. (UK) vol.7 (1955) p.85-182 ] B.R. Nag [ Electron Transport in Compound Semiconductors (Springer-Verlag, Berlin, 1975) ] J.D. Wiley [ Semicond. Semimet. (USA) vol. 10 Eds R.K. Willardson, A.C. Beer (Academic Press, New York, 1975) p.91-174 ] O.V. Emel'yanenko, T.S. Lagunova, D.N. Nasledov [ Sov. Phys.-SolidState (USA) vol.2 (1960) p. 176-80] Sh.M. Gasanli, O.V. Emeryanenko, V.K. Ergakov, F.P. Kesamanly, T.S. Lagunova, D.N. Nasledov [ Sov. Phys.-Semicond. (USA) vol.5 (1972) p. 1641-4 ] D.E. Hill [ Phys. Rev. (USA) vol. 133 (1964) p.A866-72 ] D.E. Hill [J Appl. Phys. (USA) vol.41 (1970) p.1815-8 ] F.D. Rosi, D. Meyerhofer, R.V. Jensen [ J. Appl. Phys. (USA) vol.31 (1960) p. 1105-8 ] F.E. Rosztoczy, F. Ermanis, I. Hayashi, B. Schwartz [ J. Appl. Phys. (USA) vol.41 (1970) p.264-70] J. Vilms, J.P. Garrett [ Solid-State Electron. (UK) vol. 15 (1972) p.443-55 ] D.C. Look [ in Properties of Gallium Arsenide, 2nd Edition, EMIS Datareviews Series No.2 (INSPEC, IEE, 1990), ch.6 p. 101-2 ] D. H. Zhang, K. Radhakrishnan, S.F. Yoon, H. M. Li [ J. Vac. Sci. Technol. A (USA) vol. 12 (1994) p. 1120-3] M. Ilegems [ J. Appl. Phys. (USA) vol.48 (1977) p. 1278-87 ] J.C. Bean, R. Dingle [ Appl. Phys. Lett. (USA) vol.35 (1979) p.925-7 ] K. Ploog, A. Fischer, H. Kunzel [J Electrochem. Soc. (USA) vol.128 (1981) p.400-10 ] C. Hilsum [ Electron. Lett. (UK) vol. 10 (1974) p.259-60 ] S.A. Stockman, G.E. Hofler, J.N. Baillargeon, K.C. Hsieh, K.Y. Cheng, G.E. Stillman [J. Appl. Phys. (USA) vol.72 (1992) p.981-7 ] RW. Glew [ J. Cryst. Growth (Netherlands) vol.68 (1984) p.44-7 ] M. Konagai, T. Yamada, T. Akatsuka, K. Saito, E. Tokumitsu, K. Takahashi [ J. Cryst. Growth (Netherlands) vol.98 (1989) p. 167-73 ] P.M. Enquist [ Appl. Phys. Lett. (USA) vol.57 (1990) p.2348-50 ] M.C. Hanna,ZH. Lu, A. Majerfeld [Appl. Phys. Lett. (USA) vol.58 (1991)p. 164-6 ] E. Nowak, H. Neumann, G. Kuhn [ Cryst. Res. Technol. (Germany) vol.24 (1989) p.K13-5 ] T. Yamada et al [ J. Cryst. Growth (Netherlands) vol.95 (1989) p. 145-9 ] W. Nakwaski [ Phys. Status Solidi A (Germany) vol. 132 (1992) p.K47-9, and references therein] SJ. Kim, Y. Kim, M.S. Kim, CK. Kim, S.K. Min, C. Lee [ J. Cryst. Growth (Netherlands) vol. 141 (1994) p.324-30 ] S.I. Kim et al [ J. Crystal Growth. (Netherlands) vol. 126 (1993) p.441-6 ] S.I. Kim, M.S. Kim, S.K. Min, C. Lee [J. Appl. Phys. (USA) vol.74 (1993) p.6128-32 ] GJ. Sullivan, M.K. Szwed, RW. Grant [ J. Electron. Mater. (USA) vol.24 (1995) p. 1-4 ] LW. Yang, P.D. Wright, V. Eu,Z.H. Lu, A. Majerfeld [J. Appl. Phys. (USA) vol.72 (1992) p.20635] K.H.Zschauer[/«5r.P^. Conf.Ser. (UK) vol.17 (1973) p.3-10] G. Weimann [ Phys. Status Solidi A (Germany) vol.53 (1979) p.K173-6 ] C.E.C. Wood, D. DeSimone, K. Singer, GW. Wicks [J. Appl. Phys. (USA) vol.53 (1982) p.4230-5] J.B. Clegg, CT. Foxon, G. Weimann [J. Appl. Phys. (USA) vol.53 (1982) p.4518-20 ] M.E. Sherwin et al [J. Electron. Mater. (USA) vol.23 (1994) p.809-18 ] S.S. Gill, BJ. Selay [ J. Electrochem. Soc. (USA) vol. 133 (1986) p.2590-6 ] LK. Naik [ J. Electrochem. Soc. (USA) vol. 134 (1987) p. 1270-5 ] T. Humer-Hager, P. Zwicknagl [ Jpn. J. Appl. Phys. (Japan) vol.27 (1988) p.428-33 ] SJ. Pearton, K.D. Cummings, G.P. Vella-Coleiro [ J. Appl. Phys. (USA) vol.58 (1985) p.3252-4] H. Ryessel, I. Ruge [ Ion Implantation (John Wiley & Sons, 1986) p.289-304 ]

[41] [42] [43] [44] [45] [46] [47] [48] [49] [50] [51] [52] [53] [54] [55] [56] [57] [58] [59] [60] [61] [62] [63] [64]

D.V. Morgan, F.H. Eisen, A. Ezis [IEEProc. I(UK) vol.128 (1981) p.109-30 ] AJ. Moll, K.M. Yu, W. Walukiewicz, W.L. Hansen, E.E. Haller [ Appl. Phys. Lett. (USA) vol.60 (1992) p.2383-5] Y.K. Yeo, Y.S. Park, F.L. Pedrotti, B.D. Choe [ J. Appl. Phys. (USA) vol.53, (1982) p.6148-53 ] GJ. vanGurp,AH. vanOtnmen,PR. Boudewijn, DP. Oosthoek, M.F.C Willemsen [Appl. Phys. (USA) vol.55 (1984) p.338-46 ] Y.K. Yeo, Y.S. Park, P.W. Yu [ J. Appl. Phys. (USA) vol.50 (1979) p.3274-81 ] O.V. Emel'yanenko, D.N. Nasledov, N.A. Urmanov [ Sov. Phys-Semicond. (USA) vol. 10 (1976) p.72-4] K. Tabatabaie-Alavi, A.N.M. Masum Choudhury, CG. Fonstad [ Appl. Phys. Lett. (USA) vol.43 (1983) p.505-7] A.C.T. Tang, BJ. Sealy, A. Rezazadeh [ Electron. Lett. (UK) vol.25 (1989) p.861-3 ] G. Landgren, W.H. van Berlo [ J. Appl. Phys. (USA) vol.63 (1988) p.2783-6 ] S.G. Liu, S.Y. Narayan [J. Electron. Mater. (USA) vol.13 (1984) p.897-911 ] J. Kasahara, K. Taira, Y. Kato, M. Arai, N. Watanabe [ Jpn. J. Appl. Phys. (Japan) vol.22 (1983) p.L373-5] N. Morris, BJ. Sealy [Nucl. Instrum. and Methods, Phys. Rev. B (Netherlands) vol.39 (1989) p.453-6 ] K.K. Patel, BJ. Sealy [ Appl. Phys. Lett. (USA) vol.48 (1986) p. 1467-9 ] Honglie Shen, Zuyao Zhou, Honglai Xu, Guanqun Xia, Schichang Zou [ Appl. Phys. Lett. (USA) vol.61 (1992) p.2093-5] J.P. Donnelly, FJ. Leonberger, CD. Bolzer [Appl. Phys. Lett. (USA) vol.28 (1976) p.706-8 ] S. Nojima, Y. Kawasaki [ Jpn. J. Appl. Phys. (Japan) vol. 17 (1978) p. 1845-50 ] M.V. McLevige, MJ Helix, K.V. Vaidyanathan, B.G. Streetman [ J. Appl. Phys. (USA) vol.48 (1977) p.3342-6] M.V. Rao, PE. Thompson, H.B. Dietrich, D.S. Simons [ J. Appl. Phys. (USA) vol.67 (1990) p.6165-70] SJ. Pearton, CR. Abernathy [ Appl. Phys. Lett. (USA) vol.55 (1989) p.678-80 ] H. Jiang, R.G. Elliman, J.S. Williams [J. Electron. Mater. (USA) vol.23 (1994) p.391-6 ] Y. Yuba, K. Gamo, K. Masuda, S. Namba [ Jpn. J. Appl. Phys. (Japan) vol. 13 (1974) p.641-4 ] S.S. Kular, BJ. Sealy, K.G. Stephens [ Electron. Lett. (UK) vol. 14 (1978) p.2-4 ] A. J. Moll, J.W. Ager III, Kin Man Yu, W. Walukiewicz, E.E. Heller [ J. Appl. Phys. (USA) vol.74 (1993) p.7118-23] B.D. Choe, Y.K. Yeo, Y.S. Park [ J. Appl. Phys. (USA) vol.51 (1980) p.4742-6 ]

3.2

Hole mobility in GaAs5 temperature dependence A.R. Adams and V.A.Wilkinson July 1995

The low field Hall mobility of holes in high purity GaAs has been measured as a function of temperature by Hill [1], by Mears and Stradling [2], and by Zschauer [3] and the data are shown in FIGURE 1. They observed a room temperature mobility of 400 cm2/ Vs at 300 K rising to about 9000 cm2/ Vs at 77 K. In his excellent review of the mobility of holes in III-V compounds Wiley [4] concludes that a temperature dependence of the form H(T) = 400(300/T) 23 cm 2 /Vs

(1)

Hall mobility (cm2/Vs)

gives the best agreement with experiment in agreement with Rosi et al [5].

Temperature (K) FIGURE 1. The temperature dependence of the Hall mobility for high-purity p-GaAs. (Reproduced from reference [4] with the permission of Academic Press, original data from references [1-3].)

Due to the complex nature of the valence band a full theoretical treatment including band warping, non-parabolicity and the presence of light holes is very difficult. Wiley concludes however, that a reasonable fit to experiment can be achieved assuming non-polar optical and acoustic deformation potential scattering, and scattering by polar optical phonons.

Adding the influence of impurities and the temperature dependence of ionized impurity scattering, given by the Brooks-Herring theory assuming Matthiessen's rule, Blakemore [6] concludes that, close to room temperature, the Hall mobility may be written in the form (IH(T)

= l/(2.5xlO"3 x (T/300)23 + 4 x K)"21 N1 (300/T)15) cm2/ Vs

(2)

where N1 is the concentration of immobile charged scattering centres («p0, the free hole density) for weakly compensated material. The 'drift' or 'conductivity' mobility, \iD, which is normally required in device analysis, is related to the Hall mobility, (iH, by the expression HD=HJA

(3)

where r is the Hall scattering factor. Mears and Stradling [2] measured the apparent change in \iH caused by increasing the magnetic field to 4T. Assuming they reached the high field limit where r approaches 1.00 their results give values of r « 1.25 at 300 K and r « 1.5 at 77 K for Hall mobility measurements at low magnetic fields. We can therefore conclude that the drift mobility of holes in GaAs rises from HD = 320 cm2/ Vs at 300 K to ILi0 = 6000 cm 2 /Vs at 77 K This is less fast than the Hall mobility and writing HD = 320 (300/T)22 cm2/ Vs

(4)

appears appropriate for pure material. In heavily doped GaAs the minority carrier mobility of holes can be very different from that found for majority carriers. Lovejoy et al [7] studying the temperature dependence of minority holes in n+ GaAs report a 1/T dependence with a mobility of 235 cm2/ Vs at 295 K rising to 1015 cm2/ Vs at 83 K. REFERENCES [1] [2] [3] [4] [5] [6] [7]

D.E. Hill [J Appl. Phys. (USA) vol.41 (1970) p.1815 ] A.L. Mears, R.A. Stradling [J. Phys. C (UK) vol.4 no.l (1971) p. 122-6 ] K.H. Zschauer [Proa 4th Int. Symp. GaAs and related compounds, Boulder, Co., USA, Sept 1972 (Inst. Phys., London, UK, 1973) p.3-10 ] J.D. Wiley [ Semicond Semimet. (USA) vol. 10, Eds R.K. Willardson, A.C. Beer (Academic, New York, 1975) ch. l,p.91] F.D. Rosi, D. Meyerhofer, RV. Jensen [ J Appl Phys (USA) vol.31 (196) p. 1105 ] J.S. Blakemore [J Appl. Phys. (USA) vol.53 no. 10 (1982) p.R123-81 ] M.L. Lovejoy, M.R. Melloch, M.S. Lundstrom [ J. Electron. Mater. (USA) vol.23, no.7 (1994) p.669 ]

3.3

Hole mobility in GaAs, pressure dependence A.R. Adams May 1990

Adams et al [1] observed that the Hall mobility of holes in p-type GaAs grown by LPE on semiinsulating substrates increases linearly with pressure at a rate in the range 0-0.8 GPa given by: d [ M h /^ 0 ]/CJP = O^l(GPa)- 1

(1)

where |i h is the hole mobility at pressure P and \iho is the hole mobility at atmospheric pressure. This increase contrasts with the observed decrease in electron mobility with pressure [2]. Analysis of the hole mobility is difficult due to the complexities of the valence band. The light hole effective mass will increase with pressure in a similar way to the electron effective mass. However, because of its relatively large density of states, most of the holes are in the heavy hole band and therefore changes in the light hole band have little effect on the average hole mobility. The heavy hole effective mass itself is relatively insensitive to pressure. The dominant effects therefore appear to be: (a)

The energy of the polar phonons increases with pressure. This decreases their density and hence decreases polar phonon scattering.

(b)

An increase in the density of GaAs with pressure which decreases the acoustic phonon scattering.

Combining these effects using Matthiessen's rule, Adams et al [1] were able to obtain reasonable agreement with experiment. REFERENCES [1] [2]

A.R. Adams,L.G. Shantharama [PhysicaB&C(Netherlands) vol. 139/140 (1986)p.419-422] D. Lancefield, A.R. Adams, BJ. Gunney [Appl Phys. Lett. (USA) vol.45 no. 10 (1984) p. 1121-3 ]

3.4

Carbon doping of GaAs SA. Stockman August 1995

A

INTRODUCTION

A decade ago, carbon in epitaxial III-V materials was commonly viewed as an undesired contaminant. Today, however, carbon has gained wide acceptance as a preferred p-type dopant in GaAs, AlGaAs, and InGaAs. The primary motivation for carbon-doping of GaAs is that C has a very low atomic diffusion coefficient relative to the column II acceptors Be, Zn, and Mg. In addition, very high p-type doping levels (p > 1020 cm"3) may be easily achieved while maintaining excellent material quality. These properties make C an ideal choice for doping in many modern electronic and optoelectronic device applications, where precise control over dopant profiles is critical. This Datareview is intended as a review of the development and current status of carbon doping of GaAs. The following sections highlight the epitaxial growth, electrical properties, and thermal stability of carbon-doped GaAs, together with a discussion of device applications. Other relevant issues briefly reviewed here include hydrogen passivation of carbon acceptors and the extension of carbon doping to other III-V materials. B

EPITAXIAL GROWTH

Carbon is an important residual impurity in all types of bulk and epitaxial growth of GaAs. Here we focus on the intentional introduction of carbon in GaAs and AlxGa1^As grown by metalorganic chemical vapour deposition (MOCVD) and molecular beam epitaxy (MBE). Bl

MOCVD

In early studies of the growth of GaAs by MOCVD, the purity of the films was limited by the purity of the Ga and As precursors [1], typically trimethylgallium (TMGa, Ga(CH3)3) and arsine (AsH3). These impurities included residual donors and acceptors such as Si, Ge, and Zn. Carbon was also identified as one of the dominant residual acceptor impurities in GaAs, with acceptor concentrations NA ranging from 1014 to 1016 cm"3. Early work focussed on the role of hydrocarbon impurities in the TMGa as a potential source of carbon [2]. However, as the source material purity was improved subsequent studies identified the methyl radicals of TMGa as the source of carbon contamination [3,4]. The unintentional incorporation of C from TMGa during MOCVD growth of GaAs and AlGaAs has been studied thoroughly by a number of researchers [5,6]. The incorporation of C is a function of substrate temperature, V/III (AsH/TMGa) ratio, substrate orientation, and reactor pressure. The study of these effects has led to an improved understanding of the decomposition of TMGa and the epitaxial growth mechanism, as well as the carbon incorporation mechanism. In simplified terms, the growth is a reaction between Ga(CH3)3 and AsH3 to form GaAs and CH4, with the highly stable CH4 being removed in the gas phase. Carbon incorporation results from an incomplete reaction. For example, a low V/III ratio leads to a shortage of atomic H necessary

for formation of CH4, resulting in an excess of surface CH3 which further decomposes and leaves C to incorporate on an As site. Carbon incorporation is also insensitive to the substitution of He or N2 for H2 as a carrier gas, and the addition of CH4 to the growth ambient [7]. Very low carbon incorporation (< 1013 cm"3) is observed when triethylgallium (TEGa, Ga(C2H5)3) is used as the Ga source, due to differences in the TEGa decomposition route [5]. The differences in growth chemistry between TMGa and TEGa may be exploited to achieve heavy C-doping in GaAs grown by atomic layer epitaxy (ALE) and associated MOCVD-based techniques in which the surface chemistry is carefully controlled [8]. The unintentional incorporation of C acceptors is much more efficient in AlxGa1^As than in GaAs when TMGa and/or TMAl are used as sources [9]. The carbon incorporation originates from the methyl precursors, and increases superlinearly with Al content (x). In fact, a carbon concentration and free hole concentration well in excess of 1 x 1018 cm"3 can be achieved for x > 0.75 without intentional addition of a carbon doping source [9,10]. The increased carbon incorporation efficiency in AlGaAs can be explained by the fact that the Al-C bond is significantly stronger that that of Ga-C, leading to a reduced C desorption rate. The average Ga-CH3 bond strength in TMGa is -59 kcal/mol, compared with -66 kcal/mol for the Al-CH3 bond in TMAl [H]. Trimethylarsenic (TMAs, As(CH3)3) was originally intended as a safer alternative to highly toxic AsH3 in the growth of GaAs, but a high degree of C contamination was commonly observed [12]. This fact was exploited by Kuech et al [13] who used TMGa and a mixture OfAsH3 and TMAs to achieve intentional, controlled p-type carbon doping of GaAs and AlGaAs. They demonstrated high electrical activation for carbon levels of greater than 1 x 1019 cm"3, together with a low deep level trap concentration and abrupt doping profiles. Controlled C doping and abrupt profiles in GaAs and AlGaAs were also achieved by Cunningham et al using CCl4 as the dopant source [14,15]. The carbon incorporation was shown to be strongly dependent on the V/in ratio and growth temperature in a manner qualitatively similar to the trends observed for residual carbon contamination. Carbon doping levels well in excess of 1 x 1020 cm"3 have been achieved [16] and this C-doping technique has become widely used for a variety of device applications. This work was also extended to include C doping using a variety of halomethane sources (CHyX4.y, where X = Cl, Br, or I, and y = 0 to 3) [17], and an incorporation mechanism based on competition between halomethane decomposition and desorption from the crystal surface during growth was proposed [15,17]. Unfortunately, CCl4 will not be available in the near future due to its ozone-depleting properties and the international agreement to ban its use [18]. B2

MBE

Carbon doping of GaAs has in recent years been demonstrated using a variety of methods by the associated high-vacuum epitaxial growth techniques of molecular beam epitaxy (MBE, utilizing solid elemental sources), gas-source molecular beam epitaxy (GSMBE, an extension of MBE where gaseous sources are used for column V or dopant sources), and metalorganic molecular beam epitaxy (MOMBE, another extension of MBE where organometallic precursors are used for column HI sources). Early attempts at growth of high-purity GaAs by MBE were limited in part by the incorporation of C acceptors (-1O14 cm"3) originating from impurities in elemental sources and reactor components. Intentional p-type doping was achieved using elemental Be as

a source, since no convenient elemental C source was available. Growth of GaAs by MOMBE using TMGa and As4 was found to result in extremely efficient C incorporation and high p-type conductivity, with p ~1 x 1018 to 1 x 1019 cm"3 [19]. It was further noted that introduction of H2 or substitution OfAsH3 for elemental arsenic resulted in a significant reduction in C incorporation. In addition, the growth of GaAs using TEGa as a Ga precursor can result in extremely low (< 1014 cm"3) levels of C incorporation in GaAs [20]. These results are suggestive of a C incorporation mechanism similar to that of MOCVD. These observations led to the use of TMGa as a convenient carbon doping source, in combination with TEGa or elemental Ga as the primary Ga source, in MOMBE growth of GaAs and AlGaAs. Weyers et al [21] demonstrated the ability to controllably cover the doping range between p - 1 x 1014 cm"3 (growth using TEGa) and p -1 x 10 20cm "3 (growth using TMGa) by varying the TMGa/TEGa ratio. This technique was later extended by other authors [22,23] to achieve extremely low-resistivity p-type material with p > 1 x io 21 cm"3. Solid-source carbon doping of GaAs was first demonstrated by Malik et al., who used a resistively-heated high-purity graphite filament to evaporate elemental C at a controlled rate [24]. They achieved doping levels in excess of 1 x io 20 cm"3, with good agreement between Hall effect and SIMS data suggesting low compensation. Halomethane sources such as CCl4 and CBr4 have also been used for C-doping by GSMBE [25,26]. A very high doping efficiency was observed for CBr4, due to the low CBr4 decomposition temperature resulting from the weak chemical bond between C and Br. C

ELECTRICAL PROPERTIES

Although low atomic diflEusivity is typically cited as the motivation for the use of C as a p-type dopant, the ultimate usefulness of C-doped GaAs is determined by its electronic transport characteristics. This is especially true for demanding applications such as the base layer of an HBT, where very high p-type doping (p > IO19 cm"3) and good electron transport are simultaneously required. Here we briefly review the majority carrier (hole) and minority carrier (electron) properties of C-doped GaAs. Cl

Majority-Carrier Transport

Temperature-dependent van der Pauw, Hall effect data can be very useful in gaining an understanding of the mechanisms dominating hole transport in p-type semiconductors [27]. Data obtained from C-doped GaAs is summarized in FIGURES 1 and 2 [28]. Samples A-D were grown by MOCVD using CCl4 as a dopant source, and samples E-G were grown by MOMBE with C originating from TMGa. SIMS and room-temperature Hall data are in good agreement, suggesting minimal compensation. The mobility for the least heavily doped sample (sample A, p(300 K) = 2.0 x IO17 cm"3) increases as the temperature is decreased from 300 K to about 80 K, as shown in FIGURE l(a), due to a decrease in phonon scattering. The hole concentration decreases as the temperature is lowered due to freeze-out of holes onto acceptor sites, as can be seen in FIGURE 2(a). As the temperature is lowered below 80 K, the mobility decreases due to increased ionized impurity scattering and the fact that hopping conduction becomes dominant. The freeze-out of carriers and the decrease in mobility result in high resistivity material for T > 80 K as shown in FIGURE 2(b).

Mobility (cm2/Vs)

Temperature (K)

Mobility (crr»2/Vs)

(a)

Temperature (K) (b)

Resistivity (Q»cm)

Hole concentration (cm 3 )

FIGURE 1. Hole mobility as afiinctionof temperature for C-doped GaAs: (a) samples grown by MOCVD (300 K hole concentrations are p(A) = 2.0 x 1017 cm"3, p(B) = 9.7 x 1017 cm"3, p(C) = 4.4 x 1018 cm"3, p(D) = 1.6 x 1019 cm"3), and (b) samples grown by MOMBE(p(E) = 6.5 x 1019crn3,p(F) = 8.4 x 10 19 crn 3 ,p(G)= 1.4 x 1020Cm'3). From reference [28].

FIGURE 2. Hole concentration (a), and resistivity (b) as a function of reciprocal temperature for C-doped GaAs grown by MOCVD (samples A-D) and grown by MOMBE (samples E-G). The dashed line in (a) is the effective density of states for holes in GaAs. From reference [28].

Phonon scattering is also important for temperatures greater than 100 K for sample B (p = 9.7 x 1017 cm"3), and for T < 70 K ionized impurity scattering is the dominant mobility-limiting mechanism. The hole concentration reaches a minimum of -5 x 1017 cm"3 at 70 K, and then appears to increase at lower temperatures, indicating that the carbon acceptors are close enough for impurity band conduction to occur. The mobility for sample B becomes nearly constant for T < 20 K as the sample becomes degenerate. For sample C (p = 4.4 x 1018 cm"3), both phonon scattering and ionized impurity scattering are important at temperatures above 100 K. No carrier freeze-out is observed for this sample, indicating that the impurity band and valence bands have effectively merged to form a continuum of states. This can also be seen in room temperature photoluminescence data, which indicate that the effective bandgap is reduced by about 25 meV as the doping is increased in a very narrow range from 1.5 x 1018 to 4.4 x io18 cm"3. This change in energy gap is approximately equal to the carbon acceptor ionization energy of 26 meV. For T < 100 K, the sample is degenerate, and the ionized impurity scattering-limited mobility becomes nearly constant. This 'degenerate conduction' where the ionized impurity scattering is independent of T, is similar to the behaviour seen in metals. Degenerate conduction is dominant over the entire temperature range studied when p > 1019 cm"3, as observed for sample D (p = 1.6 x io 19 cm"3). At low temperatures, the mobility is dominated by ionized impurity scattering and is independent of temperature. At higher temperatures, the total mobility may be roughly estimated by combining the effects of ionized impurity and phonon scattering according to Matthiessen's rule to obtain l/^i * 1/Hn + 1/^t(T), where |i J T ) * Ml(297K/T)a

(1)

Through variable temperature Hall effect analysis of high-purity p-type MBE-grown GaAs where C is the dominant residual acceptor, it has been determined that (I1 -435 cm2/ Vs and a -2.37 [28]. An attempt was made to fit the data from sample D using the measured mobility at 4.2 K as (Li11, and the values of (I1 and a given above. The calculated mobility was found to be about 15% higher than that measured between 200 K and 300 K. FIGURE l(b) shows the Hall mobility as a function of temperature for several C-doped GaAs layers grown by MOMBE (samples E, F, and G). Carrier freeze-out is not observed in these highly degenerate samples, as seen in FIGURE 2(a). This behaviour is similar to that observed in the most heavily doped MOCVD-grown sample. The expression for total mobility described above shows good agreement with the measured data when using a ~1.5 (nondegenerate optical and acoustic phonon scattering) and values for ^ 1 that range from 250 cm2/ Vs when p = 1 x io 19 cm"3 to 150 cm2/ Vs when p = 1.4 * IO20 cm"3. While Matthiesen's rule is expected to provide only a crude approximation to the total mobility, the trends described above demonstrate that the conduction is dominated by heavily screened ionized impurity scattering at low temperatures and is influenced by phonon scattering at higher temperatures. A comparison of room temperature hole mobilities obtained with C, Zn, and Be doping is shown in FIGURE 3. The solid line represents the best fit to the data taken for C-doped GaAs as part of the study described above [28]. Results from other studies of C-doped GaAs are in close agreement with this curve for layers grown by either MOCVD or MOMBE [16,22,29], indicating that the maximum achievable mobilities for the two techniques are similar. The hole mobilities

C MOCVD Enquist C MOCVD Hanna

Mobility (cm2/Vs)

C MOMBE Konagai Be MBE Konagai Be GSMBE this study Zn MOCVD Glew

Hole concentration (cm"3)

FIGURE 3. The 300 K hole mobility as a function of 300 K hole concentration for p-type GaAs grown by several techniques. The solid curve represents the fit to data taken for C-doped GaAs in [28].

reported in the literature for Be-doped GaAs grown by MBE vary somewhat, but they are always at least 20% lower than the mobilities attainable using C for p > 2 x 1018 cm"3. Similarly, Zndoped GaAs grown by MOCVD exhibits hole mobilities which are nearly equal to those using Be and C for p < 2 x 1018 cm"3, but are also significantly lower than for C-doped GaAs when p > 2 x 1018 cm"3. This may be related to higher compensation in GaAs heavily doped with Be or Zn. C2

Minority Carrier Transport

The electron lifetime, Tn, in heavily-doped p-type GaAs is typically sensitive to the specific growth technique and conditions used, due to the fact that non-radiative recombination centres (deeplevel defects and impurities) often play an important role. For p > 1019 cm"3, radiative recombination, non-radiative recombination via defects, and Auger recombination all play a role, resulting in Xn of less than 1 ns. These short lifetimes also make transport characteristics difficult to determine accurately. A wide variety of techniques have been used to measure Tn, |Hn (electron mobility, related to diffusivity Dn by Dn/jin = kT/q), and Ln (electron diffusion length, Ln = (DnTn)172) in p-type GaAs, and several of these methods have been recently applied to C-doped material. The first measurement OfLn in heavily C-doped GaAs was reported by Saito et al for MOMBEgrown material with p = 5 x 1019 cm"3 [30]. The electron beam induced current (EBIC) and optical collection efficiency for a p-n junction were used to estimate L n - 0.3 jxm. This value is in good agreement with data for Be-doped GaAs. The current gain of AlGaAs/GaAs HBTs with a C-doped base (p = 2.4 x 1019 cm"3) was also used to estimate Tn - 0.18 ns, in agreement with data for devices with a Be-doped base [31].

Time-resolved photoluminescence (TRPL) measurements on heavily C-doped GaAs layers using a double-heterostructure for carrier confinement have been performed to obtain a more direct measure of Tn. A minority carrier lifetime of 0.09 ns was measured for p = 5 x 1019 cm"3 in GaAs grown by MOCVD [32]. The carbon in this layer originated from TMGa, and doping was achieved by growing at a V/III ratio near unity. TRPL measurements for MOMBE-grown layers have indicated long lifetimes for carbon-doping levels as high as 7 x 1019 cm"3, where Tn = 0.08 ns [33]. However, Tn degraded rapidly for further increasing doping levels, with Tn < 3 ps for p > 2 x 1020cm"3. Extensive characterization of minority carrier transport in heavily C-doped GaAs using the zerofield time-of-flight (ZFTOF) technique has been reported by Colomb et al [34] This technique, which consists of measuring the transient photovoltage generated in a diode illuminated by a picosecond laser pulse, may be used to estimate both Tn and Dn (and thus |i n and Ln may be calculated). A summary of data for Tn in p-GaAs measured by ZFTOF, together with other techniques, is shown in FIGURE 4. Significant scatter in the reported data is evident in FIGURE 4, but no clear difference between C, Be, and Zn emerges for 1018 < p < 1020 cm"3. The electron mobility is also comparable for p-type GaAs doped with C, Be, and Zn, and increases substantially as the temperature is decreased to 77 K. A summary of electron diffusion length data for p-GaAs is shown in FIGURE 5, and good agreement between C, Be, and Zn is again evident [34]. Two potentially important exceptions have been noted. One exception is C-doped GaAs grown by MBE using graphite as the doping source. In this case, an estimate for Tn of only 0.01 ns and Ln < 0.1 urn was made for p = 1.5 x 1019 cm"3. Another observation is that while Tn is not adversely affected by the presence of hydrogen in C-doped GaAs grown by MOCVD, \in may be significantly reduced, resulting in a short diffusion length Ln [34,35].

Electron Lifetime (ns)

C This Work (U of I) C This Work (Epitronics) Be This Work (HP)

m

C Ashizawa et al. C Makimoto et al. C Benchimol et al. Be Benchimol et al. Be Love joy et.al. Be Tiwari et.al. Be lto et.al. (700 C) Be lto et.al. (550 C) Ge Casey et al. Zn Ohkubo et al. Zn lkeda et al. Casey et al. (theoretical)

Hole Concentration (cm'3) FIGURE 4. Minority electron lifetime data available in the literature for p-type GaAs. From [34].

Diffusion Length (^ m)

C This Work (U of I) C This Work (Epitronics) Be This Work (HP) C Saito et.al. Be Nathan et.al. Be Tiwari et.al. Be Lievin et.al. Be Lovejoy et.al. Ge Casey et al. Zn Ohkubo et al.

Hole Concentration (cm3) FIGURE 5. Electron diffusion length data available in the literature for p-type GaAs. From [34].

D

STABILITY OF CARBON-DOPED GaAs

Carbon is an amphoteric dopant in the III-V materials, meaning that it could act as an acceptor if incorporated substitutionally on the column V (As) sublattice or as a donor if it resides on the column HI (Ga) sublattice. This raises a series of issues regarding compensation and stability in C-doped GaAs. In addition, the atomic diffusivity of C in GaAs is lower than for column II acceptors, which have a different diffusion mechanism. These topics and other related issues are described in this section. Dl

Lattice Site Location

Although C is potentially amphoteric in GaAs, no clear experimental data showing evidence of C incorporation on a Ga site has been reported. Residual carbon incorporation in bulk and epitaxial GaAs grown by a variety of techniques always shows acceptor behaviour [36,37]. In GaAs intentionally doped with C by MOCVD or MBE-based techniques, good agreement between p (as measured by van der Pauw Hall effect or C-V) and [C] (as measured by SIMS) is in general observed for [C] < 1 * 1019 cm"3. However, a discrepancy between [C] and p sometimes exists for [C] > 1019 cm"3 [16,38]. A number of possible explanations, including self-compensation by substitutional or interstitial C donors, compensation by donor-like GaAs point defects, and C precipitation have been proposed. However, no direct evidence of C donors or C precipitates has been reported for as-grown samples. In addition, de Lyon et al [38] have shown continuously increasing lattice contraction for increasing C content for [C] well above 1020 cm"3, in agreement with Vegard's law assuming a C covalent bonding radius of 0.77 A. This suggests that the C is entirely substitutional. It is now known, as described below, that unintentional hydrogen passivation of C acceptors is the

dominant mode of C deactivation when the GaAs growth takes place in an H-containing ambient. Direct evidence for interstitial C in MOCVD-grown GaAs with p > 5 x 1019 cm"3 has also been reported by Hofler and Hsieh [39], who used nuclear reaction analysis in conjunction with channelling Rutherford backscattering spectrometry. D2

Carbon Diffusion in GaAs

The primary advantage of C as a p-type dopant in GaAs is its low thermal diffusion coefficient, Dc, which allows growth of structures with abrupt and stable doping profiles. A summary of reported values for Dc is given in TABLE 1 [8,13,23,30,40-42]. The apparent discrepancies between the various reports are probably due to subtle differences in crystal growth and annealing conditions, as well as experimental error resulting from the fact that the C redistribution is minimal and is difficult to quantify. However, all reports are consistent in that they report values for difliisivity of C which are a factor of 10 to 100 lower than for the column II acceptors Be, Zn, and Mg. Cunningham et al [40] observed a significant background doping dependence for C diffusion, and also reported a weak dependence on surface encapsulation and ambient As4 pressure. They speculated that the C diffusion mechanism involved diffusion on the As sublattice via As vacancies incorporated during epitaxial growth. Later work by You et al [42] demonstrated a strong dependence of Dc on the ambient As4 pressure, suggesting that C diffusion proceeds via an interstitial-substitutional (kickout) mechanism involving interstitial As. The rate of Al-Ga interdiffusion on the column III sublattice in AlGaAs/GaAs superlattices has also been reported to be strongly affected by the background C doping level [42,43]. D3

Thermal Stability

High temperature processing steps ranging from contact alloying (300 - 5000C) to impurity induced layer disordering and epitaxial regrowth (600 - 9000C) are necessary for fabrication of many electronic and optoelectronic devices. In addition, the fact that C has a low atomic diffiisivity makes it an especially attractive dopant for devices which require particularly severe thermal treatments in processing. Thus the ability of heavily C-doped epitaxial layers to maintain desirable electrical characteristics such as low resistivity and long electron diffusion lengths throughout high-temperature post-growth treatments is critical. Thermal limitations of highly C-doped GaAs have been studied in detail by a number of researchers. Abernathy et al [23] reported that in MOMBE-grown GaAs with [C] = 2 x 1020 cm"3, a 9000C, 30 s anneal resulted in a decrease in the hole concentration by a factor of three, with a decrease in the net strain in the layer. They speculated that site-switching of C from As to Ga sites or C precipitation could be responsible. However, subsequent infrared absorption studies failed to detect local vibrational modes associated with C on a Ga site [44]. Similar work on MOCVD-grown GaAs [45,46] showed that reductions in p, |n, lattice mismatch, and photoluminescence intensity resulted from thermal annealing above 7000C only if [C] > 6 x 1019 cm"3, and additional models based on C switching to interstitial sites or the formation of misfit dislocations due to strain were proposed [45]. However, addition of In to eliminate the strain caused by heavy C doping did not result in improved stability [47]. A detailed study of annealing behaviour of MOCVD-grown GaAs by Hofler et al [41] showed that the concentration of non-

TABLE 1. Data for the diffusion coefficient of C in GaAsfromthe literature. Growth Technique

Carbon Source

Carbon Concentration

Background Doping

Annealing Conditions

Dc(cm7s)

Reference

MOCVD

TMGa, TMAl

9.5x 10 ll cm* 2 (sheet concentration)

undoped

1 hr, 8000C

2xlO' 16

Kobayashi et al [8]

MOCVD

TMAs

undoped

1 hr, 9200C, AsH3 ambient

<; IxIO" 16

Kuechetal[13]

undoped

4hr,900°C

6xlO" 15

Saito et al [30]

undoped

24 hr, 8250C, various ambients

IxIO' 16 2.3xl0" 16
Cunningham et al [40]

19

2xl0 cm" MOMBE

TMGa

MOCVD

CCl4

3

lxl0 1 9 cm 3 5xl0 18 cm- 3

n+ MOMBE

TMGa

3xl0 20 cm- 3

undoped

30 s, 9000C

MOCVD

CCl4

1 xlO 19 to 5xl0 20 cm- 3

undoped

36 hr, 825°C

MOCVD

CCl4

3xl0 19 cm- 3

undoped

12 hr, 9600C, As4 ambient 12hr,960°C,Ga-rich 12 hr,825°C, As4 ambient 12hr,825°C,Ga-rich

<; IxIO" 16 5xlO' 16 8xl0'15 <6xlO" 16 1.6xlO' 16 <4xlO" 1 7

Abernathy et al [23] HOfleretal[41] You et al [42]

substitutional C increased after annealing at T > 6000C for samples with [C] > 5 x 1019 cm'3. They suggested that 5 x 1019 cm"3 represented a solid solubility limit for C in GaAs, and that C could move into isolated interstitial sites or precipitates. Direct evidence of C precipitates has been observed in the Raman spectra of GaAs with [C] = 6 x 1020 cm"3 after annealing at 8500C, suggesting that precipitate formation may be the dominant mechanism of degradation upon thermal annealing [48]. E

DEVICE APPLICATIONS

Over the past five years, carbon has become widely used in industry as a p-type dopant in a number of high-performance electronic and optoelectronic device applications. The proliferation of heterojunction bipolar transistor (HBT) technology into high-power and microwave electronics, in particular, has benefitted greatly from the development of C as an alternative p-type dopant to Be and Zn. Some of the motivating factors and key issues in carbon doping for HBTs and other device applications are briefly reviewed here. El

HBTs

In GaAs-based HBTs, the p-type base region is made thin (< 1000 A) to keep the base transit time small, and heavy doping (p > 1019 cm"3) is required to keep the external base resistance low. The cutoff frequency for unity current gain (ft) is often limited in part by the base transit time, while the maximum frequency of oscillation for unity power gain (^12x) is typically limited by the Rt,Cc time constant, where R4, is the external base resistance and Cc is the base-collector junction capacitance [49]. Thus, the ability to grow thin p-type layers of very low resistivity is important to the HBT process, and is critical in determining the high-frequency performance of such devices. Other considerations for the base region include the minority carrier transport across the base and control over the location and abruptness of the emitter-base junction. In MBE-grown HBTs where Be is used as the base dopant, control over the Be profile is one of the primary challenges [50]. Displacement of the emitter-base junction into the wide-gap emitter (typically Al0 3Ga0 7As or In 0 5Ga 0 5P) due to diffusion of the base dopant during growth or subsequent thermal processing can reduce the emitter injection efficiency. In addition, the primary mode of device degradation involves Be redistribution during device operation at high current density [51]. The diffusion of Zn and Mg is so rapid and difficult to control that MOCVD growth of reliable high-performance HBTs using Zn or Mg as the base dopant is widely considered impractical. The first MBE-grown HBTs utilizing C as a base dopant were grown by Malik et al using a graphite filament as the C source [24]. They pointed out that the low diffiisivity of C at high doping levels (p = 1 x 1019 cm"3) resulted in negligible movement in the emitter-base junction during high-temperature growth of the AlGaAs emitter, in contrast to Be-doped devices. Makimoto et al [31] grew AlGaAs/GaAs HBTs by MOCVD using C from TMGa to dope the base (p = 2.4 x 1019 cm"3). They also demonstrated excellent DC device characteristics and control over junction position which is difficult to achieve using Zn or Mg by MOCVD. The use of TMGa and CCl4 for C-doping in HBTs grown by MOMBE and MOCVD [52], respectively, soon became popular because high doping and excellent material quality could be easily achieved under relatively standard growth conditions. Hanson et al [53] demonstrated that

a C-doping setback of only 15 - 25 A was necessary to accommodate the CCl4 memory effect and redistribution in MOCVD-grown InGaP/GaAs HBTs with p = 2.5 x 1019 cm"3 in the base. Stateof-the-art microwave device and integrated circuit results have been achieved in both the AlGaAs/GaAs and InGaP/GaAs materials systems using C as the base dopant, with p as high as 1020 cm"3. Today, some MBE-based HBT manufacturers continue to rely on Be as a base dopant while others have switched to C. In addition, MOCVD-based production has now become practical due to the development of C doping as an alternative to Zn and Mg. Carbon-doped HBTs have been reported to exhibit improved device reliability compared to Beand Zn-doped devices [51,54,55]. This has been attributed to a significantly lower degree of recombination-enhanced dopant migration in C-doped HBTs, resulting in more stable junction position and injection characteristics. However, recent reports suggest that other issues unique to C-doped HBTs, such as unintentional hydrogen passivation, may affect the operating stability of HBTs under high current stress [56]. It has also been reported that In co-doping to eliminate the strain in the C-doped base may result in improved device reliability [57]. E2

Other Devices

The extremely high doping levels that can be attained make C-doped GaAs an ideal choice for low-resistance and non-alloyed p-type ohmic contacts to a host of III-V semiconductor devices. Non-alloyed Ti/Pt/Au contacts to GaAs with p = 7 x 1019 cm"3 exhibited a specific contact resistivity of 4 x 10"7 Q cm2, [16] while a contact resistance of 7 x 10"8Q cm2 has been measured for non-alloyed Mo/Au on C-doped GaAs with p = 1 x io 21 cm"3 [22]. A shallow C implant may also be used to create a highly C-doped contact region, though incomplete electrical activation of C limits the achievable doping density to the mid 1019 cm"3 range [58]. Carbon-doping of AlGaAs is also useful in a number of device applications. In fact, the first demonstration of C-doping in a device structure was the p-doping of the Al075Ga0 25 As confining layer of an MOCVD-grown quantum well laser structure using C incorporation from methyl precursors [10]. The low difflxsivity of C makes it the preferred choice for devices requiring ptype delta-doping, and excellent hole transport characteristics have been observed in modulation doped heterostructures [59]. Another important application for C is the p-type AlGaAs/GaAs distributed Bragg reflector (DBR) mirrors commonly used in vertical cavity surface emitting lasers (VCSELs). Tight control over the p-type doping is required in order to minimize the series resistance of these structures. Kopf et al [60] have shown that the problems of dopant segregation and diffusion encountered using Be can be solved using C in MBE-grown structures. F

HYDROGEN PASSIVATION

It is well established that shallow donors and acceptors in GaAs, including C, can be passivated by hydrogen [61]. The solubility and diffixsivity of H in GaAs are highest in heavily-doped p-type layers and the efficiency of acceptor passivation is dependent on the acceptor species, with C being the most efficiently passivated [62]. In addition, unintentional incorporation of hydrogen in III-V materials during crystal growth and subsequent device processing has been reported in InP and other DI-V materials in which dopant-H complexes are efficiently formed [63-65]. These factors combine to make the case of heavily C-doped GaAs (p > 1019 cm"3) for device applications one in which unintentional H incorporation during epitaxial growth is an especially important issue.

Acceptor passivation due to incorporation of hydrogen during growth of C-doped GaAs was first reported for layers grown by MOMBE [66]. Infrared absorption measurements performed on GaAs grown by MOMBE using TMGa and AsH3 showed evidence of stable H-C complexes [67]. This work was followed by several reports of H in C-doped GaAs grown by MOCVD [32,41,6870]. These studies showed that a significant fraction (> 50%) of C acceptors could be unintentionally passivated in as-grown GaAs, and that post-growth annealing at temperatures as low as -4000C could be used to reactivate the acceptors such that p = [C]. It was speculated that H was incorporated from TMG and/or AsH3 during epitaxial growth. Woodhouse et al reported partial passivation of C-doped GaAs upon annealing in an AsH3ZHe ambient and speculated that H could be incorporated into MOCVD-grown GaAs from AsH3 or H2 during post-growth cooling [71].

Concentration (atoms/cm )

The importance of the cooling ambient in the post-MOCVD-growth cooling stage was investigated by Stockman et al. [72,73], who showed that the final degree of H passivation was determined primarily by the amount OfAsH3 present during cooling from -500 0 C to -400 0 C, and by the total C concentration. The location of the C-doped layer in a device structure is also important, as illustrated in FIGURE 6, which shows SIMS depth profiles for an n-p-n-p structure grown at 6250 C as-grown and after annealing at -400 0 C in N2 for 5 min. In the as-grown case (FIGURE 6(a)), [H] is significantly higher in the upper C-doped layer than in the buried C-doped layer. Hydrogen in the buried layer was incorporated during growth, while additional H entered the surface layer during the cool-down. The n-type GaAs layer blocks diffusion of H into the buried p-type layer during the cool-down. This is consistent with the finding that the degree of

Concentration (atoms/cm )

Depth (y.m)

Depth (\im)

FIGURE 6. Atomic concentration of H, C and Si vs. depth in a GaAs n-p-n-p structure (a) as-grown, and (b) annealed at ~ 4000C in N2 for 5 minutes. In the as-grown case, (a), [H] is significantly higher in the upper C-doped layer than in the buried C-doped layer. After annealing, (b), [H] in the upper layer is reduced to near the SIMS resolution limit, while [H] is unchanged in the buried layer. From [73].

passivation is significantly lower in the C-doped base region of InGaP/GaAs HBTs than in identically grown C-doped GaAs layers where the surface is exposed during cool-down [72]. FIGURE 6(b) shows that [H] is reduced in the upper C-doped layer due to out-diffusion of H from the surface during annealing. The H concentration is unchanged, however, in the buried Cdoped layer. Thus, the n-type layer can block out-diffusion of H incorporated during growth. This trapping of H may be related to the built-in fields of the p-n junctions which could inhibit out-diffusion of positively ionized hydrogen (H+). Although both donors and acceptors in GaAs, AlGaAs, and InGaAs are easily passivated by H in certain plasma processing steps, significant unintentional H passivation in a non-plasma environment which contains H appears to be unique to heavily C-doped layers. Recent studies also suggest that H may in some cases behave as a fast-diffusing isolated donor in C-doped GaAs [74]. This susceptibility of C-doped layers to H incorporation has implications not only in epitaxial growth, but also in device processing and reliability [65]. G

CONCLUSION

Carbon acceptor concentrations in GaAs as high as 1019 to 1021 cm"3, well in excess of the equilibrium solid solubility limit, may be achieved using the non-equilibrium epitaxial growth techniques of MOCVD and MBE. These layers also possess excellent majority and minority carrier transport characteristics. Carbon exhibits an atomic diffixsivity in GaAs which is 10 - 100 times lower than for the column II acceptors, although high-temperature processing may still be limited by C precipitation when [C] > 5 x 1019 cm'3. C-doping is now widely used in many device applications, and the development of C-doping has enabled recent advances in HBT technology. Much of the current (and future) work on C-doping revolves around obtaining a better understanding of reliability-related issues, and extending the present understanding of C as a dopant to include the P-, Sb-, and N-based HI-V materials. ACKNOWLEDGEMENTS The author would like to thank Professor G.E. Stillman for guidance and support, as well as friends and former colleagues AW. Hanson, CM. Colomb, M.T. Fresina, N.F. Gardner, QJ. Hartman, S.L. Jackson, and G.E. Hofler for collaborative work on C in III-V materials at the University of Illinois. REFERENCES [1] [2] [3] [4] [5] [6] [7] [8]

See, for example, G.B. Stringfellow [ Organometallic Vapor-Phase Epitaxy (Academic Press, 1989) and references therein ] K.L. Hess, P.D. Dapkus, H.M. Manasevit, T.S. Low, BJ. Skromme, G.E. Stillman [J. Electron. Mater. (USA) vol. 11 (1982) p. 1115 ] T.F. Kuech, E. Veuhoff [ J Cryst. Growth (Netherlands) vol.68 (1984) p. 148 ] J. van de Ven, H.G. Schoot, LJ. Giling [ J Appl. Phys. (USA) vol.60 (1986) p. 1648 ] T.F. Kuech, J.M. Redwing [ J. Cryst. Growth (Netherlands) vol. 145 (1994) p.382 and references therein ] M. Kondo, T. Tanahashi [ J. Cryst. Growth (Netherlands) vol. 145 (1994) p.390 ] T.F. Kuech, GJ. Scilla, F. Cardone [J Cryst. Growth (Netherlands) vol.93 (1988) p.550 ] N. Kobayashi, T. Makimoto, Y. Horikoshi [Appl. Phys. Lett. (USA) vol.50 (1987) p. 1435 ]

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Next Page

[44] [45] [46] [47] [48] [49] [50] [51] [52] [53] [54] [55] [56] [57] [58] [59] [60] [61] [62] [63] [64] [65] [66] [67] [68] [69] [70] [71] [72] [73] [74]

CR. Abemathy, SJ. Pearton, M.O. Manasreh, D.W. Fischer, D.N. Talwar [ Appl. Phys. Lett. (USA) vol.57 (1990) p.294 ] M.C. Hanna, A. Majerfeld, D.M. Szmyd [Appl. Phys. Lett. (USA) vol.59 (1991) p.2001 ] K. Watanabe H. Yamazaki [ Appl. Phys. Lett. (USA) vol.59 (1991) p.434 ] P.M. Enquist [ J. Appl. Phys. (USA) vol.71 (1992) p.704 ] A.J. Moll, E.E. Haller, J.W. Ager III, W. Walukiewicz [ Appl. Phys. Lett. (USA) vol.65 (1994) p. 1145] H. Kroemer [ Proc. IEEE (USA) vol.70 (1982) p. 13 ] D.L. Miller, P.M. Asbeck [J Appl. Phys. (USA) vol.57 (1985) p.1816 ] K.P. Roenker [Microelectron. Reliab. (UK) vol.35 (1995) p.713 and references therein ] B.T. Cunningham, G.E. Stillman, G.S. Jackson [Appl. Phys. Lett. (USA) vol.56 (1990) p.361] A.W. Hanson, S.A. Stockman, G.E. Stillman [IEEEElectron Dev. Lett. (USA) vol. 14 (1993) p.25 ] F. Ren et al [ Appl. Phys. Lett. (USA) vol.59 (1991) p.3613 ] T. Ahmad, A.A. Rezazadeh, S.S. Gill [ Electron. Lett. (UK) vol.29 (1993) p. 1725 ] CR. Abernathy et al [ J. Cryst. Growth (Netherlands) vol. 136 (1994) p. 11 ] T. Nittono, N. Watanabe, H. Ito, H. Sugahara, K. Nagata, O. Nakajima [ Jpn. J. Appl. Phys. (Japan) vol.33 (1994) p.6129] A.J. Moll, J.W. Ager, III K.M. Yu, W. Walukiewicz, E.E. Haller [ J. Appl. Phys. (USA) vol.74 (1993)p.7118] T. Makimoto T. Kobayashi [ Jpn. J. Appl. Phys. (Japan) vol.32 (1993) L1300 ] R.F. Kopf, E.F. Schubert, S.W. Downey, A.B. Emerson [Appl. Phys. Lett. (USA) vol.61 (1992) p. 1820] N. Panet al [Appl. Phys. Lett. (USA) vol.51 (1987) p.596 ] I. Szafranek G.E. Stillman [ J. Appl. Phys. (USA) vol.68 (1990) p.3554 ] S. Cole, J.S. Evans, MJ. Harlow, A.W. Nelson, S. Wong [ Electron. Lett. (UK) vol.24 (1988) p.929] G.R. Antell et al [ Appl. Phys. Lett. (USA) vol.53 (1988) p.758 ] see, for example, SJ. Pearton [ Hydrogen in Compound Semiconductors fTrans Tech Publications, 1994) and references therein ] R. Iga, H. Sugiura, T. Yamada, K. Wada [Appl. Phys. Lett. (USA) vol.55 (1989) p.451 ] D.M. Kozuch, M. Stavola, SJ. Pearton, CR. Abernathy, J. Lopata [ Appl. Phys. Lett. (USA) vol.57 (1990)p.2561] K. Watanabe, H. Yamazaki [ Appl. Phys. Lett. (USA) vol.60 (1992) p.847 ] J. Wagner et al [ Phys. Rev. B (USA) vol.45 (1992) p.9120 ] W.S. Hobson, SJ. Pearton, D.M. Kozuch, M. Stavola [Appl. Phys. Lett. (USA) vol.60 (1992) p.3259] K. Woodhouse, R.C. Newman, TJ. de Lyon, J.M. Woodall, GJ. Scilla, F. Cardone [ Semicond. Sci. Technol. (UK) vol.6 (1991) p.330 ] S.A. Stockman et al [ J. Electron. Mater. (USA) vol.21 (1992) p. 1111 ] S.A. Stockman, A.W. Hanson, S.L. Jackson, J.E. Baker, G.E. Stillman [ Appl. Phys. Lett. (USA) vol.62 (1993) p. 1248] H. Fushimi, K.Wada [ J. Cryst. Growth (Netherlands) vol. 145 (1994) p.420 ]

3.5

Carrier concentration dependence of the hole mobility in GaAs Previous Page

M.C. Hanna and A. Majerfeld July 1995

A

INTRODUCTION

In this Datareview, we present the dependence of the room temperature hole mobility in epitaxial GaAs on carrier concentration for the shallow acceptors C, Be, Zn, Mg and Ge, which are commonly used to produce p-type GaAs. Since the publication of the 2nd Edition of this book [1,2], a large amount of new data has become available, including new dopant atoms for concentrations up to 1021 cm"3 for various growth techniques, as high hole concentrations are routinely used in many devices. There are several existing reviews [3-6] which include discussions of experimental and theoretical aspects of hole mobility in GaAs. For each acceptor species, we give for the first time a comprehensive compilation of the experimental hole mobilities in GaAs which puts into perspective the results obtained by different growth techniques. We also discuss factors that affect the mobility values, such as compensation. By displaying the data in a graphical form, a comparative analysis of the large amount of information existing at the present time can be made. As detailed data for the total acceptor (N^ and donor (ND) concentrations and the Hall r factor are generally not reported, the compiled hole mobilities and concentrations summarized here are the concentrations and mobilities as measured by the Hall effect, which is the most commonly used measurement technique (see [7] for a thorough discussion of the Hall measurement technique). The Hall r factor, which is the ratio between the Hall and drift mobilities, is generally taken to be equal to one (r = 1) due to difficulties in measuring and calculating the Hall factor for hole transport. Theoretical predictions [8-11] of the Hall factor in p-GaAs at 300 K fall within the range of 1.25 to 2.1. This means that the reported Hall mobility will be greater than the hole drift mobility and that the actual hole concentration will be greater than the Hall hole concentration. The Hall hole concentration, p, is usually reported instead of the total acceptor concentration. Compensation by donors or passivation of the acceptors often cause the measured value of p to be less than NA. Compensation can be important, particularly at low and high doping concentrations, and depends upon the growth conditions [12], annealing [13], precipitates [14], etc. The mobility in heavily doped material is determined primarily by ionized impurity scattering, which depends upon the total ionized impurity concentration, N1 = NA + ND. To determine the compensation ratio and the total acceptor concentration in p-type GaAs requires additional characterization such as temperature dependent Hall measurements [8,15,16] and SIMS measurements of the impurity species in the sample. Recent progress [5,9-11,17,18] in the theory of hole transport can provide assistance in extracting the true hole concentration and the compensation ratio from mobility vs. concentration and mobility vs. temperature data. A representative sample of 300 K hole mobilities versus the hole concentration in GaAs produced by a number of growth techniques for the acceptors C, Be, Zn, Mg and Ge are shown in FIGURES 1, 2, 3 and 4 respectively. Included in FIGURE 1 are several data points from high purity undoped p-type GaAs grown by MOVPE [53,54], MBE [55] and LPE [56] in which the dominant acceptor is believed to be C.

B

EXPERIMENTAL DATA

Bl

Carbon Doping

Hole Mobility (cm 2 /Vs)

The use of C as a p-type dopant in GaAs has become widespread in the past five years (see Datareview 3.4 on carbon doping in this book) due to its nearly ideal dopant properties, such as the ability to achieve high carrier concentrations with little compensation and a low diffusion coefficient. Efficient and practical carbon doping sources are available for both metal-organic vapour phase epitaxy (MOVPE) (CCl4 [19-23], CBr4 [24], TMG [25]) and molecular beam epitaxy (MBE), and related techniques such as metal organic MBE (MOMBE) (TMG [26-29], graphite filament [30,31], CCl4 [32], CBr4 [33,34]). Ion implantation [35,36] has also been employed to produce carbon doped p-type GaAs. The most heavily doped p-type GaAs reported to date has been achieved using carbon [29] had a resistivity of p = 1.9 x 10"4 Q cm, and a mobility \i = 22 cm2/ Vs and p = 1.5 * 1021 cm"3.

REF. • O • D V • A • A O + *

Dopant Source

Growth Tech.

[211 CCK APMOVPE [261 TMG MOMBE (451 CCM APMOVPE [31 I Graphite MBE 1291 TMG MOMBE (251 TMG LPMOVPE (201 CCI4 LPMOVPE [221 CCI4 LPMOVPE 1321 CCI4 MBE 1341 CBr4 MBE (231 CCI4 LPMOVPE [53-563 Undoped

CARBON Hole Concentration (cm 3 )

FIGURE 1. Experimental 300 K hole mobilities as a function of hole concentration for carbon doping. For the MOVPE grown material, AP stands for atmospheric pressure, LP stands for low pressure. The solid and dashed lines are the mobilities predicted by the formulas of EQN (1) and (2), respectively, which track the highest reported mobilities for all p-type dopants.

B2

Other Dopants

Beryllium is widely used in MBE [29,31,37,38] and related growth systems as a p-type dopant and has also been successfully used in MOVPE [39,40] to produce p-type GaAs. Zinc is primarily used in MOVPE [41-46] from either trimethylzinc or triethylzinc sources and in the production of bulk p-type GaAs ingots. Germanium is used as an acceptor in liquid phase epitaxy (LPE) [15,47] grown GaAs. Magnesium has been investigated for p-type doping of GaAs in LPE [48], MOCVD [44,49,50] and MBE [51] but it is not commonly used in practice due to its high vapour pressure, large diffusion coefficient and the tendency for Mg compounds to strongly absorb to reactor components and thereby prevent the realization of sharp doping profiles [52].

Hole Mobility (cm 2 /Vs)

REF.

Dopant Source

Growth Tech.

D V

[31J (291

Be Be

MBE MBE

• • O A O •

(39J 1401 (37) (32) (341 (38]

DMBe OEBe Be Be Be Be

MOVPE MOVPE MBE MBE MBE MBE

Hole Concentration (cm 3 )

Hole Mobility (cm2/Vs)

FIGURE 2. Experimental 300 K hole mobilities as a function of hole concentration for beryllium doping. The solid and dashed lines are the mobilities predicted by the formulas of EQN (1) and (2), respectively.

REF. • [45J • 146] • 143] A 14U • t42]

Dopant Source DEZd DEZn DMZn DMZn DEZn

Growth Tech. MOVPE MOVPE MOVPE MOVPE MOVPE

Hole Concentration (cm"3) FIGURE 3. Experimental 300 K hole mobilities as a function of hole concentration for zinc doping. The solid and dashed lines are the mobilities predicted by the formulas of EQN (1) and (2), respectively.

Hole Mobility (cm2/Vs)

REF, O 1511 • 1491 • [44J • 150] O 1481 A [15) V (4 71

Dopant Source Mg Cp2Mg Cp2Mg Cp2Mg Mg Ge Ge

Growth Tech. MBE MOVPE MOVPE MOVPE LPE LPE LPE

MAGNESIUM GERMANIUM Hole Concentration (cm 3 )

FIGURE 4. Experimental 300 K hole mobilities as a function of hole concentration for magnesium and germanium doping. The solid and dashed lines are the mobilities predicted by the formulas of EQN (1) and (2), respectively.

As can be seen in the data of FIGURES 1-4, the hole mobilities for different acceptor species have nearly the same dependence on hole concentration regardless of growth technique or acceptor species. Carbon doping provides the highest mobility in very highly doped material where p > 1 x 1019 cm"3 and also an overall better consistency among several growth techniques. There is a fairly large range of reported mobility values for a given hole concentration, which we assign to differences in the total ionized impurity content of the samples. Several empirical formulas [5759] have been proposed to describe the concentration dependence of the hole mobility. Wiley [3] showed that combining the ionized impurity mobility calculated from the Brooks-Herring formula with a lattice-limited mobility of 400 cm2/ Vs provided a reasonable approximation of the 300 K mobility data. In each figure, we have plotted the mobility given by this approximation, EQN (1), up to p = 5xlO18 cm"3, using an assumed lattice-limited mobility of 450 cm2/ Vs: \i = (1/450 + 1.0xl0"21p(ln(l+b) - b/O+b)))"1

(1)

where b = 3>
(2)

fits the highest reported mobilities for p > 5 x 1018 cm"3 reasonably well. The formulas in EQN (1) and (2) provide an approximate representation of the maximum 300 K hole mobilities that have been reported to date. The experimental mobilities in the range 1019 - 10 20 cm "3 are in reasonable agreement with theoretical computations taking properly into account the interaction between the heavy and light hole bands for impurity scattering [60]. For example, for a Hall hole concentration

of 1020 cm"3 the experimental Hall mobility is around 70 cm 2I Vs, whereas the theoretical value is ~ 100 cm2/ Vs, which is acceptable considering that constant effective masses for the nonparabolic light and heavy hole bands were used [60].

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3.6

Minority hole mobility in GaAs M.L. Lovejoy, M.R. Melloch and M.S. Lundstrom September 1995

A

INTRODUCTION

Minority carrier mobility is a critical parameter for accurate device modelling of all bipolar devices including heterojunction bipolar transistors (HBTs)3 solar cells and photodiodes. Most work in characterizing minority mobilities in GaAs has focused on minority electron mobility [1] because of the high electron mobility in GaAs and resulting high-performance devices. Devices with performance limited by hole transport have been expected to be low performance due to the high effective mass, hence low mobility, of majority holes in GaAs. However, in heavily-doped n-GaAs which is found in many devices including HBTs, the 300 K minority hole mobility has been shown to be higher than majority hole mobility in comparably doped p-GaAs which significantly impacts device design. This is exemplified by recent reports of Pnp AlGaAs/GaAs HBTs [2,3] with performance that is better than predictions based on majority hole mobility. In addition to very different doping dependences of minority/majority hole mobility in n-GaAs, very different temperature dependence of minority hole mobility has been shown as compared to majority hole mobility. In this Datareview the minority hole mobility dependences on both doping level and temperature are reviewed. B

THEORETICAL RESULTS

The scattering mechanisms that are present in a material system determine the minority and majority carrier mobilities. Many mechanisms have been characterized within the framework of Fermi's Golden Rule and the calculated scattering rates have been used in investigations of the influence of each mechanism on carrier mobility. Important mechanisms for majority carriers in GaAs include ionized impurity scattering, polar and acoustic lattice scattering and carrier-carrier scattering. Scattering rates for most of these mechanisms are documented in many textbooks (see for example [4]). Additional scattering mechanisms must be considered for minority carriers in degenerately doped GaAs. The first mechanism to be discussed is the scattering of minority carriers via a Coulomb interaction with a single majority carrier. The second is the interaction of a minority carrier and the collective charge oscillations (plasmon oscillations) that occur in a degenerate majority carrier gas. Theory for a gaseous plasma has been used to describe these oscillations. A rigorous treatment of minority carrier-single majority carrier scattering includes the majority carrier distribution function and results in an untractable formulation. Walukiewicz et al [5] used an approximation in calculations of electron mobilities in GaAs that was suggested by Ehrenreich [6] where, due to the greater effective mass of holes relative to electrons, the electron scattering was described by a screened impurity scattering rate. This approximation may be appropriate for minority electrons but is not appropriate for minority holes. Lowney and Bennett extended the work of Walukiewicz et al by considering different formulations of screened impurity scattering. They included effects due to plasmon scattering, and included a correction to the approximation

JLX (CITI2/VS)

of treating minority-single majority carrier scattering as an elastic scattering event with all majority carriers [7]. The screened impurity scattering rate calculated with the Born approximation was compared with results using partial wave analysis, which unlike the Born treatment is quantitatively different for scattering by like-charged carriers and scattering by oppositely-charged carriers. As shown in FIGURE 1, the agreement between techniques is reasonable for minority holes; the agreement is excellent for minority electrons [7]. While Lowney and Bennett found plasmon scattering to be strong for minority electrons in heavily doped p-GaAs it was found to be negligible for minority holes in n-GaAs.

Log N A (cm"3) FIGURE 1. Theoretical calculation of minority-hole mobility in n-GaAs vs. doping for different dominant scattering mechanisms showing that an increase in minority carrier mobility can occur with heavy doping [7].

An important result that was shown by Lowney and Bennett is that the treatment of the minoritysingle majority carrier scattering mechanism as ionized impurity scattering with each majority carrier is not correct because the carriers are fermions obeying the Pauli exclusion principle. In addition to the momentum transfer between the carrier systems, a small energy exchange occurs since the process is an inelastic scattering event. In degenerately doped material where the Fermi level is deep in the majority band, as shown in FIGURE 2, all states below the Fermi energy are filled.

Wavevector

Occupancy

FIGURE 2. Dispersion curve and occupancy diagram showing that majority carriers below the Fermi level in degenerately doped GaAs are forbidden to participate in first order inelastic scattering events since all states within a small energy range are folly occupied (from [10]).

Because the majority carriers in states below the Fermi level do not have states within a small energy range in which to scatter they are forbidden to participate in minority-majority scattering. As shown in FIGURE 1, when this correction is made, the minority hole mobility increases as the material becomes more degenerate since more majority carriers are prohibited from participating in minority carrier scattering. This increasing behaviour contrasts with majority hole mobility which monotonically decreases with increasing doping level. This treatment of single minority-majority carrier interaction is a first approximation of minoritymajority carrier scattering. However, in highly non-equilibrium systems with a large applied field, this treatment does not account for the highly energetic majority carriers. For these conditions, the average drift momentum of majority carriers is directed diametrically opposite to that of minority carrier drift. Majority-minority scattering results in a net momentum transfer to minority carriers that will impede the minority carrier drift; consequently, minority carrier mobilities that are lower than the zero-field mobility can be observed. This effect is expected theoretically and has been observed experimentally for minority electrons in GaAs [8-10]. Recent theoretical studies of p-i-n diode currents show that the effect may be great for silicon devices [11,12]. However, due to the greater effective mass of minority holes in n-GaAs as compared to electrons this effect is expected to be small for minority holes. Nevertheless one must be careful to distinguish between low-field mobility that is related to the diffusivity by the Einstein relation and high-field mobilities which are reduced from the low-field value by minority-majority carrier scattering. Another mechanism that can influence measurements of minority carrier mobilities is photon recycling. Photon recycling is the absorption of a photon emitted by a recombination event occurring in the same semiconductor. This phenomenon has two important effects on minority

carrier transport. First, recombination and subsequent creation of an electron-hole pair at another location effectively represents a transport mechanism in addition to the processes of diffusion and drift. Second, recombination and subsequent absorption effectively enhances minority carrier lifetime beyond that of the radiative lifetime [13,14]. These effects can become important in minority carrier mobility measurements. C

MINORITY HOLE MOBILITY MEASUREMENT TECHNIQUES

In the last few years, several techniques have been employed to measure minority carrier mobility in GaAs. Techniques such as one-sided diode I-V characterization and the photo-Hall effect technique have been used for minority carrier mobility measurements in moderately-doped GaAs. For measurements of degenerately-doped GaAs, techniques that have been applied are the conventional time-of-flight technique, the unity gain transistor cutoff-frequency technique (fT) and the zero-field time-of-flight technique (ZFTOF) [1,10,15,16]. A significant number of research groups have applied these techniques to minority electron measurements; however, only a few of the techniques have been applied to minority hole measurements. Chuang et al applied the one-sided diode I-V characterization technique to the measurement of minority hole mobility in n-GaAs doped to 1.6 x io 16 cm"3 [17]. Slater et al applied the fT technique to the measurement of minority hole mobility in n+-GaAs doped to 4 x io 18 cm"3 [18]. This technique has also been used to measure the temperature dependence of minority electrons in GaAs [19], but it has not been applied to the temperature dependence of minority holes in GaAs. Lovejoy et al used the ZFTOF technique to measure the doping dependence of minority hole mobility in n-GaAs doped from 1 xlO 1 7 cm"3 to 2 x 10 18 cm "3 [20,21] and to measure the temperature dependence of minority hole mobility in heavily doped n-GaAs from 77 K to room temperature [22,23]. The latter three techniques actually measure the minority hole zero-field diffusivity under low-level injection and the low field mobility is calculated with the non-degenerate Einstein relation. D

MINORITY HOLE MOBILITY DOPING DEPENDENCE

Minority hole mobility exhibits a very different doping dependence as compared to majority hole mobility. In FIGURE 3 hole mobilities vs. doping concentration for both minority holes and majority holes are shown. The majority carrier mobility data are Hall mobilities while the minority carrier data are drift mobilities; the two mobilities are related by the Hall factor which approaches unity as the material becomes degenerate [24]. Data at low concentrations are limited but more data are available for degenerately doped n-GaAs. Minority hole mobilities measured with the fT [18] and the ZFTOF [15,20,21] techniques are in good agreement. Below mid-1017 cm"3 the minority mobility is comparable to majority Hall mobility. At higher doping levels the minority mobility exceeds the majority hole mobility. The difference is as great as a factor of two at the highest doping level for which data are available. Also shown in FIGURE 3 is the theoretical curve of Lowney and Bennett [7] which agrees qualitatively with the measured data showing the same trend but underestimates the mobility over the doping range for which measured data are available. This theoretical curve, also shown in FIGURE 1, is from partial wave analysis with an approximation for the Pauli principle included. It is the inclusion of the approximation that causes the mobility to rise above mid-1017 cm"3 and level off at mid-1018 cm"3. This result of high minority hole

Hole Mobility (cm2/V-sec)

ZFTOF fT 1-sided diode Theory Majority Hole

Doping Concentration (cm3) FIGURE 3. Hole mobility vs. doping. Markers are measurements made by diode I-V [17], ZFTOF [20,21] and fT techniques [18]. Also shown are the minority hole mobility theoretical [7] and typical majority Hall mobility data. At high doping levels, the minority hole mobility is nearly twice the majority hole mobility.

mobility has important consequences on device design and optimization because the minority hole mobility is significantly higher than the majority Hall mobility. E

MEVORITY HOLE MOBILITY TEMPERATURE DEPENDENCE

The temperature dependence of minority hole mobility in heavily-doped GaAs also differs greatly from that of majority hole mobility in comparably doped p-GaAs. In FIGURE 4 the temperature dependence of minority hole mobility in n+-GaAs doped to 1.8 x 1018 cm"3 measured from 77 K to 300 K by the ZFTOF technique [22,23] is shown. In the temperature range investigated, the minority hole mobility has an approximately inverse temperature dependence that varied from 1015 cm2/ Vs at 80 K to 235 cm 1I Vs at 300 K. Also shown in FIGURE 4 is the majority electron Hall mobility measured in the n+-GaAs sample in which the minority hole mobility was measured and the majority hole mobility in comparably doped p-GaAs (1.5 x 1018 cm"3). Little temperature dependence is exhibited by either majority carrier Hall mobility. The very different temperature dependent characteristics show that different dominant scattering mechanisms determine the mobilities of the majority carriers and the minority holes. This same trend has been shown in a comparison of measured majority carrier mobilities [25] and the temperature dependence of minority electron mobility [19].

Mobility (cm2/V-sec)

Minority Hole Majority Electron

Majority Hole

Temperature (K) FIGURE 4. Mobility vs.temperaturefor minority holes in 1.8 x 1018 cm"3 n+-GaAs measured with the ZFTOF technique [22,23], for majority electrons in the n+-GaAs and for majority holes in p-GaAs doped to 1.5 * 1018 cm'3 [25]. A 1/T temperature dependence is found for minority hole mobility while majority Hall mobilities are roughly constant.

A possible explanation for the minority carrier mobility temperature dependence is majority carrier freeze-out which would reduce ionized impurity scattering; however, it has been shown that freeze-out is not the origin of the temperature dependence of minority holes [25]. An explanation that has been presented parallels the discussion of minority carrier mobility increase with increased doping due to increased degree of degeneracy (see Section D on theory). The difference is that the increased degree of degeneracy of majority carriers occurs due to decreasing temperature which reduces the number of majority carriers that can participate in scattering events with the minority carriers [25]. This may be the origin of the different temperature dependence of minority as compared to majority carriers. F

CONCLUSION

Minority hole mobility has very different doping and temperature dependences in GaAs. Below mid-1017 cm"3 doped n-GaAs, the minority hole mobility is comparable to the majority mobility while at higher doping levels the minority mobility exceeds the majority hole mobility. At the highest doping level, mid-1018 cm"3, the minority hole mobility is roughly twice that of majority holes. This result of higher minority hole mobility was observed in silicon as well [26]. Theoretical work has shown that degeneracy effects limit the scattering rate of majority electrons with minority holes and may be the origin of the higher minority hole mobility. It is the heavily doped regime that is important for devices such as Pnp-HBTs and accurate mobility data are needed for device design and optimization. Prior to the availability of these measured minority hole data, predictions of Pnp-HBT performance based on majority hole mobility predicted poor performance if difficult fabrication schemes to achieve ballistic transport were not invoked [27]; however, these data show that higher performance is possible due to the higher minority hole mobility.

Minority and majority hole mobility temperature dependences are also quite different in GaAs. Minority hole mobility exhibits a roughly inverse temperature dependence while majority hole mobility in comparably doped GaAs is nearly temperature independent from 77 K to 300 K. Majority carrier freeze-out, which reduces both majority-minority carrier and ionized impurity scattering, has been shown not to be responsible for the inverse temperature minority hole mobility dependence. It has been suggested that the difference in minority and majority carrier mobility temperature dependences is due to an increased degree of degeneracy of minority carriers with decreased temperature, which decreases majority-minority carrier scattering. More work is needed to understand this phenomenon. The impact on device design of these large differences in minority and majority hole mobilities is great. To realize accurate device design and optimization, the correct doping dependences of minority mobilities, as well as the correct temperature dependences, must be incorporated in device modelling tools. ACKNOWLEDGEMENTS This work was supported by the Department of Energy under contract #DE-AC04-94AL85000. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] II1] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21]

E.S. Harmon, M.L. Lovejoy, M.S. Lundstrom, M. R. Melloch [ Datareview in this book: 2.13 Minority electron mobility in doped GaAs ] D.G. Hill, T.S. Kim, H.Q. Tserng [ Dig. 51st Device Research Conf. (USA) (1993) p.IIIA-6 ] D.B. Slater Jr., P.M. Enquist, J.A. Hutchby, A.S. Morris, R J. Trew [ /JEEE Electron Device Lett (USAJvol 15(1994)p.91] M.S. Lundstrom [ Fundamentals of Carrier Transport. (Addison-Wesley, 1990) ] W. Walukiewicz, J. Lagowski, L. Jastrzebski, H.C. Gatos [ J. Appl Phys. (USA) vol.50 (1979) p.5040 ] H. Ehrenreich [ J. Phys. Chem. Solids (USA) vol.8 (1959) p. 130 ] J.R. Lowney, H.S. Bennett [ J. Appl. Phys. (USA) vol.69 (1991) p.7102 ] W.P. Dumke [ Solid-State Electron. (USA) vol.28 (1985) p. 183 ] RA. Hopfel, J. Shah, PA. Wolff, A.C. Gossard [ Phys. Rev. Lett. (USA) vol.56 (1986) p.2736 ] M.L. Lovejoy [ Minority and Majority Carrier Transport Characterization in Compound Semiconductors. (AIP Press of the American Institute of Physics, 1995) ] T.T. Mnatsakanov, B.N. Gresserov, L.I. Pomortseva [ Solid-State Electron. (UK) vol.38 (1995) p.225 ] D.E. Kane, R.M. Swanson [ IEEE Trans. Electron. Devices (USA) vol.40 (1993) p.1496 ] W. van Roosbroeck, W. Shockley [ Phys. Rev. (USA) vol.94 (1954) p.1558 ] O. von Roos [ J. Appl. Phys. (USA) vol.54 (1983) p.1390 ] M.L. Lovejoy [ PhD, Purdue University, 1992) ] M.L. Lovejoy, M.R. Melloch, RK. Ahrenkiel, M.S. Lundstrom [ Solid-State Electron. (USA) vol.35 (1992)p.251] H.L. Chuang, M.E. Klausmeier-Brown, M.R. Melloch, M.S. Lundstrom [J. Appl. Phys. (USA) vol.66 (1989) p.273] D.B. Slater Jr., P.M. Enquist, F.E. Najjar, M.Y. Chen, J.A. Hutchby, A.S. Morris, RJ. Trew [ IEEE Electron. Device Lett. (USA) vol.12 (1991) p.54 ] K. Beyzavi, K. Lee, D.M. Kim, M.I. Nathan, K. Wrenner, S.L. Wright [ Appl. Phys. Lett. (USA) vol.58 (1991) p. 1268] M.L. Lovejoy, B.M. Keyes, M.R. Melloch, RK Ahrenkiel, M.S. Lundstrom [ Proc. 8th Int. Symp. On GaAs and Related Compounds, Seattle, WA, USA, 9-12 Sept. 1991 (IOP, Bristol UK, 1992) ] M.L. Lovejoy, M.R. Melloch, M.S. Lundstrom, RK. Ahrenkiel [Appl. Phys. Lett. (USA) vol.61 (1992)p.2683]

[22] [23] [24] [25] [26] [27]

M.L. Lovejqy, M.R. Melloch, M.S. Lundstrom, B.M. Keyes, R.K. Ahrenkiel [ in Proc. Electronic Materials Conf., Santa Barbara, CA, USA (The Minerals, Metals and Materials Society, 1993) ] M.L. Lovejoy, M.R. Melloch, M.S. Lundstrom, B.M. Keyes, R K. Ahrenkiel [ J. Electron. Mater. (USA) vol.23 (1994) p.6693 ] D.C. Look [ Electrical Characterization of GaAs Materials and Devices (John Wiley and Son, Chichester, UK, 1989)] M.L. Lovejoy, M.R. Melloch, M. S. Lundstrom [Appl Phys. Lett. (USA) vol.67 (1995) p. 1101 ] J. del Alamo, S. Swirhun, R. M. Swanson [ IEDM Technical Digest (USA) (1985) p.290 ] G.-B. Gao, DJ. Roulston, H. Morko? [ Solid-State Electron. (UK) vol.33 (1990) p. 1209 ]

3.7

Theoretical hole mobility curves D.C. Look and J.R. Sizelove November 1995

A

INTRODUCTION

The measurement of hole mobilities in GaAs by the Hall effect is, in principal, no more difficult than that of electron mobilities. However, the analysis of hole mobilities is far more involved, because electron transport occurs in a single, nearly spherical energy band, whereas hole transport involves two rather warped energy bands, the light- and heavy-hole bands. Furthermore, these bands are degenerate at small wave vectors, so that interband scattering transitions must be considered. The correct calculation of hole mobility in a polar semiconductor such as GaAs involves a numerical solution of coupled Boltzmann equations, in which all relevant scattering mechanisms are included, with realistic wave functions. Such an 'exact' calculation has evidently never been performed; however, several efforts have at least included a large share of the necessary theoretical complexity [1-4]. Wiley [5] has given an excellent discussion of the various problems involved. B

EFFECTIVE MOBILITY

An effective drift mobility \i can be assigned to the total hole concentration, p = p h + P*, i.e.,

H- ^

^ Ph+P*

(1)

where h denotes heavy holes, and i light holes. EQN (1) is, strictly speaking, valid only for decoupled bands (no interband scattering); however, much of the inaccuracy can be normalized out by fitting the unknown scattering constants (e.g., the acoustic deformation potential) to experimental u vs. T curves. A more serious problem is that most experimental mobilities are determined by the Hall effect, and in the spirit of EQN (1), the Hall mobility u H is given by

MH

_ W h + WPt ~ MhPh + M«p«

, \r)

n

where rh = \im /\ih and rc = \ijJ\Lv The quantities rh and rc are known as 'Hall factors', and denote the ratios of Hall to drift mobilities in a given band. From EQNs (1) and (2), a two-band Hall factor would be given by

r

r - ^ - ( P . * P.) , ^ *

(HhPh

* ''"'"' „ +

№

+

WVT

(3)

In GaAs5 single-band Hall factors at 300 K and 77 K are calculated to be in the range 1.0- 1.4 for the electron band [6] and also for each individual hole band. However, even if rh = rc = 1 for a single band, the combined r for two bands can be different from 1, according to EQN (3). Under typical conditions, several calculations suggest that the theoretical r for holes can be 2 or even greater [2,3,7], and this possibility introduces great uncertainty into comparisons between theory and experiment. For high-mobility electrons in GaAs, r can be experimentally determined by using the relationship r = nH (B=O) /Ji n (B=*) where B is the magnetic field strength. In this case, B = °° means that nB»10 8 , where \i is in cm2/ Vs, and B is in Gauss. This condition is difficult to achieve for holes in GaAs, so very few measurements of r have been attempted in p-type GaAs. However, recently we have carried out a comparison of 296 K Hall data, and electrochemical capacitance-voltage data, in both n-type and p-type GaAs molecular beam epitaxial layers, with n ranging from 2 x 1016 to 8 x 1017 cm"3, and p ranging from 5 x 1016 to 5 x 1019 cm"3, respectively [8], In this case, r = n ^ / nHall, and the experimental values of r turned out to be 1.00 ± 0.08 for all samples. The values of r for electrons were an average of 1.07 below the expected theoretical values [6], and the r values for holes, if multiplied by this factor to effect a rough normalization, ranged from 0.99 to 1.13, certainly far below the expected two-band theoretical values determined from EQN (3). Thus, in light of this surprising experimental result, we consider that EQN (1) will more accurately describe the Hall mobility in p-type GaAs than EQN (2). Also, we reiterate that much of the potential inaccuracy in EQN (1) can be normalized out by fitting to the experimental \i vs. T data for a pure GaAs sample. To calculate the hole mobilities in GaAs, we follow the procedure discussed in [7]. The Boltzmann transport equation is solved for each band by the Rode iterative technique. Both intraband and interband transitions are included, as induced by the following scattering mechanisms: ionized-impurity (Brooks-Herring), acoustic-mode piezoelectric potential, acoustic-mode deformation potential, optical-mode polarization potential, and optical-mode deformation potential (nonpolar). The concentration of ionized impurities in p-type samples is 2ND + p, so that N D is an unknown. The other two unknowns are Eac, the acoustic deformation potential, and Enpo, the nonpolar optical deformation potential. By using data for pure GaAs, [9], in which ND and NA are known from a p vs. T analysis, we can force EQN (1) to give the values n(300 K) = 400 and n(77 K) = 10,000 cm2/ Vs by setting Eac = 8.4 eV, and Enpo = 9.5 eV. As a comparison, Scholz [4] assumes that Eac = 5.6 eV and then fits his more exact theory to experimental data with Enpo =4.2 eV. Other workers get results in the same range, but care must be taken to ensure that these phenomenological terms have the same meaning when making comparisons between various theoretical treatments. C

MOBILITY CURVES

Curves of ^i vs. NA at 300 K and 77 K for various compensation ratios ND / NA are shown in FIGURES 1 and 2. Here, it is assumed that ND+ = ND. The hole concentration is given by p = NA~ - N D . For shallow acceptors, NA" ~ NA. For the following reasons, the curves must be

considered only approximate. (1) The Brooks-Herring ionized- impurity scattering, which assumes the Born approximation, can be inaccurate by at least 20% over the given temperature and concentration ranges, if p-type Si is taken as an example [10]. (2) At the higher concentrations, multi-ion scattering will become important, and long-range potential fluctuations due to acceptor inhomogeneity should be considered [H]. (3) The Boltzmann equations for each band were solved individually, rather than as a coupled pair, and the resulting individual mobilities were then mixed according to EQN (1). (4) The Hall factor r was assumed to be unity, a fact which surprisingly seems to have experimental verification, within 10 - 20%. (5) Band anisotropies were ignored, and in fact, at the higher concentrations, severe band tailing might be expected. However, these potential inaccuracies are partially minimized by fitting the theory to experimental data for pure GaAs at 77 and 300 K.

|i (cm2/Vs)

A convenient way to use FIGURES 1 and 2 is as follows: (1) measure p and \i by the Hall effect; (2) assume NA = p, and plot the point on the appropriate figure; read off N 1 ZNA; (3) get a new NA from NA = p + ND , and replot the point, getting a new N0ZNA . This process may be iterated further, but the major utility of these figures is probably to simply judge whether the compensation

Log N A (cm-3) FIGURE 1. Hall mobility \i vs. compensation ratio ND / NA for holes in GaAs at 300 K. It is assumed that the Hall factor, r, is unity.

(i (cm2/Vs)

Log N A (cm"3) FIGURE 2. Hall mobility vs. compensation ratio N D / N A for holes in GaAs at 77 K. It is assumed that the Hall factor, r, is unity.

is high or low, rather than to determine accurate values of NA and ND . If a measured u falls well below the whole set of curves, it is probable that inhomogeneity of the current flow, rather than microscopic scattering, is causing the low . REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] II1]

M. Costato, C. Jacoboni, L. Reggiani [ Phys. Status Solidi B (Germany) vol.52 (1972) p.461 ] D. Kranzer [ J. Phys. C Solid State Phys. (UK) vol.6 (1973) p.2967 ] K. Takeda, N. Matsumoto, A. Taguchi, H. Taki, E. Ohta, M. Sakata [ Phys. Rev. B (USA) vol.32 (1984)p.ll01 ] R. Scholz [ J. Appl. Phys. (USA) vol.77 (1995) p.3232 ] J.D. Wiley [ Semiconductors and Semimetals (Academic Press, New York, 1975) ch.2 ] D.C. Look [ Electrical Characterization of GaAs Materials and Devices (John Wiley, New York, 1989) p.58] D.C. Look, D.K. Lorance, J.R. Sizelove, CE. Stutz, K.R. Evans, D.W. Whitson [ J. Appl Phys. (USA) vol.71 (1992) p.260] D.C. Look, CE. Stutz [ unpublished ] I. Szafranek, M.A. Piano, S.L. Jackson, S.A. Stockman, G.E. Stillman [Abstracts of the 1990 Electronics Materials Conference (TMS, Warrendale, PA, 1990) p. 14 ] J.R. Meyer, FJ. Bartoli [Phys. Rev. B (USA) vol.23 (1981) p.5413 ] J.R. Meyer, F.J. Bartoli [ Phys. Rev. B (USA) vol.31 (1985) p.2353 ]

3.8

Hole lifetimes in n-type GaAs G.B. Lush, M.R. Melloch and M.S. Lundstrom October 1995

A

EVTRODUCTION

Enough work has now been done in n-type GaAs that there are reference data for all growth techniques in use today. While the hole lifetimes in n-type GaAs are not as well behaved as the electron lifetimes in p-type GaAs3 the data are very consistent despite being collected over a period of 20 years. This Datareview begins with a discussion of the three recombination mechanisms which influence the lifetimes in n-type and p-type GaAs. It then continues with a review of the present state of understanding through a presentation of the data reported in the literature. Recently reported data from samples grown by metal-organic chemical vapour deposition (MOCVD) show almost complete domination by radiative recombination and photon recycling. These data fill what was a void in that no comprehensive study had been done in n-type GaAs grown by MOCVD. B

MECHANISMS OF RECOMBINATION: INTRODUCTION

The lifetime observed in any semiconductor is a combination of the lifetimes of all the recombination mechanisms present in the material. In GaAs the lifetime can be represented as 1 T 1

1

1

T

1

T l

r

U

SRH

T X

^

'

Auger

where x is the minority carrier lifetime, and the terms on the right side represent the three major recombination mechanisms in GaAs (n-type or p-type): radiative recombination, Shockley-Reed Hall (SRH) or thermal recombination, and Auger recombination. It is typically necessary to know not only x but the behaviours of xSRH, xr and xAugcr as functions of doping concentration in order to properly characterize, model, and design GaAs devices. As a first step toward interpreting the data that appear in this review, we must learn to anticipate the expected behaviour of individual mechanisms as functions of doping by first discussing the models we use to describe the recombination mechanisms. C

MECHANISMS OF RECOMBINATION: RADIATIVE

Radiative recombination is the most important recombination mechanism in GaAs both because it is often the dominant mechanism and because GaAs is an 'optical' semiconductor, used for light-emitting diodes and diode lasers [I]. Radiative recombination, which involves the emission of a photon of light during the electronic transition, is an efficient recombination mechanism in GaAs because GaAs is a direct bandgap semiconductor [2].

The equation used to describe the radiative recombination rate is R1. = Bpn which, under low injection conditions, simplifies to R1 = BN4Ap for an n-type semiconductor. In these equations, Nd is the donor concentration, R1. is the rate of radiative recombination, p and n are the hole and electron concentrations, respectively, Ap is the excess hole concentration and B is a constant known as the Einstein coefficient [3]. In general, recombination in semiconductors is described by R = Ap/t so the radiative recombination lifetime can be written as

'•' m;

<2>

At first glance, EQN (2) suggests that the radiative lifetime will be inversely proportional to the electron concentration, but the values of B reported for n-type GaAs vary over two orders of magnitude and B does not seem to be constant with doping [3-5]. B can also be computed theoretically using the detailed-balance method [6] which takes advantage of the fact that the total generation rate in a semiconductor in equilibrium must balance the total recombination rate. B is therefore computed from OO

Bn12

= f rbb(hv)a(hv)vg(hv)d(hv)

(3)

O

where Fbb(hv) is the black body spectrum, a(hv) is the absorption coefficient, and vg(hv) is the group velocity. The limitation of this technique is that it requires knowledge of a(hv) and of ^ in non-degenerately doped semiconductors. Deducing the radiative lifetime in a semiconductor is made more difficult by the presence of what has come to be known as photon recycling [7-10]. Photon recycling is the re-absorption of photons emitted during radiative recombination events. These re-absorbed photons generate new electron-hole pairs with new lifetimes, increasing the lifetime observed experimentally. The photon recycling effect can be lumped into a single term called the recycling cofactor, (j)r [7], such that

*• - jh

w

where F is the probability that an average photon is re-absorbed before escaping the device in question. F is a strong function of the geometry of the device, the device boundary conditions, and of the material absorption coefficient, a(hv). The oc(hv) in turn is a strong function of dopant type and doping concentration [11,12]. Calculations of 4>r have been made for selected conditions, but the results have been inconsistent [7,8,13] and have limited portability, and it is computationally intensive to do it correctly [9].

The recycling co-factor is used as a multiplying factor for the radiative lifetime so that EQN (1) can be re-written as

1 = J - + _L. + X

<Mr

+

1

SRH

+

1 (5)

W

The values for (J)n which can range from a little greater than 1 to possibly as high as 50 for specialized devices [7,14], represent the number of times re-emission/re-absorption events occur before a photon escapes the device. Photon recycling has the effect, therefore, of increasing the effective radiative lifetime with strong implications for material that is dominated by radiative recombination. D

MECHANISMS OF RECOMBINATION: SHOCKLEY-READ-HALL

SRH recombination [15,16] often dominates in n-type GaAs by an amount depending on dopant, film growth techniques, and doping concentration. The SRH mechanism is a two-step recombination process involving mid-gap energy levels or recombination centres. In n-type semiconductors, the recombination centre is typically filled with an electron in equilibrium, and is ready to capture minority holes; in p-type semiconductors the recombination centre is typically empty (has a hole) and is ready for minority electron capture. There is no simple doping dependence for TSHR except that TSHR tends to become smaller with increased dopant concentration for a given growth system. E

MECHANISMS OF RECOMBINATION: AUGER

Auger recombination [17-21] is only important at very high carrier concentrations in semiconductors and is typically not important in GaAs devices except in those such as lasers where high injection levels are present [19,21]. During an Auger recombination event, the energy lost by the electron is absorbed by either another electron or another hole. It is a three carrier event. The rate of Auger recombination is written as R

Auger =

C

n n ( P n ~ *k)

+

C

pP(P* " ^)

(6)

where Cn and Cp are rate constants. Under low injection conditions, only one of these terms will be important, depending on dopant type, and we can write

W = T-bl W

= — ^

n type GaAs

~

p-typ e

GaAs

(7a) (Tb)

Assuming Cn and Cp to be constants, Auger recombination has a very explicit doping dependence. Since Cn and Cp are dependent on band structure, however, and band structure changes at the high

doping or injection concentrations where Auger recombination is important, it is unlikely that Cn and Cp are constants. Attempts to compute Cn and Cp theoretically or to deduce them from experiments have yielded values which vary by well over an order of magnitude [22,17]. Some of these values are inconsistent with recent data. F

TECHNIQUES FOR LIFETIME MEASUREMENT

There is a myriad of techniques to measure lifetime in GaAs that have been reported in the literature. Most of these use heterostructure devices because they offer carrier confinement and keep surface recombination to a minimum. Casey et al used a pn junction with an AlGaAs heteroface passivation layer to perform short-circuit photo-current measurements to deduce diffusion lengths in GaAs [23]. Ettenberg and Kressel observed the electroluminescent decay of heterojunction diodes [24]. Hwang measured steady-state luminescence spectra on unpassivated GaAs to extract the diffusion lengths in n-type GaAs [25]. Because all of these techniques include carrier diffusion, the results of the measurements depend on knowing the diffiisivity of the minority carriers. A more appropriate measurement would measure the lifetime directly in a device in which diffusion is nullified; such a situation can exist in double heterostructures (DHs). Many workers [8,26-31] have used DHs to great advantage to observe the photoluminescence decay of GaAs layers. Under certain conditions which are typically easily met in a DH, a uniform excess carrier concentration is created, eliminating diffusion. Under these conditions the decay constant of the DH can be written as [32] 1

1

1

l

DH

2S w

bulk

where xbulk is the bulk lifetime described by EQN (1), S is the interface recombination velocity, and w is the width of the active GaAs layer. If several DHs of the same doping but varying widths are probed, a plot of l/t DH versus 2/w will yield S as the slope and l/TbuJk as the intercept. In practice, however, photon recycling can complicate the analysis because (J)1. is also a function of the DH thickness (in a thicker DH emitted photons are more likely to be re-absorbed before leaving the device). EQN (8) can be re-written as

-L T

DH

=

_L_ 4>,(d)T r

+

J_

+

tSRH

_ J _ + 2S tAuger

W

which shows explicitly the thickness dependence of c|)r. Lush et al [33] and Nelson and Sobers [34,35] observed non-linear plots of 1/TDH versus 2/w and both attributed this curvature to photon recycling. G

n-TYPE MELT-GROWN OR LPE GaAs

FIGURE 1 shows data reported for lifetimes in n-type GaAs grown by liquid phase epitaxy and melt-grown GaAs. The measurements were made using a wide variety of techniques. Casey et al used short-circuit photo-current measurements [23], Puhlmann et al measured the electron beam induced current [22], and Hwang observed the photoluminescent decay of melt-grown, unpassivated wafers [4,25,36]. All these methods yielded the diffusion length rather than lifetime,

Hole Lifetime (ns)

Casey Hwang Puhlmann Radiative Limit

Doping Concentration (cm'3) FIGURE 1. Hole lifetimes for melt-grown GaAs and GaAs grown by LPE.

and the first two did not distinguish between bulk and interface recombination. FIGURE 1 displays the remarkably consistent data reported by these studies which span 20 years. One can see that the lifetime data seem to saturate at around 20 nanoseconds for lower electron concentrations; this is believed to be caused by background impurities incorporated during growth [36]. The solid line is the supposed radiative limit assuming B = 2.0 x 10"10 cmVs, the number often reported for p-type GaAs. It is called the radiative limit because radiative recombination is an intrinsic mechanism independent of growth technique, and the observed lifetime should not exceed the radiative lifetime. But B is not known for n-type GaAs, it is not known to be constant with doping concentration, and photon recycling can make the radiative lifetime appear longer. Comparison of observed lifetimes to a theoretical xr based on EQN (2) is therefore unreasonable. It is unlikely that B is a constant in n-type GaAs because of heavy doping effects, most notably the Burstein shift [37]. FIGURE 1 clearly shows that we know neither B nor whether photon recycling is being observed because some of the lifetimes observed are longer than the supposed radiative limit, but many are shorter. For Na > 1018 cm"3, the lifetime decreases approximately linearly with increasing electron concentration before it drops more rapidly at the highest concentrations. The initial decrease in lifetime occurs as radiative recombination becomes more competitive with SRH recombination. At the highest concentrations, it was found that SRH recombination began to dominate once again. No evidence of Auger recombination has been found in n-type GaAs [4,29]. The predominance of SRH recombination in these GaAs films has made estimation of the intrinsic recombination parameters difficult; deduced values of the radiative lifetime vary by an order of magnitude. Nonetheless, these data make a good reference of lifetimes to be expected for melt-grown or LPE-grown, n-type GaAs. They are not representative of what is to be expected of n-type GaAs grown by more modern techniques such as metal-organic chemical vapour deposition, MOCVD.

H

n-TYPE GaAs GROWN BY MOCVD

Hole Lifetime (ns)

Lush et al reported the only comprehensive study of lifetimes in n-type GaAs grown by MOCVD [14,33]. The electron concentration ranged from 1.3 x io 17 cm"3 to 3.8 x io 18 cm"3, and the material quality was so good that radiative recombination dominated the lifetimes for most of the samples. Five DHs ranging in width from 0.25 to 10 |im were grown for each of six electron concentrations for this study and the PL decays were observed. FIGURE 2 displays the lifetime data reported in the study; the solid line is the theoretical estimate for the radiative lifetime, computed with the commonly cited value of B = 2.0 x IO"10 cm3/s. Each symbol in the figure represents one of the DHs. All of the DHs exhibited lifetimes longer than this radiative estimate.

Doping Concentration (cm "3 ) FIGURE 2. Hole lifetimes for GaAs grown by MOCVD.

Interface recombination was shown to be negligible in these DHs so the lifetimes observed were truly representative of the bulk lifetime in n-type GaAs. Detailed studies of the intensity dependence of the PL decays showed that for Na < IO18 cm3, the lifetimes were dominated by radiative recombination. For Na > IO18 cm3, SRH recombination became increasingly prominent, but radiative recombination was still important. Much evidence for photon recycling was presented, and recycling was used to explain the dependence of the lifetimes on DH thickness. No evidence of Auger recombination was observed, but an upper limit of Cn = 1.6 x 10"29 (see EQN (6)) was deduced by attributing all the recombination in a particular DH to Auger recombination. The actual value of Cn must be much smaller than this. Lifetimes exceeding 1.0 ^s were observed in two of the DHs doped such that Na = 1.3 x l o17 cm'3, after the substrate was removed by selective etching [14]. These ultra-long lifetimes were a result of enhanced photon recycling by improved photon confinement. Combined, these studies show that the lifetimes to be expected will depend significantly on the device geometry and boundary conditions for any electron concentration for which radiative recombination is important.

I

HOW TO USE THESE DATA

To properly model recombination in n-type GaAs requires detailed simulation including photon recycling because the significance of this effect on the effective lifetime will vary for each device structure. However, there is enough data available to make good estimates according to the device thickness and the growth technique employed. For n-type GaAs grown by LPE or for melt-grown GaAs, the data from FIGURE 1 should be referenced regardless of device structure since non-radiative recombination limited the effects of recycling. For GaAs grown by MOCVD, the data from FIGURE 2 should be referenced with the appropriate device thickness and boundary conditions taken into account, n-type GaAs grown by molecular beam epitaxy has not been studied extensively, but some data [38] indicate that the lifetimes fall somewhere between the LPE/melt-grown and MOCVD materials. These MBE data are from samples grown in a chamber which has produced record efficiency solar cells [39] so the material quality should be the best MBE can produce. J

ULTRA-LONG LIFETIMES

Much work has been done to measure the level of purity possible in GaAs (p-type and n-type) and to study fundamental recombination processes, by studying unintentionally doped GaAs layers. Ultra-long lifetimes (we can arbitrarily define ultra-long lifetimes in GaAs as any x >1 jxs) have been observed by many workers [40,41,31] in unintentionally doped GaAs, but these data must be studied with care because of modulation doping and some non-exponential decays. Lush et al [14] observed lifetimes longer than \\xs in n-type membranes, formed by selectively etching the substrate, and Nelson and Sobers observed a PL decay longer than one microsecond in low doped, LPE p-type GaAs [8]. K

CONCLUSION

Much work has been done to characterize lifetimes in n-type GaAs, and there are sufficient data that estimates of expected lifetime can be made knowing the doping and growth technique used. Many authors have made comparisons between observed lifetimes and the expected radiative lifetime assuming B = 2.0 x 1O~10 cm3/s but there is no evidence to support that value in n-type GaAs and no evidence to support a B coefficient that is constant with electron concentration. More work is needed to deduce B, and the recent data for lifetimes [33], 1I162Dn [42], a(hv) [11], and diffiisivity [43] indicate that new efforts at computing B theoretically and from experiment are warranted. REFERENCES [1] [2] [3] [4] [5] [6] [7] [8] [9]

J. Gowar [ Optical Communication Systems (Prentice Hall, London, 1984) ] S.M. Sze [Physics of Semiconductor Devices fJohn Wiley and Sons, Inc, New York, USA, 1981) ] H.C. Casey Jr., Frank Stern [ J. Appl Phys. (USA) vol.47 no.2 (1976) p.631 ] CJ. Hwang [ Phys. Rev. B (USA) vol.6 no.4 (1972) p. 1355 ] R. N. Hall [ Proc. IEEE (USA) vol. 106B (1960) p.923 ] W. Van Roosbroeck, W. Shockley [ Phys. Rev. (USA) vol.94 no.6 (1954) p. 1558 ] P. Asbeck [ J. Appl Phys. (USA) vol.48 no.2 (1977) p.820 ] RJ. Nelson, R.G. Sobers [ J. Appl. Phys. (USA) vol.49 no. 12 (1978) p.6103 ] T. Kuriyama, T. Kamiya, H. Yanai [ Jpn. J. Appl Phys. (Japan) vol. 16 no.3 (1977) p.465 ]

[10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33] [34] [35] [36] [37] [38] [39] [40] [41] [42] [43]

O. Von Roos [ J. Appl. Phys. (USA) vol.54 no.3 (1983) p. 1390 ] G.B. Lush, H.F. MacMillan, S. Asher, MR. Melloch, M.S. Lundstrom [ J. Appl. Phys. (USA) vol.74 (1992) p.4694] H.C. Casey Jr., D.D. Sell, K.W. Wecht [ J. Appl. Phys. (USA) vol.46 no. 1 (1975) p.250 ] R.K. Ahrenkiel et al [ Appl. Phys. Lett. (USA) vol.55 no. 11 (1989) p. 1088 ] G.B. Lush, D.H. Levi, H.F. MacMillan, R.K. Ahrenkiel, M.R. Melloch, M.S. Lundstrom [ Appl. Phys. Lett (USA), vol.61 (1992) p.2440 ] RN. Hall [ Phys. Rev. (USA) vol.87 no.5 (1952) p.387 ] W. Shockley, W.T. Read Jr., [ Phys. Rev. (USA) vol.87 no.5 (1952) p.835 ] A. Haug [ J. Phys. C. (UK) vol. 16 (1983) p.4159-72 ] L.R. Weisberg [ J. Appl. Phys. (USA) vol.39 no. 13 (1968) p.6096 ] M. Takeshima [ J. Appl. Phys. (USA) vol.58 no. 10 (1985) p.3846 ] M. Takeshima [ Phys. Rev. B (USA) vol.28 no.4 (1983) p.2039 ] L. Jastrzebski, J. Lagowski, H.C. Gatos, W. Walukiewicz [ Gallium Arsenide and Related Compounds (Institute of Physics, London, 1979) p.437 ] N. Puhlmann, G. Oelgart, V. Gottschalch, R. Nemitz [ Semicond. Sd. Technol. (UK) vol.6 (1991) p.181] H.C. Casey Jr., B. I. Miller, E. Pinkas [ J. Appl. Phys (USA), vol.44 no.3 (1973) p. 1281 ] M. Ettenberg, H. Kressel [ J. Appl. Phys. (USA) vol.47 no.4 (1976) p. 1538 ] CJ. Hwang [ J. Appl Phys. (USA) vol.40 no.9 (1969) p.3731 ] G.W. t'Hooft, C. Van Opdrop, H. Veenvliet, A.T. Vink [J. Cryst. Growth (Netherlands) vol.55 (1981)p.l73-82] G. W. t'Hooft [Appl. Phys. Lett. (USA) vol.39 no.5 (1981) p.389 ] L.W. Molenkamp, H.F.J. vanrt Blik [ J. Appl. Phys., vol.64, no.8 (1988) p.4253 ] DZ. Garbuzov [ Semiconductor Optoelectronics Ed M.A. Herman (Wiley, USA, 1980) p.305-43] D.Z. Garbuzov [ Semiconductor Physics Eds M.M. Tuchkevich, V. Y. Frenkelpp, (Consultants Bureau, 1986) p.5386 ] J.M. Olson, RK. Ahrenkiel, DJ. Dunlavy, B. Keyes, A.E. Kibbler [ Appl. Phys. Lett. (USA) vol.55 no. 12 (1989) p. 1208] A. Many, Y. Goldstein, N.B. Grover [ Semiconductor Surfaces (John Wiley and Sons, New York, 1965)] G.B. Lush et al [ J. Appl. Phys. (USA) vol.72 no.4 (1992) p. 1436 ] RJ. Nelson, RG. Sobers [ Appl. Phys. Lett. (USA) vol.32 no. 11 (1978) p.761 ] RJ. Nelson [ J. Vac. Sci. Technol. (USA) vol. 15 no.4 ( 1978) p. 1475 ] C. J. Hwang [ J. Appl. Phys. (USA) vol.42 no. 11 (1971) p.4408 ] E. Burstein [ Phys. Rev. (USA) vol.93 (1954) p.632 ] M.R. Melloch [ Unpublished data from double heterostuctures grown by MBE and measured by photoluminescence decay ] S. Tobin et al [ IEEE Trans. Electron Devices (USA) vol.37 (1990) p.469 ] LW. Molenkamp, G.L.M. Kampschoer, W. de Lange, J.W.F.M. Maes, PJ. Roksnoer [ Appl. Phys. Lett. (USA) vol.54 no.20 (1989) p. 1992 ] D. J. Wolford et al [ J. Vac. Sci. Technol. B (USA) vol.9 no.4 (1991) p.2369 ] E.S. Harmon, M.R. Melloch, M.S. Lundstrom [ Appl. Phys. Lett. (USA) vol.64 (1994) p.502 ] E.S. Harmon et al [ Datareview in this book: 2.13 Minority electron mobility in doped GaAs ]

CHAPTER 4 BAND STRUCTURE AND CARRIER IONISATION 4.1 4.2 4.3 4.4 4.5 4.6 4.7 4.8 4.9

Band structure of GaAs: overview Direct bandgaps of GaAs, temperature dependence Direct bandgap of GaAs, pressure dependence Optical properties of amorphous GaAs Intra- and intervalley deformation potentials for electrons in GaAs Intravalley deformation potentials for holes in GaAs Electron effective mass and hole effective mass in GaAs, pressure dependence Effective bandgap narrowing in doped GaAs Carrier ionisation coefficients of GaAs

4.1

Band structure of GaAs: overview E.P. O'Reilly January 1996

A

INTRODUCTION

GaAs is the most widely studied of the III-V semiconductors. In this Datareview, we provide an overview of the main features of its band structure. GaAs has a direct bandgap, with both the valence band maximum and the conduction band minimum at the centre of the Brillouin zone (F point). The conduction band minimum is doubly degenerate (of F 6 symmetry), and is separated by an energy E0 from the valence band maximum which is four-fold degenerate (F 8 symmetry). Away from the zone centre, the highest valence bands separate into two sets of bands, referred to as the heavy-hole (HH) and light-hole (LH) bands, because of their large and small effective masses respectively. The next set of valence bands, referred to as the spin-split-off (SO) bands, also have their maximum at F, where they are doubly degenerate (of F 7 symmetry), separated by an energy A0 from the F 8 states.

Energy (eV)

FIGURE 1 shows the band structure of GaAs, as calculated by Chelikowsky and Cohen [1] using the empirical pseudopotential method.

Wave vector k FIGURE 1. Band structure of GaAs, calculated along high symmetry directions (from [I]).

The most important directions in the Brillouin zone lie along the [111] A direction, the [001] A direction and also the [110] S direction, which reach the edge of the first Brillouin zone at the L, X and K points respectively. While the conduction band minimum is situated at F, higher sets of minima are found at L and near X (about 10% away from the zone boundary along A) which are important for optical and electronic properties. We present information in Section B on the energies of states at high symmetry points in the Brillouin zone (F, X and L), focussing in particular on the lowest conduction band and the three highest sets of valence bands but including also details of higher-lying conduction states. Section C presents details of carrier effective masses at the F conduction band minimum and for the heavy- and light-hole valence bands. B

SYMMETRY POINT ENERGIES

The energies of interband transitions between valence and conduction states in GaAs and their temperature dependence have been measured by a wide variety of optical techniques, including photoluminescence (for the E0 gap), absorption, reflectance, and several reflectance-modulation techniques, including piezoreflectance, magnetoreflectance, thermoreflectance, electroreflectance and wavelength-modulated reflectance. A review of the E0 data is presented by Adams [2] and of the optical spectra as a function of photon energy by Nolte [3], while an extensive list of experimental references is included in, for example, Lautenschlager et al [4]. These optical techniques provide information on the separation of valence and conduction states which are associated with the same wavevector (often F, X or L) [5]. The relative energies of states at different wavevectors must then be found by other techniques. The valence band dispersion along A, A and S has been measured by angle-resolved photoemission, with an energy resolution of approximately 0.2 eV [6], while the dispersion of the lowest conduction bands [7] and of higher conduction bands along the S [8] and A [9] directions has also been measured using photoemission techniques. The energy separation between the lowest conduction states at F and X has been accurately measured by studying the variation of direct and indirect recombination transition energies as a function of hydrostatic pressure [10]. The energy of the lowest conduction L states has been determined from Schottky barrier electroreflectance [11] and from a careful analysis of Hall effect measurements at elevated temperatures [12,13]. The energy separation between any two states at F, X or L can in general be independently measured by at least two different paths which are in good agreement with each other: e.g. the F-X separation at 4 K in the lowest conduction band, X6 - F6, is found directly and accurately from high pressure photoluminescence data to be 0.490 eV [10] or by combining optical transition [14] and photoemission data to be 0.55 eV [6]. There have been many theoretical calculations of the band structure of GaAs, including both empirical and first principles methods. The non-local empirical pseudopotential method [1,15] gives symmetry point energies in the highest valence and lowest conduction bands to within about 0.2 eV while tight-binding parameters have been developed which describe well the valence band structure and the lowest conduction band, e.g. [16]. Ab-initio calculations have been performed using density functional theory (DFT) with the local density approximation (LDA) for exchange and correlation [17]. This is formulated to give exact ground state energies but underestimates the energies of excited (= conduction band) states. To go beyond the LDA ground-state theory, the quasiparticle band structure of GaAs has been calculated using the GW approximation [18-

20]. Displacing the LDA conduction bands rigidly upwards by 0.8 eV gives a band structure within 0.1 eV of the GW calculations in the bands closest to the gap. We present in TABLE 1 the experimentally determined low-temperature (T close to 0 K) symmetry point energies of GaAs, as well as the energies calculated using the non-local empirical pseudopotential method (NEPM) [1] and the GW approximation [18,19]. The valence band maximum is chosen as the zero of energy. The GW ab-initio calculations give valence and conduction symmetry point energies near the band gap which are within 0.1 eV of experiment. The overall good agreement between the GW and experimental data is discussed further in refs [9] and [20], both of which include plots of the calculated and experimental dispersion along the A and A directions. TABLE 1. Energies of symmetry points in GaAs band structure (in eV). Symmetry point

Experimental energy

NEPM[I]

GW

T6v T7v IV IV T7c T88

-13.1 [6] - 0.341 [21]a 0.0 1.519 [2] 4.72[6,14]a 4.72 [6,14]a

-12.55 - 0.35 0.0 1.51 4.55 471

-12.69[2O] (-0.34)b [18] 0.0 b [18] 1.47 [18] 4.52 [18] 4.52 [18]

X6. X6. X6. X7v X60 Xzfi

-10.75 [6] -6.70 [6] - 2.80 [6]a -2.80[6]a 2.01 [10] 2.58 [6,14]

-9.83 -6.88 - 2.99 -2.89 2.03 2M

-10.27 [20] -7.16[2O] (-2.80)b [18] -2.73 b [18] 2.08 [18] 2.30 [18]

L6v L6v L6v L45v L6I L60

-11.24 [6] - 6.70 [6] -1.30[6]a -1.30[6]a 1.84 [10,11] 5.45 [9]

-10.60 - 6.84 -1.42 -1.20 1.82 5.47

-11.02[2O] -6.91 [20] (-1.32)b[18] -1.11b [18] 1.82 [18] 5.41 [18]

L^

I 5.45 [9]

I 5.52

| 5.41 [18]

a)

The spin-orbit splittings between the highest valence bands were not determined by Chiang et al [6]. They have been measured by other authors using reflectance techniques. Aspnes and Studna [21] determined from electroreflectance measurements that L4 5v - L6v = 0.22 eV and that X7v - X6, = 0.076 eV.

b)

The GW calculations did not include spin-orbit interaction. The values in brackets indicate where the theoretical results were modified to include experimentally determined splittings [18].

The temperature dependence of the bandgaps is associated theoretically with the dilation of the lattice and with electron-phonon interactions. It has been postulated that relative valence band energies do not vary with temperature [22]. The bandgaps between the valence band maximum and the lowest conduction states at F, X and L all decrease with temperature and their temperature dependence can be fitted by the Varshni equation [23], whereby

E(T) = E(T=O K) - - ^ T +b

(1)

The Varshni coefficients a and b for the lowest F, L and X conduction band states relative to the top of the valence band are shown in TABLE 2. The direct gap (T) data, is reviewed by Wilkinson and Adams [2], while the temperature dependence of the L gap is presumed to follow that of the E1 optical transition. The temperature dependence of the X gap has not been measured directly but was inferred by Aspnes [11] from a comparison with data for the indirect gap semiconductor GaP. TABLE 2. Varshni coefficients for the F, L and X conduction-band minima of GaAs relative to the top of the valence band. aClO^eVK-1)

b(K)

r [24,4]

5.405, 5.5

204,225

LJ4]

12,

205

X[Il]

6£5

204

C

BAND EDGE EFFECTIVE MASSES

Cl

Electron Mass

A detailed comparative survey of the known experimental and theoretical values of the electron and heavy-hole effective mass in GaAs has been given by Nakwaski [25]. The conduction band dispersion very close to the band minimum at F is parabolic and isotropic with the effective mass Inn at 4 K measured to be 0.0665 Tn0 [26] where mo is the free-electron mass. The band structure becomes nonparabolic and anisotropic at higher energies. The dispersion at higher energies can be described using a k.p perturbation theory which includes 14 bands (the 6 highest valence and 8 lowest conduction bands) [27]. The values of the k.p parameters have been determined experimentally by Sigg et al [28] from cyclotron resonance data. The conduction band dispersion E(k) within 50 meV of the band edge can be described by an expression up to fourth order in k which is of the form E(k) = ^ !

+

Ak 4

+

B(Icx2ICy2

+

k ^

+ Ic22Icx2)

n

±c/k 2CkxX2 + kX 2 + KX) - ^K2K2

(2)

where the vector k has components (Icx, ky, Ic2) and the parameters A3 B and C are all negative and can be determined from a 14-band k.p calculation [27]. The nonparabolicity parameters have been determined by Malcher et al [29] as A B C

= = =

-2107 eV (A)4 -2288 eV (A)4 -27.57 eV (A)3

The conduction band nonparabolicity can generally be ignored when using effective mass theory to calculate the lowest confined electron state in GaAs quantum well structures but has a significant influence on the energies of higher confined states [30] and must also be included to account for the observed enhancement and strong nonparabolicity of the electron effective mass in the plane of the quantum well [31]. Because the value of the electron effective mass depends from k.p theory on the bandgap, it decreases with increasing temperature, so that Inn = 0.063 is the recommended value at room temperature [25]. C2

Heavy and Light-Hole Masses

The heavy- and light-hole valence bands are anisotropic in III-V semiconductors at the valence band maximum. This effect is most pronounced for the HH band which has a strong directionally dependent effective mass, with a larger mass along the (111) direction compared to the (001) direction, In11(111) > Inn(OO 1). The bulk valence band effective masses are often defined in terms of the Luttinger y parameters [32], Yi, y2>anc* Y3 with IV(OOl) VO11) V(OOl) V(Hl)

= = = =

Yi-2y 2 YI-2Y 3 Yi+ 2Y2 Yi + 2Y3

The Y values were originally determined by magneto-optical or cyclotron resonance experiments [33-35] in which an exciton mass or cyclotron hole mass is measured. These latter quantities result from averaging band effective masses in the plane perpendicular to the applied magnetic field. More direct measurements have since been undertaken using photoluminescence excitation spectroscopy [36] and resonant electron Raman scattering [37] to determine hole confinement energies and hence hole effective masses in GaAs/GaAlAs quantum well structures grown on [001], [111] and [310] substrates. The Y parameters measured by the two techniques are in good agreement with each other. Molenkamp et al [36] deduce Yi = 6.790, Y2 = 1.924 and Y3 = 2.681, while Shanabrook et al [37] find Yi = 6.8, y2= 1.9 and y3= 2.73. These values are consistent with the results of recent hot-electron luminescence measurements which probed the directional dependence of the heavy-hole dispersion [38], and are to be preferred over the earlier values, in particular for applications to quantum well and low dimensional structures. REFERENCES [1]

J.R. Chelikowsky, M.L. Cohen [ Phys. Rev. B (USA) vol.14 (1976) p.556-82 ]

[2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33] [34] [35] [36] [37] [38]

V. A. Wilkinson, A.R. Adams [ Datareview in this book: 4.2 Direct bandgaps of GaAs, temperature dependence ] D.D. Nolte [ Datareviews 5.1 and 5.2 in this book] P. Lautenschlager, M.Garriga, S.Logothetidis, M. Cardona [ Phys. Rev. B (USA) vol.35 (1987) p.9174-89] M. Alouani, L. Brey, N.E. Christensen [ Phys. Rev. B (USA) vol.37 (1988) p. 1167-79 ] T.C. Chiang, J.A. Knapp, M. Aono, D.E. Eastman [ Phys. Rev. B (USA) vol.21 (1980) p.3513-22 ] D. Straub, M. Skibowski, FJ. Himpsel [ Phys. Rev. B (USA) vol.32 (1985) p.5237-44 ] G.P. Williams, F. Cerrina, G.J. Lapeyre, J.R. Anderson, RJ. Smith, J. Hermanson [ Phys. Rev. B (USA) vol.34 (1986) p.5548-57 ] J.E. Ortega, FJ. Himpsel [ Phys. Rev. B (USA) vol.47 (1993) p.2130-7 ] DJ. Wolford, J.A. Bradley [ Solid State Commun. (USA) vol.53 (1985) p. 1069-76 ] D.E. Aspnes [ Phys. Rev. B (USA) vol. 14 (1976) p.5331-43 ] A.K. Saxena [ J. Phys. C (UK) vol. 13 (1980) p.4323-34 ] HJ. Lee, LY. Juravel, J.C. Woolley, AJ. Springthorpe [Phys. Rev. B (USA) vol.21 (1980) p.65969] D.E. Aspnes, CG. Olson, D.W. Lynch [ Phys. Rev. B (USA) vol. 12 (1975) p.2527-38 ] K.C. Pandey, J.C. Phillips [ Phys. Rev. B (USA) vol.9 (1974) p. 1552-9 ] P. Vogl, HJ. Hjalmarson, J.D. Dow [ J. Phys. Chem. Solids (UK) vol.44 (1983) p.365-78 ] GB. Bachelet,N.E. Christensen [Phys. Rev. B (USA) vol.31 (1985)p.879-87 ] R.W. Godby, M. Schluter, LJ. Sham [Phys. Rev. B (USA) vol.35 (1987)p.4170-1 ] X. Zhu, S.G. Louie [ Phys. Rev. B (USA) vol.43 (1991) p. 14142-56 ] M. Rohlfing, P. Kruger, J. Pollmann [ Phys. Rev. B (USA) vol.48 (1993) p. 17791-805 ] D.E. Aspnes, A.A. Studna [ Phys. Rev. B (USA) vol.7 (1973) p.4605-25 ] D. Auvergne, J. Camassel, H. Mathieu, M. Cardona [ Phys. Rev. B (USA) vol.9 (1974) p.5168-77 ] Y.P. Varshni [ Physica (Netherlands) vol.34 (1967) p. 149 ] CD. Thurmond [ J. Electrochem. Soc. (USA), vol. 122 (1975) p. 1133-41] W. Nakwaski [ Physica B (Netherlands) vol.210 (1995) p. 1-25 ] GE. Stillman, DM. Larsen, CM. Wolfe, R.C. Brandt [ Solid State Commun. (USA) vol.9 (1971) p.2245 ] M. Braun, U. Rossler [ J. Phys. C (UK) vol. 18 (1985) p.3365-77 ] H. Sigg, J.A.AJ. Perenboom, P. Pfeffer, W. Zawadski [ Solid State Commun. (USA) vol.61 (1987) p.685-90 ] F. Malcher, G. Lommer, U. Rossler [ Superlattices Microstruct. (UK) vol.2 (1986) p.267-72 ] A.T. Meney, B. Gonul, E.P. O'Reilly [ Phys. Rev. B (USA) vol.50 (1994) p. 10893-904 ] U. Ekenberg [ Phys. Rev. B (USA) vol.40 (1989) p.7714-26 ] J.M. Luttinger [ Phys. Rev. (USA) vol. 102 (1956) p. 1030 ] Q.H.F.Vrehen [ J. Phys. Chem. Solids (UK) vol.29 (1968) p. 129 ] M.S. Skolnick, A.K. Jain, R A Stradling, J. Leotin, J.C. Ousset, S. Askenazy [ J. Phys. C (UK) vol.9 (1976)p.2809-21] K. Hess, D. Bimberg, N.O. Lipari, J.U. Fishbach, M. Altarelli [ Proc. 13th Int. Conf. on Physics of Semiconductors, Ed. F.CFumi (North-Holland, Amsterdam, 1976) p. 142 ] LW. Molenkamp, R. Eppenga, G.W. 't Hooft, P. Dawson, CT. Foxon, KJ. Moore [ Phys. Rev. B (USA) vol.38 (1988) p.4314-17 ] B.V. Shanabrook, OJ. Glembocki, D.A. Broido, W.I. Wang [ Superlattices Microstruct. (UK) vol.5 (1989) p.503-6] W. Hackenberg, H.P. Hughes [ Phys. Rev. B (USA) vol. 49 (1994) p.7990-9 ]

4.2

Direct bandgaps of GaAs9 temperature dependence V.A. Wilkinson and A.R. Adams July 1995

A INTRODUCTION We review here the effects of temperature on the direct gaps of GaAs and recommend values for use in analysis. The direct gaps decrease with increasing temperature, due partly to thermal expansion but also to electron-phonon interactions, with the latter effect dominating at temperatures above -100 K. The temperature variation of the bandgap is commonly described by the empirical relation proposed by Varshni [1] E(T) = E(O)-aT 2 /(T +P)

(1)

where a and P are material-dependent constants. More recently the semi-empirical Bose-Einstein expression has been proposed [2,3] E(T) = E 8 - aB(l + 2/{exp (0/T) -1})

(2)

where aB represents the strength of the electron-phonon interaction and 0 corresponds to the average phonon frequency.

B TEMPERATURE DEPENDENCE IN GaAs The temperature dependence of the most important direct bandgaps in GaAs has been studied by many authors using a wide range of optical techniques. These have included absorption [4-6], reflectivity [7,8], electroreflectance [9], piezoreflectivity [10], wavelength-modulated reflectance [11], thermoreflectance [12] and ellipsometry [13]. The effect has also received considerable theoretical attention (see for example [14-18]). Lautenschlager et al [13] used ellipsometry over the range 20 to 750 K to study the temperature dependence of the direct gaps E0, E0H-A0, E1, E1H-A1, E 0 ' and E 2. Their paper also provides an excellent review to which readers are referred for a detailed discussion and bibliography. The best fits to EQN (1) and EQN (2) obtained by Lautenschlager et al [13] agree with most other workers and are given in TABLE 1. Grilli et al [19] have made a high-precision determination of the temperature dependence OfE0, between 2 and 280 K, using photoluminescence. Their paper includes a discussion of the limitations of EQNs (1) and (2) to accurately describe experimental data for GaAs, to which readers are referred.

TABLE 1. Values for the fitting parameters defined in E Q N s ( I ) and (2), taken from Lautenschlageretal [13]. Numbers in parentheses indicate error margins. E(O) (eV)

a 10-4 (eV/K)

p (K)

EB (eV)

aB (meV)

9 (K)

En

1.517(8)

5.5(1.3)

225(174)

1.571(23)

57(29)

240(102)

E o +A o

1.851(5)

3.5(4)

225(fixed)

1.907(9)

58(7)

240(fixed)

E,

3.041(3)

7.2(2)

205(31)

3.125(9)

91(11)

274(30)

E0'

4.509(8)

4.0(7)

241(177)

4.563(21)

59(26)

323(119)

E,

I 5.133(21)

| 5.161(33)

| 38(33)

I 6.6(4)

| 43(66)

| 114(95)

The spin orbit splitting at the L point, A1, was found to be constant with temperature within experimental accuracy. For accurate fitting to E0, note that the free exciton, which may dominate optical properties, has a binding energy of a few meV. Sell [20] reports a value of 4.2 meV for epitaxial GaAs at 2 K. C CONCLUSION TABLE 1 gives useful values for determining the direct gaps of GaAs as a function of temperature. For highly accurate values refer to comments given in [19]. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] II1] [12] [13] [14] [15] [16] [17] [18] [19] [20]

Y.P. Varshni [ Physica (Netherlands) vol.34 (1967) p. 149 ] L. Vina, S. Logothetidis, M. Cardona [ Phys. Rev. B (USA) vol.30 (1984) p. 1979 ] S. Logothetidis, L. Vina, M. Cardona [ Phys. Rev. B (USA) vol.31 (1985) p.947 ] M.D. Sturge [ Phys. Rev. (USA) vol. 127 (1962) p.768 ] M.B. Panish, H.C. Casey Jr. [ J. Appl Phys. (USA) vol.40 (1969) p. 163 ] M. Cardona, G. Harbeke [ J. Appl. Phys. (USA) vol.34 (1963) p.813 ] D.D. Sell [ Proc. I lth Int. Conf. on Physics of Semiconductors, Warsaw, 1972, Ed. M. Miasek (Elsevier, Amsterdam and PWN-Polish Scientific Publishers, Warsaw, 1972) ] A.G. Thompson, J.C. Woolley, M. Rubenstein [ Can. J. Phys. (Canada) vol.44 (1966) p.2927 ] T. Nishino, M. Okuyama, Y. Hamakawa [ J. Phys. Chem. Solids (UK) vol.30 (1969) p.2671 ] J. Camassel, D. Auvergne, H. Mathieu [ J. Appl. Phys. (USA) vol.46 (1975) p.2683 ] R.R.L. Zucca, Y.R. Shen [ Phys. Rev. B (USA) vol. 1 (1970) p.2668 ] E. Matatagui, A.G. Thompson, M. Cardona [ Phys. Rev. (USA) vol. 176 (1968) p.950 ] P. Lautenschlager, M. Garriga, S. Logothetidis, M. Cardona [ Phys. Rev. B (USA) vol.35 (1987) p.9174] P.B. Allen, M. Cardona [ Phys. Rev. B (USA) vol.23 (1981) p. 1495 vol.24 (1981) p.7479 ] P.B. Allen, M. Cardona [ Phys. Rev. B (USA) vol.27 (1983) p.4760 ] P. Lautenschlager, P.B. Allen, M. Cardona [Phys. Rev. B (USA) vol.31 (1985) p.2163 ] P. Lautenschlager, P.B. Allen, M. Cardona [ Phys. Rev. B (USA) vol.33 (1986) p.5501 ] CK. Kim, P. Lautenschlager, M. Cardona [ Solid State Commun. (USA) vol.59 (1986) p.797 ] E. Grilli, M. Guzzi, R. Zamboni [ Phys. Rev. B (USA) vol. 45 (1992) p. 1638 ] D.D. Sell [ Phys. Rev. B (USA) vol.6 (1972) p.3750 ]

4.3

Direct bandgap OfGaAs9 pressure dependence V.A. Wilkinson and A.R. Adams July 1995

A INTRODUCTION The direct bandgap of GaAs increases with hydrostatic pressure throughout the Brillouin zone. This can be regarded as the separation of the bonding (valence band) and the anti-bonding (conduction band) energy levels as the atoms are brought closer together. Optical measurements of the band structure can be made using diamond-anvil-cell techniques up to 17 GPa when a structural phase transition occurs [I]. B DIRECT GAP, E0 Goni et al [2] have tracked E0 up to the phase transition pressure using optical absorption at 300 K. In common with all crystalline semiconductors, the pressure dependence for GaAs is generally described by a quadratic form: E0 = E0(O) + bP + cP2

(1)

Goni et al obtain the following coefficients from their fitting: E0(O) = 1.43 ± 0.01 eV b = (10.8 ± 0.3) x 10"2 eV / GPa c =-(14 ±2) x 10" 4 eV/GPa 2 Similar results using the same technique at 77 K have been obtained by Kobayashi [3]. The sublinearity arises from the decrease in compressibility with increasing pressure. Goni et al [2] showed that E0 varied linearly with density D: dE0/d(lnD) = 8.5eV

(2)

Wolford and Bradley [4], reporting on photoluminescence measurements at liquid helium temperatures, give a linear coefficient of: dEo/dP = (10.73 ± 0.05) x 10"2 eV / GPa

(3)

for pressures up to ~4 GPa4 Perlin et al [5] have conducted detailed experiments on GaAs quantum wells and deduce a linear pressure coefficient of 11.6 x 10"2 eV / GPa for bulk GaAs in the low pressure range 0-1.5 GPa.

C

INDIRECT GAP AND CROSSOVER PRESSURES

The valence band broadens with pressure more quickly than E2 increases and hence the indirect gap from the valence-band maximum to the conduction-band X-minima decreases with pressure. Goni et al [2] obtained a value for this pressure coefficient of-1.35 x 10"2 eV/GPa and observed that the F and X conduction-band minima cross at 4.2 ± 0.2 GPa at 300 K. Wolford and Bradley [4] give the coefficient: dEr_x/dP = -1.34 x 10"2 eV / GPa

(4)

with a F-X crossing at 4.13 GPa. D

CONCLUSION

A linear coefficient of 10.73 x 10"2 eV / GPa can describe the pressure dependence OfE0 at low pressures. For high pressures a non-linear description should be used. Sly and Dunstan [6] have discussed inadequacies in the quadratic form. Comparing the room temperature coefficients of Goni et al with the ~4 K coefficients of Wolford and Bradley we may conclude that the pressure coefficients are insensitive to temperature. From the results of Goni et al [7] we may also conclude that the spin orbit splitting, A0, is insensitive to pressure. REFERENCES [1] [2] [3] [4] [5] [6] [7]

M. Baublitz, A.L. Ruoff [ J. Appl Phys. (USA) vol.53 (1982) p.6179 ] A.R. Goni, K. Strossner, K. Syassen, M. Cardona [ Phys. Rev. B (USA) vol.36 (1987) p. 1581 ] T. Kobayashi [ Semicond. Sd. Technol. (UK) (1989) p.248 ] DJ. Wolford, J.A. Bradley [ Solid State Commun. (USA) vol.53 (1985) p. 1069 ] P. Perlin, W. Trzeciakowski, E. Litwin-Staszewska, J. Muszalski, M. Micovic [ Semicond. Sci. Technol. (UK) vol.9 (1994) p.2239] A.D. Prins, J.L. Sly, DJ. Dunstan [ Phys. Status Solidi B (USA) vol. 198 (to be published 1 Nov. 1996)] A.R. Goni, K. Syassen, K. Strossner, M. Cardona [ Semicond. Sci. Technol. (UK) vol.4 (1989) p.246 ]

4.4

Optical properties of amorphous GaAs R. Muni and N. Pinto March 1996

A

INTRODUCTION

Tetracoordinated III-V compound semiconductors can be obtained in the form of amorphous (a III-V) thin films by different deposition techniques. Several efforts have been made to investigate, both experimentally and theoretically, this class of amorphous material for application in optoelectronic devices operating in the visible range. These amorphous compounds retain the tetrahedral coordination of the crystalline structure. However, defects related to the occurrence of unsatisfied bonds (dangling-bonds) and deviations from the normal heteropolar bonding, i.e. wrong-bonds or bonds between atoms of the same type, may be present. Additional problems arising from non-stoichiometry of the deposited thin films explain why amorphous compounds have been less investigated than the elemental amorphous materials. Al

Growth

Several deposition techniques have been used to deposit a III-V films. Flash evaporation of a crystalline powder gives films whose composition depends strongly on the evaporation conditions (the grain size and crucible temperature). Any excess of the more volatile group V element can be reduced or eliminated by adjusting the crucible temperature or by increasing the substrate temperature during the deposition [1,2]. Cathode sputtering has also been widely used to prepare a III-V films, generally from a crystalline target of GaAs. Although the different sputtering rates of the two constituents generate nonstoichiometric samples, stoichiometric films can be obtained by careful control of the deposition conditions [3-5]. Other different deposition techniques like molecular beam epitaxy (MBE) [7] and plasma deposition [8,9] have been used. The density deficit of optimized amorphous GaAs films with respect to the corresponding crystal is small: from a few % to about 10%. A2

Dielectric Properties

The real, e l3 and the imaginary part, e2, of the complex dielectric constant of stoichiometric aGaAs are functions of "hco, and both show broad peaks characteristic of amorphous materials when compared to their crystalline counterparts. The behaviour of the e2 curve corresponds to that expected from the Perm model, e2 ^ (ho - Eg)"1/2 with Eg (the average bond strength) being smaller in the amorphous than in the crystalline state. In the absorption edge region, e2 follows the well known empirical law :hco2e2« (hco - E0)2 for amorphous tetracoordinated semiconductors, where E0 is the extrapolated optical gap. In this context, E0 represents the energy separation between extended states in the valence and conduction bands. E0 has a different meaning from Eg, the

average separation between the two bands. B

OPTICAL CONSTANTS

Bl

Flash Evaporated a-GaAs

Nearly stoichiometric (within 1%) a-GaAs films deposited by flash evaporation of crystalline powder under ultra high vacuum [10] show a change of the optical gap, E0, from (1.0- 1.1) eV to (1.15 -1.20) eV when annealed at 570 K, while the refractive index at 0.5 eV is lowered from (3.75-3.85) to (3.65-3.70). The crystallization temperature of a-GaAs seems to be about 300 0 C. Stoichiometric flash evaporated a-GaAs films deposited at 100 K were annealed at room temperature or at 515 K [2]. The optical constants were found to change if annealing is carried out between these temperatures and saturate to a 'stabilized' configuration of the structure with n = 3.6 (hv = 0.5 eV) and E0 = 1.15 eV. Annealing seems to have less effect on the gap than hydrogenation. Structural (dangling-bond type) and chemical (wrong-bond type) defects both introduce states in the pseudo-gap, or at the band edges. Hydrogen eliminates defects responsible for localized states in the pseudo-gap and modifies the sample structure. Films containing As excess do not present substantial differences from stoichiometric samples. Electrical conductivity measurements as a function of T found that E0 = 0.55 eV, and suggested that the distribution of localized states introduced by unsatisfied bonds seems to be structureless across the pseudo-gap [2]. Wrong-bonds seem not to be present. Photoabsorption spectra near the As and Ga L3 edges in a-GaAs indicate that only Ga s-states are found at the bottom of the conduction band. In other words, disorder does not induce noticeable mixing of As and Ga states close to the conduction band edge [H]. B2

Sputtered a-GaAs

RF sputtered hydrogenated a-GaAs films [3] show a shift of the room temperature absorption edge with hydrogen pressure, Pn, from « 0.9 eV (pH = 0 torr ) to « 1.42 eV (pH = 5.2 x 10"3 torr) both computed for an absorption coefficient 104 cm"1. These films were deposited at a substrate temperature near 40 C, because above 250 C they were partially crystalline. The refractive index of both unhydrogenated and hydrogenated films as a function of photon energy presents a maximum value with the highest n « 4.2 - 4.4 at hv = 2.0 - 2.17 eV [4]. The absorption coefficient, a, and the optical gap, E0, vary when (Pn/ p tot) increases from 0% to 20%. In fact a, decreases rapidly from 1.1 x 105 cm"1 to 2 x 104 cm -1 at hv « 1.7 eV, while E0 varies from * 1.05 eV to a constant value « 1.45 eV. This behaviour reflects that of the hydrogen content, CH, which increases from 0 to 77 mol %, as (p H/ p tot) varies from 0% to 40%. C H is estimated by nuclear analysis assuming perfect stoichiometry of the films. Using Mo co-sputtering [12] films are doped with a nominal 3% Mo. The optical gap is reduced

by 0.1 eV and an increased absorption tail is introduced. The photoacoustic technique was used [13] to measure the refractive index in the range 900 -1800 nm at room temperature. An increase of n towards shorter wavelengths was found by increasing the hydrogen content (0% < CH < 40 %). Thus, n varies from « 3.6 to « 5.6 according to X and CH. Stoichiometric films can be obtained at a substrate temperature of T s = 290 0 C [14]. At lower T s an As excess is found whilst at higher T s a Ga excess occurs. Bond-angle disorder in amorphous material and/or wrong-bond defects in compound semiconductors can cause a reduction of the energy gap and a modification of the optical constants For near stoichiometric films (within 3-4%), an increase of the deposition temperature from room temperature to 3000C causes an increase of the slope of a vs. hv curves and the widening of the optical gap, E0, from (0.9 -1.0) eV to 1.55 eV [14-16]. The direct effect of stoichiometry on the optical properties of the films is more difficult to find. A variation in the composition ratio (CM I C02) from 1.22 to 1.05 slightly shifts the absorption curve towards higher energies whilst barely affecting the optical gap [6]. On the contrary, the absorption coefficient decreases rapidly at lower energies after hydrogenation in stoichiometric or As-rich samples. The refractive index is of the order of 3.2 (X = 1.5 ^m) [15] for unhydrogenated and sputtered materials, lower than the value of 3.8 obtained by flash evaporation [2]. B3

Hydrogenated a-GaAs

Because of its successful use in ex-Si, hydrogenation of amorphous material during preparation has been generally adopted. The advantage of hydrogen incorporation in the film is the strong reduction of the band tails due to a saturation of dangling bonds by atomic hydrogen. Only 5% of atomic hydrogen in the film causes a strong increase in the Tauc optical gap from 0.9 -1.0 eV (C n = 0%) to 1.5 -1.6 eV (C11= 5%) [5,14,15]. The refractive index of unhydrogenated amorphous GaAs is higher than that of crystalline material, exhibiting a strong variation with hydrogenation. The index of refraction was computed [6] from reflectivity measurements in the spectral range from 0.4 to 1.5 eV. The refractive index at 0.5 eV is of the order of 3.7 (CH = 0%) and decreases with increasing hydrogen content. This difference increases with photon energy as the fundamental absorption edge is approached corresponding to increasing CH. C

URBACH ENERGY

In all disordered materials, the allowed energy bands are not sharp but have band tails in the energy-gap. The exponential absorption edge, also known as the Urbach edge, is an important parameter for characterizing amorphous semiconductors and has implications for the performance of electronic devices. The Urbach edge energy, E n , is obtained by fitting the exponential edge to the relation : a oc exp(hv/Eu) This absorption region is due to transitions between valence tail band states and extended conduction band states. The Urbach energy depends on the deposition method and other parameters in a similar way to the other optical parameters. It has been reported that in unhydrogenated material [14], a rise of the growth temperature from

30 0 C to 300 0 C has the effect of decreasing E1; from 140 meV to 94 meV. However, the opposite behaviour has also been reported [16]: E1; rises from 50 meV (T s = 14°C) to 74 meV (T s = 223 0 C) before a more rapid increase to 120 meV (T s = 2400C). For unhydrogenated and sputtered a-GaAs, E n has been found to depend on the quantity (W/p)1/2, where W is the RF power and p the argon sputtering pressure [5], with a rapid decrease of E n being measured as a function of (W/p)1/2 from 260 meV to 145 meV, followed by a saturation. The presence of hydrogen in the deposition chamber reduces the value OfEn from about 140 150 meV (unhydrogenated samples) to a constant value of 90 - 95 meV when the H percentage in the sputtering mixture reaches « 10% [5]. Values OfE n even for the best hydrogenated amorphous GaAs are about one order of magnitude higher than those of crystalline GaAs [17].

Values of E1J measured in flash-evaporated samples prepared at room temperature range between (105 ± 5) meV in as-deposited and (110 ± 5) meV in annealed (T a = 250 0 C) material [18]. The optical absorption at hv < 0.5 eV, determined using both spectrophotometry and photothermal deflection spectroscopy, showed that the Urbach energy, Eu = (110±3) me V, is insensitive to annealing. It follows that since the a-GaAs films are essentially chemically ordered the defect states must be attributed to As and Ga dangling-bonds. D

PHONONS

Dl

Infrared Absorption Studies

The absorption coefficient due to absorption by phonons was measured in the range 30 - 600 cm"1 in 40 |iim thick unhydrogenated sputtered a-GaAs films by Stimets et al [19]. A shift was found in the main phonon frequency from 273 cm"1 (c-GaAs) to 254 cm"1. In addition, two other peaks, not found in c-GaAs, occur at lower frequencies («140 cm"1 and * 90 cm"1). A peak was also found at about twice the main frequency but it is not clear if it should be attributed to oxygen impurities or to two-phonon processes. Infrared spectroscopy [6,14,20] of RF sputtered hydrogenated a-GaAs reveals two hydrogen related broad bands at 510 - 540 cm"1 (wagging) and 1430 - 1460 cm"1 (near stretching). H atoms are in bridging positions between two Ga atoms and these bands amount to most of the hydrogen induced infrared absorption. Other sharp lines are interpreted in terms of Ga-H and As-H modes. At high hydrogen dilution of the sputtering mixture, (pH/ptot) ^ 15%, other peaks are observed and assigned to As-H and As-H2 bonds [14]. These experimental data are supported by theoretical modelling of the infrared absorption in hydrogenated or deuterated a-GaAs [21]. However, many other modes in the low frequency region were predicted but have not been reported in the literature. Unhydrogenated sputtered [6], MBE [7], and implanted [22,23] material gives similar results showing a Ga-As TO normal mode absorption peak at « 250 cm"1 (c-GaAs:268 cm"1). Two more broad bands centred at 70 cm"1 and 150 cm"1 are sometimes found.

D2

Raman Scattering

Zallen [23] discusses the finite-size effects on the optical properties of nano- and micro-crystalline GaAs compared to cc-GaAs. Raman shifts were measured in unhydrogenated sputtered thin films [24] as functions of the deposition parameters: argon pressure, RF power and substrate temperature. Both TO and LO phonons were found allowing the proportions of amorphous and crystalline components to be evaluated as a function of the deposition parameters. At low T s (20 0 C), for W = 200 W and pM > 2.5 Pa the Raman spectra peak appears at 250 ± 2 cm"1 with an FWHM « 5 5 cm"1, showing that the amorphous phase is about 90 - 98% of the whole. On increasing Ts, the LO peak shifts about 14-15 cm"1 and the TO about 16-17 cm"1, indicating that the amorphous phase decreases to 30 - 40 % . The influence of increasing the sputtering pressure, pAr, is towards almost completely (99%) amorphous films: crystalline LO and TO peaks are broadened and shifted toward lower frequencies E

PHOTOCONDUCTIVITY

Room temperature photoconductivity of RF sputtered hydrogenated cc-GaAs under illumination of « 1015 photons / cm2 s at hv = 1.95 eV increases by a factor of 5 when the hydrogen pressure rises from 0 to 5.2 x 10"3 torr [3]. The photoconductivity spectral edges shift to higher energy with hydrogenation in a consistent way with the absorption edges. The spectral response of photoconductivity has been measured in RF sputtered Ct-GaAs films, with and without hydrogen [25]. The peak values of the normalized spectral photoconductivity (hv « 2 eV) plotted as a function of hydrogen pressure, pH, increase when (p H / p tot ) < 0.1 Pa and then tends to a constant value. The highest measured value of the normalized photocurrent is iph^ » 8 x 10"9. The maximum value of the ratio (oph / O6^) is 204 in samples deposited at T s = 20 0 C with about 25% hydrogen within the deposition chamber and an RF power of 200 W. The recombination lifetime of photogenerated carriers near the maximum of the photoresponse is « 10"9S, several orders of magnitude lower than in the best a-Si. These results are interpreted in terms of the possible effects that hydrogenation can induce in a compound semiconductor: a less disordered network, a reduction of the band tail depth, a decrease of bond angle fluctuations, but an ineffectiveness of hydrogen on the density of wrong bonds. An anomalous photovoltaic effect and negative photoconductivity were observed in obliquely DC sputtered unhydrogenated thin films [26] deposited at T s < 200 0 C (compared to the crystallization temperature of 4000C). Both effects can be explained if photocarriers with negative mobilities are present in the extended states. In the presence of carrier density gradients, the diffusion of the photocarriers in the localized and extended states can lead to a photovoltage of several hundred volts in a single photocell with a length of 10 mm. F

OTHER DEPOSITION METHODS

Matsumoto et al [27] found in MBE grown a-GaAs that the refactive index n (A = 1.5 \xm) approaches that of C-GaAs (n = 3.4) monotonically as the composition ratio (r = As/Ga) approaches stoichiometry at r = 1. For films strongly As rich (r = 1.8), n approached a value of 4.6. The optical absorption coefficient shifts towards higher energies at increasing values of r. The Tauc optical gap, E0, decreases with r going from 1 eV (r = 1) to about 0 . 7 e V ( r = 1 . 8 ) and is always smaller than that of a bulk crystal by «0.4 eV, indicating the role of disorder on E0.

Segui et al [28] use a mixture of arsine (AsH3) and trimethylgallium (TMG) for plasma deposition of ot-GaAs. The incorporation of hydrogen depends on the ratio (pMm IP1-M0 ) and leads to a decrease of the absorption coefficient as a function of hv, when the ratio is « 10. Aguir et al [9] prepared a-GaAs films by sputtering a gallium target and decomposing arsine in an RF glow discharge. The optical gap is constant (« 1.35 eV) for a gallium content < 40% and decreases rapidly before stoichiometry is reached. For films containing a gallium excess, a is not modified by plasma hydrogenation. The absorption coefficient for stoichiometric and As rich samples decreases faster when using Ar rather than H2 as diluent. Excess hydrogen tends to deteriorate the properties of these films. G

CONCLUSION

This Datareview has shown that problems encountered in the preparation and characterization of compound a-III-V semiconductors are more complex than those found in elemental cc-Si. Because properties are dependent on stoichiometry, accurate control of the deposition method is necessary to result in well characterized material. Results presented in the literature are rarely directly comparable, due to different methods of deposition and a lack of systematics in collecting experimental results as functions of deposition conditions. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] [II] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22]

S.K. Barthwal, K.L. Chopra [ Phys. Status Solidi A (Germany) vol.36 (1976) p.345 ] A. Gheorghiu, MX. Theye [ Philos. Mag. B (UK) vol.44 (1981) p.285 ] W. Paul et al [ Proc. 7th Int. Conf. on Amorphous and Liquid Semicond., Edinburgh, UK, 1977 (Stevenson, Dundee, 1977) p.467 ] L. Alimoussa, H. Carchano, J.P.Thomas [ J. Phys. (France) vol.42 (1981) p.C4-683 ] R. Murri, L. Schiavulli, N. Pinto, T. Ligonzo [ J. Non-Cryst. Solids (Netherlands) vol. 139 (1992) p.60] M.I. Manssor, E.A. Davis [ J Phys. Condens. Matter. (UK) vol.2 (1990) p.8063 ] G. Monnom et al [ J. Non-Cryst. Solids (Netherlands) vol.83 (1986) p.91 ] J. Bandet et al [ Philos. Mag. B (UK) vol.58 (1988) p.645 ] K. Aguir, M. Hadidou, P. Lauque, B. Despax [ J. Non-Cryst. Solids (Netherlands) vol. 113 (1989) p.231] A. Gheorghiu, M.L. Theye [J Non-Cryst. Solids (Netherlands) vol.35-36 (1980) p.397 ] S. Guita et al [ J. Phys. (France) vol.47 (1986) p.C8-427 ] S. Najar, K. Seedek, F. Lalande, H. Carchano [ Mater. Res. Soc. Symp. Proc. (USA) vol.47 (1985)p.283] T. Flohr, R. Helbig [ J Non-Cryst. Solids (Netherlands) vol.88 (1986) p.94 ] H. Carchano, K. Seedek, J.L. Seguin [Proc. Euroforum - New Energies, Saarbrucken, Austria (H.S. Stephens and Associates, Brussels, 1988) vol.3 (1988) p. 197 ] A. Carbone et al [Nuovo Cimento (Italy) vol.13 (1991) p.571 ] S. H. Bakert et al [ J. Phys. C Condens. Matter. (UK) vol.5 (1993) p.519 ] S. R. Johnson, T. Tiedje [ J Appl. Phys. (USA) vol.78 (1995) p.5609 ] M. L. Theye et al [ Philos. Mag. B (UK) vol.52 (1985) p.325 ] RW. Stimets et al [Proc. 5th Int. Conf. Amorph. LiquidSemicond, Garmisch, Germany (Taylor and Francis, London, 1974) p. 1239 ] Z.P.Wang, L. Ley, M. Cardona [ Phys. Rev. B (USA) vol.26 (1982) p.3249 ] B.K. Ghosh, B.K. Agrawal [J Phys. C, Solid State Phys. (UK) vol.19 (1986) p.7157 ] R. Zallen et al [ J Non-Cryst. Solids (Netherlands) vol. 114 (1989) p.795 ]

[23] [24] [25] [26] [27] [28]

R. Zallen [J. Non-Cryst. Solids (Netherlands) vol.141 (1992) p.227 ] LD. Desnica et al [ J. Non-Cryst. Solids (Netherlands) vol. 170 (1994) p.263 ] U. Coscia et al [ J. Non-Cryst. Solids (Netherlands) vol. 194 (1996) p. 103 ] H. Reuter, H. Schmitt [ J. Appl. Phys. (USA) vol.77 (1995) p.3209 ] N. Matsumoto, K. Kumabe [ JpnJ. Appl. Phys. (Japan) vol. 19 (1980) p. 1583 ] Y. Segui, F. Carrere, A. Bui [ Thin Solid Films (Switzerland) vol.92 (1982) p.303 ]

4.5

Intra- and in ten alley deformation potentials for electrons in GaAs S. Zollner and M. Cardona September 1996

A

DEFORMATION POTENTIALS

A deformation potential (DP) is a quantity proportional to a matrix element of an operator belonging to a crystal deformation (which may be caused by pressure, strain, or a phonon displacement) between a final and initial electron or hole state [I]. It usually has the units of eV or eV/A. The DP is called an intravalley DP if the wave vectors of the final and initial states are within the same electron valley; otherwise, it is called an intervalley DP. Intravalley DPs cause shifts, splittings, and/or intravalley scattering of carriers, whereas intervalley DPs are responsible for the scattering of electrons to a different valley. We note that an intravalley DP can cause an interband splitting, if several valleys (i.e., L or X) are degenerate. A DP is called an intraband DP if the band indices of the final and initial states are the same; otherwise, it is called an interband DP. The DP for an optical gap is the difference of the absolute DPs for the valence and conduction band states and describes the change of the gap due to pressure, strain, or an optical phonon. The bulk of this review deals with intraband intravalley and intervalley DPs for the conduction band of GaAs and for the shifts and splittings of optical gaps. We also mention briefly two-phonon DPs. The intravalley DPs for holes in the F and L valence bands are discussed in the following Datareview by Adachi [2]. Deformation potentials relate shifts and splittings to the elements of the 3><3 strain tensor e defined by f' = (1+e)?, where r1 is the strained and r the unstrained coordinate [3,4]. It is convenient (but not necessary) to break up an arbitrary strain into its irreducible components, which are (for a cubic crystal) hydrostatic pressure (with F 1 symmetry), [100] strain (Fi2), and [111] strain (F15), see [5]. An experiment, on the other hand, usually applies a stress defined by a 3x3 stress tensor X [4,5]. The strain and stress tensors are related through the 6x6 compliance tensor S [4], which has three independent components S11, S12, and S14. Yu and Cardona [5] list the strain tensors for [100] and [111] uniaxial stress. The important conversion factors are: Hydrostatic pressure: - P O O X=

0 ,0

-P

0 , Tre = 3e H = -3P(S 11 + 2S12)

0 -Pj

Tensile uniaxial stress along [100]: X 0 O^ X = O O O ,

Io o o,

Tre = 3eH = -3P(S11

+

2S12X), e^ -

Cyy

= 3es = (S11 - S12)X

Tensileuniaxial stress along [111]: 1 1 1

X =I

1 1 1,

Tre = 3eH = (S11 + 2S12)X, e^ = -U 44 X

U 1 lj Different conversion factors apply to the biaxial strain found in pseudomorphically strained layers (e.g., AlAs on GaAs or Ge on Si): see [74]. We note that there is also a different convention, where the strain becomes a 6-dimensional array [5]. In this case, E4=S44XzG. B

INTRAVALLEY DEFORMATION POTENTIALS FOR ELECTRONS IN GaAs

Bl

T-Point

The absolute shift of the conduction band minimum under pressure (or strain) is equal to the absolute deformation potential a times the trace of the strain tensor e: AECB(r) = a T r e

(1)

In principle, the same a gives the coupling of the electrons with the longitudinal acoustic phonons [6]. However, while the coupling to acoustic phonons is uniquely defined, that to a strain is not for a crystal extending to infinity: an absolute average potential cannot be defined nor its derivative versus strain. In a finite crystal, such a potential can be defined, but depends on the surface orientation (electron affinity, photothreshold) and hence cannot be used to calculate electron-LO phonon coupling constants. Methods to calculate absolute hydrostatic DPs which can be used to determine electron-phonon coupling constants (i.e., including screening) have been given [7-13]. Such calculations yield very small values of a for holes at the top of the valence band (less than 1 or 2 eV). Hence, it is a good approximation to set a for electrons equal to the DP of the corresponding gap [14,15] which can be obtained easily by measuring luminescence, reflectivity, or absorption under pressure [16-20] (or in pseudomorphically strained quantum wells [21]) or from band structure calculations of solids under pressure [20,22-25]. We therefore list in TABLE 1 both the deformation potentials for the gap and the absolute deformation potentials for the conduction band at F, which should be used in the calculation of electron-LA phonon interactions. Some methods, e.g., transport measurements, can only determine the magnitude of the DP, but not its sign. The absolute deformation potentials have been calculated using (now obsolete) empirical pseudopotentials [26] (EPM), ab initio pseudopotentials (PSP) based on the local densityfunctional theory (LDA) [8-10,13] and with the M y relativistic linear-muffin-tin orbitals (LMTO) method [7,27] within the atomic sphere approximation (ASA). The most probable value lies between 7 and 9 eV (see TABLE 1). The experimental values were determined from transport measurements [28-38] on high-mobility bulk GaAs samples and 2D-GaAs/AlGaAs heterostructures, free-electron absorption [39,40], deep-level transient spectroscopy (DLTS), luminescence of transition metal impurities in GaAs and doping of GaAs with Te [43] or Si [44]. They fall into two groups around 7-9 eV and 13-14 eV, but the values depend somewhat on the model used for fitting the data [45-48], for example on the type of wave functions, the screening mechanism, the phonon modes in the AlGaAs barrier [37], or the value of the piezo-electric

constant used in the simulation. Since the DP for deep levels may differ from that of the band, it may be difficult to derive absolute DPs from the studies of d-impurities [49]. Not all available experimental data for a are given in TABLE 1. The coupling of electrons OfF1 symmetry to optical phonons is forbidden by group theory, at least at or close to F. This is also true for their coupling to a shear, i.e., to transverse acoustic phonons or to [111] or [100] strain. Therefore, the optical and transverse acoustic deformation potential constants are zero. Interband deformation potentials for [lll]-strain with F15 symmetry connect the lowest conduction band at F, i.e., F1, with the p-bonding valence band and the p-antibonding conduction band at F with F15 symmetry [50]. These have been calculated with the EPM and tight binding (LCAO) methods [26] and are important for spin-relaxation phenomena [51] and the recently observed piezooptical activity of GaAs [52]. TABLE 1. Deformation potentials for electrons in GaAs at F, in eV. a -7.0 -7.99 -8.4±0.9 -8.7±0.3 -8.4±0.2 -9.8 -9.9 -8.1 -8.3±1 -9.0±l -8.34 -7.3 -8.8 -8.47 _±7 ±12 ±(13.5±0.5) ±(9.5±0.5) ±8.5 _±8 ±17.5 ±(11±1) ±13.5 -9.3 -7.7 ±15.7 ±(12±1) -8.5 _±7 ±(11.5±0.5)

I

Method Deformation potential for the band gap

absorption photoluminescence piezo-electroreflectance absorption reflectivity absorption empirical pseudopotential method (EPM) ab initio pseudopotentials ab initio pseudopotentials linear-muffin-tin orbitals method (LMTQ) tight binding Absolute deformation potential for the conduction band ab initio pseudopotentials LMTQ (screened) ab initio pseudopotentials transport transport transport transport transport transport transport transport (2D) transport (2D) d impurity levels d impurity levels infrared absorption transport (2D) doping with Si transport (2D) I transport (2D)

|

Reference [14,15] [21] [16] [17] [20] [19] [22] [23] [8,9] [24] [25]

[8,9] [7] [13] [15,28,33] [46] [45] [3JJ [47] [48] [32] [30] [29] [41] [42] [40] [34] [44] [35] 1 [36]

B2

L-Point

The four degenerate conduction band states at L split under [111] (but not [100]) uniaxial strain into a triplet and a singlet. The interband splittings, determined by the shear deformation potentials S d and H11 in the Herring-Vogt notation [53], have a 3:-l ratio. S u is also denoted by E2 [54]. An overview of the different notations has been given by Kane [55]. The associated interband shift for valley n for an arbitrary strain (relative to the unstrained energy level) is given by [9,54]

AEfl = E 2 n | e - i T r e j

ft

(2)

where e is the strain tensor and the unit vector fi is the direction of the valley. The hydrostatic deformation potential a for L-electrons, see EQN (1), is also denoted by E1 and given by a = E

i = }

S

u

+ S

d

(3)

It is derived from the average shift of all four states [53]. As in the case of the conduction band at F, it is a reasonable approximation to set a equal to the deformation potential for the indirect gap. Then, a can be obtained from band structure calculations under hydrostatic pressure [23,24,56]. An experimental value of a = 3 eV was derived from photoreflectance studies of GaAsZAlxGa^xAs quantum wells under high pressure [57]. S u has been calculated using ab initio pseudopotentials [9], LMTO [24], and EPM [58]. Calculations of Hd and S u turn out to be sensitive to the internal displacement parameter £ which describes the change of the basis vectors with strain [9]. Therefore, the accuracy of the atomic sphere-approximation (ASA) is limited for this case. Experimental data for S11 are available from piezoresistance [59] and hot-electron photoluminescence measurements [60], see TABLE 2. TABLE 2. Deformation potentials for electrons at the L-point in GaAs, in units of eV. S11 = E 2

(j 5

a = E1

14.26 19.6±3 14.5±1.5

18.5

±92 -3 -2J -33 -3J

±17

-20 ^32 -21.7

Method

Reference

ab initio pseudopotentials piezoresistance hot-electron photoluminescence

[9] [59] [60]

transport photoreflectance dielectric theory linear-muffin-tin orbitals method (LMTQ-ASA) ab initio pseudopotentials empirical pseudopotentials (EPM) empirical pseudopotentials (EPM) empirical pseudopotentials (EPM)

[64] [57] [56] [24] [23] [58] [63] [61]

The interaction between the electrons at L and optical phonons at F is given by the optical deformation potential d£, which is often written as

D = D1K = O* a

W

o

where Z0 is the lattice constant [61,62]. d^ has been calculated using the EPM [58,61,63] and has also been determined from transport measurements [64], see TABLE 2. B3

X-Point

The deformation potentials Hd and S u (or a, E1, and E2) for the three X-valleys are defined just as in the case of the L-valleys, see EQNS (1) to (3). However, for the X-point, the interband shifts are nonzero for [100] strain and vanish for [111] strain. Experimental values for E2 have been obtained using hot-electron photoluminescence [60] and 2D-transport under uniaxial stress [65], for a with absorption measurements of the indirect gap at X [17]. AlAs can be studied using timeresolved photoluminescence, since it is indirect [66,67]. The splitting of the X-valleys in GaAs under [100] strain, i.e., E2, has been calculated with pseudopotentials [9]; the shifts of the indirect gap at X under hydrostatic pressure relative to the valence band top, a, with dielectric theory [56]; pseudopotentials [9,23,59]; and LMTO [24]. See TABLE 3. TABLE 3. Deformation potentials for electrons at the X-point in GaAs, in units of eV. S11=E2 8.61 6.5±1 5.8±0.1

a = E1

±93 6.3 9.6±1.8 +1.7 +O6 +L6 +1.7 -0.09

B4

Method ab initio pseudopotentials hot-electron photoluminescence time-resolved photoluminescence (AlAs) transport empirical pseudopotentials (EPM) transport (2D, IV-characteristics) absorption dielectric theory linear-muffin-tin orbitals method (LMTO-ASA) ab initio pseudopotentials ab initio pseudopotentials

Reference [9] [60] [66] [64] [59] [65] [17] [56] [24] [23] [13]

Lines of High Symmetry

Deformation potentials can be defined at any point in the Brillouin zone, not just at points of high symmetry. The corresponding shifts can be measured, for example, using angle-resolved photoemission [68]. The dispersion of DPs in the conduction and valence bands of GaAs, calculated using empirical pseudopotentials, is shown in [58,63,71]. EPM results compare favourably with LMTO results; see [69]. C

DEFORMATION POTENTIALS FOR INTERBAND TRANSITIONS IN GaAs

Cl

F-Point (E0 Transition) and Indirect Transitions

The deformation potentials for the E0 gap are similar to the absolute deformation potentials for F electrons. Therefore, they have been discussed in Section Bl. Also, the DPs for indirect

transitions between F and X or L are about the same as the absolute DPs for the electrons at X and L listed in the previous section. C2

L-Point (E1 and E^A 1 Transitions)

Since optical interband transitions near L and along most of the A direction produce strong peaks in the dielectric function (the so-called E1 and E1-^A1 transitions), it is also of interest to know the effect of hydrostatic pressure, uniaxial strain (acoustic phonons), and of optical phonons at F on these transitions. The shifts OfE1 and E^A 1 due to hydrostatic pressure are described by the hydrostatic deformation potential (a) - D11 / ^J; see TABLE 4. TABLE 4. Hydrostatic deformation potentials for E1 and E1H-A1 transitions in GaAs, in eV. a=D Vt/3 -4.3±0.4 -6.9±0.2 -8.3 -5,5 -5.82 -4.8±0.5

Method piezo-reflectance reflectivity empirical pseudopotential method (EPM) empirical pseudopotential method (EPM) linear-muffin-tin orbitals method (LMTO-ASA) ellipsometry under uniaxial stress

Reference [16] [20] [16] [22] [20] [72]

We note that there are three contributions to the splittings of the E1 and E1^A1 doublet due to uniaxial strain [71]: 1.

The four degenerate <111> directions split under [111] (but not [100]) strain (interband splitting).

2.

The doubly degenerate valence band at L (neglecting spin orbit splittings) splits under [111] and [100] strain (intraband splitting). If spin-orbit splittings are included, the (apparent) value OfA1 will change.

3.

Also, there are spin-exchange terms, which we neglect here, since they are small.

Diagrams showing the shifts and splittings of the different bands are given in [72,73]. In the previous (2nd) edition of this book, we discussed interband and intraband splittings in our Datareview. In this edition, the intraband splittings are discussed in the Datareview by Adachi [2], since they are related to the hole deformation potentials D35 and D33. The optical deformation potentials d^and d35o are also given by Adachi [2]. A [ I I l ] strain shifts the [111] -direction from the three (1 TT)-directions in a manner similar to that in EQN (2), with the energy shifts in a 3:-l ratio. The corresponding deformation potentials are usually given in Kane's notation [55] as, D15 whereby D15 must be plugged into EQN (2) through the substitution [55] E2 = & D15 2 2 x

(5)

to find the strain splittings of the E1 interband transitions. D15 has been determined experimentally using electroreflectance [16,73] and ellipsometry [71] under uniaxial stress and calculated using the EPM [16,22,58]; see TABLE 5. The doubly-degenerate valence bands at L (or along A) are split by the [111] strain. A [100] strain does not remove the degeneracy of the four L-points, but it splits the doubly-degenerate valence band maximum at L. These intraband splittings are described by the deformation potentials D35 and D33, see [2]. The resulting eigenvalues of the effective strain Hamiltonian (i.e., the energies of the E1 and E1^A1 transitions under strain) are given in [71,72,74]. They are presumed to be correct in this work as well as in [72-74], but there are misprints in [16,71]. For [100] stress, these eigenvalues are [16,73,74]

E1, E 1+ A 1 = E10

+

J • (AE^) 2 ,

^ l + AEH ± ^

where AEH = /3D 1 C n , A E ^ * = ^/6D3 e s

In the large-shear approximation, the intraband shear splittings are much larger than the spin-orbit splitting (AE s m t r a » A1). This approximation is appropriate for Si [72]. For GaAs, it is usually feasible to use the small-shear approximation (AE s m t r a « A1), see [16,71], instead of the accurate expression above. For [111] stress, the E1 and E1^A1 transitions split into a singlet and a triplet. The energies of the singlet are [16,72,73] E18 = E10 + AE n + A E 8 ^ ,

(E1 + A1)8 = E10 + A1 + AEH + A E 8 ^ ,

where the hydrostatic shift AE n is defined just like for [100] stress and the interband shift of the singlet is A E ^ = fiD^e^. The triplet energies in the small-shear approximation are [16,71]

E18 1

= E10+ 1

/A-pinterV2

AEH - I AE8*" H

3

^1 . ^1, - ^1 . «, . _ H

S

1

r^-L, ^

^8

T

^L-L

where the intraband shear splitting is AE^* = V^Di5exy The analogous results for the large-shear case are [71]: „0 Aj-* I A T - into" • A r> intra ES ! = E 1 + AEH - - AE8 ± AE8

TABLE 5. Acoustic deformation potentials D 1 5 describing the interband splittings of the E 1 and E1-^A1 transitions due to a [111] shear; s e e E Q N (5). D15 9.2±1.0 8.5±0.8 12.0±0.7 7.9±0.5 8.3 5.5 12.4 9

C3

Method piezo-electroreflectance (77 K) piezo-electroreflectance (300 K) ellipsometry under uniaxial stress (300 K) EPM(L) EPM(L) EPM (A) EPM(L) I EPM(A)

Reference \\6] [73] [71] [16] [22] [22] [58Jl] I [58]

X-Point

The X5 - X1 gap at the X-point is also split by a [100] strain (but not by a [111] strain) in a manner somewhat similar to the splitting of the E r gap by [111] strain. The interband splitting of the [100] valley from [010] and [001] is given by EQN (2), where E2 in Kane's notation [55] is replaced by

Values OfD13 = -10.4 eV have been obtained with the EPM method [58]. The intraband splittings in the valence band at X are discussed in [2]. D

INTERVALLEY DEFORMATION POTENTIALS FOR ELECTRONS IN GaAs

Dl

Introduction and Notation

A scattering process of an electron by a phonon is called an intervalley scattering (IVS) process, when it connects two different minima in the conduction band or maxima in the valence band (e.g., r -> L, F - X, L - X, X - X', L - L', etc). In GaAs, the wavelength of the scattered phonon is of the order of the lattice constant, i.e., the crystal momentum transfer Q is larger than about one sixth of a reciprocal lattice vector. According to Eqn (3.6.7) of ConwelFs book [62], the matrix element for this process is proportional to the intervalley deformation potential D (IDP, in units of eV/A). The expression also contains the intervalley phonon energy Q and the mass M=Vp of the primitive cell (i.e., the mass of one Ga plus one As atom) [75]. The selection rules for IVS processes have been given by Birman, Lax, and Loudon [76]. They are discussed in [77] in the light of new arguments about the symmetries of electrons and phonons at the X- and L-points. Selection rules along lines of high symmetry are given in [75]. The conditions for these selection rules to be valid are usually not fulfilled, since energy conservation rules out scattering processes between electrons exactly at high-symmetry points. Nevertheless, it is generally assumed that the matrix elements for intervalley transitions are nearly independent of the phonon wave vector Q. Therefore, the matrix elements can be integrated over all possible

final states in a spherical energy band resulting in ConwelFs expression for the IVS time x:

I = X

Mvn 2m v3/2

(7)

[(N + 1) y'E - AE - Q + N^E - AE + o]

^7C^2pQ

where N v is the number of (final) valleys, m v the effective mass of the final valley, p the density of the crystal, N the occupation number of the intervalley phonon, E the energy of the electron in the initial valley, and AE the energy separation between the two valleys. See eqn (3.6.14) in [62] and also [77,78]. D2

Theoretical Results

At least five independent calculations of intervalley deformation potentials for GaAs using empirical band structure methods have appeared in the literature [75,77,79-83]. They are based on parameterized lattice dynamical models for the phonon eigenvectors (note that for the LO and LA phonons at X the eigenvectors are fixed by symmetry) and empirical pseudopotential or tightbinding electron wave functions. More recently, two ab initio calculations have been performed [84-87]. Their results are given in TABLE 6. The agreement between the different methods is rather good. The IDPs given here are for the screened case [79,80]. Values calculated from the EPM using local and nonlocal pseudopotentials are in good agreement [82,83]. There is, however, some disagreement on how the form factors should be extrapolated to q = 0 [75,83,88]. IDPs for hole-phonon scattering have also been reported [84-86]. TABLE 6. Calculated intervalley deformation potentials for GaAs in units of eV/A.

r-L

I r - xIL^L

LA+LO 2.6±0.3

LQ

I

X -» X '

LA+LO

2.8±0.1

1.2

LO

I

4J)

3A

_ L ~* X

I TA

TA

08

1Ref•

I LA+LO

0.8

I TQ

1.8

1.8

5.0±0.5

3.2

_37 JA 3.86

[77]

[80] [81]

4^0 34 4£7

1 4.134

1.0 3_8 036

[

I

6,3

1

~

0.0

I

Z9 I

1

1

3.7 2.8 2.32

1

| 2.2 I

1

[82] [83] [84]

[85]

Herbert [79] has found that the IDP depends on the phonon wave vector Q. Detailed plots of the dispersion of the IDPs are given in [75,89]. There it has been shown that the transverse acoustic phonon needs to be included for F - X scattering, whereas the assumption of k-independent matrix elements gives smaller (but still significant) errors for F - L scattering. Because of this Qdependence, IDPs extracted from experiments are effective values depending on the experimental conditions, like the bath temperature [90], the direction of the electric field, or the incident laser energy [78,91,92]. D3

Experimental Results

Estimates of IDPs have been obtained from a number of different optical and electrical measurements. The evaluation of the data is usually difficult, as several different scattering mechanisms (carrier-carrier, intervalley, intravalley, impurity, and possibly alloy scattering [93])

are possible and compete with each other. Usually, a numerical model based on rate equations or Monte-Carlo simulations is fitted to the data. There are almost as many sets of IDPs as there are experiments, disagreeing with each other by more than one order of magnitude. We have therefore selected only a few papers which we believe to give the most reliable results. The IDP for F - L scattering has been determined by various recent optical experiments; see TABLE 7. These include subpicosecond luminescence spectroscopy [94,95], hot-electron photoluminescence [96-106], time-resolved Raman spectroscopy [107], broadening of the direct exciton under hydrostatic pressure [108,109], and infrared four-wave mixing [HO]. Mirlin and coworkers [60] have used the method of hot-electron luminescence under uniaxial stress to determine the L - L' IDP relative to the F - L. A better picture of intervalley scattering has evolved recently, since Monte Carlo simulations [111-119] are better able to keep track of the many particles generated in femtosecond laser pump and probe and four-wave mixing experiments [90,95,120-125]. An important issue is also the inclusion of the full band structure for electrons and phonons [126] and the Q-dependence of the electron-phonon matrix elements [78,83]. The values obtained by Collins and Yu [127] for the F - L process using non-equilibrium phonon spectroscopy have been criticized in [128]. F - L IDPs obtained from Monte-Carlo [33,81,129-133] and Green's function [134,135] simulations of electrical measurements [136-138] are in the same range of values, but the velocityfield curves are not very sensitive to this parameter [117]. Similar values are listed in [139] from an analysis of impact ionization coefficients. Combined optical and electrical methods (transient grating and noise experiments) have been used in [140], resulting in a lower F - L I D P . The IDP for F - X scattering is generally assumed to be larger than the F - L IDP. Little experimental evidence exists about IDPs for scattering between other valleys, see TABLE 7. TABLE 7. Experimentally determined absolute values of IDPs in GaAs, in units of eV/ A. T-L 6.5±1.5

I

T-X

I

L-L'

I

X-L

I

X-X

3.5

QOOj

7±2 8±1 9±2 _>5

15±3

[107] [60,96,97] [106] [116]

5±1

_9

__I1I2]

7±1 <1.5 3.25 JO JO 2.85 1.8 J

IUl

2.75±0.2

10 10 10 10 8

10 10 5 5

5 9 3J6 1 1.8...10

7 9 10 10 10

J

[UO] [127] [81] [33,131,135] [129,130] [134] [132] [139]

[140] 10±l 4.8±0.3

5±0.5

I Reference [94]

1

[108] [109]

I

I

1

I [118]

When experimental IDPs are given in the literature, it is generally assumed that the band structure of a valley can be described in an effective mass approximation. This is certainly not the case for the X-valley with its camelback structure [141,142], where f and g processes (as in silicon) are possible. Also, because of the camelback structure it is not clear if there are 3 or 6 X-valleys. E

ELECTRON-TWO PHONON INTERACTIONS

Processes in which an electron scatters through creation or annihilation of two phonons in a single (renormalized) vertex can be important in transport phenomena [143,144] and in second-order Raman scattering (i.e., by two phonons). The interaction constant is usually represented by a deformation potential D1, which describes the second derivative of the interaction energy with respect to the phonon displacement, and multiplied by the square of the lattice constant in order to obtain the dimensions of energy (eV). From resonant Raman scattering deformation potentials corresponding to the E0 and E1 gaps have been obtained for GaAs, (see [70] Table 2.11). They represent the effect of combinations of two phonons OfF1 symmetry for the DP D 1 (for the E0 and E1 gaps), or OfF15 symmetry for D15 (E0 gap) or for D530 (E1 gap). F

CONCLUSIONS

An empirical rule due to W. Paul [145] states that DPs for the same type of process are similar for most semiconductors. Therefore, the DPs listed in this Datareview can be used (with reasonable accuracy) for materials similar to GaAs, such as Ge, InP, AlAs, and, possibly, even Si. By the same token, if we do not list a particular DP for GaAs, values for similar semiconductor materials can be used. A few DPs, such as those describing the shifts of the direct gap with hydrostatic pressure, are known fairly accurately. Most DPs, however, particularly absolute DPs and intervalley DPs, are difficult to calculate or measure. Therefore, there is considerable scatter in the literature. Uncertainties of 50% or more are common. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] II1] [12] [13] [14] [15]

For a general introduction, see G.L. Bir, GE. Pikus [ Symmetry and strain-induced effects in semiconductors (Wiley, New York, 1974) ] S. Adachi [ Datareview in this book: 4.6 Intravalley deformation potentials for holes in GaAs ] L.D. Landau, E.M. Lifshitz [ Theory of Elasticity (Pergarnon, New York, 1986) ] J.F. Nye [ Physical Properties of Crystals (Clarendon, Oxford, 1985) ] P. Y. Yu, M. Cardona [ Fundamentals of Semiconductors (Springer, Berlin, 1996) ] J. Bardeen, W. Shockley [ Phys. Rev. (USA) vol.80 (1950) p.72-80 ] M. Cardona, N.E. Christensen [ Phys. Rev. B (USA) vol.35 (1987), p.6182-94; vol.36 (1987) p.2906 (E) ] CG. van de Walle, R.M. Martin [ Phys. Rev. Lett. (USA) vol.62 (1989), p.2028-31 ] C G vande Walle [Phys. Rev. B (USA) vol.39 (1989)p. 1871-83 ] R.D. King-Smith, RJ. Needs [ J. Phys., Condensed Matter vol.2 (1990) p.3431-44 ] R. Resta, L. Colombo, S. Baroni [ Phys. Rev. B (USA) vol.41 (1990) p. 12358-61 ] A. Franceschetti, S.-H. Wei, A. Zunger [ Phys. Rev. B (USA) vol.50 (1994) p.7797-801 RA. Casali [ Solid State Commun. (USA) 1996 (in press) ] H Ehrenreich [ Phys. Rev. (USA) vol. 120 (1960), p. 1951-63 ] D.L. Rode [ Phys. Rev. B (USA) vol.2 (1970) p. 10 12-24 ]

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4.6

Intravalley deformation potentials for holes in GaAs

Previous Page

S. Adachi July 1995

A

INTRODUCTION

The deformation potential (DP) is an important electron-lattice (phonon) parameter related to many physical phenomena in semiconductors. A review of the importance of the DPs and their definition and methods of calculation was given by Bir and Pikus [I]3 Blacha et al [2] and Cardona and Zollner [3], We discuss here the intravalley DPs for holes in the F and L-valence bands of GaAs. B

T POINT

Bl

Hydrostatic and Shear Deformation Potentials

The strengths of the hole scattering mechanisms are determined essentially by the valence-band DPs, namely a, b and d (Pikus-Bir's notation) [1,4]. The parameter a is the hydrostatic-pressure DP, and b and d are the shear DPs. Although the shear DPs b and d have been measured for many crystals, it is difficult to obtain values for the hydrostatic potential a since most experiments measure changes in energy gaps and their related effects rather than absolute shifts of the band edges [5,6]. Lawaetz [7] has, however, proposed a theoretical expression based on the dielectric band theory of Phillips [8] which allows a to be estimated with reasonable accuracy. TABLE 1 summarizes the theoretical and experimental DPs a, b and d reported for GaAs. TABLE 1. Valence-band DPs a, b and d at the T point in GaAs (in eV). Hydrostatic

Shear

Remark

Ref

a

b

d

Theoretical

-4J

-21

-42

(EPM)

[2]

-3J)

-4J6

(LCAQ)

[2]

-23

-2/7

(LCAQ)

[2]

-L43

-4J54

(LMTQ)

[9]

-4J8

(LMTO)

[9]

(TBA)

[10]

(LMTO)

[11]

A3 -L6

AA6

-2J9

-4/77

(TBA)

[12]

AJO

A23

(AIPC)

[13]

(AIPC)

[6]

AA

-12

I

I

|

QT)

I [14]

TABLEl. (Continued). Hydrostatic

Shear

a

Remark

b

d

Ref

Theoretical:

-17

(EPM)

[15]

Experimental: -1.96 ± 0.1

-5.4 ±0.3

(PL)

[16]

-1.7 ±0.2

-4.4 ± 0.6

(PZR)

[17]

-2.0 ± 0.2

-6.0 ± 0.4

(ER)

[18]

-L7S

-£55

(PZR)

[19]

-1.66 ±0.1

-4.52 ±0.25

(ER)

[20]

-LO

(CT)

[21]

•0.7±1.0

(DLTS)

[5]

-2.00 ±0.2

1

-4.43 ±0.6

|

(PR)

|

[22]

EPM=empirical pseudopotential method, LCAO=linear combination of atomic orbitals, LMTO=linear muffin-tin-orbital method, TBA=tight-binding approximation, AIPC=ab initio pseudopotential calculation, DT=dielectric theory, PL=photoluminescence, PZR=piezoreflectance, ER=electroreflectance, CT=carrier transport, DLTS=deep-level transient spectroscopy, PR=photoreflectance.

B2

Acoustic-Mode Deformation Potentials

Lawaetz [7] defined an effective acoustic-mode DP, 3 eff , appropriate for low-field transport in p-type materials with Ge-like valence bands. This parameter can be written as C

11/2

i

where C, and Ct are spherically averaged elastic coefficients given by C1 = 1(3C 1 1 + 2C12

c

t

= )(Cn

-

c

i2

+

+

4C44)

3C

44>

(2a)

(2b)

Although the deformation potential Heff is a somewhat phenomenological average over the longitudinal and transverse acoustic phonon modes, it is widely used to model the performance of semiconductor devices and is particularly useful for device engineers. The deformation constant H^ is related to the phenomenological acoustic DP, Each, proposed by

Wiley and DiDomenico [23,24] by the equation

Ei, = M . ^

(3)

where P = C1ICx. If we use a = -2.7 eV, b = -1.7 eV, d = -4.55 eV and a set of the elastic stiffness values C11 = 11.885 C12= 5.38 and C44= 5.94xlO n dyn/ cm2, we obtain H eff = 6.7 eV and E ach = 3.6 eV for GaAs [14,25]. An estimate in [24] also finds E ach = 4 eV. A time-resolved transport analysis of highly nonequilibrium photoexcited carrier-phonon systems in GaAs was presented by Potz and Kocevar [26]. The theoretical model used by them was in qualitative agreement with hitherto unexplained experimental results from time-resolved transmission spectroscopy. The E ach of 3.5 eV was obtained from this analysis [26]. Ruhle et al [27] carried out time-resolved photoluminescence experiments in the picosecond regime to reveal the density dependence of the cooling of a photogenerated electron-hole plasma in AlxGa^xAs with x = 0 - 0.44. The DP scattering in this experiment determined the energy-loss rate at lower temperatures (-10 K) and showed a pronounced dependence on x. They found a variation of E ach from 4.8 ± 1 eV (x = 0; GaAs) to 9 ± 1.5 eV (x = 0.4). B3

Optical Deformation Potential

The phonons which usually dominate in the scattering probability are the long-wavelength optical phonons. These produce a short-range potential in the crystal which shifts the electronic band states. In polar semiconductors, like GaAs, the phonons are also accompanied by a long-range macroscopic electric field which produces additional scattering. The shifts of the electronic band states per unit ionic displacement associated with the long-wavelength optical phonons are called optical DP, do. The deduction of do from either low-field transport (see [28,42]) or resonant Raman data [29,30] is quite involved. Several theoretical calculations of d0 were carried out using various calculation methods [2,9,28,31-36]. In TABLE 2, we list the results of the theoretical and experimental do for GaAs. C

L POINT

Cl

Shear Deformation Potentials

The valence bands at the L point are split in the absence of spin-orbit coupling by the <111> stress. The intraband splitting can now be written with Kane's DP notation D35 as AEVL = 4(|) 1/2 D 3 5 eij

(4)

TABLE 2. Optical DP d0 at the T-valence band of GaAs (in e V). (I2

Remark

Ref

29J

(EPM)

[28]

37X)

(LCAO)

[28]

25j9

(BOA)

[3T]

364

(EPM)

[2]

3^3

(LCAO)

[2]

209

(LCAO)

[2]

1^5

(LMTO)

[9]

17^

(LMTO)

[32]

2^0

(LMTO)

[33]

3£0

(AIPC)

[34]

37J

(AIPC)

[35]

29^3

(LMTO)

[36]

41

Exper

(See [281)

274

Exper

[42]

48

I

Exper (RRS)

|

[29]

EPM=empirical pseudopotential method, LCAO=linear combination of atomic orbitals, BOA=Born-Oppenheimer approximation, LMTO=linear muflFin-tin orbital method, AIPC=ab initio pseudopotential calculation, RRS=resonant Raman scattering.

where e^ is the off-diagonal component of the <111> strain [37]. Similarly, the <001> strain splits the valence bands equally for all L valleys by

AE,L =2(I)^ 3 C 1 1 -1(CJ0 + eg]

(5)

In TABLE 3, we summarize the DPs D35 and D33 for GaAs [3,18,20]. The experimental data were obtained from the electroreflectance spectra with static uniaxial compression along the [001], [110] and [111] directions. C2

Optical Deformation Potentials

The optical DP d0 in Section B3 is determined by the splitting of the F-valence-band states produced by the long-wavelength optical phonons (F15 phonons). Similarly, the optical DP dlov (dloso) can be determined by the shift of the L-valence-band states produced by the F15-phonon displacement which is decomposed into the L1 and L3 representation of the k-group C3v at the L point [28]. The L1 part of the phonons causes a shift of the L-valence bands (L45-valence band d ^ and Lg-valence band Cd1080)). In the approximation that the strain dependence of the spin-orbit interaction is neglected, dlov is equal to dloso [28].

TABLE 3. Shear DPs D35 and D33 at the L point (A direction) of GaAs (in eV). DP

Theory (Ref)

Experiment (Ref)

D3^

4.5(L) [3]

8.5 [18]

3.7 (A) [3]

0±0.5 [20] -6.4±1.5f411

D,

3

2.4 [18] 3.45±0.3 [20] 3.2±0.3 [20]

I

I

-5±1.5[41]

If the spin-orbit interaction is neglected altogether, the L3 part of the phonon deformation splits the doubly degenerate valence-band edge. (This splitting is caused by the optical DP d3o. In the neighbourhood of the F point (i.e., k-0), the following two relations should be satisfied [38]: d3o = / 2 d o ,

dlov = -do

(6)

We list in TABLE 4 the theoretical [28,31,34,36] and experimental [39] L-valence-band optical DPs dlov, dloso and d3o for GaAs. We understand from this table that the difference |dlov-dloso| is very small. The experimental d3o was determined from resonant Raman scattering by Cardona [39]. He also reported the DP ratio d3o/(dloc-dlov) = -2, where dloc is the optical DP in the Lconduction band. TABLE 4. Optical D P s d lo v , dlo80 and d3o for the upper valence and spin-orbit-splitoff valence bands at the L point of GaAs (in eV). dlov

dj£

d^

Remark

Ref

-11.1

AhS

403

Theory (EPM)

[28]

-37.2

-rri

50.13

Theory (AIPC)

[34]

4L6

Theory (LMTO)

[36]

33

Theory (BOA)

[31]

45

Experiment (RRS)

[39]

-7.7

35.5

Theory (EPM)

[40]

-15.5

40.5

Theory (EPM)

[43]

-10.2

EPM=empirical pseudopotential method, AEPC=ab initio pseudopotential calculation, LMTO=linear muffin-tin orbital method, BOA=Born-Oppenheimer approximation, RRS=resonant Raman scattering.

D

SHEAR DEFORMATION POTENTIALS AT THE X POINT

The intravalley splitting of the X5 valence band is determined by the DP d3 according to the expression [37].

AE x = ^d3Ti3

where T13 = (2E22 - e^ - « g /

ft

(7)

The value OfD3 = 6.7 eV has been obtained with the EPM method for GaAs [40]. E

CONCLUSION

We have discussed the intravalley DPs for holes in the F- and L-valence bands of GaAs. The existing experimental data from various sources do not show good agreement. This is because those measurements are indirect and require a significant amount of analysis, interpretation and assumption. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] II1] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30]

G.L. Bir, G.E. Pikus [ Symmetry and Strain-Induced Effects in Semiconductors (Wiley, New York, 1974)] A. Blacha, H. Presting, M. Cardona [ Phys. Status Solidi B (Germany) vol. 126 (1984) p. 11-36] M. Cardona, S. Zollner [ in this Volume, Datareview 4.5} G.E. Pikus, G.L. Bir [Sov. Phys.-SolidState (USA) vol. 1 (1959) p. 136-8 ] D.D. Nolte, W. Walukiewicz, E.E. Haller [ Phys. Rev. Lett. (USA) vol.59 (1987) p.501-4 ] CG. Van de Walle, R.M. Martin [ Phys. Rev. Lett. (USA) vol.62 (1989) p.2028-31 ] P. Lawaetz [ Phys. Rev. (USA) vol. 174 (1968) p. 867, also see J.D. Wiley [ Solid State Commun. (USA) vol.8 (1970) p. 1865-8 ] J.C. Phillips [ Phys. Rev. Lett. (USA) vol.20 (1968) p.550-3 ] N.E. Christensen [ Phys. Rev. B (USA) vol.30 (1984) p.5753-65 ] E.P. O'Reilly [ Semicond. Sd. Technol (UK) vol. 1 (1986) p. 128-32 ] M. Cardona, N.E. Christensen [Phys. Rev. B (USA) vol.35 (1987) p.6182-94; erratum, ibid, vol.36 (1987)p.2906] C. Priester, G. Allan, M. Lannoo [ Phys. Rev. B (USA) vol.37 (1988) p.8519-22 ] CG. Vande Walle [Phys. Rev. B (USA) vol.39 (1989) p. 1871-83 ] S. Adachi [ GaAs and Related Materials: Bulk Semiconducting and Superlattice Properties (World Scientific, Singapore, 1994) ] M. Silver, W. Batty, A. Ghiti, E.P. O'Reilly [ Phys. Rev. B (USA) vol.46 (1992) p.6781-8 ] R.N. Bhargava, M.I. Nathan [ Phys. Rev. (USA) vol. 161 (1967) p.695-8 ] I. Balslev [ Solid State Commun. (USA) vol.5 (1967) p.315-7 ] RH. Pollak, M. Cardona [Phys. Rev. (USA) vol.172 (1968)p.816-37 ] A. Gavini, M. Cardona [ Phys. Rev. B (USA) vol. 1 (1970) p.672-82 ] M. Chandrasekhar, F.H. Pollak [ Phys. Rev. B (USA) vol. 15 (1977) p.2127-44 ] W. Walukiewicz [ J. Appl. Phys. (USA) vol.59 (1986) p.3577-9 ] H. Qiang, F.H. Pollak, G. Hickman [ Solid State Commun. (USA) vol.76 (1990) p.1087-91 ] J.D. Wiley, M. DiDomenico Jr. [ Phys. Rev. B (USA) vol.2 (1970) p.427-33 ] J.D. Wiley [ in Semicond. Semimet. vol. 10, Eds RK. Willardson, AC. Beer (Academic Press, New York, 1975) p.91-174] S. Adachi [ J. Appl. Phys. (USA) vol.58 (1985) p.Rl-R29 ] W. Potz, P. Kocevar [ Phys. Rev. B (USA) vol.28 (1983) p.7040-7 ] W.W. Ruhle, K. Leo, E. Bauser [ Phys. Rev. B (USA) vol.40 (1989) p. 1756-61 ] W. Potz, P. Vogl [ Phys. Rev. B (USA) vol.24 (1981) p.2025-37 ] M.H. Grimsditch, D. Olego, M. Cardona [ Phys. Rev. B (USA) vol.20 (1979) p. 1758-61 ] V.I. Gavrilenko, C. Trallero-Giner, M. Cardona, E. Bauser [ Solid State Commun. (USA) vol.67 (1988) p.459-63]

[31 ] [32] [33] [34] [35] [36] [37] [38] [39] [40] [41] [42] [43]

AJ. Hernandez-Cabrera, J. Sanchez-Dehesa, C. Tejedor [ J. Phys. C Solid State Phys. (UK) vol. 16 (1983) p.2251-9] N.E. Christensen, S. Satpathy, Z. Pawlowska [ Phys. Rev. B (USA) vol.36 (1987) p. 1032-50 ] L. Brey, N.E. Christensen, M. Cardona [ Phys. Rev. B (USA) vol.36 (1987) p.2638-44 ] B.-S. Wang, Z-Q. Gu, J.-Q. Wang, M.-F. Li [ Phys. Rev. B (USA) vol.39 (1989) p. 12789-95 ]; see also comment by M. Cardona, N.E. Christensen [Phys. Rev. B (USA), vol.41 (1990) p.5407-8 ] Z-Q. Gu, J.-F. Li, J.-Q. Wang, B.-S. Wang [Phys. Rev. B (USA) vol.41 (1990)p.8333-9 ] R.-Z. Wang, S.-H. Ke, M.-C. Huang [ J. Phys., Condens. Matter (UK) vol.4 (1992) p.6735-42 ] E.O. Kane [ Phys. Rev. (USA) vol. 178 (1969) p. 1368-98 ] M.A. Renucci, J.B. Renucci, R. Zeyher, M. Cardona [Phys. Rev. B (USA) vol. 10 (1974) p.4309-23] M. Cardona [ in Light Scattering in Solids II, Eds M. Cardona, G. Guntherodt (Springer-Verlag, Berlin, 1982) p. 19-178] H. Presting [ PhD Thesis, Univ. Stuttgart (Germany) (1985) unpublished] P. Etehegoin, J. Kircher, M. Cardona, C. Grein, E. Bustarret [Phys. Rev. B (USA) vol.46 (1992) p. 15139-49] R. Scholz [ J. Appl. Phys. (USA) vol.77 (1995) p.3232-42 ] R. Trommer [ PhD Thesis, Univ. Stuttgart (Germany) (1977) unpublished]

4.7

Electron effective mass and hole effective mass in GaAs, pressure dependence V.A. Wilkinson and A.R. Adams July 1995

A INTRODUCTION Effective masses are influenced by the separation of the bands, which will be modified by the application of hydrostatic pressure. We review below the available data for the pressure dependence of the effective electron and hole masses. B

ELECTRONEFFECTIVEMASSES

The zone-centre electron effective mass in GaAs is almost directly proportional to the direct energy-gap, E0, which increases with pressure as described in [I]. Shantharama et al [2] studied high purity GaAs grown by LPE on semi-insulating substrates. They determined the effective mass from magnetophonon oscillations in magnetic fields up to 9 T while applying hydrostatic pressure up to 20 kbar. By also measuring the photoconductive response of the crystal they were able to determine E0 and m*e as pressure was applied. Thus they were able to compare their results directly with k.p theory. After correcting for non-parabolicity and polaron enhancement effects, best fit to the effective mass was obtained by dm*e / dE0 = (0.045 ± 0.002) mo eV 1

(1)

Taking an average value of the pressure dependence of E0 of 10.73 x 10"2 eV/ GPa [1] and the atmospheric value of m*e = 0.0667 m0

(2)

leads to the pressure dependence of m*e given by d[ln m*e] /dP = 7.2 x 10"2 GPa 1

C

(3)

HOLEEFFECTIVEMASSES

Since the heavy-hole band is highly anisotropic and non-parabolic any change in mass with pressure will depend on the direction considered, the method of measurement and the effective energy of the holes. No direct measurement has yet been made of the pressure dependence of the effective mass of holes in GaAs. However, Alekseeva et al [3] have calculated dm*h/dP, for the [111] direction, from the valence band parameters deduced from the pseudopotential calculations of Neumann et al [4] and give

d[ln m*h] /dP = -5 x 10"3 GPa"1

(4)

More recently Adams et al [5] have studied the pressure dependence of the hole mobility from which they deduce an effective averaged change in the hole mass given by d[ln m*h] /dP = -1.0 x 10"3 GPa 1

(5)

From k.p theory one would expect the light-hole effective mass to vary with pressure at a similar rate to the electron effective mass. Readers are referred to the related Datareview 3.3 in this book, Hole mobility in GaAs, pressure dependence [6]. REFERENCES [1] [2] [3] [4] [5] [6]

V.A. Wilkinson, A.R. Adams [ Datareview in this book: 4.3 Direct bandgap of GaAs, pressure dependence ] L.G. Shantharama, A.R. Adams, CN. Ahmad, RJ. Nicholas [ J. Phys. C (UK) vol. 17 no.25 (1984) p.4429-42 ] Z.M. Alekseeva, A.P. Vyatkin, G.F. Karavaev, N.P. Krivortov [ Phys. Status Solidi B (Germany) vol.88 (1978) p.321] H. Neumann, I. Topol, K.R. Schulze, E. Hess [ Phys. Status Solidi B (Germany) vol.56 (1973) p.K55 ] A.R Adams,L.G. Shantharama [PhysicaB&C(Netherlands) vol. 139nos. 1-3 (1986)p.419 ] A.R. Adams [ Datareview in this book: 3.3HoIe mobility in GaAs, pressure dependence ]

4.8

Effective bandgap narrowing in doped GaAs M.S. Lundstrom, E.S. Harmon and M.R. Melloch August 1995

A

INTRODUCTION

GaAs bipolar devices such as solar cells and heterojunction bipolar transistors often contain regions with heavy impurity doping. It has long been known that heavy impurity doping alters the band structure of a semiconductor thereby altering its optical properties (e.g. [I]). Work on Si devices has demonstrated that the corresponding electrical effects can be substantial [2]. In this Datareview, we examine the doping dependence of the equilibrium np product in GaAs. In a heavily doped semiconductor, the free carrier - free carrier interactions lead to a bandgap renormalization, and interactions of the free carriers with randomly located dopants can produce band tails. Several workers have reported band structure vs. doping calculations which have been used to evaluate the doping-dependent n0 p0 product [3,4], written as

n oPo = n£ = I V ^

(1)

where ^02 is the equilibrium np product in a lightly doped semiconductor, and AG is termed the 'apparent', 'effective,' or 'electrical' bandgap [5]. It is important to note that AG is not the true bandgap shrinkage because heavy impurity doping can perturb the density of states, and the nopo product can be influenced by Fermi-Dirac statistics. EQN (1) is a defining relation for AG, which includes actual bandgap shrinkage, band distortion, and Fermi-Dirac statistics. From the optical absorption edge, one deduces an 'optical' bandgap, which is different from the true bandgap because optical transitions must occur from filled states in the valence band to empty states in the conduction band, so the optical gap depends on the location of the Fermi level. AG is an important material parameter because one can usually obtain a reliable description of the electrical performance of bipolar devices from knowledge of AG alone, without knowing about the detailed shape of the energy bands [5,6]. The relation of AG to the actual and optical bandgaps is discussed in [7]. For device modelling, knowledge of AG as a function of impurity doping is required. B

MEASUREMENT TECHNIQUES

In principle, it would be preferable to extract the band shape of a heavily doped semiconductor by optical techniques, then use the band shape to evaluate n^2 for device modelling. In practice, experience has shown that the most reliable results are obtained from electrical measurements of devices themselves. Measurements of effective bandgap narrowing in GaAs have been reported by using IV [8] and CV [9] measurements of diodes and from the IV analysis of bipolar transistors [10-12]. The transistor data are the most reliable because the device cleanly separates current components. Heavy doping effects can be studied by analyzing transistors with various base doping densities. The collector current consists almost entirely of the current injected into the heavily doped base, which is strongly influenced by AG.

The DC collector current of a homojunction (or graded junction HBT) is given by j

= o

°

^^(eq NBWB

V

»

/ K r

- l)

(2)

where the effective intrinsic carrier concentration, n^, is the quantity of interest. In EQN (2), WB is the base width, NB the base doping density, and DB the diffusion coefficient for minority carriers in the base. The doping density and base width of the device may be extracted from secondary ion mass spectroscopy measurements, but to determine n^ from the measured collector current, the minority carrier mobility must be known. It is well-known that the minority carrier mobilities in Si can differ significantly from the majority carrier mobilities in comparably doped materials [13], and the differences are even larger for GaAs [12,14]. This topic is the subject of two related Datareviews which present data for the doping dependent minority electron and hole mobilities in GaAs [15,16]. The parameter a in EQN (2) accounts for the possibility of non-diffusive carrier transport in thin bases. Because the minority electron mobility actually increases in very heavily doped material [14,12], the base width can become comparable to the mean-free-path, and Fick's law can lose validity. The parameter a may be evaluated from a solution to the Boltzmann equation [17] or obtained with high accuracy from an analytical treatment [18] as

, . 1V

W

.

(3)

where kT UR

=

\

o—:

(4)

is the Richardson velocity. In brief, the measurement procedure was as follows. Homojunction transistors with various base doping densities were fabricated and the I c vs. | VBE | characterized to extract the quantity, a DB, as a function of base doping density. Using measured values of the minority carrier diffusion coefficients, the non-diffusive transport parameter, a , was evaluated, which permitted ^ 6 to be extracted. Both npn and pnp transistors were used in order to characterize effective bandgap narrowing in both n- and p-type GaAs [H]. C

RESULTS

FIGURE 1 plots the effective intrinsic carrier concentration vs. doping density for both n- and p-type GaAs as obtained from transistor-based measurements. The results show significant heavy doping effects. For p-type GaAs, n^2 increases by a factor of * 7 - 8, or AG « 50 meV. This value

p-GaAs n-GaAs Bennett & Lowney parametric fits

Doping Density (cm 3 ) FIGURE 1. The measured (points) and theoretical (solid lines) effective intrinsic carrier concentration ratios in nand p-type GaAs. Also shown (as dashed lines) are simple empirical fits based on EQN (5). (Reprinted with permission from [12].

is about one-half of the corresponding value in p+-Si [2]. This difference is expected because of the small conduction band effective mass for GaAs [3]. Above « 3 x 1019 cm"3, the apparent bandgap shrinkage begins to decrease because of degeneracy effects which push the Fermi level into the valence band and effectively widen the bandgap. Although the magnitude of the effective bandgap shrinkage is considerably less than in Si, injected currents in bipolar devices, which scale as n^2, will be increased by a factor of « 7 - 8 . FIGURE 1 also shows results for n+-GaAs. In n-type GaAs, the apparent bandgap shrinkage is much smaller because the competing effects of bandgap shrinkage and degeneracy are dominated by degeneracy effects which lower r^2 The result is that, in contrast to Si, there is a considerable difference between n- and p-type material. This difference can be used to advantage in devices, for example the GaAs pseudo-HBT [19]. Also shown in FIGURE 1 are theoretical results of many body calculations as reported by Bennett and Lowney [3]. There are undoubtedly sizeable error bars on the measured results (perhaps 20 - 30%), but the agreement between theory and experiment is good. For device modelling purposes, it is desirable to have an empirical fit to the results. If we express the apparent bandgap shrinkage as

A G = AN 1 / 3

+

kT In [F1/2(EF)] - E F

(5)

We find A = 2.55 x 1O"8 eV (p-GaAs)

(6a)

A = 3.23 x 10"8 eV (n-GaAs)

(6b)

In these expressions, F,/2 is the Fermi-Dirac integral of order 1A, and E F is the Fermi level referenced to the majority carrier band edge. These empirical fits are shown in FIGURE 1 as dashed lines. D

CONCLUSION

Effective bandgap shrinkage alters the np product in GaAs. For P+-GaAs, n^2 can be « 7 - 8 times higher than its value in lightly doped GaAs. Although smaller than the corresponding results for P+-Si, the resulting factor of « 7 - 8 enhancement in injection currents can have important consequences for devices. In contrast to Si, there is a significant difference in the apparent bandgap shrinkage in n- and p-type GaAs. Available data agrees well with theoretical predictions. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] [II] [12] [13] [14] [15] [16] [ 17] [18] [19]

H.C. Casey Jr., F. Stern [ J. Appl. Phys. (USA) vol.47 (1976) p.631-43 ] J. W. Slotboom, H.C. de Graaff [ Solid State Electron. (UK) vol. 19 (1976) p.857-62 ] H.S. Bennett, J.R. Lowney [ J. Appl. Phys. (USA) vol.62 (1987) p.521-7 ] R.A. Abram [ in Semicond. Semimet. vol.39, Eds. R.A. Ahrenkiel, M.S. Lundstrom, (Academic Press, New York, 1993) p.259-3173 ] CM. van Vliet [ IEEE Trans. Electron Dev. (USA) vol.40 (1993) p. 1140-47 ] M.S. Lundstrom, RJ. Schwartz, J.L. Gray [ Solid-State Electron. (UK) vol.24 (1981) p. 195-202 ] J.A. del Alamo, J. Wagner [ J. Appl. Phys. (USA) vol.63 (1988) p.425-9 ] M.E. Klausmeier-Brown, M.S. Lundstrom, M.R. Melloch, S.P. Tobin [ Appl. Phys. Lett. (USA) vol.52 (1988) p.2255-7 ] P. van Mieghem, R.P. Mertons, G. Borghs, R. van Overstraeten [ Phys. Rev. Lett. (USA) vol.41 p.5952-9 ] M.E. Klausmeier-Brown, M.R. Melloch, M.S. Lundstrom [Appl. Phys. Lett. (USA) vol.56 (1990) p. 160-2] M.S. Lundstrom [ in Semicond. Semimet. vol.39, Eds. R A Ahrenkiel, M.S. Lundstrom, (Academic Press, New York, 1993) p. 193-258 ] E.S. Harmon, M.R. Melloch, M.S. Lundstrom [ Appl. Phys. Lett. (USA) vol.64 (1994) p.502-4 ] H.S. Bennett [ Solid-State Electron. (UK) vol.26 (1983) p. 1157-66 ] D.M. Kim et al [Appl. Phys. Lett. (USA) vol.62 (1993) p.861-3 ] E. S. Harmon et al [ Datareview in this book: 2.13 Minority electron mobility in doped GaAs ] M.L. Lovejoy et al [ Datareview in this book: 3.8 Hole lifetimes in n-type GaAs ] A.A. Grinberg and S. Luryi [ Solid-State Electron. (UK) vol.35 (1992) p. 1299-309 ] S. Tanaka, M.S. Lundstrom [ Solid-State Electron. (UK) vol.37 (1994) p.401-10 ] P.E. Dodd, M.S. Lundstrom, M.R. Melloch [IEEEElectron Dev. Lett. (USA) vol.12. (1991) p.629-31 ]

4.9

Carrier ionisation coefficients of GaAs J.P.R. David and G.E. Stillman July 1996

A

INTRODUCTION

When a high electric field is applied across a semiconductor, electrons (and holes) can gain sufficient energy to promote an electron from the valence band into the conduction band. This ionisation process results in extra carriers which in turn can also be accelerated by the field resulting in an avalanche multiplication process. The minimum energy required for an ionisation process to occur is called the threshold energy. The reciprocal of the average distance travelled by a carrier in the direction of the electric field before an ionisation event is defined as the ionisation coefficient or rate. Electron and hole ionisation coefficients, usually referred to as a and P respectively, are given in units of cm"1. These ionisation coefficients are important since they determine the breakdown voltage in devices such as HBTs and FETs thereby limiting their performance. In other devices such as avalanche photodiodes they determine the avalanche multiplication characteristics. Experimental results of a, P are usually expressed as a, P = A exp(-(B/E)m) cm"1 where A5B are constants, E is the electric field in volts/cm and m is a power term typically between 1 and 2. A detailed review of the physical processes underlying impact ionisation in semiconductors such as the effect of band-structure and scattering is given by Stillman and Wolfe [1] and more recently by Capasso [2]. The most popular technique for determining a and p is from photomultiplication measurements. Pure electron and hole primary currents are created in Schottky diodes or epitaxially grown junctions by optical illumination and the increase in photocurrent as a function of reverse bias voltage is measured. With information about the electric field profile across the diode (usually obtained from capacitance-voltage measurements) and electron and hole initiated photomultiplication characteristics, a and P can be derived as described in [1,2] for abrupt one sided junctions, reach-through structures and for p-i-n structures. There is a wide spread in the early experimental data of a and p in GaAs. This is due to errors in the measurement technique or the use of inappropriate device structures. The main sources of error in determining accurate ionisation coefficients are as follows: i)

Lack of pure electron and hole multiplication characteristics. This seems to be a particular problem in the earlier Schottky barrier structures used.

ii)

Mixed injection arising due to Franz-Keldysh electroabsorption.

iii)

Incorrectly accounting for the variation in primary photocurrent.

iv)

Inaccurate knowledge of the electric field profile across the avalanching region.

v)

Not correcting for the 'dead space', the distance carriers need to gain the ionisation threshold energy. This is a particular problem in heavily doped or short avalanching structures.

B

CARRIER IONISATION COEFFICIENTS, ELECTRIC FIELD DEPENDENCE

Logan et al [3] were the first to report on the measurements of ionisation coefficients in GaAs using diffused p-n GaAs junctions. Photomultiplication measurements were undertaken using white light and monochromatic light. Assuming that a = P, they determined the ionisation coefficients over the electric field range 950 kV/cm to 1500 kV/cm as -1.3 x 104 cm"1 to 2.6 x 105 cm"1, respectively. Kressel and Kupsky [4] undertook photomultiplication measurements in thick p+- n - n+ layers formed by vapour deposition. They also assumed that a = P in these thicker avalanching structures and the electric field range covered was lower between 280 kV/cm and 360 kV/cm. Their results disagreed with those of Logan et al [3] and suggested ionisation coefficients that were almost an order of magnitude larger. Although several other groups attempted to measure ionisation coefficients in GaAs after this, the assumption that a = P was always made. Stillman et al [5] summarised this data and were the first to show that a and P were unequal from photomultiplication measurements on semi-transparent Pt Schottky barrier diodes on n-type material. Using different wavelength light to inject 'pure' electrons and holes, they found that P > a over the electric field range ~243kV/cm to 333kV/cm. Pearsall et al [6] used LPE grown p-n junctions on (100), (110) and (111) orientations (described more fully in Section C). The (100) junction showed P > a at low electric fields and a «P at higher fields. Ito et al [7] deduced a and p from photomultiplication measurements on a (100) GaAs p-n diffused diode with a novel crater mesa structure. Different wavelengths of light focused at the top and bottom of the mesa were used to provide the pure carrier injection. Over the electric field range 400 kV/cm to 600 kV/cm the results obtained could be expressed as a function of electric field E in V/cm as ex = 5.6 x 106 exp (-2.41 x 106/E) cm"1 P = 1.5 x 106exp(-1.57 x 106ZE)Cm"1 suggesting that P < a. Ando and Kanbe [8] were the first to suggest that a > P in GaAs. They deduced this from photomultiplication measurements and excess noise measurements on diffused p+- n mesa diodes. These measurements are described more fully in Section E. The ionisation coefficient ratio PZa, K6^ was also found to be nearly constant, independent of electric field strength over the electric field range 240 kVZcm to 560 kVZcm. Using this constant value of K^, the expressions for the ionisation coefficients were given as

a = 1.1 x io 7 exp(-2.2 x 1O6ZE) cm'1 P = 0.5 a The most comprehensive and accurate set of measurements of ionisation coefficients in bulk GaAs to date has been carried out by Bulman et al [9]. Forty-four devices from eight separate p+- n LPE grown structures covering a n-type doping concentration of between 1.1* 1017 cm"3 to 2.2 x io 15 cm'3 and thicknesses of 4.5 |nm - 8 |im were used in this study. Great care was taken to ensure pure carrier injection for the photomultiplication measurements and an etch-stop layer was used to allow the substrate to be removed accurately. These structures also eliminated errors due to mixed injection of electrons and holes from Franz-Keldysh electroabsorption. The normalised photomultiplication characteristics were corrected for small changes in the primary photocurrent and a simple form of dead space correction was utilised to obtain the final a and P coefficients. Their results are given as a function of electric field E by; a = 1.899 x io 5 exp(-(5.75 x 105/E)182) cm"1 P = 2.215 x io 5 exp(-(6.57 x 105/E)L75) cm"1

Ionisation Coefficients (cm" )

and are valid over the electric field range 222 kV/cm to 625 kV/cm. FIGURE 1 shows these equations graphically.

Inverse Electric Field (x10~6 cm V" 1 ) FIGURE 1. Electron (bold) and hole (normal) ionisation coefficients for GaAs from Bulman et al [9] (solid lines) and Millidge et al [11] (dotted lines).

Further evidence of the a/p ratio from these layers was obtained when excess noise measurements were undertaken. With pure electron injection, the a/p ratio obtained was 1.7 and with pure hole injection the p/a ratio was 0.6 as would be expected. These results corroborate the a,p values

derived from the photomultiplication measurements. In recent years attempts have been made to extend the electric field range over which these measurements have been made. Conventional techniques of measuring small values of multiplication mean measuring small changes to photocurrents which is difficult. However, by monitoring the base hole current and utilising the gain of an n-p-n AlGaAs/GaAs heterojunction bipolar transistor operating in the common base configuration, Canali et al [10] have shown that very low multiplication values can be accurately measured. This technique makes the assumption that a > p and at low multiplication values secondary ionisation events are insignificant. Their results agreed well with the data of Bulman et al over the electric field range of 250 kV/cm to 220 kV/cm and extended the low electric field region to 176 kV/cm where a values as low as lcm"1 were obtained. Millidge et al [11] deduced the ionisation coefficients at high electric fields in a series of thin p-i-n structures with the thickness of the T regions varying from 0.4 jim down to 0.05 fim by measuring the breakdown voltage. They calculated the ionisation coefficients by solving the Boltzmann equation and this showed larger values of a and P than the extrapolation of the data of Bulman et al. Breakdown voltages calculated using these higher ionisation coefficients gave good agreement with experimental results. For electric fields between 500 kV/cm and 1600 kV/cm a more accurate expression for the ionisation coefficients than an extrapolation of Bulman et al was given as: a = 2.3 x io 6 exp(-(4.08 x 106/E)a634) cm"1 P = 2.66 x 107 exp(-(3.29 x 1O7ZE)044) cm"1 From FIGURE 1 we can see that below 600 kV/cm there is close agreement between these ionisation coefficients and those of Bulman et al [H]. However, at higher electric fields they increase at a faster rate and P is found to be larger than a. Plimmer et al [12] undertook similar measurements on a series of GaAs p-i-n structures with T region thicknesses from 1 |um to 250 nm. In spite of the large tunnelling dark currents in the very thin structures, photomultiplication characteristics were obtained with pure carrier injection, using lock-in measurement techniques. To corroborate these measurements, several complementary n-i-p structures were also measured. At low electric fields in the thicker structures, good agreement with the results of Bulman et al [9] was obtained as shown in FIGURE 2. However, at low electric fields for T region thicknesses of 0.1 \xm and less, a and P were found to be significantly lower. However, they increased at a faster rate with increasing electric field than that predicted by Bulman et al as shown in FIGURE 2. At the highest electric fields close to breakdown in these thin T region structures the ionisation coefficients were in close agreement with those deduced by Millidge et al [11]. The thinnest 250 nm thick structure allowed ionisation coefficients to be measured up to an electric field of 2000 kV/cm. At electric fields above 600 kV/cm, oc and P were seen to be very similar in magnitude. The reason for the T region thickness affecting the ionisation coefficients was attributed to the 'dead-space', the distance required by the carrier to achieve the threshold energy, which becomes increasingly important in thinner structures.

Ionisation coefficients (cm )

open symbols - a closed symbols - (3

6

1

Inverse electric field (x1Cf cm V" ) FIGURE 2. Ionisation coefficients for GaAs samples with different V region thicknesses, from Plimmer et al [12]. Solid lines are the data of Bulman et al [9]. a ~ p for structures thinner than 0.1 ^m.

C

CARRIER IONISATION DEPENDENCE

COEFFICIENTS

IN

GaAs,

ORIENTATION

The first group to consider the effects of the orientation dependence on impact ionisation coefficients of GaAs was Pearsall et al [6]. They calculated the ionisation threshold energy and its location in the Brilliouin zone for electrons and holes in <100>, <110> and <111> directions. Experimental measurements were carried out on LPE grown p+- n junctions. Chopped monochromatic light was used for the photoexcitation and the photocurrent was detected by conventional lock-in techniques. Their results showed that P, the hole ionisation coefficient, was similar for all three orientations. However there was some variation in a, the electron ionisation coefficient. For the <111> direction a was smaller than P over the entire electric field range investigated of 476 kV/cm to 313 kVZcm, whereas for the <110> direction P was smaller than a over the electric field range 556 kV/cm to 333 kV/cm. In the <100> direction intermediate behaviour was observed with a > P at higher fields but P > a at the lower electric fields. Lee and Sze [13] fitted the a and p data of Pearsall et al [1] to obtain expressions as a function of electric field (V/cm) as follows: a <100> = 9.12 x 104 exp (-(4.77 * 105ZE))348Cm"1 P <100> = 3.47 x 106 exp (-(2.18 x 106ZE))100Cm"1 a <110> = 2.19 x io 7 exp (-(2.95 x 106ZE))100Cm"1 P <110> = 3.47 x 106 exp (-(2.27 x 106ZE))100cm"1 a <111> = 7.76 x io 4 exp (-(4.45 x 105ZE))691 cm"1 P<111> = 6.31 x IO6 exp (-(2.31 x 106ZE))100Cm"1

Using these expressions, Lee and Sze [13] calculated the breakdown voltage (V B ) for p + - n junctions in the three orientations. They found that for background doping concentrations of approximately 1016 cm"3, V B is independent of orientation. For lower doping levels the <111> orientation had the largest VB while for higher doping levels the <100> orientation had the lowest VB. No new work has been done in this area since the original results of Pearsall et al [6]. However, more recent detailed work has been carried out on the orientation dependence of ionisation rates in InP [14] and in silicon [15]. In neither of these material systems was any anisotropy of the ionisation rates seen with orientation. A carrier trying to achieve the threshold for ionisation undergoes a considerable number of momentum randomising phonon scattering events. Therefore the direction of the electric field should not significantly alter the measured ionisation coefficients. The differences in a and p with orientation observed by Pearsall et al [6] are relatively small except at low electric fields and may be due to small uncertainties in the electric field and determination of the primary photocurrent. D

CARRIER IONISATION DEPENDENCE

COEFFICIENTS

IN

GaAs,

TEMPERATURE

The effects of temperature on the ionisation coefficients were investigated by Pearsall et al [6] in the <100> direction and by Capasso et al [16] in the <110> and <111> directions. Electron and hole multiplication measurements were made between 293 K and 523 K for the <100> devices and between 293 K and 473 K for the other two orientations. For the <100> orientation, the a/p ratio varied from <1 at temperatures below 473 K to >1 at temperatures above 473 K for the entire electric field range considered. The <111> orientation was found to have a similar temperature dependence to that of the <100> while the <110> orientation had an a/p ratio > 1. In all orientations, a and p decreased with temperature, a consequence of increased phonon scattering. The most comprehensive study of the temperature variation of ionisation coefficients in GaAs to date has been carried out by Robbins [17]. Three LPE grown p + - n - n+ wafers with active layer thicknesses of 4.5 \xm - 6.6 \im and n-type doping concentrations of 2.2 - 6.3 x 1015 cm"3 were used and the temperature range covered was 80 K - 465 K. The results for a and p as a function of electric field in V/cm can be summarised as follows a (cm1)

P (cm'1)

4.39 x 10 5 exp-(7.31 x IQ5TE)12

4.00 x IQ5 exp-(8.12 * IQ5TE)12

120

2.01 xlQ 5 exp-(5.11 XlQ5TE)16

5.89 x IQ5 exp-(9.24 x IQ5TE)12

160

3.27 x 10 5 exp-(6.73 x IQ5TE)14

3.95 x IQ5 exp-(7.59 x IQ5TE)14

T(K) JO

6

6

11

6

6

10

210

1.61 x 10 exp-(1.29x IQ TE)

3.71 x 10 exp-(1.89 x IQ TE)

250

2.88 x 10 6 exp-(1.75x IQ6TE)10

1.5Qx 10 6 cxp-(1.27 x IQ6TE)12

300

7.94 x 10 5 exp-(1.01 x IQ6TE)13

1.11 x 1 0 6 e x p - ( U 3 x IQ6TE)13

350

2.31 x 10 5 exp-(6.31 x IQ5TE)18

4.05 x 10 5 exp-(8.16 x IQ5TE)16

390

3.4Qx 10 5 exp-(7.31 x IQ5TE)17

2.54 x 10 5 exp-(6.71 x IQ5TE)20

430

I 3.31 x IQ5 exp-(7.32 x IQ5TE)17

| 7.20 x 10 6 exp-(2.50 x IQ6TE)10

_ _

T(K)

cc (cm"1)

p (cm'1)

465

5.79 x IQ5 exp-(8.55 x IQ5ZE)16

5.33 x 10 6 exp-(2.47 x IQ6ZE)10

These results show that the decrease in a and P with increasing temperature is much more pronounced for lower values of electric field and temperature. At the highest electric fields and temperatures there is little variation in the ionisation coefficients with temperature. These results were verified by studying the variation of the breakdown voltage with temperature. The breakdown voltage increased with temperature for the lower doped devices (i.e. lower electric fields) and at low temperatures but became almost constant for the highly doped devices (i.e. higher electric fields) and at high temperatures. For all temperatures the value of the a/p ratio was found to be similar to the room temperature value. E

CARRIER IONISATION COEFFICIENTS IN GaAs, DOPING DEPENDENCE

Law and Lee [18] were the first to study explicitly the effects of the doping density on the electron and hole ionisation coefficients. The effect of changing the doping is to change the electric field profile across the avalanching distance and as such is usually undertaken to allow a wide electric field range to be investigated. The doping range of 2.94 x 1015 cm"3 to 5.5 x 1016 cm"3 with a corresponding electric field range of 220 kV/cm to 470 kV/cm was covered by four Pt-Schottky GaAs diodes and one p+- n epilayer. Light from a mercury arc lamp and a 633 nm He-Ne laser was used to create the photocurrents which were detected by conventional lock-in techniques. The doping in the layers was adjusted slightly to ensure that p, the hole ionisation coefficient, was fitted by the same expression for all layers, as follows P = 4.10 x 105 exp(-(6.34 x 1O5ZE))20 cm"1 where E is the electric field in V/cm. a, the electron ionisation coefficient obtained for these dopings was then given by the following expressions: a (low doping)

= 1.97 x 105 exp(-(5.58 x 1O5ZE))20 cm"1

a (medium doping)

= 4.54 x 105 exp(-(6.28 x 105ZE))20 cm"1

a (high doping)

= 1.83 x 105 exp(-(5.79 x 105ZE))20 cm"1

where again E is the electric field in VZcm. Thus at low doping values aZP > 1 whereas at high doping aZP < 1. This behaviour is explained by the transit time of electrons in the high field region being similar to the band to band scattering time; electrons in the heavily doped structures transit the high field region before they can ionise. Ando and Kanbe [8] carried out noise measurements on (100) GaAs p+- n mesa diodes with a doping range of 6 x io 15 cm"3 to 9 x io 16 cm"3. The layers were grown by VPE and the p+-n junction was formed by zinc difliision. Initially photomultiplication measurements were undertaken by illuminating the p+ layer with a 633 nm laser and a 0.81 ^m LED. The photomultiplication curve obtained with 633 nm light was always larger than that obtained with

0.81 |Lim light at the same electric field, suggesting that a > p. Excess noise measurements were undertaken on these diodes which showed that the a/p ratio was nearly constant at 0.5 for all the dopings considered. Utilising the pure electron photomultiplication characteristics and the a/p ratio, the following expression for the ionisation rates could be derived, independent of the doping density: a = 1.1 x io 7 exp(-2.2 x 106ZE) cm"1 P = 0.5 a Bulman et al [9] undertook measurements on several specially fabricated (100) GaAs p-n junction diodes covering the doping range 2.2 x 1015 cm"3 to 1.1 x 1017 cm"3, corresponding to an electric field range of 220 kV/cm to 625 kV/cm. Although the uncorrected data appeared to show that the more heavily doped structures had slightly lower ionisation coefficients at a given electric field, when the 'dead-space' correction was made, a and P could be fitted smoothly across the entire electric field range with the equations given previously in Section B. Plimmer et al [12] also showed that 'dead-space' effects can alter the inferred ionisation coefficients from photomultiplication characteristics and the trends with doping seen in [18] may be in part due to dead-space corrections not being applied. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] II1] [12] [13] [14] [15] [16] [17] [18]

GE. Stillman, CM. Wolfe [ Semicond. Semimet (USA) Eds R.K. Willardson and A.C. Beer, vol. 12 (Academic Press, New York, 1977) ] F. Capasso [ Semicond. Semimet (USA) Eds R.K. Willardson and A.C. Beer, vol.22, Part D (Academic Press, New York, 1985) ] RA. Logan, A.G. Chynoweth, B.G. Cohen [ Phys. Rev. (USA) vol. 128 no.6 (1962) p.2518 ] H. Kressel, G. Kupsky [ Int. J. Electron. (UK) vol.20 no.6 (1966) p.535 ] G.E. Stillman, CM. Wolfe, JA. Rossi, A.G. Foyt [Appl. Phys. Lett. (USA) vol.24 no. 10 (1974) p.471 ] T.P. Pearsall, F. Capasso, RE. Nahory, MA. Pollack, J.R. Chelikowsky [ Solid-State Electron. (UK) vol.21 (1978) p.297-302 ] M. Ito, S. Kagawa, T. Kaneda, T. Yamaoka [ J. Appl. Phys. (USA) vol.49 no.8 (1976) p.4607 ] H. Ando, H. Kanbe [ Solid-State Electron. (UK) vol.24 (1981) p.629-34 ] G.E. Bulman, V.M. Robbins, GE. Stillman [ IEEE Trans. Electron. Devices (USA) vol.ED-32 no.ll(1985)p.2454-66] C Canali et al [ IEEE Electron Device Lett. (USA) vol. 15 no.9 (1994) p.354 ] S. Millidge, D. C Herbert, M. Kane, G.W. Smith, D.R. Wight [ Semicond. Sci. Technol. (UK) vol. 10(1995)p.344-47] SA. Plimmer, J.P.R David, T.W. Lee [ IEEE Trans. Electron. Devices (USA) (in press) ] M.H. Lee, S.M. Sze [ Solid-State Electron. (UK) vol.23 (1980) p. 1007-9 ] N. Tabatabaie, V.M. Robbins, N. Pan, GE. Stillman [Appl. Phys. Lett. (USA) vol.46 no.2 (1985) p. 182-4] V.M. Robbins, T. Wang, K.F. Brennan, K. Hess, G.E. Stillman [ J. Appl. Phys. (USA) vol.58 no. 12 (1985)p.4614-7] F. Capasso, R.E. Nahory, MA. Pollack [ Electron. Lett. (UK) vol. 14 no.4 (1979) p. 117 ] V.M. Robbins [ PhD Thesis, University of Illinois at Urbana-Champaign, January 1988 ] H.D. Law, CA. Lee [ Solid-State Electron. (UK) vol.21 (1978) p.331 ]

CHAPTER 5 OPTICAL FUNCTIONS OF GALLIUM ARSENIDE 5.1 5.2

Optical functions of GaAs Table of GaAs optical functions at 3 00 K

5.1

Optical functions of GaAs D.D. Nolte October 1995

A

INTRODUCTION

There are several pairs of analytic functions that each uniquely describe the optical properties of any material. These functions include the real and imaginary parts of the complex dielectric function, the reflectance and the reflected phase, and the real and imaginary parts of the complex refractive index. The complete definition of one pair defines all the other pairs. However, for optical investigations the complex refractive index is most convenient for use. The complex refractive index is defined by n((o) = n(co) + i k(co) = ^e(co)

(1)

where the complex dielectric function is expressed as

8(G)) = E1(G)) + i C2(G))

(2)

The real part of the complex refractive index n(G)) is called the real refractive index and the imaginary part k(G)) is called the extinction coefficient. Once these are defined, the components of the dielectric function can be obtained through the relations E1(G)) = n(G))2 - k(G))2 E2(G)) = 2n(G)).k(co)

^

One additional optical function that is convenient to use is the absorption coefficient, which determines the decay of irradiance by 1(Z) = IOe""2

(4)

where the absorption coefficient is given by .

2k((o)co

«(0)) =

V I

/

(5)

B

ANALYTIC PROPERTIES

Bl

Kramers-Kronig Relationships

By defining one component of a complex pair of analytic functions for all frequencies the other component is uniquely defined through the Kramers-Kronig relationship. Examples of these relationships are given by ei(w) - , . i pf- " ^ do,' Jo 2 * to* ~ (O , . 2(0 D r ~ E1(O)O . ,

e,(to) = - — P/

dor

n(o>)-l = £ p f 71

^

Jo

7

^c

Jo

InR(CO)^lPf0(o>) = JL pf2TT


a/ 2 / x 4(O _ r oc(co) = P/

o/2

_

<**& dlnR

JO

*• '

O2

dfa)/

( ^ In 1 ^7 + "! dco'

dco

a) - co

where P stands for the principal part, and the last relationship is for reflectance R(co) and the reflected phase 6(o>). The Kramers-Kronig relationships are useful when only one function of a pair is known or measurable. For instance, the optical functions of semiconductors are usually obtained by measuring the reflectance R(o>), calculating the phase 0(co) through EQN (8), and from these deducing the refractive index and extinction coefficient. The optical functions obtained in this way can be checked for self consistency by substituting them into EQN (7). B2

Sum Rules

In addition to the Kramers-Kronig relationships, additional checks on the self-consistency of a measured or derived pair of optical functions can be made using sum rules. The sum rule for the imaginary part of the dielectric function is ^

COe2(O)) dco = ^ - COp

(9)

where cop2 = ne2/me0 is the squared plasma frequency for an electron density n. A second sum rule

provides a check on the real part of the dielectric function, given by f~ G)3Je1(G)) - 1) C2(G)) dG) = -^- G)p4

(10)

Higher-order sum rules can also be defined [1-3]. Using the Kramers-Kronig relationships with the sum rules provides stringent tests of the fidelity of measured or derived optical functions [4]. C

GaAs OPTICAL FUNCTIONS

Cl

Static and Far Infra-Red Dielectric Functions

The static dielectric constant e(0) has a temperature dependence given by [5] 6(0)

=12.79 (1 + 1.0 x 10"4T) (H) = 13.17 at T = 295 K

which is valid for the temperature range 100 K < T < 600 K. The infra-red dielectric constant e(°°) is related to the static dielectric constant through the Lydane-Sachs-Teller relation

02)

f\ - M

where Cx^0 and 0)TO are the LO and TO optical phonon angular frequencies. The temperature dependence of the infra-red dielectric constant was extrapolated from the Lydane-Sachs-Teller relation to give [6] e(«>)

= 1 0 . 6 0 ( 1 + 9 x 10"3T) (13) = 10.88 at T = 300 K

Direct experimental data obtained near room temperature [7] yield a temperature dependence given by e(oo)

= 10.88 + 1 x io-3 (T - 293)

(14)

for the temperature range 15°C
1 +

"LO " ^ 7 0 (O70 - a)2 + iwF

(15)

Experimental dielectric values are derived from reflectance data and fit to this equation. The parameters describing GaAs at room temperature with an electron density of 1.8 x 1016 cm"3 are O)x = 33.23 meV, G>L = 36.21 meV and T = 0.23 meV [9], which exhibits smaller damping than earlier data with F = 0.30 meV [10, 11]. C2

Near-Infra-Red Optical Functions (80 meV - 1.3 eV)

An assesment of the mid-infrared regime from 0.1 eV to 1.3 eV by Blakemore [12], based on previously published data, leads to the suggested form for the refractive index given by ^ 7R

n(hv) -

7.10 + ^

1 07

— — 1 - 0 . 1 8 0 (hv)2 (30.08 hv)2 - 1

(16)

for room temperature, where hv is in units of eV. At low temperature, for the same photon energy range, the expression is given by n(hv) = \|

7.13 + — 1 - 0.175 (hv)2 (29.6 hv)2 - 1

(17)

for 0 < T < 30 K. These expressions agree with the asymptotic forms given for the longwavelength and short-wavelength limits [6] . The temperature dependence of the refractive index in the mid-infrared range from 60 meV to 0.25 eV has been measured and analyzed [13,14] with a temperature dependence at room temperature given by

P = £ £ = 4.5*10" K"

(18)

The more complicated temperature dependence for the energy range from 0.5 eV to 1 eV has been measured by McCaulley [15] with a temperature coefficient at room temperature given by P(hv) = (5.7 - 0.5 hv + 2.2 (hv)2) x 10"5 K"1

(19)

with hv in units of eV. C3

Band-Edge Optical Functions (1.3 eV - 1.5 eV)

The optical absorption [16,17] and refractive index [18-20] near the fundamental absorption edge of GaAs are important for photonic applications. The optical functions in this energy range are extremely sensitive to temperature and concentration because of the excitonic properties of the fundamental absorption edge. Therefore, it is difficult to assign intrinsic values to the optical functions in this energy range, although values have been obtained on high-purity samples [17, 19]. The strongest perturbation is the temperature dependence of the bandgap, which is given by the Varshni equation [21]

E 0 (T) = 1.519 0

5A

x

1Q4

T 2 eV T + 204

(20) K '

C4 Visible and Ultraviolet Optical Functions (1.5 eV - 1 0 eV) The optical functions of GaAs in the range of visible and ultraviolet frequencies have been measured by Aspnes and Studna for the range 1.5 eV to 6 eV [22]. They maintained careful sample surface quality by directly assessing the surface conditions during ellipsometric measurements of surface reflectance. The range from 6 eV to 22 eV was first measured by Philipp and Ehrenreich [23]. However, these early measurements do not join continuously with the more recent data of Aspnes, nor do they match with new data measured above 10 eV [24,25]. A remeasurement of the optical functions in the energy range from 6 eV to 10 eV would therefore be a significant new contribution. The temperature dependence of the optical functions from 1.24 eV to 5 eV has been obtained for temperatures from 220 C to 650 0 C [26,27]. C5

X-Ray Optical Functions (10 eV - 1000 eV)

The optical functions of GaAs in the X-ray region from 22 eV to 150 eV were originally measured by Cardona et al [28]. More recent data has been obtained by Henke et al [24] from 10 eV to 10,000 eV, and by Windt [25] from 10 eV to 1000 eV. D

CONCLUSION

The available data on the optical functions of GaAs at room temperature have been reviewed. Significant new contributions to the quantitative measurement of these functions have been made in the past decade. There have especially been significant contributions to the dispersion properties in the infrared transparency region because of the importance of this spectral range for optoelectronic applications. Significant contributions have also been made in the soft X-ray spectral regime, based on the increasing availability and importance of synchrotron light sources. However, improvements on previous data are still required. The most striking gap in improved measurements occurs from 6 eV to 10 eV. This is an important spectral range because the real part of the dielectric function becomes negative in this range. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] II1]

A. Villani, A.H. Zimerman [ Phys. Lett. A (USA) vol.44 (1973) p.295 ] R. Grundler [ Phys. Status Solidi B (Germany) vol. 115 (1983) p.K147 ] G. Leveque [ Phys. Rev. B (USA) vol.34 (1986) p.5070 ] H.W. Ellis, J.R. Stevenson [ J. Appl Phys. (USA) vol.46 (1975) p.3066 ] T. Lu, GH. Glover, K.S. Champlin [Appl Phys. Lett. (USA) vol.13 (1968) p.404 ] J.S. Blakemore [ J. Appl Phys. (USA) vol.53 (1982) p.R123 ] M. Bertolotti et al [ J. Opt. Soc. Am. B (USA) vol.7 (1990) p.918 ] E. D. Palik [ in Handbook of Optical Constants, vol.429 Ed. E.D. Palik, (Academic, New York, 1985)] H.R. Chandrasekar, A.K. Ramdas [ Phys. Rev. B (USA) vol.21 (1980) p. 1511 ] R.T. Holm, J.W. Gibson, E.D. Palik [ J. Appl Phys. (USA) vol.48 (1977) p.212 ] A.H. Kachare, W.G. Spitzer, F.K. Euler, A. Kahan [ J. Appl. Phys. (USA) vol.45 (1974) p.2938 ]

[12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28]

J.S. Blakemore [ J Appl. Phys. (USA) vol.62 (1987) p.4528 ] D.T.F. Marple [ J. Appl Phys.(USA) vol.35 (1964) p.1241 ] P.Y. Yu, M. Cardona [ Phys. Rev. B (USA) vol.2 (1970) p.3193 ] J.A. McCaulley, V.M. Donnelly, M. Vemon, I. Taha [ Phys. Rev. B (USA) vol.49 (1994) p.7408 ] M.D. Sturge [ Phys. Rev. (USA) vol. 127 (1962) p.768 ] H.C. Casey, D.D. Sell, K.W. Wecht [J Appl. Phys. (USA) vol.46 (1975) p.250 ] H.C. Casey, D.D. Sell, M.B. Panish [ Appl. Phys. Lett. (USA) vol.24 (1974) p.63 ] D.D. Sell, H.C. Casey, K.W. Wecht [ J Appl. Phys. (USA) vol.45 (1974) p.2650 ] CD. Thurmond [J Electrochem. Soc. (USA) vol.122 (1975) p.1133 ] Y.P. Varshni [ Physica (Netherlands) vol.34 (1967) p.149 ] D.E. Aspnes, A.A. Studna [ Phys. Rev. B (USA) vol.27 (1983) p.985 ] H.R. Philipp, H. Ehrenreich [ Phys. Rev. (USA) vol. 129 (1963) p. 129 ] B.L. Henke, J.C. Davis, E.M. Gullikson, R.C.C. Perera [ Lawrence Berkeley Laboratory (1988) LBL-26259 ] D.L. Windt [ Appl. Opt. (USA) vol.30 (1991) p. 15 ] H. Yao, P.G. Snyder, J.A. Woollam [J Appl. Phys. (USA) vol.70 (1991) p.3261 ] G.N. Maracas, CH. Kuo, S. Anand, R. Droopad [ J. Appl. Phys. (USA) vol.77 (1995) p. 1701 ] M. Cardona, W. Gudat, B. Sonntag, P. Y. Yu [ in Proc. 10th Int. Conf. Phys. Semicond. (USA) vol.209 (US AEC, 1970) ]

5.2

Table ofGaAs optical functions at 300K D.D. Nolte October 1995

k

a (cm1)

Ref

Energy (eV)

n

0.005

3.601

[1]

0.010

3.622

[1]

0.0150

3.666

0.001370

2.080

[1]

0.0200

3.752

0.002780

5.643

[1]

0.0250

3.947

0.007140

18.10

[1]

0.0300

4.680

0.03975

120.9

[1]

0.0310

5.156

0.07569

237.9

[1]

0.0320

6.235

0.2047

664.1

[1]

0.0325

7.588

0.4703

1550

[1]

0.0330

12.33

2.391

8000

[1]

0.0332

17.32

8.503

2.865e+04

[1]

0.0334

6.996

6.599

2.235e+O4

[1]

0.0335

4.253

5.311

1.804e-K)4

[1]

0.0340

1.120

2.868

9885

[1]

0.0350

0.4409

1.403

4979

[1]

0.0360

0.4877

0.5187

1893

[1]

0.0362

0.6679

0.3273

1202

[1]

0.0370

1.554

0.08782

329.4

[1]

0.0380

2.065

0.04124

158.9

[1]

0.0400

2.518

0.1673

67.83

[1]

0.0450

2.904

0.004731

21.58

[1]

0.0500

3.044

0.002183

11.06

[1]

0.0550

3.116

0.001239

6.907

[1]

0.0600

3.159

0.0007895

4.803

[1]

0.0650

3.187

0.0005422

3.573

[1]

Energy (eV)

n

k

a (cm"1)

Ref

0.0700

3.207

0.0003922

2.783

[1]

0.0800

3.233

0.0002284

1.853

[1]

0.0800

3.236

[2]

0.0900

3.252

[2]

0.100

2.262

[2]

0.150

3.285

[2]

0.200

3.294

[2]

0.250

3.300

[2]

0.300

3.304

[2]

0.400

3.313

[2]

0.500

3.324

[2]

0.600

3.337

[2]

0.700

3.353

[2]

0.800

3.372

[2]

0.900

3.395

[2]

1.00

3.422

[2]

1.10

3.454

[2]

1.20

3.493

[2]

1.30

3.540

[2]

1.30

3.535

[3,4]

1.31

3.541

[3,4]

1.32

3.546

[3,4]

1.33

3.553

[3,4]

1.34

3.558

[3,4]

1.35

3.566

[3,4]

1.36

3.573

10.00

[3,4]

1.37

3.581

20.00

[3,4]

1.38

3.590

30.00

[3,4]

1.39

3.600

0.000

90.00

[3,4]

1.40

3.611

0.0010

210.0

[3,4]

1.41

3.628

0.0030

460.0

[3,4]

Energy (eV)

n

k

a (cm"1)

Ref

1.42

3.646

0.0260

4000

[3,4]

1.43

3.640

0.0560

8100

[3,4]

1.44

3.635

0.0530

7700

[3,4]

1.45

3.634

0.0590

8700

[3,4]

1.46

3.634

0.0620

9100

[3,4]

1.47

3.635

0.0600

9800

[3,4]

1.48

3.636

0.0680

1.02e+O4

[3,4]

1.49

3.638

0.0710

1.070e-K)4

[3,4]

1.50

3.640

0.0740

1.120e+04

[3,4]

1.50

3.666

0.0800

1.221e+04

[5]

1.60

3.700

0.0910

1.483e-K)4

[5]

1.70

3.742

0.1120

1.928e+04

[5]

1.80

3.785

0.1510

2.749e-K)4

[5]

1.90

3.826

0.1790

3.445e+04

[5]

2.00

3.878

0.2110

4.279e+O4

[5]

2.10

3.904

0.2400

5.115e+04

[5]

2.20

4.013

0.2760

6.146e-K)4

[5]

2.30

4.100

0.3200

7.456e-K)4

[5]

2.40

4.205

0.3710

9.034e+04

[5]

2.50

4.333

0.4410

IAlIe-H)S

[5]

2.60

4.492

0.5390

1.420e+05

[5]

2.70

4.694

0.6960

1.905e+05

[5]

2.80

4.959

0.9910

2.813e+05

[5]

2.90

5.052

1.721

5.057e-K)5

[5]

3.00

4.509

1.948

5.925e-H)5

[5]

3.10

4.373

2.146

6.742e+05

[5]

3.20

3.938

2.288

7.422e-K)5

[5]

3.30

3.709

2.162

7.231e+05

[5]

3.40

3.596

2.076

7.153e+05

[5]

3.50

3.531

2.013

7.142e+05

[5]

3.60

3.495

1.965

7.171e+05

[5]

_J

Energy (eV)

n

k

a (cm"1)

Ref

3.70

3.485

1.931

7.241e+05

[5]

3.80

3.501

1.909

7.353e+05

[S]

3.90

3.538

1.904

7.526e-K)5

[5]

4.00

3.601

1.920

7.786e-K)5

[5]

4.10

3.692

1.969

8.182e+05

[5]

4.20

3.810

2.069

8.809e+05

[5]

4.30

3.939

2.260

9.849e-K)5

[5]

4.40

4.015

2.563

1.143e-K)6

[5]

4.50

3.913

2.919

1.331e-K)6

[5]

4.60

3.769

3.169

1.478e-K)6

[5]

4.70

3.598

3.452

1.664e-K)6

[5]

4.80

3.342

3.770

1.834e-K)6

[5]

4.90

2.890

4.047

2.010e+06

[5]

5.00

2.273

4.084

2.070e-K)6

[5]

5.10

1.802

3.795

1.982e4O6

[5]

5.20

1.599

3.484

1.836e-K)6

[5]

5.30

1.499

3.255

1.749e-K)6

[5]

5.40

1.430

3.079

1.685e4O6

[5]

5.50

1.383

2.936

1.637e-K)6

[5]

5.60

1.349

2.815

1.598e-K)6

[5]

5.70

1.325

2.710

1.566e-K)6

[5]

5.80

1.311

2.625

1.543e-K)6

[5]

5.90

1.288

2.557

1.529e-K)6

[5]

6.00

1.264

2.472

1.503e-K)6

[5]

6.00

1.395

2.048

1.245e-K)6

[6]

6.20

1.424

1.976

1.241e-K)6

[6]

6.60

1.247

2.047

1.369e+06

[6]

7.00

1.063

1.838

1.304e-K)6

[6]

8.00

0.899

1.435

1.163e-K)6

[6]

9.00

0.901

1.136

1.036e-K)6

[6]

9.52

0.582

0.785

7.571e-K)5

[7]

Energy (eV)

n

k

a (cm1)

Ref

9.97

0.656

0.893

9.022e-K)5

[7]

10.3

0.655

0.729

7.633e+O5

[7]

10.6

0.776

0.845

9.091e+05

[7]

10.8

0.691

0.737

8.035e+05

[7]

10.9

0.697

0.717

7.934e-H)5

[7]

11.4

0.700

0.649

7.508e-K)5

[7]

11.6

0.706

0.649

7.642e-K)5

[7]

11.8

0.716

0.646

7.741e+05

[7]

11.9

0.692

0.619

7.481e40.5

[7]

12.0

0.701

0.612

7.456e-K)5

[7]

12.1

0.712

0.656

8.032e+05

[7]

12.3

0.714

0.595

7.414e-K)5

[7]

12.5

0.704

0.585

7.431e-K)5

[7]

12.9

0.727

0.557

7.258e+05

[7]

13.3

0.709

0.529

7.129e-K)5

[7]

13.5

0.704

0.500

6.828e+05

[7]

13.5

0.703

0.491

6.727e-K)5

[7]

14.1

0.704

0.516

7.381e+05

[7]

14.9

0.692

0.391

5.885e-K)5

[7]

16.6

0.730

0.259

4.355e-K)5

[7]

16.7

0.726

0.215

3.631e-K)5

[7]

16.8

0.739

0.230

3.925e-K)5

[7]

17.3

0.723

0.242

4.230e-K)5

[7]

18.5

0.761

0.219

4.096e-K)5

[7]

20.1

0.819

0.172

3.505e-K)5

[7]

21.2

0.822

0.171

3.676e-K)5

[7]

22.3

0.820

0.204

4.616e-K)5

[7]

23.0

0.835

0.154

3.588e-K)5

[7]

23.1

0.837

0.155

3.625e-K)5

[7]

25.3

0.843

0.111

2.848e-K)5

[7]

25.6

0.843

0.011

2.848e-K)5

[7]

Energy (eV)

n

k

a (cm"1)

Ref

26.9

0.872

0.0899

2.451e-K)5

[7]

27.3

0.875

0.0839

2.320e+05

U]

27.8

0.880

0.0739

2.080e+05

U]

28.8

0.882

0.0820

2.394e-K)5

U]

30.6

0.9034

0.0610

1.888e+O5

U]

32.7

0.9130

0.0601

1.990e-K)5

U]

34.8

0.9258

0.0433

1.524e-K)5

U]

35.1

0.9260

0.0532

1.893e+05

U]

35.3

0.9222

0.0503

1.798e+05

U]

37.5

0.9397

0.0402

1.525e+O5

U]

37.9

0.9420

0.0436

1.675e+05

U]

38.9

0.9457

0.0493

1.941e+05

U]

39.5

0.9474

0.0410

1.640e+05

U]

40.8

0.9520

0.0532

2.200e-K)5

U]

43.8

0.9563

0.0803

3.561e+O5

U]

46.4

0.9581

0.0712

3.349e+O5

U]

48.4

0.9627

0.0749

3.670e-K)5

U]

51.0

0.9626

0.0497

2.569e+05

U]

52.2

0.9640

0.0563

2.979e+05

U]

54.4

0.9660

0.0550

3.030e-K)5

U]

56.9

0.9703

0.0567

3.267e-K)5

U]

72.3

0.9825

0.0460

3.371e+O5

U]

91.5

0.9805

0.0434

4.023e-K)5

U]

96.5

0.9890

0.0311

3.035e-K)5

U]

109

0.9801

0.0376

4.143e+05

U]

114

0.9831

0.0327

3.778e+O5

U]

133

0.9818

0.02460

3.308e-K)5

U]

151

0.9828

0.01720

2.631e+05

U]

172

0.9863

0.02330

4.054e-K)5

U]

183

0.9867

0.01320

2.450e-K)5

U]

193

0.9971

0.01140

2.224e-K)5

U]

Energy (eV)

n

k

a (cm"1)

Ref

237

0.9903

0.00630

1.512e+05

[7]

260

0.9914

0.00570

1.502e+05

[7]

277

0.9918

0.00472

1.326e+05

[7]

284

0.9922

0.00504

1.452e-K)5

[7]

312

0.9943

0.00372

1.175e-K)5

[7]

395

0.9954

0.00190

7.610e+04

[7]

447

0.9961

0.00137

6.196e-K)4

[7]

500

0.9970

0.000888

4.501e+04

[7]

511

0.9971

0.000791

4.097e+04

[7]

525

0.9973

0.000712

3.786e-K)4

[7]

556

0.9974

0.000763

4.299e+04

[7]

573

0.9978

0.000650

3.773e-K)4

[7]

615

0.9981

0.000543

3.385e-K)4

[7]

638

0.9981

0.000518

3.345e-K)4

[7]

678

0.9984

0.000401

2.753e+04

[7]

705

0.9986

0.000464

3.313e+O4

[7]

743

0.9987

O.OOO339

2.551e+04

[7]

776

0.9987

0.000230

1.809e-K)4

[7]

811

0.9990

0.000272

2.235e+04

[7]

852

0.9990

0.000192

1.656e+04

[7]

930

0.9992

0.000131

1.234e-K)4

[7]

REFERENCES [1] [2] [3] [4] [5] [6] [7]

H.R. Chandrasekar, A.K. Ramdas [Phys. Rev. B (USA) vol.21 (1980) p. 1511 ] J.S. Blakemore [ J. Appl. Phys. (USA) vol.62 (1987) p.4528 ] H.C. Casey, D.D. Sell, K.W. Wecht [J. Appl. Phys. (USA) vol.46 (1975) p.250 ] D.D. Sell, H.C. Casey, K.W. Wecht [ J. Appl. Phys. (USA) vol.45 (1974) p.2650 ] D.E. Aspnes, A.A. Studna [ Phys. Rev. B (USA) vol.27 (1983) p.985 ] H.R. Philipp, E. Ehrenreich, [ Phys. Rev. (USA) vol. 129 (1963) p. 129 ] D.L. Windt [Appl. Opt. (USA) vol.30 (1991) p.15 ]

CHAPTER 6 ELECTRO-OPTIC PROPERTIES OF GALLIUM ARSENIDE 6.1 6.2 6.3

Electro-optical coefficients of GaAs Franz-Keldysh effect in GaAs Piezoelectric properties of GaAs

6.1

Electro-optical coefficients of GaAs DJ. Robbins January 1996

A

INTRODUCTION

There has been considerable interest recently in the electro-optic properties of semiconducting materials, both in theoretical (parametrised) models [1-3] and in experimental measurement [4-7] as a result of the development of phase modulators utilising this effect and the development of electro-optic sampling as a device assessment tool. In zinc-blende structures the F43m symmetry allows only one independent non-vanishing tensor component for the linear electro-optic tensor [8]. This is the case for GaAs [1] and the linear electro-optic (Pockels) effect, which is characterised by a single coefficient X23x = r41. The quadratic electro-optic (Kerr) effect, which is discussed in some detail in [3], is usually negligible in bulk GaAs (~ 10"21 - 10"20 (m/V)2), and will not be discussed further here, but in lower symmetry quantum well and superlattice systems the electro-optic coefficients can be considerably enhanced over the bulk [9]. B

LINEAR ELECTRO-OPTIC COEFFICIENT

Experimental measurements of r41 are often performed at low frequencies compared to acoustic resonator frequencies so that it is the undamped or constant stress value rT41 which is measured; rT41 is identical to rs41 + r p41 [2] where 1^41 is the clamped or constant strain value and the difference between rT41 and 1^41 is P41, the piezoelectrically induced electro-optic coefficient. This distinction is important because it appears that the large discrepancies between many early measurements of the electro-optic coefficient may have been caused by residual strain in the samples. Additional problems are caused by multiple reflections, and the effects of photoconductivity [4]. Older work, much of which is reviewed in [5], yields a scatter of values between -1.0 x 10'12 m/V and -1.6 x 10"12 m/V for a wavelength range 1.0 |um to 10.6 ^m. Theoretical predictions, which are in satisfactory agreement with experiment, suggest that there should be little wavelength dependence within this range, but show that r41 begins to increase steeply towards the band edge [10]. Papers by Suzuki and Tada [5] and Sugie and Tada [4] contain the most comprehensive and reliable studies to date. These were carried out on high resistivity GaAs:Cr samples which had a sufficiently high absorption coefficient to render multiple reflections negligible. Their results, which are in good agreement with measurements due to Walsh [11], are shown in TABLE 1. A more recent measurement on single crystal GaAs [12] finds rT41 = 1.3 * 10"12 m/V ± 10%. Values of rs41 in TABLE 1 were obtained from rT41 and measurements of the elasto-optic coefficient p 4 4 .

TABLEl Ref

Wavelength (microns)

rT41 (10-12m/V)

Refractive index

1^ 4 1 (1012m/V)

_[6]

09

3.57

-1.6

-

_[4]

1.064

3.48

-1.17

-L33

J5]

U5

J4]

L21

-1.43 ±0.07 3.43

-

-1.25

-L41

J4]

L31

3.41

-1.28

-L46

J4]

L50

3.38

-1.36

-L53

_[5]

^39

-1.24 zb 0.04

-

J5]

1O6

-1.51 ±0.05

-

[12]

I 10.6

I

I -1.3 ±0.13

1

-1.9±0.19

Indirect support for a value approximately equal to -1.5 x 10"12m/V at 3.39 \im arises from the first-order Raman coefficient which is related to r41 and to the second harmonic generation coefficient d41 [13]. Generally, results from 1.3 \im phase modulators in GaAs are consistent with a value of r41 of approximately -1.5 * 10"12m/V but do not provide a precise measurement. In electro-optic sampling a relative measurement is usually made. C

TEMPERATURE DEPENDENCE

Measurements of the temperature dependence of r41 [6,7] indicate a linear variation in the range 100-300 K, with: rT41 = r41 (100 K)[I + (3.2 ± 0.5) x 10"3 x (T-IOO)]

(1)

where T is the temperature [6]. In [6] values of n3 r41 for the energy range 1.2 eV to 1.45 eV are given. At 1 ^m r41 is approximately equal to (1.16 ± 0.09) x 10"12 m/V in good agreement with [4]. A useful relation to older work is :

'«

r41

(esu)

"T

^

(SI)

<2>

REFERENCES [1] [2] [3] [4] [5] [6] [7] [8]

Chun-Ching Shih, A. Yariv [ J. Phys. C (UK) vol. 15 (1982) p. 825-46 ] S. Adachi, K. Oe [ J. Appl. Phys. (USA) vol. 56 (1984) p. 74-80 ] S. Adachi, K. Oe [ J. Appl. Phys. (USA) vol. 56 (1984) p. 1499-504 ] S. Sugie, K. Tada [ Jpn. J. Appl. Phys. (Japan) vol. 15 (1976) p. 421-31 ] N. Suzuki, K. Tada [ Jpn. J. Appl. Phys. I (Japan) vol. 23 (1984) p. 1011-6 ] Yu.V. Shaldin, D.A. Belogurov [ Sov. J. Quantum Electron. (USA) vol. 6 (1976) p. 897-901 ] T. G. Okroashvili [ Opt. Spektrosk. (USSR) vol.47 (1979) p. 442-3 ] F.A. Hopf, G.I. Stegeman [ Applied Classical Electrodynamics: Nonlinear Optics, vol.2 (Wiley, New York, 1986)]

[9] [10] [11] [12] [13]

A. Sa'ar, N. Kuze, J. Feng, I. Grave, A. Yariv [Appl. Phys. Lett. (USA) vol.61 (1992) p. 1263-5 ] A. Hernandez-Cabrera, C. Tejedor, F. Meseguer [J. Appl. Phys. (USA) vol. 58 (1985) p.4666-9 ] T. E. Walsh [ RCA Rev. (USA) vol.27 (1966) p.323 ] GL. Herrit, H.E. Reedy [Mater. Res. Soc.Symp. Proc. (USA) vol. 152 (1989) p. 169-74 ] S. C. Varshney, A.A. Gundjian [ J. Appl. Phys. (USA) vol.52 (1981) p.6301-5 ]

6.2

Franz-Keldysh effect in GaAs DJ. Robbins January 1996

A

INTRODUCTION

The Franz-Keldysh effect is an electric field induced change of the complex dielectric constant at optical energies close to the bandgap energy. It has two parts, electroabsorption and electrorefraction, i.e. the absorption change and the real index change associated with it, which are related by the Kramers-Kronig integral. The Franz-Keldysh effect has been studied widely in various guises, particularly in modulation spectroscopy which has been extensively used to establish the band structure of semiconductors. Electro-reflectance detects the electro-refractive effect (see for example a number of articles in [I]), while in photo-reflectance surface fields are modulated by optically generated carriers. A more up-to-date review of such modulation methods is given in [2]. However, only a few direct measurements of the absorption edge shift with applied field have been made. Since the applied field is generally obtained by making measurements within a depletion region one experimental difficulty is that of ensuring a uniform field. As a result the absorption coefficient, a, is often given as a function of voltage so that the results remain device-specific. There are exceptions, however, where the use of thick high quality, low doped (up to 1015 cm"3) epitaxial layers has ensured a highly uniform field and low built-in field at zero applied bias [3-5]. B

ELECTRO-ABSORPTION

In the absence of comprehensive measured data, a theoretical approach is valuable. Although theoretical expressions for absorption in the band edge region due to Tharmalingham and Callaway [6] have been widely used in the past, it is now clear as a result of more recent work [4,5] that an expression due to Rees [7], which rather than providing absolute values relates the absorption at two different fields, provides a better means of extrapolating from the relatively well known zero field absorption data to the case in an electric field. The absorption coefficient for a photon of energy, E, by GaAs under an electric field, F, is given by OO

a(E,F) = C I o(E - y, O) Ai (-Cy) dy

(1)

- O O

where Ai is the Airy function and the field, F, is given in V/m. The Airy function is defined as OO

Ai(x) = —

f exp (i(u3/3 + ux))du

(2)

-OO

and the field dependence is contained within C, where

(h2e2F2J

W

In this expression, \i* is the reduced effective mass. Tabulated values for the zero field absorption coefficient a (E, 0) [5] are given in TABLE 1, and the above convolution integral can readily be performed using this zero bias data to generate the absorption coefficient in an electric field. For GaAs: c

.

1.897x10» p 2/3

V

'

v*/

The above expression compares favourably with experimental data from references [3] and [5] over a field range 0 - 4 x 105 V/cm. In fact C depends only on (|u*)1/33 which changes only weakly across the III-V alloys so that the same expression is also useful for other materials [4]. This approach has also been used recently by Leeson and Payne [8], who have analysed the data of [3,4] in some detail. At energies above the absorption edge the change in absorption with field Aa (E, F) exhibits the well known Franz-Keldysh oscillations. The variation in period of the oscillation as a function of field is well characterised and is discussed by a number of authors: for GaAs see for example Silberstein and Pollack [9] and Bobylev et al [10], but more general data for <x(E, F) does not appear to be available. Electroabsorption also exhibits a polarisation dependence because of the valence band anisotropy induced by the field, which causes light polarised along the field to be absorbed more strongly than light polarised normal to the field. A study in AlGaAs/GaAs heterostructures is given by Reinhart [H]. However, the electro-absorption constant is related to the applied voltage rather than field. C

ELECTROREFRACTION

The electrorefractive effect displays a quadratic field dependence for small fields but saturates at high field. A number of theoretical analyses have been published based upon a Kramers-Kronig integration of the electro-absorption change [12], but there are few direct measurements available for GaAs. In [4] average values of An = 2 x 10"5 to 6 x 10"5 were found for fields in the range 2.6 x lo 4 to 5.2 x lO4 V/cm at photon energies 20 - 40 meV below the band edge, and these values are fairly consistent with theoretical estimates [4,12].

TABLE 1. Zero field GaAs absorption. a (cm'1)

(Bandgap energy - E) (meV)

4.9 x IQ4

-640

3.7 x IQ4

-480

4

2.8 x IQ

-345 4

2.12 x IQ

-275

1.89 x IQ4

-225

4

-175

1.66 x IQ 4

1.5 x IQ

-125 4

1.29 x IQ

4

-75

1.11 x IQ

-35

9.35 x IQ3

J^

9.1 x IQ3

-8 3

9.35 x IQ 3

9.3 x IQ

-5 -2

8.7 x IQ3

j)

6.0 x IQ3

2

4.1 x IQ3

Jk5

3

3.5 x IQ

j> 3

1.63 x IQ

10

2

7.7 x IQ

15

3.9 x IP2

^20

_86

JO

J2

_50

JL8

_75

_L6

100

2.0

I 125

REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] II1] [12]

D.F. Blossey, P. Handler, B.O. Seraphin, D.E Aspnes, N. Bottka [ in Semicond. Semimet. vol.9, Eds R.K. Willardson, A.C. Beer (Academic Press, New York, 1974) ] H. Shen, M. Dutta [J. Appl Phys. (USA) vol.78 (1995) p.2151-76 ] G.E. Stillman, CM. Wolfe, CO. Boder, J.A. Rossi [ Appl Phys. Lett. (USA) vol.28 (1976) p.544 ] T.E. Van Eck, LM. Walpita, W.S.C. Chang, H.H. Weider [ Appl Phys. Lett. (USA) vol.48 (1986) p.451-3] D.R. Wight et al [ IEEProc. J (UK) vol. 135 (1988) p.39-44 ] K. Tharmalingham [ Phys. Rev. (USA) vol.130 (1963) p.2204-6 ]; J. Callaway [ J. Phys. Chem. Solids (UK) vol.29 (1968) p. 143-53 ] H.D. Rees [ J Phys. Chem. Solids (UK) vol.29 (1968) p. 143-53 ] M.S. Leeson, F.P. Payne [IEEProc. J, Optoelectron. (UK) vol.141 (1994) p.257-64 ] RP. Silberstein, F.H. Pollack [ J. Vac. Sci. Technol (USA) vol. 17 no.5 (1980) p. 1052-6 ] B A Bobylev, A.F. Kravchenko, A.S. Terekhov [Sov.-Phys. Semicond. (USA) vol.7 (1974) p.1381 ] F.K. Reinhart [Appl Phys. Lett. (USA) vol.22 (1973) p.372-4 ] B.R. Bennett, R.A. Soref [IEEE J. Quantum Electron. (USA) vol.23 (1987) p.2159-66 ]

6.3

Piezoelectric properties of GaAs DJ. Robbins January 1996

A

INTRODUCTION

The piezoelectric effect arises from a charge polarisation induced by the application of a stress in crystals which lack a centre of inversion symmetry. The III-V compounds are the simplest crystals with this property. However, the cubic (F43m) symmetry allows only one independent, non-vanishing tensor component [I]. The magnitude of the piezoelectric effect is described by the piezoelectric constants: e14, d14, h14 and g14, which are related by the standard constitutive equations, and the terminology used varies somewhat according to the context within which the effect is being studied. For example, d is often referred to as the charge coefficient and g as the voltage coefficient. The piezoelectric stress constant, e14, relates the induced polarisation (or displacement) to the strain, or conversely stress to field. Likewise, the piezoelectric modulus, d14, relates the strain to the associated electric field or conversely the displacement to the stress. For a full discussion the reader is referred to [1] and other standard texts on piezoelectricity. Since only the (1,4) component survives, it requires shear stresses to generate a piezoelectric response in GaAs grown in the <100> direction widely favoured by epitaxy. The effect is largest along the <111> axes, and recently a number of groups have begun the study of piezoelectric effects in strained layer InGaAs/GaAs quantum wells grown on the <111> axis (see for example [2]). B

PIEZOELECTRIC PROPERTIES

The piezoelectric constants of GaAs were reliably measured many years ago by ArIt and Quadflieg [3]. As discussed in [3], the piezoelectric effect arises from a combination of effects: ionic polarisation, change in ionicity and electronic polarisation. The sign of e14 is negative if the crystal is expanded in the <111> direction, when the A-faces become negatively charged. Using a Hall effect method at room temperature on GaAs samples of fairly high conductivity, they found that e14= -0.16 C/m2 ± 10%, and hence d14 = -2.7* 10-12m/V (e14 = d14.c44 where C44 is the elastic stiffness coefficient which is ~ 6 x 1010 N/m2 in GaAs). The quoted accuracy includes estimates of the measurement error and variability between samples of different origin. The remaining piezoelectric constants can also be derived from e14, and are g14 = 2.4x 10-2m2/Cand h 14 =14.5 x 108VAn (g14 = d14/e and g14 = d14/e, where e is the dielectric function).

The sign of e14 is negative if the crystal is expanded in the <111> direction, when the A-faces become negatively charged. This sign convention is significant when comparing with other cubic crystals such as the II- VIs where e14 is positive for this sense of the charge polarisation. However, in other circumstances authors tend to report only the magnitude of e. Much more recently, Chin [4] has studied the Hall mobility in high purity MBE grown samples, of residual electron density ~ 3 x 1013 cm"3, in the temperature range 30 - 70 K where acoustic phonon piezoelectric scattering dominates. A best fit to the mobility data in this temperature range was obtained using e14 = 0.21C/m2. A closely related system, strained layer InxGa1^AsZGaAs quantum wells, has been studied by optical spectroscopy [5,6]. Analysis of the data, which is by no means straightforward, suggests a value for e14 in the ternary alloy for low In-contents in the range x = 0.07 to 0.23 which is about 70% that which would be obtained by simple linear interpolation between accepted values for GaAs and InAs. In [6] it is found that e14 =(0.115- 0.2324x) C/m2 in this range. REFERENCES [1] [2] [3] [4] [5] [6]

J. F. Nye [ Physical Properties of Crystals (Clarendon, London, 1959) ] D.L. Smith, C. Mailhiot [ J. Appl. Phys. (USA) vol.53 (1988) p.2717-9 ] G. ArIt, P. Quadflieg [ Phys. Status Solidi (Germany) vol.25 (1968) p.323-30 ] V.W.L. Chin [SolidState Electron. (UK) vol.37 (1994) p. 1345-7 ] R.A. Hogg et al [ Phys. Rev. B (USA) vol.48 (1993) p.8491-4 ] J.L. Sanchez-Rojas, A. Sacedon, F. Gonzalez-Sanz, E. Calleja, E. Munoz [Appl Phys. Lett. (USA) vol.65 (1994) p.2042-4]

CHAPTER 7 INFRARED ABSORPTION AND ENERGY LEVELS DUE TO IMPURITIES 7.1 7.2 7.3 7.4 7.5

IR absorption bands due to localised vibrational modes (LVM) of impurities in bulk and epitaxial GaAs IR absorption due to free carriers in GaAs Electronic absorption bands of impurities and defects in GaAs Energy levels due to transition metals in GaAs DX centres in GaAs

7.1

IR absorption bands due to localised vibrational modes (LVM) of impurities in bulk and epitaxial GaAs R. Murray July 1995

A

INTRODUCTION

'Light' impurities in GaAs are able to vibrate at frequencies greater than the maximum lattice frequency (-285 cm"1). Consequently, they are able to interact only weakly with their neighbours and the mode is localised. The motion of the impurity induces a dipole moment which can couple to light leading to absorption in the far infrared. Before the advent of Fourier transform infrared (FTIR) spectrometers LVM absorption measurements were limited to relatively thick (~ 1 mm) bulk samples. Epitaxial layers only a few mm thick and even d-layers are now investigated routinely using FTIR. LVM features are narrow since the lifetimes in the excited states are relatively long due to the weak coupling to neighbours and are superimposed on broad intrinsic two-phonon features. LVM frequencies of isolated impurities in bulk and epitaxial GaAs are contained in TABLE 1. B

COMPENSATION

If the impurity is isoelectronic then LVM measurements are readily obtained by subtracting the two-phonon background using an undoped control sample of similar thickness. However, if the impurity is electrically active there is strong free carrier absorption which can mask the LVM features. Samples have been compensated by (a) in-diffusion of lithium or copper although this has the disadvantage that Li- and Cu- complexes form which complicate the analysis (see TABLE 2) or (b) irradiation by fast (~2 MeV) electrons which create deep intrinsic electron and hole traps. The latter technique although more commonly used does result in some cases in the generation of irradiation related complexes (see TABLE 3). Many boron-related complexes, B(3)B(IO), have been reported in irradiated and annealed material [49] although the structure of these is unknown. Recently, there has been considerable interest in the behaviour of hydrogen in III-V and group IV semiconductors (see Datareview 10.8, 16.6 in this book). Hydrogen may be introduced by implantation or from a plasma and is found to compensate (passivate) shallow donors and acceptors resulting in new H-related LVM features in the mid-infrared region (see TABLE 4). In carbon-doped material grown by MOMBE the presence of these features in as-grown material results in partial compensation which may be removed by annealing. Hydrogen also modifies the LVM frequency of the passivated impurity (see TABLE 5). C

CALIBRATION

LVM features may be calibrated providing an alternative method of determining the impurity concentration [X], when other quantitative techniques such as SIMS or Hall measurements are available. Then the integrated absorption is given by:

/

a(E)dE =

'L J C Mimpn 2

JLVM

(i) '

v

where n is the refractive index, M^1, the mass of the impurity and r|, the 'apparent charge', is the constant of proportionality. r| takes values between 0.5 and 3. LVM absorption is therefore weak. Calibration factors for Be, B, C, Si and Al are listed in TABLE 6. Although more data exist for C there is some dispute regarding the ratio of the calibration factors at 300 K and 77 K. The reader is referred to [70] for some discussion on this point. Cl

Site Specific Nature of the Impurity

Absorption Coefficient (cm1)

The LVM frequency depends primarily on its mass and to a lesser extent on the masses of its nearest neighbours [52]. Impurities occupying Ga lattice sites have 75As neighbours while impurities occupying As lattice sites are surrounded by different combinations Of69Ga or 71Ga neighbours with abundances of 60% and 40% respectively, giving a spread of LVM frequencies and fine structure in the absorption feature. Such fine spectroscopy detail has been invaluable in understanding the compensation processes which occur with amphoteric impurities such as Si (see FIGURE 1).

Wavenumbers (cm 1) FIGURE 1. LVM spectrum (4 K) obtained from a heavily Si-doped Bridgman crystal. Absorption from 28Si(Ga), 29 Si(Ga) and 30Si(Ga) donors is evident at 384, 379 and 373 cm'1 respectively. The Si(As) LVM line at 399 cm"1 shows the fine structure exhibited by all 'light' impurities substituting for As atoms [3]. Other features include Si(Ga)-Si(As) donor-acceptor pairs at 393 cm'1 and Si-X and Si-Y defects around 365 cm'1 [30]. The background two-phonon absorption has been subtracted.

TABLEl . Vibrational mode frequencies of substitutional impurities in GaAs. Impurity 1

H H

1

2

D D 6 Li(Ga) 7 Li(Ga) 9 Be(Ga) 10 B(Ga) 11 B(Ga) 10 B(As) 11 B(As) 12 C(As) 13 C(As) 14 N(As) 16 O 18 O 24 Mg(Ga) 25 Mg(Ga) 26 Mg(Ga) 27 Al(Ga) 28 Si(Ga) 29 Si(Ga) 30 Si(Ga) 28 Si(As) 30 Si(As) 31 P 2

0

LVM frequency (cm-1) 1832 1998, 1838 1765,1711 1318 1442, 1321, 1291 482 450 482 540 517 628 601 582.8,1157° 561.8 480 836 790 331 326 322 362 384 379 373 398 389 355

Comments

Ref.

implanted implanted

[1] [2]

implanted implanted Li-diffused Li-diffused shallow acceptor LEC, residual LEC, residual Ga-rich material Ga-rich material LEC, residual MOMBE Bulk, implanted N-implanted

[1] [2] [3] [3] [4,51 [6-8] [6-81 [9-12] [9-12] [6,8,13-18] [6,13,17] [19] [20-23] [20-23] [24]

shallow acceptor shallow acceptor shallow acceptor iso-electronic major n-dopant major n-dopant major n-dopant shallow acceptor shallow acceptor iso-electronic

T241 [24] [8,25,26] [3,7,26-30] [7,30] [7,30] [8,11,14,15,26,27,30] [3,32] [251

Second harmonic TABLE 2. Vibrational frequencies of substitutional complexes. Impurity complex 6 Li-Mn(Ga) 7 Li-Mn(Ga) 6 Li-Cd(Ga) 7 Li-Cd(Ga) 6 Li-Zn(Ga) 7 Li-Zn(Ga) 6 Li-Te(Ga) 7 Li-Te(Ga) 6 Li-Si(Ga) 7

Li-Si(Ga)

LVM frequency (cm1) 391,413,419 365, 386, 391 377,401,423 354, 375, 395 361,385,404,433 340,361,378,405 419,510 391,475 374, 379,405, 476, 480 487 374, 379, 405,438,447 454

Comments Li-diffused Li-diffused Li-diffused Li-diffused Li-diffused Li-diffused Li-diffused Li-diffused Li-diffused

Ref. [331 [33] [33,34] [33,34] [33] [331 [33,34] [33,34] [3,35]

Li-diffused

[3,35]

Table 2. (Continued) Impurity complex 6 Li-Mg(Ga) 7

Li-Mg(Ga)

6

Li complexes

7

Li complexes

9

Be(Ga)-28Si(Ga)

28

Si(Ga)-11B(As) Si(Ga)-10B(As)

LVM frequency (cm"1) 320, 338, 350 392,404,419 318,338,349 367, 377 392 351,389,406 410,451 328, 364, 379 383,421 388.1,388.8 349,570.9,661

28

352, 596,684.8

28

393, 464 456, 389 457,389 449, 384 373, 403 374, 376, 399 378, 382, 395

Si(Ga)-28Si(As) 28 Si(Ga)-30Si(As) 30 Si(Ga)-28Si(As) 30 Si(Ga)-30Si(As) 28 Si(Ga)-Ge(As) 28 Si(Ga)-Cu(Ga) 28 Si(Ga)-Zn(Ga) 28

Si(Ga)-X Si(Ga)-X 28 Si(Ga)-Y 11 B(As)-Ge(Ga)

369 359 367 582, 587?

10

608,614?

30

B(As)-Ge(Ga)

11

566?

10

591?

11

580.7,606.5, 622.6 605,633, 650

B(As)-Sn(Ga) B(As)-Sn(Ga) B(As)-Te(As)

10

B(As)-Te(As)

11

B(As)-Se(As)

10

B(As)-Se(As)

576.4, 609.4 621.7 601,636.5,649

Mg(Ga)-Se(Ga)

335

Mg(Ga)-Te(Ga)

337, 350

Comments Li-diffused

Ref. [24]

Li-diffused

[24]

Li-diffiised

[34,36,37]

Li-diffused

[34,36,37]

doubly-doped Si+Be doubly-doped B+Si doubly-doped B+Si heavily Si-doped heavily Si-doped heavily Si-doped heavily Si-doped doubly-doped Cu-diffiised Si doped doubly-doped Si+Zn heavily Si-doped heavily Si-doped heavily Si-doped doubly-doped Si+Ge doubly-doped Si+Ge doubly-doped Si+Sn doubly-doped Si+Sn doubly-doped B+Te doubly-doped B+Te doubly-doped B+Se doubly-doped B+Se Mg-doped Li diffused Mg-doped Li diffused

[38] [9,14,40,43,44] [9,14,40,43,44] [3,29,31,32,45] [3,32] [3,32] [3,321 [46,47] F261 [42,48] [29-32,38] [32] [29-32,38] [46] [46] [46] [46] [9,40-42] [9,40-42] [9,40-42] [9,40-42] [24] [24]

TABLE 3. LVM frequencies of irradiation induced complexes. Impurity complex 11 B(I) 10 B(I) 12 C(I)

LVM frequency (cm1) 371,641,763 387,669, 796 577,606

Comments

Ref. [8,50,51] [8,49-51] [9,11,511

TABLE 4. LVM frequencies of H modes in H-related complexes. Impurity complex 1 H-Be(Ga) 2 H-Be(Ga) 1 H-Si(Ga) 2 H-Si(Ga) iH-Si(As) 2H-Si(As) 1

H-Zn(Ga) H-Zn(Ga) 1 H-12C(As) 2 H-12C(As) 1 H-13C(As) 2 H-13C(As) 1 H-(C(As^2 2

1

H-Mg(Ga) 1 H-Ge 1 H-Sn 2 H-Sn transition metal complexes

LVM frequency (cm'1) 2037.1** 1471.2** 896*, 1717** 637*, 1248** 972.2*, 2094.7** 703.5*, 1514.5** 2146.9** 1549.1** 2635.2** 1968.6** 2628.5** 1958.3** 2636,2643,2651 2688 2144** 2010** 746.6*, 1327.8** 967.7** 2001.0,2009.5, 2011.4,2020.6, 2024.3,2051.8

Comments

CBE CBE CBE CBE MOMBE C-complexes

Ref. [5,53,54] [5,53,54] [53,55] [53,551 [56,571 [56,57] [58,59] [58,59] [60,61] [60,61] [60,61] [60,61] [62] [631 [63,64] [64,65] [64,65] [66]

* denotes wag mode ** denotes stretch mode TABLE 5. Modified LVM modes of H-impurity complexes Impurity 9 Be-1H 9 Be-2H 12 C(As)-1H 13 C(As)-1H 12 C(As)-2H 13 C(As)-2H 12 C(As)-1H 13 C(As)-1H 28 Si-1H 28 Si-2H

LVM frequency (cm'1) 555.7 553.6 452.7 437.8 440.2 426.9 562.6 547.6 409.95 409.45

Comments

A1 mode A1 mode A1 mode A1 mode E mode E mode

Ref. [5,531 [5,53] [60,61] [60,61] [60,61] [60,61] [60,61] [60,61] [53,551 [53,551

TABLE 6. Calibration factors for selected LVM features. Impurity

Calibration factor

Comments

Ref. [4] [5] [67,68]

(XlQ 16 Cm- 3 ) 9

Be

6.5(77K) 5.7 (4 K) 5.0 (4 K)

B, '0B [_C

6.0 (77 K) 2.4 (77 K) Lj 1.3±0.3

Hall Hall Hall/SIMS/ Raman SSMS/AAS SSMS/AAS Hall Hall

1.1&JQ.2 0.8±0.2

Hall Hall

[71] [72]

11

[8] [8] [69] [70]

0,95±0.29

SMS

[73]

__AJ 28 Si(Ga) Si(Ga)ZSi(As)

3.1±0.4(77K) 5.3 (77 K) 6.0±1.0(4K)

SSMS/AAS SSMS Hall/SIMS

[8] [28]

Si(Ga)ZSi(As)

5.0±0.4 (4 K)

Hall

PO]

Si(As)

6.8 (77 K) 7.0

SSMS/AAS Hall MBE and LPE

[8] [74]

Si-X

I

2.7±1.0(4K)

I

Hall

I

[301

Calibration corresponds to an integrated absorption of 1 cm'2 at 300 K unless otherwise indicated. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] II1] [12] [13] [14] [15] [16] [17] [18] [19]

R.C. Newman, J. Woodhead [Radiat. Effects (UK) vol.53 (1980) p.41 ] J. Tatarkiewicz, A. Krol, A. Breitschwerdt, M. Cardona [Phys. Status Solidi (Germany) vol.140 (1987)p.367] W.M. Theis, W.G. Spitzer [ J. Appl. Phys. (USA) vol.56 (1984) p.890 ] K. Laithwaite, R.C. Newman, P.D. Greene [ J. Phys. C (UK) vol.8 (1975) p.L77 ] P.S. Nandhra et al [ Semicond Sa. Technol. (UK) vol.3 (1988) p.356 ] RC. Newman, F. Thompson, M. Hyliands, R.F. Peart [Solid State Commun. (USA) vol. 10 (1972) p.505] F. Thompson, R.C. Newman [ J. Phys. C (UK) vol.5 (1972) p. 1999 ] M.R. Brozel, J.B. Clegg, R.C. Newman [ J. Phys. D (UK) vol. 11 (1978) p. 1331 ] F. Thompson, S.R. Morrison, RC. Newman [Inst. Phys. Conf. Ser. (UK) vol 16 (1973)p.371 ] S.R. Morrison, R.C. Newman, F. Thompson [ J. Phys. C (UK) vol.7 (1974) p. 633 ] J. Woodhead, R.C. Newman, I. Grant, D. Rumsby, R.M. Ware [ J. Phys. C (UK) vol. 16 (1983) p.5523 ] G.A. Gledhill, R.C. Newman, J. Woodhead [ J. Phys. C (UK) vol. 17 (1984) p. L307 ] M.R. Brozel, R.C. Newman [ J. Phys. C (UK) vol. 11 (1978) p.3135 ] W.M. Theis, K.K. Bajaj, CW. Litton, W.G. Spitzer [Appl. Phys. Lett. (USA) vol.41 (1982) p.70 ] RS. Leigh, R.C. Newman [ J. Phys. C (UK) vol. 15 (1982) p.L1045 ] KLWoodhouse et al [ Semicond. Sci. Technol. (UK) vol.6 (1991) p.330 ] K.Woodhouse et al [ J Cryst. Growth (Netherlands) vol. 120 (1992) p.323 ] L.Z. Zhang, YC. Du,B.R. Zhang, Y.H. Wang,B.C. Ma,GG. Gin [J. Phys., Condens. Matter (UK) vol.1 (1989) p.4025] A.A. Kachare, W.G. Spitzer, A. Kahan, F.K. Euler, T.A. Whatley [ J. Appl. Phys. (USA) vol.44

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7.2

IR absorption due to free carriers in GaAs W. Szuszkiewicz April 1996

A

INTRODUCTION

Al

General Remarks

When free carriers are present in a semiconductor optical properties can differ significantly from those of the intrinsic material. Some physical phenomena (e.g. free-carrier absorption or intervalence band optical transitions) simply cannot occur without free carriers, while others, like interband absorption, are modified on changing the free carrier concentration. The change in the optical absorption mechanisms produces a change in the corresponding absorption coefficient value (and hence in the extinction coefficient). Due to the Kramers-Kronig relations the refractive index changes also. Finally both the real and imaginary parts of the complex dielectric function e(a)) differ from those of the intrinsic semiconductor even far from the spectral region corresponding to the absorption mechanism under consideration. The free-carrier-related absorption mechanisms are important over a wide energy range starting from zero up to energies above the value corresponding to the energy gap. Many review papers have considered the infrared optical properties of GaAs, including free-carrier absorption (see, for example, [1-5]). This Datareview is limited to experimental data obtained by various techniques for optical studies as well as results of theoretical investigations concerning the following absorption mechanisms: (a) (b) (c)

Free-carrier absorption Inter-conduction band absorption Inter-valence band absorption

Section B gives a brief overview of effects in the far- and middle-infrared (including the reststrahlen) region of the spectrum resulting from mechanism (a). Sections C and D are devoted to mechanisms (b) and (c), respectively, and describe effects in the mid-infrared part of the spectrum. A2

Plasmon-Phonon Coupling

When free carriers are present in a semiconductor, one of the important parameters is the plasma frequency G)p related to the carrier concentration N by the formula: 2 *>P

4 TT e 2 N =

; m

(D e

where m* is the free carrier effective mass and e is the dielectric constant of the material under investigation. For a nonparabolic and/or warped band, m* is equal to the effective mass value averaged over those energies in the region where the Fermi-Dirac distribution function has a nonzero derivative value. In the degeneracy case m* corresponds simply to the effective mass value at the Fermi level mF*. The dielectric constant should be taken as equal to e0 when the plasma frequency is below the lattice excitation frequency range (reststrahlen region) or as equal to E00 when the plasma frequency is above that region. The approximate value a)p for n-GaAs can be obtained from the formula: o)p ( c m 1 )

= 3.4 x i o 7 ^N(cm 3)

(2)

o)p corresponds to the plasma energy, Ep, determined roughly from the simplified equations: Ep2 = 2.0 x 10"21Ncm3eV2

(3)

in the case of n-type GaAs and Ep2 = 3.0 x 10- 22 Ncm 3 eV 2

(4)

in the case of p-type material. EQNs (2) and (3) do not take into account the energy dependence of the carrier effective mass which can be important for high free electron concentration regions. When the plasma frequency is close to the frequency of the lattice modes strong coupling between plasmon and phonon modes occurs. As a result instead of two separate LO-phonon and 'pure' plasma modes, two longitudinal modes having mixed phonon-plasmon character arise. Roughly speaking the mode with a dominant plasma-like character can be observed close to a)p only in the spectral range 'far' from the TO and LO phonon frequencies on both sides of the reststrahlen band. Many papers have reported Raman scattering studies of n-GaAs [6-8] although only in [7] has the nonparabolicity of the conduction band been taken into account. The wave-vector dependence of the LO-phonon-plasmon coupling was analysed in [9-11] although the analysis has been limited to the parabolic conduction band approximation. Raman scattering measurements enable one to determine the free carrier concentration for ultrahigh n-doped GaAs from the coupled mode frequency values [12]. Raman scattering was also recently used to estimate free carrier concentrations in GaAs oval defects by the analysis of TO, coupled and uncoupled LO phonon mode intensities [13]. The plasmon-phonon coupling in p-type material differs both qualitatively and quantitatively from that in n-type GaAs. For instance, the finite lifetime of collective modes plays a dominant role leading to a single damped (plasmon - LO-phonon) coupled mode. This strong damping results mainly from the larger effective masses of free holes compared to free electrons. For a proper description of the concentration dependence of the coupled mode energy both the extrinsic plasmon damping and the additional intrinsic damping provided by the inter-valence band transitions should be taken into account (see Section D). The interaction of free holes with optical phonons in p-GaAs has been investigated for a very wide hole concentration range from 1.0 x 1016 cm'3 to 1.4 xio 21 cm"3 by Raman scattering measurements [14-22]. The position of the peak in

the coupled mode spectra [19,20] and its intensity [21] has been proposed as a nondestructive means of determining the hole concentration. The optical hole mobility can also be obtained from line-shape analysis of the coupled mode [22]. However, this value is about two times underestimated in comparison to the Hall mobility. B

FREE-CARRIER ABSORPTION

Bl

Free-Carrier Absorption in n-GaAs

Bl. 1 Absorption mechanisms Free-carrier absorption results from the optical excitations of electrons occupying states in the partially filled energy minimum of the conduction band, in the vicinity of the F point of the Brillouin zone, when carriers remain in the same minimum. Two kinds of excitations contribute to this absorption: individual excitations and collective excitations. In the case of individual excitation an electron absorbing a photon must also be scattered by some crystal imperfection in order to conserve both energy and wave vector. There are three main contributions to such free-carrier absorption, due to the interaction of the electron with: (a) (b) (c)

screened ionized impurities (Coulomb interaction), screened optical phonons (polar interaction), acoustic phonons (deformation interaction).

After optical excitation into a higher state of the same energy minimum the electron can thermalize, losing its energy to the crystal lattice. Collective excitations correspond to the excitation of oscillations (the generation of plasmons) of the free-electron plasma by light. This is normally a forbidden absorption process which is allowed when ionized impurities or defects are present in a semiconductor. Such absorption is called 'free-carrier absorption by photon-ionized impurity-plasmon processes' in the literature. These oscillations are damped by plasmon decay giving rise to individual electron excitations. Free-carrier absorption due to individual electron excitations in n-GaAs, which results in relatively high values of absorption coefficient and has the greatest practical impact, is the best known. At frequencies much higher than a)p, free-carrier absorption related to individual electron excitations has been measured for values smaller than 1000 cm"1 [23-27]. Experimental data taken close to (op, corresponding to absorption coefficient values exceeding 3000 cm'1, can be found in [28,29]. The general theory of free-carrier absorption in n-GaAs can be found in [30,31], while the theory of free-carrier absorption in the frequency range much above cop is given in [32,33]. There exist many reviews (see [1-5]), and useful condensed information can be found in [34]. Free-carrier absorption due to individual electron excitations occurs in a wide spectral range both below and above Ep; the absorption coefficient exceeds 2500 cm"1 at the plasma energy and strongly decreases with photon energy. For electron concentrations exceeding 1.0 x 1018 cm"3' mechanism (a) is dominant at room temperature; mechanism (b) prevails at lower concentration.

B 1.2

Calculations

For electron concentrations lower than 3.Ox io18 cm"3 and at frequencies significantly higher than 0)p the free-carrier absorption, a, is a complicated function of the wavelength, A, and has the form: a(A) = A A7/2 + B A5/2 + C A3/2

(5)

where these three terms correspond to the three absorption mechanisms (a), (b) and (c), respectively. A, B and C are constants. Theoretical values of the three main contributions to the total free-carrier absorption coefficient value a at A = 10 \xm and over a wide carrier concentration range at room temperature have been tabulated in [33] (see TABLE 1). In particular, the idea of the optical determination of the compensation ratio by means of free-carrier absorption was given in [33]. This idea has been developed in [35], where it was demonstrated that both the electron concentration and the compensation ratio can be determined from the value of the free-carrier absorption at any wavelength between 8 ^m and 12 \im. Theoretical values of A and B can be found in [4]. Recently it has been pointed out that the above method may overestimate the compensation ratio for electron concentrations exceeding 1.0 * 1017 cm "3 [36] or 1.0 x 10 18 cm "3 [37] because the defect interaction leading to dipole or dislocation formation was not taken into account [33,35]. The influence of dipole scattering on the freecarrier absorption in GaAs was calculated in [37] and free-carrier absorption induced by dislocation scattering mechanisms was studied both theoretically and experimentally in [38-40]. It should be stressed that in [37] and [38-40], respectively, the free-carrier absorption mechanisms under consideration are not significant. This is because the expected principal effect related to the creation of extended defects results simply from the reduction of the carrier density [36-40]. In a wide frequency range lower than G)p, where the photon energy is small in comparison to the energy of the electrons , the classical Drude-Zener theory becomes valid: (X(A) = A1 A2

(6)

where A1 is a new constant [30]. The free-carrier absorption due to individual excitations saturates for GaAs in the far-infrared, at least at the spectral range below 7 meV (60 cm"1) [41,42]. Some information about other, less important, scattering mechanisms can be found in the literature. The contribution of piezoelectric scattering has been estimated in [33], and the electronelectron scattering in [43]. The influence of the non-parabolicity of the conduction band and the short-range components of the impurity potential on the free-carrier absorption observed in highly n-doped GaAs was discussed in [43,44] and [25,26,43], respectively.

TABLE 1. Theoretical values of three principal contributions to the total free-carrier absorption coefficient value a for A=IO ^m at room temperature taken from [33]. The compensation ratio K has been assumed to be equal to zero. In the case of non-zero compensation ratio K one should multiply the absorption coefficient value (a) by the value (1 + K)/(l-K). Electron concentration /cm-3x

B1.3

Absorption coefficient (in cm'1) due to scattering by: impurities (a)

optical phonons (b)

acoustic phonons (c)

1.OxIO16 1.5xlO16 2.OxIO16 3.0xl0 16 4.OxIO16 5.OxIO16 6.OxIO16 7.OxIO16 8.OxIO16 9.OxIO16

0.0036 0.0080 0.0142 0.032 0.056 0.087 0.125 0.169 0.220 0.277

0.341 0.510 0.679 1.10 1.35 1.68 2.00 2.32 2.65 2.97

0.025 0.038 0.050 0.075 0.100 0.126 0.152 0.178 0.203 0.229

LOxIO17 1.5xl017 2.OxIO17 3.0xl0 17 4.OxIO17 5.0xl0 17 6.OxIO17 7.OxIO17 8.OxIO17 9.OxIO17

0.340 0.746 1.29 2.79 4.82 7.26 10.2 13.5 17.2 21.3

3.28 4.84 6.35 9.26 12.1 14.8 17.5 20.1 22.7 25.3

0.255 0.387 0.521 0.796 1.08 1.34 1.68 1.99 2.32 2.65

1.OxIO18 1.5xlO18 2.OxIO18 3.0xl0 18 4.OxIO18 5.OxIO18

25.8 53.4 88.8 178 284 400

27.9 40.5 52.8 76.6 98.6 120

3.00 4.85 6.91 11.6 16.9 22.8

Other effects

General expressions describing the influence of quantising magnetic fields on free-carrier absorption due to acoustic phonon scattering in n-GaAs are given in [45,46] where the nonparabolicity of the conduction band has been taken into account. Results of numerical calculations performed at a wavelength of 10.6 ^m and electron concentration of 1.73 x 1015 cm"3 demonstrated that the absorption coefficient oscillates with the magnetic field. These oscillations diminish with decreasing temperature and vanish at high magnetic fields while the absorption coefficient decreases with decreasing temperature for any value of the magnetic field. The influence of the free-carrier absorption on the refractive index at important near-infrared wavelengths used for integrated optics and optical-probing applications was discussed in [47]. The contribution to the free-carrier absorption coefficient by collective excitations could be significant in the narrow spectral region just above 0)p (plasmon-phonon mode). Such absorption takes place from Ep up to no more than about 1.4E p , the frequency range depending both on

temperature and electron concentration. This absorption gives rise to a hump in the wavelength dependence of the free-carrier absorption, its height being equal to at least 15% of the absorption 'background' due to all individual excitations [5]. Free-carrier absorption due to collective excitations has been experimentally demonstrated in [28,29]; corresponding values of the total free-carrier absorption coefficient were higher than 3000 cm"1. It was also shown in the case of thin films below a few microns that even for pure material nominally forbidden absorption of light due to plasmon excitations occurs for oblique incidence. This effect results from charges trapped at the film interface [48]. B2

Free-Carrier Absorption in p-GaAs

Free-carrier absorption in p-GaAs with hole concentration up to 1.5 x 1020 cm"3 was measured and theoretically calculated in [49]. For these calculations the frequency dependent conductivity resulting from the absorbed power was used. It has been demonstrated that the relevant absorption coefficient value is high enough (20 000 cm"1 for a hole concentration of 4.5 x 1019 cm"3) to be important for optical devices. The theoretical analysis of free-carrier absorption in pGaAs by scattering from acoustic phonons, polar optical phonons and ionized impurities is given in [50]. In the same paper the influence of this absorption on the refractive index value is also discussed. B3

The Plasma Edge

Free-carrier absorption is responsible for a strong modification of the reflectivity coefficient R in the spectral region close to cop giving rise to the plasma edge. R is described by the following formula: (n - I) 2 + k 2 R = -^ '(n + I) 2 + k 2 o

(7) U)

where n is the refractive index and k is the extinction coefficient, related to the absorption coefficient a and wavelength X by: , k

ct X

= T^

<8

>

When the 'pure' plasma frequency is far from the spectral range of optical phonons the exact position of the plasma edge depends on the energy of the coupled plasmon-phonon mode. From the position of the plasma minimum on the infrared reflectivity curve the free electron concentration for n-GaAs can be determined with high accuracy. When the pure plasma frequency is in the vicinity of, or corresponds to, the phonon modes the analysis of the reflectivity curve is much more complicated and should be done numerically. The reflectivity spectra of n-GaAs were analysed in detail both experimentally and theoretically in [7,51]. The theoretical analysis of the shape of the reflectivity curve (neglecting possible coupling with the optical phonons) can be found in [4,52]. In these papers the concentration dependence of the plasma minimum energy position in n-GaAs is described by the formula:

log [N(Cm-3)] = 3.8 log [O)1n(Cm"1)] + 4.8

(9)

where 0)m is the plasma minimum frequency. A comparison of the reflectivity spectra for n-type, intrinsic and p-type GaAs can be found in [53]. C

INTER-CONDUCTION BAND ABSORPTION

Inter-conduction band absorption corresponds to electron transitions from the lowest minimum of the conduction band (T point of the Brillouin zone) to a subsidiary minimum. More than 90% of the relevant absorption coefficient is due to the contribution arising from electron scattering by acoustic phonons [54]. This absorption can be observed in the energy range from about 0.4 eV (3200 cm"1) up to 0.9 eV (7200 cm -1 ); reported values of the absorption coefficient after the sharp increase remain nearly constant in a wide energy range [23,24,26,55]. For example, a does not exceed 30 cm"1 for an electron concentration as high as 5.0 x 1018 cm"3 and strongly decreases with decreasing concentration. However, for a given electron concentration, in a large part of the energy region mentioned above, absorption (b) is more important than (a). Absorption (b) has been analysed theoretically in [54,44], with the theoretical curve being compared with earlier experimental data taken from the literature. Useful figures can be found in [34] and comments in [47,56]. It should be mentioned that some authors call the absorption related to interconduction band transitions 'free-carrier absorption'. D

INTER-VALENCE BAND ABSORPTION

Inter-valence band absorption results from optical transitions between subbands of the valence band. There exist three bands of this absorption due to transitions: (a) from the light hole to the heavy hole valence band (peak position close to 0.15 eV), (b) from the spin-orbit split-off band to the light hole valence band (peak position 0.31 eV) and (c) from the spin-orbit split-offband to the heavy hole valence band (peak position 0.42 eV). Although their position is only weakly temperature dependent, their form changes drastically with temperature. Inter-valence band absorption has been analysed both experimentally and theoretically in [57-59]; a theoretical analysis can be found in [60]. It was demonstrated that for p-GaAs with a hole concentration of about 4.5 x 1019 cm"3 the absorption coefficient could be as high as 2 x 104 cm"1 [49]. It has been also estimated that the absorption coefficients for inter-valence band transitions are, for a given hole concentration, 20 times higher than corresponding values for n-GaAs with the same carrier concentration [61]. E

CONCLUSION

When free carriers are present in GaAs the optical properties of the material change in comparison with those of the intrinsic crystal. These changes are not limited only to the spectral region corresponding to the particular free carrier related absorption mechanism, but due to KramersKronig relations can be noticeable or significant in the whole energy region from zero to energies exceeding the bandgap. The analysis of the optical constants, resulting from infrared transmission and/or reflectivity measurements, enables one to estimate the free carrier concentration in n-GaAs. From the analysis of the Raman scattering spectra the free carrier concentration can be estimated in both n-type and p-type materials.

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H.Y. Fan [ in Optical Properties of IU-V Compounds Eds R.K.Willardson, A.C.Beer, Semicond. Semimet. vol.3 (Academic Press, New York, 1967) ch.9 p.405-19 ] B. Jensen [ Ann. Phys. (USA) vol.80 (1973) p.284-360 ] J.S. Blakemore [ J. Appl Phys. (USA) vol.53 no. 10 (1982) p.R.123-81 ] B. Jensen [ in Handbook of Optical Constants of Solids Ed. E.D.Palik (Academic Press, New York, 1985) ch.9 p. 169-88] W. Szuszkiewicz [ inProperties of Gallium Arsenide, 2nd Edition, EMIS Datareviews Series no.2 (INSPEC, IEE, 1990) ch. 10 p. 180-4 ] A. Mooradian, A.L. McWhorter [ Phys. Rev. Lett. (USA) vol. 19 no. 15 (1967) p.849-52 ] H.R. Chandrasekhar, A.K. Ramdas [Phys. Rev. B (USA) vol.21 no.4 (1980) p.1511-5 ] Huade Yao, A. Compaan [ Appl. Phys. Lett. (USA) vol.57 no.2 (1990) p. 147-9 ] K. Murase, S. Katayama [Phys. Rev. Lett. (USA) vol.33 no.25 (1974)p.l481-4 ] U. Nowak, W. Richter, G. Sachs [Phys. Status Solidi B (Germany) vol. 108 no. 1 (1981) p. 131-43 ] CP. Chang, K.F. Pai, CS. Fang, W.S. Tse [ Phys. Status Solidi B (Germany) vol. 154 no. 1 (1989) p.135-41] M. Ramsteiner, J. Wagne, P. Hiesinger, K. Kohler, U. Rossler [ J. Appl. Phys. (USA) vol.73 no. 10 (1993)p.5023-6] P.S. Dobal, H.D. Bist, S.K. Mehta, RK. Jain [ J. Appl. Phys. (USA) vol.77 no.8 (1995) p.3934-7 ] D. Olego, M. Cardona [ Phys. Rev. B (USA) vol.23 no. 12 (1981) p.6592-602 ] D. Olego, M. Cardona [ Phys. Rev. B (USA) vol.24 no. 12 (1981) p.7217-32 ] R Fukusawa, M. Wakaki, K. Ohta, H. Okumura [ Jpn. J. Appl. Phys. (Japan) vol.25 no.4 (1986) p.652-3 ] T. Yuasa, M. Ishii [ Phys. Rev. B (USA) vol.35 no.8 (1987) p.3962-70 ] M. Gargouri, B. Prevot, C. Schwab [ J. Appl. Phys. (USA) vol.62 no.9 (1987) p.3902-11 ] K. Wan et al [ J. Appl. Phys. (USA) vol.63 no. 11 (1988) p.5598-600 ] K. Wan, J.F. Young [ Phys. Rev. B (USA) vol.41 no. 15 (1990) p. 10772-9 ] A. Mlayah, R. Carles, G. Landa, E. Bedel, A. Munoz-Yague [ J. Appl. Phys. (USA) vol.69 no.7 (1991)p.4064-70] R. Fukasawa, S. Perkowitz [ Phys. Rev. B (USA) vol.50 no. 19 (1994) p. 14119-24 ] W.G. Spitzer, JM. Whelan [Phys. Rev. (USA) vol. 114no. 1 (1959)p.59-63 ] MG. Milvidskii, V.B Osvenskii, E.P. Rashevskaya, T.G. Yugova [Sov. Phys.-Solid State (USA) vol.7 no. 11 (1966) p.2784] E.P. Rashevskaya, V.I. Fistul [ Sov. Phys.-Solid State (USA) vol.9 no.6 (1967) p. 1443 ] E.P. Rashevskaya, V.I. Fistul [ Sov. Phys.-Solid State (USA) vol.9 no. 12 (1967) p.2849 ] K. Osamura, Y. Murakami [ Jpn. J. Appl. Phys. (Japan) vol. 11 no.3 (1972) p.365-71 ] Vu Hai Son, K. Karpierz, W. Szuszkiewicz [AdaPhys. Pol. A (Poland) vol.73 no.3 (1988) p.353-6 ] W. Szuszkiewicz, K. Karpierz, Vu Hai Son [ Phys. Scr. (Sweden) vol.37 (1988) p.836-9 ] P. Kleinert, M. Giehler [ Phys. Status Solidi B (Germany) vol. 136 no.2 (1986) p.763-77 ] W. Bardyszewski, W. Szuszkiewicz, Vu Hai Son [ Ada Phys. Pol. A (Poland) vol.75 no. 1 (1989) p.47-50 ] E. Haga, H. Kimura [ J. Phys. Soc. Jpn. (Japan) vol. 19 no.5 (1964) p.658-69 ] W. Walukiewicz, L. Lagowski, L. Jastrzebski, M. Lichtensteiger, H.C Gatos [ J. Appl. Phys. (USA) vol.50 no.2 (1979) p.899-908 ] A.S. Jordan [J. Appl. Phys. (USA) vol.51 no.4 (1980) p.2218-27 ] L. Jastrzebski, L. Lagowski, W. Walukiewicz, H.C. Gatos [J. Appl. Phys. (USA) vol.51 no.4 (1980) p.2301-3 ] D. Wruck, A. Knauer [ Phys. Status Solidi A (Germany) vol.107 no.l (1988) p.321-8 ] H. Boudriot, H.A. Schneider [ Phys. Status Solidi B (Germany) vol. 154 no.2 (1989) p.K175-8 ] A. V. Bazhenov, L.L. Krasilnikova [ Sov. Phys.-Solid State (USA) vol.28 no. 1 (1986) ]

[39] [40] [41] [42] [43] [44] [45] [46] [47] [48] [49] [50] [51] [52] [53] [54] [55] [56] [57] [58] [59] [60] [61]

D. Vignaud, J.L. Farvacque [ Phys. Status Solidi B (Germany) vol. 156 no.2 (1989) p.717-23] D. Vignaud, J.L. Farvacque [J Appl. Phys. (USA)vol.67 no.l (1990) p.281-6 ] S. Perkowitz [J Phys. Chem. Solids (UK) vol.32 no.10 (1971) p.2267-74 ] H.C. Huang, S. Yee, M. Soma [ J. Appl. Phys. (USA) vol.67 no.3 (1990) p. 1497-503 ] P.Pfefifer,IGorczyca,W.Zawadzki[5o//fi?5toteCo/M»7M«. (USA) vol.51 no.3 (1984) p.179-83] Sunh-Huei Lin, Yi-Chen Cheng [ Chin. J. Phys. (USA) vol. 19 no.2,3 (1981) p.76-86 ] Chhi-Chong Wu, Jensan Tsai, Chau-Jy Lin [ J. Appl. Phys. (USA) vol.68 no.6 (1990) p.3024-6 ] Chhi-Chong Wu, Jensan Tsai, Chau-Jy Lin [ Phys. Rev. B (USA) vol.43 no.9 (1991) p.7328-31 ] H.C. Huang, S. Yee, M. Soma [ J. Appl. Phys. (USA) vol.67 no.3 (1990) p. 1497-503 ] D. Kirillov, D. Liu, Shang-Lin Weng [Appl. Phys. Lett. (USA) vol.55 no.21 (1989) p.2199-201 ] M.L. Huberman, A. Ksendzov, A. Larson, R. Terhune, J. Maserjian [ Phys. Rev. B (USA) vol.44 no.3 (1991) p. 1128-33] H.C. Huang, S. Yee [ J. Appl. Phys. (USA) vol.70 no.2 (1991) p.925-9 ] R.T. Holm, J.W. Gibson, E.D. Palik [ J Appl. Phys. (USA) vol.48 no. 1 (1977) p.212-23 ] B. Jensen [ Phys. Status Solidi B (Germany) vol.86 no. 1 (1978) p.291-301 ] J.A.A. Engelbrecht, LG. Lee, D.J.L. Venter [ Infrared Phys. (UK) vol.27 no. 1 (1987) p.57-62 ] E. Haga, H. Kimura [ J. Phys. Soc. Jpn. (Japan) vol. 19 no.9 (1964) p. 1596-606 ] I. Balslev [ Phys. Rev. (USA) vol. 173 no.3 (1968) p.762-6 ] H. KaIt [ Springer Ser. Solid-State Sci. (Germany) vol. 120 (1996) ch.4 p. 125-57 ] R. Braunstein [ J. Phys. Chem. Solids (UK) vol.8 (1959) p.280-2 ] R. Braunstein, E.O. Kane [ J. Phys. Chem. Solids (UK) vol.23 (1962) p. 1423-31 ] CH. Henry, RA. Logan, F.R Merritt, J.P. Luongo [ IEEE J. Quantum Electron. (USA) vol.QE-19 no.6 (1983) p.947-52] M. Takeshima [ Phys. Rev. B (USA) vol.32 no. 12 (1985) p.8066-70 ] Yu.I. Ukhanov [ in Optical Properties of Semiconductors (Nauka, Moscow, 1977) in Russian ]

7.3

Electronic absorption bands of impurities and defects in GaAs A.M. Hennel Updated by M.R. Brozel March 1996

A

INTRODUCTION

The electronic absorption bands of impurities and defects in semiconductors can be divided into two major groups - intracentre transitions and band-to-level photoionization transitions. The first group should manifest itself only in absorption, whereas the second can be observed in photoelectric measurements as well as in absorption. However, there are some important exceptions to this simple rule: 1)

excited states of localized centres degenerate with semiconductor bands,

2)

excited states of shallow (hydrogen-like) centres undergoing thermal ionization to semiconductor bands,

which are intracentre mechanisms with associated electrical conduction. B

INTRACENTRE TRANSITIONS INVOLVING HYDROGEN-LIKE IMPURITIES AND DEFECTS

In the case of shallow impurities (C, Mg, Zn and Si) transitions to excited states described by effective mass theory have been observed as photoconductivity low temperature peaks detected in high quality epitaxial layers. Data for manganese and gallium antisite defects were obtained from absorption measurements. Ga^ and Ga^" correspond to neutral and negatively charged states of the gallium antisite double acceptor. TABLEl Impurity or defect

Reference

Ground state energy (me V)

Transition elements from the ground state to:

IS 3 7 2 (F 8 + )

2P3/2 TO

2P5/2 TO

2P5/2 TO

C

26.9

15.2

19.4

21.3

[1]

Mg

28.7

17.1

21.1

23.1

[1]

Zn

30.6

19.4

23.2

25.0

[1]

Si

34.8

27.3

29.1

[1]

78

71.3

73.1

[2]

102.5

104.2

[3]

172.1

180.7

[4]

Mn

108(2) 203

98.5

C

EVTRACENTRE ABSORPTION OF LOCALIZED CENTRES

The most important among optically active localized centres in semiconductors (sometimes called deep impurities) are transition metal (TM) impurities and some lattice defects. Typically, intra 3d-shell electronic transitions of TM atoms consist of sharp zero-phonon lines (ZPL) observed only at low temperatures and/or associated broad bands. These broad bands are much less dependent on the temperature. In TABLE 2 positions of the ZPL are given in wave numbers (cm"1). However, energies of the broad band maxima (band max.) are expressed in electron volts (eV). Complexes are marked in square brackets. TABLE 2 Impurity

Ti3+(Sd1) Ti2+(3d2) V3+(3d2)

Ground state

2

E

3

A2

3

A2

Excited states

band max (eV)

4565.6 4589.4

0.64

[5]

T1(F) T1(P)

no ZPL no ZPL

0.66 1.01

[5]

T2

5957.85 5968.05 5968.25 7333 8131 10773

T2

3 3

1

E 3 T1(F) 1 A V2+(Sd3)

2

E

2

5

T2

0.79 1.1

[6,7] [6] [6,8] [6]

T,(?) T,(?)

NoZPL NoZPL

0.68 1.03

[7,8]

E

6620

0.9

[9,10]

2

Cr2+(3d4)

Reference

ZPL (cm1) 2

3

Transition energy

5

[Cr-V*.]

6770

[H]

Fe2+(3d6)

5

E

5

3002 2988 2962 2979

[12,13]

Fe3+(3d5)

6

A1

%

3057

[30,31]

Co 2+ Qd 7 )

4

% 4 T1(F) 2 E(G) 4 T1(P)

4035 no ZPL no ZPL no ZPL 11317

A2

T2

3885

[Co-Te] Ni 2+ (3d 8 )

3

T1(F)

3

T2 T1(P) 3 A2

no ZPL 8634 no ZPL

2

4615

3

Ni + (3d 9 )

0.87 and 0.94

2

T2

E

[14,15] [13,16] [14] [14] [16] [15]

0.55 1.10 1.22

[17] [18,19]

[Ni-S]

4427

[18]

[Ni-Se]

4410

[18]

[Ni-Te]

4369

[18]

[Ni-Si]

4699

[18]

[Ni-Ge]

4740

[18]

[Ni-Sn]

4627

[18]

Nb3+(4d2)

3

1

6416

[20]

EL2

1

1

8090

[21]

D

A2(?)

A1C?)

A1C?)

T2C?)

PHOTOIONIZATION ABSORPTION SPECTRA

Photoionization spectra of defects and impurities consist of very broad bands which can be characterized by an absorption edge and sometimes by broad band maxima. The energy positions of absorption edges roughly correspond to band-to-level energy distances. Sometimes these broad bands start with weak absorption lines which correspond to transitions from the 3+ state into the 2+ state with a loosely bound hole. In GaAs these types of lines were observed in the case of iron around 0.49 eV and 0.87 eV [30,32]. TABLE 3. Valence band to level transitions Impurity

Absorption energy range (eV)

Reference

Ti3+Cd1)

1.25-1.5

[5]

V3+(d2)

1.35-1.5

[6,7,8]

Cr4+(d2)

0.4-?

[22]

Cr3+Cd3)

0.8-1.5

[23]

Mn3+(d5+h)

0.1-1.5

[3]

Fe3+Cd5)

0.46-1.5

[24,25]

Co3+Cd6)

0.14-1.5

[26]

[Co-Te]

0.25-1.5

[15]

Ni3+Cd7)

0.4-?

[17]

Ni2+Cd8)

1.1-1.5

[17]

Cu

0.15-1.5

[27,28]

Ag

0.23-1.5

J28]

TABLE 4. Level-to-conduction band transitions Defect

Energy range

Reference

EL2

0.8-1.5

[29]

REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] [II] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32]

R.F. Kirkman, R.A. Stradling, PJ. Lin-Chung [J. Phys. C (UK) vol. 11 (1978) p.419 ] K.E. Elliot, D.E. Holmes, R.T. Chen, CG. Kirkpatrick [ Appl. Phys. Lett. (USA) vol.40 (1982) p.898] R.A. Chapman, W.G. Hutchinson [ Phys. Rev. Lett. (USA) vol. 18 (1967) p.443 ] WJ. Moore, B.V. Shanabrook, T.A. Kennedy [ Semi-insulating HI-VMaterials, Kah-nee-ta 1984, Eds D.C.Look, J.S.Blakemore (Shiva, Nantwich, England, 1984) p.453 ] A.M. Hennel et al [ Phys. Rev. B (USA) vol.35 (1986) p.7353 ] B. Clerjaud et al [ J. Appl. Phys. (USA) vol.58 (1985) p.4207 ] W. Ulrici, K. Friedland, L. Eaves, D.P. Halliday [ Phys. Status Solidi B (Germany) vol. 131(1985) p.719] A.M. Hennel, CD. Brandt, K.Y. Ko, J. Lagowski, H.C. Gatos [ J. Appl. Phys. (USA) vol.62 (1987) p.163] B. Clerjaud, A.M. Hennel, G. Martinez [ Solid State Commun. (USA) vol.33 (1980) p.983 ] P.J. Williams et al [ J. Phys. C (UK) vol. 15 (1982) p. 1337 ] E.C. Lightowlers, M.O. Henry, CM. Penchina [ Inst. Phys. Conf. Ser. (UK) vol.43 (1979) p.307] G.K. Ippolitowa, E.M. Omel'yanowski [ Sov. Phys. Semicond. (USA) vol.9 (1975) p. 156 ] J.M. Baranowski, J.W. Allen, G.L. Pearson [ Phys. Rev. (USA) vol. 160 (1967) p.627 ] H. Ennen, U. Kauftnann, J. Schneider [ Solid State Commun. (USA) vol.34 (1980) p.603 ] B. Deveaud et al [ J. Phys. C (UK) vol.19 (1986) p. 1251 ] A.M. Hennel, S.M. Uba [ J. Phys. C (UK) vol.11 (1978) p.4565 ] W. Ulrici et al [Mater. Sci. Forum (Switzerland) vol. 10-12 (1986) p. 1251 ] H. Ennen, U. Kaufmann, J. Schneider [Appl. Phys. Lett. (USA) vol.38 (1981) p.355 ] W Drozdzewicz et al [ Phys. Rev. B (USA) vol.29 (1984) p.2438 ] S. Gabilliet, V. Thomas, J.P. Peyrade, J.Barrau, CA. Bates [ Phys. Lett. A (Netherlands) vol. 119 (1986) p. 197] M. Kaminska, M. Skowronski, J. Lagowski, J.M. Parsey, H.C. Gatos [ Appl. Phys. Lett. (USA) vol.43 (1983) p.302 ] W. Ulrici, P. Kleinert [ Phys. Status Solidi B (Germany) vol. 129 (1985) p.339 ] G. Martinez, A.M. Hennel, W. Szuszkiewicz, M. Balkanski, B. Clerjaud [ Phys. Rev. B (USA) vol.23 (1981)p.3920] E.M. Omel'yanowski, L.Ya. Pervova, E.P. Rashevskaya, N.N. Solov'ev, V.I. Fistul* [ Soviet Phys.-Semicond. (USA) vol.4 (1970) p.316 ] M. Klevermanet al [J. Appl. Phys. (USA) vol.54 (1983) p.814 ] J.M. Baranowski, M. Grynberg, E.M. Magerramov [Phys. Status Solidi B (Germany) vol.50 (1972) p.433] F. Willmann, M. Blatte, HJ. Queisser, J. Treusch [ Solid State Commun. (USA) vol.9 (1971) p.2281] M. Blatte, F. Willmann [ Opt. Commun. (Netherlands) vol.4 (1971) p.178 ] G.M. Martin [ Appl Phys. Lett. (USA) vol.39 (1981) p.747 ] K. Pressel, G. Ruckert,K. Thornke.A. Domer [Mater. Sci. Forum (USA) vol.83-87 (1992) p.695 ] K. Pressel, G. Ruckert,A. Dorner,K. Thonke[ Phys. Rev. B. (USA) vol.46 (1992) p. 13171 ] A.M. Hennel, A. Wysmolek,R Bozek,D. Cote, C Naud [Mater. Sci. Forum (USA) vol.83-87 (1992) p. 729 ]

7.4

Energy levels due to transition metals in GaAs A.M. Hennel September 1996

A

ENERGY LEVELS OF TRANSITION METAL IMPURITIES IN GaAs

The energy levels of transition metal impurities are collected in TABLE 1. Energy is given either relatively to the top of the valence band (Ev) or to the bottom of the conduction band (Ec). From the electrical point of view there are three kinds of transition metal (TM) levels in GaAs. At the lowest energy is located a donor level (D) between the 4+ and 3+ charge states. Then there is an acceptor level (A) between the 3+ and 2+ charge states, and at the highest energy is located a double acceptor (AA) between the 2+ and 1+ charge states. Typically two of these levels will exist in the GaAs energy gap. The TM levels were observed and identified with the help of many different techniques - optical absorption (OA), temperature dependent Hall effect (TDH), photoluminescence (PL), photoluminescence excitation (PLE), deep level transient spectroscopy (DLTS), Hall effect under hydrostatic pressure (HHP) and others. The pressure coefficients of some levels have been measured. Their values are also given in TABLE 1 (in meV/GPa) relatively to the valence band (VB) or the conduction band (CB). TABLEl Impurity

Energy

Identification

Experimental technique

Pressure coefficient (meV/GPa)

Ref.

Ti

Ev + 0.6eV

D

DLTS

-87 ± 25(CB)

[1-3]

Ti

Ec- 0.2 eV

A

DLTS

-116 ± 12(CB)

[1-3]

V

E c -0.15eV

A

DLTS

-116 ± 12(CB)

[2,4]

Cr

E v +0.32eV

D

TDH

Cr

Ev + 0.74eV

A

OA

30 ±7(VB)

[6]

Cr

Ec + 0.05eV

AA

HHP

-63 ± 5(CB)

[7]

Mn

E v + 0.1 I e V

A

OA, PL

-63 ± 2(VB)

[8-11]

Fe

Ev + 0.5eV

A

DLTS, PLE

[12,13]

Co

Ev + 0.14eV

A

OA

[14]

Co

Ec + 0.1 IeV

AA

HHP

Co-Te complex

Ev + 0.25 eV

A

PLE

[16]

Ni

Ev + 0.2eV

A

TDH

[17]

Ni

Ec - 0.40 eV

AA

DLTS

-155±16(CB) -136±12(CB)

[18-20]

Cu

Ev + 0.15eV

A

OA, PL

3 ± 2(VB)

[21-24]

[5]

-75 ± 5(CB)

[15]

TABLEl. Continued Impurity

Energy

Identification

Experimental

Pressure coefficient

technique

(meV/GPa)

Ref.

Cu

Ev + 0.4eV

AA?

DLTS

[25,26]

Ag

Ev + 0.24eV

A

DLTS

Au

Ev + 0.40eV

A

DLTS

[27]

W

Ev + 0.2eV

D?

DLTS

[29]

W

Ev + 0.65eV

A?

DLTS

[29]

Ta

E c -0.27eV

A?

DLTs

[29]

5 ± 2(VB)

[27,28]

REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] [II] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29]

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7.5

DX centres in GaAs D.K. Maude May 1996

A

INTRODUCTION

DX centres exist in all n-type III-V semiconductors investigated to date and at a concentration comparable to the donor doping level irrespective of whether the donor is from group IV or VI of the periodic table [I]. The large difference between the thermal (~0.1 eV) and optical (-1 eV) ionisation energies suggests that the donor level undergoes a large lattice relaxation (LLR) [2]. These features, together with the relatively large thermal barrier to capture (-0.3 eV), give rise to the well known persistent photoconductivity (PPC) effect at low temperatures. Carriers freezeout onto DX centres at low temperatures but can be persistently restored to the conduction band upon illumination to photoionise the DX centre. Much of the interest which has been focused on the DX centre is due to the highly undesirable effects this deep centre has on the performance of devices based on the GaAs/(AlGa)As system [3]. B

MICROSCOPIC MODEL

Historically, the DX centre was thought to be a complex in which D is the donor impurity and X is an arsenic vacancy. This model subsequently had to be abandoned as arsenic vacancies cannot be present in sufficient concentration to explain the PPC effect observed in heavily doped GaAs. This observation is confirmed by local vibrational mode measurements which indicate that Si is present as a simple substitutional donor rather than a complex in n+ GaAs [4]. This is further

FIGURE 1. Microscopic model of the DX centre (after Chadi and Chang [7]) showing the configuration of atoms for the defect related to a donor atom from group IV (silicon) and group VI (sulphur).

supported by EXAFS measurements which indicate no perturbation of the simple substitutional donor environment when the DX centre is ionised [5]. With the disappearance of the arsenic vacancy, the LLR of the donor atom was also questioned and a small lattice relaxation (SLR) model was proposed [6] in which the DX centre is identified as an effective mass like level associated with higher (L or X) conduction band minima. The obvious weakness of the SLR model was that the PPC effect could only be explained by invoking the orthogonality of the F, X and L wavefiinctions, an unlikely explanation given that the DX centre is a deep and highly localised level (in real space) and thus delocalised in k-space. Based upon the results of a local density approximation using an 18 atom supercell, Chadi and Chang [7] have proposed that the DX centre is a negative-U centre in which the relaxed configuration corresponds to the capture of two electrons and to the breaking of a donor-arsenic bond (FIGURE 1). Yamaguchi and co-workers [8] have performed a similar calculation but with a larger 64 atom supercell and conclude that the single electron A1(T4) level is the correct groundstate. However, there is now an overwhelming experimental support for the negative-U model for the DX centre [9]. TABLE 1. Summary of chemical trends of DX centre properties in GaAs. Donor

Ec(eV)

Ee(eV)

Et(eV) DX°

Et(eV) DX-

Tc(K)

Eio(eV)

Si

0.22a 0.31 b

0.30a 0.32c 0.33g

0.26h

0.44-0.49h

120-140

l.l c

Se

90

Te

0.23c

0.39h

0.66h

90

0.6f

Sn

0.16C

0.28h

0.42-0.49h

60

0.8e'f

Ge

0.28c

0.1 l h 0.43h

40d

0.255h

Pb S

0.39c

Ec = thermal capture barrier, Ee = thermal emission barrier, Et = ionisation energy, Eio = photoionisation threshold, Tc = critical temperature for PPC quenching. a measured at P = 29 kbars [23] b from [13] c from [9] d measured at P = 15 kbars [9] e value for AlGaAs [24,25] f value for AlGaAs [26] g from[l] h recalculated from FIGURE 3.

C

EMISSION AND CAPTURE

The emission and capture processes are best understood by referring to the configurational coordinate diagram in FIGURE 2 for both the positive and negative-U models for DX. The

horizontal axis represents the lattice position of the donor impurity while the vertical axis shows the total energy (lattice plus electronic). The curve labelled U r depicts the energy of the system with the electron in the conduction band while the curve UDX depicts the energy of the system after capture. For capture and emission from DX the system must surmount the energetic barriers indicated by Ec (capture) and E e (emission). Capture takes place because the energy of DX depends on the position of the atoms surrounding it. Prior to capture DX is resonant with the Fconduction band. As the lattice vibrates, the energy of DX moves up and down. For sufficiently large vibrations, the level overlaps with the occupied vibronic states in the F-conduction band and electron capture can occur. After capture, the lattice near DX relaxes and the energy of DX is decreased. Immediately after capture, the lattice is displaced far from the new equilibrium position and there will be a violent lattice vibration at DX. This vibration is rapidly damped and the energy propagates away from DX as lattice phonons. Hence, the capture process is described as nonradiative capture by multi-phonon emission [10]. Optical ionisation (Ei0) results in a vertical transition (conservation of momentum) and hence the large lattice relaxation model provides a natural explanation for the large Stokes shift exhibited by DX. posltive-U

negative-U

FIGURE 2. Configuration co-ordinate diagram for the positive and negative-U models for DX. The horizontal axis shows the lattice position while the vertical axis shows the total energy of the system. (Redrawn from P M Mooney, Semicond ScL Technol Bl, vol.6 (1991) ).

D

FERMI-LEVEL PINNING

In GaAs the DX centre is resonant with the F-conduction band and DX centre effects in GaAs are therefore extremely limited in normal circumstances. Nevertheless the DX centre can act to pin the Fermi level at very high doping levels [11] and therefore limit the highest achievable carrier concentration to around 1.1 x 1019 cm"3 for Si-doping and 1.8 * 1019 cm"3 for Sn-doping. Fermi level pinning has also been observed in 6-doped GaAs [12]. In a similar way to increasing the Al composition in AlxGa1^xAs, the DX centre can be made to emerge into the gap by applying hydrostatic pressure to GaAs. For example the Si-DX centre emerges into the forbidden gap at around 24 kbars [13]. E

LOCAL ENVIRONMENT

Unlike the ternary (Al3Ga)As system, in GaAs there is no effect of the local environment (number of Al close neighbours) and hence there is a single DX level associated with each n-type dopant [14]. The presence of an Al nearest neighbour atom lowers the energy of the DX centre and thus,

in addition to any band structure effects, the DX centre in GaAs is generally at higher energy than in (Al5Ga)As. This is particularly important for the group VI dopants which as substitutional donors occupy the group V sites and therefore have the possibility to have Al as a nearest neighbour. A striking example is Te-DX, which is reported to be -100 meV higher in energy in GaAs than in (Al3Ga)As [15]. For AlGaAs the emission energy is independent of the Al composition (band structure). Nevertheless, the emission barrier in AlGaAs is systematically - 0 . 1 eV higher than in GaAs [15] which can be attributed to a local environment effect (one or more Al second nearest neighbours) [16]. F

CHEMICAL SHIFT

The DX centre has previously been reported to shift to higher energy with increasing doping density, an effect which has been tentatively attributed to bandgap renormalisation [18]. However Wilamowski and co-workers [19] have pointed out that the apparent dependence on doping level can be explained without invoking bandgap renormalisation if a large Coulomb broadening of the DX level of around 25 meV is assumed. The problem with pressure experiments is that the energetic position of DX has to be determined by extrapolating to zero pressure. Neglecting the Gaussian broadening of the DX energy level can lead to a significant underestimation of the experimentally deduced pressure coefficient used by extrapolation to find the energy of DX at zero pressure. This explains the widely fluctuating values for the pressure coefficient of DX to be found in the literature ( 5 - 1 2 meV/kbar). FIGURE 3 shows the single electron energy of the DX centre (the relevant energy for Fermi-level pinning) for a number of dopants as a function of doping density recalculated from the data in the figure assuming a Gaussian broadening (FWHM) of-30 - 60 meV. Within experimental error, the energies of the Si and Sn DX centres are independent of the doping density. The doping density dependence cannot be determined for the Te3 Pb and Ge DX centres as there is insufficient data. The Sn-, Si- and Pb-DX centres all have similar energies, i.e. the chemical shift is small. This is not the case for the Te-DX which is much shallower and Ge-DX which is much deeper. In Sedoped GaAs the DX centre has not been observed, despite pressure measurements up to 30 kbar indicating that the Se-DX is shallower than even the Te-DX [16]. G

DX CENTRE PARAMETERS

Deep level transient spectroscopy (DLTS) under hydrostatic pressure has been widely used to measure the capture and emission barriers of DX centres. These results together with other relevant parameters for DX centres in GaAs are summarised in TABLE 1. In AlGaAs the emission energy is independent of the Al composition [20] and is therefore also expected to be independent of applied pressure. This is not the case for the capture barrier which in AlGaAs follows the L-minima. It is standard practice to quote the capture barrier with respect to the Lminima in order to allow a comparison between different chemical species independently of the Al composition. This does not however signify that the capture occurs via an L-related state but can be understood in terms of an average over all the conduction band minima in the Brillouin zone [21]. The ionisation energy (thermal depth) of DX has been recalculated from published transport data using Fermi-Dirac statistics [22] assuming that the population of DX is cfrozen-in' and a temperature corresponding to the critical temperature Tc for PPC quenching. The energies for the

Energy (D)^ ) [meVl

Te(Wisnievvskietal) Sn (Maude etal) Si (Tachikawa) Si (Maudeetal) Si(Theisetal) Pb (Willke etal) Ge (Fujisawa et al)

FIGURE 3. Energy of the DX centre (positive-U) as a function of doping density, recalculated from the data in the references indicated in the legend, assuming a Gaussian broadening of the DX level and with a pressure coefficient of -10.8 ± 1.0 meV/ kbar with respect to the P-minimum.

positive and negative-U models were calculated by fitting to the carrier concentration versus pressure data reported in the literature. When a Gaussian broadening of the DX level is included the energy of DX is found to be independent of the doping density and the data is well fitted by a single pressure coefficient for all chemical species of -10.8 ± 1.0 meV / kbar (positive-U) or -16.0 ±1.0 meV/kbar (negative-U) measured with respect to the F-conduction band minimum. H

CONCLUSION

The DX centre in GaAs has little impact on physical properties except under extreme conditions of very high doping levels or for large applied hydrostatic pressures. Nevertheless, from a fundamental physics point of view, the GaAs system has been important for understanding the origin and nature of DX, and in particular disentangling the effects of band structure, local environment and chemical shifts in the more complicated AlGaAs system where the DX centre can seriously impair device performance.

REFERENCES [1]

[2]

For reviews see: P.M. Mooney [ J. Appl. Phys. (USA) vol.67 (1990) p.Rl ]; J.C. Bourgoin (Ed) [ Semicond. Sci. Technol (UK) vol.6 no. 10B (1991) ]; J.C. Bourgoin Ed [ Physics of the DX centre in AlGaAs alloys (Sci-Tech Pub., Vaduz, Liechtenstein, 1990) ]; E Munoz Merin (Ed) [ DX centers-donors in AlGaAs and related compounds, defect and diffusion forum vol. 108 (Scitech Pub. Ltd., Switzerland, 1994) ]; M. Saito, A. Oshiyama [Mod. Phys. Lett. B (Singapore) vol.7 (1993) p. 1567-84] D.V. Lang, RA. Logan [Phys. Rev. Lett. (USA) vol.39 (1977) p.635 ]; RJ. Nelson [Appl. Phys.

[3]

[4] [5] [6] [7]

[8] [9]

[10]

[11] [12] [13] [14] [15] [16]

[17] [18] [19] [20] [21] [22] [23] [24] [25] [26]

Lett. (USA) vol.31 (1977) p.351 ] T.N. Theis, B.D. Parker, P.M. Solomon, S.L. Wright [ Appl. Phys. Lett. (USA) vol.49 (1986) p. 1542 ]; J.R. Kirtley, T.N. Theis, P.M. Mooney, S.L. Wright [ J. Appl. Phys. (USA) vol.63 (1988) p.1541] L. Eaves etai[Inst. Phys. Con/. Ser. (UK) no.91 (1988)p.355 ] T. Kitano, M. Mizuta [Jpn. J. Appl. Phys. (Japan) vol.26 (1987) p. 1806 ]; T.M. Hayes et al [ J. Electron. Mater. (USA) vol.18 (1988) p.207] H.P. Hjalmarson, TJ. Dnimmond [Appl. Phys. Lett. (USA) vol.48 (1996) p.656 ]; J.P.M. Ansems [Semicond. Sci. Technol. (UK) vol.2(1986)p. 1 ] DJ. Chadi, KJ. Chang [ Phys. Rev. Lett. (USA) vol.61 (1988) p.873 ]; similar calculations have been performed by J. Dabrowski, M. Scheffler [ Mater. Sci. Forum (USA) vol.83-87 (1992) p.735 ] who also concluded a negative-U ground state for DX. E. Yamaguchi [Jpn. J. Appl. Phys. (Japan) vol.25 (1986) p.L643 ]; E. Yamaguchi, K. Shiraishi, T. Ohno [Jpn. J. Phys. Soc. (Japan) vol.60 (1991) p.3093 ] J.A. WoIk et al [ Phys. Rev. Lett. (USA) vol.66 (1991) p.774 ]; D.L. Williamson, P. Gibart [Semicond. Sci. Technol. (UK) vol.6 (1991) p.B70]; M. Baj, L.H. Dmowski, T. Slupinski [ Phys. Rev. Lett. (USA) vol.71 (1993) p.L388 ]; U. Wilke et al [ Proc. 22nd Int. Conf. on Physics of Semiconductors, Vancouver, 1994, Ed. DJ. Lockwood (World Scientific, Singapore) vol.3 p.2295 ] D.L. Williamson, P. Gibart [ Semicond. Sci. Technol. (UK) vol.6 No. 10b (1991) p.B70 ]; M. Baj, L.H. Dmowski, T. Slupinski [Phys. Rev. Lett. (USA) vol.71 (1993) p.3529 ]; T. Fujisawa, J. Yoshino,H. Kukimoto [ Jpn. J Appl. Phys. (Japan) vol.29 (1990) p.L388 ]; U. Willke,M.L. Fille, D.K. Maude, J.C. Portal, P. Gibart [Proc. 22nd Int. Conf. on the physics of semiconductors, Vancouver, Canada, Ed DJ. Lockwood vol.3 (World Scientific, Singapore, 1994) vol.3 p.2295 ] D.K. Maude, L. Eaves, TJ. Foster, J.C. Portal [ Phys. Rev. Lett. (USA) vol.62 (1989) p. 1922 ]; T.N. Theis, P.M. Mooney, S.L. Wright [Phys. Rev. Lett. (USA) vol.60 (1988) p.361 ] S. Arscottetal[ Semicond. Sci. Technol. (UK) vol.7 (1992) p.620 ] M. Mizuta, M. Tachikawa, H. Kukimoto, S. Minomura [ Jpn. J. Appl. Phys. (Japan) vol.24 (1985)p.L143] E. Calleja et al [ Phys. Rev. Lett. (USA) vol.56 (1990) p.934 ]; L. Dobaczewski et al [ J. Appl. Phys. (USA) vol.78 (1992) p.2468 ] P. Wisniewskietalt&ffHcoMrf. Sci. Technol. (UK) vol.6 no. 10b (1991) p.B146] O. Kumagai,H. Kawai,Y. Mori, K. Kaneko [Appl. Phys. Lett. (USA) vol.45 (1984) p. 1322]; T. Fujisawa, J. Yoshino, H. Kukimoto [J. Cryst. Growth (Netherlands) vol.98 (1989) p.243 ]; M. Mizuta, M. Tachikawa, H. Kukimoto, S. Minomura [ Jpn. J. Appl. Phys. (Japan) vol.24 (1985)p.L143];P.K. Bhattacharya, A. Majerfeld, A.K. Saxena [ Inst. Phys. Conf. Ser. (UK) vol.45 (1979) p. 199] P.M. Mooney [ J Appl. Phys. (USA) vol.67 (1990) p.Rl ] D.K Maude, L. Eaves, TJ. Foster, J.C. Portal [ Phys. Rev. Lett. (USA) vol.62 (1989) p. 1922 ] Z Wilamowski et al [ Semicond. Sci. Technol. (UK) vol.6 (1991) p.B34 ] E. CaIIeJa5I. Izupra [DXcenters-donors in AlGaAs and related compounds, defect and diffusion forum vol. 108 Ed. E. Munoz Merin (Scitec Pub. Ltd., Switzerland, 1994) p.75-95 ] DJ. Chadi, KJ. Chang [ Phys. Rev. B (USA) vol.39 (1989) p. 10366 ] T. Suski et al [ Phys. Rev. B (USA) vol.40 (1989) p.4012 ] M.F. Li et al [Appl. Phys. Lett. (USA) vol.51(1987) p.349 ] D.V. Lang [ in Deep Centres in Semiconductors, Ed. S.T. Pantelides (Gordon & Breach, New York, USA, 1986) p.489] D.V. Lang, R.A. Logan [Inst. Phys. Conf. Ser. (UK) no.43 (1979)p.433 ] G.A. Northrop, P.M. Mooney [ J. Electron. Mater. (USA) vol.20 (1991) p. 13 ]

CHAPTER 8 PHOTOCONDUCTIVITY SPECTRA 8.1 8.2

IR photoconductivity spectra of SI bulk GaAs Far-IR photoconductivity spectra of shallow donors in GaAs epilayers

8.1

IR photoconductivity spectra of SI bulk GaAs D.C. Look January 1990

A

INTRODUCTION

We will define semi-insulating (SI) GaAs as material with a resistivity of greater than 107 ohm cm. The only known impurities or defects which produce such resistivities are due to Fermi level pinning by the CrGa2+/CrGa3+ transition and EL2 (probably related to AsGa3+/AsGa4+ transition). Thus, we will discuss material dominated by one or both of these imperfections. Photoconductivity (PC) can occur by a direct transition of an electron or hole from a bound imperfection state to the relevant band, or by an indirect transition to a higher-lying state and then to a band. The former process is characterized by a 'threshold' in the PC spectrum, and the latter process by a 'peak'. B

Cr-DOPED SI GaAs

To summarize the data for Cr-doped SI GaAs, thresholds are observed at 0.5 and 0.7 eV, varying with temperature, and peaks at 0.82 and 0.88 eV, independent of temperature [1-12]. The threshold at 0.7 eV is attributed to the transition Cr2+ - Cr3+ + e" (CB)

(1)

while the peaks at 0.82 and 0.88 eV are associated with the intracentre process Cr2+(5T2) - Cr2+(5E) - Cr 3+ + e" (CB)

(2)

The magnitude of the PC for Cr-doped SI GaAs at 295 K ranges from about 10"8 to 10 "7 mho/cm at the 0.88 eV peak, and from 10"7 to 10"6 mho/cm in the bandgap region for a photon flux of 5 x 1014 cm"2 sec"1 [8]. The 77 K response in the peak region can be about two orders of magnitude higher [6]. Carrier lifetimes are typically 10"9 to 10"8 sec at 295 K [1] but can be much longer at low temperatures due to carrier trapping. Spatial fluctuations of potential can influence this process [16]. The PC can also be quenched by applying a second light source. At 82 K the quenching can be nearly 100% with a 1.1 eV source [6]. C

UNDOPED SI GaAs (CONTROLLED BY EL2)

The exact identity of EL2 remains uncertain, with current models ranging from the isolated antisite AsGa to an antisite complexed with one or more other defects (but not impurities) [16-18,24]. In fact, it is clear that there are several species of mid-gap centres which have most, but not all, electrical and optical properties in common. For simplicity, we will describe here only the mid-gap centre that is predominant in Bridgman and liquid-encapsulated Czochralski SI GaAs. This centre, defined as EL2, has a 4 K PC threshold at 0.82 eV, a zero-phonon peak at 1.039 eV, phonon-replicas separated by 11 meV, and a broad peak at 1.17 eV [14,19]. The peaks, also seen

in optical absorption, are believed to result from the following intracentre transition: AsGa3+ (1A1) - AsGa3+ (1T2)

(3)

where the 1A1 level is at Ec - 0.82 eV (the threshold) and the 1T2 level is 1.039 eV above the 1A1 level, i.e., resonant with the conduction band. In order to contribute to PC, the electron in the 1T2 state must transfer to the conduction band. The most obvious way for this to happen would be through the Fano coupling between resonant states, but another possibility is an Auger process, in which some of the energy lost in the decay back to the ground state is absorbed by electrons in shallow donor levels, which are then excited to the conduction band and increase the PC. The latter mechanism is suggested by the fact that the zero-phonon peak is stronger when sweeping downward in light energy, in which case more of the shallow donors would have trapped electrons. However, there still is some controversy regarding the interpretation of the PC and its relationship to optical absorption [20]. Some very unusual low-temperature (T < 140 K) quenching effects are observed with optical absorption, PC, EPR, photoluminescence and other experimental techniques when samples containing EL2 are irradiated with high-intensity white light, or monochromatic light centred around 1.1 eV. These are summarized in recent reviews (see, for example, [18] and [24]), but basically the broad peak is quenched by the secondary light in both optical absorption and PC, while the zero-phonon peak is quenched only in optical absorption, but is actually enhanced in PC [21]. These phenomena can be explained by excitation to an excited, metastable state EL2*. There is some controversy over whether or not EL2* is electrically active, i.e., whether or not it can compensate shallow acceptors such as C [18,22]. An argument against its electrical activity is that persistent hole photoconductivity can be observed [22]. The mechanism suggested for this effect is that: (1) electrons are excited from EL2 to EL2* or to the conduction band, leaving neutral EL2* or EL2+, respectively; (2) electrons are also excited from the valence band to EL2+ which can then create more EL2* by step (1). Eventually, all of the EL2 is converted to EL2*. If the electron on the EL2* is unable to compensate the C and the other acceptors (i.e., if the EL2* is electrically inactive), then a persistent hole conductivity will result when the light is turned off, since holes can be thermally excited from the now neutralized acceptors to the valence band. Other photocurrent quenching effects are discussed in [23]. REFERENCES [1] [2] [3] [4] [5] [6] [7] [8]

B.R. Holeman, C. Hilsum [J Phys. Chem. Solids (UK) vol.22 (1961) p.19 ] G.A. Allen [ J Phys. D (UK) vol. 1 (1968) p.593 ] D.R. Heath, P.R. Selway, CC. Tooke [ J Phys. D (UK) vol. 1 (1968) p.28 ] E.M. Omeiyanovskii, L.Ya Pervova, E.P. Rashevskaya, V.I. Fistul1 [ Sov. Phys.-Semicond. (USA) vol.5 (1971) p.484] A.A. Gutkin, A.A. Lebedev, G.N. Talalakin, T.A. Shaposhnikova [ Sov. Phys.-Semicond. (USA) vol.6 no.6 (1972) p.928-32] A.L. Lin, RH. Bube [J Appl. Phys. (USA) vol.47 no.5 (1976) p.1859-67 ] E.M. Omel'yanovskii, A.N. Pantyukhov, L.Ya. Pervova, V.I. Fistul', Yu.A. Visil'ev [ Sov. Phys. Semicond. (USA) vol.9 no.10 (1976) p.1267-9 ] D.C. Look [ Solid State Commun. (USA) vol.24 no.12 (1977) p.825-8 ]

[9] [10]

[11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24]

L. Eaves, PJ. Williams [ J. Phys. C (UK) vol. 12 (1979) p.L725-8 ] L. Eaves, T. Englert, T. Instone, C. Uihlein, PJ. Williams, H.C. Wright [ Proc. Conf. Semi-Insulating IH-V Materials, Nottingham, Ed. GJ. Rees (Shiva Publishing, UK, 1980) p. 145-53] L. Eaves, PJ. Williams, C. Uihlein [J. Phys. C (UK) vol.14 no.23 (1981) p.L693-7 ] J. Jimenez, M.A. Gonzalez, J.A. DeSaja, J. Bonnafe [ J. Mater. Sci. (UK) vol.19 no.4 (1984) p. 1207-19] A.L. Lin, E. Omel'yanovskii, R.H. Bube [ J Appl. Phys. (USA) vol.47 no.5 (1976) p. 1852-8 ] N. Tsukada, T. Kikuta, K. Ishida [ Jpn. J. Appl. Phys. 2 (Japan) vol.24 no.5 (1985) p.L302-4 ] B. Pistoulet, P. Girard, G. Hamamdjian [J. Appl. Phys. (USA) vol.56 no.8 (1984) p.2275-83] M. Kaminska, M. Skowronski, W. Kuszko [ Phys. Rev. Lett. (USA) vol.55 no.20 (1985) p.2204 ] M. Taniguchi, T. Ikoma [ J Appl. Phys. (USA) vol.54 no.II (1983) p.6448-51 ] M.O. Manasreh, D.W. Fischer, W.C. Mitchel [ Phys. Status Solidi B (Germany) vol.154 no.l (1989) p . l l ] N. Tsukada, T. Kikuta, K. Ishida [ Jpn. J. Appl. Phys. (Japan) vol.25 no.3 (1986) p.L196-7 ] J. Lagowski, MS Kowronski, H.C. Gatos [ Jpn. J. Appl. Phys. (Japan) vol.25 no.3 (1986) P.L194-5] N. Tsukada, T. Kikuta, K. Ishida [ Semi-Insulating IU-V Materials, Hakone, Japan, Ed H. Kukimoto, S. Miyazawa (Ohmsha, Japan, 1986) p.367 ] J.C. Parker, R. Bray [ Phys. Rev. B (USA) vol.38 no.5 (1988) p.3610-13 ] J. Jiminez, A. Alvarez, M.A. Gonzalez, J. Bonnafe, J.A. de Saja [ Jpn. J. Appl. Phys. I (Japan) vol.27 no.10 (1988) p. 1841-4] J.C. Bourgoin, HJ. von Bardeleben, D. Stievenard [J Appl. Phys. (USA) vol.64 no.9 (1988) p.R65]

8.2

Far-IR photoconductivity spectra of shallow donors in GaAs epilayers RA. Stradling (Updated by M.R. Brozel) September 1996

A

GENERAL REMARKS

In GaAs the shallow donor energies are small and rather well described by hydrogenic theory so that the chemical shifts of the ground-state energies of the individual donor species amount to only a few per cent or about 0.1 meV (~ 1 cm"1). Far-infrared spectroscopy of the transitions between bound states provides a very high-resolution technique for studying chemical shifts and identifying residual contaminants. A number of mechanisms, notably the 'photothermal' effect [1,2], enable bound-to-bound transitions to be detected by photoconductivity, with the result that very high sensitivity can be achieved. Consequently very thin high-purity samples can be easily studied. (Stillman et al [2] also provide an excellent review of early work on photoconductivity from donors in GaAs.) Some pioneering Soviet work, particularly on the line broadening mechanisms, is reviewed in [3]. The residual impurity concentrations are usually in the 1014 to 1015 cm"3 range, even in the best material presently available. The resultant line widths are such that the central cell structure is barely resolved, even in the highest-purity samples. A variety of techniques can be employed to narrow the line including the following: (i)

The application of a magnetic field narrows the lines through the shrinkage of the electronic wavefiinction. This has the additional beneficial effect of increasing the centralcell splittings.

(ii)

Illumination with bandgap radiation generates electrons and holes which are captured and neutralise a proportion of the ionised donor and acceptor sites, thereby reducing the random electric field which broadens the lines [1,4]. However, because donor-acceptor pair recombination is very rapid in direct-gap semiconductors, a significant line narrowing by intrinsic radiation is only found in the highest-purity samples.

(iii)

A reduction in temperature between 4 and 2 K can frequently narrow the lines, even though the total number of ionised sites does not change significantly in this temperature range. Correlation effects arise because electrons in compensated material shun donor sites located near to ionised acceptors at the lowest temperatures [5,6] and the lines are frequently found to narrow on reducing the temperature below 4 K because of the increasingly dipolar character of the electric fields [7]. When this line narrowing is found, a further substantial improvement is often observed on introducing intrinsic radiation.

(iv)

The application of hydrostatic pressure increases the bandgap and hence also the effective mass, donor binding energy and chemical shifts, thereby also improving the resolution. The application of pressure of 8.5 kbar also produces an anticrossing of the deepest lying

of the shallow donors with states associated with higher-order conduction band minima [8,9]. At higher pressures the donors are lost from the experimental spectrum because of the decrease in oscillator strength, thereby simplifying the identification of the remaining donors. In addition, a remarkable line narrowing is sometimes found in the region of crossover. (v)

A line narrowing can be observed when the sample is biased close to electric-field induced breakdown [10]. The increase in concentration of free carriers screens the random electric fields from neighbouring impurities leading to a reduction in linewidth, particularly in compensated samples. However this narrowing is accompanied by a loss in sensitivity of the photoconductive mode of detection because of the increased concentration of free carriers.

Care must be taken in interpreting the far infrared spectrum because all the fine structure does not arise necessarily from different donor species. In addition, the overlap with adjacent, incompletely resolved lines can act to shift the line positions by an appreciable fraction of the observed linewidth. There also exists the possibility of distortion of individual line shapes by such effects as 'channel spectra' (i.e. standing-wave resonances within the sample) and polarisation phenomena produced inadvertently by the various optical components and the sample itself. There can be unresolved substructure due to non-parabolicity effects arising from 'the spin-up and spin-down' impurity states having slightly different transition energies. While permitting very high sensitivity, the use of photoconductivity for detection introduces other possibilities for the distortion of the lines. It has been noted that the line positions observed in photoconductivity and in absorption may differ by up to half a linewidth when tunnelling between adjacent impurity sites rather than the normal photothermal mechanism is responsible for the generation of the photoconductivity signal [H]. It should also be noted that any photoconductive line represents a convolution of absorption energy relaxation and recombination processes. In addition, the use of intrinsic light, while being beneficial in reducing the number of ionised sites and hence the linewidth, introduces a further complication. The current paths will certainly be altered by the presence of photoexcited carriers, thereby increasing the contributions from regions close to the surface. In an inhomogeneous sample this can change the relative heights of individual components. Recently a 'notch effect' has been reported for strong lines where the centre of some lines appears to invert. The mechanism involves the change of dielectric function close to a strong absorption peak. Great care must be taken to avoid misidentifying a 'notched' line as two central-cell components [12]. Important pointers to the validity of interpreting the component peaks of any Is - 2p transitions in terms of an equivalent number of donor species lie within the tracking of the components with measuring frequency or magnetic field and the reproducibility of individual donor lines from sample to sample and of different Is - 2p lines in the same sample. As the chemical shifts are quite small in GaAs, effects arising from central cell structure must be clearly distinguished from those caused by band non-parabolicity and anisotropy. Band nonparabolicity will cause differences in the energies for spin-up and spin-down transitions which can produce a doubling of each donor peak [13,14]. This can be mistaken for central cell structure. Band anisotropy can cause shifts of peak positions when the magnetic field is applied along

different crystal directions. These shifts are typically 50 mT at a transition energy of 20 meV for the <100> compared with the <111> orientation [13,15]. Consequently care must be taken in assigning particular spectral lines to individual donors that the magnetic field had been aligned along the appropriate direction. The so-called 'signature curve' of energy against magnetic field has been studied in detail by Afsar et al [16] for a number of different donors. The shallow donors can exist also in a negatively charged state (D" ion) and extensive structure due to D" ions has been reported in VPE and MOCVD samples by Armistead et al [ 17,18]. Small differences are found which are attributable to differing chemical shifts from the different donor species introduced by the two techniques. The effects of temperature, optical excitation compensation and DC electric field are studied [18,19]. A particular novelty of D- spectra in GaAs is that spectra can be followed in the high magnetic field regime where the dimensionless parameter characterising the field (i.e. hf(c)/2R*) is greater than unity. With high purity samples and strong magnetic fields, the infrared spectrum is particularly rich. There has been much interest recently in identifying the higher order transitions to bound states deep into the conduction band [20-24].

B

LPE GaAs

Liquid phase epitaxy (LPE) can produce material with liquid nitrogen mobilities as good as that of material grown by VPE. Sulphur (X2) is generally the dominant residual donor [25,26]. In contrast to VPE the donor line attributed to silicon (XI), if detected at all, appears at concentrations an order of magnitude lower and donor X3 has not been observed as a contaminant in undoped material. Two other donor lines are observed regularly in high purity material. The first of these has been identified as Sn [25,26], and the second is thought to be Pb [26-28]. In some nominally undoped material the Sn line has been reported to be as strong as that due to S(X2) [29], whereas in other samples the Sn line was always considerably weaker than X2. However, the Pb line varied from being markedly stronger than Sn and comparable with X2 to being an order of magnitude weaker than X2 and even weaker than Xl (Si), although comparable with Sn. C

VPE GaAs

Samples prepared by both the hydride and the chloride VPE processes generally show the same set of donors Xl, X2 and X3, if the material has been prepared without intentional doping. As shown by Skromme et al [30], reducing the Ga/As ratio enhances the intensities of Xl and X3 with respect to X2 which is consistent with the interpretation of Xl as substitutional Si and with X3 as a group IV donor (carbon or germanium). X3 is now generally accepted as arising from substitutional germanium. X2 was initially identified as Si. However, when material with liquid nitrogen mobilities in excess of 105 cm2 /Vs was produced by the MOCVD growth technique, it was found that significant amounts of the X2 donor were not introduced with this technique. Consequently, back-doping experiments could be performed with sulphur which produced a strong X2 peak without substantially broadening or shifting the other components in the central cell spectrum. Skromme et al [30] found that sulphur was always the dominant donor with

hydride material. With the conventional chloride process, Ozekhi et al [31] found that either silicon or sulphur could become the most abundant donor. These workers also studied samples in which N2 was employed as the carrier gas instead OfH2. These samples showed no sign of Xl. Ozekhi et al [31] also suggested that the initial assignment of Xl as associated with a gallium vacancy and X2 as silicon was erroneous, as was subsequently confirmed by the back-doping experiments with MBE and MOCVD material. Skromme et al [30] find the same residual donors (Si and S) and acceptors (Zn and C) in InP samples produced in the same hydride growth system and suggest that common sources for the contaminants are involved. Much of the early donor identification was performed with material grown with the AsCl3:Ga:H2 growth technique and back-doped by this method with Ge, Sn, Pb, Se and Te [32]. The positions of the Sn and Se donors agreed with the results of Stradling et al [25] for LPE material. The chemical shifts derived for the Ge and Pb donors were similar to the results of Cooke et al [28], but differences of the order of 0.02 meV in the reported values probably arise from broadening effects. The order of the chemical shifts and their magnitude at zero magnetic field are believed by Low et al [29] and Armistead et al [33] to be as follows: Pb (0.03 meV); Si (0.05 meV); Se (0.06 meV); Sn (0.07 meV); S (0.1 meV); Te (0.14 meV); Ge (0.18 meV). It should be noted however that the shallowest donor identified in [33] and [34] as lead has also been identified as Te [34]. Extremely sharp lines have been reported by Golubev et al [26] for samples grown by the chloride method. The linewidths for the ls-2p transition were in the range 1-2 jxeV (i.e. approximately 0.01 cm"1). Anomalous spin behaviour was found for the germanium donors. D

MOCVD GaAs

The first reported photoconductive spectrum for MOCVD material was with a relatively impure sample (liquid nitrogen mobility of 50,000 cm2/ Vs) which nevertheless showed central cell structure consisting of two peaks whose amplitudes were of the ratio of 10 to 7 which were well resolved because the two donor species involved happened to be the deepest and shallowest of the shallow donor contaminants frequently found in epitaxial material [28]. The shallower donor reported [28] is thought to be Pb or Te [34]. The deepest and slightly more abundant X3 (in the nomenclature of Stillman [12]) was originally identified as carbon but may be germanium in view of its appearance when neutron transmutation is used to introduce donors (see Section F). Hess et al [35], however, note that X3 is frequently found as the dominant donor in samples where photoluminescence shows only carbon and occasionally Zn acceptors and no significant concentrations of germanium acceptors. They also comment that the absence of germanium acceptors could arise from the high As to Ga ratios employed in growth which could force the amphoteric group IV impurities to take up the Ga site and act as acceptors. By the same argument, however, concentrations of gallium vacancies and arsenic antisite defects are also expected and the possibility that X3 is a complex involving these defects cannot be completely excluded, although germanium is the commonly accepted identification. With a series of samples of different origin X3 was found to be the dominant donor in all cases, with only traces of Pb together with Si, S and Sn [29]. In extremely high mobility layers obtained by fractional distillation of both the AsH3 and trimethylgallium, Hess et al [35] show that donor X3 accounts for more than 90% of the total shallow donors present.

This control of the donor species present in high purity MOCVD enabled Low et al [36] to show that donor X2 was associated with sulphur which is generally only present as a minority donor in MOCVD material. Sulphur doping was accomplished by introducing H2S into the gas stream. It should be noted, however, that the back-doped samples produced very broad photothermal spectra with the central cell structure of the ls-2p" line shifted somewhat to lower energy compared with the normal positions of the X2 and X3 donors and the peaks on the ls-2p+ transition shifted correspondingly to higher energy by the interaction with charged impurities as discussed by Larsen. Low et al [36] also point out that in an earlier report with a VPE sample the S donor lay between the X2 and X3 donors and almost certainly arose because of distortion of the X2 line shape by the 'notch effect', which arises with strong lines and results from a variation of the optical constants in the region of such lines as discussed by Stillman et al [12]. Holmes et al [37] found comparable amounts of X3, X2, Se and Si with a range of MOCVD and MOMBE samples and concluded that the contaminating donors were being introduced from the source gases. E

MBE GaAs

The very low donor background achievable in the present generation of MBE reactors has provided the clue to the identity of the 'XI' donor species [2] almost invariably found in VPE grown material. Previously this donor was thought to be sulphur, but by intentional doping of ptype MBE material, Xl has been positively identified as silicon by Low et al [38]. This group also reports the presence of a background of sulphur at a level of 2 x 1014 cm"3 in layers grown in a Varian Gen II with no other intentional dopant than silicon and an elemental As source. With material grown using an elemental As source and using Sn as the n-type dopant, Si is present, but at a lower concentration than S; whereas if AsH3 is cracked in a SiO2 furnace to provide As, the amplitude of the Si peak increases with increasing furnace temperature. Traces of Pb were found in several samples by Low et al [29,38,39]. Kuchar et al [40] find only the Si intentionally introduced in a high purity sample grown by a Gen II but report sulphur to be dominant in a sample produced by an earlier model of MBE reactor (Varian 360). The latter result was confirmed by Armistead et al [41] who also found sulphur to be the most abundant donor in material produced by a similar reactor without intentional doping. Initial results obtained from a sample grown by MOMBE showed virtually identical spectra to samples grown by MOCVD using the same source material [37], i.e. a high proportion of germanium, silicon and sulphur. In contrast samples grown by MBE with solid sources invariably had only traces of germanium present and no silicon unless the silicon dopant was introduced deliberately. Normally MBE material grown without intentional doping is p-type due to carbon contamination. Occasionally when the CO2 background in the reactor is low, the samples are found to be n-type. Selenium is sometimes found in such samples. The chemical shifts of selenium and sulphur are very similar and consequently selenium cannot be resolved when the samples are deliberately doped with silicon. The quality of material which can be produced by MBE techniques is such that shallow donor spectra can now be obtained with fine structure due to different donor states in thin quantum well material. The donor energies are modified in quantum wells because of:

(i) (ii)

the quantum confinement effect which deepens the donor energies; and the change in symmetry of the wavefimctions imposed by the potential barrier at the interface.

For a donor located at the interface in a single heterojunction the ground state wavefiinction is predominantly 2p rather than 2s because of symmetry considerations. The binding energy is therefore reduced by a factor of approximately four in the absence of a magnetic field. Initial results [42] show three distinct donor lines in a magnetic field. Two of these are thought to be due to donors located at the centre and edges of the well. The origin of the third peak is unclear, but could arise from a chemically-shifted donor. The quality of MBE material has now improved sufficiently that excited states of higher order than 2p are observable in quantum well structures. In multiple quantum well samples where the central 50 A region of the 150 A wells was lightly doped with silicon, eight different magnetooptical transitions were observed and identified [43]. The quantum confinement increased the binding energy by a factor of two.

F

GaAs DOPED BY NEUTRON TRANSMUTATION

Neutron transmutation doping (NTD) of semiconductors provides a convenient method for introducing donor impurities in a controlled manner. The technique involves bombarding the crystal with a flux of thermal neutrons. A small percentage of the lattice atoms capture a neutron and decay into shallow donors. This doping method is well understood in the elemental semiconductors, Si and Ge, and is now applied on a large scale in the semiconductor industry. NTD doping of GaAs by thermal neutrons was first reported by Mirianoshvilli and Nanobashvili [44]. Thermal neutrons are captured by a small proportion of the GaAs lattice nuclei, depending on the relative abundance of the various isotopes of Ga and As and their different capture crosssections. The irradiated atoms decay from an excited state to form a stable donor atom: i.e. a lattice atom is transmuted to a donor impurity. The radioactive decay rate probabilities and relative donor yields are well documented (Vesaghi [45]); 40% of the donors formed are 70Ge, 72 Ge and 60% are 76Se, due to the different thermal neutron capture cross-sections of the respective lattice atoms. After irradiation, a considerable amount of crystalline damage occurs and the samples are found to have very high resistivities. Thermal annealing of the radiation damaged crystal is required in order to activate a high proportion of the transmuted atoms as electrically active shallow donors. FIR photoconductivity from shallow donors in NTD GaAs has been reported by Stoelinga et al [46] and by Afsar et al [47] and also has been studied by Najda et al [48]. All groups agree that the intensity in the region of the X3 donor is considerably enhanced and attribute this to the introduction of germanium. Se is either not observed or appears less strongly than would be expected from the respective cross-sections for neutron capture of Ge and Se. Najda et al [48] used hydrostatic pressure to improve the resolution of the lines and concluded that a significant proportion of the selenium atoms remained in complexes after annealing.

G

PASSIVATION OF Si DONORS IN GaAs BY A HYDROGEN PLASMA

The neutralisation of the donors in GaAs caused by exposure to a hydrogen plasma has been studied by photothermal ionization spectroscopy [49]. Typical conditions for the exposure were a sample temperature of 2500C, time of 30 minutes, a pressure of 750 mtorr and a power density of 0.4 W/cm2 The samples were grown by MBE and back doped with small amounts of silicon. Traces of sulphur and germanium were also present as contaminants. The hydrogenation had a dramatic effect on the electrical properties of the thinnest sample (4 microns), dropping the carrier concentration by a factor of three and increasing the liquid nitrogen mobility by almost a factor of two. Exposure to the plasma had a progressively smaller effect on the other samples as the thickness increased. The linewidth of the ls-2p transition for the silicon donor decreased from 0.58 cm"1 to 0.26 cm"1 although the germanium and sulphur donors were not significantly affected. H

UPDATE FOR THIS EDITION

There has been little published in the period between the previous (2nd) Edition and the current one. The application of hydrostatic pressure GaAs: Ge showed a giant PC effect from a localized donor state with A1 symmetry (in addition to DX effects). The observed shallow-deep donor instability is attributed to the change of the short-range component of the donor potential due to the interactions with acoustic phonons [50]. An improved model of photoconductivity spectra which accounts for the contribution of many excitonic levels and two broadening mechanisms has been presented, together with the variation of free-exciton properties with increasing epitaxial layer thickness [51]. The photoconductivity of lightly doped GaAs epitaxial layers of different thicknesses has been reconsidered using an adoption of the model by DeVore [52]. The model involves only excitonic and interband transitions and the carrier lifetime is assumed to be shortened by recombination at the surface of the sample. The modelling and the comparison with photoluminescence spectra have shown that the free exciton transitions appear as a dip in the PC spectra. The position of the main peak does not correspond to a resonance; it is caused by an enhancement of carrier recombination, due to an increase of the absorption coefficient. By calculating PC line shapes, it was found that there are two thickness ranges for the signature of the free exciton transition: it should appear as a peak in GaAs layers thinner than about 2 fim and as a dip in thicker ones [52]. Far-infrared photoconductivity measurements on n-GaAs using a far-infrared laser and a Fouriertransform spectrometer have been reported, extending previous measurements to higher magnetic fields (14 T) and higher laser energies (330 cm"1). Results in those regions of the spectra where there are resonant polaron effects on transitions to metastable states are presented. To describe the field dependence of the transition energies, a new variational-type approach for the calculation of the energies of the metastable states has been developed. As such an approach yields analytical expressions for the wave functions, the effects of resonant polaron interactions on the transition energies can be calculated. Experimental and theoretical transition energies are compared for the more important transitions to metastable states, including resonant polaron effects, and very good agreement is obtained even though the theoretical model is very simple [53]. The low temperature dynamics of far-infrared (FIR) excited electrons bound to shallow donors

have been studied in n-GaAs. In a magnetic field of 3.6 T electrons are excited from the ground state to the 2p+1 donor state with a short pulse of 118.8 um radiation and the photoconductive response is monitored with sub-nanosecond resolution. The observed response can be described with a simple three level rate equation model, yielding a minimum value of 5 x 109 s"1 for the ionization rate of a shallow donor in the 2p+1 bound state [54]. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] II1] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28]

S.M. Kogan, B.I. Sedunov [ Sov. Phys.-SolidState (USA) vol.8 (1966) p. 1898 ] GE. Stillman, CM. Wolfe, J.O. Dimmock [ Semicond. Semimet. (USA) vol. 12 (Academic Press, New York, 1977) p. 169] L.V. Berman, Sh.M. Kogan [ Sov. Phys.-Semicond. (USA) vol.21 no.9 (1987) p.933-43 ] D.M. Larsen [ Phys. Rev. B (USA) vol.8 (1973) p.535 ] J. Golka [ Phys. Rev. B (USA) vol.8 no.8 (1973) p.3895-901 ] D.M. Larsen [ Phys. Rev. B (USA) vol. 11 no. 10 (1975) p.3904-9 ] J.Golka, J.Trylski, M.S.Skolnick, RA. Stradling, Y.Couder [ Solid State Commun. (USA) vol.22 no.l0(1977)p.623-6] R.A. Stradling [ Festkorperprobleme (Germany) vol.XXV (Friedr. Vieweg, Wiesbaden, 1985) p .591 ] Z. Wasilewski, R.A. Stradling [ Semicond. Sci. Technol. (UK) vol. 1 no.4 (1986) p.264-74 ] OZAlekperov, V.G.Golubev, V.I. Ivanov-Omskii [Sov. Tech. Phys. Lett. (USA) vol.9no.6 (1983) p.319-20] A.C. Carter, GP. Carver, RJ. Nicholas, J.C. Portal, R.A. Stradling [ Solid State Commun. (USA) vol.24 no.l (1977) p.55-60] GE. Stillman, T.S. Low, B. Lee [ Solid State Commun. (USA) vol.53 no.12 (1985) p.1041-8 ] V.G Golubev, V.I. Ivanov-Omskii, LG Minervin, A.V. Osutin, D.G Polyakov [ Sov. Phys.-JETP (USA) vol.61 (1985) p. 1214] V.G. Golubev, V.I. Ivanov-Omskii, A.V. Osutin, D.G. Polyakov [ Sov. Phys.-Semicond. (USA) vol.21 no.l (1987) p.18-21] H. Sigg, JAAJ. Perenboom, P. Pfeffer, W. Zawadzki [ Solid State Commun. (USA) vol.61 no. 11 (1987)p.685-90] M.N. Afsar, KJ. Button, GL. McCoy [ lnst. Phys. Conf. Ser. (UK) no.56 (1980) p.547-55 ] CJ. Armistead, S.P. Najda, RA. Stradling, J.C.Maan [ Solid State Commun. (USA) vol.53 no. 12 (1985) p. 1109-14] S.P. Najda, CJ. Armistead, C. Trager, RA. Stradling [ Semicond. Sci. Technol. (UK) vol.4 no.6 (1989)p.439-54] C. Trager, D.A. Cowan, RA. Stradling [ Physica B&C (Netherlands) vol. 134 (1985) p.250-4 ] CJ. Armistead, RA. Stradling, Z. Wasilewski [ Semicond. Sci. Technol. (UK) vol.4 no.7 (1989) p.557-64 ] RT. Grimes, M.B. Stanaway, J.M. Chamberlain, M. Henini, O.H. Hughes [ Semicond. Sci. Technol. (UK) vol.4 no.7 (1989) p.548-52 ] S.N. Holmes, P.D. Wang, RA. Stradling [Semicond. Sci. Technol. (UK) vol.5 no.2 (1990) p. 1439] A. Labrujere, T.O. Klassen, W.T. Wenekebach, CT. Foxon [ Int. J. InfraredMillim. Waves (USA) vol.9 (1988) p. 1057] H.P.Wagner, W.Pretu" [ Solid State Comnun. (USA) vol.66 no.4 (1988) p.367-70 ] RA. Stradling, L. Eaves, RA. Hoult, N. Miura, P.E. Simmonds [ lnst. Phys. Conf. Ser. (UK) no.l7(1972)p.65] V.G Golubev,Yu.V. Zhilyaev, V.I. Ivanov-Omskii, GR. Markaryan, A.V. Osutin, V.E. Chelnokov [ Sov. Phys.-Semicond. (USA) vol.21 no. 10 (1987) p. 1074-7 ] CM. Wolfe, GE. Stillman, D.M. Korn [ lnst. Phys. Conf. Ser. (UK) no.33b (1977) p. 120-8 ] RA. Cooke, RA. Hoult, RF. Kirkman, RA. Stradling [ J. Phys. D (UK) vol. 11 no.9 (1978) p.945-

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53] T.S. Low, G.E. Stillman, CM. Wolfe [ Inst. Phys. Conf. Ser. (UK) no.63 (1982) p. 143-8 ] B.J. Skromme, T.S. Low, TJ. Roth, G.E. Stillman, J.K. Kennedy, J.K. Abrokwah [ J. Electron. Mater. (USA) vol.12 no.2 (1983) p.433-57 ] M. Ozekhi, K. Kitahara, K. Nakai, A. Shibatomi, K. Dazai [ Jpn. J. Appl. Phys. (Japan) vol. 16 no.9 (1977) p. 1617-22] CM. Wolfe, G.E. Stillman [Appl. Phys. Lett. (USA) vol.27 (1976) p.564-7 ] CJ. Armistead, P. Knowles, S.P. Najda, R.A. Stradling [ J. Phys. C (UK) vol. 17 no.35 (1984) p.6415-34] S.S. Bose, B. Lee M.H. Kim, G.E. Stillman [Appl. Phys. Lett. (USA) vol.51 no. 12 (1987) p.937-9 ] K.L. Hess, P.D. Dapkus, H.M. Manasevit, T.S. Low, B.J. Skromme, G.E. Stillman [ J. Electron. Mater. (USA) vol. 11 no.6 (1982) p. 1115-37 ] T.S. Low, G.E. Stillman, T. Nakanisi, T. Udagawa, CM. Wolfe [Appl. Phys. Lett. (USA) vol.41 no.2 (1982) p. 183-5] S. Holmes et al [ Semicond. Sci. Technol. (UK) vol.4 no.9 (1989) p.782-90 ] T.S. Low,H. Morkoc, AR. Calawa [Appl. Phys. Lett. (USA) vol.40no.7(1982)p.611-13 ] T.S. Low, G.E. Stillman, D.M. Collins, CM. Wolfe, S.Tiwari, L.F. Eastman [Appl. Phys. Lett. (USA) vol.40 no. 12 (1982) p. 1034-6 ] F. Kuchar, R. Meisels, G. Weimann, H. Burkhard [ Appl. Phys. A (Germany) vol.33 no.2 (1984) p.83-5] CJ. Armistead et al [ Lect. Notes Phys. (Germany) vol. 177 (1983) p.289 ] N.C. Jarosik, B.D. McCombe, B.V. Shanabrook, J. Comas, J. Ralston, G. Wicks [ Phys. Rev. Lett. (USA) vol.54 no. 12 (1985) p. 1283-6 ] R.T. Grimes et al [ Semicond Sci. Technol. (UK) vol. 5 (1990) p.305-7 ] S.M. Mirianoshvilli, D.I. Nanobashvili [ Sov. Phys.-Semicond. (USA) vol.4 (1971) p. 1612 ] M.A.Vesaghi [ Phys. Rev. B (USA) vol.25 no.8 (1982) p.5436-50 ] J.H.M. Stoelinga, D.M. Larsen, W. Walukiewicz, R.L. Aggarwal, CO. Bozler [ J. Phys. Chem. Solids (UK) vol.39 no.8 (1978) p.873-7 ] M.N. Afsar, KJ. Button, G.L. McCoy [ Inst. Phys. Conf. Ser. (UK) no.56 (1981) p.547-55 ] S.P. Najda, S. Holmes, R.A. Stradling, F. Kuchar [ Semicond. Sci. Technol. (UK) vol.4 no.9 (1989) p.791-6] N. Panet al [Appl. Phys. Lett. (USA) vol.50 no.25 (1987) p.1832-4 ] C Skierbiszewiski, W. Jantsch, Z. Wilamowski, K. Lubke, T. Suski [ Phys. Rev. B, Condens. Matter (USA) vol.52 no.20 (1995) p.R14312-15 ] Wu Fengmei et al [ Phys. Status Solidi B (Germany) vol. 186 no. 1 (1994) p. 133-41 ] M. Parenteau, Fengmei Wu, A. Jorio, S. Carlone[J. Appl. Phys. (USA) vol.77 no. 10 (1995) p.5185-90] P.W. Barmby et al [ J. Phys., Condens. Matter. (UK) vol.6 no.36 (1994) p.7867-77 ] J. Burghoorn, TO. Klaassen, WTh. Wenckebach [ Semicond. Sci. Technol. (UK) vol.9 no. 1 (1994)p.30-4]

CHAPTER 9 PHOTOLUMINESCENCE SPECTRA OF GALLIUM ARSENIDE 9.1 9.2 9.3 9.4 9.5 9.6 9.7 9.8 9.9 9.10 9.11

Photoluminescence spectra of undoped SI GaAs Photoluminescence spectra of undoped LEC p-type GaAs Photoluminescence spectra of LPE GaAs Photoluminescence spectra of VPE GaAs Photoluminescence spectra of MOVPE (MOCVD) GaAs Photoluminescence spectra of MBE GaAs Photoluminescence spectra of group II atoms in GaAs Photoluminescence spectra of group IV atoms in GaAs Photoluminescence spectra of group VI shallow donors in GaAs Photoluminescence spectra of transition and rare earth metals in GaAs Room temperature photoluminescence mapping of GaAs substrates and epitaxial layers

9.1

Photoluminescence spectra of undoped SI GaAs B. Hamilton (Updated by M.R. Brozel) September 1996

A

INTRODUCTION

Low temperature (usually 4.2 K) luminescence spectra of semi-insulating (SI) liquid encapsulated Czochralski (LEC) grown GaAs are used in characterising this material, which is employed as an integrated circuit substrate. In particular, they are important for determining the nature of the deep intrinsic defect EL2, situated 0.75 eV below the conduction band [I]. The compensation mechanism between EL2 and a shallow acceptor, probably residual C, is thought to be responsible for the SI properties of undoped GaAs. EL2 related photoluminescence (PL) intensity has been found to be greatest in As-rich samples [2-4]. B

0.65 eV PL BAND

There is a well studied PL band at 0.65 eV, also referred to as the 0.635 or 0.63 eV band. Refs [5-7] attribute the band to oxygen on As sites or an oxygen-related complex. Later work [8-15] associates this directly with the EL2 defect. However, Samuelson et al [16] throw into question the direct relationship between this band and EL2. PL spectra of samples ([EL2] - 2 x 1016 cm"3) before and after quenching by persistent light irradiation suggest that EL2 acts only as an intermediate step in the excitation process. PL, infrared absorption and resistivity investigations [14] of a large number of LEC SI GaAs wafers having either 0.65 eV or 0.80 eV as the dominant band revealed that the 0.65 eV dominant crystals have a Fermi level close to the conduction band (CB), while for the 0.80 eV dominant crystals it is further from the CB. C

0.68 eV PL BAND

More recently a PL band at 0.68 eV has been studied [17-21]. Measurements of PL temperature dependence [19] show the 0.68 eV band to be due to a radiative transition between the EL2 centre and the valence band. The presence of both normal and metastable states of EL2 was inferred from the occurrence of the persistent PL quenching (fatigue) effect. Tajima [21] also correlated this peak with EL2. A scanning electron microscope cathodoluminescence study of undoped SI LEC GaAs ingots by Warwick and Brown [22] has revealed the spatial distribution of the 0.68 eV emission to a resolution of 3 microns. The authors conclude from the emission image and other experimental evidence that polygonised dislocation arrays (cell walls) attract or getter, rather than generate, the EL2 defect centres responsible for the emission.

D

0.80 eV PL BAND

The 0.68 eV emission has been accompanied by a broad band at 0.77 eV with peak energy 0.775 eV and a half width of 0.25 eV [20]. Windcheif et al [23] designated this peak at 0.8 eV and observed shoulders between 0.6 and 0.7 eV. It was associated with the AsGa antisite defect by comparison with the EPR signal of the defect. The onset of a band at 1.0 eV corresponds to the second ionisation energy of the Ga^ double donor. Yu et al [20] concluded that possible origins of the emission were deep centres at about 0.45 or 1.07 eV from the CB edge, assigned to a doubly ionised anti-site donor AsGa2+ at Ec - 1.0 eV and a donor at Ec - 0.42 eV respectively. The 0.80 eV PL band has been variously associated with the AsGa antisite defect [23], the VGa defect or its complex [24] and the EL2+ charge state [14]. E

SURFACE DAMAGE

Surface damage by saw cutting, mechanical polishing and scribing has been found to cause a new PL band at 1.4 eV in the low temperature (10 K) spectrum, regardless of growth method, doping or conductivity type [25]. The band results from mechanically induced radiative centres at or near the surface, although deep radiative and nonradiative centres may also be created. The damage extends 10-20 microns into the crystal and anneals out at or above 673 K. The damaged layer has to be removed by chemical polishing before use in devices. The polariton band in GaAs is influenced by surface conditions. A study by Maciaszek et al [26] has shown that the surface quality of SI wafers can be gauged from the intensity of this band. F

NEW PL TECHNIQUE

Tajima [27] has developed a new PL method called twofold excitation modulated PL spectroscopy. A modulation technique is used to analyse variations induced by below-gap as well as above-gap excitation. This is useful for characterising deep levels. Using the technique the author determined that high concentrations of nonradiative deep centres (1 x 1016 cm"3) are present in LEC but not in boat grown crystals. G

ACCEPTORS IN SI WAFERS

The semi-insulating behaviour of undoped (EL2 controlled) material also depends on control of the background acceptor level. Selective pair spectroscopy has been used to investigate the acceptor content of this material, especially carbon and zinc. It was found [26] that carbon acceptor activity dominated the seed end of ingots, whilst the ratio of carbon to zinc was comparable at the tail end. Resistivity was correlated with acceptor activity. Some evidence has been found that carbon free to bound luminescence near to 1.49 eV is quenched in undoped SI material which has been subject to near IR, EL2 bleaching excitation [28]. The effect was attributed to perturbation of the competing non-radiative recombination efficiency.

H

EL2 CHARGE STATE INFLUENCE ON PL SPECTRA

Selective excitation of EL2 photoluminescence has been used to show that the 0.63 eV band and the 0.68 eV band are a measure of the neutral and ionised EL2 density, respectively [29]. It was shown that both bands exhibit a U shaped profile across the diameter of a wafer. The residual donor concentration was found to be an inverted U shape. A recent study of the fine structure of the 0.61 eV (probably the same as the 0.63 eV) line under uniaxial stress and magnetic field has concluded that this band results from the EL2 centre with full Td symmetry [30]. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] [II] [12] [13] [14] [15] [16] [17]

[18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30]

G.M. Martin, J.P. Farges, G. Jacob, J.P. Hallais, G. Poiblaud [ J. Appl. Phys. (USA) vol.51 no.5 (1980) p.2840-52 ] T. Kikuta, K. Terashima, K. Ishida [ Jpn. J. Appl. Phys. (Japan) vol.22 no.8 (1983) p.L541-3 ] D.E. Holmes et al [ IEEE Trans. Electron Devices (USA) vol.ED-29 (1982) p. 1045 ] J. Lagowski et al [ Appl. Phys. Lett. (USA) vol.40 no.4 (1982) p.342-4 ] W.I. Turner, G.D. Pettit, N.G. Ainslie [ J. Appl. Phys. (USA) vol.34 (1963) p.3274 ] B. Deveaud, P.N. Favennec [ Inst. Phys. Con/. Ser. (UK) no.45 (1979) p.492-500 ] P.W. Yu, D.C. Walters [ Appl. Phys. Lett. (USA) vol.41 no.9 (1982) p.863-5 ] A. Mircea-Roussel, A. Makram-Ebeid [ Appl. Phys. Lett. (USA) vol.38 no. 12 (1981) p. 1007 ] P. Leyral, G. Guillot [Proc. 2nd Conf. on Semi-Insulating Hl-VMaterials, Eds S. Makram-Ebeid, B. Tuck (Shiva, Nantwich, England, 1982) p. 166-71 ] P. Leyral, G. Vincent, A. Nouailhat, G. Guillot [ Solid State Commun. (USA) vol.42 no.l (1982) p.67-9] B.V. Shanabrook, P.B. Klein, EM. Swiggard, S.G. Bishop [ J. Appl. Phys. (USA) vol.54 no.l (1983)p.336-40] M. Levinson [ Phys. Rev. B (USA) vol.28 no.6 (1983) p.3660-2 ] M. Tajima [ Jpn. J. Appl. Phys. 2 (Japan) vol.23 no.9 (1984) p.L690-3 ] T. Kikuta, H. Emori, T. Fukuda, K. Ishida [ Extended Abstracts 16th Conf. on Solid State Devices and Materials, Kobe, Japan, 1984 (Business Centre for Acad. Soc. Japan, 1984) p. 173-6 ] B.V. Shanabrook, P.B. Klein, S.G. Bishop [ Physica B&C (Netherlands) vol. 117&118 pt.l (1983) p. 173-5] L. Samuelson, P. Omling, H.G. Grimmeiss [Appl. Phys. Lett. (USA) vol.45 no.5 (1984) p.521 ] M.S. Skolnick, D.A.O. Hope, B. Cockayne [ Proc. Conf. Semi-InsulatingIH-VMaterials, Kahnee-ta, OR, USA, Apr 1984, Eds D.C. Look, J.S. Blakemore (Shiva, Nantwich, England, 1984) p.446-52 ] T. Kikuta, T. Katsumata, T. Obokata, K. Ishida [ Inst. Phys. Conf. Ser. (UK) no. 74 (1985) p.47] P.W. Yu [ Appl. Phys. Lett. (USA) vol.44 no.3 (1984) p.330-2 ] P.W. Yu [ Phys. Rev. B (USA) vol.29 no.4 (1984) p.2283-5 ] M. Tajima [Appl. Phys. Lett. (USA) vol.46 no.5 (1985) p.484-6 ] CA. Warwick, G.T. Brown [ Appl. Phys. Lett. (USA) vol.46 no.6 (1985) p.574-6 ] J. Windcheif et al [Appl. Phys. A (Germany) vol.30 no.l (1983) p.47-9 ] J. Lagowski,HC. Gatos, J.M. Parsey,K. Wada,M. Kaminska, W. Walukiewicz [Appl. Phys. Lett. (USA) vol.40 no.4 (1982) p.342-4 ] V. Swaminathan, M.S. Young, R. Caruso [ J. Appl. Phys. (USA) vol.57 no.4 (1985) p. 1387-90 ] M. Maciaszek et al [ Can. J. Phys. (Canada) vol.67 no.4 (1989) p.384-8 ] M. Tajima [Proc. 13thInt. Conf. on Defects in Semiconductors, Coronado, CA, USA, Aug 1984 (Metallurg. Soc. AIME, 1985) p.997-1003 ] J. Jimenez, A. Alvarez, J. Bonnafe, L.I. Murin [ Phys. Rev. B (USA) vol.39 no. 18 (1989) p. 13310] M. Tajima, T. lino [ Jpn. J. Appl. Phys. 2 (Japan) vol.28 no.5 (1989) p.L841-4 ] MK. Nissen, A. Villemaire, ML. W. Thewalt [ Phys. Rev. Lett. (USA) vol.67 no. 1 (1991) p. 112]

9.2

Photoluminescence spectra of undoped LEC p-type GaAs EMIS Group (Updated by M.R. Brozel) September 1996

p-type material is formed if the melt stoichiometry becomes Ga rich during liquid encapsulated Czochralski (LEC) growth of GaAs [I]. The photoluminescence spectrum at 4.2 K contains two dominant emission bands at about 1.49 eV and 1.441 eV [2] (FIGURE 1). The former consists of two structures at 1.493 and 1.490 eV which are due to free electron to neutral-acceptor and neutral-donor to neutral-acceptor pair transitions respectively. The neutral acceptor is CM. The 1.441 eV band has two phonon replicas due to emission of one and two LO phonons separated by 36 and 72 meV from the ZPL. The band is interpreted as the unresolved neutraldonor neutral-acceptor pair and free-electron to neutral-acceptor transition involving a deep acceptor, the former dominant at temperatures up to about 30 K and the latter dominant above this temperature. Ga As LEC Grown

Relative Emission Intensity

p-Type 1*4.2 K

Relative Emission Intensity

GaAs LEC Grown p-Type T=2K

Low Excitation

FIGURE 1. (a) typical PL spectrum obtained at T = 4.2 K from undoped p-type crystals. The near-intrinsic region emissions (E1n), C^-related emission at ~ 1.49 eV and the 1.441 eV band due to -77 meV deep acceptor are seen, (b) The emission lines at the near intrinsic region. The peak position becomes clear with the change of excitation as shown for the 1.5102 eV line. (Taken from [2].)

A study of the temperature dependence of this emission band between 2 and and 300 K gave an activation energy of 77 ± 2 meV for the acceptor. The acceptor involved has been suggested to arise from the cation antisite double acceptor Ga^ [3]. The first and second ionisation energies for this centre were found to be 77 and 230 meV from the valence band edge. The centre is favoured by a high concentration of As vacancies as a result of Ga-rich growth. Hetzler et al [4] carried out selective excitation photoluminescence and electronic Raman scattering on LEC samples grown from As deficient melts, by which they identified two excited states of the acceptor, which they refer to as having an activation energy of 78 meV. The excitation energies measured are 62.9 and 66.9 meV above the ground states. It was thought that they correspond to transitions from the Is 2 ground state to the split Is 1 ^s 1 excited states of a double acceptor. Bishop et al [5] observed the 78 meV acceptor level in samples grown from stoichiometric melts which suggests that a Ga-rich melt is not a necessary condition for its occurrence. Near the intrinsic region of the PL spectrum weak emission lines are observed at 1.5117, 1.5124, 1.5102 and 1.5071 eV [2], and 1.283 eV [5]. The 1.5124 eV line is assigned as the unresolved emission of an exciton bound to C^. The 1.5071 eV band is always present with the 1.441 eV line, possibly being due to an exciton bound to the Ga^ acceptor. The relative intensities of the 1.443 and 1.284 eV lines of samples from two different suppliers have been found to be identical for a wide range of excitation intensity [3,5]. Elliot [6] observed a strong PL band at about 1.32 eV in a Si-compensated sample from a Ga-rich melt. Bishop et al [5] studied the same band in the same sample and found it to be clearly distinguishable from the 1.283 eV band of Yu et al [3] by virtue of its low temperature peak energy, its strongly diminishing intensity with increasing temperature and its decay characteristics. A more recent study of the 1.32 eV band in conjunction with electron irradiation and infrared absorption measurements does not support the view that the emission results from the Ga"^ to Ga2"^ transition. It is proposed that the acceptor is a [Ga^-A2"] effective mass donor complex whose activation energy is approximately 197 meV [7]. Renolds et al [6] had previously reported six sharp lines between 1.508 and 1.515 eV which were assigned to donor-acceptor complexes associated with an antisite defect. Two of these lines correspond to the 1.5117 and 1.5102 eV emissions. In conclusion, the main emission bands are at 1.49 and 1.44 eV, and p-type conductivity in undoped LEC GaAs is controlled by the cation antisite double acceptor Ga^ along with C^. REFERENCES [1] [2] [3] [4] [5] [6] [7]

D.E. Holmes et al [ Appl. Phys. Lett. (USA) vol.40 (1982) p.46 ] P.W. Yu, D.C. Reynolds [ J. Appl Phys. (USA) vol.53 no.2 (1982) p. 1263-5 ] P.W. Yu et al [Phys. Rev. Lett (USA) vol.41 (1982) p.531 ] S.R. Hetzler et al [Appl. Phys. Lett. (USA) vol.44 no.8 (1984) p.793-5 ] S.G. Bishop, B.V. Shanabrook, WJ. Moore [ J. Appl. Phys. (USA) vol.56 no.6 (1984) p. 1785 ] K.R. Elliott [Appl. Phys. Lett. (USA) vol.42 (1983) p.274 ] P.W. Yu, D.W. Fischer, J.R. Sizdove [Semicond. Sci. Technol (UK) vol.7 no.4 (1992) p.556]

9.3

Photoluminescence spectra of LPE GaAs B. Hamilton (Updated by M.R. Brozel) August 1996

A

INTRODUCTION

The photoluminescence (PL) spectra of liquid phase epitaxy (LPE) layers is an important aid to the detection of impurities and defect levels as well as to the assessment of crystalline quality. B

EXCITON PEAKS

The dominant excitonic emissions from high purity LPE layers (N d + Na £ 1015 cm"3) [3,4] are given in TABLE 1. TABLEl Peak (eV)

Assignment

1.5153

(F, X)n^1

free exciton, n=l state of upper polariton branch

c.1.515

(F, X)

free exciton, lower polariton branch

1.515-1.5146

(D0, X)*

excited states of (D°,X) or excitons bound to complexes involving neutral donors

1.5141

(D0, h) and/or

(D0, XU

exciton bound to neutral donor or valence band to neutral donor

1.5133

(D+, X)

exciton bound to ionized donor

1.5128 1.5124 1.5122

J = 1/2 J = 3/2 J = 5/2

exciton bound to neutral acceptor

1.5097 (weaker)

(D°,X)n=2

two electron transition of (D°,X) with donor left in the n=2 state

Ref [3] shows how the relative intensity of the (D0, X) and (A0, X) lines varies with doping type and compensation.

C

IMPURITY RELATED PEAKS

Silicon and carbon are the main residual impurities in undoped LPE layers [1,5,6]. The following lines measured at low excitation density are commonly present in high purity layers:

TABLE 2 Energy (eV)

Assignment

Reference

1.4931

conduction band to C^

[1,4,7-9]

1.4891

neutral donor to C^

[4]

1.4851

CB to Si^

[4,5,10,11]

1.4814

neutral donor to Si^

[4]

Other impurities (as well as Si and C) were identified by Skromme et al [4] (FIGURE 1) using variable temperature PL and photothermal ionisation spectroscopy. Most samples were grown in a graphite slider-boat. The following PL lines were identified using low excitation intensities of 2.4 mW/cm2 ( CB = conduction band ): TABLE 3 Donor species

Energy (eV)

Assignment

Mg

1.4914

CB to neutral acceptor (7 K)

Mg

1.4875

neutral donor to neutral acceptor (2 K)

Ge

1.4781

CB to neutral acceptor (7 K)

Ge

1.4746

neutral donor to neutral acceptor (2 K)

S is the dominant donor in LPE material, with smaller quantities of Sn, Pb and Si donors [4]. Mn and Cu acceptors at about 1.406 and 1.355 eV respectively [4] are also present with their associated phonon replicas. An exciton bound to a neutral tin acceptor centre is seen in samples containing large amounts of tin. Since the conduction band to acceptor (or free-to-bound (FB)) and donor acceptor pair transitions (DAP) are only accurate to about ±0.5 meV and about ±1 meV respectively, depending on factors such as impurity concentration, 'strain' and, for DAP transitions, excitation power density [2], unambiguous peak identification is only reliable if FB and DAP transitions are clearly distinguished. This is normally accomplished by varying the sample temperature [1,4] and excitation power density [2]. Taniguchi and Ikoma [21] observed an oxygen related peak at 0.64 eV with foil width at half maximum of 120 meV as previously reported [22-24]. Morkoc et al [11] report a 25 K Zn peak at 1.489 eV for samples grown in a high purity Spectrosil boat, though zinc is not normally seen in LPE layers. D

PL DETECTION OF SUBSTRATE/LPE INTERFACE DEGRADATION

A degraded layer (conducting or SI) can occur at the substrate/LPE interface. This is formed at the substrate surface during the 30 minutes or so of heat treatment in a hydrogen gas flow needed prior to moving the substrate into the liquid. This can adversely affect device performance by,

for example, causing a depletion zone to penetrate into an n-type epilayer. The degraded layer can be detected by PL emission arising from the As vacancies which are caused by the evaporation of As atoms. Koschel et al [9] reported such emission at 1.412 eV, with a phonon replica at 1.376 eV, for a measurement temperature of 4.2 K. Undoped n-GaAs epilayers with a carrier concentration of 7 * 1016 cm"3 and a thickness of 0.5 micron were grown on a virgin SI substrate. The line is attributed to a [C^ - V^] complex. The layer and the associated PL were virtually removed by etching with an unsaturated (Ga + GaAs) solution before LPE growth. This is known as the 'Ga etch-back' technique [12,13]. Otsubo et al [14] found a degradation related peak at 1.40 eV in the 77 K PL of variously doped substrates and epilayers. The peak was found to be concentration dependent being strong for n-type material and weak for p-type. For times up to 30 minutes and temperatures of about 700 - 900 0 C, degraded layers are formed at the substrate/LPE layer interface (typically 5 - 1 0 microns thick) and the thickness is found to be proportional to the square root of the heating time. For investigations of this phenomenon see refs [15-18,27]. It has been reported that the point defect content, especially in the interfacial regions is dependent on the dopant species of the substrate, and on the thermal cycling [28]. Tellurium doping of the substrates, produced a much lower defect concentration than tin doping. It was concluded that the out-diflusion of defects from the substrate was related to movement of gallium vacancies, or defects complexed with gallium vacancies.

Band~to-Accep1or Region P =2.4mW/cm2

Intensity

Exciton Region (1.6°K) P L r O.24W/cm2

Wavelength (A) Wavelength (A)

FIGURE 1. (A) PL spectrum of LPE GaAs in the band-to-acceptor region; (B) PL spectrum of LPE GaAs in the exciton region (from [4]).

E

SOLUTION BAKEOUT INVESTIGATED BY PL

Baking of LPE Ga solutions markedly reduces the residual impurities in the undoped epilayers produced [5,17-20]. Garrido et al [5] obtained PL spectra for 8000C baking times of up to 48

hours. These, together with Hall effect measurements, strongly suggest a decrease in epilayer contamination by C, and probably Si, as the heating time increases. Source and ambient materials, rather than the graphite crucible, appear to have been responsible for the C contamination. F

OTHER ASPECTS

PL has been used as a tool for studying LPE layer electronic surface properties [25]. The possibility of using PL changes under high-intensity excitation to assess the quality OfAl2O3 facet coatings on LPE GaAs has been examined [26]. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] II1] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28]

DJ. Ashen, PJ. Dean, D.TJ. Hurle, J.B. Mullin, A.M. White, P.D. Greene [ J. Phys. Chem. Solids (UK) vol.36 no.10 (1975) p.1041-53 ] E.W. Williams, H.B. Bebb [ Semicond. Semimet. (USA) vol.8 (Academic Press, New York, 1972) p.324] U. Heim, P. Hiesinger [ Phys. Status Solidi B (East Germany) vol.66 no.2 (1974) p.461-70 ] BJ. Skromme, T.S. Low, G.E. Stillman [ Inst. Phys. Conf. Ser. (UK) no.65 (1983) p.485-492 ] L. Garrido, J.L. Castano, J. Piqueras [J. Appl. Phys. (USA) vol.56 no.2 (1984) p.569-71 ] J.L. Castano, J. Piqueras [ J. Appl. Phys. (USA) vol.54 no.6 (1983) p.3422-6 ] G.B. Stringfellow, W. Koschel, F. Briones, J. Gladstone, G. Patterson [Appl. Phys. Lett. (USA) vol.39 no.8 (1981) p.581-2] T. Chen, B. Sun [ Chin. Phys. (USA) vol.2 no.3 (1982) p.643-7 ] W.H. Koschel, S.G. Bishop, B.D. McCombe, W.Y. Lum, H.H. Wieder [ Inst. Phys. Conf. Ser. (UK) no.33a (1977) p.98-104 ] D.W. Kisker, H. Tews, W. Rehm [ J. Appl. Phys. (USA) vol.54 no.3 (1983) p. 1332-6 ] H. Morkoc, L.F. Eastman, D. Woodard [ Thin Solid Films (Switzerland) vol.71 no.2 (1980) p.2458] L.M.F. Kaufinann, K. Heime, W.-G. Burchard [ J. Cryst. Growth (Netherlands) vol.34 no.2 (1976) p.289-97 ] P. Nordquist, H. Lessoff, E.M. Swiggard [Mater. Res. Bull. (USA) vol. 11 (1976) p.939 ] M. Otsubo, H. Miki, S. Mitsui [ Jpn. J. Appl. Phys. (Japan) vol. 16 no. 11 (1977) p. 1957-66 ] M. Bugajski, E. Gawronska, J. Petryk, L. Szymanski [Electron Technol. (Poland) vol.10 no.3 (1977)p.61-7] J.S. Harris, Y. Nannichi, G.L. Pearson [ J. Appl. Phys. (USA) vol.40 (1969) p.4575 ] E. Munoz, W.L. Snyder, J.L. Moll [Appl. Phys. Lett. (USA) vol.16 (1970) p.262 ] T. Itoh, M. Takeuchi [ Jpn. J. Appl. Phys. (Japan) vol. 16 no.2 (1977) p.227-32 ] H. Mike, M. Otsubo [ Jpn. J. Appl. Phys. (Japan) vol. 10 (1971) p.509 ] L. Lin, Z. Fang, B. Zhou, S. Zhu, X. Xiang, R. Wu [J. Cryst. Growth (Netherlands) vol.56 no.3 (1982)p.533-40] M. Taniguchi, T. Ikoma [ Inst. Phys. Conf. Ser. (UK) no.65 (1963) p.65-70 ] WJ. Turner, G.D. Petit, N.G. Ainsley [ J. Appl. Phys. (USA) vol.34 (1963) p.3274 ] P.W. Yu [ Solid State Commun. (USA) vol.32 (1979) p. 1111 ] A. Mircea-Roussel, A. Makram-Ebeid [Appl. Phys. Lett. (USA) vol.38 no. 12 (1981) p. 1007-9 ] K. Mettler [ Appl. Phys. (West Germany) vol. 12 no.l (1977) p.75-82 ] V. Swaminathan, W.C. Dautremont-Smith, P.J.Anthony [Mater. Lett. (Netherlands) vol.2 no.3 (1984) p. 179-83] CK. Kim, RM. Malbon, Y.S. Park [ Inst. Phys. Conf Ser. (UK) no.45 (1979) p.305-14 ] A.S. Bruk, A.V. Govorkov, M.G. MilVidisku, E.V. Popova, AA. Shlenskii [ Sov. Phys.-Semicond. (USA) vol.22 no. 10 (1988) p. 1133-5 ]

9.4

Photoluminescence spectra of VPE GaAs B. Hamilton (Updated by M.R. Brozel) September 1996

A

INTRODUCTION

There are two vapour phase epitaxial (VPE) techniques: hydride VPE (starting materials Ga, HCl, AsH3) and chloride VPE (starting materials AsCl3, H2 and Ga). Both methods involve the same reactants at the growth surface, though the hydride process offers more flexibility of the As/Ga ratio. Low temperature (about 4 K) PL provides a way of identifying (radiative) impurities and defects from their characteristic emission lines, as well as providing an idea of crystalline quality and overall impurity content from peak widths. B

EXCITONIC PEAKS

High purity (impurity concentration up to 1015 cm"3) and strain free GaAs grown by VPE exhibits some (or all) of the following lines as observed by Heim and Hiesinger [I]. TABLEl Assignment

Peak e V 1.5153

(F, X)

n=l state of free exciton, upper polariton branch

1.515

(F, X)

free exciton - lower polariton branch

1.5146-1.515

(D0, X)

excited states of exciton bound to neutral donor

1.5141

(D0, X)

exciton bound to neutral donor

1.5133

(D+, X)

exciton bound to ionised donor and/or valence band to neutral donor transition (D°,h)

1.5128

(A°,X)J=1/2[weak]:

1.5124

(A°, X)J=3/2 [dominant]:

1.5122

(A°,X)J=5/2:

1.5108

'two-electron' transitions of a free exciton recombining near a neutral donor

1.5101-1.5014

'two-electron' transitions of (D0, X)*

1.5097

'two-electron transitions of (D0, X), donor left in n=2 state

1.5089

'two-electron' transitions of (D°, C), with donor left in n=3 state

excitons bound to neutral acceptor with differing total angular momentum states [2]

VPE layers showing these transitions are discussed in refs [3-6,17]. The weak 'two-electron' transitions - only present in the purest materials (with Na + Nd £ 1014 cm"3) - have been used to identify different donors (which have unresolved n = 1 states); see ref [7]. The so-called 'twohole' acceptor exciton transitions [3,8-10] involving n = 2 states of neutral acceptor levels are

used for accurate acceptor identification, often to confirm assignments made with the conduction band (free electrons) to neutral acceptor (bound hole) FB transitions - see following sections. Excitons bound to deep oxygen donors [3,11] have also been observed at 1.4885 eV, using high excitation powers, and a weak line attributed to excitons bound to a complex involving copper has also been observed at 1.50275 eV [3,12,14,15]. C

FREE-TO-BOUND TRANSITIONS

(FB)

AND

DONOR-TO-ACCEPTOR

PAIR

(DAP)

At low excitation power densities and approximately 4 K the following transitions can be seen in VPE material (the neutral acceptor involved is given in each case) : TABLE 2 Peak e V

Assignment

Reference

1.4932

FBC^

[3]

1.4915

FBBe 0 ,

[8]

1.4911

FBMg 08

[8]

1.4894

FBZn 0 ,

[3]

1.4892

DAPCA,

[3]

1.4854

DAPZn 0 ,

[3]

1.4850

FBSi^

[8]

1.4848

FBCd 0 ,

[8]

1.4816

DAPSIA,

[3]

1.4782-1.479

FBGeA8

[3,8]

1.4745

DAPGe^

[3]

1.406

FB Mn-related

[3,13]

1.356

FBCu-related

[3,14,15]

1.22

complex of a donor associated with a gallium vacancy

[16]

Phonon replicas of the above peaks, especially the deeper acceptors (Mn5Cu), are often seen [3,16], the LO replica 36.5 meV below the original peak, and the TA replica about 10 meV below the original peak [5]. D

SHALLOW ACCEPTORS

The peak positions for the FB emissions depend somewhat on total concentration of neutral and/or ionised impurities, strain and substrate problems (see [8]) but for fairly high purity samples (Na + Nd £ 1015 cm"3) the positions should be accurate to ±0.05 meV. The DAP peak positions depend implicitly on the donor and acceptor concentrations [16,18] but for high purity samples

should be accurate to within about 1 meV. It can be seen that peaks often overlap and variable temperature and excitation power density measurements have to be made to identify transitions [3,8,16]. In undoped material, the residual impurities (which arise from a number of sources; see refs [3,4,17,19]), their degree of incorporation and for the case of amphoteric group IV impurities, whether they form n- or p-type centres, have been studied as a function of growth conditions [3,4,20]. Both hydride and chloride techniques have C and Zn as dominant acceptor impurities with germanium and silicon sometimes present in trace quantities. Germanium and silicon incorporate predominantly as donors (together with the usually predominant S donor). Selective pair luminescence has been used as a sensitive probe for measuring the background shallow acceptor impurity species in VPE material. Carbon, zinc and silicon acceptors were identified. Excited state information on acceptor levels was also obtained. E

DEEPER ACCEPTORS

Cu and Mn and their complexes have been seen in layers grown by both VPE techniques, and probably arise from out-diffusion from the substrate or from stainless steel in the growth reactor pipework or gas cylinders [4]. As discussed in detail by Williams and Bebb [16] a broad band at about 1.2 eV is often seen in heavily n-doped GaAs materials and is associated with a complex centre involving a gallium vacancy. This band has distinct properties [16] and it has been observed in undoped VPE material grown on heavily n-doped substrates [3]. F

OTHERASPECTS

PL has been used to probe the VPE grown active/buffer layer interface of FET like structures [22] and to observe changes in impurity content of a buffer layer subject to implantation and annealing [23]. Refs [3,4] show how the relative heights of PL peaks due to DAP or FB transitions may be used to estimate relative amounts of different acceptors in the GaAs. For further discussion of the semi-quantitative use of PL see [24]. G

MAGNETOLUMINESCENCE

The question of the chemical identity of donor impurities in GaAs has always been a difficult one, because of the extremely small exciton binding energy characteristic of low electron effective mass systems. Even in high purity epitaxial layers all donors contribute to the single donor bound exciton band. This problem has now been overcome in ultra high quality VPE using the combination of resonant excitation and strong magnetic field [25]. This has the effect of sharpening the excitonic spectra, and splits the s and p states. Chemical identification is made using the two electron satellite components, which result when the exciton decays leaving the donor in an excited state. The Is - 2p splittings show strong central cell effects characteristic of chemical species. The major donors S, Si, Ge and Te were identified. Sn or Se was also observed, but discrimination proved difficult.

At a field of 0.9 tesla, the following Is - 2p splittings were measured: Si S Ge

4.64 meV 4. 73 meV 4.87 meV

REFERENCES [I] [2] [3] [4] [5]

[6] [7] [8] [9] [10] II1] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25]

U. Heim, P. Hiesinger [ Phys. Status Solidi B (Germany) vol.66 no.2 (1974) p.461-70 ] A.M. White, PJ. Dean, B. Day [ J. Phys. C (UK) vol.7 no.7 (1974) p. 1400-11 ] B.J. Skromme, T.S. Low, TJ. Roth, G.E. Stillman, J.K. Kennedy, J.K. Abrokwah [J Electron. Mater. (USA) vol.12 no.2 (1983) p.433-57 ] J.K. Abrokwah, J.N. Peck, RA. Walterson, G.E. Stillman, T.S. Low, B.J. Skromme [ J. Electron. Mater. (USA) vol. 12 no.4 (1983) p.681-99 ] Landolt-Bornstein [ Numerical Data and Functional Relationships in Science and Technology (New Series) vol.7 sub-volume A, Physics of Group FV Elements and III-V Compounds (SpringerVerlag, New York, 1982) p.234 ] LA. Dorrity, D. Hewins, J.D. Medland [ Inst. Phys. Conf. Ser. (UK) no.74 (1985) p.211 ] RJ. Almassy, DC. Reynolds, CW. Litton, K.K. Bajaj, GL. McCoy [ Solid State Commun. (USA) vol.38 no.ll (1981)p.1053-6] DJ. Ashen, PJ.Dean et al [J Phys. Chem. Solids (UK) vol.36 no.10 (1975) p.1041-53 ] P.W. Bonn, R. Bhat, T.D. Harris [ Anal. Chem. (USA) vol.56 (1984) p.58-62 ] A.M.White et al [ J Phys. C (UK) vol.6 no. 11 (1973) p.L243-6 ] A.M. White et al [ Proc. 12th Int. Conf. on Phys. of Semiconductors (1974) p.381 ] F. Williman, D. Bimberg, M. Blatte [ Phys. Rev. B (USA) vol.7 no.6 (1973) p.2473-80 ] W. Schairer, M. Schmidt [ Phys. Rev. B (USA) vol. 10 no.6 (1974) p.2501-6 ] Z.G. Wang, H.P. Gislason, B. Monemar [ J. Appl. Phys. (USA) vol.58 no.l (1985) p.230-9 ] H.P. Gislason, Z.G. Wang, B. Monemar [ J. Appl. Phys. (USA) vol.58 no.l (1985) p.240-7 ] E.W. Williams, H.B. Bebb [ Semicond. Semimet. (USA) vol.8 (Academic Press, 1972) p.321 ] Lin Lan-Ying, Lin Yao-Wang et al [ J. Cryst. Growth (Netherlands) vol.56 no.2 (1982) p.344-9] R. Dingle [ Phys. Rev. (USA) vol. 184 no.3 (15 Aug 1969) p.788 ] CM. Wolfe, G.E. Stillman, E.B. Owens [ J. Electrochem. Soc. (USA) vol. 117 no.l (1971) p. 129130] T.S. Low, BJ. Skromme, G.E. Stillman [ Inst. Phys. Conf. Ser. (UK) no.65 (1983) p.515-22 ] T. Nishino, Y. Fujiwara, Y. Hamakawa [ Inst. Phys. Conf. Ser. (UK) no.65 (1983) p.71-8 ] RJ. Almassy, D C Reynolds, CW. Litten, KK Baja, D.C Look [ J Electron. Mater. (USA) vol.7 no.2 (1978) p.263-77] P.K. Bhattacharya, J.K. Rhee, S.J.T. Owen, JG. Yu,K.K. Smith,RY. Koyama [ J. Appl. Phys. (USA) vol.52 no. 12 (1981) p.7224-31 ] PJ. Dean [ Prog. Cryst. Growth Charact. (UK) vol.5 nos. 1-2 (1982) p.88-174 ] S.S. Bose, B. Lee, M.H. Kim, G.E. Stillman [ Appl. Phys. Lett. (USA) vol.51 no. 12 (1987) p.937]

9.5

Photoluminescence spectra of MOVPE (MOCVD) GaAs B. Hamilton (Updated by M.R. Brozel) August 1996

A

INTRODUCTION

Interest in metalorganic chemical vapour deposition (MOCVD) stems from the need to produce thin layers and abrupt interfaces suitable for use in high electron mobility transistors, superlattices and quantum well structures. Low temperature (less than about 10 K) photoluminescence (PL) spectra are an effective means of detecting impurities and defects in the layers produced. B

EXCITON PEAKS (c. 1.51 eV)

High purity (i.e. Na + Nd £ 1016 cm"3) MOCVD films, like all epitaxial films, give highly structured narrow emission lines arising mainly from free exciton polaritons and excitons which may have become bound to an impurity site. Swaminathan et al [1] obtained the following MOCVD exciton peaks close to the band edge (1.5194 eV) at 12 K (the film was grown at 589°C, 800 rpm, with a VfLII ratio of 8 and a carrier concentration of 1015 cm"3 on a semi-insulating substrate): TABLEl PL line

Probable identification

(eV)

(X = exciton)

1.5156

free

X (upper polariton branch)

1.5149-1.5145

excited states of donor X (and including free X, lower polariton branch)

1.5142

X bound to neutral donor

1.5136

X bound to ionised donor, and donor to VB

1.5126

X bound to neutral acceptor

1.5119

defect X

1.5112

defect X

The authors point to the similarity of the positions and widths of the above lines to all those of epitaxially produced films [2-6]. It should be noted, however, that these positions differ slightly (probably due to strain in the specimens) from the generally accepted values given in ref [3O]. Roth et al [7] report 12 lines in the range 1.5113 to 1.5056 eV for comparable MOCVD samples and suspect that these are closely related to similar lines in MBE films [8,31]. C

CARBON PEAKS (c. 1.49 eV)

Carbon is the dominant background impurity in MOCVD GaAs prepared using TMG

(trimethylgallium, Ga(CH3)3) as the Ga source, but was not revealed in the PL spectra of samples produced using TEG (triethylgallium (Ga(C2H5)3) by Bhat et al [9]. The C probably originates from the adsorption of hydrocarbon radicals onto surface As atoms [10], these radicals being due to fragmentation of the organometallic compounds. The amount of C incorporated depends on growth temperature, gas phase stoichiometry during growth, operating pressure and substrate orientation (see below). C could be incorporated as a donor, CGa, or as an acceptor, CM; it is found [32,35] that C^ predominates. Where the 'X3 donor' is identified as Ge rather than C, carbon is not expected to incorporate as a donor. The following transitions are generally observed at about 4 K using low excitation power densities: 1.4935 (± 0.0005) eV

FBiC^

1.4895 (± 0.001) eV

( D ^ C J ; DAP : C^. Exact position is dependent on concentration and excitation.

See refs [1,3,9,12]. Refs [3,9,35] contain a good discussion of observed variations in these peak positions and how to separate out the often overlapping peaks due to zinc. Kuech et al [10] conclude from PL spectra at 77 K that the contamination level, which is proportional to the ratio of the integrated intensity of the FB(C) peak to the conduction band to valence band peak [3,12,13], is reduced when the growth temperature is lowered e.g. by a factor of 2 when the temperature is lowered from 700 to 6000C for a fixed AsH3/TMG ratio of 40. An increase in the AsH3/TMG ratio decreased the C content, as reported previously [1,13-16]. Layers on the [111] As surface had the highest C concentrations and layers on the [111] Ga surface had the lowest C content. Takagishi and Mori [11,17] have determined from PL that the C concentration falls as the operating pressure is increased over the range 3 x 10"3 to 75 torr. For an AsH3ATMG ratio of 75 there was a conductivity-type conversion at 0.5 torr. At pressures below this, p-type layers resulted from C contamination, while for high pressures n-type layers with low C content were obtained. This 'cross-over' effect from n-type to p-type is characteristic of MOCVD growth and the exact AsH3/TMG ratio at which cross-over occurs depends strongly on source purity, growth temperature and growth rate (small effect if sources are pure) as well as overall operating pressure [14]. Blaauw et al [18] found from PL spectra that C incorporation was enhanced when TMAs (trimethylarsenic, As(CH3)3) was used instead of arsine (AsH3). Mori and Takagishi [19] investigated the influence of extra hydrogen carrier gas in TMG/arsine low pressure MOCVD and found that this increased the C incorporation for samples grown at 1.5 torr. Films grown at constant gas velocity had high resistivity and poor PL features. D

SILICON PEAKS (c. 1.485 eV)

Si is observed as a residual impurity in MOCVD since it is often present in the organometallic

source material AsH3. Si could also come from the SiC coated susceptors or interactions with the growth vessel. Note that Si has been used as a substrate for MOCVD GaAs [20] and as a dopant [21,26]. It can be present as a donor [14] Si03 which is its predominant incorporation site, and/or an acceptor [18,32]. Blaauw et al [18] report a prominent Si peak at 1.4852 eV when using the adduct TMG - TMAs (containing 9 ppm Si according to the manufacturer's analysis) with arsine. The peak which appears to be associated with Si on As sites and which was previously identified by Ashen et al [3], was not present in samples grown simply from TMG and arsine. Dapkus et al [14] found residual Si^ by PL and SiGa by far-IR photoconductivity analysis of high purity MOCVD GaAs with a total impurity content (including C and Zn) of 5 x 1014 cm"3 and an electron mobility of 125,000 cm 2 /Vs at 77 K. Okamoto et al [21] have observed a broad-band peak (1.355 eV) in addition to the band edge peak in the 77 K PL of Si-doped GaAs, which they assign to a self-activated complex of [VGaSi0J: 1.485 eV FB 1.482 eV DAP

involving Si^ - see refs [18,3] involving Si^ - see ref [12] and note that exact position is influenced by concentration and excitation power density.

See also [19] for discussion on probable variations in peak positions and separations of overlapping peaks. Degenerately Si doped MOVPE layers have been studied using PL [36]. At doping levels of less than 1 x io 18 cm"3 DAP luminescence dominated; above this value, free electron to acceptor transitions dominated. Electron concentrations of up to 2 x 1019 cm"3 were studied. Two systematic trends were observed for the highly degenerate samples. Some layers exhibited a significant spectral shift, and consistent Moss Burstein shift, characteristic of heavy doping; others produced only small shifts. The existence of resonant donor levels in the conduction band, causing Fermi level pinning, was invoked to explain these data. E

Zn, Mg, Mn, Ge AND Cu PEAKS

Bhat et al [9] point out that volatile Mg compounds (e.g. dialkylmagnesium) may be present in the organometallic starting materials. ZnAs2 is used as a starting material in the manufacture of high purity arsine and can also be present in TMG. This explains why zinc is, with C, the dominant acceptor in MOCVD GaAs [32,14,35]. There is evidence that lowering the growth temperature enhances Zn incorporation [22,23]. See the discussion following carbon peak positions for variations in the above peak values. Germanium, which is probably a contaminant in arsine sources, is incorporated in MOCVD GaAs predominantly as a donor [32], but under the right growth conditions is present as an acceptor

TABLE 2 Peak energy (eV)

Assignment and references

1.4893

FB involving Zn 0 , [1,3,7,9,12,28,32]

1.4854

DAP involving Zn011 [1,7,9,12,28]

1.4785

FB involving Ge^ [3,12,32]

1.4745

DAP involving Ge^ [ 12,32]

1.4911

FB involving Mg011 [1,3,7,9] and note the differing peak assignments for FB peak from [1]

1.487

DAP involving Mg011 [9]

An Mg peak at 1.4911 eV is suspected by Swaminathan et al [I]. Bhat et al [9] observe PL lines which they suspect are due to Mg or Zn. Both TMG and TEG derived samples were used. Bhat et al [9] point out that volatile Mg compounds (e.g. trialkylmagnesium) may be present in the organometallic starting materials [33]. Blaauw et al [18] and, much earlier, Lee and Anderson [24] report a line at 1.4089 eV due to Mn on a Ga site. Blaauw et al [18] also detected a PL line at 1.3578 eV due to Cu on Ga sites, previously reported in [12,25]. The authors determined that these impurities were due to the diffusion of Mn and Cu from the SI substrate. Recently, PL peaks in MOCVD GaAs doped with oxygen have been reported. In [37], O was introduced using dimethylaluminium methoxide. Peaks at Eg - 72 meV, Eg - 97 meV, Eg - 150 meV, Eg - 410 meV and Eg - 510 meV are interpreted as DAP transitions involving O. In [38], the ethoxide was used. Complex behaviour of emissions at 0.8 eV and 1.1 eV were associated with the high O concentrations (~ 1018 cm"3) introduced. F

OTHERASPECTS

Stress (caused by lattice mismatch) induced changes in the PL intensity and energy for MOCVD growth of GaAs on a Si substrate with GaAs/GaAlAs buffer layers are reported in ref [20]. PL spectra of atomic layer epitaxy (ALE) of 100 cycle MOCVD samples are reported by Bedair et al [27]. A sharp PL peak with a full width at half maximum of 11 meV was observed. Selectively excited PL of the (A, X)n=1 lines for the unambiguous identification of acceptors is reported by Bohn et al [28]. See also refs [3,12]. Non-linearity in the band-to-band recombination PL under high density picosecond photoexcitation was investigated by Lehmen and Ballantyne [29]. A broad luminescent peak at about 1.2 eV (at 2 K) appearing with increasing sulphur doping and attributed to [S^-V 0 J is observed in MOCVD GaAs and thought to be responsible for compensation in n-doped material [34].

REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] [II] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33] [34] [35] [36] [37] [38]

V. Swaminathan, D.L.Van Harejn, J.L. Zilko, P.Y. Lu, N.E. Schumaker [ J. Appl. Phys. (USA) vol.57 no. 12 (1985) p.5349-53 ] D.D. Sell, S.E. Stokowski, R. Dingle, J.V. DiLorenzo [Phys. Rev. B (USA) vol.7 no. 10 (1973) p.4568-86 ] DJ. Ashen et al [ J. Phys. Chem. Solids (UK) vol.36 no. 10 (1975) p. 1041-53 ] R. Dingle, C. Weisbuch, H.L. Stoimer, H. Morkoc, A.Y. Cho [ Appl. Phys. Lett. (USA) vol.40 no.6 (1982)p.507-10] H. Temkin, J.C.M. Hwang [Appl. Phys. Lett. (USA) vol.42 no.2 (1983) p.178-80 ] M. Heilblum, E.E. Mendez, L. Osterling [ J. Appl. Phys. (USA) vol.54 no. 12 (1983) p.6982-8 ] A.P. Roth, R.G. Goodchild, S. Charbonneau, D.F. Williams [ J. Appl. Phys. (USA) vol.54 no.6 (1983)p.3427-30] H. Kunzel, K. Ploog [Appl. Phys. Lett. (USA) vol.37 no.4 (1980) p.416-18 ] R Bhat, P. O'Connor, H. Temkin, R. Dingle, V.G. Keramidas [ Inst. Phys. Conf. Ser. (UK) no.63 (1982)p.l01-6] T.F. Kuech, E. Veuhoff [ J Cryst. Growth (Netherlands) vol.68 no.l (1984) p. 148-56 ] S. Takagishi, H. Mori [ Jpn. J. Appl. Phys. 2 (Japan) vol.22 no. 12 (1983) p.L795-7 ] BJ. Skromme, T.S. Low, TJ. Roth, G.E. Stillman [ J. Electron. Mater. (USA) vol. 12 no.2 (1983) p.433-57 ] K. Mohammed, J.L. Merz, D. Kasemset [Appl. Phys. Lett. (USA) vol.43 no.l (1983) p.103-5 ] P.D. Dapkus, H.M. Manasevit, K.L. Hess, T.S. Low, G.E. Stillman [J. Cryst. Growth (Netherlands) vol.55 no.l (1981) p. 10-23] J.R. Shealy et al [ Inst. Phys. Conf. Ser. (UK) no.65 (1983) p. 109-16 ] M. O. Watanabe, K. Morizuka, M. Mashita, Y. Ashizawa, Y. Zohta [ Jpn. J. Appl. Phys. 2 (Japan) vol.23 no.2 (1984) p.L103-5 ] S. Takagishi, H. Mori [ Jpn. J. Appl. Phys. 2 (Japan) vol.23 no.2 (1984) p.L100-2 ] C. Blaauw, C. Miner, B. Emmerstorfer, AJ. Springthorpe, M. Gallant [ Can. J. Phys. (Canada) vol.63 no.6 (1985) p.664-9] H. Mori, S. Takagishi [ Jpn. J Appl. Phys. 2 (Japan) vol.23 no. 12 (1984) p.L877-9 ] M. Akiyama, Y. Kawarada, K. Kaminishi [ J. Cryst. Growth (Netherlands) vol.68 no.l (1984) p.21-6] K. Okamoto, S. Onozawa, T. Imai [J Appl. Phys. (USA) vol.56 no.10 (1984) p.2993-5 ] SJ. Bass [ J. Cryst. Growth (Netherlands) vol.31 no.l (1975) p. 172-8 ] HM. Manasevit, A.C. Thorsen [ J. Electrochem. Soc. (USA) vol. 119 (1972)p.99] T.C. Lee, W.W. Anderson [ Solid State Commun. (USA) vol.2 (1964) p.265 ] HJ. Queisser, CS. Fuller [ J Appl. Phys. (USA) vol.37 (1966) p.4895 ] E.Veuhoff, T.F. Kuech, B.S. Meyerson [ J Electrochem. Soc. (USA) vol. 132 no.8 (1985) p. 1958] S.M. Bedair, M.A. Tischler, T. Katsuyama, N.A. El-Masry [ Appl. Phys. Lett. (USA) vol.47 no.l (1985) p.51-3] P.W. Bohn, T.D. Harris, R Bhat, H.M. Cox [Appl. Spectrosc. (USA) vol.38 no.3 (1984) p.417-22] A. Von Lehmen, J.M. Ballantyne [ J Appl. Phys. (USA) vol.5.8 no.2 (1985) p.958-62 ] U. Heim, P. Hiesigner [ Phys. Status Solidi B (Germany) vol.66 no.2 (1974) p.461-70 ] A.C. Beye, G. Neu [ J Appl. Phys. (USA) vol.58 no.9 (1985) p.3549-55 ] T.S. Low, BJ. Skromme, G.E. Stillman [ Inst. Phys. Conf. Ser. (UK) no.5 (1983) p.515-22 ] B.E1. Jani et al [Rev. Phys. Appl. (France) vol.19 no.l (1984) p.7-15 ] L. Samuelson, P. Omling, H. Titze, H.G. Grimmeiss [J. Cryst. Growth (Netherlands) vol.55 no.l (1981)p.l64-72] K.L. Hess et al [ J Electron. Mater. (USA) vol. 11 no.6 (1982) p. 1115-37 ] T. Lideikis, G. Treideris [ Semicond. ScL Technol. (UK) vol.4 no. 11 (1989) p.938-42 ] Y. Park, M. Skowronski [ J. Appl. Phys. (USA) vol.75 no.5 (1994) p.2640-3 ] J.M. Ryan, J.W. Huang, T.F. Kuech, K.L. Bray [ J. Appl. Phys. (USA) vol.76 no.2 (1994) p. 1175]

9.6

Photoluminescence spectra of MBE GaAs B. Hamilton (Updated by M.R. Brozel) August 1996

A

INTRODUCTION

Photoluminescence (PL) spectra of molecular beam epitaxy (MBE) grown layers at low temperature (about 4 K) are useful in characterising the layers and determining their quality. The positions of the normally excitonic peaks seen in pure material (up to 1015 cm"3 of impurities) are those originally measured by Heim and Hiesinger [1] and confirmed for MBE material in [8]. TABLEl Peak (eV)

Assignment

1.5153

(F, X)

n-1 state of free exciton, upper polariton branch

1.515

(F 3 X)

lower polariton branch

1.515-1.5146

(D°,X)*

excited states of neutral donor bound exciton

1.5141

(D 0 , X)

neutral donor bound exciton

1.5133

(D + , X)

and (D°, h) ionised donor bound exciton and/or neutral donor to valence band

1.5128 1.5124 1.5122

J = 1/2 J = 3/2 J = 5/2

In addition, up to 60 individual features have been observed [15] in the spectral region 1.50 to 1.511 eV. These are known as KP lines after Kunzel and Ploog who first reported 16 sharp lines in this region for samples grown at a substrate temperature of 5300C [3]. The dominant impurity incorporated as an acceptor is C^, from CO contamination of the growth chamber, and from the substrate. This gives rise to peaks measured at low excitations [8,36]: 1.4932 (2 K) 1.4892 (2 K)

free

electron to bound hole at CM acceptor - FB(C) donor bound electron to acceptor bound hole at CM- DAP(C)

Note, these peak positions are subject to fluctuations of about 0.5 meV (see [33]) and the DAP peak positions depend strongly on excitation power density and doping [34]. If doping is such that FB and DAP peaks overlap and are of approximately the same intensity then their positions are shifted toward one another i.e. the energy difference between them decreases. The issue of residual impurities becomes more important as MBE technology strives for higher quality layers. A study of residual impurities, using many techniques, including photoluminescence, has been made by Stanway et al [36]. The implications of varying source geometry were investigated and improved source configurations suggested. The principal residual

impurity was found to be carbon; concentrations as low as 8.7 x 1012 cm"3 were measured with donor impurities silicon, sulphur and selenium also observed. Kerr et al [37] have investigated the variation of the background carbon level obtained with solid arsenic sources from different suppliers. A carbon concentration ranging from 1 x 10 14 to2 x 1016 cm"3 was obtained. A broad PL band is often seen in undoped or lightly doped MBE GaAs growing peaks between 1.47 and 1.49 eV [10,11,4] including a peak at 1.479 eV. The KP lines have been the subject of much study. Recent papers [33,34] indicate that they can be split into two groups: the 'g peak' at 1.5111 eV (highest energy emission, shorter lifetime (~ 1.2 ns), unpolarised) which is present under the widest range of growth conditions [13,34] and dopants (n- and p-type), and the emissions h - v, present mainly in p-type material (longer lifetimes (~ 2.4 ns) and indicating definite polarisation properties). The origin of this emission is still uncertain but the latest suggestions are that the cg peak' is due to an exciton bound to a neutral 'acceptor complex defect' - involving the major residual carbon acceptor and an isoelectronic defect [33,34]. The 'h - v' emissions are due to exciton recombinations at axially oriented complex defects acting as isoelectronic centres [34]. These isoelectronic impurities may arise from the fact that indium (used to mount substrate to molybdenum block) and boron nitride for the crucible containing the Ga, As and dopant species are present in the growth system. These latest results cast considerable doubt on earlier theories [12,15]. In general KP lines are not observed in non-MBE samples, although lines in the same spectral region emitted by MOCVD layers [20,21] may be related. Neither are they always emitted by MBE samples; they are often absent when As2 molecular beams are used [3], or when substrate temperatures above 600 0 C are used [3]. In attempting to relate the KP lines to the electrical characteristics Rao et al [10] performed low temperature PL measurements on various high quality MBE layers (undoped n- or p-type, Sndoped n-type, and Be-doped p-type) and found that for most types of layer the 1.5109 and 1.5046 eV lines were prominent, but much more so for p-type than for n-type samples. Recently, more complex forms of spectroscopy including magneto-optical and resonant excitation techniques have been applied to the KP lines [2,38,39]. In particular, analyses of two hole spectra have been made. These measurements greatly strengthen the assignment of the KP lines to acceptor species, with exciton localisation energies following Hund's rule. Carbon implantation into ultra-pure MBE GaAs has demonstrated that C related g lines exist and it is suggested that acceptor pair carbon-vacancy centres are responsible, i.e. (C^-V^) [50]. B

1.47 - 1.49 eV REGION

The luminescence in the 1.47 and 1.49 eV region first measured by Briones and Collins has also generated considerable research activity. In their original study [4] the authors show a correlation of the intensity and energy of the luminescence at 1.491 and 1.473 eV with that of the exciton-like KP lines at 1.5109 and 1.5046 eV respectively. These are attributed to complex centres of the type C-V^ and C-VGa respectively. Koschel et al [23] suggested that a simple vacancy, or

carbon-related defect centre differing from a simple carbon acceptor, may be responsible, while Jeong et al [24] proposed interstitial carbon impurities. Akimoto et al [5] found that intensity of KP lines increased as the amount of CO in the background increased and As4/Ga beam flux ratio decreased. They also carried out Fourier transform infrared spectroscopy and inferred the existence of C-O stretching vibrations of polynuclear carbonyls. They suggest that the KP luminescence is due to excitons bound to carbonyls incorporated on As sites. From the correlation between the KP lines and those in the 1.47 - 1.49 eV region they suggest that luminescence in the latter may be due to a C impurity whose acceptor level is shifted by oxygen. Contour et al [7] also adopted a correlation approach. They performed PL excitation measurements and correlated the 1.5109 eV line, which they attribute to a defect-bound-exciton, with emission at 1.495 eV, thought to result from a free-to-bound acceptor transition. Rao et al [10] also looked for correlations between the KP lines and those in the 1.47 - 1.49 eV region. They found exceptions to Hayne's rule [25] upon which the correlations of Briones and Collins [4] and others are founded. Lines at 1.468, 1.475, 1.48 and 1.496 eV were detected. Some of the lines in this region have now been unambiguously associated with the acceptor-like defects which give rise to the excitonic KP transitions [39]. These lines in fact result from ground state splitting of these acceptors. Returning to the near bandgap part of the PL spectrum there is a doublet peak at about 1.516 eV. Hopfield [26] and Pekar [27] suggested that exciton-polaritons were responsible (coupling of photons and free excitons). Recently, Koteles et al [28] clearly demonstrated that enhanced scattering of exciton-polaritons at the exciton energy is the cause of the exciton-polariton lineshape observed for pure undoped MBE samples. Bloss et al [29] found that the dip between the two peaks increased as the thickness of the undoped layers, controlled by etching, increased since for a given density of scattering centres the distance an exciton-polariton travels before reaching the sample surface is governed by the layer thickness, and for thicker samples this is greater than the exciton-polariton scattering length. C

UP CONVERSION

Up-conversion or Anti-Stokes emission is observed in MBE and other high purity epitaxial layers [30]. Near band gap emission is generated by below bandgap excitation. The authors explain this as a two-step optical excitation of an electron-hole pair via a deep centre: Step 1 Step 2

electron excited from VB to deep centre electron transfers from deep centre to CB

An electron is thus created in the CB and a hole in the VB while the deep centre remains unchanged. This mechanism depends on native defects, not on intentionally introduced impurities, and could be a useful means of studying deep centres. A PL band of peak energy equal to (Eg+A) has been reported but although its intensity seems to correlate with the quality of the MBE layer, no definitive explanation has been given [42].

D

SILICON PEAKS

Silicon incorporates primarily as a donor, SiGa [35], but under the right growth conditions can occupy As sites giving rise to peaks at - 4.2 K [11] 1.485 eV 1.4813 eV

FB DAP

[ S i J - ref [33] [ S i J - ref [36]

These peaks are subject to the same variations as described for FB and DAP peaks due to C. Erickson et al [22] investigate the effect of substrate temperature and As pressure on Si- and Cacceptor peaks at 1.483 eV (SiGa°, Si^ 0 ) and 1.490 eV (D0, CM) to show the nature of Si incorporation in MBE GaAs grown using As2 beams. The PL spectra together with carrier concentration measurements show that the Si dopant exhibits an enhanced acceptor character as the substrate temperature is increased from 570 to 6800C at a given As pressure, while for a fixed substrate temperature of 6800C the PL peaks suggest that there may be an optimum As2 pressure for minimising the incorporation of non-radiative defects and Si acceptors in films grown at this substrate temperature. Mendez et al [31], using an As4 beam, studied the incorporation of Si as a dopant in GaAs as a function of substrate temperature and As/Ga flux ratio. They too found that the acceptor character of Si increased with the substrate temperature. Also, the Si selfcompensation decreased as the flux ratio was raised from 2 to 6. Metze et al [32] investigated MBE: Si films grown at low substrate temperature and growth rates. PL spectra and electrical measurements showed a big improvement in quality over films grown at normal temperatures and rates (5700C, 1 micron/hr). In particular, the PL peaks listed in TABLE 2 were observed. TABLE 2 Substrate temperature 0 C

Growth rate microns/hr

PL peaks observed 1.514 eV

1.495 eV

570

1

Yes

Yes

450

1.1

No

Yes

450

0.2

Yes

Yes

380

0.02

Yes

Yes

The 1.514 eV peak was assigned to bound exciton recombination and the 1.495 eV to carbon. Heavily Si doped MBE GaAs has been studied [51] and found to yield three PL lines at 1250 nm, 1200 nm and at very high Si concentrations (> 1019 cm"3), 1000 nm [50]. X-ray spectroscopy using quasi forbidden transitions suggests that SiGa complexes may play a role in generating these bands, and that at the highest Si concentrations, Si clustering is involved. In GaAs grown using cracked AsH3 as the As source, rather than the normal solid As, Mg/Be

peaks are dominant: 1.4917 eV FB due to MgGa or BeGa 1.4880 eV DAP due to MgGa or BeGa E

measured at 6.2 K and low excitation power

BORON PEAKS

There has been one report of B doping of MBE GaAs. The PL spectrum of GaAs doped only with B exhibits a single peak (B1) at 1.443 eV. However, co-doping with Si to lift the Fermienergy reveals a further peak (B2) at 1.328 eV while moving the B 1 peak to 1.446 eV. In addition, the improved luminescence efficiency that also results by co-doping reveals one and two phonon replicas OfB1 and a single phonon replica OfB2. B 1 and B 2 are acceptor peaks where analysis yields activation energies of 71 - 72 meV (B1) and 188 meV (B2). B 1 is assigned to B^07" while B 2 may be due to [B^ - SiGa]. The authors note that these energies are distinctly different from the 77 meV and 203 meV levels seen in Ga-rich LEC GaAs [40]. F

HYDROGENATION OF MBE GaAs

Impurity passivation by hydrogenation is a field of growing interest. Representative data and references can be found in the paper by Pavesi et al [41], who studied the effect in both MBE GaAs and AlGaAs. A redistribution of luminescence strength from the donor bound exciton into the acceptor bound exciton is typically observed in the binary, due to donor passivation. A more complex behaviour was associated with acceptor passivation [42] of Si doped p-type GaAs; p-type conductivity resulted from the growth conditions used. It was observed that acceptors could be re-activated at low temperatures by excitation with above bandgap light. The effect disappeared at higher temperatures. G

INDIUM DOPED MBE GaAs

In doping can have profound effects on the electrical behaviour of GaAs wafers. The fundamental processes involved have been studied by indium doping of MBE GaAs layers, over a wide range of indium concentrations. The photoluminescence signal was found to be a sensitive test of the influence of indium on the quality of the layers. It was concluded that In concentrations of around 1000 ppm produced optimum quality layers [43]. H

LOW TEMPERATURE MBE (LTMBE) GaAs

As-grown LTMBE GaAs exhibits extremely low PL efficiency, a result of the very small minority carrier lifetime. However, there have been reports of PL on this material. PL from hot carriers in material grown at 200 0 C has been reported. Band-to-band radiative recombination depends strongly on equilibrium carrier concentrations and hence the electron and hole quasi Fermi energies as the material is excited. Results are interpreted in terms of ultrafast carrier recombination, intervalley transfer and thermalization [44]. An intrinsic feature resulting in sharp-line spectra with a no-phonon line at 1.467 eV has been interpreted as being due to a defect of C3V symmetry, possibly to [VGa - AsJ pairs [45]. The same group later reported annealing effects on this PL. Three no-phonon lines seen in as-grown

material are ascribed to C3V defects of the following types: [As1 - VGa], [As1 - CM] and [As1 - AsGa]. Heat treatments between 3000C and 5000C increase the concentration OfAs1. This is consistent with increases in the PL lines associated with [As1 - VM] and [As1 - AsGa], [46]. Deep centre luminescence in LTMBE GaAs grown between 200 and 300 0 C was found to depend on the growth temperature and the As4/Ga ratio. A dominant emission at 0.68 eV was ascribed to recombination via EL2. A second emission line at 1.1 eV is related to VGa. Layers grown between 325 and 4000C show a single transition at 0.75 - 0.81 eV (called the 0.8 eV line). This is attributed to the [As1 - VGa] complex. This centre exhibits a Franck-Condon shift of 0.34 eV and a thermal ionization energy of 0.36 eV and may be assigned to the EL6 electronic level [47]. The PL excitation spectra at 0.8 eV due to [As1 - V 0 J was investigated in material grown between 325 and 4000C. Oscillating structure that was observed in the spectra was interpreted as resulting from the splitting of the heavy- and light-hole valence bands, [48]. The 0.68 eV line associated in [47] with EL2 has been discussed in more detail [49]. Like the EL2 centre the 0.68 eV luminescence can be photoquenched. It can also be photo-enhanced. However, not all the 'EL2 centres' can be photoquenched in LTMBE GaAs. The concentration that is photo-quenchable decreases with increasing growth temperature. The authors interpret this in terms of a model where As03 centres exist under relaxed or strained conditions. The stress present in low-temperature layers is responsible for the unrelaxed condition and consequently for the lack of photo-quenchability. The photo-enhancement is attributed to the hole photoionization of AsGa+ and to a high concentration of the compensating VGa [49]. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] [II] [12] [13] [14] [15]

U. Heim, P. Hiesinger [ Phys. Status Solidi B (Germany) vol.66 (1974) p.461-70 ] M.S. Skolnick, T.D. Harris, CW. Tu, T.M. Brennan, M.D. Sturge [Appl Phys. Lett. (USA) vol.46 no.4(1985)p.427-9] H. Kunzel, K. Ploog [Inst Phys. Conf. Ser. (UK) no.56 (1981) p.519-28 ] F. Briones, D.M. Collins [ J Electron. Mater. (USA) vol. 11 no.4 (1982) p.847-66 ] K. Akimoto, M. Dohsen, M. Arai, N. Watanabe [Appl. Phys. Lett. (USA) vol.45 no.9 (1984) p.922-4 ] G.B. Scott, G. Duggan, P. Dawson, G. Weimann [J Appl. Phys. (USA) vol.52 no. 11 (1981) p.6888] J.P. Contour, G. Neu, M. Leroux, C. Chaix, B. Levesque, P. Etienne [ J. Vac. Sci. Technol B (USA) vol.1 no.3 (1983) p.811-15 ] M. Heilblum, E.E. Mendez, L. Osterling [ J. Appl. Phys. (USA) vol.54 no. 12 (1983) p.6982-8 ] T. Temkin, J.C.H. Hwang [ Appl. Phys. Lett. (USA) vol.42 no.2 (1983) p. 178-80 ] E.V.K. Rao, F. Alexandre, J.M. Masson, M. Allovon, L. Goldstein [ J Appl. Phys. (USA) vol.57 no.2 (1985) p.503-8] BJ. Skromme, G.E. Stillman, A.R. Calawa, G.M. Metze [ Appl. Phys. Lett. (USA) vol.44 no.2 (1984)p.240-2] D.P. Halliday, L. Eaves, P. Dawson [ Proc 13th Int. Conf. on Defects in Semiconductors, Coronado, USA, Aug 1984 (Metallurgical Soc. of AIME, 1985) p. 1005-11 ] PJ. Dobson, GB. Scott, J.H. Neave, B.A. Joyce [ Solid State Commun. (USA) vol.43 no. 12 (1982) p.917-19] P.K. Bhattacharya, H.-J. Buhlmann, J. Ilegems, J.I. Staehli [J. Appl. Phys. (USA) vol.53 no.9 (1982)p.6391-8] D.C.Reynolds et al [ Solid State Commun . (USA) vol.52 no.7 (1984) p.685-8 ]

[16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33] [34] [35] [36] [37] [38] [39] [40] [41] [42] [43] [44] [45] [46] [47] [48] [49] [50] [51]

W.I. Wang [Appl. Phys. Lett. (USA) vol.44 no.12 (1984) p.1149-51] CJ. Hwang [ Phys. Rev. B (USA) vol.8 no.2 (1973) p.646-52 ] G.C. Osbourne, D.L. Smith [ Phys. Rev. B. (USA) vol.20 no.4 (1979) p. 1556-61 ] L. Eaves, D.P. Halliday [ J. Phys. C (UK) vol. 17 no.27 (1984) p.L705-9 ] A.P. Roth, S. Charbonneau, R.G. Goodchild [ J. Appl. Phys. (USA) vol.54 no.9 (1983) p.5350-8] A.P. Roth, R.G. Goodchild, S. Charbonneau, D.F. Williams [ J. Appl. Phys. (USA) vol.54 no.6 (1983)p.3427-30] L.P. Erickson et al [ J. Appl. Phys. (USA) vol.56 no.8 (1984) p.2231-5 ] W.H. Koschel, RS. Smith, P. Hiesinger [ J. Electrochem. Soc. (USA) vol.128 no.6 (1981) p.1336] M. Jeong, J. Shirafuji, Y. Inuishi [Jpn. J. Appl. Phys. (Japan) vol.20 no.4 (1981) p.795-6 ] J.R. Haynes [ Phys. Rev. Lett. (USA) vol.4 (1961) p.361] J.J.Hopfield [ Phys. Rev. (USA) vol. 112 (1958) p. 1555 ] S.I. Pekar [ Sov. Phys. JETP (USA) vol.6 (1985) p.685 ] E. S. Koteles, J.P. Salerno, W. Bloss, E.M. Brody [ 17th Int. Conf. on the Physics of Semiconductors, San Francisco, USA, 6-10 Aug 1984 (Springer-Verlag, New York, 1985) ] W.L. Bloss, E.S. Koteles, E.M. Brody, BJ. Sowell, J.P. Salerno, J.V. Gormley [ Solid State Commun. (USA) vol.54 no.l (1985) p. 103-6 ] L.G. Quagliano, H. Nather [ Appl. Phys. Lett. (USA) vol.45 no.5 (1984) p.555-7 ] E.E. Mendez, M. Heiblum, R. Fischer, J. Klem, R.E. Thorn, H. Morkoc [ J. Appl. Phys. (USA) vol.54 no.7 (1983) p.4202-4] G. Metze, A.R Calawa, J.G. Mavroides [ J. Vac. Sci. Technol. B (USA) vol. 1 no.2 (1983) p. 166] DJ. Ashen, PJ. Dean etal [J. Phys. Chem. Solids (UK) vol.36 no. 10 (1975) p. 1041-53 ] R. Dingle [ Phys. Rev. (USA) vol. 184 no.3 (1969) p.788 ] T.S. Low, G.E. Stillman, D.M. Collins, CM. Wolfe, S. Tiwari, L.F. Eastman [Appl. Phys. Lett. (USA) vol.40 no. 12 (1982) p. 1034-6 ] M.B. Stanway et al [ Inst. Phys. Conf. Ser. (UK) no.95 (1989) p.295-300 ] T.M. Kerr, C.E.C. Wood, S.M. Newstead, J.D. Wilcox [ J. Appl. Phys. (USA) vol.65 no . 7 (1989) p.2673 ] M.S. Skolnick et al [ Proc. 18th Int. Conf. Physics of Semicond., Ed O. Engstrom (World Scientific, Singapore, 1987) p. 1389 ] S. Charbonneau, W.G. McMullan, M.O.Henry, M.L. W. Thewalt [Mater. Res. Soc. Symp. Proc. (USA) vol.104 (1988) p.549-54 ] S.K. Brierley, H.T. Hendriks, W.E. Hoke, PJ. Lemonias, DG. Weir [Appl. Phys. Lett. (USA) vol.63 no.6 (1993) p.812-14 ] L. Pavesi, D. Martin, F.K. Reinhart [ Appl. Phys. Lett. (USA) vol.55 no.5 (1989) p.475-7 ] I. Szafranek, S.S. Bose, G.E. Stillman [Phys. Appl. Phys. Lett. (USA) vol.55 no. 12 (1989) p. 12057] Yu.Yu. Bacherikov [ Status Solidi A (Germany) vol. 140 no.2 (1993) p.559-66 ] H.M. van Driel, X.-Q. Zhou, W.W. Ruhle, J. Kuhl, K. Ploog [Appl. Phys. Lett. (USA) vol.60 no.l8(1992)p.2246-8] P.W. Yu, D.C. Reynolds, CE. Stutz [Appl. Phys. Lett. (USA) vol.61 no.12 (1992) p.1432-4 ] P.W. Yu, D.N. Talwar, CE. Stutz [ Appl. Phys. Lett. (USA) vol.62 no.21 (1993) p.2608-10 ] P.W. Yu, GD. Robinson, J.R. Sizelove, CE. Stutz [ Phys. Rev. B, condens. Matter (USA) vol.49 no.7 (1994) p.4689] P.W. Yu, CE. Stutz [ Solid State Commun. (USA) vol.89 no.3 (1994) p.293-6 ] P.W. Yu, M.A. Capano, A.T. D'Agostino, CE. Stutz [ Phys. Rev. B, Cond. Matter. (USA) vol.49 no.23 (1994) p. 16398-402] I. Fujimoto,N. Kamata,H. Katahama,Y. Shakuda,K. Kobayashi,T. Suzuki [Inst. Phys. Conf. Ser. (UK) vol.91 (1988) p.247 ] S. Shigetomi, Y. Makita, A.C. Beye, A. Yamada,N. Ohnishi, T. Matumori[7. Appl. Phys. (USA) vol.63 no.3 (1991) p. 1613 ]

9.7

Photoluminescence spectra of group II atoms in GaAs B. Hamilton (Updated by M.R. Brozel) September 1996

A

SPECTROSCOPIC FEATURES

Shallow impurity photoluminescence (PL) diagnostics is normally carried out at cryogenic temperatures (below 10 K) when particles are thermally stabilised. Three principal transitions are commonly studied: bound excitonic decay [BE]5 free-to-bound [FB] and donor-acceptor pair [DAP] transitions [I]. Excitonic decay at neutral donors or acceptors provides the sharpest signatures and is favoured for impurity identification especially in low-doped systems. In particular, two hole replica BE emission [2] has been used to chemically identify group II dopants. FB transitions are also extremely useful for dopant identification [3]. B

CHEMICAL IDENTIFICATION

Group II impurities form acceptor states by substituting for Ga on the group III sublattice. The main transitions used for dopant identification are listed in the table. Enhanced characterisation of acceptor states Zn (C and Ge) from excited state analysis are reviewed by Reynolds et al [6]. Several comprehensive studies of many shallow impurity transitions including group II elements exist [3-5,7] (FIGURE 1). The upper concentration limit for precise chemical identification of the group II elements is limited to about 1016 cm'3 by impurity screening. High resolution data for group II elements, including acceptor excited state energies which can be obtained from two hole replica spectra for excitonic decay at neutral acceptors, and exciton lifetime can be found in [8]. C

GROUP H DOPING

The commonest group II element used as a p-type dopant in LEC, LPE, VPE and MOCVD growth technologies is Zn. Zinc doping by diffusion is also of major importance. The luminescence characteristics of Zn substituted on a Ga site are well understood. Cl

Zinc

A rather broad emission band has been observed in Zn doped material (especially at relatively high Zn concentration) close to 1.4 eV [9]. This band has been associated with the zinc-arsenic vacancy [Zn03 - V^ nearest neighbour pair. The defect has been investigated using high energy (2.2 MeV) electron irradiation and annealing in order to probe defect reactions associated with this Zn related band [1O]. Results indicate that V^ production leads to enhancement of the band (measured at 77 K with hv = 1.37 eV) through pairing with substitutional Zn. Anneal temperatures of 2000C (30 min) optimised the defect production as monitored by the PL intensity. At higher temperatures, mobile, radiation-induced defect interactions and thermal dissociation of

the [Zn03 - VM] pairs quench the luminescence. The energy gap of GaAs can be made indirect by the application of hydrostatic pressure; the F 1 P 8 gap converting to X' - F 8 at around 40 kbar. All impurity related PL transitions reduce in efficiency at around this pressure [H]. The effect of hydrostatic pressure on heavily Zn doped GaAs PL spectra [12] showed the direct-indirect transition but with luminescence persisting up to 96 kbar in very heavily Zn doped (p = 9 x 1019 cm"3) material, due to the high free hole concentration [13]. At this pressure, the broad emission band peaked at approximately 2.4 eV at 120 K. The pressure coefficient for indirect gap (X' - F8) was found to be x = - (1.8 ± 0.6) x 10"3 eV/kbar. Indirect luminescence was also observed. Uniaxial stress has been shown to produce a short wavelength peak on the broad near edge emission measured in heavily Zn-doped GaAs (p above 5 x 1018 cm"3). Polarisation studies [14] have demonstrated that this is due to valence band splitting. Zinc related hot luminescence has been observed [15,16] with an emission threshold at 1.800 ± 0.3 eV. The observed emission was explained using hot hole and hot electron models. A theoretical study of such emission has recently been performed [17]. Zinc has been used as a dopant in MBE growth, the low sticking coefficient associated with neutral Zn species being partially overcome by using a Zn+ ion source. Photoluminescence spectra indicate that p-type behaviour comparable with LPE can be achieved [18]. On deuteration the spectrum reveals a broad band between 1.30 and 1.45 eV possibly associated with [Zn03 - D] complexes [19].

FIGURE 1. The evolution of the free-to-bound band FB from the donor acceptor pair band P with increasing temperature in VPE GaAs, which contains only one significant residual shallow acceptor Zn (taken from [7]).

C2

Beryllium

Beryllium doping is commonly used in MBE growth of GaAs. The increase of PL intensity within

the Be emission bands has been observed to be sublinear with hole concentration [20]. The reason for this may be that a fraction of the incorporated Be atoms (possibly interstitial) play a role in forming non-radiative shunt paths [21]. Hot luminescence has been observed in Be doping MBE material [22]. A hot electron mechanism was used to explain the spectral emission and to estimate the L-F splitting of the conduction band to be 320 ± 4 meV. C3

Calcium

Photoluminescence of Ca implanted GaAs has been reported [23] to produce several lines, including exciton decay at neutral Ca acceptors (Ca0, X), electron to Ca acceptor (e, Ca) and a donor to acceptor (D, Ca) line. These are listed below, TABLE 1. In addition to these expected emission lines a line (actually a sharp doublet) labelled the SM line is located on the low energy side of the bound exciton line. The SM line is thought to be due to exciton decay at a Ca impurity located at a site which is different from the normal substitutional one. The Ca acceptor also produces a g(gca) line resembling the family of g lines seen in MBE material and implying that interactions between the Ca impurities and vacancies occur. A spectroscopic binding energy of 28 meV was deduced for the Ca acceptor. TABLE 1. PL lines from Ca ion-implanted into GaAs. eV

Assignment

Comments

1.5065

c

a novel emission at high [Ca]

1.5127

(Ca°, X)

1.5112

g

1.4913

(D, C) and (e, Ca)

SM' possibly (X, Cadeep°)

bound exciton

1.481

a broad emission

1.488

a shoulder feature

C4

Other Group II Elements

These are recorded in TABLE 2. TABLE 2. Spectroscopic features of group II impurities. Element

PL feature (e V)

T(K)

Ref

1.4923

2

[3]

1.4915

5

[7]

1.5

[7]

5

[7]

FB Be

BE(IS)

1.5124 _Mg

1.4911

THR (2S)

1.4926

Next Page

1.5124 Ca

1.4913

Sr

produces no extra PL emission lines

Zn

1.4888

[7] 5 1.4904

1.5

[7]

5 1.5123

1.461

[7] [23]

1.4848

Hg

1.5

1.5127

1.5122 Cd

1.4922

1.4869

1.5

[7] [24]

Note: for the excitonic features measured at high resolution at temperatures of below 2 K doublet centre energies are normally quoted. REFERENCES [I] [2] [3] [4] [5] [6] [7]

[8] [9] [10] II1] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24]

P.J. Dean [ Prog. Cryst. Growth & Charact. (UK) vol.5 (1982) p.89 ] P.J. Dean, DJ. Cuthbert, D.G. Thomas, R.T. Lynch [ Phys. Rev. Lett. (USA) vol. 18 (1967) p. 122] K. Akimoto et al [ Appl. Phys. Lett. (USA) vol.45 (1984) p.922 ] A.M. White, P.J. Dean, L.L. Taylor, RC. Clarke, DJ. Ashen [ J. Phys. C (UK) vol.5 (1972) p. 1727] A.M. White, P.J. Dean, DJ. Ashen, J.B. Mullin, B. Day, P.D. Greene [ J. Phys. C (UK) vol.6 (1973) p.L243 ] D.C. Reynolds, KJ. Bajaj, CW. Litton [ Solid State Comm. (USA) vol.53 no.12 (1985) p.1061 ] DJ. Ashen, P.J. Dean, DJ.T. Hurle, T.B. Mullins, A.M. White, P.D. Green [ J. Phys. & Chem. Solids (UK) vol.36 (1975) p. 1041]; see also P.W. Yu, M.Y. Yen, CE. Stutz [ J. Appl. Phys. (USA) vol.75 no.5 (1994) p.2628-32 ] Landolt-Bornstein [ vol. 17 (1982) p.224-30 ] CJ. Hwang [ Phys. Rev. (USA) vol. 180 (1969) p.827 ] K.D. Glinchuck, A.V. Prokhorovich, N.S. Zayats [Phys. Status Solidi (Germany) (1984) p.625 ] P. Yu, B. Welber [ Solid State Comm. (USA) vol.25 (1978) p.209 ] D. Dego, M. Cordona, H. Muller [ J. Phys. Soc. Jpn. (Japan) vol.47 (1976) p.631 ] H.C. Casey, F. Stern [ J. Appl. Phys. (USA) vol.47 (1976) p.631 ] N.S. Averkiev, SA. Obukhov, A.A. Rogachev, N.A. Rud [ Sov. Phys. Solid State (USA) vol.25 (1983) p. 193] D. Olego, M. Cardona [ Phys. Rev. B (USA) vol.22 (1980) p. 1905 ] D.N. Mirlin, LY. Karlik, L.P. Nikitin, I.I. Reshina, V.F. Sapega [ JETP Lett. (USA) vol.32 (1980) P-31] LA. Merkulov, A.V. Rodina [ Semicond. (USA) vol.28 no.7 (1994) p.720-4 ] J.C Bean, R. Dingle [ Appl. Phys. Lett. (USA) vol.35 (1979) p.925 ] P. De Mierry, M. Stutzmann [ Phys. Rev. B; Condens. Matter. (USA) vol.46 no.20 (1992) p. 13142-51] M. Illegems [ J. Appl. Phys. (USA) vol.48 (1977) p. 1278 ] N. Duhamal, P. Henoc, F. Alexandre, E.V.K. Rao [Appl. Phys. Lett. (USA) vol.39 (1981) p.49 ] E.A. Imhoff, M.I. Bell, R.A. Forman [ Solid State Comm. (USA) yol.54 (1985) p.845 ] Shen Hong-Lie et al [ Appl. Phys. Lett. (USA) vol.65 no. 11 (1994) p. 1427-9 ] W.P. Gillin, BJ. Sealy [J. Appl. Phys. (USA) vol.71 no.4 (1992) p.2021-2 ]

9.8

Photoluminescence spectra of group IV atoms in GaAs Previous Page

B. Hamilton (Updated by M.R. Brozel) September 1996

A

INTRODUCTION

Group IV elements can in principle form either donors or acceptors in GaAs by substitution into the Ga or As sublattices respectively (so called amphoteric behaviour). In practice the behaviour of group IV dopants depends critically on the growth or doping conditions and thermodynamics so that generalisation can be difficult. For example, Si tends to exhibit donor activity SiGa in most materials but under certain conditions shows acceptor activity in LPE GaAs [I]. Carbon favours acceptor activity [2] in LPE and MBE material and also in VPE material which is grown via the AsCl3 route [3]. Donor related transitions are difficult to resolve in all but the highest purity samples, i.e. at the residual level. High resolution spectroscopy and magnetic field splitting are needed. Several group IV donor states have been resolved using two electron transitions [4,5] and magnetic splitting [6,31] (FIGURE 1).

RELATIVE INTENSITY

Acceptor states, with larger binding energies and central cell corrections, are more readily distinguishable [7] using excitonic or free to bound transitions.

ENERGY (cV) FIGURE 1. Shallow donors identified in three different GaAs samples from the 2 P., transition at a magnetic field of 3 - 6 kG. In the solid curve, residual Si, S and Ge donors are seen in p-type GaAs. The dashed curve shows the Si donor in an intentionally Si doped sample. The dot-dashed curve shows S donors and Si donors in another Si doped sample. The inset shows the transitions that occur to the 2 P-1 state (taken from [4]).

B

CARBON AS AN IMPURITY

The chief dopant role of the carbon acceptor is as a background contaminant in many major growth technologies [2,3] causing compensation in n-type material and degradation of mobility. The potentially complex role of C in photoluminescence transitions has been revealed by C ion implantation into pure GaAs grown by MBE [8]. These authors report seven transitions related to carbon including the 'g' line (1.5110 eV) which forms part of the sharp defect bound exciton luminescence commonly observed in MBE material [9]. High resolution data for the C neutral bound exciton (A0, X) J = 3/2 and J = 5/2 transitions are reported in [6]. A transition resulting from recombination between a donor excited (n = 2) state and C (Zn and Si) acceptors has been observed between the normal free to bound (FB) and donor-acceptor pair (DAP) transitions [10]. C

CARBON AS A DOPANT

Dilute carbon implants into GaAs produced a g line 1.8 meV lower in energy than the A°-X line [49]; this work also produced evidence for a series of g lines (the [g-g] family) which may result from centres involving both C and O. The use of C as a dopant, especially for growing p+ layers in devices like HBTs, has led to several PL investigations of C doped GaAs [32-41] including the change in Eg as a function of carbon concentration. D

SILICON

The predominant electrical activity of Si in GaAs is that of a donor SiGa. High resolution measurements of transitions from magnetically split excited states of excitons bound to Si donors have been made [6]. Such measurements are only possible at very low donor concentrations and normally Si is observed in the collective donor bound exciton peak in this spectral region. In addition to SiGa, Si^ and Si in a variety of complexed forms have been observed by photoluminescence. Early work on Si doping of LPE material showed that Si could be made to favour the As site and that a variety of Si associated acceptor levels seemed to form [1], though at the relatively high doping levels produced, the luminescence bands were too broad to make definite assignments. Luminescence from what is thought to be the acceptor Si^ at 1.485 eV has been reported for MBE material [H]. Silicon is an important n-type dopant in MBE growth. Recent work has shown extremely high quality spectra from lightly Si doped MBE layers [12]. These authors concluded that much residual dopant (Si and C) incorporation originated in the dopant sources. They further concluded that longer wavelength emission observed in MBE material (1.466 - 1.482 eV) was due to DAP recombination involving four rather deep, but unidentified acceptors. Silicon is also the major n-type dopant in ion implantation (into SI GaAs) technology. Photoluminescence has been used [13] to measure the ratio OfSi03ISiA8 and hence the degree of

n-type activation for various rapid anneal schedules following implantation. The increase of FB luminescence, from the Si^ acceptor, with anneal temperature demonstrates the siting preference of Si on the As site at higher temperatures. Uniaxial stress splitting of free exciton luminescence and an accompanying shift of donor bound exciton lines has been observed in Si implanted GaAs [14]. The complexes [ V03-Si^] (1.35 eV) ^ d [VAs-Si0J (1.23 eV) together with SiGa and Si^ have been observed in annealed Si implanted SI material [15]. The ratio of the various defect species was shown to depend on the composition of capping material and its effectiveness as a barrier to Ga out-diffusion during anneal. A gallium vacancy related band (probably [VGa-SiGa]) has been observed in Si doped MOCVD GaAs [15]. The intensity of the band varied with arsine molar ratio in the growth reactor. Similar bands have been observed in a variety of Si doped GaAs materials technologies, and similar luminescence is associated with other donor species [16]. In heavily Si doped material the role of conduction band tail states and compensation is important [1,17] and a recent analysis argues that the luminescence band shape and its excitation dependence can be explained by competition between valence band to tail transitions and transitions between tail states originating from both bands [18]. The complex nature of compensation caused by site competition between Si02 and Si^ and its influence on DAP transitions is discussed in [19]. Time resolved luminescence has been used to elucidate the dynamics of radiative recombination in heavily Si doped GaAs [20]. Recently, PL has been used to investigate doping concentration effects [41], effects of annealing [42], the influence of hydrogen in degenerate materials [43] and Si delta-doped GaAs [44]. In particular, excitation spectroscopy [42] using photoluminescence detection has been used to help clarify the nature of the defect centres involved in the two deep luminescence bands associated with silicon. The absorption process leading to both deep Si related bands exhibited a strong peak near 1.49 eV which has been associated with the [Si^-Si^] defect, strengthening the link with Si and native defects. E

GERMANIUM

Neutron transmutation doping (NTD) is used to controllably dope material with Ge (Ga transmuted to Ge) and Se (As transmuted to Se). Photoluminescence has been used to study the formation of donor species and acceptor species caused by site switching e.g. GeGa to Ge^ (acceptor). There is strong evidence for this acceptor incorporation in [2] monitored through the intensity of DAP and Ge A0X luminescence. Further evidence for site switching based on PL analysis is given in [21]. Changes in luminescence behaviour of NTD doped GaAs which relate to the properties of the neutron irradiation are discussed in [22]. Magnetic splitting of the Ge neutral bound exciton is reported in [6]. Germanium doping in LPE AlxGa1^xAs with x varying from 0 to 0.32 is reported in [23]. The material remained p-type and the shift of Ge related peaks showed that the ionisation at the Ge acceptor depends strongly on x. Luminescence from heavily Ge doped material is reported in [24] which demonstrates that

compensation is important in determining band shape. Ion implantation of Ge into GaAs has been reported in [25]. Luminescence spectra supported the view that significant damage remained after annealing, a peak at 1.360 eV (T = 6 K) being generally present after annealing and thought to be linked with Ga vacancies [26]. The epilayer substrate interface in Ge doped LPE material has been investigated in [27]. It was found that a defect line at 1.413 eV thought to be associated with carbon-arsenic vacancies [28] was reduced, probably by Ge occupation of the vacancy site. Recent PL investigations have considered effects of the application of a magnetic field on LEC growth of GaAsGe [45] and the complexing behaviour of In and Sb with Ge in p-type GaAsGe epitaxial layer [46]. F

TIN

Rather deep and broad bands are observed for Sn doped GaAs for various growth technologies. These bands have been associated with Ga vacancies in the case of both Sn and other donors [16]. The effect of excitation level, frequency response and thermal quenching have been used [29] to investigate the kinetics of a 1.39 eV band in Sn doped LPE material. This author discounted donor-acceptor transitions as a mechanism for this band, but invoked intra-centre or free to bound transitions. A broad PL band near 1.2 eV has been investigated by stress splitting. Results are interpreted in terms of the re-alignment Of[Sn03 - VGa] complexes [47,48]. Free to bound transitions were used to model the lineshape of Sn related luminescence in [30]. Luminescence decay and temperature evolution were used to show that conduction band to deep acceptor transitions fit the observed characteristics and to explain peak shift with doping. Group IV associated PL transitions are given in TABLE 1. TABLE 1. Group IV PL transitions. Element

Peak energy (eV)

Comments

Reference

C

1.4935

FB transition measured at 5 K

U]

1.5124

(A0, X) measured at 1.5K 1S doublet centre

U]

1.4850

FB measured at 5 K

U]

Si

Ge

0

1.5123

(A , X) measured at 1.5K 1S doublet centre

U]

1.4790

FB measured at 5 K

U]

Peak energy (eV)

Comments

Reference

1.5126

(A0, X) measured at 1.5K 1S doublet centre

[7]

1.349

FB measured at 5 K

[7]

1.5067

(A0, X) measured at 1.5K 1S doublet centre

[7]

1.4843

FB highly resolved in MBE material

[H]

[V 03 -SiJ

1.35

measured at 12 K

[14]

[V 04 -SiJ

1.23

measured at 12 K

C associated

1.5100

excitonic, measured at 2 K identical with g line in MBE GaAs [9]

[8]

[V 01 -Sn 0 J

1.2

stress splitting

[47,48]

Element

Sn

Complexes:

REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] II1] [12] [13] [14] [15] [16] [17] [18] [19]

H. Kressel, J.U. Dunse, H. Nelson, F.Z. Hawrylo [ J. Appl. Phys. (USA) vol.39 (1968) p.2006 ] T.S. Low, M.H. Kim, B. Lee, B.J. Skromme, T.R. Lepkowski, CE. Stillman [ J. Electron. Mater. (USA) vol.14 (1985) p.477] M. Ozeki, K. Kitahara, K. Naki, K. Dazai, S. Okawe, O. Ruyzar [ Jpn. J. Appl. Phys. (Japan) vol.16 (1977) p. 1617] D.C. Reynolds, CW. Litton, E.B. Smith, P.W. Yu, K.K. Bajaj [ Solid State Commun. (USA) vol.42 (1982) p. 827 ] RJ. Almassy, D.C Reynolds, CW. Litton, K.K. Bajaj, CL. Macoy [ Solid State Commun. (USA) vol.38 (1981) p. 1053] D.C Reynolds, P.C Colter, CW. Litton, E.B. Smith [ J. Appl. Phys. (USA) vol.55 (1984) p. 1610] DJ. Ashen, PJ. Dean, DTJ. Hurle, J.O. Mullin, A.M. White, P.D. Greene [ J. Phys. & Chem. Solids (UK) vol.36 (1975) p. 1041 ] Y. Makita et al [Nucl. Instrum. & Methods Phys. Res. B (Netherlands) vol.7/8 (1985) p.433 ] H.Kunzel,K.Ploog[/«5r. Phys. Con/ Ser. (UK) no.56 (1981) p.519] BJ. Skromme, CE. Stillman [ Phys. Rev. B (USA) vol.29 (1984) p. 1982 ] H. Kunzel, K. Ploog [Appl. Phys. Lett. (USA) vol.37 (1980) p.416 ] BJ. Skromme et al [ J. Appl. Phys. (USA) vol.58 (1985) p.4685 ] S. Tiku, W.M. Duncan [ J. Electrochem. Soc. (USA) vol. 132 (1985) p.2237 ] YJ. Chan, M.S. Lin, T.P. Chen [ J. Appl. Phys. (USA) vol.58 (1985) p.545 ] K. Okamoto, S. Onozawa, T. Imai [ J. Appl. Phys. (USA) vol.56 (1984) p.2293 ] E.W. Williams [ Phys. Rev. (USA) vol. 168 (1968) p.922 ] V.A. Vilkotskii, DS. Domanevski, S.V. Zhokovets, M.V. Prokopenya [ Sov. Phys. Semicond. (USA) vol.18 (1984) p. 1368] V.L. Korolev, V.V. Rossin, V.G. Sidorov, YuK. Shalabutov [ Sov. Phys. Semicond. (USA) vol. 19 (1985)p.325] H. Takano, T. Kamijoh, M. Sakota [ J. Lumin. (Netherlands) vol.31/32 (1984) p.442 ]

[20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33] [34] [35] [36] [37] [38] [39] [40] [41] [42] [43] [44] [45] [46] [47] [48] [49]

E.O. Gobbel, W. Graudszus [ Phys. Rev. Lett. (USA) vol.48 (1982) p. 1277 ] J. Carrido, J.L. Castano, J. Piqueras [ J. Appl. Phys. (USA) vol.57 (1985) p.2186 ] L.I. Kolenski et al [ Sov. Phys.-Semicond. (USA) vol. 19 (1985) p.742 ] Y.R. Yuan, K. Mohammed, J.L. Merz [ J. Appl. Phys. (USA) vol.57 (1985) p.2986 ] V.F. Kovalenko, V.A. Krasnov, Yu.E. Karanchuk j Sov. Phys.-Semicond (USA) vol.18 (1984) p.575] S.S. Chan, CT. Marcyk, B.C. Streetmand [ J. Electron. Mater. (USA) vol.10 (1981) p.213 ] P.K. Chatterjee, K.V. Vaidyanathan, M.S. Durschlag, B.C. Streetman [ Solid State Commun. (USA) vol.17 (1975) p. 1421] W.Y. Lum, A.R. Clawson, D.I. Elder, H.H. Wieder [ J. Appl. Phys. (USA) vol.49 (1978) p.3333] W.H. Koschel et al [ Inst. Phys. Conf. Ser. (UK) no.33a (1977) p.98 ] CB. Norris [ J. Appl. Phys. (USA) vol.48 (1977) p.4706 ] S. Zemon, M.O. Vasell, C Lambert, R.H. Bartram [ J. Appl. Phys. (USA) vol.53 (1982) p.3347 ] S. Zemon, G. Lambert [ J. Appl. Phys. (USA) vol.70 no.9 (1991) p.4909-18 ] S. Shigetomi, Y. Makita, A.C. Beye, A. Yamada,N. Ohnishi, T. Matumori[7. Appl. Phys. (USA) vol.69 no.3 (1991) p. 1613-17] Lei Wang, B.J. Aitchison, N.M. Haegel [ Appl. Phys. Lett. (USA) vol.60 no.9 (1992) p. 111-13 ] Seong-Il Kim, Moo-Sung Kim, Yong Kim, Kyung Sook Eom, Suk-Ki Min, Choochon Lee [ J. Appl. Phys. (USA) vol.73 no.9 (USA) p.4703-5 ] Seong-Il Kim, Yong Kim, Min Suk Lee, Moo-Sung Kim, Suk-Ki Min, Choochon Lee [ Solid State Commun. (USA) vol.88 no.9 (1993) p.743-6 ] Seong-Il Kim, Moo-Sung Kim, Suk-Ki Min, Choochon Lee [ J. Appl. Phys. (USA) vol.74 no. 10 (1993)p.6128-32] Z.H. Lu, M.C Hanna, A. Majerfeld [Appl. Phys. Lett. (USA) vol.64 no. 1 (1994) p.88-90 ] Z.H. Lu, A. Majerfeld [ J. Appl. Phys. (USA) vol.75 no.5 (1994) p.2648-51 ] Lei Wang, N.M. Haegel, J.R Lowney [ Phys. Rev. B, Condens. Matter (USA) vol.49 no. 16 (1994) p. 10976-85] Horng Dar Chen, Ming Shiann Feng, Po An Chen, Kun Oman Lin, Janne Wha Wu [ Jpn. J. Appl. Phys. (Japan) pt. 1 vol.33 no.4A (1994) p. 1920-7 ] M. Suezawa, A. Kasuya, Y. Nishina, K. Sumino[J. Appl. Phys. (USA) vol.73 no.6 (1993) p.3035-40 ] M. Suezawa, A. Kasuya, Y. Nishina, K. SuminofJ. Appl. Phys. (USA) vol.76 no.2 (1994) p. 1164-8] M. Capizzi,V. Emiliani, A. Frova,F. Sarto [ Phys. Rev. B, Cond. Matter. (USA) vol.47 no.8 (1993)p.4301-6] J.C.M. Henningetal[ Sew/com* 5c/. Technol. (UK) vol.6 no. 11 (1991) p. 1079-87] J. Kang,K. Hoshikawa,M. Tajima,T. Fukuda[J. Cryst. Growth (Netherlands) vol.135 no.3-4 (1994)p.623-8] K.S. Zhuravlev, A.V. Katkov [ Sov. Phys.-Semicond. (USA) vol.25 no. 1 (1991) p.51-4 ] A.A. Gutkin, M.A. Reshchikov, V.R Sosnovskii [ Sov. Phys.-Semicond. (USA) vol.27 no.9 (1993) p.838-43 ] A.A. Gutkin, M.A. Reshchikov, V.R. Sosnivskii [ Sov. Phys.-Semicond. (USA) vol.27 no.9 (1993) p.844-8 ] Y. Makita et al [ Appl. Phys. Lett. (USA) vol.47 (1985) p.623 ]

9.9

Photohiminescence spectra of group VI shallow donors in GaAs B. Hamilton (Updated by M.R. Brozel) September 1996

Group VI elements form shallow donor states in GaAs by substitution on the As sublattice. Donor states in GaAs (and other low effective mass direct energy gap semiconductors) are extremely difficult to identify using photoluminescence because their bonding energies are small (around 6 meV) as are the chemical shifts associated with individual impurities [I]. Attempts to chemically identify Group VI elements using sharp excitonic features are restricted to carefully doped high purity samples, with donor concentrations below approximately 1015 cm"3; above this concentration value impurity screening leads to line broadening, and the observation of a single donor bound exciton line. Donor species have been implicated in the appearance of broad bands near 1.2 eV in melt grown and epitaxial n-type GaAs [2]. These bands may come from donor-gallium vacancy [D - VGa] complexes, and can be described by a multiphonon model on the strongly coupled limit with a coupling constant S approximately equal to 15. The bands are most readily observed in quite heavily doped material (approx. 1018 cm"3). Details are given in TABLE 1. The effect of doping level on the photoluminescence characteristics of S doped LPE material shows that the emission intensity increases with S concentration up to about [S] = 3 x 1018 cm"3 [3]. The half width of the S donor luminescence also increases. Above this S concentration the luminescence efficiency drops and this was attributed to microprecipitation of S. Sulphur associated deep level luminescence was also observed at ~ 1 nm. A comparison of luminescence efficiency for S and Se doping in VPE material shows [4] that the integrated intensity for both dopants increases linearly with electron concentration. However, Se doped material did not drop from linear variation until around ~ 1018 cm"3, whereas the efficiency of S doped crystals began to degrade at about 5 * 1017 cm"3. Direct evidence for microprecipitation of dopant was given in [4]. The influence of compensation (by Ge) on the band edge emission due to heavy Te doping indicates that a so-called 'mobile' band is produced which shifts towards lower energies with increasing excitation power [5]. Recombination via tail states was rejected as a mechanism for this band, in favour of a donor-acceptor complex model. A more detailed model by the same authors is discussed in [6]. Luminescence peaks have also been related to Te complexed with copper and with lattice vacancies. In [7], [Te^-CuoJ pairs were thought to be responsible for emission near 1.28 eV; and in [8,9] [V03 - Te^ - VAJ for emission at 1.38 eV (all measurements at 77 K). Electron irradiation and annealing was used to study the growth and complexing kinetics of these defects. Hydrostatic pressure was found to modify the properties of Te associated luminescence in a way which is dictated by the pressure dependence of the Fermi-level [8].

High resolution photoluminescence has been used [9] to resolve bound exciton luminescence from residual Te doping in VPE material. Luminescence transitions from excited (n = 2) non-rigid stator states of the exciton were observed at 1.510216 eV. Technique developments aimed at resolving donor states in GaAs using PL are reviewed in [11]. A more recent study on heavily Te-doped LPE GaAs reports further effects due to band filling, band tailing and bandgap reduction [10]. TABLE 1. Group VI related luminescence bands. element

peak energy, eV

comments

refs

S

1.197

[12]

Se

1.224

Te

1.216

These bands are commonly observed in heavily doped material (eg, [S] above 1018 cm"3) and have been linked to donorgallium vacancy pairs. They are strongly phonon coupled. Peak energies quoted here correspond to T approximately 80 K.

[Te^-Cu 0 J

1.28

Measured at 77 K

[7]

[V01-Te^-VxJ

1.35

Measured at 77 K

[8,9]

[12]

[12]

REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] II1] [12]

DJ. Ashen, PJ. Dean, DTJ. Hurle, J.B. Mullin, A.M. White [ J. Phys. & Chem. Solids (UK) vol.16 (1975) p.1041] T. Instone, L. Eaves [ J. Phys. C (UK) vol. 11 (1978) p.L771] M. Ishi, T. Tanaka, W. Susaki [ J. Cryst. Growth (Netherlands) vol.46 (1979) p.265 ] L. Hollan, M. Boulou, J.P. Chare [ J. Electron. Mater. (USA) vol. 10 (1981) p. 193 ] B.G. Arnaudov et al [ Sov. Phys. Semicond. (USA) vol. 17 (1983) p.378 ] B.G. Arnaudov et al [ J. Phys. C (UK) vol. 17 (1984) p.331 ] K.D. Glinchuck, A.V. Prokorovich, N.S. Zayats [ Phys. Status Solidi A (Germany) vol.72 (1982) p.715] D. Olego, M. Cardona, H. Muller [ Phys. Rev. B (USA) vol.22 (1980) p.894 ] P.C. Colter, D.C. Reynolds, CW. Litton, E.B. Smith [ Solid State Commun. (USA) vol.45 (1983) p.375] Chyuan-Wei Chen, Meng-Chyi Wu, Shoei-Chyuan Lu, Chung-Chi Chang [ Jpn. J. Appl. Phys. (Japan) vol.32 no.6A (1993) p.2725-30 ] D.C. Reynolds, K.K. Bajaj, CW. Litton [ Solid State Commun. (USA) vol.53 (1985) p. 1061 ] E.W. Williams [ Phys. Rev. (USA) vol. 163 (1968) p.922 ]

9.10 Photoluminescence spectra of transition and rare earth metals in GaAs A.M. Hennel and M.R. Brozel October 1989 Updated July 1996

A

INTRODUCTION

Transition metal atoms have important properties in GaAs. When introduced intentionally they produce deep electron traps which can compensate shallow donors and result in high resistivity material. The use of Cr for the compensation of residual shallow Si donors in HB GaAs was almost universal some years ago to produce SI GaAs. Many other transition elements produce high resistivity GaAs. They are also important as they are contaminants often introduced into GaAs wafers at the processing stage especially if high temperatures are involved: in general, they are very rapid difiusers. Small concentrations so introduced can have disproportionately adverse effects on devices. Photoluminescence (PL) is a powerful technique to detect these impurities at low concentrations. B

TITANIUM (Ti)

Titanium impurities in GaAs reveal a characteristic luminescence band between 2.6 and 2.2 microns (0.48 - 0.56 eV) with a sharp zero-phonon line (ZPL) at 4566 cm"1 [I]. This band is currently assigned to the intracentre 2T2 - 2E transition of the neutral acceptor charge state Ti3+ (3d1 electronic configuration) [2]. The absence of any shift or splitting of the corresponding absorption ZPL in a magnetic field [1,3] put some doubts on this interpretation. However, Zeeman effect measurements of analogous spectra in Ti-doped GaP [4] and InP [5] confirm the above model. Bl

Titanium-related Complexes

In GaAs crystals implanted with titanium another luminescence band between 2.15 and 1.85 microns (0.58 - 0.68 eV) with two ZPLs at 5365 and 5368 cm'1 was observed [6]. This structure depends upon annealing [7] and probably corresponds to titanium-related complexes. C

VANADIUM (V)

Vanadium impurities in GaAs reveal a characteristic luminescence band between 1860 nm and 1680 nm (0.67 - 0.74 eV) with two sharp ZPLs (zero-phonon lines) at 5958 cm"1 (cold line) and 5968 cm"1 (hot line) [8]. From uniaxial stress [9] and magnetic field [10,11] measurements of the ZPLs it was found that this band corresponds to the intracentre 3T2 - 3A2 transition of the neutral acceptor charge state V3+ (3d2 electronic configuration). Recent PLE measurements have reported a Jahn-Teller distortion of the 3T2 state resulting in splitting of this transition into a triplet [12].

Cl

Pressure Dependence

A hydrostatic pressure coefficient of the vanadium band is equal to 0.3 ± 0.1 meV/kbar [9] which corresponds to the increase of the crystal field strength under pressure. C2

Decay Time

Decay time of this luminescence band depends on temperature and a kind of excitation source and can be described by the formula: T(T) = t(o) [1 + C exp(-Ea/kT) J"1 Under argon laser excitation T(O) = 125 JIS, C = 1.8 xlO6, Ea = 0.29 eV; under YAG laser excitation T(O) = 86 \is, C = 25, Ea = 0.025 eV [13-15]. D

CHROMIUM (Cr)

Chromium impurity reveals a number of luminescence bands in GaAs attributed to different charge states of isolated chromium and to chromium related complexes. Dl

0.57 eV Band

In semi-insulating and p-type crystals there is a luminescence band between 3.1 and 1.85 microns (0.40 - 0.67 eV) with a zero-phonon line (ZPL) at 5370 cm"1 and a maximum at 0.57 eV [16,17]. This band is either attributed to the internal spin-forbidden transition 2E(G) - 4T1(F) within the neutral chromium impurity charge state Cr3+ (3d3) [16,17], or to a recombination of a single ionized chromium charge state with a hole: Cr2+(Sd4) + h+(VB) - Cr3+(Sd3) + hv [18,19] Measurements of a decay time of this luminescence band can be described by the formula: T(T) = T(O) [1 + C expt-EykT)]"1 with T(O) = 0.9 jis, C = 5 x 104, Ea = 0.2 eV [20]. This fast decay of the luminescence is in favour of the second model. D2

0.62 eV Band

In n-type crystals there is a luminescence band between 2.7 and 1.65 microns (0.50 - 0.75 eV) with a maximum at 0.62 eV and without any ZPL [18,21]. This band is attributed to a recombination of a neutral ionized chromium charge state with an electron: Cr3+(Sd3) + e(CB) - Cr2+(Sd4) + hv

D3

Internal Cr2+ Transitions

The single ionized chromium charge state Cr2+(Sd4) has one excited state 5E, which is degenerate with the GaAs conduction band. Due to this fact the internal transition 5E - 5T2 within the 3d4 configuration can be observed (together with the 0.62 eV band) only either under 1.32 micron YAG laser excitation or with above bandgap excitation under hydrostatic pressure higher than 0.8 GPa. This fact can be explained by the degeneracy lifting between CB and the 5E state. The Cr2+ band has a ZPL at 6620 cm"1 and a maximum at about 0.79 eV [17]. A hydrostatic pressure coefficient of this band is equal to 1.7 ± 0.2 meV/kbar [17] which corresponds to the increase of the crystal field strength under pressure. After further increase of pressure the intensity of this band reaches a maximum around 2 - 3 GPa and then quickly decreases between 3 and 5 GPa. In the same pressure region some new emission features were observed between 0.7 and 0.9 eV which are connected with a crossing of the excited 5E state with a low spin state (either 1A1 or 3T1) [22]. D4

Chromium-Arsenic Vacancy Complex

A characteristic chromium-related emission spectrum can be observed in high resistivity crystals between 1.77 and 1.48 microns (0.7-0.84 eV) with a maximum at 0.8 eV and a very rich ZPL structure (about 20 lines) at 6770 cm"1 [16,23]. The trigonal (C3v) structure of the centre was shown from magnetic field and uniaxial stress measurements of this ZPL structure [24-26]. From the thermal annealing experiments, it was found that the emission centre is a complex chromiumarsenic vacancy [27]. Consequently, this band is interpreted as the internal transition within the Cr2+^d4) configuration in C3v symmetry [28]. The decay time of this luminescence is equal to 0.7 ^s at 4 K [29]. D5

Chromium-Tellurium Complex

In crystals doped with chromium and tellurium additional luminescence bands can be observed. The most characteristic is a rich ZPL structure at 6810 cm"1 [28], interpreted as the internal transition within the 3d4 configuration of the Cr-Te complex. Magnetic field measurements have confirmed this model [30]. D6

Chromium-Selenium Complex

In crystals doped with chromium and selenium a rich ZPL structure around 6750 cm"1 was found and interpreted as the internal transition within the 3d4 configuration of the Cr-Se complex [31]. This model has been confirmed by diffusion profile measurements. D7

Chromium-Indium Complex

In GaAs: Cr crystals strongly doped with indium two new ZPLs at 6907 and 7041 cm"1 were observed. These lines are tentatively attributed to a [Cr - V^ - In] complex [32].

D8

Other Chromium-related Complexes

Several sharp luminescence lines at 4630, 4646, 4710, 4767 and 4823 cm"1 were observed in heattreated semi-insulating Cr-doped GaAs [33]. These lines probably correspond to some unidentified chromium complexes. E

MANGANESE (Mn)

Manganese impurity forms an acceptor level at about 0.11 eV above the top of the GaAs valence band [34]. Its luminescence spectrum is due to the recombination of electrons from shallow donors or the conduction band. El

Temperature Dependence

The Mn luminescence spectrum in GaAs is located between 880 nm and 920 nm (1.35 - 1.41 eV) [35]. At 4 K only donor-acceptor transitions with a zero-phonon line (ZPL) at about 1.403 eV can be observed; at higher temperatures (above 20 K) band to acceptor transitions with a ZPL at about 1.406 eV are dominant. E2

Pressure Dependence

A hydrostatic pressure coefficient of the manganese band is equal to 1.2 ± 0.2 meV/kbar [36,37] which corresponds to the increase of the Mn acceptor energy relative to the GaAs valence band. F

IRON (Fe)

Iron in GaAs reveals a characteristic luminescence band between 4.13 and 3.26 microns (0.30 0.38 eV) with five ZPLs at 3059, 3002, 2988, 2979 and 2962 cm"1 [35-37]. Four equidistant ZPLs (except the 3059 cm"1 line) correspond to the intracentre 5T2 - 5E transition of the ionized acceptor charge state Fe2+ (3d6 electronic configuration). This characteristic spectrum was observed in absorption and emission in other III-V materials [38]. Uniaxial stress measurements of these spectra performed in GaP [39] and InP [40] also confirm this interpretation. Fl

Decay Time

The decay time of this luminescence band can be described by the formula: T(T) = T(O) [1 + C exp(-EJkT)Yl where x(o) = 10 jis, C = 180, Ea = 0.05 eV [20]. F2

The 3057 cm 1 Line

The 3057 cm'1 line was originally related to an unidentified iron-related complex [36,37], but is now interpreted as the transition 4T1 to 6A1 of FeGa3+ (3d5) [72]. The decay time of this luminescence is 1.9 ± 0.3 ms [73].

G

COBALT (Co)

Cobalt in GaAs reveals a characteristic luminescence band between 2.8 microns and 2.5 microns (0.44 - 0.50 eV) with a ZPL at 4035 cm"1 [41]. This band is interpreted as the intracentre 4T2 - 4A2 transition of the ionized acceptor charge state Co2+ (3d7 electronic configuration) [42]. This interpretation is supported by uniaxial stress and magnetic field measurements of cobalt luminescence spectra in GaP and InP [38]. Gl

Cobalt-related Complexes

In GaAs doped with cobalt and tellurium it is also possible to observe a ZPL at 3885 cm"1 with a cold line at 3867 cm"1. This spectrum is interpreted as corresponding to the Co-Te complex [41]. Two other ZPLs attributed to unidentified complexes Co-X (3983 cm"1) and Co-Y (3545 cm"1) were also observed [41]. H

NICKEL (Ni)

Nickel impurities in n-type GaAs produce a characteristic luminescence band between 2.3 microns and 2 microns (0.54 - 0.62 eV) with several sharp ZPLs [42]. The shape of the band and the number of ZPLs depend on shallow donors in the crystal. Hl

Identification

From uniaxial stress and magnetic field measurements of a nickel-related absorption ZPL located at 0.57 eV it was found that this ZPL corresponds to the intracentre 2T2 -> 2E transition of the doubly ionized acceptor charge state Ni+ (3d9 electronic configuration) [43]. However, there is no emission spectrum corresponding to this specific ZPL, because the excited 2E state is degenerate with the GaAs conduction band. Observation of the luminescence ZPLs is clearly correlated with different shallow donors in the crystal [42]. This fact allows one to assign them to complexes involving donors and the Ni in the Ni+(3d9) charge state (TABLE 1). TABLEl

H2

Complex

Energy (cm1)

Ni-S complex

4427

Ni-Se complex

4409.8

Ni-Te complex

4369

Ni-Si complex

4699.2

Ni-Ge complex

4739.6

Ni-Sn complex

4629.8

Decay Time

Decay time of the luminescence band due to the Ni-Sn complexes can be described by the

formula: T(T) = T(O) [1 + C expC-E/kT)]"1 where T(O) = 4.2 us, C = 15, Ea = 0.015 eV [20]. I

COPPER (Cu)

Copper forms an acceptor level at about 0.15 eV above the top of the GaAs valence band [44]. Its luminescence spectrum due to the recombination of electrons either from shallow donors or the conduction band can be observed between 910 and 960 nm (1.29 - 1.36 eV) with a strong ZPL at about 1.36 eV [45,46]. II

Pressure Dependence

The hydrostatic pressure coefficient of this band is equal to 3 ± 2 meV/GPa [47], which corresponds to the increase of the Cu acceptor energy relative to the GaAs valence band. 12

Identification and Other Bands

Although there is a consensus that the 1.36 eV luminescence band is due to Cu, there is no definite identification of this luminescence band. The same situation exists for other luminescence bands attributed to copper observed at 1.04 eV and 1.25 eV [48,49]. J

SILVER (Ag)

Silver forms an acceptor level at about 0.24 eV above the top of the GaAs valence band [50]. A luminescence spectrum due to the recombination between shallow donors and Ag acceptors can be observed at 4 K between 1060 nm and 960 nm (1.17 - 1.29 eV) [51]. Due to a strong phonon coupling to longitudinal-optical (LO) phonons a zero-phonon line (ZPL) of this band cannot be seen directly. A numerical simulation of the band shape enabled the ZPL to be located at about 1.282 eV [51]. Jl

Pressure Dependence

The hydrostatic pressure coefficient of the silver band is equal to 0.5 ± 0.2 meV/kbar [52] which corresponds to the increase of the Ag acceptor energy relative to the GaAs valence band. K

GOLD (Au)

Au results in the appearance of photoluminescence bands with maxima near 1.07, 1.425 and 1.46 eV. The profile of the 1.07 eV band was approximated by assuming an interaction of holes localized at the luminescence centres with LO and TA phonons. The zero-phonon transition energy (± 0.005 eV) and the constants of the interaction with LO and TA phonons (1.0 ±0.1 and 2.9 ± 0.3 eV, respectively) have been determined. The energy of the zero-phonon line agrees well with the energy of the Ev + 0.4 eV level observed in GaAs: Au by electrical methods.

Kl

Pressure Dependence

The behaviour of the 1.07 eV band under uniaxial pressure is qualitatively similar to the behaviour of the photoluminescence of CuGa° and Ag080 suggesting that the centre responsible for this band is a doubly charged AuGa° acceptor whose environment is distorted by the Jahn-Teller effect so that the complex [AuGa° - 4As] has tetragonal symmetry. Quantitative deviations from the behaviour of the CuGa° photoluminescence is associated with higher values of the Jahn-Teller stabilization energy of AuGa° and of its ratio to the spin-orbit splitting [53]. L

TUNGSTEN (W)

Tungsten atoms implanted in GaAs result in two characteristic luminescence bands. After implantation in n-type crystals a band between 1.82 microns and 1.75 microns (0.68 - 0.71 eV) with a zero-phonon line (ZPL) at 5694 cm'1 was observed [54,55]. After implantation in p-type crystals a band between 1.90 microns and 1.82 microns (0.65-0.68 eV) with a ZPL at 5468 cm'1 was observed [54,55]. Ll

Identification

There is no definite identification of these luminescence bands. Their Fermi-level position dependence suggests that the first one corresponds to an intracentre transition of the ionized acceptor charge state W2+ (5d4 electronic configuration), and the second corresponds to an intracentre transition of the neutral acceptor charge state W3+ (5d3 electronic configuration) [56]. The spin state of the 5d configuration is unknown; theoretical calculations predict a low-spin state for the 5d4 shell [57]. M

NIOBIUM (Nb)

Niobium in GaAs reveals a characteristic luminescence band between 1.70 microns and 1.55 microns (0.73 - 0.80 eV) with a zero-phonon line at 6416.4 cm"1 [54,58]. Ml

Identification

There is no final identification of this emission band. From magnetic field measurements of the ZPL it was suggested that this band corresponds to the intracentre 3T2 - 3A2 transition of the neutral acceptor charge state Nb3+ (4d2 electronic configuration) [59]. However, from subsequent magnetic field and uniaxial stress measurements it was concluded that this band corresponds to the intracentre 1A1 - 3A2 transition of the same Nb2+ (4d2) charge state [60]. N

TANTALUM (Ta)

Tantalum atoms implanted in n-type and p-type GaAs result in a characteristic luminescence band between 2.07 microns and 1.94 microns (0.60 - 0.64 eV) with a zero-phonon line at 5160 cm"1 [54,61]. Nl

Identification

There is no definite identification of this band. However in analogy to niobium it was suggested

[60,62] that this band corresponds to the intracentre 1A1 - 3A2 transition of the neutral acceptor charge state Ta3+ (5d2 electronic configuration). 0

RARE EARTH ELEMENTS

Rare earth impurities in GaAs reveal characteristic luminescence bands with multiple ZPLs between 1.6 and 0.9 microns [63]. 01

Identification

All these bands are interpreted as intracentre transitions within trivalent RE ions. However, observed spectra are slightly dependent on doping technique as well as growth and annealing conditions. These facts suggest that a number of different RE centres in GaAs may be present. The band positions (in cm'1 and microns), and symmetries of electron states are collected in TABLE 2. TABLE 2 Impurity

cm"1

X (^m)

Ground state

Excited state

Reference

Pr3+(4f2)

6200*

1.61

4

3

[64,71]

U3+(5f3)or

6246

1.60

Er3+(4fn)

6500*

1.54

4

I l5/2

4

I1372

[66,70]

Nd 3+ (4f 3 )

7000*

1.43

4

I1372

4

F3/2

[67,68]

Pr3+(4f2)

7650*

1.31

3

H5

1

G4

Tm3+(4f12)

8100*

1.23

3

H6

3

H5

[64]

Nd3+(4f3)

9000*

1.11

4

I1172

4

F372

[67]

Pr3+(4f2)

9700*

1.03

3

H4

1

Yb3+(4f13)

10000*

1.00

2

F772

Nd 3+ (4f 3 )

11000*

091

4

I272

H4

F3

[65]

U4+(5f2)

[64,71]

G4

[64,71]

2

F572

[69]

4

F32

[67]

* Indicates more than one emission line.

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9.11 Room temperature photoluminescence mapping of GaAs substrates and epitaxial layers CJ. Miner and CJ.L. Moore August 1996

A

INTRODUCTION

Photoluminescence (PL) mapping of GaAs has been used for over 10 years to characterize the uniformity of substrates and epitaxial layers. The development of the method and its applications over this period can be followed in [1-16]. A PL map is a record of the spatial and spectral distribution of the photons produced through recombination of electrons and holes created by above bandgap illumination of the sample. The PL signal contains information about the material's bandgap energy, doping level and the presence of electrically active defects. Since room temperature PL mapping is a relatively simple and non-destructive test, it has become one of the primary tools in the development of improved growth techniques. In addition, room temperature PL is highly sensitive to surface recombination, which can be a bane to the researcher studying the bulk properties of the sample or a boon to those investigating surface passivation. In this Datareview, the interpretation of room temperature PL measurements, experimental considerations and typical results are discussed. BASIC PRINCIPLES

Bl

Doping Measurements Using PL Spectra

PL Intensity

B

Wavelength (nm) FIGURE 1. Typical room temperature photoluminescence spectra for n- and p-type GaAs showing the effect of doping.

Unlike the low temperature PL spectra discussed in Datareviews 9.1 to 9.9 in this book, the room temperature PL spectra of GaAs are usually dominated by a single broad peak. Examples of room temperature PL spectra for GaAs samples with a range of doping levels are shown in FIGURE 1. A review of the general theory of spectral shape may be found in [ 17,18]. In the limit of low excitation power density and doping levels below mid 1017 cm"3, the PL spectra of GaAs peak at -870 nm, corresponds to the donor to valence band transition [Eg - E d = 1.423 eV]. Under these conditions, the PL peak width is limited by thermal broadening to slightly more than kT (typically -22 nm at 300 K RT). As the n-type doping level is increased beyond mid 1017 cm"3, the dominant PL transition is between electron states near the Fermi energy and the valence band maximum. Since the Fermi energy lies above the conduction band minimum at high n-type doping levels, the PL peak wavelength drops below that corresponding to the bandgap energy. This is the well known Burstein-Moss effect. Similar changes to spectral shape occur at very high excitation power densities when excessive numbers of photo-generated carriers mimic the behaviour of high n-type doping. In contrast, the formation of an acceptor impurity band in p-type GaAs, centred 30 meV above the valence band maximum, shifts the PL peak to higher wavelengths as doping increases.

p-type GaAs

PL Peak Energy (eV)

PL Peak Wavelength (nm)

Room temperature PL peak positions for both n-type and p-type GaAs are shown in FIGURE 2.

n-type GaAs

Carrier Concentration (cm"3) FIGURE 2. Variation in the room temperature photoluminescence peak position for n- and p-type GaAs as a function of carrier concentration. The data describes both bulk doped substrates and epitaxial layers where the carrier concentration was measured by Hall effect. The curves are similar to those acquired by Cusano [19] for room temperature cathodoluminescence peak positions.

In both n-type and p-type samples, the range of states that can participate in a radiative transition increases with doping and so the PL peak broadens with increasing carrier concentration, FIGURE 3. Thus, variations in the PL spectral width or peak position for GaAs samples can be used to determine changes in doping above mid 1017 cm"3.

PL fwhm (nm)

n-type GaAs

p-type GaAs

Carrier Concentration (cm"3) FIGURE 3. Variation in the room temperature photoluminescence peak full width half maximum (fwhm) for n- and ptype GaAs as a function of carrier concentration. The data describes both bulk doped substrates and epitaxial layers where the carrier concentration was measured by Hall effect. Similar data for n-type GaAs were reported by Casey and Kaiser [20].

B2

Doping Measurements Using PL Intensity

Since the recombination rate is limited by the number of free carriers [18], doping is also expected to influence the PL peak intensity. Under low excitation density conditions in moderately doped materials, where the density of photo-generated electron hole pairs is less than the density of majority carriers but significantly greater than the density of competing non-radiative recombination paths, the PL peak intensity is linearly dependent on doping level and excitation power density [21]. This is observed for both n-type and p-type GaAs (FIGURE 4). Thus PL intensity mapping is a more sensitive measure of doping than either peak position or peak width and can be applied to device quality GaAs with doping levels between 1015 and mid 1017 cm"3 [5]. At higher dopant levels, non-radiative recombination at the numerous defects that are formed as the dopant solubility limit is approached usually causes the PL intensity to decrease precipitously. Because the onset of significant defect formation is dependent on the growth conditions, the use of PL intensity to map doping levels is not recommended for doping levels above 1018 cm"3. At doping levels below 1015 cm"3, the variation of background impurities and crystal defects, especially those that act as non-radiative recombination centres, will significantly affect PL

PL Peak Intensity (arbitrary units)

p-type epilayers n-type MBE epilayers n-type CBE epilayers n-type substrates

Carrier Concentration (cm"3) FIGURE 4. Variation in the room temperature photoluminescence peak intensity for n- and p-type GaAs as a function of carrier concentration. The data describes both bulk doped substrates and epitaxial layers where the carrier concentration was measured by Hall effect.

intensity. Under those conditions, maps of PL intensity will contain signatures of both doping and defect variations. By reducing the excitation power density, the contrast due to defects can be enhanced [22]. At the other extreme of very high excitation conditions, where the density of photo-generated electron hole pairs is greater than the density of majority carriers, the PL intensity is independent of doping and varies as the square of excitation power density. The PL intensity maps taken under these conditions are generally featureless and of little value for doping or defect mapping. B3

Defect Monitoring Using PL Intensity

PL intensity maps are often acquired specifically to detect defects. Once the variations in PL intensity due to carrier concentration are taken into account, the reduction in PL intensity can be used as a sensitive measure of deep level concentration. For example, after calibrating the PL intensity with independent secondary ion mass spectrometry measurements as shown in FIGURE 5, researchers in our laboratory were able to use the PL intensity as a non-destructive monitor of oxygen incorporation in epitaxial Al03Ga0^7As. In a similar fashion, a correlation was established between InGaAs photodiode leakage current and the reduction of PL intensity measured before processing [23]. Both are controlled by the density of non-radiative recombination centres. The photoluminescence mapping of semi-insulating GaAs represents a special case of defect monitoring under low doping conditions. GaAs is made semi-insulating by the careful balance of shallow acceptors, shallow donors and a significant density of mid-gap levels. Even under moderate excitation conditions, the density of deep levels will exceed the photo-generated carrier density by orders of magnitude. Under these conditions it can be shown that the PL intensity will

PL Peak Intensity (arbitrary units)

Oxygen Concentration (cm"3) FIGURE 5. Variation in the room temperature photoluminescence peak intensity for n-type Al03Ga07As epitaxial layers grown by chemical beam epitaxy as a function of oxygen impurity concentration. The oxygen concentration was determined by secondary ion mass spectrometry.

depend inversely upon the non-radiative deep level concentration and on excitation power density to the three halves power [7]. PL intensity maps will show individual dislocations as dark spots usually imbedded in brighter regions many tens of microns wide. The PL bright zones are believed to be created by defect gettering at the dislocation cores which reduces the non-radiative recombination centre density in the surrounding regions [1,2,24]. B4

Surface Recombination Effects

In the foregoing discussion the influence of surface or interface recombination has been ignored. Surface recombination in GaAs is not negligible, and indeed, variations in PL intensity have been used as a sensitive measure of surface passivation [25-27]. For the purposes of mapping doping variations using PL intensity, it is sufficient to ensure that the surface and interface conditions are uniform so that variations across the sample can be correctly interpreted. The PL contrast associated with defects, on the other hand, is affected both by the uniformity and the absolute degree of surface passivation [8,10,28]. Interpretation of sample-to-sample defect variations is therefore more problematic unless measures are taken to create a stable reproducible surface, for example, through a variety of solution [8,29] or plasma based passivation processes [10,30] or by the deposition of an optically transparent coating with low interface recombination velocity (e.g. AlxGa^xAs [8]). B5

PL Transients

A further complication to the PL mapping of lightly or undoped GaAs is the transient associated with changes in the surface passivation induced by exposure to the PL excitation source [8,10,31].

Time constants of hundreds of ms or longer have been observed. Current models for this phenomenon include photo-induced oxidation [31] and photo-enhanced As interstitial migration [32]. These two models address the semi-permanent decrease in PL intensity observed in air exposed samples during PL measurements. Removal of the upper 100 - 300 nm thick surface layer restores the original PL intensity [8,10]. The excitation wavelength and power density strongly affect the magnitude of the decrease [10]. Local heating and interactions with long time constant surface states also can cause the PL intensity to drop with prolonged exposure but unlike the semi-permanent transients just described, the PL intensity recovers once the excitation source is turned off. Both kinds of transients make the absolute PL intensity for air exposed samples difficult to reproduce. The relative change in PL intensity across a sample, however, is still a reliable semi-quantitative measure of doping or defect density. C

EXPERIMENTAL CONSIDERATIONS

There are two major types of PL maps: those where the pixels are collected simultaneously in the form of an image and those built up by rastering the illumination source over the sample and detecting the PL from each point sequentially. The former type of system; tends to be used for panchromatic PL imaging of small areas with high spatial resolution (e.g. [12]), whereas the latter are more suitable for spectrally resolved spatial mapping of whole wafers (e.g. [6]). In both cases a microscope can be used to increase spatial resolution. More recently, new nanoscale PL mapping methods have been derived from scanning probe microscopies and near field optical microscopies [33-35]. A fibre tip or other form of waveguide probe is rastered over the sample to either excite or collect the PL signal. The requirements for the sample, the excitation sources, the collection optics and detection components that are used in the two major PL mapping approaches are discussed in the following sections. Cl

Sample Preparation

As discussed, the sample suitable for PL mapping must have surfaces with uniform recombination velocities. This can be achieved in practice by exposing the as-received samples to the same set of chemical steps, typically consisting of degreasing followed by native oxide removal and oxide regrowth. Some researchers have precisely controlled oxide regrowth by water washing [29], plasma exposure [30] or UV ozone treatments. Sulphur passivation can also be used [25-28], although the effect is not stable with prolonged air exposure. Researchers who use thicker passivation coatings such as AlGaAs epitaxial cap layers must be aware that optical interference can change spectral shape or the excitation and PL collection efficiencies. Other factors in material preparation relate to the optical properties of the sample. The surface must be flat enough for all areas in the map to be equally in focus. The smaller field of view in the imaging systems reduces the severity of this restriction. Rough surfaces are to be avoided in either case, since the amount of excitation light absorbed in the sample and the amount of PL that is extracted can be changed locally by non-uniform scattering. C2

Excitation Sources

A variety of excitation sources can be used for PL mapping, so long as photons with energy in excess of the room temperature bandgap of GaAs are produced. These include filtered broad band sources, HeNe, Ar, HeCd and diode lasers. The penetration depth of the illumination varies for

the different excitation sources (TABLE 1). TABLE 1. Comparison of typical PL excitation laser sources. Laser

Wavelength (nm)

Energy (eV)

Penetration depth (iim)

HeCd

_488

L96

0A2

Ar ion

_514

_2_41

OM

HeNe

_633

_L96

026

AlGaAs

L72O

I 1.72

1 0.50

The most widely used are the HeNe lasers emitting at 633 nm since these typically have excellent temporal and pointing stability, good reliability and a well defined beam shape. The PL imaging systems are less sensitive to temporal fluctuations in excitation source intensity, but require more exact spatial filtering to achieve uniform illumination. Incoherent light sources such as wavelength filtered broad band sources are easier to work with in these systems, although they usually deliver less power. In scanning systems, lasers are more commonly used. The Ar and HeCd lasers are less suitable because their emission wavelengths lie near the resonance peak for enhanced photo-induced surface degradation [10]. The excitation power density is a critical parameter in determining the PL contrast mechanism and spectral peak shape in either type of PL mapper. Excessive excitation power densities can induce artifacts due to local heating or band filling. In theory, the total excitation reaching the sample and the effective spot size can be measured. In practice, however, the uncertainty in the actual distribution of light across the beam makes this calculation approximate at best. A more practical approach is to measure the PL spectra for a range of excitation powers to confirm the expected power dependences. C3

Collection and Detection Components

The RT PL peak wavelength of GaAs is well matched to both commercially available IR enhanced photo-multiplier tubes and solid state detectors such as silicon or germanium photodiodes. The latter have the advantage of shorter response time that makes them the preferred choice for scanning systems. In the special case of deep level PL mapping, where the excitation and detection wavelengths are very different, separate optical paths for illumination and PL collection allow each to be optimized [16]. The choices for spectrally analyzing the PL signal include scanning monochromators, fixed monochromators with optical multichannel analyzers and a selection of narrow pass filters. Given the relatively broad spectral peaks in RT PL, the transfer efficiency of the detection path is more important than spectral resolution. Spectral resolution of a few nanometres is adequate. D

TYPICAL APPLICATIONS

In this section, the principles of PL mapping will be demonstrated using two typical materials: ntype and undoped semi-insulating GaAs substrates. For easy comparison, one of the experimental configurations discussed above is used for all examples. As described in more detail in [6], the PL mapping system used is a spectrally resolved scanning PL instrument, using a 633 nm HeNe laser,

a lens of numerical aperture = 0.4 and a silicon detector. The excitation power delivered to the sample is 80 ^W and the spot size is 12 |im in diameter. From power dependence studies the excitation power density is confirmed to be low. Dl

PL Mapping of n-type GaAs Substrates

FIGURES 6, 7 and 8 show PL intensity maps of Si-doped n-type GaAs substrates grown by vertical gradient freeze (VGF)5 vertical Bridgman and a vertical boat growth technique, respectively. In each case, the whole 75 mm diameter wafer and a square region from the centre of the wafer are shown. The nominal doping level for each sample was 1018cm"3.

FIGURE 6. Whole wafer and higher resolution centre area PL intensity maps of silicon doped n-type GaAs grown by the vertical gradient freeze method. The wafer diameter is 75 mm in this case and for FIGURES 7-11.

For the VGF wafer shown in FIGURE 6, the bright regions were correlated with a two fold increase in carrier concentration as measured by Hall effect. This is consistent with the data in FIGURE 4. The PL bright core and streaks along the diagonals are seed artifacts typical of Sidoped GaAs grown by VGF and have been studied in detail by Toda et al [15]. These workers use high spatial resolution capacitance-voltage profiling to show that the carrier concentration is increased in the bright core and diagonal streaks. Individual dislocations may be seen in the detailed portion of FIGURE 6 as dark spots within the bright regions. KOH etch pit maps confirn the identification of the PL dark spots as dislocations. Unlike defect etching, C-V profiling or Hall effect measurements, PL mapping is non-destructive. The dislocation density in the vertical Bridgman sample (FIGURE 7) is 104 cm"3 and is two orders of magnitude higher than in the other two. The dislocations are clearly seen in the detailed map. They are arranged in a cellular pattern similar to that seen in semi-insulating GaAs. Such a PL map could be used for dislocation counting. Using an extrapolation from the VGF results, it is proposed that the enhanced photoluminescence along the dislocation cell walls is due to local

FIGURE 7. Whole wafer and higher resolution centre area PL intensity maps of silicon doped n-type GaAs grown by the vertical Bridgman method.

doping variations. Doping variations may occur even if the dopant concentration is uniform, since the point defect density surrounding the dislocations alters which lattice site is occupied by the silicon atoms and hence whether the silicon acts as a donor or acceptor. Neither the Hall nor the C-V methods (used with the VGF samples) would have sufficient spatial resolution to confirm this hypothesis. The dominant macroscale feature of the vertical boat grown n-type GaAs sample illustrated in FIGURE 8 is a roughly circular pattern of variations in PL intensity. They are caused by doping striations that often appear in bulk grown samples. A PL map is a rapid, non-destructive alternative to delineation etching for the detection of doping striations. In addition, the isolated dislocations found only at the centre of the sample can be identified easily. D2

Mapping of Semi-insulating GaAs

FIGURES 9, 10 and 11 compare semi-insulating GaAs prepared by ingot annealed liquid encapsulated Czochralski (LEC), wafer annealed LEC and VGF methods. All samples were commercial grade with resistivity in the range of 1 - 5 x 107 Qcm. Those chosen for illustration are typical of each class. The full wafer map of FIGURE 9 shows the four fold symmetry typical of Czochralski growth where, due to the interaction between thermal stress and the symmetry properties of the crystal, dislocations nucleate and propagate preferentially along certain crystal directions. The highest PL comes from regions of high dislocation density. This is somewhat counter-intuitive until it is realized that the PL signal is dominated by enhanced emission from the dislocation cell boundaries

FIGURE 8. Whole wafer and higher resolution centre area PL intensity maps of silicon doped n-type GaAs grown by the vertical boat method.

FIGURE 9. Whole wafer and higher resolution centre area PL intensity maps of undoped semi-insulating GaAs grown by LEC.

seen in the detailed map. The dislocation cores are not resolvable in this image, although similar features have been conclusively identified using a variety of techniques [24]. As in the case of ntype GaAs, it is widely believed that the material surrounding the dislocations appears relatively bright in PL due to the gettering of a non-radiative point defect, possibly As related, to the dislocation cores. Wafer annealing of as-grown material reduces the contrast within the cells and leads to overall homogenization of the luminescence and electrical properties [36]. The improved uniformity has been shown to correlate with reduced microscale resistivity variations and improved FET circuit performance [37]. A typical example is shown in FIGURE 10. The four fold symmetry is lost in the whole wafer map and the difference in PL intensity between the boundaries and the interiors of the dislocation cells is much reduced. The fabrication of GaAs circuits often involves thermal cycles with similar temperatures to those used in wafer annealing, and so it is not unexpected that an analogous improvement in microscale uniformity occurs [38]. FIGURE 11 shows a typical semi-insulating VGF wafer. The dislocations are few enough (<103 cm"3) that the cells may be detected on the whole wafer map. In such cases, millimeter scale resolution resistivity measurements show good correlation with the PL maps [38]. The cell boundaries of the VGF boundary are more complex than those in FIGURE 9 or 10. The well defined regions of intermediate PL brightness are found within the dislocation free cell interiors. They are likely to have an origin in common with the bubbles of micro-precipitates seen in annealed material [24]. Like the dislocations, the micro-precipitates getter the non-radiative recombination centres in their immediate vicinity and thereby increase PL intensity. Clearly the three types of semi-insulating material have different defect spatial distributions. Work

FIGURE 10. Whole wafer and higher resolution centre area PL intensity maps of wafer annealed undoped semiinsulating GaAs grown by LEC.

FIGURE 11. Whole wafer and higher resolution centre area PL intensity maps of undoped semi-insulating GaAs grown by the vertical gradient freeze method.

is on-going in several laboratories to elucidate what impact the defects may have on device performance. E

CONCLUSION

PL mapping has matured since its introduction over ten years ago. Room temperature PL mapping has emerged as a useful tool to support the development of improved semiconductor materials. The basic principles and practical pitfalls of RT PL mapping are now more thoroughly understood so that accurate inferences about material properties can be made from PL data. Nondestructive measurements of GaAs doping level can be made using either the PL intensity or spectral shape data: the preferred approach depends on the sample doping level. In addition, PL intensity mapping has yielded valuable insights into the origins of defects in substrates and epitaxial layers. Work is on-going to improve methods for extracting quantitative information concerning defect densities from PL intensity maps and to understand the implications for device performance. REFERENCES [1] [2] [3] [4] [5]

HJ. Hovel, D. Guidotti [ IEEE Trans. Electron Devices (USA) vol.ED-32 (1985) p.2331 ] A.T. Hunter [Appl Phys. Lett (USA) vol.47 (1985) p.715 ] W. Wettling, J. Windscheif [ Appl Phys. A (USA) vol.40 (1985) p.191 ] S.K. Krawczyk, M. Garrigues, H. Bouredoucen [ J. Appl Phys (USA) vol.60 (1986) p.392 ] J. Windscheif, W. Wettling [ 1987 Defect Recognition and Image Processing in IH-VCompounds, Ed. E.R.Weber (Elsevier, Amsterdam, Netherlands, 1987) p. 195 ]

[6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28]

[29] [30] [31] [32] [33] [34] [35] [36] [37] [38]

C.J.L. Moore, CJ. Miner [ J. Cryst. Growth (Netherlands) vol. 103 (1987) p.21 ] H. Ch. Alt, H. Mullerborn, G. Packeiser [ Semi-insulatingIU-VMaterials, Toronto, 1990, Eds A. Milnes, CJ. Miner (Adam Hilger, Bristol, England, 1990) p.309 ] H.J. Hovel [ Semicond. Sci. Technol. (UK) vol.7 (1992) p.Al ] T.W. Steiner, M.L.W. Thewalt [ Semicond. Sci. Technol. (UK) vol.7 (1992) p.A16 ] G.E. Carver [ Semicond. Sci. Technol. (UK) vol.7 (1992) p.A53 ] CJ. Miner, A. Harrison, R. Clayton [Semi-insulatingIU-VMaterials, Ixtapa, 1992, Eds. CJ. Miner, W. Ford, E.R. Weber (IOP, Bristol, England, 1992) p. 189 ] W. Jantz, M. Baeumler, Z.M. Wang, J. Windscheif [ Mater. Res. Soc. Symp. Proc. (USA) vol.325 (1994)p.4O9] W. Jantz, M. Baeumler, J. Forker, P. Hiesinger, Z.M. Wang [ Semi-insulating III-VMaterials, Warsaw, 1994 , Ed. M. Godlewski (World Scientific, Singapore, 1994) p.79 ] J. Forker et al [ Inst. Phys. Conf. Ser. (UK) vol. 145 (1995) p.37 ] R. Toda, M. Warashina, M. Tajima [ Mater. Sci. Forum (Switzerland) vol. 196-201 (1995) p. 1785] M. Tajima [ Inst. Phys. Conf. Ser. (UK) vol. 149 (1995) p.243 ] M. Gershenson [ Semicond. Semimet. vol.2 (Academic Press, New York, 1972) p.289 ] J. Pankove [ Optical Process in Semiconductors (Dover, USA, 1971) p. 107 ] D.A. Cusano [ Solid State Commun. (USA) vol.2 (1964) p.353 ] H.C. Casey, R.H. Kaiser [ J. Electrochem. Soc. (USA) vol. 114 (1967) p. 149 ] R.A. Smith [ Semiconductors (Cambridge University Press, Cambridge, England, 1978) p.264 ] CJ. Miner, D.G. Knight, B. Watt, C. Blaauw, N. Puetz [ IEEELEOS' 90 Conference Proceedings (1990)p.615] G.D. Knight, CJ. Miner, AJ. SpringThorpe [ Mater. Res. Soc. Symp. Proc. (USA) vol. 184 (1990) p.157] J.L. Weyher et al [ Semicond. Sci. Technol. (UK) vol.7 (1992) p.A45 ] BJ. Skromme, CJ. Sandroff, E. Yablonovitch, T. Gmitter [ Appl. Phys. Lett. (USA) vol.51 (1987) p.24] T. Saitoh, H. Hasegawa [ Mater. Sci. Forum (Switzerland) vols. 185-188 (1995) p.53 ] CS. Liu, J.F. Kauffinan [ Appl. Phys. Lett. (USA) vol.66 (1995) p.3504 ] S.K. Krawczyk, M. Garrigues, A. Khoukh, C. Lallemand, K. Schohe [ Semi-insulating IU-V Materials, Malmo, 1988, Eds G. Grossmann, L. Ledebo (Adam Hilger, Bristol, England, 1988) p.555 ] R. Toda, M. Tajima [ J. Cryst. Growth (Netherlands) vol. 103 (1987) p.28 ] R.A. Gottscho, BL. Preppemau, SJ. Pearton, AJ. Emerson, KP. Giapis [ J. Appl. Phys. (USA) vol.68 (1990) p.440] T. Suzuki, M. Ogawa [ Appl. Phys. Lett. (USA) vol.31 (1977) p.473 ] D. Guidotti, HJ. Hovel [ Appl. Phys. Lett. (USA) vol.53 (1988) p. 1411 ] H. Hess, E. Betzig [ Semiconductor Characterization, Eds W.M. Bullis, D.G. Seiler, A.C. Diebold (AIP Press, Woodbury, NY, USA, 1996) p 395 ] L. Samuleson, J. Lindahl, L. Montelius, M.-E. Pistol [ Inst. Phys. Conf Ser. (UK) vol. 135 (1994) P-51] JP. Fillard [Inst. Phys. Conf. Ser. (UK) vol. 149(1995)p. 195 ] M. Mori, G. Kano, T. Inoue, H. Shimakura, H. Yamamoto, O. Oda [ Semi-insulating III-V Materials, Toronto, 1990, Eds A. Milnes, CJ. Miner (Adam Hilger, Bristol, UK, 1990) p. 155 ] O. Oda et al [ Semicond. Sci. Technol. (UK) vol.7 (1992) p.A215 ] CJ. Miner, J.A. Zorzi, M. Young, K. Borg [J. Cryst. Growth (Netherlands) in press (1996) ]

CHAPTER 10 DEFECTS, DEEP LEVELS AND THEIR DETECTION 10.1 10.2 10.3 10.4 10.5 10.6 10.7 10.8 10.9

Structure of native defects in undoped SI LEC GaAs Energy levels and fundamental properties of the EL2 defect in GaAs Electrical techniques for the measurement of deep state properties Defect densities in melt-grown GaAs (a review) Deep states in as-grown epitaxial GaAs Defect energy levels in electron and neutron irradiated photon damaged and ion implanted GaAs Analysis of LEC GaAs by near-IR mapping Passivation of defects in GaAs by hydrogenation Transmission electron microscopy of GaAs

10.1 Structure of native defects in undoped SI LEC GaAs S. Yasuami August 1995

A

INTRODUCTION

Native defects are classified on the basis of their dimensionality: extended or linear defects such as dislocations are one-dimensional, micro-precipitates are three-dimensional and point defects are considered to be zero-dimensional. In undoped semi-insulating (SI) liquid-encapsulated Czochralski-grown (LEC) GaAs crystals, both dislocations and lineage are extended defects commonly observed. Microprecipitates decorating dislocations have been found. Anti-site, vacancy, and complex defects have been reported to be present. The EL2 defect is best known for its important role realizing the semi-insulating nature of GaAs, although its identity remains controversial. B

EXTENDED DEFECTS

Undoped SI LEC GaAs crystals contain a high density of dislocations. The majority of dislocations constitute bundles to form cellular walls towards the core of the crystal and lineage extending from the core towards the periphery. The cell wall is a complex agglomeration of dislocations with different directions and Burgers vectors and polygonizes the crystal. The lineage is composed of dislocations with similar directions and Burgers vectors and forms a low angle tilt boundary [I]. Dislocations are significant because they create so-called Cottrell atmospheres and are considered to act as sinks for point defects and redistribute them [2]. The non-uniformity of infrared (IR) absorption [3], luminescence efficiency of cathodoluminescence [4,5] and photoluminescence [6] have been ascribed to the redistribution of the point defects (non-radiative centres) associated with dislocations. C

MICRODEFECTS

Microprecipitates accompanying dislocations have been observed for Cr-doped [7] and undoped [8] crystals by transmission electron microscopy (TEM). The precipitates were identified by electron diffraction to be single crystals of elemental hexagonal arsenic [9]. As-rich matrix microprecipitates were observed for doped and undoped LEC and horizontal Bridgman (HB) crystals. Their size and density are 0.1 - 0.5 \xm and 107- 1011 cm"3, respectively. X-ray diffuse scattering analysis suggested that each microprecipitate is accompanied by point defects in the associated strain field [10]. Polycrystalline insular particles of about 1000 A in size were observed with a low density (108 109 cm"3) [11]. Amorphous precipitates with 20 to 30 A diameter have also been observed using the lattice imaging technique of TEM [12].

D

POINT DEFECTS

The stoichiometry-dependent [13] deep donor EL2 [14] defect is the most important of all, because its compensation of acceptor impurities, mainly carbon, makes the crystal semi-insulating. Vacancy-type defects are considered to be the origin of reverse contrast defects [15]. Dl

The EL2 Defect

Although several microscopic models have been proposed for the defect, no consensus has yet been reached. In the following subsections, these models with the techniques used for deriving the conclusions are briefly reviewed. Recently two particular models have become popular: the isolated arsenic antisite (AsGa) and the antisite-interstitial pair defect along the [111] direction (As0J1-As1) [16]. There are both pros and cons of the measurements and analyses supporting these results, mainly the observation of the zero-phonon line (ZPL) and optically detected electron nuclear double resonance (ODENDOR) and optically detected electron spin resonance (ODESR). Refer to more comprehensive review articles for details [16-18] Dl.1

Isolated AsGa model

Dl.1.1 ESR The simplest model for the EL2 defect is the isolated arsenic antisite: AsGa. The ESR spectrum for AsGa+ was identified with the EL2 defect [19,20].

Dl.1.2 ZPL The defect that gives rise to the zero-phonon line (ZPL) at 1.039 eV in the intracentre IR absorption spectra of the EL2 defect was assumed to be the neutral charge state of the arsenic antisite: AsGa° [21]. Kaminska et al [22] proposed that the EL2 defect has tetrahedral symmetry (Td) and is As03, based on the splitting of ZPL due to the A1-T2 transition under a uniaxial stress and a magnetic field. This experiment was later confirmed by Bergman et al [23] and by Trautman et al [24]. Dl. 1.3 Photoluminescence Nissen et al studied the fine structure in the 0.61 eV photoluminescence band from the EL2 defect under a uniaxial stress and a magnetic field. There was no deviation from the Td symmetry; supporting the isolated AsGa model [25]. Dl.1.4 Other issues on this model Theoretical calculations supporting this model have been reported [26-28]. The isolated AsGa defect can exhibit metastability, the AsGa becoming an arsenic interstitial Asf leaving a gallium vacancy V03 behind, and making a complex As1 -VGa in the [111] direction. However, the energy barrier is lower than the experimental value by a factor of 2 [28]. The metastable VGa was detected by positron annihilation and its density linearly depended on the EL2 density [29]. These results are in good agreement with the isolated AsGa model.

On the other hand, a question remains as to whether the ZPL experiment would be sensitive enough to detect the C3v component of the potential caused by the nearby As * Baraff[30,31] calculated the effects on the stress splitting on the assumption that the EL2 defect has the structure determined from the ODENDOR spectra. The C3v component of the potential could cause a splitting of the ZPL that was far larger than the detectivity limit. However, none was detected. Another challenge to the ZPL experiment was made by Skowronski [32]. He suggested that it detects the hydrogenic effective mass states but not the stress splitting. The experimentally obtained final-state energy of the ZPL-assumed transition A1-T2 is immediately below the L band minimum. Von Bardeleben and Bourgoin [33] pointed out that the energy state assumed to be due to the ZPL exactly traced that of the L band minimum under pressure in the experiment carried out by Baj and Dreszer [34]. Lannoo gave further support in line with this scheme [35]. Additional experimental evidence supporting the hydrogenic interpretation was given; another defect showed the same ZPL and phonon replica [36]. D 1.2

ASG11-AS1 pair

model

D 1.2.1 Ballistic phonon transmission Culbertson et al [37] measured the transmission of ballistic phonons through a sample containing EL2 defects, which was cycled between the ground and metastable state. The EL2 transition from the ground to metastable state reduced the scattering from defects with C3v symmetry, implying that this was the symmetry of the EL2 ground state. This is consistent with the pair model. However, the conversion from the ground to the metastable state generates holes and there is no guarantee that the reduction in scattering was actually caused by the EL2 defect and not by other defects that captured the holes. Dl.2.2 Photoquenching One of the prominent properties of the EL2 defect is photoquenching. When the crystal is irradiated with sub-bandgap light, the defect becomes inactive and no longer compensates acceptors. The original property of the defect recovers after being heated above 140 K [16]. Levinson and Kafalas [38] studied transient photocapacitance using polarized light under uniaxial stress. The observed transient behaviour was understood to consist of two processes. They interpreted this as: (i) Td-symmetry defect undergoes photoionization transition and (ii) a pair defect is involved in the following metastable transition. Dl.2.3 Emission rates Dobaczewski studied emission rates from the EL2 and other defects varying the electric field strength for different directions. A unique field-strength dependence of the EL2 defect for the [111] direction was ascribed to the axial, namely pair, nature of the defect [39]. Dl.2.4 ODENDOR The energy levels and optical properties in the diamagnetic state of the EL2 defect were detected by 0DEND0R and ODESR [40,41]. A spectrum with Td-symmetry was obtained for the

unquenchable defect (AsGa), while one with C3v-symmetry was obtained for the quenchable form (AsGa-ASi) of the defect. All results obtained by ESR and DLTS on the EL2 defect can be understood if the stable state of the defect corresponds to As1 in the second-neighbour position of AsGa, while the metastable state corresponds to As1 in the first-neighbour position [42]. Dl.3

The complex models

The isolated AsGa and the EL2 defect have the same or very similar ESR spectra but different photoquenching behaviours [43]. The EL2 defect is generated and the EL2 defect is modified by dislocations and clusters introduced by neutron transmutation doping (NTD). Assuming that both defects are identical, their photoquenching and thermal annealing behaviour were compared [44]. The EL2 defect was identified as a complex defect composed of three point defects. Quenching phenomena in tunnelling assisted hopping conduction below 125 K were observed for the NTD crystal. The possibility of EL2 defects being formed from AsGa, As1, V^ and VGa was suggested [45]. D1.4

Other models

The ESR spectrum was identified with a superposition of spectra, and only AsGa-V03-V^ was believed to be the probable configuration for the EL2 defect [46]. Two peaks were observed in the photoquenching data of the normal state of the EL2 defect and a complex structure in the photoinduced-recovery spectra coincided with the calculated energy levels of VAs. The EL2 defect was hence considered to be associated with a complex defect involving VM [47,48]. In the stable state of the EL2 defect, both vacancies are on one side of AsGa, while the metastable state corresponds to an AsGa that separates the two vacancies [49]. The other models are: the EL2 defect is a complex of the AsGa-AsGa type, or of the AsGa-ASj type [50]. It was also suggested that the EL2 defect may represent complex As atom clusters [51]. D2

The Ga^ Antisite Defect

This defect has not yet been identified by the ESR technique. Nevertheless, experimental data already available on electro-physical and optical properties of samples prepared by different methods indicate the presence of a doubly charged acceptor, most probably Ga^ [52,53]. D3

Vacancy-type Defect

The defect assumed to be V^ was detected by positron annihilation measurements in undoped n-type crystals containing an EL2 concentration of 2 x 1016 cm "3 [54]. According to optically induced positron annihilation measurements, its transition to the negative state is consistent with 'reverse contrast' IR absorption. The absorption contrast changes when the sample is cooled down lower than 100 K with illumination of light a few tens of millivolts lower than the bandgap energy [15]. The defect is considered to act as the dominant non-radiative recombination centre in bulk SI GaAs [55]. This defect undergoes photoquenching and thermally recovers similarly to

the EL2 defect. It has been suggested by one positron annihilation group that for impurity concentrations less than 1 x 10 17 cm"3, divacancies dominate; for those larger than 4.5 x 1017 cm"3 the monovacancy complex Asoa-Voa is observed, while between them, both YM and VGa are estimated to be present [56]. E

CONCLUSION

The EL2 defect is most important among those defects found in undoped SI LEC crystals, since it makes the crystal semi-insulating. A great number of experimental and theoretical studies have been published concerned with its identification. There are two models that have gained credibility to date. However, they are incompatible. It is high time that the controversy is reconciled and the structure of the EL2 defect is established [16].

FURTHER READING For more comprehensive reviews, see refs [16-18,57,58]. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] II1] [12] [13] [14] [15] [16] [17] [18] [19] [20]

J. Matsui, T. Kitano, T. Kamejima, T. Ishikawa [ Gallium Arsenide and Related Compounds, Biarritz, France, 26-28 Sept 1984 (Adam Hilger, Bristol, 1985) p. 101 ] A.H. Cottrell [ Dislocations and Plastic Flow in Crystals (Clarendon Press, Oxford, 1953) ch.II p.53] M R Brozel, I. Grant, R.M. Ware, DJ. Stirland, M.S. Skolnick [ J Appl Phys. (USA) vol.56 (1984) p. 1109] A.K. Chin, A.R. Von Meida, R Caruso [ J Electrochem. Soc. (USA) vol. 129 (1982) p.2386 ] T. Kamejima, F. Shimura, Y. Matsumoto, H. Watanabe J. Matsui [ Jpn. J. Appl Phys. 2 (Japan) vol.21 (1982) p.L721] K. Kitahara, M. Ozeki, A. Shibatomi [Jpn. J Appl Phys. 1 (Japan) vol.23 (1984) p.207 ] A.G. Cullis, P.D. Augustus, DJ. Stirland [ J Appl Phys. (USA) vol.51 (1980) p.2556 ] B-T. Lee, R Gronsky, E.D. Bourret [J Cryst. Growth (Netherlands) vol.96 (1989) p.333 ] DJ. Stirland, P.D. Augustus, M.R Brozel, EJ. Foulkes [ Semi-Insulating IU-VMaterials (Shiva, Nantwich, 1984) p.91] L.A. Charniy, A.N. Morozov, K.D. Scherbachov, V.T. Bublik [ J Cryst. Growth (Netherlands) vol.116 (1992) p.369] J-P.Cornier, M. Duseaux, J-P.Chevalier [ Inst. Phys. Conf. Ser. (UK) no.74 (1985) p.95 ] F.A. Ponce, F-C. Wang, R Hiskes [ Semi-InsulatingIU-VMaterials (Shiva, Nantwich, 1984) p.68 ] D.E. Holmes, RT. Chen, K.R Elliot, CG. Kirkpatrick [ Appl Phys. Lett. (USA) vol.40 (1982) p.46] G.M. Martin, A. Mittonneau, A. Mircea [ Electron. Lett. (UK) vol. 13 (1977) p. 191 ] M.S. Skolnick, LJ. Reed, A.D. Pitt [Appl Phys. Lett. (USA) vol.44 (1984) p.447 ] G.A. Baraff [ Semi-Insulating IU-VMaterials (Inst. Phys. Publ., Bristol, 1993) p. 11 ] M.R Brozel, C. Corbel [ J Phys. IV1 Colloq. (France) vol.5 (Jan. 1995) p.63 ] U. Kaufmann [ in FestkorperproblemeU/Advances in Solid State Physics vol.29 Ed. U. Rossler (Vieweg, Braunschwieg, 1989) p. 183 ] K. Elliot, RT. Chen, S.G. Greenbaum, RJ. Wagner [ Appl. Phys. Lett. (USA) vol.44 (1984) p.907 ] M. Baeumler, U. Kaufmann, J. Windscheif [ Appl Phys. Lett. (USA) vol.46 (1985) p.781]

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N. Tsukada, T. Kikuta, K. Ishida [ Jpn. J. Appl. Phys. 2 (Japan) vol.24 (1985) p.L689] M. Kaminska, M. Skowronski, W. Kuszko [ Phys. Rev. Lett. (USA) vol.55 (1985) p.2204 ] K. Bergman, P.Omling, L. Samuelson, H.G. Grimmeiss [ Semi-InsulatingIH-VMaterials (Adam Hilger, Bristol, 1988) p.397 ] P.Trautman, J.P.Walczak, J.M. Baranowski [ Phys. Rev. B (USA) vol.41 (1990) p.3074 ] M.K. Nissen, A. Villemaire, M.L.W. Thewalt [ Phys. Rev. B (USA) vol.67 (1991) p. 112 ] J. Dabrowski, M. Scheffler [ Phys. Rev. Lett. (USA) vol.60 (1988) p.2183 ] D.J. Chadi, KJ. Chang [ Phys. Rev. Lett. (USA) vol.60 (1988) p.2187 ] E. Kaxiras, K.C. Pandey [ Phys. Rev. B (USA) vol.40 (1989) p.8020 ] K. Saarinen, S. Kuisma, P.Hautojarvi, C. Corbel, C. Le Berre [ Phys. Rev. B (USA) vol.49 (1994) p.8005 ] G.A. Baraff [ Phys. Rev. Lett. (USA) vol.62 (1989) p.2156 ] G.A. Baraff [ Phys. Rev. B (USA) vol.40 (1989) p. 1030 ] M. Skowronski [ Defects in Electronic Materials (The Materials Research Society, Pittsburgh, 1988)p.4O5] HJ. von Bardeleben, J.C. Bourgoin [ Impurities, Defect and Diffusion in Semiconductors (The Materials Research Society, Pittsburgh, 1990) p.865 ] M. Baj, P.Dreszer [Mater. Sci. Forum (Switzerland) vol.38-41 (1989) p. 101 ] M. Lannoo, C. Delerue, G. Allan [ Defects in Semiconductors 16 (Trans Tech Publications, Zurich, 1992) p.865 ] J-M. Spaeth, K. Krambrock [ 20th Int. Conf. on the Phys. ofSemicond. vol.1 (World Scientific, Singapore, 1990)p.491] J.C. Culbertson, U. Strom, S.A. Wolf [ Phys. Rev. B (USA) vol.36 (1987) p.2962 ] M. Levinson, J. A. Kafalas [ Phys. Rev. B (USA) vol.35 (1987) p.9383 ] L. Dobaczewski [ Materials Science Forum 38-41 (Trans Tech Publications, Aedermannsdorf,1988) p. 113 ] B.K. Meyer, J-M. Spaeth [ J. Phys. C (UK) vol. 18 (1985) p.L99 ] B.K. Meyer, D.M. Hofinan, J.R. Niklas, J-M. Spaeth [ Phys. Rev. B (USA) vol.36 (1987) p. 1332 ] H. J. Von Bardeleben, D. Stievenard, D. Deresmes, A. Huber, J. C. Bourgoin [ Phys. Rev. B (USA) vol.34 (1986) p.7192] M.O. Manasreh, P.F. McDonald, S.A. Kivlighn, J.T. Minton, B.C. Covington [ Solid State Commun. (USA) vol.65 (1988) p. 1267 ] M.O. Manasreh, D.W. Fischer [ Phys. Rev. B (USA) vol.39 (1989) p.3239 ] M. Satoh, K. Kuriyama [ Phys. Rev. B (USA) vol.40 (1989) p.3473 ] G. Wang, Y. Zou, S. Benakki, A. Goltzene, C. Schwab [ J. Appl. Phys. (USA) vol.63 (1988) p.2595] M.O. Manasreh, D.W. Fischer [ Phys. Rev. B (USA) vol.39 (1989) p. 13001 ] M.O. Manasreh, D.W. Fischer [ Phys. Rev. B (USA) vol.40 (1989) p. 11756 ] J.F. Wagner, J.A. Van Vechten [ Phys. Rev. B (USA) vol.35 (1987) p.2330 ] A.N. Georgobiani, LM. Tiginyanu [ Sov. Phys.-Semicond. (USA) vol.22 (1988) p. 1 ] T. Dcoma, Y. Mochizuki [ Jpn. J. Appl. Phys. 2 (Japan) vol.24 (1985) p.L935 ] Ph.W. Yu, W.C. Mitchel, M.G. Mier, S.S. Li, W.L. Wang [ Appl. Phys.. Lett. (USA) vol.41 (1982) p.532] K.R. Elliott [ Appl. Phys. Lett. (USA) vol.42 (1983) p.274 ] G. Dlubek, A. Dlubek, R. Krause, O. Bruemmer, K. Friedland, R. Rentzsch [ Phys. Status Solidi A (Germany) vol.106 (1988) p.419 ] S. Tuzemen, L. Breivik, M. R. Brozel [ Semicond. Sci. Technol. (UK) vol.7 (1992) p.A36-40 ] S. Dannefaer, D. Kerr [J. Appl. Phys. (USA) vol.60 (1986) p.591 ] M.O. Manasreh, D.W. Fischer, W. C. Mitchell [ Phys. Status Solidi B (Germany) vol. 154 (1989)

[58]

J.C. Bourgoin, H. J. von Bardeleben [ J. Appl. Phys. (USA) vol.64 (1988) p.R65 ]

[24] [25] [26] [27] [28] [29] [30] [31] [32] [33] [34] [35] [36] [37] [38] [39] [40] [41] [42] [43] [44] [45] [46] [47] [48] [49] [50] [51] [52] [53] [54]

P-Il]

10.2 Energy levels and fundamental properties of the EL2 defect in GaAs J.M. Baranowski and P. Trautman July 1995

A

INTRODUCTION

EL2 is the dominant native point defect in bulk GaAs crystals grown by both liquid-encapsulated Czochralski (LEC) and horizontal Bridgman techniques. The interest in EL2 comes from its technical importance and from purely scientific reasons. The presence of EL2 in GaAs crystals makes possible the fabrication of semi-insulating (SI) GaAs material, that is crystals of high resistivity («108 Q cm) with a free carrier concentration of about 106 cm"3 at room temperature. Such a low carrier concentration cannot be achieved solely by purification. The semi-insulating property of undoped GaAs results from the pinning of the Fermi level by the midgap donor level of EL2. Semi-insulating GaAs serves as a substrate material for high-speed digital and microwave monolithic integrated circuits. The scientific interest in EL2 comes mainly from its rare property of having an excited metastable state. Illumination of the crystal at low temperature induces this transformation when the energy levels, optical absorption band and other properties characteristic in the normal state disappear. We will discuss EL2 in the framework of the normal (ground) state being the isolated arsenicantisite AsGa and the metastable state being the tightly bound gallium vacancy arsenic interstitial V^-ASi defect pair of trigonal C3v symmetry. B

OCCURRENCE OF THE EL2 DEFECT IN GaAs

The name EL2, electron level number 2, was given [1] in 1977 to a deep trap having characteristic electron emission parameters Ea (activation energy) and aa (emission cross section) determined by deep level transient spectroscopy (DLTS) experiments. These parameters for EL2 are Ea = 0.82 eV and oa« 10-13cm2. The capture cross section for electrons On was directly measured [2] in a large temperature range (50 - 270 K) and is given by On (cm2) = 5 x 10"19 + 6 x 10"15 x exp (-0.066)/kT

(1)

Thus it is strongly temperature dependent and, in order to determine the true value of the EL2 thermal ionization energy E1, one has to correct the DLTS activation energy by subtracting the activation energy of On [3]. This gives E1= 0.82-0.07 eV = 0.75 eV, i.e., a value equal to the ionization energy obtained from the Hall effect on SI GaAs [4]. EL2 is commonly present at a concentration of about 1016 cm"3 in GaAs crystals grown by both LEC and Bridgman methods. However, EL2 was directly detected in 1966 [5] at concentrations around 1016 cm"3 in measurements of capacitance of Schottky barriers produced on n-type GaAs.

Semi-insulating GaAs is now commonly obtained by the LEC technique without intentional doping [6]. The concentration of EL2 determined by measurements of optical absorption in crystals grown from pyrolytic boron nitride crucibles, starting with a charge of high purity Ga and As, and using dry B2O3 encapsulant, was found to increase from 5 x 1015 to 1.7 x 1016 cm"3 as the As atom fraction in the melt increases from 0.48 to 0.52 [6]. A similar dependence was observed in crystals grown by the horizontal Bridgman technique; the concentration of EL2 determined by DLTS increasing from 1016 to 3 x 1016 cm"3 as the As source temperature increases from 613 to 619°C [7]. The EL2 defect is also present in GaAs layers grown by vapour phase epitaxy (VPE) [8-10]. Its concentration was observed to be roughly proportional to the As/Ga ratio in the vapour phase during growth ranging from 2 x 1013 cm"3 for an As/Ga ratio of 1/3 to 4 x 1014 cm"3 for an As/Ga ratio of 12. The fact that EL2 is a native defect is confirmed by a recent observation of EL2 in GaAs layers grown by molecular beam epitaxy (MBE) at low substrate temperatures (LT GaAs). Measurements of near infrared absorption demonstrated the presence of EL2 in concentrations of about 1020 cm"3 in layers grown at 200 0 C [H]. The concentration of EL2 in these layers is many orders of magnitude higher than the concentration of any chemical impurity. C

OPTICAL ABSORPTION DUE TO THE EL2 DEFECT EV GaAs

ABSORPTION COEFFICIENT

(cm"1)

The near infrared absorption spectrum due to EL2 was reported by Martin [12]. Two years later, Kaminska et al [13] found that the central band of this absorption begins with fine structure consisting of a no-phonon line at 1.039 eV (8378 cm"1) followed by phonon replicas separated by about 11 meV (FIGURES 1 and 2).

PHOTON ENERGY

(eV)

FIGURE 1. Near infrared absorption spectra at 10 K in semi-insulating GaAs containing the EL2 defect in a concentration of about 2 x 1016 cm'3. Curve a was recorded after cooling the sample in the dark. Curve b was obtained after illumination of the sample with white light for 10 minutes.

The absorption due to EL2 begins at 0.8 eV with photoionization transitions from the midgap donor level of EL2 to the F valley of the conduction band (see curve a in FIGURE 1). At 1.1 eV, there is onset of transitions to the L valleys of the conduction band resulting in a rapid increase of the slope. A band of intracentre absorption beginning at 1.039 eV with the no-phonon line, and extending to about 1.3 eV, is superimposed on this photoionization background. The most characteristic feature of the optical absorption is that it can be quenched by illumination of the crystal at low temperature [12]. The absorption disappears completely after illumination with white light for about 10 min (see curve b in FIGURE 1). FIGURES 1 and 2 are commonly observed in n-type and undoped semi-insulating GaAs and spectra are attributed to the neutral charge state of EL2. When EL2 is partially quenched, the whole absorption band is quenched, by the same factor. The magnitude of the no-phonon line is proportional to the magnitude of the entire band. The magnitudes of curve a and the no-phonon line are proportional to the concentration of EL2 determined by DLTS in n-type GaAs [12,14].

ABSORPTION COEFFICIENT

(cm- 1 )

The EL2 defect, being a double donor, may exist also in two other charge states: singly positively charged (+) and doubly positively charged (++) states. The optical properties of EL2+ and EL2++ are less well known than those of EL2°. For EL2++, only transitions from the valence band to the empty +/-H- level of EL2 are expected to occur and these were observed by Skowronski [15]. For EL2+, three kinds of optical transitions can occur: (i) transitions from the valence band to the empty 0/+ level of EL2, (ii) transitions from the filled +/++ level of EL2 to the conduction band, and (iii) intracentre A 1 -T 2 transitions. The magnitudes of the relative contributions of these transitions as observed by magnetic circular dichroism (MCD) are a matter of controversy [16-

PHOTON ENERGY

(eV)

FIGURE 2. Fine structure of the intracentre absorption due to the EL2 defect in GaAs. The fine structure consists of a no-phonon line at 1.039 eV and its phonon replicas. This spectrum was measured in the same sample as that shown in FIGUREl.

20]. The rare occurrence of GaAs crystals with most EL2 in the EL2+ state is one of the reasons that the optical properties of EL2+ are poorly understood. D

MICROSCOPIC IDENTIFICATION OF THE EL2 DEFECT IN GaAs

Dl

Discovery of the AsGa Antisite Defect by EPR Spectroscopy

In 1980, the arsenic antisite was discovered in as-grown semi-insulating GaAs by submillimetre electron paramagnetic resonance (EPR) spectroscopy [21]. The corresponding spectrum consists of four, nearly equally spaced lines and is isotropic. The four lines were interpreted as originating from the hyperfine interaction of an unpaired electron with a central nucleus of spin 1 = 3/2. The spectrum can be described by the spin Hamiltonian H = g^ B B.S + AS.I

(2)

The experimentally determined [21] parameters are: g = 2.04 ± 0.01 and A = 0.090 ± 0.001 cm"1. As the arsenic isotope 75As is 100% abundant and has a nuclear spin I = 3/2 the observed spectrum may be due to the positively charged arsenic antisite AsGa+. This identification was confirmed by comparison, with appropriate scaling, of the hyperfine interaction parameter A with that for the phosphorus antisite PGa in GaP, where super-hyperfine interactions with ligands is resolved allowing unambiguous identification of the defect. The identification of the AsGa with EL2 was postulated on the basis of a correlation between the intensity of the EPR spectrum and the concentration of carbon acceptors [22]. It was based on the following argument: if EL2 and As03 were not the same defect, then the EPR signal would be independent of the changes in residual acceptor content, since these would change the occupation of the EL2 defect, the compensating centre, but not the occupation of AsGa. The identity of EL2 with As^ was further supported by the similarity between the enhancement and quenching spectra of the AsGa+ EPR signal and the photoionization spectra of EL2 [23]. The presence of the AsGa+ EPR signal in undoped semi-insulating GaAs and its ability to be enhanced and quenched by illumination of the crystal indicates that the 0/+ level of AsGa coincides, within a few meV, with the 0/+ midgap level of EL2. Indeed, if this were not the case, then all the AsGa antisites would be in the same charge state and either the entire EPR signal or no signal at all would be observed before illumination. The coincidence of the levels of EL2 and AsGa is a strong argument for the identity of the two defects. D2

Piezospectroscopic Determination of the Symmetry of EL2

The large width of the EPR lines due to AsGa+ did not allow the unambiguous determination that EL2 is due to the isolated arsenic antisite. The symmetry of EL2 in the normal state was determined by splitting the no-phonon line (also termed zero phonon line (ZPL)) of EL2 under uniaxial stress applied in high symmetry crystallographic directions [24]. These results indicated that the no-phonon line is due to the A 1 -T 2 transition in a centre of tetrahedral Td symmetry. Therefore, EL2 has to be a simple isolated point defect, of not lower than tetrahedral symmetry. Combining this determination with previous experimental and theoretical data [24] identified the EL2 defect with the isolated arsenic antisite.

ZPL ENERGY (cm" 1 )

STRESS

(MPa)

FIGURE 3. The splitting of the 8378 cm1 zero phonon line (ZPL) of the EL2 intracentre absorption under [110] stress. The symbol E[hkl] indicates that incident light was polarized in the [hkl] crystallographic direction. The points represent experimental results. The solid lines are theoretical fits to the experimental points.

On the other hand, on the basis of EPR and DLTS investigations performed on GaAs samples subjected to electron irradiation and heat treatments, von Bardeleben et al [25,26] suggested that EL2 is a complex of an AsGa antisite and an arsenic interstitial As1. This suggestion found strong experimental support in optically detected electron nuclear double resonance (ODENDOR) [27,28]. These conclusions [24] were questioned by Figielski and Wosinski [29]. Therefore, an independent uniaxial stress experiment on the no-phonon line of EL2 was performed [3O]. The 8378 cm"1 no-phonon line of EL2 was observed to split into 2, 2, and 3 components under a uniaxial stress applied along the [100], [111], and [110] crystallographic directions, respectively [30]. Exactly one stress split component of the no-phonon line was observed in each polarization including three nonequivalent polarizations in the case of [110] stress. This indicates unequivocally that the no-phonon line of EL2 is due to the A 1 -T 2 (or A 2 -T 1 ) electric dipole transition in a defect of tetrahedral Td symmetry. Uniaxial stress experiments cannot distinguish between A 1 -T 2 and A 2 -T 1 transitions, but we will refer to the A 1 -T 2 transition as it is in agreement with theoretical predictions for the AsGa defect. The results of [30] are essentially in agreement with those of [24], the main difference being that the viewing directions for [110] stress in Figure 1 of [24] are interchanged as was pointed out in [29,31]. The excited T2 state, transitions to which give rise to the no-phonon line, is triply orbitally degenerate. This is clearly seen from the splitting pattern of the no-phonon line under [110] stress, in which case exactly one component is observed for light polarized along each of the three principal axes ([HO], [001], and [110]) of the crystal stressed in the [110] direction (see FIGURE 3). Threefold orbital degeneracy does not exist in defects of lower than Td symmetry. Therefore,

the EL2 defect has to have unperturbed tetrahedral Td symmetry. Clear confirmation of the Td symmetry of the EL2 defect was provided by Nissen et al [32] from measurements of a sharp luminescence line at 0.7028 eV under uniaxial stress. This line remained unresolved under [100], [110], and [111] stress up to 500 MPa [32], the only effect being a hydrostatic shift of 30 meV/GPa. The luminescence is due to transitions of an electron from an A1 symmetry hydrogenic state bound under the F point minimum of the conduction band to the midgap A1 state of EL2. The lack of any splitting of the 0.7028 eV line under uniaxial stress strongly supports the isolated AsGa model of EL2, since any complex of lower than Td symmetry should exhibit a splitting due to removal of the orientational degeneracy. D3

Electronic Structure of the AsGa Antisite Defect

Native point defects in GaAs, and in particular the arsenic antisite, have been intensively studied by theoretical methods [33-38]. In a simple defect-molecule model, the arsenic antisite can be described as a defect formed by filling a perfect undistorted gallium vacancy with an arsenic atom. Qualitatively, the electronic structure of the defect results from the interaction of the atomic s and p orbitals of the central As atom with the localized states of the vacancy. There are four dangling orbitals U1, U2, U3, and U4 of the nearest neighbours pointing towards the vacant site. In the Td symmetry of the undistorted vacancy, one may construct four molecular orbitals from them: the lower in energy belongs to the ax representation, the three other linear combinations lying higher in energy having t2 symmetry (see FIGURE 4). Hybridization of these orbitals with the s and p states of the central arsenic atom results in

VQa

A s G a As atom

FIGURE 4. Single-particle energies and interaction of a neutral gallium vacancy in GaAs and an As atom that results in an As08 antisite defect. The s and p orbitals of the As atom hybridize with the Si1 and X2 states of the vacancy forming bonding (b.) and antibonding (a.) states of the arsenic antisite defect.

formation of two ax and two t2 states of the arsenic antisite. The lower in energy Z1 and t2 states have bonding character and are completely filled with electrons forming sp3 bonds between the central As atom and its ligands. The ax antibonding state lies approximately in the middle of the energy gap and can hold zero, one or two electrons depending on the position of the Fermi level. It holds two electrons in the neutral charge state of the defect. The antibonding t2 state is resonant with the conduction band and is empty for the ground state of the defect in any charge state. Therefore, the arsenic antisite defect is expected to act as a double donor. The positions of the donor levels were determined for the arsenic antisite by photo-EPR experiments [39] and for EL2 employing DLTS and photocapacitance spectroscopies [40-43] (see FIGURE 5). The consistent level positions determined by the different experimental methods supports the identity of AsGa with the EL2 defect. The two donor levels located at Ev + 0.54 eV (+/++) and Ec - 0.75 eV (0/+) both disappear after the transformation of EL2 into the metastable state. The electronic structure of AsGa is in agreement with the optical absorption spectrum of EL2 in the neutral charge state. The photoionization transitions are due to excitation of an electron from the midgap at state to the conduction band. The intracentre transitions result from the excitation of an electron from the midgap a2 state to the resonant t2 state. The symmetry of the electronic states of AsGa is in agreement with the symmetry of electronic states determined by piezospectroscopic experiments, supporting the identification of the EL2 defect with the AsGa antisite.

Conduction band

Valence band FIGURE 5. Energy levels of the EL2 defect, or the arsenic antisite, determined by photo-EPR, DLTS, and photocapacitance spectroscopies. The As08 defect being a double donor has three (0, +, -H-) charge states in the energy gap-

D4

ODENDOR Investigations of AsGa

There is a long-standing controversy between the above attribution of EL2 to the isolated AsGa defect and results of optically detected electron nuclear double resonance (ODENDOR)

measurements, which attribute EL2 to the arsenic antisite and arsenic interstitial complex As03-As1 of trigonal C3v symmetry. The investigations of EL2 by ODENDOR became possible after the discovery of the magnetic circular dichroism (MCD) signal due to AsGa+ [16]. A large number of ODENDOR lines due to AsGa+ were detected in the frequency range 10-140 MHz [44-46]. The angular dependence of 0DEND0R lines is analyzed by guessing an appropriate spin Hamiltonian and fitting its eigenvalues to the experimental line positions. Initially, the presence of both a fully symmetrical AsGa defect and a perturbed AsGa defect in as-grown SI GaAs was deduced by analysis of the 0DEND0R data [45]. Improved data and interpretation indicated that the entire 0DEND0R spectrum is due to the AS 03 -AS 1 pair of trigonal C3v symmetry [20,27,28,46]. It was determined that the arsenic interstitial is singly positively charged and located at about two bond lengths from AsGa in the [111] antibonding direction. The As1 interstitial cannot be neutral since it would be paramagnetic, inconsistent with the observed EPR spectrum. These conclusions are in conflict with those reported in Section D2. However, recently Spaeth et al admitted that it is not possible to say anything definite about the location and symmetry of the As interstitial because of the low intensity of the 0DEND0R spectra and the possibility that the super-hyperfine interactions of the As interstitial could be confused with those of a displaced nearest As ligand [47]. Therefore, the presence of an As interstitial in the structure of the EL2 defect remains disputable. It should be stated that the weakly bound As0J1-As1 defect pair model is not consistent with the high thermal stability of the EL2 defect. All properties of EL2 remain unchanged even after annealing the crystal to 8500C followed by rapid or slow cooling to room temperature, although thermal destruction of EL2 was reported to take place above 100O0C [48,49]. D5

Observation of AsGa by Scanning Tunnelling Microscopy

The arsenic antisite As^ defect was recently observed by scanning tunnelling microscopy (STM) in GaAs layers grown by MBE at low temperature (LT) [5O]. Spectroscopy of these defects revealed an intense band of midgap states centred at about 0.5 eV above the maximum of the valence band maximum. Four types of defect images were observed and interpreted as arising from a single type of defect located at different depths below the (110) cleavage surface. The concentration of As^ defects observed by STM is about 1020 cm"3, the same as the concentration of EL2 defects observed in LT GaAs [H]. The STM image of the As^ defect consists of a central spot due to the defect core and two satellites located about 15 A from the core along [lT2]and [T 12] surface directions. These satellites most likely arise from tails of the antisite wave function. All the arsenic antisites observed by STM have the (110) symmetry plane, the only symmetry operation present on the (110) surface, and it is concluded that the defects have tetrahedral Td symmetry in the bulk. Therefore, STM in LT GaAs confirms once more that EL2 has unperturbed tetrahedral symmetry. E

METASTABILITY OF THE EL2 DEFECT IN GaAs

El

Photocapacitance Studies of the Metastability of EL2

Metastability is commonly considered to be a fingerprint of the EL2 defect, but not all metastable phenomena in GaAs are directly related to the EL2 defect. Therefore, it requires careful consideration to decide if an observed metastable effect gives information on EL2 itself or is due

to other defects. This is particularly important in semi-insulating GaAs because in this material the transformation of EL2 into the metastable state changes the position of the Fermi level. Persistent quenching of photoconductivity of high-resistivity GaAs, evidently due to the metastability of EL2, was reported in 1976 [51]. At the low temperature of 82 K, the crystal was converted from a high photosensitivity state with n-type photoconductivity to a low photosensitivity state with p-type photoconductivity by exposure to photons in the 1.0 - 1.3 eV energy range. Recovery occurred during heating of the crystal at about 120 K. The observed range of photon energies inducing quenching and the temperature of recovery are characteristic oftheEL2 defect. Systematic studies of the metastability of EL2 began in 1977 after photocapacitance quenching of Schottky barriers produced on n-type GaAs had been discovered [52,53]. At low temperature («77 K) and for photon energies around 1.15 eV, a peculiar behaviour of transient photocapacitance was observed [52,53] shown schematically by curve a in FIGURE 6. The photocapacitance transient begins with filling the EL2 traps with electrons by forward biasing the Schottky diode then reverse bias is applied in the dark. After switching the light on, a fast increase of capacitance is observed followed by a decrease to a final value that is nearly equal to the initial value. The fall time is about 100 times longer than the rise time. A usual photocapacitance transient is shown by curve b in FIGURE 6.

TIME (s) FIGURE 6. Schematic photocapacitance transients due to the EL2 level in a Schottky barrier diode produced on n-type GaAs. The light is switched on at time t=0. Curves of type a correspond to transients measured at low temperature (T<110 K) for photon energy lying in the band of EL2 photoquenching (1 eV120 K) for any photon energy inducing photoionization of EL2 and at low temperatures for photon energies lying outside the photoquenching band. The two curves are computer plots of solutions of appropriate differential equations and closely resemble the experimental results.

The peculiar photocapacitance transient shown by curve a in FIGURE 6 is explained by the transformation of EL2 from the normal photo-active state into a passive metastable state EL2* in which the defect has the same charge state as in the normal state. EL2* must have the same charge state as EL2°, since the value of capacitance after completion of photoquenching, when EL2° is completely converted to EL2*, is the same as the initial value. The normal photo-active state can be recovered by heating the sample. There are two distinct recovery processes. The first one is observed in the absence of free electrons, that is when the diode is reverse polarized. In this case, the recovery rate is dependent only on the temperature and is given by EQN (3) [54] r± (s"1) = 8.6 x io 11 exp (-0.34 /kT) (3) The recovery rate is greatly accelerated when free electrons are present, that is when forward bias is applied. In this case, it is proportional to the concentration n of free electrons and is given by EQN (4) [54] rn (s"1) = 1.9 x 10"14 cm2 x nVth x exp (-0.107 /kT) (4) where w± is the thermal velocity of electrons in the conduction band, equal to «2 x 107 cm/s at 77 K. This recovery is also thermally activated, but with a smaller activation energy of 0.107 eV. The total recovery rate may be written as the sum of r^ and rn, r (s"1) = 8.6 x 1011 exp (-0.34 /kT) + 1.9 x 10"14 x nVth x exp (-0.107 /kT) E2

(5)

Optical Studies of the Metastability of EL2

In optical measurements the metastability of EL2 manifests itself as a quenching of the near infrared absorption (see FIGURE 1). The absorption can be recovered by heating the crystal (see FIGURE 7) [55]. The recovery of EL2 from the metastable state occurs at «50 K in n-type GaAs and at «130 K in semi-insulating GaAs. Analysis of the rate of EL2 recovery as a function of temperature by means of Arrhenius plots for several n-type and semi-insulating GaAs crystals allows the recovery rate to be expressed by EQN (6) [55] r (s"1) = 1.7 x io 12 exp (-0.36 /kT) + 1.6 x 10"9 x n x exp (-0.085 /kT)

(6)

Similar results were obtained by Fuchs and Dischler [56]. EQN (6) obtained from the recovery of absorption is consistent with EQN (5) obtained from the recovery of photocapacitance. Differences between them are in part due to experimental errors, e.g. the calibration of temperature, and in part due to different experimental conditions. In photocapacitance experiments, measurements are performed in a thin layer of semiconductor in which a gradient of free electron concentration is present and faster recovery is measured, i.e., the recovery is investigated at higher temperature than in optical experiments. A very important property of the photoquenching process is the spectral dependence of the optical cross section a* for the bleaching transitions. Vincent, Bois and Chantre [57] determined from photocapacitance measurements that a* is in the form of a Gaussian band extending between about 1.0 and 1.3 eV. This covers the same energy range as the intracentre absorption band observed in optical absorption measurements.

ABSORPTION

n—type

TEMPERATURE (K) FIGURE 7. Thermal recovery of absorption due to the EL2 defect measured at a photon energy of 1.2 eV for n-type (n« 1016 cm-3) and semi-insulating (SI) GaAs during slow («2 K/min) heating of the crystal after previous quenching of the absorption with white light at 10 K.

Metastability of EL2 was also observed as a quenching of the luminescence bands peaked at «0.62 eV and «0.68 eV [58-60]. The excitation spectrum of quenching of luminescence extends over the same 1.0 - 1.3 eV energy range and the luminescence returns to its initial value after heating the crystal above * 120 K. E3

Studies of the Metastable Properties of EL2 under Hydrostatic Pressure

Detailed studies of the EL2 defect and its metastability under hydrostatic pressure were made by Baj and Dreszer [61-65]. The most important of their findings is that EL2* has an acceptor level EL2*"/0 lying «15 meV above the bottom of the conduction band and that this level moves down under hydrostatic pressure entering the bandgap at about 200 MPa [63]. The properties of EL2 are significantly altered when a moderate hydrostatic pressure (-300 MPa) is applied. Then EL2 can be effectively photorecovered and in n-type GaAs a broad absorption band due to the negative charge state of EL2* appears. These properties are unlike those of EL2 under atmospheric pressure. These changes may be the basis for the apparently different properties of EL2 in GaAs irradiated with high energy particles and in GaAs grown at low temperature, since the local strain fields present in these materials may induce changes in the properties of EL2 similar to those due to the application of hydrostatic pressure. E4

Theoretical Model of the Metastability of EL2

When the identification of EL2 with the isolated AsGa was proposed, this defect was considered

to be too simple to account for the metastable properties [66,67]. The turning-point in the search for a model of EL2 metastability was the suggestion of Baranowski et al [68] that the transformation of EL2 into the metastable state is connected with a displacement of the arsenic antisite to an interstitial position. Thus, the transformation more closely resembles the photodissociation of AsGa into a gallium vacancy and an arsenic interstitial As Oa -V Oa - As1 than the lattice relaxation usually accompanying the change of electronic state of a defect. This new idea found strong support in theoretical calculations performed by Dabrowski and Scheffler [69,70], Chadi and Chang [71] and Kaxiras and Pandey [72]. These calculations showed that the arsenic antisite defect indeed has an excited metastable configuration in which the total energy of the defect is only about 0.2 eV higher than in the ground state. During the transformation into the metastable state the central As atom breaks one of its four As-As bonds and is displaced by approximately 1.3 A along the antibonding [111] direction to an interstitial position (see FIGURE 8). Therefore, the metastable state of AsGa is the tightly bound gallium vacancy-arsenic interstitial V04-As1 defect pair. As a result of the transformation, two As atoms change their coordination from fourfold to threefold. The stability of the metastable configuration is a consequence of the ability of group V atoms to be threefold bonded as in the case of arsine AsH3 or various forms of solid arsenic. In consequence of the transformation the symmetry OfAs03 is reduced from tetrahedral Td to trigonal C3v. The metastable state V03-As1 is fourfold orientationally degenerate because the transformation into the metastable state may take place in four ways by breaking any one of the four bonds between the central As atom and its ligands. The transformation of AsOa into the metastable state is initiated by a Jahn-Teller distortion following the excitation of AsOa to the electronically degenerate T2 state. From this distorted excited state the defect may reach the metastable state by tunnelling. The theoretical values of the energy barrier between the metastable and the normal configuration of AsGa strongly depend on the particular method of calculation and lie in a rather broad range from 0.16 to 0.92 eV, i.e., from less than half to more than twice the experimental value of 0.36 eV.

As Ga Td

V Ga Asj

C3v

FIGURE 8. Local atomic structure of the undistorted arsenic antisite As03 and the gallium vacancy arsenic interstitial V03-As1 defect pair considered to be the excited metastable state OfAs08 responsible for the metastability of EL2.

E5

Experimental Determination of the Symmetry of EL2 in the Metastable State

When the transformation of AsGa into V01-As1 was proposed as the mechanism of EL2 metastability, the only experimental support was that it successfully explained the metastability of the isolated arsenic antisite. The metastable state of EL2 could not be studied by conventional methods because of the lack of absorption, EPR, and luminescence signals. The process of thermally activated recovery of EL2 from its metastable state is one of the few experimentally observed phenomena in which the metastable state is involved. Study of the thermal recovery of optical absorption under uniaxial stress applied along the [111] and [100] crystallographic directions was performed by Trautman and Baranowski [73,74]. The idea of the experiment was that if EL2* has a symmetry lower than tetrahedral, then EL2* centres having different orientations with respect to the stress may recover at different temperatures, i.e., the recovery step (see FIGURE 7) may split under the uniaxial stress. From the dependence of splitting on the direction of the stress in the crystal the symmetry of the metastable state can be determined. FIGURE 9 shows the thermal recovery of EL2 absorption in SI GaAs under [111] and [100] stresses of selected values from 0 to 600 MPa. The recovery step starting at about 120 K evidently splits into two components for large values of [111] stress but no splitting is observed under [100] stress. This difference is more evident after computation of numerical derivatives of the recovery curves with respect to temperature, as is shown in FIGURE 9.

ABSORPTION COEFFICIENT (cm-0

(cm-i) ABSORPTION COEFFICIENT

stress

stress

TEMPERATURE (K)

TEMPERATURE (K)

stress

da/dT

(cm-VK)

da/dT (cm-VK)

stress

TEMPERATURE

(K)

TEMPERATURE (K)

FIGURE 9. Thermal recovery of the absorption due to the EL2 defect in semi-insulating GaAs measured during slow heating of the crystal under uniaxial stress applied in the [ 111 ] and [ 100] directions (top) and computed first derivatives of these curves with respect to temperature (bottom). Curves for different values of the stress are shifted vertically for clarity. The abrupt increase of absorption at the end of the curve measured under [111] stress of 600 MPa was caused by a fracture of the sample. The recovery step splits into two components under [111] stress but no splitting is observed under [100] stress; this indicates that the metastable state of EL2 has trigonal C3v symmetry.

The observed splitting of the recovery under [111] stress is clear evidence for an orientational degeneracy of EL2*, i.e., EL2 * has lower than tetrahedral (noncubic) symmetry. If the symmetry of EL2* was tetrahedral, then all of them would be equivalent with respect to the stress and no splitting of the recovery step would be possible. The EL2* defects are equivalent with respect to the [100] stress because there is no splitting or broadening of the recovery step up to 600 MPa. Among defects having orientational degeneracy only those having trigonal C3v symmetry do not split under the [100] stress. The recovery step splits into two components under [111] stress as expected for a defect of C3v symmetry. Therefore, EL2* has trigonal C3v symmetry. Trautman and Baranowski [73,74] have also studied the effect of different methods of excitation of EL2 to the metastable state on the splitting of the recovery step in semi-insulating GaAs. It was observed that the same EL2 defect in the metastable state can be oriented in different directions depending on the method of transformation of EL2 into the metastable state. Namely, EL2* defects can be oriented mostly along the [111] direction by excitation with light polarized parallel to this direction and they can be oriented completely aslant the [111] direction by excitation under a [111] stress of 500 MPa. It is much easier to understand the ability of EL2* defects to assume different orientations assuming that normal EL2 has perfect tetrahedral symmetry because the orientation of EL2* is expected to be determined by the polarization of light only. Therefore, the investigation of recovery of absorption due to EL2 under uniaxial stress [73,74] not only allowed the determination of the C3v symmetry of EL2* but also supports the identification of the EL2 defect with a defect of unperturbed tetrahedral symmetry.

E6

Observation of a Vacancy in the Structure of EL2 in the Metastable State by Positron Annihilation Spectroscopy

It was observed by positron annihilation spectroscopy that vacancies appear in undoped semiinsulating GaAs after the transformation of EL2 from its normal into the metastable state [75,76]. In a positron annihilation experiment, the lifetime of positrons in the investigated crystal is measured. Positrons being positively charged tend to avoid positively charged atomic cores and are trapped by vacancies where a lattice atom is missing. When a positron is trapped by a vacancy, the average electron density experienced by the positron is lower than that when it is delocalized in the bulk of the crystal and, consequently, the lifetime of the positron is increased. Therefore, positron lifetime spectroscopy is a technique sensitive to the presence of vacancies. When EL2 is in the normal state, the positron lifetime in semi-insulating GaAs is observed to be nearly temperature independent in the range 25 - 300 K and equal to the value of 230 - 235 ps reported for positrons delocalized in semi-insulating GaAs [75,76]. This means that before illumination practically no positron trapping at vacancies is observed. After illumination that induces transformation of EL2 into the metastable state, the positron lifetime at 25 K is observed to have increased by 2 to 6 ps depending on the sample [75,76]. This increase is persistent and the temperature of annealing that results in the return of the lifetime to its initial value is exactly the same as that of recovery of the absorption due to the normal state of EL2. The range of photon energies inducing the increase of the positron lifetime is the same as that inducing the transformation of EL2 into the metastable state. Altogether these experimental findings indicate that vacancies are created in semi-insulating GaAs as the result of the transformation of EL2 into the metastable state. Therefore, a vacancy must be involved in the atomic structure of EL2*. F

CONCLUSION

Much progress has been made in the understanding of the electrical, optical, and structural properties of the EL2 defect in GaAs since 1977 [52,53]. It is now well established that EL2 is due to the arsenic antisite AsGa and the nature of its metastability is well understood. Recent experimental findings indicate that the metastable state of EL2 has trigonal C3v symmetry and that it contains a lattice vacancy in its atomic structure. These findings confirm the theoretical model of metastability of the arsenic antisite, or EL2, which attributes the metastability to the optically inducible transformation OfAs04 into a tightly bound gallium-vacancy-arsenic-interstitial V021-As1 defect pair. The metastable properties of EL2 are changed when hydrostatic or uniaxial pressure is present. This effect may explain why the properties of EL2 in GaAs irradiated with high energy particles or in GaAs grown by MBE at low substrate temperatures (LT GaAs) are different from those in bulk as-grown GaAs, because in irradiated and in LT GaAs large strain fields are present. The isolated AsGa antisite is the best model of the EL2 defect in undoped asgrown GaAs. REFERENCES [1] [2]

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10.3 Electrical techniques for the measurement of deep state properties Previous Page A.R. Peaker October 1996

A

INTRODUCTION

The purpose of this Datareview is to provide an outline of the techniques which are commonly used to characterise deep defect states in terms of their electrical properties. A general overview and an outline description of each method is given with information on its principal limitations and references from which more detailed information can be obtained. Techniques can be broadly divided into space charge methods and 'bulk' methods. This is an important distinction in relation to a technique's capabilities. In general, space charge methods have much greater sensitivity but have the disadvantage that the measured parameters are determined in the depletion field, a factor which may result in significantly different values being observed for some properties (e.g., reduced thermal activation energies due to the Poole-Frenkel effect). In contrast, 'bulk' techniques (such as Hall measurements) provide parameter values in the neutral (low field) region of the material but for these techniques to be applicable, the concentration of the deep state must be sufficient to have a measurable effect on the position of the Fermi level at some accessible temperature. In order to achieve a deep state dependent Fermi level, prior to the widespread use of space charge techniques, it was common practice to prepare samples for deep state measurements by compensating the shallow donors or acceptors. This was sometimes done by diffusion of a known species of deep state (copper, for example) or by irradiation. Inevitably, the subsequent analysis was difficult and no data observed by these 'compensation' methods is included in the review in this chapter. In 1966 Williams [1] published the first paper describing space charge measurements of deep states in GaAs. In this work he used an electrolytic Schottky barrier. Most of the work on deep state properties reported in the subsequent sections of this book has been undertaken on p-n junctions or metal-semiconductor barriers. All techniques for the measurement of the electric parameter of deep states can be regarded as having three stages. These are: 1. 2. 3.

The occupancy of the state is set The occupancy is perturbed The change in occupancy is measured directly or indirectly

It is the variants on these three stages which distinguish the methods of measurement. For example, in deep level transient spectroscopy (DLTS) the occupancy is set by applying zero bias to the normally reverse biased depletion region. The shift of the Fermi level will tend to result in states becoming occupied by majority carriers. The diode is then restored to reverse bias and at a sufficiently high temperature the majority carriers are thermally emitted from the deep state. This sequence provides a transient perturbation of the occupancy which is monitored by observing the small signal capacitance of the diode and using a simple analogue signal processing technique

to evaluate the resultant capacitance transient. Almost all the data reported in Datareviews on defect states in this book have been obtained using techniques which result from the use of thermal emission to characterise the state. There are, however, avast number of variants. In 1970, Sah [2] compiled a very comprehensive review of space charge techniques for deep level measurements. An extended and updated compilation of techniques by Blood and Orton appeared in 1992 [3]. Reference can be made to these publications for a more detailed analysis than it has been possible to include in this Datareview. B

OCCUPANCY

FIGURE 1 shows an overview of some of the principal charge exchange mechanisms involving deep states. The arrows denote the direction of electron transfer while the symbol C represents the capture rate and e the emission rate. The subscript (p) indicates a hole process and (n) an electron process. (A) represents carrier generation, (B) electron trapping and e-emission, (C) hole trapping and re-emission and (D) recombination.

FIGURE 1. Principal charge exchange mechanisms associated with deep states.

The symbols en and ep represent electron and hole emission rates with units of s"1 while Cn and cp represent the electron and hole capture coefficients with units of cm3 s"1. This follows the original notation of Sah et al [3] but it should be emphasised that some authors use Cn and cp to represent capture rates which have units of s'1 and in our notation these are equal to ncn. We use the symbol Cn to represent this quantity, i.e., Cn = ncn. At equilibrium taking nT as the number of traps occupied by electrons, then the number of empty traps is NT - nT. Consequently, (cn + e p )(N T -n T ) = (en + cpn)nT

(1)

This gives a steady state value for the occupation factor fT for the level, defined by fT = nT/NT: cn +e f

T = e

n

+

C

= P P

E +

C

n"

(2) +

%

Cn and cp are often written in terms of capture cross sections On and op which have units of cm2

Cn = ° n V t h

(3)

where v± is the thermal velocity of electrons. EQN (2) is the basis of all deep level measurements and is true irrespective of whether the emission processes are stimulated optically or thermally, although it is normal practice to distinguish between optical and thermal processes with the appropriate superscript, e° or e\ Essentially we force a change in fT by changing the trap environment thermally, optically or simply by expanding or collapsing the depletion region so as to shift the Fermi level at the trap location. In the most commonly used case we try to establish an initial condition of fT = 1 and then observe the small signal capacitance change of the depletion region AC as fT - 0. In general this change will be a time dependent exponential function so that for the emission of electrons in n-type material or holes in p-type (majority carrier emission) ACd) = A C 0 - 1 0 ( I - e * )

(4)

and for the emission of holes in n-type and the emission of electrons in p-type (minority carrier emission) AC(t) = AC(total) e "T

(5)

The time constant x provides information which we can use to investigate the properties of the deep state. The next section considers this in the cases of most interest to us: thermal emission and carrier capture. We then describe the commonly used measurement techniques. C

THERMAL EMISSION

By far the most common measurements of the deep state parameters involve using thermal energy to change the occupancy of the centre. If we consider the case where a defect is fully electron occupied and is emitting its electrons to the conduction band, then the thermal emission rate is proportional to a Boltzmann factor: ^ = AnCXpC-E1ZkT)

(6)

An is a property of the trap-semiconductor system and we will refer to Ea as the activation energy. This form of relationship means that the emission rate changes very rapidly with temperature. If the log of en, or more usually for reasons discussed below, (log en)/T2 is plotted against 1/T, Ea can be obtained from the gradient of the line. The value Ea is sometimes referred to as the trap energy, meaning its energetic displacement from the band. This is not strictly correct and although the measurement is very convenient experimentally, its detailed physical interpretation is fraught with difficulty. In reality our plot of (log en)/T2 vs 1/T on its own merely gives us a convenient fingerprint of the defect. This is not necessarily unique to a particular species but by comparison with published emission measurements taken from known defects in the same semiconductor, this fingerprint can often provide an 'identification' when combined with other information. An example of such an Arrhenius plot is shown in FIGURE 2.

FIGURE 2. Arrhenius plot of the electron emission rate of a defect in silicon. The slope of the plot is 381 meV (Ea) and the intercept is 1.06 x 106 s"1 k'2 at T = °°. This is equal to the prefactor Ax.

EQN (6) relating emission rate to temperature is often written in the form en = gvmNconexp(-Ea/kT)

(7)

where Nc is the density of states and g the degeneracy of the electron state, the other terms are as previously defined. This relationship is derived by applying the principle of detailed balance to the emission and capture processes. The fundamental justification for this is somewhat dubious as the derivation can only be applied for near equilibrium conditions for a single reversible process. If, for example, the capture proceeds through an intermediate state or states, EQN (7) is erroneous. There is some evidence that such capture processes do occur. However, an indisputable deficiency with EQN (7) is a mathematical one because we have assumed that the terms Ea and On are independent of temperature. We know that the vth Nc product varies as T2 and this has already been alluded to in terms of a correction to the determination of Ea by plotting log en/T2 against 1/T. More importantly both a and E x (the real trap energy depth) are, in general, functions of temperature. As will be discussed later, it is possible to measure a directly and determine its temperature dependence. Hence if a increases with increasing temperature with an activation energy, E0 then ED (the thermal energy of the state) is taken to be E a - E o where E a is the slope of the T2 corrected Arrhenius plot. This makes the tacit assumption that any dependency of a on T can be described by an exponential relationship. If this dependence is due to an energetic barrier which can be surmounted thermally this is accurate. However, it is far from universally true but because of the usually narrow range of temperature over which o(T) is measured, a power law or other dependency can be approximated (knowingly or otherwise) by the exponential function. This very widely used calculation evades the question of what Ea really is in physical terms because the true energetic separation of the deep state from the band is likely to increase with decreasing temperature. Consequently, it is more appropriate to consider the emission as a thermodynamic process under constant pressure. For the case of electron emission from a state we can describe the process in terms of the change in the Gibbs free energy of ionisation AGn, in the enthalpy AHn and the entropy ASn. In general:

ET - AGn = AHn - TASn

(8)

en = gVfc N0 an exp (ASn/k) exp (-AHnZkT)

(9)

so that at a specific temperature:

ENERGY (meV)

If none of the prefactors is temperature dependent then the slope of the Arrhenius plot is AHn the enthalpy change. However, a much more likely scenario is shown in FIGURE 3 where the Gibbs free energy is changing with temperature. For the simplest case where the capture cross section is invariant with temperature or we have made due allowance for its dependency when interpreting the Arrhenius plot (by subtracting or adding its activation energy to Ea) we will measure AH. This enthalpy change is equal to the energy given by the T = O intercept of the tangent to the AG curve taken at a point equal to the measurement temperature. This is the value that an Arrhenius plot gives us and in order to derive the 'true' value of ET we need to know AS.

true value

TEMPERATURE (K)

FIGURE 3. A representation of the relationships between the energies measured using thermal emission techniques and the thermodynamic parameters for a hypothetical deep state. E0 is the energy derived from the Arrhenius plot, T1n the mean of the measurement temperatures, AG the change in Gibbs free energy, AS the change in entropy and AH the change in enthalpy.

D

CAPTURE CROSS SECTIONS . . . REAL AND APPARENT

In practice, the physical interpretation of the energy we measure is of concern primarily when comparing energies obtained from optical measurements with the thermal energies and most importantly in the derivation of capture cross sections from the emission data. EQN (9) might lead us to believe that we can obtain a value of On from the Arrhenius plot of the emission data using the intercept at 1/T = O. Such values are variously called Gn (x = °°), the apparent cross section or the intercept cross section. The Gibbs free energy is generally a function of temperature and for such a calculation to be meaningful we must know ET (T1n) quite precisely, i.e., it is not sufficient to know Ea and E00. A small error in ET(Tm) will result in a massive error in Gn because ET is within the exponential term. As Ea > E 1 the value of Gn derived from an Arrhenius plot using Ea is generally larger than the real value, indeed, in some cases, there are discrepancies of four orders of magnitude between the 'intercept' cross section and the directly measured value. As the capture cross section is the most important parameter in assessing the

Shockley-Read-Hall recombination and hence the effect of defects on minority carrier lifetime it is rather important not to rely on the value of cross sections derived from the Arrhenius plot. The process of a defect capturing a carrier is conceptually more straightforward than emission, although the details of the process such as its temperature dependence, the screening of long range potentials and excitonic effects make the simplistic model of a geometric cross section look somewhat inadequate. We merely assume that the defect has a sphere of influence and ascribe it a cross sectional area when considering it in two dimensions. We assume that any carrier intercepting that sphere will be 'captured'. If we take an empty defect state and allow it to capture electrons we can observe the capture process by monitoring the change in occupancy of the state and so derive a capture cross-section. This process will proceed with a rate Cn given by: Cn = o n V t h N T n

(10)

This gives rise to an exponential transient with a time constant equal to Cn"1 N x . There are several experimental techniques available for this measurement, two of which are widely used: DLTS with variable filling pulse width which measures majority carrier cross sections and MCTS which measures minority carrier cross-sections. These techniques will be described in Section E. The measurement of capture cross sections enables the Shockley-Hall-Read kinetic model to be applied and although the measurement is normally carried out in an environment in which only one type of carrier is present, the results have been very successfully used to calculate minority carrier lifetimes. The recombination process in FIGURE 1 occurs in bulk material (not the depletion region) and we can see that two processes are involved. These are (a) the capture of a hole and (b) the capture of an electron. EQN (10) demonstrates that the capture rate depends on the number of carriers present. Electrons have been considered in EQN (10) but the case for holes is identical if the hole parameters are substituted for the electron values. In the bulk of p-type material the hole concentration is high and so unless op « On the hole capture rate will be faster than electron capture, i.e., the state will tend to be hole occupied. Consequently, as we need both processes to occur for recombination, the rate limiting step will be minority carrier capture, in this case, electron capture. Consequently for low excitation levels in p-type material we can write a very simple expression for the minority carrier lifetime:

^ n = KV 1 HN x ] 1

(11)

A similar expression can be written for n-type material depending on the hole cross section and concentration.

E

CAPACITANCE OR CURRENT

Although the change in occupancy of the deep state in GaAs can be detected in many ways this is almost always done by a measurement of depletion capacitance or by the detection of the current produced by the emitted carriers. In a few techniques, changes in conductance are measured.

For the space charge methods, some generalisations can be made in contrasting current and capacitance techniques. In general, capacitance techniques are more favoured because in the case of a one sided junction (Schottky barriers, p+n and n+p junctions) a distinction can be drawn between hole and electron emission simply by observing the sign of the capacitance change. There are, however, other important distinctions. The magnitude of capacitance charge induced by a unit point charge depends on its spatial location within the depletion region. Charge on the junction plane p-n or Schottky has no effect on the capacitance while charge introduced at the depletion region edge has a maximum effect. The weighting of the AC calculation for intermediate locations is linearly graded. In the case of current measurements, a released majority carrier makes an equal contribution to the current irrespective of the location of the trap from which it originated. Perhaps the most confusing feature of current techniques is that fundamentally the peak height for a given trap concentration is a function of the emission rate: this is not the case for capacitance. Apart from the need to take this into account when calculating trap concentrations, it also causes a shift of the peak (in temperature scanned experiments) to higher temperatures. This is often a source of misinterpretation in the literature particularly in comparison of trap 'fingerprints' obtained by different methods. F

TRANSIENT SPACE CHARGE METHODS . . . DLTS AND RELATED TECHNIQUES

DLTS was first described by Lang in 1974 [4]. Although conceptually similar methods were in use prior to this date, DLTS had a simple elegance in the form of its output which endeared it to many users. The technique produces a sequence of peaks when a function relating to time dependence of capacitance change is plotted against temperature. At its simplest, each peak represents a deep state. In DLTS the traps are filled with carriers by applying a zero or forward bias pulse to the device and then observing the transient after the device is switched into reverse bias, usually between 2 and 10 V. The pulse sequence and resultant capacitance transient are shown in FIGURE 4. zero bias

time (t)

minority carrier transient

majority carrier trap

increasing temperature (T)

FIGURE 4. Capacitance transients resulting from applying a voltage pulse to a diode containing deep states.

In FIGURE 4 it can be seen from the voltage plot at the top of the diagram that the device is switched into forward bias for a short period of time (typically 1 ms). This collapses the depletion region and allows the trap to fill from the majority carrier population. When the voltage is switched into reverse bias again the captured majority carriers have reduced the net space charge and the capacitance is lower. If the temperature is sufficiently high the majority carriers are emitted and the net space charge increases which we detect as an increase in capacitance. As predicted from EQN (9) the rate at which this happens is strongly dependent on temperature as shown in the lower part of the diagram. If we apply a 'fill' pulse which forward biases the p-n junction to such an extent that a substantial amount of current flows, minority carriers are injected. Their concentration will always be less than that of the majority carriers but if the minority carrier capture cross-section is comparable with or larger than the majority carrier crosssection then we will capture minority carriers at the defect and may observe a transient of the opposite sense as shown by the dotted line. In DLTS the exponential change of capacitance, which is the emission transient, is sampled at two points in time, X1 and t2, which are typically in the range 0.1 ms - 1 s. The amplitude difference between the two samples provides the DLTS signal which is then plotted against the sample temperature. The reasoning behind this is shown in the lower part of FIGURE 4. The times X1 and ^ are the times after the end of the reverse bias pulse. If the exponential is very slow (as at the bottom of the figure) almost no difference between the values of the samples exists. If the exponential is very fast (as at the top) again no difference exists. At intermediate values of time constant C(tl) - C(t2) is non-zero and for given values of tj and t2 there is a time constant that will give a maximum output. Because of the Boltzmann factor the time constant changes rapidly with temperature so if we alter the sample temperature the plot of C(tl) - C(t2) against T will peak when the emission rate matches the system time constant, i.e. when %l = x = (I1 -12) In(Vt2)

(12)

Many other ways of analysing the exponential transients have been proposed, some of which are discussed later in this Datareview. However, one method which is widely used is to replace the boxcar with a lock-in [5]. Traps with different emission rates will produce peaks at different temperatures and so if we have several different trap species present, each will produce a characteristic peak. Indeed this is one of the main advantages of this method, namely that it is 'spectroscopic' i.e. it can scan a range of trap depths. An example of this is shown in FIGURE 5 where the DLTS scan reveals several different species of deep state. To produce an activation plot the temperature must be scanned at different values of tx and X2, or different 'rate windows' in DLTS jargon. Such a family of peaks is shown in FIGURE 5. The Arrhenius plot shown in FIGURE 2 is derived from the positions of the peaks at the various rate windows. The variation of emission rate with temperature which the family of DLTS curves reveal is shown on a narrow range by the width of a single peak. It is, of course, the variation of emission rate with temperature which gives rise to the peak width and indeed to its characteristic shape and so in principle the peak width above could be used to derive the activation energy. However the peak width is not exclusively a function of the trap parameters. It depends on the DLTS system, in particular the X7)X1 ratio, the sampling width etc and so these must be taken into account if the activation energy is to be derived from the width. In general the method is not favoured due to the narrow temperature range sampled and the uncertainties associated with the degree of broadening due to physical effects which may make

FIGURE 5. A family of DLTS peaks taken at six different 'rate windows' i.e. settings of ^ and I1. These have been chosen to give a peak at emission rates in the range 20-1000 s 1 . The highest rate window always appears at the highest temperature. Note the width of these near ideal peaks and the characteristic shape, i.e. the high temperature edge is steeper than that on the low temperature side. These are the raw data from which the Arrhenius plot shown in FIGURE 2 has been produced.

the emission transient non-exponential, such as field dependence and high trap densities. It is not only the position of the peaks that provides information; their magnitude is directly related to the amplitude of the experimental emission transient and so contains the information necessary to extract the concentration of the deep state. There are, however, a number of situations which occur which can introduce significant inaccuracies. By far the most common is the use of inappropriate bias conditions. If the filling pulse is too short or of inadequate magnitude, not all the traps will fill and so the observed capacitance change is reduced, underestimating the trap concentration. Rather less obvious is the effect of using small values of reverse bias. This is often done in order to reduce the effect of the electric field in the depletion region which tends to increase the emission rate and hence distort the 'signature' of the defect. This is particularly severe for donors in heavily doped material. If a low reverse bias is used the transition region between the space charge region and the bulk is a significant part of the depletion region and so must be taken into account when calculating the concentration. The effect is temperature dependent and results in the high temperature peaks being larger than those at lower temperatures [6]. Errors in the computation of carrier concentration also occur for the case where the trap concentration is high i.e. N x > IND - NA 1/10. This means that the approximation given in EQN (2) must be replaced by a more exact relationship. Even more significant is the fact that the depletion width changes during the emission transient. For the case where both the shallow dopant concentration and the trap concentration are uniform, allowance can be made for this analytically. However, a much more general solution is to modify the DLTS system so that the capacitance is maintained constant during the transient by a feedback network and the variation in voltage necessary to do this is monitored and analysed. The method is known as constant

capacitance DLTS [6]. The limiting case of high trap concentration is that of semi-insulating material. In this case manipulation of the Fermi level cannot be used to change the initial occupancy of the deep states. A widely used method of characterising deep states in such material is to apply a pulse of above bandgap light to the sample. This perturbs the equilibrium occupancy of the deep states and the emission of carriers is observable as a current transient. It is not possible to determine trap concentrations using this technique nor is it possible to distinguish between hole and electron emission. The technique and its variants are discussed in detail by Look [7] and are referred to by different groups as PICTS (photo-induced current transient spectroscopy), PITS (photoinduced transient spectroscopy) and OTCS (optical transient current spectroscopy). The task of obtaining meaningful results from samples with non-uniform doping and deep state population gradients is probably one of the most difficult of all deep level measurements. The deep state population can be profiled by holding the filling pulse constant and varying the reverse bias or by holding the reverse bias constant and changing the amplitude of the fill pulse. Both are quite simple to do experimentally, the problems arising in the data analysis. Simple calculations do not account for the Debye tail and it was shown many years ago that this resulted in very misleading results [8,9]. Calculations taking full account of the Debye tail are tedious but analytical approximations have now enabled this to be done on desk top computers with good accuracy for the general case of variable NT and IND - NA I. This is of crucial importance for the case of the assessment of damage near steep ion implantation profiles. For the case where lN D - NA1 is substantially constant a hardware solution exists referred to as double DLTS [1O]. DDLTS subtracts the DLTS signals obtained from consecutive transients with slightly different filling pulses, essentially subtracting the Debye tail effects. The problem of profiling deep states through a band discontinuity (i.e. a heterojunction) has not as yet been solved in a rigorous way, although solutions exist for specific cases. One of the major problems in relation to DLTS is the lack of resolution of the technique. The peaks tend to be very broad and, as has been mentioned previously, this is a major factor in ascribing definitive identities to thermal emission fingerprints. The origin of the broadening is manyfold. Lang's box car technique is extremely crude in signal processing terms. In 1978, Hodgart [11] proposed using a somewhat more complex signal analysis method based on a weighting function. He showed that much greater selectivity could be obtained. This treatment was considerably extended in 1981 by Crowell and Alipanahi [12], who presented an analytical approach based on filter theory. Their detailed treatment showed quite clearly that selectivity carried with it a substantial noise penalty. They presented details of several schemes in which the trade-off between selectivity and noise could be optimised. In 1987, the subject was revisited by Nolte and Haller [13] on the basis of attempting to achieve the ultimate in energy resolution. Implementation of these systems has proved to be particularly difficult. Issues related to noise have already been mentioned but from a mathematical viewpoint, perhaps the most difficult issue to be resolved being the fact that the zero base line of the exponential decay is not known. Recently significant advances have been made in the practical implementation of high resolution systems by employing more advanced techniques. Because this work was initially based on the inverse Laplace transform, they have become known generically (and somewhat erroneously) as 'Laplace transform DLTS' [14]. Truly remarkable improvements in resolution have been achieved in the study of many materials, including GaAs [15]. Dobaczewski et al [14] point out that the signal processing issue is only one of a number of factors in achieving reduced line widths

of the DLTS signal. The Laplace DLTS methods are isothermal techniques, that is they analyse the range of rates at a specific temperature and provide a spectral plot of a processed capacitance signal against emission rate rather than against temperature. In order to produce an Arrhenius plot, the experiment is repeated at a number of temperatures. This eliminates the line broadening due to the shift of parameters with temperature which they consider to be of comparable significance to the signal processing broadening in conventional DLTS. T=370K

amplitude (arb. units)

EL2 in GaAs

emission rate (s y) FIGURE 6. Laplace DLTS spectra of the defect EL2 taken on a range of GaAs samples. In conventional DLTS the defect is represented by a single peak [14].

All the methods considered so far have effected the change of occupancy by shifting the Fermi level to achieve majority carrier occupancy and then examined the majority carrier emission. It is quite feasible to undertake similar experiments but filling with minority carriers. Lang [4] describes such a method in his early papers on DLTS and in this case, minority carrier occupancy was achieved by passing a forward current through a p-n junction. Minority carriers were injected into the bulk and if traps had a minority carrier cross section comparable with the majority cross section, the occupancy would be different from the case of a zero bias filling pulse. It is also possible to perturb the occupancy optically and a number of techniques have been used to this end. Optical methods can be used to excite carriers across the gap and create minority carriers. In this case, the wavelengths of the exciting light must be slightly above bandgap. Alternatively, sub bandgap light can be used to directly change he occupancy of the state according to its optical cross section. The first application of the generation of minority carriers by above bandgap light as a technique for manipulating the occupancy of deep states was described by Hamilton et al [16] and was developed into minority carrier transient spectroscopy (MCTS) by Brunwin et al [17]. This is a very convenient method for indirect gap semiconductors but requires great care in its application to GaAs because of the steepness of the absorption edge which does, of course, change with temperature. The change in occupancy resulting from sub-bandgap optical excitation was first described by Mitonneau et al and is referred to as optical DLTS (ODLTS) [18]. G

CAPTURE CROSS-SECTION

In the DLTS technique we fill the majority carrier traps with a pulse to zero bias or a small forward bias which enables them to capture carriers from the bulk population. If the pulse is very short it is possible that the traps will not have time to fill completely. In fact if we consider

electron traps, the extent to which they fill will depend on the trap cross-section On , the pulse length t, and the majority carrier population n, so that the number of traps occupied at time t is nT(t), then nT(oo) - nT(t) _J^ TW

=

AC , - ACm „ (t)

—/ \

A/^

n T (oo)

AC ( o o )

=

_ r

L v

n th

y

J

where the thermal velocity V^ = (3 kT/m*)1/2 - 107 cm s'1 at room temperature, So as the filling pulse length is reduced a point is reached where the trap fails to fill completely and the amplitude of the emission transient and hence the DLTS peak height is reduced. The usual procedure is to set the temperature so that the DLTS output is at a peak and then change the filling pulse width by known amounts. The slope of the plot of log[ (AC(oo) - AC(t)) / AC(oo) ] against t will then give the majority capture cross section directly. The interesting point about this method is that the capture is from the bulk, i.e. in a field free region, whereas the emission is observed in the space charge layer. Unfortunately, it is often not possible to make the pulse length as short as may be required. The technique requires considerable care for pulses less than 100 ns and it becomes increasingly difficult as the pulse length is reduced and virtually impossible at less than 2 ns. This is because of the necessity for impedance matching and the device and system capacitance. If a forward bias pulse is applied so that current flows then, in a p-n junction, minority carriers are injected and may be trapped and then emitted. If a p+n or an n+p device is used then a rough estimate of the minority carrier cross-section can be made from a calculation of the injected hole density. A much better method is to use a separate injecting junction [19] or optical excitation. In both cases, the diode under test is maintained at constant reverse bias and the 'filling' pulse applied to the subsidiary diode structure or to the optical source. These techniques fill the state in the depletion field which may introduce uncertainties of its own but very importantly permit the minority carrier population to be adjusted to a level which permits convenient filling times for the state. In this way extremely large cross-sections can be measured accurately. The optical method has found widespread use as it can be applied to simple structures and in particular to Schottky diodes made with thin semi-transparent metal layers. By using a light source of variable intensity with a DLTS system it is possible to undertake a selective scan surveying for states likely to be powerful recombination centres [20]. H

OTHER TECHNIQUES

The methods described so far account for over 95% of recent data relating to the electrical properties of deep states in GaAs. However, for completeness, three other techniques should be mentioned: TSC (thermally stimulated current, conductivity or capacitance), Hall measurements and admittance spectroscopy. TSC experiments played a very important role in the study of deep states prior to the widespread use of DLTS. Conceptually it is very similar but relies on filling the traps at low temperature (by electrical or optical perturbation of the equilibrium carrier population) and then observing the current capacitance or conductivity as the temperature is increased. The variation of the peak position with heating rate is used to determine the activation energy. The technique is insensitive

and inaccurate compared to DLTS. A good example of its application to semi-insulating material is provided by Martin [19]. If a significant number of defect states are present they can be observed in a C-V measurement of a junction in the sense that when the majority carrier emission rate is similar to the measurement frequency, they contribute a real component to the complex admittance. The technique is valuable in detecting defect states with energy levels too near the band edge (actually too near the dopant energy level) to be measured using DLTS. A general analysis of this admittance spectroscopy technique is presented by Blood and Orton [20]. If a deep state is present in sufficiently large concentrations to significantly perturb the Fermi level at any accessable temperature, an analysis of the temperature dependent Hall confinement can be used to determine some of the deep state parameters. In all but the most favourable samples, the analysis is difficult and often ambiguous. Some aspects of the application of the techniques to GaAs are discussed by Look and Sizelove [21]. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] II1] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21]

R. Williams [ J. Appl. Phys. (USA) vol.37 (1960) p.3411 ] CT. Sah, L. Forbes, L.C. Rosier, A.F. Tasch [ Solid-State Electron. (UK) vol. 13 (1970) p.759 ]; S. Weiss, R. Kanning [ Solid-State Electron. (UK) vol.31 (1988) p. 1733 ] P. Blood, J. W. Orton [ The Characteristics of Semiconductors: Majority Carriers and Electron States (Academic Press, London, 1992) ] DJ. Lang [Appl Phys. Lett. (USA) vol.45 (1974) p.3023 ] F.D. Auret [ Rev. Sci. Instrum. (USA) vol.57 (1986) p. 1597 ] P.I. Rockett, A.R Peaker [ Electron. Lett. (UK) vol.17 (1981) p.838 ] D.C. Look [ Electrical Characterisation of GaAs Materials and Devices (Wiley, Chichester, 1989) p.211] JJ. Shian, A.L. Fahrenbruch, R.H. Bube [ Solid-State Electron. (UK) vol.30 (1987) p.513 ] P.I. Rockett, A.R. Peaker [Appl. Phys. Lett. (USA) vol.40 (1982) p.957 ] H. Lefevre, M. Schuiz [ Appl. Phys. Lett. (USA) vol.2 (1977) p.45 ] M.S. Hodgart [ Electron. Lett. (UK) vol. 14 (1978) p.388 ] CR. Crowell, S. Alipanahi [ Solid-State Electron. (UK) vol.24 (1981) p.25 ] D.D. Nolte, E.E. Haller [ J. Appl. Phys. (USA) vol.62 (1987) p.900 ] L. Dobaczewski, P. Kaczor, LD. Hawkins, A.R. Peaker [ J. Appl. Phys. (USA) vol.76 (1994) p. 194] L. Dobaczewski, P. Kaczor, A.R. Peaker [ in Defects in Semiconductors; Mater. Sci. Forum (USA)EdSU. Heinrich,W. Jantsch(Trans. Tech.,Zurich, 1993)p. 1001 ] B. Hamilton, A.R. Peaker, D.R. Wight [ J. Appl. Phys. (USA) vol.50 (1979) p.6373 ] R. Brunwin,B. Hamilton, P. Jordan, A.R. Peaker [ Electron. Lett. (UK) vol.15 (1979) p.349] A. Mitonneau, G.M. Martin, A. Mircea [ IOP Conf. Ser. (UK) no.33 (IOP, Bristol, 1978) p.73 ] G.M. Martin [ Semi-insulating IU-VMater. Ed. GJ. Rees (Shiva, Orpington, 1990) p. 13 ] P. Blood, J. W. Orton [ The Characteristics of Semiconductors: Majority Carriers and Electron States (Academic Press, London, 1992) p.492 ] D.C. Look, J.R. Sizelove [J. Appl. Phys. (USA) vol.61 (1987) p.1650 ]

10.4 Defect densities in melt-grown GaAs (a review) M.R. Brozel October 1996

A

INTRODUCTION

Melt growth of substrate material is the starting point for all GaAs technology, whether it be direct ion-implantation or epitaxial based. In this Datareview we consider the different types of melt growth technique and the defects that are introduced into the crystal by them. For the most part, we restrict ourselves to undoped, semi-insulating GaAs. Most undoped SI GaAs is still produced by the LEC method. It contains crystal defects that result in both the semi-insulating behaviour (that makes the material so attractive for many applications) and the instability and non-uniformity of substrates that have been the bane of the GaAs industry for many years. LEC GaAs has improved in quality over the past ten or so years and for this reason, older analyses of defect concentrations in LEC GaAs have been omitted from this review as they do not represent either current material or currently accepted accuracies of measurement. For similar reasons alternative methods of growth that have been proposed but subsequently found to be commercially unviable have also been omitted. In general, crystal defects can be divided into three groups: (1) (2) (3)

line defects (dislocations) volume defects such as precipitates point defects.

We consider these in this unusual order because the distribution of volume defects and point defects will be found to depend on that of dislocations. B

LINE DEFECTS

Dislocations in all commercial SI LEC GaAs exist at densities, D, between 104 and 105 cm"2 [1-3]. They are introduced because the growing crystal cools by radiation and convection via the high pressure ambient gas (usually argon) in a rather uncontrolled manner. The thermal gradients produced radially in the crystal are sufficient to cause plastic flow by the generation of slip dislocations. These dislocations slip along the available glide directions until they interact with each other. This interaction involves climb and the resultant distribution consists mainly of polygonised arrays (cell structure). It would appear that the thermal gradients are not much increased even if boules of up to 6 inch diameter are grown. As a result, the dislocation arrangements are not much worse in larger crystals. It is generally held that these dislocation densities, being high and variable across a substrate, are unacceptable. As a result several groups have attempted to reduce D by modifying these thermal gradients or by introducing isoelectronic elements such as indium which cause lattice hardening [4-6]. The only techniques now in current use are LEC, (see Datareview 16.2), and modified techniques designed to grow LEC GaAs with reduced D (see Datareview 16.4), vertical Bridgman growth and vertical gradient freeze growth (see Datareview 16.3).

In all these modern growth techniques D is reduced by over an order of magnitude. However, the majority of the dislocation arrays are still observed to be in the form of polygonised cell walls. An image of the dislocation structure in a 2 inch diameter LEC wafer is shown in FIGURE 1, where lineage (linear dislocation arrays aligned close to <110> diameters) and residual un-polygonized slip dislocations are also seen.

FIGURE 1. Dislocation distribution in a 2" diameter {001} wafer of SI, LEC GaAs as revealed by reflection X-ray topography (DJ. Stirland, private communication).

A similar distribution is found in VGF and VB GaAs although the cell sizes are generally much larger (~ 1 mm compared to 200 - 300 |um). This arrangement is in good agreement with a model of dislocation inter-reaction which allows the total strain energy of the dislocations to minimize

Vl The adverse properties of dislocations in many devices probably result more from their associated atmospheres of defects and precipitates than from the dislocations themselves (see Datareviews 16.5 and 16.6). However, dislocation climb under intense conditions of high minority carrier injection and high luminous intensity is known to cause rapid degradation of lasers [8]. C

VOLUME DEFECTS

Volume defects are, in general, observed to be spatially associated with dislocations and those that are large enough to be studied have been analysed by X-ray diffraction to be hexagonal arsenic

precipitates [9,10]. However, the presence of microprecipitates consisting of GaAs particles has been reported by some workers [11] and amorphous precipitates by others [12]. The mechanisms by which these precipitates are produced are not known and their effects on MESFETs produced by ion-implantation are also unknown. However, adverse effects of precipitates on MESFETs produced after epitaxial growth by the chloride process have been reported [13]. Certain ingot and wafer heat treatments cause a redistribution of the precipitates and hence an improvement in material homogeneity (see Datareviews 16.5 and 16.6). D

POINT DEFECTS

Point defects can be separated into two major categories: (1)

chemical impurities

(2)

native defects.

There is also the possibility of complexing between such basic entities. Dl Chemical Impurities One of the most sensitive techniques for determining impurity concentrations in semiconductors is glow discharge mass spectrometry (GDMS). A GDMS analysis of two commercial undoped GaAs crystals from both seed and tail ends is shown in TABLE 1. TABLE 1. Glow discharge mass spectrometry data (1995) (2 crystals) Concentration in atomic ppb Crystal 1

Crystal 2

Element

Seed

Tail

Seed

Tail

Li

<3

<3

<3

<5

Be

<2

<2

<2

<3

B

81

1000

110

1200

C

1100* (23)

540* (10)

1200* (7)

940*(<4)

N

39*

57*

38*

53*

O

250*

170*

220*

140*

F

<36

<21

<26

<30

Na

2

1

1

<0.6

Mg

<0.6

2

<0.6

<0.4

Al

<0.9

5

<0.7

<0.5

Si

2

3

2

1

P

29

1

28

4

S

2

5

4

4

Cl

54 t

14t

391

6t

K

<2

<2

<3

<3

Ca

<4

<4

<4

<4

Sc

<0.1

<0.1

<0.1

<0.2

Ti

<0.2

1

5

<1

V

<0.1

<0.1

<0.3

<0.4

Cr

7

3

<0.3

<1

Mn

<0.4

<0.4

<0.3

<0.4

Fe

0.4

1

<0.3

<0.2

Co

<0.3

<0.3

<0.3

<0.3

Ni

<0.3

<0.3

<0.3

<0.3

Cu

<0.5

<0.5

<0.5

<0.5

Zn

0.6

1

<0.7

1

Ga

Host

Host

Host

Host

Ge

<18

<17

<17

<25

As

Host

Host

Host

Host

Se

<7

<8

<8

<11

Br

<3

<5

<4

<7

Rb

<0.8

<0.8

<0.8

<1

Sr

<0.1

<0.2

<0.1

<0.1

Y

<0.1

<0.1

<0.1

<0.1

Zr

<0.3

<0.3

<0.3

<0.3

Nb

<0.2

<0.2

<0.1

<0.1

Mo

<0.4

<0.5

<0.7

<0.8

Ru

<2

<2

<3

<3

Rh

<2

<1

<1

<1

Pd

<0.7

<0.9

<0.8

<0.5

Ag

<8

<12

<7

<12

Cd

<2

<2

<2

<1

In

<48

<75

<40

<87

Sn

<1

<2

<2

<3

Sb

<0.3

<0.6

<0.3

<0.3

Te

<2

<2

<2

<3

I

<0.5

<0.5

<0.5

<0.8

Cs

<0.2

<0.2

<0.2

<0.3

Ba

<0.3

<0.4

<0.4

<0.4

La

<0.1

<0.1

<0.1

<0.1

Ce

<0.2

<0.3

<0.2

<0.2

Pr

<0.1

<0.1

<0.1

<0.1

Nd

<0.3

<0.4

<0.4

<0.6

Sm

<0.3

<0.3

<0.3

<0.4

Eu

<0.2

<0.2

<0.2

<0.2

Gd

<0.3

<0.3

<0.3

<0.4

Tb

<0.1

<0.1

<0.1

<0.1

Dy

<0.2

<0.3

<0.3

<0.4

Ho

<0.1

<0.1

<0.1

<0.1

Er

<0.2

<0.2

<0.2

<0.3

Tm

<0.1

<0.1

<0.1

<0.1

Yb

<0.3

<0.4

<0.3

<0.5

Lu

<0.1

<0.1

<0.1

<0.1

Hf

<0.3

<0.4

<0.3

<0.5

Ta

<170 J

<80J

<140 J

<220{

W

<0.2

<0.2

<0.3

<0.2

Re

<0.4

<0.4

<0.4

<0.5

Os

<0.6

<0.6

<0.5

<0.6

Ir

<0.3

<0.4

<0.4

<0.5

Pt

<0.5

<0.6

<0.5

<0.7

Au

<2

<1

<1

<2

Hg

<0.5

<0.6

<0.5

<0.7

Tl

<0.4

<0.5

<0.4

<0.6

Pb

<0.2

<0.3

<0.2

<0.3

Bi

<0.2

<0.2

<0.2

<0.2

Th

<0.1

<0.1

<0.1

<0.1

U

<0.1

<0.1

<0.1

<0.1

* Background gaseous contaminant in GDMS. t Residue fromsample pre-clean. } Ta signal from sample holder. ( ) [CAJ measured by LVM absorption.

As can be seen, undoped GaAs is a very pure material even in comparison with silicon. The high concentrations of boron apparently have little effect on the crystal, probably because they are isoelectronic BGa defects. The large P signal may be due to a previous measurement on a phosphide. Concentrations of the important acceptor, carbon, determined by infrared absorption due to localised vibrational modes (LVM) (see Datareview 7.1, Section C) have now been reduced to the 1014 to 1015 cm"3 range by improvements in starting material purity and growth technique. The important shallow donor impurities S and Si have been reduced in concentration to the 1014 cm"3 range for similar reasons [14-18]. D2

Native Crystal Defects and Deep Levels

Generally, native defects can only be observed if they are associated with energy levels situated relatively deep in the bandgap. They can be detected and their concentrations measured in SI material using thermally stimulated current (TSC) spectroscopy [19-21]. This technique can be extended to detect deeper traps by using light excitation: optical transient current spectroscopy (OTCS) [22]. Because of the difficulty in separating hole traps and electron traps using either TSC or OTCS another method, deep level transient spectroscopy (DLTS), has been developed [23] (also see Datareview 10.3). Unfortunately, DLTS requires that a rectifying junction be present and this is not possible when investigating semi-insulating material. Accordingly, many workers grow n-type LEC GaAs in order to perform DLTS measurements and then assume that defects present in n-type GaAs are also present in SI material [24]. All these techniques are very sensitive (the detectivity can be about 1012 cm"3 in some cases) and it is often difficult to identify levels as being due to particular chemical impurities or native crystal defects. For this reason all these deep level defects are treated in the same section. Deep levels are split into two groups: deep donor or electron levels (EL group) and deep acceptor or hole levels (HL group) according to their signature in DLTS measurements (TSC and OTCS do not distinguish the two types). TABLE 2 is compiled mostly from data by Martin (TSC and OTCS) [25], Farges et al (DLTS) [26] and Iwata et al (DLTS) [24], but other data have been included where necessary. After growth high concentrations of EL2 centres are known to be associated with dislocations, see Datareview 10.7. It is also known that under these circumstances, an opposite correlation is found for concentrations of the 'reverse contrast' defect (also, see Datareview 10.7). Such correlations result from the optical absorption from these defects that can be imaged to produce a 2 dimensional mapping of concentration distribution. It is probable, but not proven, that concentrations of other point defects are also modified by the dislocation arrays. The concentrations of many of the native point defects shown in TABLE 1 (those without known impurity-related origins) are controlled thermodynamically. The application of single or multiple heat treatments after growth, either to the entire ingot or to wafers, can be used to modify these concentrations as well as to control the behaviour of precipitates: see Datareviews 16.5 and 16.6 for more details.

TABLE 2. Label

Origin

Concentration (cm-3)

ELIl

Emission energy E (eV)

Capture cross section (cm2)

Ref

E.-0.17

3 x 10-16

[25] 14

[25]

EL17

E c -0.22

1.0 x 10

EL14

E c -0.215

5.2 x 10-16

[25]

10 14 -10 16

E c -0.35

1.5 x 10 13

[25]

2 x 1015

E c -0.35

1.5 x 10'13

[26]

4 x 1014

E c -0.35

1.5 x 10 13

[24]

10 i 4 -10 1 6

E c -0.42

io- 13

[25]

EL6

complex defect

EL5

14

below 2 x 10

E c -0.42

io-

10 13 -10 15

Ec - 0.575

1.2 x 10"13

[25]

below 2 x 1014

E c - 0.575

1.2 x 10- 13

[26]

4 x 10 13 - 1 x 1014

E c -0.575

1.2 x I O 1 3

[24]

5 xlO15-3 xio16

Ec - 0.825

1.2 XlO' 1 3

[25,29]

2 x 10 1 3 -6x 1015

Ec - 0.825

1.2 x 10"13

[29,30]

0 - 3 x 1017 (doping dependent)

Ev +0.886

io- 14

[25]

below 2 xlO 1 4

Ev + 0.83

1.7 x IO-13

[25]

Ev + 0.83

1.7 x IO 13

[26]

Ev + 0.69

1.1 x IO-13

[24]

Ev + 0.59

3 x 1O"15

[25]

below 1015

Ev + 0.42

3 x 10"15

[25]

below 2 x 1014

Ev + 0.42

3 x IO-15

[26]

1 x 1015

Ev + 0.42

3 x 10-15

[24]

Ev + 0.35

6.4 x IO 15

[24]

Zn-associated

Ev + 0.27

1.3 x 10"14

[25]

Cu

Ev + 0.14

[25]

Ev +0.077 Ev + 0.203 (double acceptor)

[27,28, 30]

EL3

EL2*

HLl

native defect As011 or [As011 - X]

Cr

HLlO

HL9 HL3

Fe

HL4

Cu

15

1.7 x 10

HL7 HL12

13

gallium antisite** Ga^ or boron antisite** B ^

3 x 10 1 5 to3x 1016 (dependent on Garichnessofmelt)

[26]

* The DLTS signature of EL2 has been resolved into at least two components, the relative concentrations apparently being a function of crystal growth parameters [31,32]. EL2, the most important of the native deep levels in SI, LEC GaAs, gives rise to an infrared absorption band at 1 micron [33] and the measurement of this absorption is now the usual method of determining [EL2] [34-36]. For further discussion see Datareview 10.2, Makram-Ebeid [37] and Baraff [29]. ** P"tyPe material grown from Ga-rich melts. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] II1] [12] [13]

[14]

[15] [16] [17] [18] [19] [20] [21] [22] [23] [24]

[25]

A.S. Jordan, R. Caruso, A.R. Von Neida [ Bell Syst. Tech. J (USA) vol.59 (1980) p.593 ] (theoretical discussions) A.S. Jordan, A.R. Von Neida, R Caruso [ J. Cryst. Growth (Netherlands) vol.70 (1984) p.555 ] D. Rumsby, R.M. Ware, M. Whitaker [ Proc. Con/, on Semi-Insulating III-V Materials, Nottingham, April 1980 (Shiva Publishing, UK, 1980) p.59-67 ] M. Duseaux[J. Cryst Growth (Netherlands) vol.61 no.3 (1983) p.576-67] G. Jacob [Proc. Conf. on Semi-Insulating III-V Mater.,Evim, France, 1982 (Shiva Publishing, Nantwich,UK, 1982) p.2] G. Jacob[7. Cryst. Growth (Netherlands) vol.58 no.2 (1982) p.455-9] D.L. Holt [ J. Appl Phys. (USA) vol.441 no.8 (1970) p.3197 ] D.V. Lang, CH. Henry [ Solid-State Electron. (UK) vol.21 no. 11-12 (1978) p. 1519-24 ] A.G. Cullis, P.D. Augustus, DJ. Stirland [ J Appl. Phys. (USA) vol.51 no.5 (1980) p.2556-60 ] DJ. Stirland, P.D. Augustus, M.R. Brozel, EJ. Foulkes [ Proc. Conf. on Semi-Insulating III-V Materials, Kah-nee-ta, OR, USA, 24-26 Apr 1984 (Shiva Publishing, Nantwich, UK, 1984) p.91 ] J.P. Cornier, M. Duseaux, J.P. Chevalier [ Inst. Phys. Conf. Ser (UK) no.74 (1985) p.95 ] FA. Ponce, F-C. Wang, R. Hiskes [ Proc. Conf on Semi-Insulating III-V Materials, Kah-nee-ta, OR, USA, 24-26 Apr 1984 (Shiva Publishing, Nantwich, UK, 1984) p.68-75 ] H. Yamamoto, H. Shimakura, G. Kano, M. Seiwa, H. Nakata, O. Oda [ Proc. Int. Conf Sci. & Tech. Defect Control in Semiconductors, Yokohama, 1990 ]; See also H. Yamamoto, O. Oda, M. Seiwa, M. Taniguchi, H. Nakata, M. Ejima [J. Electrochem. Soc. (USA) vol.136 no. 10 (1989) p.3098-102] PJ. Pearah, R Tobin,, J.P. Tower, R.M. Ware, L. Sargent, J.S. Blakemore [ Proc. 5th Int. Conf. Semi-Insulating III-VMater., Malmo, Eds. G. Grossmann, L. Ledebo (IOP Pub. Ltd, Bristol, 1988) p. 195] L. Sargent, RS. Tang, J.S. Blakemore [Semicond. Semimet. Eds. RK. Willardson, A.C. Beer, vol.20 (Academic Press, New York, 1984) p. 159 ] S. Reichlmeier, K Lohnert, M. Baumgartner [Jpn. J. Appl. Phys. (Japan) vol.27 (1988) p.2329 ]; T. Hashizume, H. Nagabuchi [ Semicond. Sci. Technol. vol.4 (1989) p.427] R.S. Tang, L. Sargent, J.S. Blakemore [ J. Appl. Phys. (USA) vol.66 (1989) p. 1989) ] W.M. Duncan, G.H. Westphal, AJ. Purdes [ J Appl. Phys. (USA) vol.66 (1989) p.2430 ] M.G. Buehler [ Solid-State Electron. (UK) vol. 15 (1972) p.69 ] Z.Q. Fang etal [Mater. Sci. Eng. (UK) vol.B5 (1990)p.297] Z.Q. Fant et al [J. Appl. Phys. (USA) vol.69 (1991) p.69 ] C. Hurtes, M. Boulou, A. Mittonneau, D. Bois [Appl. Phys. Lett. (USA) vol.32 no. 12 (1978) p.821-3] D.V. Lang [ J Appl. Phys. (USA) vol.45 no.7 (1974) p.3023-32 ] N. Iwata, F. Hasegawa, N. Yamamoto, Y. Nannichi [ Jpn. J. Appl. Phys. Suppl. (Japan) no.22-1 (1983) p.595-6 ] (an example of this technique where a p/n junction was produced after Zn diffusion inton-GaAs). See also, same authors [ Jpn. J. Appl. Phys. (Japan) vol.22 no.8 (1983) p.L502-4 ] G.M. Martin [ Proc. Conf on Semi-Insulating III-VMaterials, Nottingham, Apr 1980 (Shiva Publishing, Orpington, UK, 1980) p. 13-28 ]

Next Page

[26]

[27] [28] [29]

[30] [31] [32] [33] [34] [35] [36] [37]

J.P. Farges, G. Jacob, C. Schemali, G.M. Martin, A. Mircea-Roussel, J. Halais [ Proc. Conf. on Semi-Insulating IU-VMaterials, Evian, France 1982 (Shiva Publishing, Nantwich, UK, 1982) p.45] K.R. Elliott, D.E. Holmes, R.T. Chen, CG. Kirkpatrick [ Appl. Phys. Lett. (USA) vol.40 no.l (1982) p.898-901 ] P. Won Yu, W.C. Mitchel, M.G. Mier, S.S. Li, W.L. Wang [Appl. Phys. Lett. (USA) vol.41 no.6 (1982) p.532-4 ]; see also K.R. Elliott [ Appl. Phys. Lett. (USA) vol.42 no.3 (1983) p.274-6 ] The identity of the As antisite and its relationship to EL2 is still unresolved. For discussions of these points see D.M. Hofinann, B.K. Meyer, F. Lohse, J.-M. Spaeth [ Phys. Rev. Lett. (USA) vol.53 no. 12 (1984)p. 1187-90]; BK. Meyer, J.-M. Spaeth, M. Scheffer [Phys. Rev. Lett. (USA) vol.52 no. 10 (1984) p.851-4 ]; for a more recent discussion, G.A. Baraff [ Proc. Conf. on Semi-Insulating IH-VMaterials, Ixtapa, Mexico (Inst. Phys. Pub., Bristol and Philadelphia, 1992) p.l 1 ] J. Woodhead, R.C. Newman, I. Grant, D. Rumsby, RM. Ware [ J. Phys. C (UK) vol. 16 no.28 (1983)p.5523-33] M. Taniguchi, T. Dcoma [ J. Appl. Phys. (USA) vol.54 no. 11 (1983) p.6448-51 ] J. Lagowski, D.G. Lin, T. Aoyama, EC. Gatos [Appl. Phys. Lett. (USA) vol.44 no.3 (1984) p.3368] G.M. Martin [ Appl. Phys. Lett. (USA) vol.39 no.9 (1981) p.747-8 ] MR. Brozel, I. Grant, RM. Ware, DJ. Stirland [Appl. Phys. Lett. (USA) vol.42 no.7 (1983) p.610-12] D.E. Holmes, RT. Chen, K.R. Elliott, CG. Kirkpatrick [ Appl. Phys. Lett. (USA) vol.43 no.3 (1983)p.305-7] M.S. Skolnick, M.R. Brozel, LJ. Reed, I. Grant, DJ. Stirland, RM. Ware [J. Electron. Mater. (USA) vol.13 no.l (1984) p. 107-25 ] S. Makram-Ebeid, P. Langlade, G.M. Martin [ Proc. Conf. on Semi-InsulatingIU-VMaterials, Kah-nee-ta, OR, USA 24-26 Apr 1984 (Shiva Publishing, Nantwich, UK, 1984) p. 184-203 ]

10.5 Deep states in as-grown epitaxial GaAs

Previous Page

A.R. Peaker May 1996

A

INTRODUCTION

A vast quantity of data exists on the energy levels and behaviour of deep defect states in gallium arsenide. Most of these data are phenomenological, i.e., they describe the behaviour of the defect in terms of an observable property. There is very little reliable information on the detailed physical structure of the defects observed in this way in GaAs. As a consequence, the assignment of the observed property (such as thermal emission behaviour for example) is often referred to by a code (e.g., ELl 1 or A) rather than a meaningful structure. In a few cases, the physical nature of the defect is known but in most cases it is at best only 'associated' with a process or the introduction of an impurity. In the majority of cases, particularly in the case of epitaxial layers, techniques which give us structural information (e.g., ESR) have not been used in conjunction with the measurement of the deep state properties or have had inadequate sensitivity. This is an important factor to remember when using the data that follow: although assignments of deep states to impurities and structural defects have been made by some authors and are given in this Datareview, the reliability of these assignments is often dubious and should be judged after reference to the original papers. Although there are some common features among layers prepared by different techniques, the deep state population depends greatly on the method of growth. This is primarily for two reasons. The first and most important relates to the native defect population. Liquid phase epitaxial material grown from a gallium melt would be expected to have fewer defects related to gallium vacancies, VGa, or arsenic-antisite defects, AsGa, than material grown from the vapour phase where an excess of arsenic was present. Similarly we would expect different impurities to be present in MBE (molecular beam epitaxy) than in MOVPE (metal organic vapour phase epitaxy) GaAs because of different source temperatures and purification techniques for the precursors. This is indeed the case and these issues are discussed under the type of epitaxy in the following sections. B

DEFECT ENERGY LEVELS IN LPE GaAs

There is a consensus that LPE GaAs is relatively trap-free compared to GaAs grown by other methods. Indeed, workers looking specifically for electron traps in material grown from gallium melts at low growth rates (around 0.1 ^m min"1) and doped only with shallow donors have found none. This means that the electron trap concentrations are certainly below 1012 cm"3 [1,2] and probably as much as an order of magnitude less than this. However, in p-type as-grown material two acceptor-like hole traps are commonly seen. These were first observed by Lang [3] and occur in roughly equal concentrations. He called them A and B and they have subsequently been investigated in considerable detail [4-14]. There is a near consensus in relation to the thermal emission characteristics:

TABLEl Name

Activation energy for hole emission

T (K) for emission rates of 100 s*1

(me V) A

400

182

JB

710

298

The states are in the 1011 to 1015 cm"3 concentration range in as-grown p-type GaAs although not in n-type material subsequently rendered p-type after diffusion with zinc [9]. It appears that this is due to some form of gettering action by the zinc because the states have been reported in undiflfijsed n-type material [10,12] with a concentration given by : [A]-[B] = Kn 0 5 cm"3

(1)

where n is the free electron concentration (cm"3) in material grown at 0.08 j^m/min. The constant K is reported by Zhou [12] as being primarily a function of growth temperature. The values reported are : TABLE 2 Growth temp ( 0 C)

740

800

850

ConstantK

1.7 x 106

9.3 * IQ6

13 x IQ6

If, however, the growth rate is increased the concentrations of the centres decrease dramatically so that at 0.24 |nm min"1 the concentration is reduced by two orders of magnitude. A model has been proposed in which the A centre is a gallium atom on an arsenic site paired with a gallium vacancy while the B centre is an arsenic atom on a gallium site paired with a gallium vacancy. However, Wang et al [6] postulated that the A and B centres were merely different charge states of the same defect, that being a gallium atom on an arsenic site. In contrast, Nouailhat et al [13] has undertaken an investigation of the centres using deep level optical spectroscopy and concluded that the states A and B were different defects. Zhou's more extensive study [12] provides additional evidence for this view as they only observe identical concentrations of the A and B centres in layers grown at 0.08 ^m min"1: at faster growth rates the concentrations were different. However, Wang et al [6] observed that in their LPE system the concentrations of A and B were below 1011 cm"3 with n = 2 x 10 16 cm "3. Only with fast cooling could measurable concentrations of A and B be produced. Other factors besides these may also affect the concentrations of these centres. Kalukhov has observed that the concentrations of A and B decrease with the addition of either In or Sb [7]. Partin [11] has measured the capture cross-sections of the A and B centres at 300 K.

TABLE 3 Name

Hole cross-section (cm2)

Activation energy (holes)

Electron cross-section

(meV)

(cm2)

A

400

1.4 xlO' 17

4.2 xio- 15

B

710

4.2 xlO"20

2.IxIO" 15

Comparable values have been obtained by Lang [3] while measurements at lower temperatures have been made by Wang et al [6]. Mines [9] has calculated that the A centre is likely to dominate recombination in n-type material at room temperature and so is of crucial importance in devices using holes as minority carriers. The overall picture of recombination in GaAs material has been discussed in detail by Peaker and Hamilton [14] and in relation to carrier lifetime in Chapters 2 and 3 of this book. Photocapacitance measurements on the A and B centres have also been undertaken [6,8,12,15]. Several other hole traps have been observed in LPE GaAs with only shallow dopants or no dopants intentionally added. A number of deep states with different properties have been observed when additional impurities have been added to the melt. TABLE 4 summarises both these data and those relating to inadvertent centres. Arrhenius plots of the emission rate of holes from these traps are given in the individual references or in a number of general reviews [8,15-26] TABLE 4 Name

Energy (meV)

T(K) at e = 100 s"1

Associated with

Reference

HBl

780

427

Cr

[20]

HL2

730

347

HB2

710

298

HS2

640

358

[25]

HSl

580

488

[25]

HB3

520

300

Fe

[2,7]

HB4

440

213

Cu

[3,21]

HL5

410

192

State A?

[22]

HB5

400

182

State A

[3,6]

HC12

270

141

Zn

[22]

Cu(a)

130

69

Cu

[3,27]

[22] State B

[3,6,H]

Martin et al present a detailed discussion of the Cu level [17] while optical characterisation of the HBl state is described in [23]. An important point is that in the cases where Fe, Cr or Cu were added to the melt the distribution coefficient was found to be very low (approx. 10"6) in comparison to typical distribution coefficients from a stoichiometric melt of the order of 10"3 at ~ 12380C. It is this fortuitous situation which enables the LPE process to produce almost trap-

free GaAs under slow cooling conditions unless large quantities of the impurities known to produce traps e.g. Cu, Fe5 Cr, are present. As mentioned previously electron traps are rarely seen. However, two cases are worthy of note. Zhou [12,28] observes electron traps at 630 and 780 meV in material produced at high growth rates (0.22 to 0.24 (im min"1). In addition Yakusheva [29] has studied LPE material grown from a bismuth melt. For material grown in the temperature range 900 to 8400C a trap with identical electron emission characteristics to EL2 is seen (Ea = 860 meV and T = 391 K for e = 100 s"1) at a concentration of about 1013 cm"3. For material grown from over 800 to 7400C the EL2 like state is present in higher concentration, 5 x 1013 cm"3, but in addition a trap with identical electron emission characteristics to EL5 is seen in comparable concentration (420 meV and T = 211 K for e = 100 s"1). Yakusheva has supported his assignment of trap identities by confirming that the absolute values and temperature dependence of the electron capture cross section are the same as those for EL2 and EL5. C

DEFECT ENERGY LEVELS IN MBE GaAs

Cl

Electron Traps

Using DLTS to study n-type MBE grown GaAs Lang [30] observed nine electron traps M0...M8. In subsequent work the levels M1...M4 have been observed consistently in MBE GaAs grown between 500 and 650 0 C whether n-type or p-type. The states occur in the 1012 to 1014 cm"3 concentration range. However, M6...M8 do not appear reproducibly and it has been proposed that they may result from surface effects [31]. Martin [16] has reported electron traps in MBE layers (ElA, EL7 and ELlO) which are similar to traps in the M series. Despite early associations with deep states observed in material grown by other techniques, it is now apparent that Ml, M3 and M4 are unique to MBE and are almost certainly related to lattice defects. Conversely, a number of levels often seen in GaAs grown by VPE techniques are not detected in MBE unless the GaAs is heat treated after growth (e.g., EL2) [32,43]. A listing of MBE states is given in TABLE 5. TABLE 5 Name

Energy (me V)

EB3LT

Cross section (cm2)

Notes (T8)

Main ref.

858

T g <490°C

[33]

EL2A

825

annealed

[32-36]

E5A

759

annealed

[32]

EL3LT

566

T g <490°C

[33,16]

M5

565

EL4

510

M4 (EB4)

480

EB6LT

405

T g <300°C

[31]

EL7

301

M3?

[16]

[30,32,35] M4? > 2 x 10"16

[16] [30,31,37]

M3(EB7)

300

1.1 x 10"16

M2

220

>4xlO" 1 7

T g <600°C

[30,31,35]

M2!

230

10" 2 0

T g <600°C

[31]

Ml (EB8)

189

> 2 x 10"16

ELlO

159

Ml?

[16]

ElO

165

annealed

[32]

MO

75

MOO

30

[30,31]

[30,31,37]

[30] low concentration

[34]

In general, concentrations of deep states decrease with increasing substrate temperature during growth (Tg). The exception to this is M2 which is only present in material grown at around 650 0 C. At lower temperatures a trap (M2!) with similar emission but very different capture cross section is present [31]. However, at very low growth temperatures (3000C) the M series disappears, but insufficient data exist to quantify the changes between 300 and 500 0 C. The trap concentrations are also dependent on the III/V flux ratio [31,35] and on whether dimeric or tetrameric arsenic is used [37]. In particular, Ml, M3 and M4 are found in much lower concentrations when the dimer is used. Heat treatment after growth increases the concentration of some traps and decreases others [36,38]. However, Chand et al [42] have recently undertaken a detailed study on the purity of the arsenic and gallium sources. They have produced GaAs with very low concentrations of deep states (in some cases undetectable) and present in a persuasive case that the M series of traps are native defect-impurity complexes. C2

Hole traps

Studies of hole traps have not been as detailed as those of electron traps. However, it seems that, unlike the case of the electron traps, defects seen in LPE and VPE material are apparently identical. TABLE 6 Name

Energy (meV)

Notes

Reference

900A

900

after 750 0 C anneal

[38]

HBl

780

associated with Cr in LPE

[33]

HL9

690

seen in VPE

[33]

HL3

590

seen in Fe diffused VPE

[31,38]

HL8

520

sameasHL3?

[18]

HB4

440

seen in Cu doped LPE

[33]

HL7

350

[18]

HB6LT

290

seen below 4900C

[33]

Although no capture cross sections have been determined in MBE material, HL3, HL8, HB4 and HBl have been measured in LPE and/or VPE [3]. All are classic acceptors with large hole cross sections and small electron cross sections. Iron, manganese and chromium have been detected in MBE GaAs using secondary ion mass spectrometry (SIMS) and their origin ascribed to the effusion sources [40]. Concentrations of these impurities were shown to decrease with increasing substrate temperature. SIMS measurements by Palmateer et al [41] indicate that substrates were a major source of manganese. C3

Recombination

Not enough measurements have been done to make definitive statements on recombination but Blood and Harris [31] suggest that the minority carrier lifetime in p-type MBE GaAs will be limited by M4 for growth temperatures below 625 0 C and by M2 above this temperature. In ntype material minority carrier lifetime at the low injection levels typical of photovoltaic devices may be limited by HB1. The general problem of recombination in GaAs has been reviewed elsewhere [14]. D

DEFECT ENERGY LEVELS IN VPE GaAs

In GaAs material grown by the arsenic trichloride-gallium or arsine-gallium-hydrogen chloride processes both hole and electron traps are commonly seen. Dl

Electron Traps

The trap present in the highest concentration is EL2, a ubiquitous defect present in melt grown and VPE material [44]. A very substantial effort has been devoted to investigating this deep state which was originally associated with oxygen. Subsequent studies have shown conclusively that EL2 is quite unrelated to oxygen [45,46] as are all other commonly observed deep states in GaAs [47]. There is now considerable evidence that EL2 is a family of slightly different states associated with the arsenic antisite [48] and probably is a complex of arsenic on a gallium site and an arsenic interstitial [49]. The structure of the EL2 defect is discussed in detail in Datareview 10.2 of this book. EL2 is of great importance in VPE GaAs technology because of its high concentration (typically up to 3 x 1014 cm"3 ) possibly resulting in semi-insulating material or in local variations of FET threshold voltages. The concentration of EL2 is influenced by dislocations [50] and by heat treatment after growth [51] but the major determining factor is the III/V ratio during growth. If a larger proportion of arsenic trichloride or arsine is used the concentration of EL2 increases [52-54]. However, the concentration is also a function of growth rate and substrate orientation [55]. Other commonly observed electron traps are EL5 traps ( N x = 1012 cm"3) characterised by emission behaviour similar to ELl (Nx = 1013 cm"3). A comprehensive listing is given in TABLE 7. The trap nomenclature is based on that used by Martin et al [55]. However, considerable care must be exercised with the alternative names quoted. The A and B centres in VPE [56] are not the same as the A and B centres seen in LPE [57] while Ruby et al [47] have named the centres they observed in VPE GaAs El, E2, E3 and E4 which appear to

correspond to ELl 1, EL5, EL3 and EL2, respectively. This is particularly confusing as the El - E5 notation has been used previously for traps induced by radiation [58]. TABLE 7 Name

Energy (meV)

T(K) at e = 100 s 1

Related to

Principal refs

EL2(A)

825

391

arsenic antisite

[55,44]

EB2 (ECl)

830

285

EL2

[57,60]

EL12 (A')

780

326

EB4 ?

[55]

EL3 (B)

575

280

reduces on anneal

[55]

EC2

480

265

Ni

[60]

ELl

430

219

EC2?

[61]

EL5(C)

420

211

reduces on anneal

[55]

EL16

370

310

[55]

EL8(D)

275

160

[56]

EL9(E)

225

134

[56]

EL2

190

113

[61]

EL3

180

104

[61]

ELIl(F)

170

122

Cu?

[56]

Energies are apparent activation energies with T2 correction measured in all cases by DLTS or its variants. Arrhenius plots of emission rate as a function of temperature are given in the references quoted. D2

Hole Traps

The hole trap population of VPE GaAs is very variable. Only HL9 seems to occur in most layers, usually in concentrations of the order of 1012 cm"3. HL2 is seen only in heat treated materials and is the same as the B state nearly always present in LPE GaAs. The other states are often undetectable. The association of HLl, HL3 and HL4 with Cr, Fe and Cu, respectively, is a result of diffusion and/or doping experiments. TABLE 8 Name

Energy (meV)

K

Notes

References

HLl

940

428

Gr

[63]

HLlO

830

361

HL2

730

347

heat treatment

[50,51]

HL9

690

309

common

[63]

HL3

590

301

Fe

[66,63]

[63]

HTl

440

222

HL4

420

225

HL6

320

156

[62] Cu

[64,65] [63]

Hasegawa et al [68] have studied the effects of the carrier gas used during growth. When they changed from helium to nitrogen the deep states associated with copper and iron increased very substantially, in the case of copper from 7 x 1013 to 5 x 1014 cm"3. They attributed the difference to impurities in the nitrogen gas itself. Other differences were also noted in particular in the detailed emission properties of EL2. D3

Recombination

In a study by Miller [52] it was shown that the minority carrier lifetime in n-type GaAs increased as the arsine to gallium chloride ratio was decreased over the range 3:1 to 1:3. The concentration of EL2 also decreased, so correlating with lifetime. Mitonneau rejects this latter hypothesis as he considers the hole cross-section of EL2 to be too small to account for the lifetime [67] and proposed HLlO as the important recombination centre. Partin et al [60] in a comprehensive study of VPE GaAs contaminated with nickel consider the state EC2 as the most likely lifetime killer in processed material. However, few of the observed states have reliable values for directly measured capture cross-sections and although it may be justified to make comments about recombination in specific layers, it is not yet possible to generalise about VPE GaAs. The overall situation is reviewed in [69]. E

DEFECT ENERGY LEVELS IN MOVPE GaAs

MOVPE gallium arsenide has a very variable deep state population but, like conventional VPE and bulk grown material, it always contains the deep electron trap EL2. This defect is associated with the arsenic antisite [70]. It is normally present in concentrations between 1012 and 1015 cm"3 dependent primarily on growth conditions [71]. Published work indicates that there are no other states which are always detectable although many are quite common. It is apparent that the concentration of many of the states depends on the source materials [72], and choice of substrate [83], and so is probably related to chemical impurities rather than lattice defects. A broad generalisation is that higher purities (both deep and shallow levels) are achieved with triethyl gallium [73] rather than the trimethyl alkyl and that the kinetics of low pressure growth results in a lower concentration of impurities being incorporated [74]. El

Electron Traps

Considerable effort has been expended in studying the effect of the V/III gas stream ratio on the trap population in MOVPE GaAs. Bhattacharya et al [75] observed a linear dependence of EL2 concentration on the vapour phase As/Ga ratio. However, Samuelson et al [76] observed a dependence on ([AsH3]/[TMG])1/2 over the As/Ga range 5 to 5Q. Zhu et al [77] gathering together data from various sources prior to 1981 showed that there were obviously many other factors influencing the EL2 concentration apart from the V/III ratio. A very comprehensive study by Watanabe et al [71] shows a clear one half power law at 720-7400C and a quarter power law at 630-6600C. Furthermore, Street et al [84] have observed an increase in the concentration of

EL2 for material subjected to rapid thermal annealing (RTA), whereas the concentration was unaffected by the introducing of indium doping (0 to 0.3% In) [85], a procedure which has been shown to reduce the dislocation density of both bulk and epitaxial material. Other factors such as shallow doping level, growth rate and substrate, all have some effect on the EL2 concentration. Systematic differences in the deep state emission Arrhenius plots were seen between the higher and lower growth temperatures, suggesting that the structure of EL2 may well vary [71]. Rapid thermal annealing (RTA) has become a common processing technique to activate dopants implanted into GaAs. Street et al [84] observe an increase in EL2 and two other electron traps with activation energies of 280 and 500 meV after RTA in the temperature range 850 to 900 0 C. Above 900 0 C these traps are accompanied by three other unidentified and unresolved electron traps at temperatures of approximately 100, 130 and 230 K (e = 4 s"1). They conclude that RTA contributes to an increase in arsenic vacancy complexes at the expense of gallium vacancy complexes. Low temperature epitaxial growth could offer advantages such as the more efficient utilisation of reactants, increase in nucleation sites, and high doping efficiency. Pande and Aina [86] have grown MOVPE layers of GaAs at low temperatures (< 4300C) by plasma-assisted growth at low pressure. A deep electron trap at Ec - 630 meV was detected at a concentration of 5 x 1013 cm"3 and is similar to a trap reported in MOVPE GaAs by Zhu et al [77]. They did not detect EL2 in their material, although analysis of the DLTS spectrum of the 0.63 eV band shows a very broad level throwing up the possibility that the EL2 level is masked by this feature. In any case the concentration of EL2 is less than 5 x 1013 cm"3. Auret et al [83] have investigated the effect that different substrate types (SI, n+ or p+) and substrate suppliers have on the defect concentration. EL2 was always present at a concentration of 1 x 1014 cm"3. Other defects were also present and their concentrations depended on substrate type and supplier. In particular electron traps labelled EO5, EO6 and EO8 were detected at high concentrations in layers grown on p+ substrates. Similarly an electron trap labelled EO2 was detected in layers grown on the n+ substrates. No electron traps other than EL2 were detected in layers grown on SI substrates, demonstrating the importance of substrate choice and efficient impurity gettering at substrate-epilayer interfaces. Defects formed by high-energy (1 MeV) electron irradiation of GaAs (El - E5) have been reported previously. NeI and Auret [87] have detected five similar levels (EBl - EB5) in layers that have electron-beam (EB) evaporated metal Schottky diodes. Their activation energies ranged from Ec - 45 meV to Ec - 560 meV. Au, Al, Pt, Cr and Ti were used, but all yielded the same defect levels which varied only in concentration as the EB current varied due to the different melting points of the metals. These results confirm that these defects are process and not material related. In TABLE 9, an electron trap which is often observed is labelled EH5. This is usually seen in the 1012 to 1014 cm"3 concentration range [72] in GaAs but is present in much higher concentrations in AlGaAs [78]. Profiles of the state in GaAs near a heterojunction with AlGaAs have shown dramatically increasing profiles with concentrations reaching 1016 cm"3 [10]. Another state, EH4, shows similar behaviour and appears as a shoulder on the EH5 DLTS peak. The energy quoted in the literature varies from 300 to 380 meV because of the problems of deconvolution. It is not clear whether these states are directly associated with aluminium, or with impurities in the

trimethyl aluminium. Allsopp et al [79] describes another state at 250 meV which exhibits similar behaviour. Zhu et al [77] have reported a number of defects which are detected near the interface region in homoepitaxy layers. Their origin is unknown but increased concentrations of inadvertent impurities have been observed near MOCVD interfaces using SIMS [80]. Another frequently seen state [76] is similar to EL14 that is observed in bulk material. Wagner et al [78] associate it with tellurium and observe that in GaAs:Te it is present typically at 0.01 of the tellurium concentration. Metastability has been reported in a variety of semiconducting materials. Buchwald et al [88] report two metastable defect configurations designated M3 (610 meV) and M4 (310 meV) in ntype GaAs grown by MOVPE. Metastable transformations between these two centres is found to be temperature induced and bias controlled. Capture cross sections are found to be temperature independent at 5.1 x 10"18 and 1.8 x 10'18 cm2. Superlattices have been used by Soga et al [89] as intermediate layers between Si substrates and GaAs epilayers. They detect two traps whose concentration decreased as the thickness of the grown GaAs layer increased. The levels were identified as EL2 (3.4 x io 14 cm"3 compared to 8.0 x 1013 cm"3 for GaAs/GaAs) and EB4 (Ec - 440 meV; 2.7 x 1014 cm"3 and not detected in GaAs/GaAs). The electron trap EB4 was induced previously when undoped VPE grown GaAs on Si-doped GaAs was deformed. In TABLE 9 the nomenclature of [81] is followed wherever possible with additions for later work. The notation Wl etc. is from Wagner [78], EW2 etc. from Watanabe [71], ETlMO from Zhu [77] and EHl etc. from Hashizume [72]. The energies quoted are apparent activation energies with T2 correction. Arrhenius plots are available in the references. TABLE 9 Name

Energy (meV)

M3

T(K)fore= 100 s"1

Notes

References

610

metastable

[88]

M4

310

metastable

[88]

EL2

785

393

EHlETlMO

[71,75,77]

ET3MO

520

230

EW2 (increase at

[71,77]

interface) ET4MO

470

231

EL4in[81]

[77]

Ell

440

264

low concentrations

[76]

EH5

380

214

W2EW3(A1)

[72,78,79]

ET6MO

370

158

interface only

[77]

ET5MO

350

107

interface

[77]

ET4

300-380

Wl (Al)

[72,79]

EH3

220

160

EL8?in[81]

[72]

EL14

190

152

W4(Te)

[76,78]

EH2

180

110

EL3in[81]

[72]

?

630

344

low temperature growth

[77,86]

EB4

470

E2

GaAs on silicon

Hole Traps

Very little work has been published on hole traps in MOVPE GaAs. The only state which is seen consistently is HL3, a level which has been associated with iron because of diffusion experiments. It is present in as-grown MOVPE material in concentrations between 2 x 1012 cm"3 and 2 x 1014 cm"3. Auret et al [83] have investigated the hole traps found in MOVPE GaAs layers grown on various substrates. The iron level HL3 is always detected and its concentration is found to increase from 2 x 1012 cm"3 for layers grown on SI substrates to > 1013 cm"3 for layers grown on n+ or p+ substrates. They also detected levels associated with Cu, HB4 and Cu(a). The concentration of HB4 increased from around 1011 cm"3 for layers grown on SI substrates to more than 1013 cm"3 on n+ and p+ substrates. The state HB4 has also been seen near interfaces [77]. A state at 340 meV has been associated with chromium through radio tracer diffusions into MOVPE but is present only in low concentrations in as-grown material [82]. TABLE 10 Name

Energy (meV)

T(K)fore= 100 s"1

Notes

References

Pl

310

254

HT4MO

[75]

HT5MO

330

176

HL5 HB5

[77]

340

150

Cr (diffused

[82]

MOVPE) P2

350

339

[75]

HT4MO

370

213

[77]

HB4

440

213

Calculated

[90]

HL3

570-590

299

HT3MO (Fe)

[71,77]

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S.C. Palmateer, WJ. Schaff, A. Galuska, J.D. Berry, L.F. Eastman [Appl. Phys. Lett. (USA) vol.42 no.2 (1983) p. 183-5] N. Chand, A.M. Sergent, J.P. van der Ziel, D.V. Lang [ J. Vac. Sci. Technol. B (USA) vol.7 no.2 (1989) p.399-404 ] A. Ito et al [ Semicond. Sci. Technol. (UK) vol.4, no.5 (1989) p.416-19] A.E. Lin, E. Omelianovski, R.H. Bube [ J. Appl. Phys. (USA) vol.47 (1976) p. 1852 ] A.M. Huber et al [J. Appl. Phys. (USA) vol.50 no.6( 1979) p.4022 ] R.H. Wallis [ Inst. Phys. Conf. Ser.(UK) vol.56 (1981) p.73 ] DS. Ruby, K. Arai, G.E. Stillman [ J. Appl. Phys. (USA) vol.51 (1976) p.3587-91 ] L Dobaczewski et al [ in Defects in Semicond., Mater. Sci. Forum (USA) vol. 143-174 Eds. H. Heinrich, W. Jantsch (TransTech, Zurich, 1993) p. 1001 ] HJ. von Bardeleben, D. Stievenard, J.C. Bourgoin, A. Huber [ Appl. Phys. Lett. (USA) vol.47 no.9 (1985)p.970-2] N. Yamamoto, F. Hasegawa, M. Onomura, Y. Nannichi [Jpn. J. Appl. Phys. (Japan) vol.24 no.5 (1985)p.L326-8] K. Saxena, N.P. Singh [ Indian J. Phys.(India) vol.59A no.2 (1985) p. 104-8 ] M.D. Miller, G.H. Olsen, M. Ettenberg [Appl. Phys. Lett. (USA) vol.311 (1977) p.538 ] M. Ozeki, J.Komeno, A.Shibatomi, S.Ohkawa [ J. Appl. Phys. (USA) vol.50 (1979) p.4808 ] Zou Yuanxi [ Inst. Phys. Conf. Ser (UK) no.63 ch.4 (1982) p. 185 ] G.M. Martin, A. Mittonneau, A.Mircea [ Electron. Lett. (UK) vol. 13 no.7 (1977) p. 191-3] A. Mircea, A. Mitonneau [ Appl. Phys. (Germany) vol.8 (1975) p. 15-21 ] D.V. Lang, RA. Logan [J Electron. Mater. (USA) vol.4 (1975) p. 1053-66 ] D.V.Lang, L.C. Kimerling, S.Y. Leung [J. Appl. Phys. (USA) vol.51 (1976) p.3587-91 ] A. Ashley, G.G. Roberts, DJ. Ashen, J.B. Mullin [ Solid State Commun. (USA) vol.20 (1976) p.61-3] D.L. Partin, J.W. Chen, A.G. Milnes, L.F. Vassamillet [ J. Appl. Phys. (USA) vol.50 no. 11 (1979) p.6845-59 ] H. Lefevre, M. Schulz [ Appl. Phys. (Germany) vol. 12(1977) p.45 ] K. Sakai, T. Ikoma [ Appl. Phys. (Germany) vol.5 (1974) p. 165-71 ] A. Mitonneau, G.M. Martin, A. Mircea [ Electron. Lett. (UK) vol.13 no.22 (1977) p.666-8] A. Mitonneau, G.M. Martin, A. Mircea [ Inst. Phys. Conf. Ser (UK) 33A (1977) p.73 ] A. Chantre [ Phys. Rev. B (USA) vol.23 (1981) p.5335 ] M. Kleverman, P. Omling, L-A. Ledebo, H.G. Grimmeiss [ J Appl. Phys. (USA) vol.54 no.2 (1983) p.814-9] A. Mitonneau [ Rev. Phys. Appl. (France) vol. 14 (1979) p. 853 ] F. Hasegawa, T. Yamata, T. Yamamoto, A. Kaokita, H. Seki [ J. Appl. Phys. (USA) vol.24 no.8 (1985) p. 1036-42] A.R. Peaker, B. Hamilton [ Chemtronics (UK) vol.3 no.4 (1988) p. 194-200 ] HJ. von Bardeleben, D. Stievenard, J.C. Bourgoin, A. Huber [ Appl. Phys. Lett. (USA) vol.47 no.9 (1985)p.970-2] M.O. Watanabe, A. Yanaka, T. Udagawa, T. Nakanishi, Y. Zohta [ Jpn. J. Appl. Phys. (Japan) vol.22 no.6 (1985) p.923-9] T. Hashizume, E. Dceda, Y. Akatsu, H. Ohno, H. Hasegawa [ Jpn. J. Appl. Phys. (Japan) vol.23 no.5 (1984)p.L296-8] N. Kobayashi, T. Fukui [ Electron. Lett. (UK) vol.20 (1984) p.887 ] T.F. Kuech, R. Potemski [Appl. Phys. Lett. (USA) vol.47 no.8 (1985) p.821-3] P.K. Bhattacharya, J.W. Ku, S.J.T. Owen [Appl. Phys. Lett. (USA) vol.36 no.4 (1980) p.304 ] L. Samuelson, P. Omling, H. Titze, H.G. Grimmeiss [ J. Cryst. Growth (Netherlands) vol.55 (1981) p. 164-72] H.-Z. Zhu, Y. Adachi, T. Ikoma [ J. Cryst. Growth (Netherlands) vol.55 no. 1 (1981) p. 154-63 ] E.E. Wagner, D.E. Mars, G. Horn, G.R. Stringfellow [ J. Appl. Phys. (USA) vol.51 (1980) p.5434] D. Allsopp, A.R Peaker, EJ. Thrush, G. Wale Evans [ J Cryst. Growth (Netherlands) vol.68 no. 1

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10.6 Defect energy levels in electron and neutron irradiated, photon damaged and ion implanted GaAs A.R. Peaker June 1996

A

INTRODUCTION

Extensive studies of electron damage in GaAs were undertaken prior to 1980. This work was seen primarily as a route to understanding defect reactions in this material. Although some issues related to electron damage are still the subject of investigation, in recent years attention has been directed more towards issues of immediate device importance. Among these are ion implantation damage related to both light and heavy ions and radiation damage in gallium arsenide solar cells, FETs and radiation detectors. The techniques of ion implantation and annealing are reviewed in Chapter 20 of this book while radiation detectors are discussed in Datareview 23.2. The substantial volume of work on electron irradiation indicates that the introduction rate and the defect energies of many of the traps produced are essentially unchanged with the type of material (LPE, VPE or bulk) and normal levels of doping. This suggests that, in relation to these traps, the influence of impurities is negligible and strongly supports the view that the defects generated in this way are predominantly intrinsic. Although there is far less evidence in relation to proton damage and ion implantation, it is believed that the situation is similar in this respect, i.e., the defects are predominantly intrinsic. B

ELECTRON DAMAGE

Lang [1] has reviewed much of the early work on carrier removal rates (compensation) in n-type material, most of which has been done using 1 MeV electrons. These studies show a linear dependence on fluence and indicate that the dominant species are acceptor-like. Indeed p-type conversion has been observed and attributed to a level 100 meV from the valence band [2], although other acceptors are also produced in high concentrations. Bl

Electron Traps

Lang et al [3] studied the orientation dependence of the defect production rate using 1 MeV electrons. He concluded that the states resulting from room temperature irradiation were associated with gallium displacement. These were EBlO, EB9 and EB6 with similar but less definitive conclusions for EB4 and EB3. These states are all annealed out at 500 K. More recent work using electron energies nearer the displacement threshold leads to the conclusion that the defects result from a single arsenic displacement (not gallium) and not from the displacement of two or more neighbouring atoms [4,5]. The discrepancy with Lang's results has been explained in terms of the recoil effects of the high energy electrons used by Lang [3] and it is almost certain that the defect reaction is indeed an arsenic displacement. Although, as stated previously, the majority of work in the deep level population does not report significant differences between material types after irradiation, recent work using 7 MeV electron damage indicates there are differences between MBE and MOVPE with high background doping [6].

A summary of the electron traps created is given in TABLE 1. TABLEl Name

Energy (meV)

T (K) for e = 100 s

Irradiation temp. K

1

Introduction

Ref.

1

rate (cm' )

EB3(E5)

900

352

300

0.1

[3]

EL2(I1)

720

383

570

1

[8]

P3(EL2?)

720

343

after anneal

-

[7]

EB4(E4)

710

320

300

0.08

[3]

12

620

348

570

0.5

[8]

13

500

275

570

below 0.01

[8]

P2(I3?)

500

298

after anneal

-

[8]

EB6(E3)

410

199

300

0.7

[3]

14(PI?)

360

204

570

0.03

[8]

Pl (EL5)

360

206

after anneal

-

[7]

15

260

172

570

0.03

[8]

E (0.23)

230

134

80

0.015

[10]

16

190

119

570

below 0.01

[8]

EB9(E2)

180

-

300

2.8

[3]

17

160

83

570

0.03

[8]

EBlO(El)

40-80

-

300

1.8

[3]

The electron energy was 1 MeV in all cases but the defect production rate is only a rough guide as the rate is sometimes dependent on flux which was not the same in all cases. However EB9 (E2) has the highest introduction rate by far and appears to be acceptor-like and tends to compensate n-type material. The energy quoted is the apparent activation energy after T2 correction. Defects EB6, EB9 and EBlO could be produced by 80 K as well as 300 K irradiation so are either primary defects or the result of interactions with species mobile at low temperature [10]. B2

Hole Traps

Less work has been done on hole traps and the results are much more variable particularly in respect of introduction rates.

TABLE 2 Name

Energy (me V)

T (K) for e = 100 s"1

Introduction rate (cm>)

Ref.

L13(H5)

940-960

410

0.48

[11,12,13]

L16(H3)

710-790

356

0.22

[11,12]

H2

370

-

-

[13]

HB6(H1,L14)

250-290

-

0.07-0.7

[6,11,14,12]

HO

90-115

-

0.5-2.4

[11,6,14]

60

0.8

[11]

50

1.6

[11]

All irradiations were at 300 K with 1 MeV electrons. A number of other shallow traps have been seen associated with very high fluences. These tend to be sample dependent [11] and are not included above. The concentration of H5 saturates with fluence at around 3 x 1015 cm"3 and is believed to be a complex with impurities in the starting material. Hl and HO appear consistently and Hl is believed to be a simple primary defect. Both HO and Hl show the same anisotropic dependence as El, E2 and E3 and so are also associated with the arsenic sublattice. B3

Annealing Behaviour

The annealing stages of defects induced by electron irradiation have been reviewed by Lang [I]. Essentially, after room temperature irradiation levels E3, E5 and the hole trap Hl have identical annealing behaviour exhibiting first order kinetics with an activation energy of 1.4 eV. El and E2 have an activation energy for annealing of 1.75 eV [16]. E4 is apparently stable. The annealing rate of the second pair of defects depends on Nd2/3 where Nd is the pre-irradiation donor level. This implies that the donor atoms are acting as sinks for the defects during the anneal phase. However, both these groups of defects exhibit recombination enhanced annealing resulting in their removal at lower temperatures when minority carriers are generated by optical or electrical excitation and subsequent recombination [17,16]. Work by Pons et al [7] showed that E4 anneals out, generating the new states Pl, P2 and P3 documented in the electron trap table. They also consider that the annealing of E2 proceeds along two concurrent paths and identify E2 as a vacancy. C

PROTON, NEUTRON AND ION IMPLANTATION

Ion implantation has become a critically important part of GaAs technology. Unfortunately, it is never possible to anneal out all the implant damage. This is in stark contrast to silicon where it is usually possible to restore the material to near perfection by optimising the implant/anneal schedules. In GaAs a significant population of residual point defects remain after all implantanneal schedules.

In addition to the creation of defects specifically associated with radiation the concentration of EL2 is often substantially increased, however it is not believed this is the dominant factor in the removal of carriers at least in the case of lighter ions [19]. The situation is however quite complex as Tan et al [20] observed that 'EL2' produced under these circumstances have rather different annealing characteristics to normal EL2. Auret et al [21] propose that in He implanted material the dominant source of compensation is the state Ea3. Some work has been done on defects produced by quite low energy beams (< 1 keV). Kosugi et al [22] observe defect formations far beyond the theoretical ion range but consisting predominantly of species that readily anneal out. In contrast, Zolper et al [23] have studied the formation of thermally stable deep states with the technological aim of providing stable semi-insulating material. They have used nitrogen implantation to achieve this aim. TABLE 3. GaAs electron traps induced by proton and heavy ion implantation. Defect label and ref.

Ea(meV)

T (K) for e=100s" 1

Irradiation and measurement technique

L4p

[24]

990

372

HB annealed, 2 meV protons, DLTS

L3Ar

[24]

960

*352

HB, 300 K, 60 keV, Ar ions, DLTS

L3p

[24]

870

343

HB, 300 K, 2 MeV protons, DLTS

B(A)

[25] [26] [27]

830 780 710

415 382 326

HB, 400 keV Se+, annealed, DLTS LPE, neutron, annealed, DLTS CZ, 200 keV Si+, annealed, DLTS

L2p

[28] [27]

680 670

339

[24]

610

307

CZ, 200 or 400 keV Se+, annealed, DLTS, 185 keV O+, similar defect produced by 1.8 MeV He+ or 400 keV protons, DLTS, broad peak HB, 300 K, 2 MeVprotons, similar defect produced with 60 keV Ar ions, DLTS

[26] [27] [29] [30] [24]

560 550 500 500 430

309 287

[26] [31] [30] [25] [32] [31] [31] [29]

380 320 300 210 200 200 130 110

262 217

LIp

208

177

LPE, neutrons, annealed, DLTS CZ, 200 keV Si+ or Se+, DLTS neutron irradiation, Hall 0.3 He and Ge implants, TSC HB, 2 MeV protons, 300 K, similar defects produced with 600 keV Ar ions, DLTS LPE, neutrons, annealed, DLTS HB, 400 keV Se+, annealed DLTS PIB, He and Ge implants, TSC HB, 400 keV Se+, annealed, DLTS 300 keV protons, DLTS HB 400 keV Se+ neutron irradiation, Hall

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10.7 Analysis of LEC GaAs by near-IR mapping M.R. Brozel July 1996

A

INTRODUCTION

The analysis of semi-insulating GaAs by near-infrared absorption imaging has developed and expanded into three major fields, defect absorption, light scattering and striation mapping. In this Datareview, we also briefly mention Franz-Keldysh absorption imaging.

B

NEAR-INFRARED ABSORPTION BY EL2 CENTRES

Classically, this has been the primary method for the mapping of EL2 centres in SI LEC GaAs. Various workers have taken up EL2 imaging as a tool and used it with other techniques to establish correlations between dislocations, high concentrations of EL2 centres, electrical data and MESFET parameters. There is now such an extensive literature on this subject that this author feels that only reference to the proceedings of three specialized conferences can do justice to the relevant workers [1-3]. However, some conclusions can be made. High values of neutral EL2 concentration are associated with dislocations in as-grown LEC GaAs [4-6]. In small research crystals where dislocations are not present, absorption due to EL2 centres is still present, although the absorption corresponds to a relatively low concentration of 5 x 101V cm3 compared with around 1.2 x 1016/ cm3 in larger dislocated crystals proving that dislocations are not necessary for the presence of EL2 centres [7]. Dislocations can act as sources of arsenic anti-site defects, AsGa, which are a precursor of EL2 centres, by non-conservative climb processes [8-10]. However, the major part of the close association of EL2 centres with dislocations is probably a result of point

FIGURE 1. Infrared absorption image due to EL2 defects. The 3" diameter {100} SI GaAs wafer is 650 microns thick. The (110) directions are close to vertical and horizontal. The image was taken at room temperature with light of 1 micron wavelength.

defect diffusion towards dislocations by reason of their strain field [11-13]. An image of EL2 distributions is shown in FIGURE 1. Dark areas correspond with high [EL2] and are associated with the presence of a high concentration of dislocations. This image is of older material (early 1980s) where EL2 defects in substrates from many suppliers were not completely homogenized by suitable post-growth heat treatments. In as-grown, commercial LEC GaAs, there is a correlation between electrical properties, such as resistivity and mobility, and dislocations [14-16], and MESFET properties, especially threshold voltage, and dislocations [17-20]. Commercial LEC GaAs is ingot-annealed resulting in a drastic improvement in the uniformity of EL2 absorption images [21-23]. Concurrently, there is a consistent improvement in electrical uniformity and MESFET uniformity [21,24-28]. However, it appears to be very difficult to render [EL2] completely uniform and cell structure, associated with polygonized dislocations, is usually still detectable. The dislocation structure is apparently unaffected by the heat treatments and this suggests that it is the presence of a non-uniform concentration of deep levels - including and probably dominantly EL2 - rather than dislocations themselves, that causes variations in MESFET properties in unannealed GaAs. After annealing, variations in the concentration of neutral EL2 centres are small and specialized equipment must be used to map concentrations [13,29,30]. Quenching of GaAs from high temperatures has been shown to reduce [EL2] [31-33] but a subsequent anneal allows [EL2] to recover [31-34]. Imaging studies have now shown that such treatments applied to entire ingots render [EL2] images very uniform, and that other parameters, such as electrical resistivity and carrier mobility, are also very uniform [35,36]. The absorption spectrum of EL2 has two components, one due to excitation of electrons into the conduction band from neutral EL2 centres, and one due to excitation of electrons from the valence band to ionized EL2+ centres. The spectral forms are quite different although they are both limited to the near-infrared. Their optical cross-sections have recently been determined [37,38] and because there exist wavelengths where either the neutral EL2 or EL2+ absorption coefficients are dominant, it is possible to selectively map the concentrations of neutral and EL2+ centres across wafers [39,40]. By using this technique it has been demonstrated that the nonuniformity previously observed in the concentration of neutral EL2 centres is a result of the variation in total [EL2] and not in the variation of acceptor concentration [40]. Concentrations of EL2+ can also be mapped using magnetic circular dichroism absorption (MCDA) [41]. However, the absorption mechanisms are different and there remains a factor of 2 between [EL2+] measured by each of these techniques [42]. Unfortunately, claims of similar accuracy are made for each technique and it remains uncertain why these calibrations are different [43]. As an assessment technique, EL2 imaging is probably of little current use as ingot and wafer anneals are now optimized, the resultant EL2 concentrations being well homogenized and in the range of 1 to 1.5 x io16 cm"3. This concentration is found in virtually all bulk SI GaAs regardless of the growth technique and wafer supplier.

C

NEAR BANDEDGE ABSORPTION IMAGING

This section is devoted to reverse contrast imaging, the imaging of near bandgap light transmitted through a cooled (< 150 K) sample. This type of mapping produces an image that has the opposite contrast to EL2 imaging, to which it is related only by using similar equipment [5,44,45]. It has been demonstrated that reverse contrast reveals dislocations and dislocation arrays with great sensitivity and in a non-destructive way [46]. An image of reverse contrast in a 500 micron thick sample at a temperature of 80 K using light of 830 nm wavelength is shown in FIGURE 2.

FIGURE 2. Reverse contrast images (top) and EL2 absorption images (bottom) of regions taken near the periphery of a wafer (showing lineage structure), left, and from the centre showing dislocation cell wall structure, right. Each image corresponds to approximately 2 mm x 2 mm. The reversal of contrast is evident.

Like FIGURE 1, dark regions are more absorbing and correspond to increased concentrations of the absorbing defect. In general, high concentrations of these defects (RC defects) are found away from dislocation arrays, the opposite of what is expected by the standard model of atmosphere creation, 'Cotterell Atmospheres', and the opposite of what is observed for EL2 distributions. However, in the central regions of larger dislocation cells there tends to be a reduction in absorption, a result that does not correlate with EL2 images. Note also that this image was obtained on material that had undergone a standard ingot anneal, a treatment that had partially homogenized [EL2] but has not been effective in smoothing out concentrations of RC defects. Observations that RC images are very similar to luminescence images obtained from the same

material [47] and that both were inhomogeneous even in material where [EL2] is practically uniform led to the idea that the same defects control both the RC absorption and the luminescence efficiency. In other words, it was suggested that the RC defect was the dominant non-radiative recombination centre in bulk SI GaAs [47], high concentrations of RC defects leading to increased absorption and reduced luminescence efficiency. This conclusion was supported by the observation that both RC absorption and the PL efficiency could be photo-quenched and recovered. The bleaching of RC absorption by photoquenching is associated with an increase in PL efficiency [48]. This is expected if the absorption and the minority carrier recombination were associated with the same charge state of the defects. The spectral form of the RC absorption, that the absorption is limited to a narrow spectral range within - 50 meV below the bandgap energy, led to the speculation that shallow acceptors were involved. However, early measurements failed to show any dependence of RC absorption on acceptor concentration [45]. However, photoconductivity measurements indicated the presence of hole conductivity in the same spectral range indicating an electron transition from the valence band to a very deep acceptor state [49-51]. In a series of measurements the presence of this deep acceptor level at about Ec-50 meV was demonstrated with photo-excited Hall effect measurements on lightly n-type GaAs. In this material, the electron concentration could be increased by the same photoquenching light that bleached the RC absorption. Moreover, the recovery to the original lower carrier concentration state occurred at the same temperature as the RC absorption recovery [52]. This measurement showed a linear relationship between the concentrations of photo-quenchable acceptors and the RC absorption coefficient and allowed a cross-section for the absorption to be estimated. The RC level is thus demonstrated to be an acceptor with an ionization energy approximately 50 meV below Ec and is photoquenchable. The subsequent use of positron annihilation showed that this electronic level is to be associated with the arsenic vacancy, V^. This measurement also used photo excitation and demonstrated that the increase in positron lifetime associated with the vacancy could be photoquenched and recovered under the same conditions as for RC absorption [53]. This result is consistent with the stoichiometry dependence of the RC absorption which is found to increase in Ga-rich material [54]. There remains the question of why high concentrations of V^ ( - I x 1015 cm"3) should exist away from dislocations in material that contains arsenic rich, EL2 centres. This was addressed in terms of the disruption to the reverse Frenkel reaction on the As sub-lattice by dislocations as the crystal cools [55]. It was argued that this causes As interstitials to be trapped by the dislocation strain field in the form of atmospheres and does not allow them to recombine with V^ uniformly distributed in the lattice. This simple model was able, in part, to explain the presence of As precipitates, EL2 centres and V^ in bulk SI GaAs, although it is acknowledged that such a simple model cannot account for all the phenomena seen in this material [55]. D

INFRARED LIGHT SCATTERING FROM MICROPRECEPITATES

This near-infrared mapping technique is sensitive to second-phases (precipitates) that appear to be ubiquitous in SI GaAs but not in doped material. The presence of micro-precipitates of diameter < 1 micron was established some time ago by TEM which identified them as hexagonal arsenic [56-59]. They are also often observed by etching [20,60,61]. Most of these precipitates are spatially associated with dislocations but are too few to be routinely examined by TEM. They

may be mapped in a 'non-destructive' way by light scattering tomography [62-64] or laser scanning microscopy [65]. In light scattering tomography (LST)3 infrared laser light scattered by precipitates is detected by a sensor adjacent to the sample but normal to the laser beam. An infrared sensitive TV camera [62-64] or a linear array of photo-diodes [66] is often used to detect the scattered light. The laser beam is scanned to produce a full 2 dimensional image in each case. Laser scanning microscopy is an extremely sensitive transmission measurement where an incident infrared laser beam is focussed to a diffraction limited spot and the transmission measured. Scattering centres within the beam reduce the transmission. The sample must be scanned to produce a raster in order for an image to be built up. Both these techniques produce well detailed images of precipitate distributions and both have observed the effects of various heat treatments on the dissolution and re-precipitation of arsenic [57,63,66-71] (FIGURE 3).

FIGURE 3. LST image from the scattering of light from precipitates (probably As) associated with dislocation cell walls in an SI GaAs wafer. There is an excellent correlation between the presence of these precipitates and dislocations in SI GaAs. The light region in the centre of the cell is due to scattering from micro-structure whose identity is unknown (courtesy of J-P. Fillard and P. Torrut-Gall, LINCS, Montpellier, France).

The use of dark field infrared microscopy has been reported to reveal these precipitates, but probably at much lower sensitivity [72] E

DOPANT UNIFORMITY IMAGING

The optical mapping of dopant striations and dopant distributions should be mentioned. This type of imaging depends on variations in the real part of the refractive index and is a result of small angle optical scattering rather than absorption. In this respect, it is a Schlieren imaging technique. These techniques have been applied to GaAs and InP [73-75]. It should be emphasized that these

techniques are suitable only for moderately doped GaAs with a dopant concentration of greater than 1017 /cm3 and in most instances the angle between the growth striations and the viewing direction is critical. Because the image is dependent on the real part of the refractive index, it is possible to image growth striations in iso-electronically doped GaAs (e.g., In-doped GaAs)! An image of a {110}, 1 mm slice cut from the shoulder region of a <001> axis, Te-doped (-10 18 cm*3) LEC GaAs ingot is shown in FIGURE 4. The seed is towards the top of the image. Recent reports on n-type GaAs suggesting that the image is due to absorption (i.e., the imaginary part of the refractive index) are probably incorrect as free carrier absorption in the spectral region where the image is obtained, near 1 micron, is very weak.

FIGURE 4. Growth striations revealed by small angle optical scattering imaging [73]. Striations representing the progress of the growth interface shape can be seen. Note the abrupt change just below the seed region where growth was restarted following a partial melt-back. The strong dark diagonal line is a fracture in the sample.

F

FRANZ-KELDYSH ABSORPTION IMAGING

The Franz-Keldysh effect results in a reduction of bandgap energy Eg on the application of an electric field (see Datareview 6.2). The change in Eg can be determined optically as amove of the bandedge absorption to lower energies. The use of optical absorption to probe the bandedge absorption will be sensitive to this change and it is possible to form an image, representing a mapping of the electric field in the material. Because the resolution limit of infrared imaging techniques is of the order of the wavelength of light, it is not possible to perform this measurement in conducting semiconductors because the depletion widths are too small. The technique has found one major application, nuclear particle detectors (see Datareview 23.2). In some particle detectors, contacts are placed on either side of a 500 micron thick SI GaAs substrate. One of the contacts is a Schottky (blocking) contact; the other is ohmic. The application of a reverse bias creates an electric field in the GaAs. This field is not uniform and results in inefficient operation of the detector.

The field can be mapped by viewing the transmittance of light at 900 nm wavelength (for a sample temperature of 300 K) through the edge of the detector, parallel to the (100) faces and orthogonal to the field [76]. The light is detected using a CCD Si video camera and it is necessary to employ image enhancement. A series of images is shown in FIGURE 5. The penetration of the field is quite unexpected being linear with respect to applied voltage with a penetration of - 1 jim/V. Further analysis reveals that within the high field area, the electric field is nearly constant at -10 6 V/m. Interpretation of this electric field distribution is beyond the scope of this Datareview, but readers can refer to [77].

FIGURE 5. A series of Franz-Keldysh absorption images (900 nm and 300 K) through a 500 jim thick detector. The Schottky contact is towards the left hand side of each image. Applied voltages are shown. The dark, high field region penetrates the detector linearly with applied voltage.

It should be noted that small changes in the temperature of the sample produce equal optical effects, i.e., a reduction in Eg. Care must be taken to control the temperature when attempting this measurement. G

CONCLUSION

Infrared imaging in its various forms remains a powerful, non-destructive tool for the assessment of GaAs materials, providing spatial information in a rapid and near-quantitative way.

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XZ. Cao, A.F. Witt [ J. Cryst. Growth (Netherlands) vol. 112 (1991) p.838 ] M.R.Brozel [Mater. ScL Monogr. (Netherlands) vol.44 (Elsevier, 1987) p.225 ] M.R.Brozel [ Inst. Phys. Conf. Ser. (UK) no.91 (1987) p. 117-20] R.L.Barns [ J. Electron. Mater. (USA) vol. 18 (1989) p.703 ] K. Berwick, M.R. Brozel [ Inst. Phys. Conf. Ser. (UK) vol.146 (1995) p.635-40] D.S. McGregor et al [Nucl. Instrum. Methods Phys. Res. A (Netherlands) vol.343 (1994) p.527]

10.8 Passivation of defects in GaAs by hydrogenation

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SJ. Pearton June 1995

A

INTRODUCTION

Atomic hydrogen is found to passivate the electrical activity of all shallow dopant impurities, both donors and acceptors, in GaAs [1,2]. Many of the common deep level states, including EL2, are also passivated. The passivation is due to formation of neutral dopant- or defect-hydrogen complexes. In p-type GaAs hydrogen complexes have characteristic vibrational bands, around 2000 cm"1 for stretching modes and 800 cm"1 for wagging modes. The dopants are reactivated by annealing at < 4000C in bulk samples, or at < 2000C in diode structures where strong electric fields act to reduce retrapping effects. Unintentional incorporation of atomic hydrogen in GaAs can occur during many of the individual processing steps needed for device fabrication. These include plasma etching or deposition, high temperature annealing, wet chemical etching, implantation and surface cleaning and passivation. Direct evidence for hydrogen permeation during these processes has been obtained in some cases using deuterated chemicals or gases and secondary ion mass spectrometry (SIMS) profiling. In all cases, near-surface shallow dopant passivation is observed as a result of formation of neutral dopant-hydrogen complexes. In general atomic hydrogen passivates both shallow donors and acceptors in all III-V materials, though the efficiency of the passivation is often much more complete for one type of impurity over the other type. This is usually ascribed to the charge state of the hydrogen and hence the position of its energy levels within the bandgap. When the atomic hydrogen is either positively (H+) or negatively (H") charged, it will undergo a strong Coulombic attraction with ionized dopant impurities of the opposite charge, so that if hydrogen is in a negative (positive) charge state in an n(p)- type semiconductor, one should observe strong pairing and efficient passivation. In GaAs it is observed that all shallow donors (SiGa, GeGa, SnGa, SAs, Se^, and Te^) and shallow acceptors (BeGa, Mg033 Zn033Si^ Ge^, CM) are passivated by atomic hydrogen and in most cases the vibrational lines of the neutral complexes have been reported. In p-type GaAs, bias-drift experiments have provided strong evidence for the existence of a donor level for hydrogen (H+ZH0). There is also evidence for an acceptor level in n-type material. B

METHODS OF INCORPORATION

Hydrogen is generally introduced deliberately either by exposure to a low-pressure plasma at 100 - 350 0 C for several hours, or by direct implantation. FIGURE 1 shows secondary ion mass spectrometry (SIMS) profiles of deuterium in p-GaAs of two different initial doping levels after exposure of the samples to a 13.56 MHz, 0.05 torr D2 plasma. Deuterium is a readily available isotope of hydrogen that diffuses at a rate of >/2 that of hydrogen due to its larger mass. The incorporation depth is quite large in both cases, and a higher concentration is achieved in the p+ material since the effective solubility is strongly dependent on the number of sites to which

DEUTERIUM CONCENTRATION (cm"3)

D PLASMA

DEPTH (/Am)

D PLASMA

DEPTH

(^m)

FIGURE 1. SIMS profiles of D in plasma-exposed GaAs of different doping levels,

hydrogen can bond (i.e. the acceptor concentration). In fact lower frequency plasmas (< 100 kHz) appear to be more effective in producing larger incorporation depths due to their higher average ion energies [3]. The H+ ions are implanted into the near-surface region beyond the zero-bias depletion depth and can diffuse into the bulk. At higher plasma drive frequencies the lower ion energies can lead to a higher apparent flux at the immediate surface, with a greater likelihood of the formation of extended hydrogen clusters which are immobile. Hydrogen implantation is widely employed in GaAs device technology for isolation or current-guiding applications. The main carrier compensation mechanism is this case is the creation of midgap states which act as traps for both electrons and holes. Generally the highest resistivity is achieved for proton implanted material that is subsequently annealed at 300 - 450 0 C, and this can lead to diffusion of hydrogen into adjoining active regions of the device where it can produce dopant passivation [4]. Hydrogen can also be incorporated during annealing in molecular gas at > 450 0 C or in AsH3 or PH3 at similar temperatures. The indiffiision of hydrogen during elevated temperature treatments can obviously be prevented by using an inert gas ambient such as He or Ar. The presence OfH2 in the annealing ambient also affects the outdiffiision of hydrogen already present in GaAs, leading to much higher apparent thermal stabilities for hydrogen retention when H2 is part of the ambient [5]C

SHALLOW DOPANT PASSIVATION

Incorporation of atomic hydrogen into GaAs causes a reduction in carrier concentration in the near-surface region, with the depth of this modification generally being inversely dependent on the initial doping level. As an example FIGURE 2 (top) shows a strong decrease in doping level

DONOR CONCENTRATION (cm 3 ) ACCEPTOR CONCENTRATION (cm 3 )

AFTER H PLASMA (1h, 175 C)

INITIAL

Ai=TER H PLASMA (1h, 175°C)

DEPTH (urn)

FIGURE 2. Depth profiles of the shallow dopant concentration in n-type, Si-doped GaAs (top) or p-type, Si-doped GaAs (bottom) after hydrogenation (1 h, 1750 C) and subsequent bias application (VR = 10 V) at elevated temperature.

of both n-type and p-type GaAs after exposure to a hydrogen plasma for 1 h at 1750 C. If these samples are subsequently heated, the initial doping profiles are recovered at -400 0 C where the neutral dopant-hydrogen complex dissociates. The hydrogen does not physically leave the GaAs until -600 0 C. If a Schottky diode structure is fabricated on the hydrogenated material and it is reversed-biased, then the dopant reactivation occurs at lower temperatures as shown in FIGURE 2 because retrapping of the charged hydrogen and the ionized dopant ion is prevented. TABLE 1 shows the dissociation energies for different dopant-hydrogen complexes in GaAs5 determined in diode structures. All of the acceptor species in GaAs5 namely Be5 Mg5 Zn5 Cd5 C and Si^5 are passivated by association with atomic hydrogen. For the Ga-site acceptors the passivation mechanism is the formation of an As-H bond, leaving the acceptor three-fold coordinated. Under the assumption that hydrogen is in a positive charge state in p-type GaAs with a donor level in the gap, one can write for the passivation of Zn acceptors ( A s - Z n 0 ^ H + - (Zn 0 4 -As-H) 0

(1)

TABLE 1. Dissociation energies (ED) and attempt frequencies (V 0 ) of dopant-H complexes in GaAs [15]. Dopant

Character

E D (eV)

Vn(S-1)

Si08

donor

1.25 ±0.05 1.2 ±0.1 1.2±0.1

2 x 1013 6 xlO 7 1.3 xlO 14

Si^

acceptor

1.45 ± 0.10

1 x 1013

Sn0,

donor

1.20 ±0.10

4 x 1013

Se^

donor

1.52 ±0.05

2 x 1013

Zn011

acceptor

1.25 ± 0.10

5 x 1013

Zn011

acceptor

1.33

10 4

C^

acceptor

1.35 ± 0.05

2 x 1013

Be 08

acceptor

1.15±0.10

1.5 x 1013

Cd08

acceptor

1.35 ±0.10

1013

Zn08

acceptor

1.3

1 x 1014

The hydrogen is in a bond-centred site under these conditions. In the special case of a Si^ acceptor in GaAs, there appears to be the formation of a Si-H bond with a hydrogen atom located between neighbouring As and Ga atoms, i.e. still in a bond-centred site. The passivation of Si^ acceptors can be written (Ga - SiJ" + H + - (Ga - Si^ - H)0

(2)

Much of what has been learned about the structures of the hydrogen-dopant complexes comes from vibrational spectroscopy of the centres. For the Ga-site acceptors the H- stretching frequencies are characteristic of As-H bonds. This supports the model with hydrogen at a bond-centred position between the acceptor and a neighbouring As. For As-site acceptors such as Site H is again suggested to be bond centred. The trend in the hydrogen stretching frequencies for the different group IV acceptors (Si, C, Sn) strongly suggests that it is attached directly to the acceptor atoms. TABLE 2 shows a compilation of the vibrational modes for dopant-hydrogen complexes in III-V materials [5,6]. The frequency of the H-stretching bands for acceptors is close to that of As-H stretching modes in molecules (2116 cm"1 for AsH3). The stretching frequencies are reduced in wavelength by a factor of -1.37 for deuterium-acceptor complexes, confirming the presence of hydrogen in these centres. Theoretical calculations generally agree on the properties of hydrogen in these complexes [8-12], with H being present as H+ in p-type material and H" in n-type GaAs, forming H2 molecules at tetrahedral sites and other defects like H2* (one hydrogen in a bond-centred position and one in an adjacent tetrahedral position). For donor-hydrogen complexes, the bonds are at lower frequencies than the stretching modes of the acceptors, indicating a relatively weak bond to the hydrogen. For the case of Si, there are distinct sidebands observed to the low frequency side of

the strong band at 1717 cm"1 whose relative intensities are consistent with the isotopic abundances Of29Si and 30Si. These sidebands indicate that the hydrogen is attached directly to the Si atom which is involved in the vibrational motion TABLE 2. Hydrogen vibrational frequencies of acceptor-hydrogen and donor-hydrogen complexes in GaAs measured near LHe temperature [5-7]. Acceptor

Frequency (cm 1 )

Character

Be 051

2037.1

stretch

Zn011

2146.9

stretch

Cd 0 8

2206.7

stretch

12

C^

2635.1

stretch

13

C^

2628.4

stretch

Si^

2094.7

stretch

Si^

972.2

Ge^

2010.3

stretch

28

Si Ga

1717.3

stretch

28

Si 0 ,

896.8

29

Si 0 ,

1716.9

stretch

30

Si 0 ,

1716.5

stretch

Sn 0 ,

1327.8

stretch

Sn 0 ,

746.6

Se^

1507.5

stretch

SeM

777.9

wag

wag

wag

wag

A wagging mode is also observed for the SiGa-H complex and has three times the intensity of the stretching band, which is consistent with the expected double degeneracy of the transverse wagging vibration. In all cases when wagging modes are observed the complexes are believed to have hydrogen in an antibonding configuration. Furthermore, the presence of a doubly degenerate mode implies trigonal symmetry or higher for the complex. Similar results are obtained for Sn-H complexes in n-type GaAs, although distinct bands for the Sn isotopes have not been resolved to prove that the hydrogen is attached directly to the Sn. It is interesting to compare the donor dependence of donor-H complexes in Si and GaAs hosts. For Si hosts there is only a small frequency shift for a change of donor, and substitution of Sb for P gives a shift in the H-stretching frequency of only 6 cm"1. This is because in Si the hydrogen is not attached directly to the donor, but to a neighbouring Si atom. The configurations for the donor-hydrogen complexes in both hosts have hydrogen at an anti-bonding site so that mode frequencies and stress characteristics are similar even though the donor frequencies are markedly different.

D

DEEP LEVEL PASSIVATION

There is a relatively high concentration of deep levels in undoped, melt grown GaAs [1,2]. A variety of these levels, most due to native defect complexes, are passivated by hydrogen, including the so-called EL2 centre which is an arsenic anti-site (AsGa). EL2 is typically present at concentrations of ~ 1016 cm"3 and acts as a deep double donor, compensating the net shallow acceptors present. The electrical activity of EL2 can be restored by annealing at -400 0 C. Not all deep traps in GaAs are passivated by hydrogen but the majority can be deactivated, including metal impurities like Cu. E

DIFFUSION OF HYDROGEN

The effective diffusivity of deuterium in GaAs is listed in TABLE 3, and shows a wide range of values even in nominally similar samples. These differences are due to variations in the pressure, drive frequency, and ionization efficiency of the plasmas used for the hydrogen incorporation [13,14]. TABLE 3. Effective diffusion coefficients for 2H in III-V materials [13-15]. Material

E a (eV)

D 0 (cm7s)

n+-GaAs

1.38

115

n+-GaAs

1.43

90

n+-GaAs

0.62

1.5x10*

n+-GaAs

2.2

2 x 10"5

Undoped GaAs

0.97

4.4 x 10"2

SIGaAs

0.83

2 x 10"2

n-GaAs

-0.50

~10"3

P+-GaAs

0.58

4.2 x 10"6

The influence of carrier type and type of plasma exposure on hydrogen diffusivity can be summarized as follows [3]. (i)

The diffusion depth is greater in p-type GaAs for a given set of plasma exposure conditions, independent of the type of plasma. The diffusivity at 250 0 C is approximately a factor of 3 faster in p-GaAs for 30 kHz plasmas.

(ii)

The diffusion depth from a 30 kHz discharge into n-type GaAs weakly increases with decreasing donor concentration, whereas in p-GaAs the diffusion appears to be enhanced by increased levels of acceptor doping.

(iii)

D diffusion is inhibited and the degree of the acceptor (donor) passivation decreases when a thin n+ region is present at the surface of p(n) material.

(iv)

Low frequency plasma exposure produces about an order of magnitude greater depth and

greater peak concentration for the same exposure time and temperature than does indirect microwave exposure, in both n- and p-type GaAs. (v)

For low frequency plasma exposure, the depth of electrical passivation in both n- and p-type material is much less than the depth of hydrogen diffusion, and the concentration of hydrogen far exceeds the donor or acceptor concentrations. However, for high frequency or microwave plasmas there appears to be a closer correlation between electrical and chemical profiles.

F

INCORPORATION DURING GROWTH AND PROCESSING

Atomic hydrogen is found to be unintentionally incorporated into GaAs, AlGaAs, InP and related compounds during gas source epitaxial growth and subsequent processing steps. During growth, the origin of the hydrogen may be the group V source (AsH3 or PH3), the group III source (e.g. trimethylgallium) or the H2 carrier gas. Boiling of the HI-V materials in water or wet chemical etching steps both lead to near surface dopant passivation which is confirmed to be due to hydrogen by secondary ion mass spectrometry profiling. Dry etching in hydrogen-containing plasmas (e.g. CH4ZH2 or CH3Cl) leads to substantial dopant passivation, and this can occur even when hydrogen is not a specific part of the plasma chemistry because of the presence of water vapour or photoresist mask erosion. Hydrogen also readily permeates into III-V materials during annealing at moderate temperatures (4500C and above) in H2-containing ambients. Hydrogen is present in virtually every step during the processing of III-V devices, as an annealing or a sintering ambient, as part of a plasma used for dry etching, or as a component of the chemicals used for conventional wet-etching or solvent cleaning. In many situations the hydrogen is an unintentional contaminant as, for example, water vapour in vacuum systems or OH species in oxides. Hydrogen is also a common contaminant of silicon nitride and is a major constituent of organic resists that are used for masking purposes during device fabrication. It is becoming more obvious that, in many situations, hydrogen present either intentionally or unintentionally may diffuse into III-V semiconductors and alter the electrically active dopant profile in the near-surface region. This is a very undesirable situation in most cases since control of the switching and transmission characteristics of a device requires close control of the electric field in the active region near the surface. The reduction in doping density will also increase the contact resistance [15]. Fl

Epitaxial Growth

Hydrogen plays a key role in the growth mechanisms for both metal organic molecular beam epitaxy and metal organic chemical vapour deposition. In MOMBE [16] the precursors are typically AsH3, PH3, and metalorganics such as methyl (CH3)- or ethyl (C2H5)-bonded Ga, Al, or In. Some of the hydrogen from these sources remains in the epitaxial layer, passivating dopants. Hydrogen plays a key role in gettering carbon in some growth reactions, thereby lowering the residual p-type doping under some conditions. ECR H2 sources have proven very effective in removing oxides and hydrocarbon species prior to epitaxial growth. In MOCVD, hydrogen is generally used as the carrier gas and although it does not have a direct role in the chemical mechanisms underlying growth, it influences the decomposition of the

metalorganic precursors and the reactor transport characteristics. On the other hand, hydrogen within the source precursor plays a critical role in the growth process suppressing carbon incorporation and under some conditions leading to dopant passivation [17]. F2

Dry and Wet Etching

TABLE 4 shows the plasma chemistries involving H2 that are used for etching GaAs. Unintentional hydrogen passivation of near-surface dopants is generally observed after etching with such mixtures. Injection of hydrogen can also occur during boiling in water and during wet chemical etching treatments [15]. TABLE 4. Dry etch chemistries for GaAs involving hydrogen [15]. Gas mixture

Typical rates

Comments

CH4ZH2

<500A/min

smooth slow etch of all IH-Vs

CI2/CH4/H2

1 nm/min(>130°C)

rapid elevated temperature etch for In-based materials

CHCI2F, CHCIF 2

<750A/min

replacement CFCs for Freon-12

HBr

<700A/min

selective for InGaAs over AlInAs

HI

<500A/min

selective room temperature etch for In-based materials

CH3I, C2H5I, C3H7I

<1000A/min

less corrosive than HI, but still have polymer deposition

CHF 3

lOOA/min

etching of dielectrics and resists from III-Vs

F3

Chemical Vapour Deposition

Substantial dopant passivation in GaAs has been reported during PECVD of SiNx using SiH4 and N2 gases and direct evidence for hydrogen incorporation was obtained from SIMS measurements [18,19]. Virtually any CVD step involves hydrogen, e.g. deposition OfSiO2 from SiH4 and O2, from the reduction by H2 of WF6 or Ti from TiCl4 and NH3 [20]. In all of these situations one can expect hydrogen passivation of near-surface dopants to occur, particularly since the deposition temperatures are relatively low, typically < 50O0C, and there is a cool-down period after deposition where the dopant-hydrogen pairs occur. REFERENCES [1] [2] [3] [4] [5]

SJ. Pearton, J.W. Corbett, T.S. Shi [ Appl. Phys. A (West Germany) vol. 43 no.3 (1987) p. 153-96 and references therein ] R. Murray, [ in Properties of GaAs, 2nd edition, EMIS Datareviews Series No.2 (INSPEC, UK 1990)] W.C. Dautremont-Smith [Mater. Res. Soc. Symp. Proc. (USA) vol. 104 (1988) p.313 ] SJ. Pearton, [Mater. Set Rep. (Netherlands) vol.4 (1990) p.313 ] M. Stavola [Mater. Sci. Forum (Switzerland) vol.148-149 (1994) p.251-80 ]

[6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20]

J. Chevallier, B. Clerjaud, B. Pajot [ in Hydrogen in Semiconductors Eds J.Pankove, N.Johnson (Academic, USA 1991) p.447-510 ] B. Pajot, A. Mil, J. Chevallier, R. Azoulay [ Semicond. Sci. Technol. (UK) vol.2 (1987) p.305 ] S.K. Estreidcher [Mater. Sci. Forum (Switzerland) vol.148-149 (1994) p.349-92 ] P.R. Briddon, R. Jones [ Phys. Rev. Lett (USA) vol 64 (1990) p.2535-7 ] R.C. Newman [ Semicond. Sci. Technol. (UK) vol.5 (1990) p.911-4 ] R. Jones, S. Oberg [ Phys. Rev. B (USA) vol.44 (1991) p.3673-7 ] L. Pavesi, P. Giannozi [ Phys. Rev. B (USA) vol.43 (1991) p.2446-51 ] J.M. Zavada, R.G. Wilson [Mater. Sci. Forum (Switzerland) vol. 148-149 (1994) p. 189-218 ] J. Chevallier [Mater. Sci. Forum (Switzerland) vol.148-149 (1994) p.219-50 ] SJ. Pearton [ Int. J. Mod. Phys. B (Singapore) vol.8 no. 10 (1994) p. 1247-1342 ] CR. Abernathy [Mater. Sci. Forum (Switzerland) vol. 148-149 (1994) p.3-26 ] W. S. Hobson [ Mater. Sci. Forum (Switzerland) vol. 148-149 (1994) p.27-60 ] J.P. DeSouza, D Sadana, H. Baratte, F. Cardone [Appl. Phys. Lett (USA) vol.57 (1990) p.l 129 ] K.L. Seaward [Appl. Phys. Lett. (USA) vol.61 (1992) p.3002-4 ] A. Katz, A Fiengold, SJ. Pearton, U.K. Chakrakarti [ Appl. Phys. Lett. (USA) vol.59 (1991) p.579-81 ]

10.9 Transmission electron microscopy of GaAs U. Bangert December 1995

A

INTRODUCTION

GaAs has increasingly been studied by transmission electron microscopy (TEM) during the past decade. Initial research concerned the microstructural properties of bulk crystals in the as grown state and after plastic deformation. The microstructural research efforts soon branched out into dopant implantation studies, followed by efforts to correlate microstructure and electrical properties. More recently hetero-epitaxial systems with GaAs as substrate (e.g. heterostructure lasers) became a focus of interest. The large majority of research articles relating to TEM and GaAs in the past 5 years deal with investigations of pseudomorphic and hetero-morphic systems based on GaAs with particular emphasis on interfaces, the behaviour of GaAs based alloys and strain relaxation mechanisms. This led on to the investigation of complex systems such as multiple quantum well stacks, superlattices and self assembled 3-dimensional nano-structures. The present Datareview will consider in detail only investigations of defects in bulk GaAs, which are predominantly or entirely based on TEM using diffraction contrast and/or phase contrast (see Sections Dl, D2 and D3). Subsequent sections are included for completeness of the subject, but are treated in lesser detail. A very good source for further references are the Microscopy of Semiconducting Materials conference proceedings, published in the Institute of Physics conference series. B

TEM TECHNIQUES

TEM gives spatially resolved information about microstructural properties, such as dislocations and precipitates. Today the notation 'conventional TEM' is used for diffraction contrast imaging, where the diffraction conditions in the crystal are set up such that apart from the transmitted electron beam, which is used for imaging, only one diffracted beam is present (two beam condition). For certain two beam conditions, line defects (dislocations) for example, are then visible as lines of dark contrast. For conditions where g.b = 0, (i.e. the scalar product of the diffraction vector g and the dislocation's Burgers vector b is zero), the dislocation line is invisible (invisibility criterion). If such a condition is achieved under at least two imaging conditions, the Burgers vector and thus the structure of the dislocation can be determined. The 'weak beam' technique is cited when the actual imaging is performed with a diffracted, rather than the transmitted beam. This technique provides improved resolution. (For details on electron diffraction theory see, for example, [I]). High resolution electron microscopy (HREM) or 'lattice imaging' provides spatial resolution on the level of atomic separations and is very useful for the study of the atomic arrangements at interfaces. In this technique the phase contrast, which arises from highly localized variations in the electron intensity due to the operation of the atomic potential and the microscope transfer functions on the electron wave function, is used [2]. Provided the sample is oriented along a zone axis, the periodic image corresponds to columns of atoms and channels between them. The atomic arrangements of dislocation cores viewed end-on can then be revealed by comparison of

experimental images with those obtained from image simulation programs [3]. Another technique, in which only electrons that have been scattered into high angles are collected, is called Z-contrast imaging [4]. It owes its name to the Rutherford scattering mechanism, the cross-section of which is sensitively dependent on the scattering angle and the atomic number (Z). This technique, when carried out in a dedicated scanning transmission electron microscope (STEM) under the application of a very fine (2 A) electron probe size, is well suited to gain elemental or chemical information with atomic resolution. For completeness, convergent beam electron diffraction (CBED) must be mentioned [5]. In this technique, diffraction patterns are taken under high electron beam convergence angles and the fine structure within the diffraction 'discs' is used to gain information about local crystal symmetries. One of its major applications is in the analysis of local strain fields in heterostructures and superlattices [6]. Other TEM-related analytical techniques frequently applied to semiconductors use X-rays for energy dispersive X-ray analysis (EDX), or light (cathodoluminescence or CL) originating from the volume which the electron beam traverses. With the help of such analytical techniques one can gain chemical information. C

TEM SPECIMEN PREPARATION

There are various ways to make electron transparent specimens. The standard way is to prepare a sample no larger than 3 mm in diameter (because this is a standard size of most TEM specimen holders) in the form of a cut or cleaved slice of crystal. For plan view, this is usually parallel to the {100} crystallographic planes, for a cross-sectional view the cleave is usually parallel to {110}. This disk is then made electron transparent to a stage at which it usually has a perforation in the middle. To achieve this a combination of mechanical grinding and polishing, chemical etching and/or argon ion-beam milling is applied. A very fast technique to prepare GaAs based samples is the 90° cleaved-wedge technique, whereby cleaved pieces of GaAs with 90° corners are directly viewed in the TEM using tilt angles of 45°. The electron transparent area is immediately close to and along the entire length of the tilted edge [7]. Precision thinning techniques use a focused ion beam (FIB) to produce a thin membrane by cutting out trenches on either side [8]. Since the ion beam can be positioned within microns, areas of very small devices (e.g. GaAs lasers) or integrated circuits can be selectively thinned for inspection in the TEM to investigate the microstructural reasons for device failure. D

TEM OBSERVATIONS IN BULK GaAs

Dl

Dislocations

TEM has helped immensely to assess the quality of GaAs bulk crystals, which have gained great interest for high speed IC applications. A very good state of the art overview of dislocations in GaAs is given by Alexander et al [9]. Many bulk studies have pointed out the close relationship between etch pit density and electrical

and optical parameters such as PL efficiency, resistivity and Hall mobility. X-ray topographs have revealed dislocation and defect arrangements over large volumes and areas in bulk GaAs crystals. However, only TEM studies reveal the defect structure in detail. As examples: in In-alloyed GaAs crystals grown by the liquid encapsulated Czochralski (LEC) method, edge dislocations usually originate near seed areas due to misfit between the In doped crystal and the undoped seed and proceed vertically. It was found that dislocation loops are predominantly part of the slip system (110)/{111}, though for n-type GaAs dislocations on the (001)/{ 110} slip system have also been reported [11,12]. Cellular dislocation structures and As rich precipitates (caused by supercooling), have been found in semi-insulating GaAs grown by the liquid encapsulated Czochralski method (SI LEC GaAs). Since these dislocations are not always part of the common slip system, climb must be an important process in GaAs. This requires vacancies, which therefore must be present in As-rich SI LEC GaAs [13]. Vacancy and interstitial loops have been reported by Ponce [14]. Other common defects are stacking faults, multiple stacking faults and microtwins, resulting from dislocation glide and subsequent slip. Stacking fault energies have been deduced from the dissociation width of the partials by several authors. Gerthsen and Carter [15] measured the stacking fault energies in GaAs with different types of doping by weak beam TEM techniques. They discuss and explain the scatter in results obtained by different techniques and by different authors. In order to study dislocation behaviour in GaAs, plastic deformation has been used by many authors. This induces 'clean' dislocations. In such experiments edge, 60° and screw dislocations were usually observed. The common Burgers vector is a/2(l 10), the glide plane is {111} and the dislocation direction (110). 60° dislocations are often dissociated into 30° and 90° Shockley partials of Burgers vector a/2(112). Screw dislocations dissociate into 30° partials [16,17], though even further dislocation reactions at the core of these partials (i.e. formation and dissociation of 60° dipoles) have been reported [18]. Dislocations in GaAs are further classified into a- and p-dislocations. 60° dislocations are either a- or P-dislocations, their dissociation partials either both of the a- or both of the P-type. The a-dislocation is either of the As-glide or the Ga-shuffle configuration, while the p-dislocation is of the As-shuffle or Ga-glide configuration [19]. Currently, there are considerable difficulties distinguishing whether dislocations are of the glide(g) or shuffle(s) type, though Gerthsen et al [19] have conducted thorough HREM studies in conjunction with SHRLI multislice image simulations [20] and the underlying model of minimizing the total free energy, from which they conclude that the majority of dissociated screw and 60° dislocations are present in the glide configuration.

FIGURE 1. End-on HREM image in <110> projection of a dissociated screw dislocation in SI-GaAs, taken at 200 kV. Left insert (top): simulation of a 30° glide set partial, right insert: simulation of a 30° shuffle partial. Both simulations at a defocus of -30 nm and 10 nm sample thickness. This HREM image shows the rare case of a 30° shuffle partial. (Courtesy of Gerthsen et al, 1989.)

D2

Plastic Behaviour and Dislocation Mobilities

Glide does not occur continuously, but by kink nucleation and migration along the dislocation line. This process is strongly affected by electron-hole recombination [21]. In a considerable number of studies attempts have been made to deduce the velocities from TEM measurements of the length and separation of dislocation segments after controlled stress application. Due to effects invoked by the thin TEM specimen foil, surface contamination and electron-beam heating and damage, it is difficult to establish 'true' values. Values for undoped GaAs cited by Alexander et al [9] for ct- and p-dislocations at 5750C are 1.2 x 10"1 cm/s and 6.3 x 10 "3 cm/s respectively. Caillard et al [22] measured in situ straining experiments between 0.1 x 1O"4 cm/s and 10"4 cm/s at 3500C, depending on the dislocation length. P-dislocations are slower by a factor of 102 to 103 than a-dislocations in n-type and undoped material. Their velocity increases in p-doped material, a-dislocations are most mobile in undoped GaAs. Since the velocity trend does not depend on the size of the dopant, core effects on the velocity can be excluded and more delocalized effects, such as Fermi level pinning, are suggested. Screw dislocations control the plasticity in uniaxially (and hydrostatically) stressed material of any dopant type. At high temperatures plastic deformation is caused by the glide of screw segments, the average velocity being largest in p-type material, followed closely by intrinsic material, and is slowest in n-type material, according to respective increasing Peierls barriers (the activation energy at 0 K) (as a result, p-type material is easier to deform at higher temperatures) [1Oa].

fault energy for all types of material being similar. They argue that this indicates that the core structure is the same in all types, and that electronic effects, i.e. the position of the Fermi level, cause the different mobilities. At low temperatures plastic deformation is dominated by an asymmetry in the velocities of the 30° a- and p-partial dislocations, which determines the crossslip ability, and which is larger in n- than in p-doped material (as a result, n-type material is easier to deform at lower temperatures) [1Ob]. In p-type material much of the plastic deformation process has been reported to be due to microtwinning, i.e. subsequent cross-slip and separation of partial dislocations on adjacent {111} planes under stress, creating trailing stacking faults. These partial dislocations originate from loops expanding under stress, made up of screw and 60° segments (see above). In p-type material 30° a-dislocations are slower than 30° P-dislocations and 30° partials are slower than 90° partials. In this way, stacking faults are created. In n- and Si-type material a's are faster than P's and therefore perfect dislocation glide can take place [24]. It is often easier to perform micro-hardness indentation than uniaxial and hydrostatic straining. In this deformation mode, perfect dislocations and microtwins are formed in materials of all dopings. Studies by Lefebvre et al [25] show that, in n-doped and SI GaAs, screw and perfect adislocations cause slip in one of the rosette arm directions, whereas microtwinning due to partial P-dislocations cause slip in the perpendicular rosette arm direction. Detailed TEM studies of micro hardness indentations in p-type, n-type and undoped GaAs and GaAlAs by Charsley et al [26] contrast the above studies in that dissociation into partials, i.e. microtwinning, is found throughout both rosette arms, one on converging and the other on diverging {111} planes. The nature of the partial would thus remain the same for any one indentation rosette: it is Ga(g) in pdoped and As(g) in n-doped material, respectively, as detailed Burgers vector analysis revealed. That means that the leading partial changes from a to p, when the doping is changed from n- to p-type. Accordingly different dopings would result in different core structure in contrast to [23]. D3

Interaction of Dislocations with Point Defects and Dopants

Dislocations are often reported to be decorated with As precipitates. This has been revealed by Moire fringe contrast, selected area diffraction and energy dispersive X-ray analysis [27]. Melts of either Ga- or As-rich stoichiometry result in crystals of substantially different properties, the former exhibiting lower dislocation densities. Fornari et al [28] explained this with the reduction of the interstitial As concentration under Ga-rich conditions. Interstitial As is thought to be the point defect that can cause micro-loops upon condensation. In several observed dislocation configurations climb must play a role. Considering the amount of interstitial As, the most likely climb mechanism is that Ga vacancies are the driving force for a climb process, during which As antisite defects are created [29]. Unfortunately it is not possible by TEM to identify point defects directly. (So, for example, the debate whether EL2 is attributed to antisite defect complexes, correlated with dislocations [30], has not been resolved by TEM, but by combination of optical microscopy and absorption.) Only the effects of point defects on dislocations e.g. dislocation climb, enhancement of kink mobilities, transformation from shuffle into glide, or the condensation of point defects and impurities into dislocation loops or precipitates are directly observable.

The effect of doping atoms can be observed in the TEM as diffraction contrast change or as Fresnel fringe contrast [31] in highly doped layers. This is explained by strain effects (i.e. overall contraction) of the doped layer, and attempts have been made to quantify the contrast and hence the dopant concentration. Zn is a widely used acceptor in GaAs and a number of diffusion studies have therefore been undertaken. The microstructural side of these contributed significantly to the understanding of the diffusion process. Combined TEM and EDX studies revealed that in the diffusion front dislocations and voids, partially filled with Ga precipitates, form as a consequence of the incorporation of Zn on substitutional sites in a kick-out mechanism. The voids are due to the subsequent emission of As into interstitial sites, thus creating As-vacancies. In As rich ambients an As rich surface region is established causing the Ga to diffuse out of the voids to the surface [32]. Si is the most widely used donor in GaAs crystals. The most common defects associated with this doping are planar precipitates, which form faulted loops [33], prismatic loops [34], or an aggregation of loops (stacking fault tetrahedra). A variety of configurations e.g. { I l l } , {110} or {113} have been reported. D4

Ion Implantation in GaAs

As soon as the potential of GaAs as device material was recognized, a large number of dopant implantation studies were performed. In many cases TEM was applied in order to establish connected defect structures, their evolution with time and the conditions for their annealing. These

FIGURE 2. Microtwins in (100) thin foils._All the partial dislocations have the same Burgers vector B6 dislocation (arrowed), (a) SI GaAs. Bright-field g = 0 2 2 . (b) p-GaAs(In). Dark field, g = 0 2 2 . (Fromref. [24] courtesy of Dr. A. Lefebvre.)

studies range from in-situ observations of defect evolution in the collision cascade [35], to the classification of damage structures depending on implant nature and dose [36-38], to the investigation of annealing and recrystallization effects [39,40]. The accumulated literature is large and represents a subset of the general defect literature. The representative selection of references given above may serve as a starting point for further reading. E

TEM OBSERVATIONS IN DEPOSITED GaAs FILMS

El

Low Temperature GaAs

GaAs layers grown by molecular beam epitaxy (MBE) at very low substrate temperatures have gained considerable interest as buffer layers for GaAs metal-semiconductor field effect transistors due to the high resistivity and therefore excellent device isolation of such films. Eaglesham et al [41] have investigated the microstructural properties of GaAs films grown from 250 0 C down to room temperature by TEM. There is a critical thickness, beyond which the initially crystalline films become amorphous, due to roughening of the growth surface. The critical layer thickness increases with increasing Ga/As ratio as well as with the growth temperature. Crystalline low temperature GaAs films were found by analytical TEM to exhibit excess As, partially in the form of As precipitates. The latter specifically form upon annealing [42,43]. As-precipitates are believed to act as Schottky barriers, and this may result in the semi-insulating properties of this material [44]. E2

GaAs on Si

Particular effort has been spent on the investigation of GaAs grown epitaxially on Si, with the purpose of monolithic co-integration of GaAs into Si circuits. The nucleation of GaAs on Si in the form of islands, which eventually coalesce, results in antiphase boundaries (APBs) [45], twins [46] and stacking faults [47]. Nucleation of threading dislocations in the epitaxial GaAs films [49] as well as 60° and edge misfit dislocations at the interface has been reported [49,50]. Recent TEM efforts concentrated on the demonstration of improvements in epitaxial GaAs films. These were achieved by reducing the defect densities e.g. by off-axis growth or high temperature anneals and/or by using strained superlattices as dislocation filters [51-54]. F

TEM OBSERVATIONS IN THE METAL/GaAs SYSTEM

Due to its use in semiconductor devices, it has been very important to find low resistance, stable ohmic contacts to GaAs. A large number of papers deal with microstructural investigations of the metal/GaAs system. Emphasis has been put on investigating the metal/semiconductor interface with regards to metal interdiffusion and the formation of metal/semiconductor phases and precipitation of intermetallic compounds. Various metals and metal-Ge alloys have been used and the composition and atomic arrangements of various intermetallic phases been determined by TEM, electron diffraction, EDX and HREM. Au-based alloys (Au/Ge/Ni) are the most universally used contacts to GaAs. They have a low resistivity due to the formation of AuGa and NiAs phases, bridging the metal and the GaAs [55], but are not very stable thermally due to the low melting point of the AuGa phases [56]. NiInW and NiSiW are more stable [57,58]. Ni acts as a

diffusion barrier to In5 and NiAs phases are crucial in determining the electrical behaviour. Also other metals have been investigated. Co for example forms a CoAs solid solution phase [59], and Pd and Rh undergo complicated interface reactions producing voids, leading to poor adhesion and formation of Schottky barriers [60]. Many metals, when deposited onto GaAs (either epitaxially or by sputter deposition techniques), form high quality Schottky barriers, and have therefore wide applications in field effect transistors. The most widely investigated system here is Al/GaAs. TEM, in particular HREM, contributed significantly to the current understanding of the atomic arrangements, dislocations and formation of AlAs phases, which cause the Schottky barrier at such interfaces [61-63]. G

TEM OBSERVATIONS IN HETEROSTRUCTURES BASED ON GaAs

Many of the fundamental microstructural properties of GaAs, for which TEM was the essential tool, are well understood now. In the past 5 years TEM of GaAs has concentrated predominantly on investigations of hetero4ayers, epitaxially grown onto GaAs. TEM has been increasingly used as a characterization tool in combination with many other techniques (e.g. sputtering ion mass spectrometry, X-ray diffraction, photoluminescence, CL and Rutherford backscattering). Some references to the fundamental aspects of the GaAs hetero-interface, for the understanding of which TEM was instrumental, will be given here. The earliest GaAs hetero-interface studied was the GaAs/GaAlAs interface, its importance being due to the rapid development and tremendous increase in the use of this system in the heterostructure laser. There are many reports of the atomic interface structure and abruptness, investigated by HREM. Monolayer abruptness was thought to be an important criterion for the achievement of desired quantum well effects. The determination of monolayer abruptness of AlAs/GaAs interfaces has been claimed by imaging along the (110) pole [64]. Thoma and Cerva [65] pointed out the ambiguity of imaging under these conditions: an interface, which appears to be 1-2 monolayers wide, imaged along another, say the (100) direction, now appears to be 5 monolayers thick. Imaging along the (100) direction is less ambiguous and chemically sensitive due to the large difference in intensity contribution from the (200) diffraction spots. Walther and Gerthsen [66] used sophisticated HREM image simulations including pattern recognition, to quantify the lattice image contrast and to extract Al concentrations with a spatial resolution of a monolayer [67]. The first interface to be studied systematically for strain or mismatch effects, was the GaAs/InGaAs interface. Numerous investigations dealt with the determination of strain [68], and the phenomenon of strain relaxation by misfit dislocation formation [69,70]. The concept of the critical layer thickness (i.e. the thickness of the epitaxial, mismatched layer up to the point at which plastic deformation by misfit dislocations occurs) developed by Matthews [71] has been proven in experiments [72], though there are reservations, and refinements have been proposed. The nature of the misfit dislocations [73], their formation, dynamics [74] and their interactions [75] have been discussed in great detail. Ways to suppress them by using dislocation filters, graded buffer layers and strain compensation have been considered. Other papers deal with epitaxial growth mechanisms, the conditions for 2-dimensional (layer-by-layer) growth and the occurrence of 3-dimensional (island) growth, in particular the mechanism of island formation of InAs on GaAs [76-78]).

The exploration of the conditions of island growth or, more generally, of 3-dimensional growth modes (e.g. on vicinal substrate surfaces, or on high index surfaces [79,80]) then led to the discovery of'step bunching' [81] or 'supersteps' [82], strain induced lateral ordering [83] or more generally the phenomenon of'self-organized' growth [84], resulting in spontaneous formation of In(Ga)As quantum wires and quantum dots. Many other GaAs based heterosystems have been investigated with TEM involvement. Particular interest has been shown in II-VI compounds on GaAs, because of their potential use for blue lasers and solar cell material [85,86]. A deeper treatment of hetero-layers cannot be given here, since it is beyond the scope of this Datareview. H

CONCLUSION

TEM investigations of bulk GaAs have been reviewed, including some references to the research activity in hetero and pseudomorphic systems based on GaAs. TEM has played a vital role in developing our understanding of the microstructural properties of GaAs. To a large extent, the knowledge gained by TEM has enabled crystal growers and electronic and opto-electronic industries to achieve their present standard in the growth of bulk material, in integrated circuit technologies and in the fabrication of advanced structures and devices. REFERENCES [I] [2] [3]

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CHAPTERIl DIFFUSION OF IMPURITIES 11.1 11.2 11.3 11.4

Diffusion of impurities in GaAs: introduction Diffusion of shallow acceptors in GaAs Diffusion of shallow donors in GaAs Cr diffusion in GaAs

11.1 Diffusion of impurities in GaAs: introduction I. Harrison August 1995

The operation of modern electronic and optoelectronic devices depends critically on the distribution of dopants inside the device. Therefore to enable the device engineers to perform their jobs correctly, they must know with some certainty the concentration profile of dopants. The redistribution of dopants may occur deliberately or accidentally. In the former case, dopants are diffused into the device structure to define a region of a certain type and concentration. An obvious example would be the diffusion of Zn into n-type GaAs to define a p-n junction. This type of diffusion is becoming less important technologically since much finer dopant control can now be achieved with modern growth techniques. As an example, with the fine control of MBE and MOCVD GaAs can be doped on the monolayer scale. This fine control brings about its own problems because accidental redistribution of the dopants is now a significant problem. One may originally dope a single GaAs monolayer but during the growth of the capping layers the dopant may diffuse out resulting in several monolayers being doped. This type of accidental redistribution has caused significant problems in the growth of HBT structures by MOCVD. This accidental redistribution of dopants not only occurs during growth but also during any subsequent processing step which requires high temperatures. To fixlly understand the processes occurring one needs an understanding of the diffusion mechanisms of the dopants. In subsequent sections the diffusion mechanisms of various dopants are discussed. The results presented for the different dopants were acquired in several different ways. The easiest method though the least reliable is the junction depth method. The dopant is diffused into a doped bulk crystal. If the dopant being studied were a donor then the bulk crystal would be p-type and if an acceptor then it would be n-type. The diffusion depth of the dopant is found by finding the position of the p-n junction. This method provides no information concerning the shape of the profile. The second method uses a radioactive tracer. An isotope of the dopant which has been rendered radioactive is diffused into the sample. The shape of the diffusion profile can be obtained by etching the diffused sample and measuring the radioactivity of the sample. All these methods and the analysis of diffusion data in semiconductors are described in more detail by Tuck [1,2]. More recently secondary ion mass spectroscopy (SIMS) has been used to study diffusion. This technique is more sensitive than radioactive tracer methods but cannot be used to study very deep diffusion and there are problems in obtaining an accurate estimation of the true concentration. REFERENCES [1] [2]

B. Tuck [ Introduction to Diffusion in Semiconductors (Peter Peregrinus, Stevenage, 1974) ] B. Tuck [ Atomic Diffusion in IH-V Semiconductors (Adam Hilger, Bristol and Philadelphia, 1988)]

11.2 Diffusion of shallow acceptors in GaAs I. Harrison August 1995

A

Zn DIFFUSION INTO GaAs

Zn diffusion into GaAs is technologically important and has been extensively researched. The diffusion coefficient of Zn depends on the concentration of zinc and often the concentration profile shows a concave section. The diffusion front is invariably abrupt. This is shown in FIGUREl. It is now accepted that Zn rapidly diffuses interstitially as a Zn1+ species, after which it then moves on the substitutional site (Zn8"), the concentration of Zn B" becoming much greater than the concentration of Zn •*. It was Longini [1] who first put forward this type of diffusion mechanism. His model assumed that the Zn1+ became Zn8" by the following quasi-chemical reaction Z*f + VGa «-> Zn8" + 2e +

(1)

This mechanism is often called the substitutional-interstitial mechanism. This name is rather confusing when the other mechanism put forward also involves substitutional and interstitial Zn, the two models differing only in the mechanism by which the Zn becomes substitutional. In the latter model, incorporation of the Zn can be written (2)

Zinc coefficient (atoms/cm 3 )

Zn1+ + Ga «--> Zn8" + IGa + 2e +

Depth (jam)

FIGURE 1. Experimental diflusion profiles for zinc in GaAs at 10000C with excess arsenic in ampoule. Diflusion times: A, 10 min; B, 90 min; C, 3 h; D, 9 h (from [4]).

This is called the 'kick-out' mechanism [2] and it is this mechanism which is currently believed to be operative although there is no overwhelming evidence to support this conclusion. If one assumes that the Zn1+ and the Zn8" are always in thermal equilibrium then the diffusion constant for the Zn8 is given by D = k.Cs2Di

(3)

where k is a constant of proportionality and depends on the stoichiometry of the crystal, D1 is the diffusion coefficient OfZn1+ and C8 is the concentration OfZn8". This relationship does not depend on the incorporation mechanism and it has been tested using the isoconcentration technique [3,4]. (Isoconcentration diffusion is a two step experiment. In the first stage of the experiment a long duration diffusion is performed in order to dope the substrate uniformly. The second diffusion is then performed using the same conditions but with radioactive Zn.) Unfortunately, the concentration profile cannot be predicted by substituting EQN (1) into Ficks law since another important effect occurs. Zn solubility in GaAs is very high and surface concentrations in excess of 1020 cm"3 are often obtained. It follows that the crystal needs to accommodate 1020cm"3 interstitial atoms. This cannot be achieved and the quasi-chemical reaction describing the Zn incorporation is no longer in equilibrium and so one of the assumptions in the derivation of EQN (1) is no longer valid. From a device point of view, there are several issues to take note of. Firstly, the diffusion of Zn into a device structure needs to be performed at a temperature less than approximately 750 0 C [5]. Above this temperature, dislocation loops and other extended defects form. (For AlGaAs compounds this transition temperature is much lower [6].) Secondly, the quality of the mask is very important. The lateral diffusion under a SiO2 mask is very large. The ratio for lateral diffusion to depth diffusion can be as high as 10 [7]. The lateral diffusion under a SiNx mask is much lower but it depends on the quality of the mask material. The optimum composition is when the SiNx is stoichiometric and has a refractive index of 2.06. In this case the lateral diffusion to depth diffusion ratio is 0.6 [8]. B

Mn DIFFUSION INTO GaAs

The diffusion of Mn, under certain conditions, is very similar to that of Zn (see Section A). It produces profiles that are not described by the complementary error function. The diffusion is enhanced in p-type material [9] and when diffused into GaAs/AlAs MQWs, it causes diffusion induced disordering [10,11]. The only systematic study of Mn into GaAs appears to be that of Seltzer [14]. In this study the surface of the GaAs was coated with Mn. To prevent As loss from the surface of the GaAs, As was added to the diffusion ampoule. A more recent study of Mndiffused GaAs [13] has shown that the surface becomes uneven after this type of diffusion. Futhermore, TEM analysis of the surface region shows that there had been a chemical reaction between the Mn and the GaAs. This chemical reaction results in grains OfMnGa2 on the surface. To prevent surface degradation, Wu and Hsieh [15] suggested the use of MnAs as the diffusion source. Since the profiles they obtained could not be described by a simple complementary error function, they did not assign diffusion coefficients to their data.

In the early 1980s, there was much discussion on the origins of 'type conversion'. There was good evidence to suggest that Mn was the culprit [14-17]. This result was however questioned by Look et al [20] who came to the conclusion that there was not enough Mn to cause the type conversion. C

Be DIFFUSION INTO GaAs

Be is a group II element and in theory should diffuse in a similar way to zinc: the Be diffuses through the crystal interstitially before moving on to the substitutional site. Early measurements [19] showed that the Be diffusion was a strong function of As pressure in the ampoule, an increase in the As overpressure resulting in a decrease in the diffusion depth. This was later confirmed by Miller et al [10] who also showed that the Be diffusion was concentration dependent. Another important similarity with zinc diffusion is the difference between the diffusion from an external source and an internal grown in source. The Be diffusion coefficient in the case of the grown in source is several orders of magnitude smaller than that of the external source case [19,21,22]. More recent results on the diffusion of Be in the active regions of an AlGaAs/GaAs laser structure have shown that the interstitial Be is charged.

Diffusion Coefficient (cmV 1 )

Many of the authors who have studied Be diffusion have defined a diffusion coefficient. Most have assumed that the diffusion can be described by Fick's law and so their results should be treated with caution since Be diffusion is a concentration dependent one. Nevertheless the diffusion coefficients do provide rough estimates and are therefore presented in FIGURE 2. The results do not include the early work of Poltoratskii and Stuchebnikov [9] since they performed in-diffiision experiments. The work of Schubert et al was performed using Be 5-doped samples.

Schubert et al (1 990) Miller et al (1985), pAs Miller etal (1985), 7PAs McLevige et al (1 97S)

FIGURE 2. An Arrhenius plot of the available Be diffiision coefficients. Since Be diffusion is concentration dependent these values should only be used as rough estimates.

D

C DIFFUSION INTO GaAs

In order to achieve high gain, high frequency GaAs/AlGaAs heterojunction bipolar transistors it is necessary to grow a highly doped base region. Initially, group II elements were used. These elements, however, have a large diffusion coefficient which leads to problems of dopant diffusion during growth or fabrication. Carbon is a group IV atom and therefore, in theory, it is amphoteric. However, in practice it produces p-type material with the C atoms occupying the As sub-lattice. Above a concentration of about 1019 cm"3 the material begins to autocompensate. Recent work [22] suggests that under high C doping high concentrations of [100]-split interstitial complexes are responsible. The diffusion of C into GaAs has been studied by several groups. Invariably the C diffusion source has been grown in and the diffusion has been studied either electrically [24] or with SIMS [25]. The analysis also assumes that the diffusion is solely governed by Fick's equation with a diffusion coefficient independent of the doping level. A summary of the diffusion coefficients obtained are shown in FIGURE 3 [25-28]. Carbon diffusion in GaAs

Diffusion Coefficient (cm 2 s~ 1 )

As rich Gc ricn

FIGURE 3. An Arrhenius plot of C diffusion coefficients.

The solid circles are the C diffusion coefficients obtained when the samples were annealed in an As rich environment. The straight line passing through this set of data is the Arrhenius fit and it is given by the equation D c = 0.11Oe

-3,2eV kT

cm2/s

(4)

The solid squares are the C diffusion coefficients obtained by You et al [24] when they used Ga rich annealing conditions. There is a large scatter on these data points and so the values should be treated with some caution. The solid straight line is again the Arrhenius fit and is given by the

following expression 4

D c = 2.8 x i(r e

-2,7eV kT

cm2/s

(5)

One point to note from FIGURE 3 is that the C diffusion coefficient is largest for As rich conditions. This means that the diffusion mechanism cannot be via the diffusion of the V^ as first suggested [25]. You et al have suggested that C diffuses by a kick-out type of mechanism, the C atom moving interstitially before moving on to the substitutional site by the following quasi-chemical reaction, [24], Cf ~

C ^ + As,,,0

(6)

On annealing the hole concentration has been reported to decrease. Although this decrease cannot be explained by the diffusion of C away from the doped areas, there have been several suggestions to account for this. (It may also be related to the autocompensation effect described earlier.) Abernathy et al [29] suggested that the C atom switches from the As sub-lattice to the Ga sub-lattice. If this explanation is correct then the effect will be the greatest when the concentration OfV03 is enhanced. As-rich annealing provides these conditions and as expected the compensation effects are maximised. However, in the light of the recent calculations of defect energies this explanation seems less likely. Hanna et al [30] and Hoke et al [31] suggest that the CAs becomes a neutral interstitial (Q0). Another explanation put forward by Hofler et al [26] suggests that the formation of C precipitates is the cause of the compensation. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] [II] [12] [13] [14] [15] [16] [17] [18]

R.L. Longini [ Solid-State Electron. (UK) vol.5 (1962) p. 127-30 ] U. Gosele, F. Morehead [ J. Appl. Phys. (USA) vol.52 no.7 (1981) p.4617-9 ] L.L. Chang, G.L. Pearson [J Appl. Phys. (USA) vol.35 (1964) p. 1960 ] M.A. Kadhim, B. Tuck [ J. Mater. Sci. (UK) vol.7 (1972) p.68-74 ] J.F. Black, E.D. Jungbluth [ J. Electrochem. Soc. (USA) vol. 114 (1967) p. 181 -7 ] I. Harrison, H.P. Ho, B. Tuck, M. Henini, O.H. Hughes [ Semicond. Sci. Technol. (UK) vol.4 (1989)p.841-6] C. Blaauw, AJ. Springthorpe, S. Dzioba, B. Emmerstorfer, [ J. Electron. Mat. (USA) vol. 13 no.2 (1984) p.251-62] W.X. Zou, GA. Vawter, LA. Merz, J.L. Coldren, [ J. Appl. Phys. (USA) vol.62 (1987) p.828-31 ] EA. Poltoratskii, V.M. Stuchebnikov [Sov. Phys. Solid-State (USA) vol.8 (1966) p.770-1 ] J.N. Miller, D.M. Collins, NJ Moll [ Appl. Phys. Lett. (USA) vol.46 (1985) p.960-2 ] EA. Skoryatina, [ Sov. Phys. Semicond. (USA) vol.20 no. 10 (1986) p. 1177-8 ] H.P. Ho, I. Harrison, N. Babaali, B. Tuck, M. Henini [J Electron. Mater. (USA) vol.20 no.9 (1991) p.649-52 ] CH. Wu, K.C. Hsieh, G.E. Hofler, N. Elzein, N. Holonyak [Appl. Phys. Lett. (USA) vol.59 no. 10 (1991) p. 1224-6] M. Seltzer [ J. Phys. Chem. Solids (UK) vol.26 (1965) p.243-50 ] CH. Wu, K.C. Hsieh [ J. Appl. Phys. (USA) vol.72 no. 12 (1992) p.5642-8 ] J. Hallais, A. Mircea-Roussel, J.P Farges, G. Poiblaud [Inst. Phys. Conf. Ser. (UK) vol.33b (1977) p.220-7 ] R. Zucca [ Inst. Phys. Conf. Ser. (UK) vol.33b (1977) p.228-35 ] J.B Clegg, J.B. Scott, J. Hallais, A. Mircea-Roussel [ J. Appl. Phys. (USA) vol.52 (1981) p.1110-12]

[19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31]

P.E. Nordquist,P.B. Klein, S.G. Bishop,P.G. Siebrmann [lnst. Phys. Conf. Ser. (UK) vol.56 (1981) p.569-78 ] D.C. Look, P.W. Yu, J.E. Ehret, Y.K. Keo, R. Kwor [ in Semi-insulating Ul-VMaterials Eds S. Makram-Edeid, B.Tuck (Shiva Publishing, Nantwich, UK) vol.52 (1982) p.372 ] E.F. Schubert, J.M. Kuo, RF. Kopf, H.S. Luftman, L.C. Hopkins, NJ. Sauer [ J. Appl. Phys. (USA) vol.67 (1990) p. 1969-79] J.E. Cunningham, TH. Chiu, A. Ourmazd, W. Jan,T.Y. Kuo[J Cryst. Growth (Netherlands) vol.105 (1990) p. 111-5] B.H. Cheong, KJ. Chang, [ Phys. Rev. B, Condens. Matter. (USA) vol.49 no.24 (1994) p. 17436 ] H.M. You et al [J Appl. Phys. (USA) vol.74 no.4 (1993) p.2450-60 ] B.T. Cunningham, LJ. Guido, J.E. Baker, J.S. Major, N. Holonyak, G.E. Stillman [ Appl. Phys. Lett. (USA) vol.55 no.7 (1989) p.687-9 ] G.E. Hofler, HJ. Hofler, N. Holonyak Jr., K.C. Hsieh, [ J. Appl. Phys. (USA) vol.72 no. 11 (1992) p.5318-24] T.H. Chiu, J.E. Cunningham, J.A. Ditzenberger, W.Y. Jan [ J. Cryst. Growth (Netherlands) vol. 107 no.l-4(1991)p.l051-2] K. Saito et al [ J. Appl. Phys. (USA) vol.64 no.8 (1988) p.3975-9 ] CR Abernathy, SJ. Pearton, R Caruso, F. Ren, J. Kovalchik [Appl. Phys. Lett. (USA) vol.55 no. 17 (1989) p. 1750-2] M.C. Hanna, A. Majerfeld, D.M. Szmydm [Appl. Phys. Lett. (USA) vol.59 (1991) p.2001 ] W.E. Hoke, PJ. Lemonias, D.G. Weir, H.T. Hendriks, G.S. Jackson [ J. Appl. Phys. (USA) vol.69 (1991)p.511 ]

11.3 Diffusion of shallow donors in GaAs I. Harrison August 1995

A

Si DIFFUSION INTO GaAs

The amphoteric nature of Si increases the number of defects that can be involved in the diffusion process. To fully understand and model the Si diffusion processes one needs a full understanding of the interaction between the different possible defects. At concentrations below 1018 cm"3, it is generally accepted that Si sits on the gallium sublattice (SiGa), although GaAs grown by MBE on high index plane substrates can yield p-type material with the Si occupancy on the arsenic sublattice (Si^). At concentrations above 1018 cm"3, the picture becomes confused. The atomic silicon concentration is no longer equal to the electron concentration and compensation occurs. There has been a significant amount of work devoted to the study of the origins of this compensation [1-5]. There is however a general acceptance that at high doping levels the concentration of Si^ starts to increase. However, this is insufficient to account for the compensation [6]. A more detailed discussion can be found in a review by Harrison [7]. It is with this uncertainty over the exact nature of and relationship between the silicon defects that one ponders over the silicon diffusion mechanisms. In the early work [8,9] the diffusion length of Si in GaAs was assessed by measuring junction depths. These results indicated that the Si penetration depended slightly on the ambient As pressure. In the mid 1980s, Greiner and Gibbons [1.0,11] diffiised Si into GaAs and obtained the concentration profiles using secondary ion mass spectroscopy (SIMS). They found that they could explain their results by assuming that the Si diffused as [Si03-SiA8] pairs. The diffusion source was a Si layer deposited on the surface which was capped with SiO2. No Si diffusion was found when the Si diffusion source was capped with Si3N4. Later it was shown that the Si diffusion depended on both the background doping level [12] and the dopant [13]. An alternative mechanism was put forward for the diffusion mechanism of Si, involving [SiGa-VGa] pairs rather than [Si03-SiA8] pairs [48]. With recent theoretical work [14,15] this must now be thought of as the most likely explanation. The Si diffusion from 5 doped samples has been studied (for example [16-18]). The group working at AT&T studied the diffusion of Si from the 5 doped spikes using CV measurements. This work implicitly assumed that the diffusion of Si was described by Fick's law using a concentration independent diffusion coefficient. The earlier results of Greiner and Gibbons [11] showed that this was not the case. This is confirmed in the SIMS study of Nutt et al [17] and Hart et al [18] who observed shoulders on both sides of the Si doping spike. These results cannot be explained by just using Fick's law. More recently, attention has become focused on the growth of GaAs on high index planes (see for example [19]). On the (111)A, (211)A and (311)A planes the Si is incorporated on the As sublattice and so acts as an acceptor, the resulting material being p-type. The silicon diffusion in these regions is significantly less than in n-type regions [20].

B

S DIFFUSION INTO GaAs

In the very early work sulphur diffusion into GaAs was studied by measuring the junction depths formed in p-type material [9,21,22] or by taking the sheet resistivity measurements [23]. To obtain a value of the interdifiusion coefficient these implicitly assume that Fick's law holds. In only two pieces of work, [24,25] both using radioactive tracer techniques, have concentration profiles been obtained and so these must be considered as more reliable. These studies show that the diffusion profiles not only depend on the temperature but also on the vapour pressures above the GaAs surface. In fact, Young and Pearson [24] found that the S diffusion coefficient increased by two orders of magnitude when the As overpressure in the ampoule was changed from 10"4 atmospheres to 0.6 atmospheres. At high As pressures the S diffusion coefficient was independent of the As pressure and was given by /2.6eV\ 2

D = 1.85 x 1 0 " e ^ k T >cm 2 /s

(1)

It should be stressed that some of the Sn diffusion profiles could not be described by the standard solution to Fick's law and so the above equation should be used with some caution. There is still some confusion about the actual diffusion mechanism. Young and Pearson [24] suggested that S exists in elemental form as well as part of the complex yGsi'SM\VGai, the complex being the mobile species. Zahari and Tuck [26] proposed a different mechanism in which arsenic vacancies were taken into account. The latter mechanism has the advantage of predicting the shape of the diffusion profile. One major problem with both of these mechanisms is that they cannot explain the solubility of S at high arsenic pressures. In this regime the solubility can be expressed as Cs = KW 8 0 1

(2)

where C8 is the S concentration, K is a constant of proportionality and W8 is the weight of sulphur added to the diffusion ampoule. C

Sn DIFFUSION INTO GaAs

There is very little information available on the diffusion of Sn in GaAs. The work of Goldstein and Keller [27] and Fane and Goss [28] indicated that Sn diffusion is a very slow process when compared to that of acceptors in the same material. The largest body of experimentation was performed by Tuck and Badawi [29] who used radioactive tracer techniques to obtain the Sn concentration profiles. They found that their profiles were adequately described by the complementary error function, and that it was meaningful to obtain values of the diffusion coefficient. They noted that the diffusion coefficient depended on the background doping of the substrate used. For Sn diffusion into undoped GaAs the diffusion coefficient at a given temperature can be obtained from D = 3.2e

3.3eV kT

cm2/s

(3)

and for the n-type substrate doped to 2 x 1018 cm"3

8

D = 9.43 x 10" e

1.9eV kT

cm2/s

(4)

Arnold et al [30] diffused tin from spin-on sources at 8000C and 900 0 C. They found that they could not fit the standard solution of Fick's law to the Sn diffusion profiles and suggested that the Sn diffusion coefficient was proportional to the square of the electron concentration. A similar result was found by Allen et al [31] who gave the following expression of the Sn diffusion coefficient 4.IeV /

D = 2 x io

D

3

e

kT

X2

-

N

cm 2 /s

(5)

Se AND Te DIFFUSION INTO GaAs

There is very little work on the diffusion of Se into GaAs. Se in-diffusion was studied by Fane and Goss [28]. Their results show that the in-diffiision process cannot be described by Fick's law. In addition the diffusion depth (16 hours at 11000C) is relatively invariant with the As overpressure. This seems to provide evidence of the formation OfGa2Se3. (From the laws of thermodynamics the degree of freedom depends on the number of distinct phases present.) IfGa2Se3 forms, it can be shown that there is no longer any degree of freedom and so increasing the amount of As in the diffusion ampoule does not alter the concentration of point defects and so one would expect the diffusion rates to be very similar [32], The work of Fane and Goss concentrated on studying the diffusion of Se at one temperature (1100 0 C). To obtain the activation energy of the diffusion coefficient one needs to perform these experiments at several different temperatures. In light of the non-Fickian diffusion profile the definition of a single diffusion coefficient is rather suspect but this data is available and is given in TABLE 1 below for both Se and Te. (For references see Sharma [33].) TABLE 1. Pre-exponential term and activation energy for diffusion of Se and Te in GaAs. Dopant

Pre-exponential (Cm2S"1)

Activation energy

Se

3 x 103

4.IeV

Te

3 x 10 1

3.5 eV

REFERENCES [1] [2] [3] [4] [5]

R.T. Chen, W.G. Spitzer [J. Electron. Mater. (USA) vol.10 (1981) p. 1085-1129 ] R. Murray et al [ J. Appl. Phys. (USA) vol.66 (1989) p.2589-96 ] W.M. Theis, W.G. Spitzer [ J. Appl. Phys. (USA) vol.56 (1984) p.890-8 ] H. Ono5 R.C. Newman [J. Appl Phys. (USA) vol.66 (1989) p.141-5 ] R. Maguire, R. Murray, R.C. Newman, RB. Beall, JJ. Harris [ Appl. Phys. Lett. (USA) vol.50 (1987) p.516-8]

[6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33]

S. Schuppler, D.L. Adler, L.N. Pfeiffer, K.W. West, E.E. Chaban, P.H. Citrin [Appl. Phys. Lett. (USA) vol.63 (1993) p.2357-9 ] I. Harrison [ J. Mater. Sci. (UK) vol.4 (1993) p. 1-28 ] G.R. Antell [ Solid State Electron. (UK) vol.8 (1965) p.943-6 ] LJ. Vieland [ J. Phys. Chem. Solids (UK) vol.21 (1961) p.318 ] M.E. Greiner, J.F. Gibbons [Appl. Phys. Lett. (USA) vol.44 no.8 (1984) p.750-2 ] M.E. Greiner, J.F. Gibbons [J. Appl. Phys. (USA) vol.57 no.12 (1985) p.5181-91 ] D.G. Deppe, N. Holonyak, F.A. Kish, J.E. Baker [ Appl. Phys. Lett. (USA) vol.50 no. 15 (1987) p.998-1000 ] D.G. Deppe, N. Holonyak, J.E. Baker [ Appl. Phys. Lett. (USA) vol.52 no.2 (1988) p. 129-31 ] J. Dabrowski, J.E. Northrup [Phys. Rev. B, Condens. Matter (USA) vol.49 no.20 (1994) p. 14286-9] B. Chen, QM. Zhang, J. Bernholc [Phys. Rev. B, Condens. Matter (USA) vol.49 no.4 (1994) p.2985-8] E.F. Schubert, T.H. Cbiu, J.E. Cunningham, B. Tell, J.B. Stark [ J. Electron. Mat. (USA) vol. 17 no.6(1988)p.527-31] H.C. Nutt, R.S. Smith, M. Towers, RK. Rees, DJ. James [ J. Appl. Phys. (USA) vol.70 no.2 (1991) p.821-6] L. Hart, MJ. Ashwin, P.F. Fewster, X. Zhang, M.R. Fahy, RC. Newman [ Semicond. Sd. Technol. (UK) vol. 10 (1995) p.32-40 ] I. Harrison, L. Pavesi, M. Henini, D. Johnston. [ J. Appl. Phys. (USA), vol.75 (1994) p.3151-7 ] M. Hirai, H. Ohnishi, K Fujita, P. Vaccaro, T. Watanabe [ J. Cryst. Growth (Netherlands) vol. 150 (1-4 PtI) (1995) p.209-13 ] RG. Frieser [ J. Electrochem. Soc. (USA) vol. 112 (1995) p.697 ] T.H. Yeh [ J. Electrochem. Soc. (USA) vol. 111 (1964) p.253 ] D.L. Kendall [ Semicond. Semimet. (USA) vol.4, Eds RK. Willardson, A.C. Beer (Academic Press, New York, 1968) p. 163] T.H. Young, G.L. Pearson [J. Phys. Chem. Solids (UK) vol.31 (1970) p.517-27 ] B. Tuck, RG. Powell, [ J. Phys. D, Appl. Phys (UK) vol. 14 no.7 (1981) p. 1317-24 ] M.D. Zahari, B Tuck [ J. Phys. D, Appl. Phys (UK) vol. 15 (1982) p. 1741-50 ] B. Goldstein, H. Keller [J. Appl. Phys. (USA) vol.32 (1961) p. 1180 ] RW. Fane, AJ. Goss [Solid State Electron. (UK) vol.6 (1963) p.383 ] B. Tuck, M.H. Badawi [ J. Phys. D, Appl. Phys (UK), vol. 11 (1978) p.2541-52 ] N. Arnold, R Schmitt, K. Heime [J. Phys. D, Appl. Phys (UK) vol.17 (1984) p.443-74 ] E.L. Allen, JJ. Murray, M.D. Deal, J.D. Plummer, K.S. Jones, W.S. Rubart, [ J. Electrochem. Soc. (USA) vol. 138(11) (1991) p.3440-3449 ] B. Tuck [ Introduction to Diffusion in Semiconductors (Peter Peregrinus, Stevenage, 1974) ] B.L. Sharma [ Diffusion Defect Data, Solid State Data A, vol.64/65 (1989) p. 1-76 ]

11.4 Cr diffusion in GaAs I. Harrison August 1995

Cr is a deep acceptor in GaAs and has been used to provide semi-insulating GaAs substrates. One major drawback with this type of substrate is the redistribution of Cr into the epi-layer during growth or subsequent device processing steps. Despite its importance there have been few systematic studies of Cr diffusion into GaAs [1,2]. The diffusion of Cr into GaAs cannot be characterised by the standard solutions to Fick's law. Close to the surface there seems to be an accumulation of Cr. This was seen in all of the profiles of Tuck et al [1] who used, in general, slightly higher temperatures than Deal et al [2]. Since the SIMS method of probing the Cr concentration is more appropriate for shallow concentrations, Deal et al [2] were able to study the Cr concentration in the first 4 \im of the profile. They observed that this build-up was unpredictable and possibly due to localised precipitates. By ignoring the 4 nm surface layer, Deal et al were able to deduce that the Cr was diffusing by the substitutional interstitial mechanism. Furthermore by assuming that the rate of vacancy generation in the bulk, by dislocation climb, was larger than the rate at which the interstitial Cr atoms are converted into substitutionals, they were able to fit their profiles using the expression

f -i ) 0erfcW C(x) = Cserfc(n) { 1 - e J where \i =

(1)

is the extrapolated surface concentration and 0 is the vacancy generation para-

2v/Dt meter which is the time required for the attainment of the vacancy equilibrium in the bulk of the GaAs. The experimental values of D, Cs and 0 from [2] are given in TABLE 1. The values for the vacancy generation parameter will depend strongly on the type and quality of the GaAs substrate. TABLE 1. Diffusion coefficients, vacancy generation parameter (0) and extrapolated Cr surface concentration (C s ) at different temperatures. Temperature ( 0 C)

D(CmV 1 )

0

Cs(cm"3)

750

1.5 x IO"10

1.3 x 105

2 x 1016

800

2.8 x 10 1 0

3.2 x 104

2 x 1016

850

4.0 x 10-10

6.6 x 103

6 x 1016

900

5.7 x 10-10

3.0 x 103

7 x 1016

950

7.0 x 10-10

1000

9.1 x 10 1 0

Out-diffusion studies have also been performed by several groups [2-5]. The Cr depth profiles are characterised by a region in which the Cr concentration has been depleted. This may or may not be accompanied by an enhancement of the Cr concentration on or near the surface. If the

surface region is ignored then one can define an out-diffusion coefficient. FIGURE 1 shows the Cr diffusion coefficients obtained by these authors. In some circumstances there is over an order of magnitude difference in the experimental results. The solid line shows the Arrhenius fit of all the data and is given by D = L I x i o 2 exp (2.IeV/ kT) cm2/s

(2)

At any given temperature, this value is smaller than the value found for the indiffiision case. This is a direct consequence of the substitutional-interstitial mechanism.

Diffusion Coefficient (cm 2 s"')

Mizutani et ai (1 982) Deal end Stevenson (1 986) Kasaharc and Watanabe (1 980) Sato (1973)

FIGURE 1. An Arrhenius plot of Cr difiusion coefficients. The solid line is a least squares fit to all the experimental data.

There is a considerable amount of data for other rapidly diffusing species including transition and noble metals in GaAs. For details of these, see [6]. REFERENCES [1] [2] [3] [4] [5] [6]

B. Tuck, G.A. Adegboyega [ J Phys. D, Appl. Phys (UK) vol. 12 (1979) p. 1895-1908 ] M.D. Deal, D.A. Stevenson [ J Appl. Phys. (USA) vol.59 no.7 (1986) p.2398-2407 ] T. Mizutani, T. Honda, S. Ishida, Y. Kawasaki [ Solid-State Electron. (UK) vol.25 no.9 (1982) p.885-91 ] J. Kasahara, N. Watanabe [ Jpn. J. Appl. Phys. (Japan) vol. 19 (1980) p.L151-4 ] Y. Sato [ Jpn. J Appl. Phys. (Japan) vol. 12 (1973) p.242 ] B. Tuck [Atomic diffusion in ULVsemiconductors (Adam Hilger, Bristol and Philadelphia, 1988) ]

CHAPTER 12 SURFACE PASSIVATION 12.1 12.2 12.3

Technical aspects of surface passivation in GaAs Fundamental aspects of surface passivation in GaAs New approaches for surface passivation in GaAs

12.1 Technical aspects of surface passivation in GaAs H. Hasegawa May 1996

A

SELECTION CRITERIA

The standard method of surface passivation is to cover the semiconductor surface by a suitable insulating film in order to terminate and passivate active dangling bonds at the semiconductor surface and to protect devices from environmental influences. A more active use of such a film is to use it in insulated gate devices. Practical criteria for selecting a passivation film include: (1) (2) (3) (4) (5) (6) (7) (8) (9)

Film density or impermeability against outside impurity atoms and ions Adherence Thermal stability Chemical stability Mechanical stress at the interface Compatibility with other processing technologies Electrical resistivity Electrical breakdown field strength Density of electronic interface states etc.

B

PASSIVATION DATA

(i)

Thermal, plasma and anodic native oxides of GaAs (see Chapter 13), which are usually amorphous mixtures OfGa2O3 and As2O3 with some inclusion of elemental As [1-3], are generally poor in the above criteria of (3), (4) and (6). Thus, silicon-based insulators such as SiO2, Si3N4 and SiOxNy produced by various CVD processes are presently used in practical devices.

(ii)

The major difficulty associated with surface passivation of GaAs is related to criterion (9). Insulator formation on GaAs either by various oxidation processes or by various insulator deposition processes results, in general, in an insulator-semiconductor (I-S) interface with a high density of interface states (surface states) within the energy gap. The interface state density generally shows a U-shaped continuous distribution with its minimum value of 1012 -10 13 cm'2 eV "l occurring at 0.4 - 0.5 eV above Ev [4]. These interface states tend to fix the position of the Fermi level at or slightly (0.1 - 0.3 eV) below midgap, being largely independent of the insulator species and formation process, the conduction type and doping of GaAs and the application of external fields. This phenomenon is called Fermi level pinning.

(iii)

Interface states at the pinned I-S interface cause various unwanted phenomena in devices. Major adverse effects due to interface states in conventional GaAs based devices include: (a)

The inability to fabricate n-channel enhancement mode MISFETs utilizing surface inversion. This results in poor gate control and instability in n-channel depletion mode MISFETs [5]

(b) (c) (d) (e) (f) (g) (h) (i) (j) (iv)

High source resistance in planar MESFETs due to surface depletion [6] Frequency dispersion of transconductance in MESFETs [7] Side-gating in MESFET integrated circuits [8] Gain reduction in HBTs due to surface recombination [9] Quantum efficiency reduction in opto-electronic devices and solar cells due to surface recombination Dark currents in photo-detectors due to surface generation Catastrophic mirror damage in lasers due to recombination enhanced interface reactions Recombination-generation noise in electronic and opto-electronic devices Long term instability phenomena in electronic and opto-electronic devices, etc.

Recent experiments on GaAs-based, near surface quantum wells [10-12] have shown that photoluminescence intensity of the quantum well is drastically reduced as the well-tosurface distance is reduced below 100 A due to interaction between confined quantum states and surface states. This indicates that surface passivation will become a more important critical technological issue for future nano-structure devices in the quantum regime due to an increase of the surface-to-volume ratio.

REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] II1] [12]

G.P. Schwartz, J.E. Griffiths, D. Distefano, GJ. Gualtieri, B. Schwartz [Appl. Phys. Lett. (USA) vol.34 (1979) p.742] K. Watanabe, M. Hashiba, T. Yamashina [ Thin Solid Films (Switzerland) vol.56 (1979) p.63 ] Y. Mizokawa, H. Iwasaki, R. Nishitani, S. Nakamura [ J. Electrochem. Soc. (USA) vol. 126 (1979) p. 1370] H. Hasegawa, T. Sawada [ Thin Solid Films (Switzerland) vol. 103 (1983) p. 119 ] [ Second Special Issue on Semiconducting HI-V Compound MIS Structures, Thin Solid Films (Switzerland) vol.103 (1983) ] N. Yokoyama, T. Ohnishi, K. Odani, H. Onodera, M. Abe [ Tech. Digest Int. Electron Device Meeting, Washington, DC, USA, 7-10 Dec. 1981 (IEEE, New York, 1981) p.80-3 ] J. Graffeuil, Z. Hadjoub, J.P. Fortea, M. Pouysegur [ Solid-State Electron. (UK) vol.29 (1986) p. 1087] H. Hasegawa, T. Kitagawa, T. Sawada, H. Ohno [ 11th Int. Symp. on Gallium Arsenide and Related Compounds, Biarritz, France, 26-28 Sept 1984 (Adam Hilger, Bristol, UK, 1985) p.521-6 ] CJ. Sandroff; RN. Nottenburg, J.C. Bishoff, R. Bhat [ Appl. Phys. Lett. (USA) vol.51 (1987) p.33] J. Moisson et al [ Phys. Rev. B (USA) vol.41 (1990) p. 12945 ] Y.-L. Chang et al [ J. Appl. Phys. (USA) vol.74 (1993) p.5144 ] Z. Sobiesierski, D.I. Westwood, D.A. Woolf, T. Fukui, H. Hasegawa [ J. Vac. Sci. Technol. B (USA)YoUl (1993)p. 1723 ]

12.2 Fundamental aspects of surface passivation in GaAs H. Hasegawa May 1996

For further progress in the technology of surface passivation, fundamental understanding of the origin of interface states and Fermi level pinning is required. Extensive studies on semiconductor surfaces and interfaces in the past ten years have revealed that Fermi level pinning is commonly observed on vacuum-semiconductor (V-S), insulator-semiconductor (I-S), metal-semiconductor (M-S) and defective semiconductor-semiconductor (S-S) interfaces (regrown interfaces, lattice mismatched heterointerfaces) of GaAs and that the pinning positions are intercorrelated among these interfaces.

model

origin of pinning

Unified Defect Model

deep level related to stoichiometry, especially, As

N ss distribution applicable and pinning interfaces position V-S, S-S, I-S, M-S

nature of references pinning

extrinsic

[1]

Ga

Missing Dimer Defect Model

acceptor due to kinks in missing dimer arrays

MIGS model

penetration of metal wave function into semiconductor

V-S

extrinsic [2] (Si-doping)

M-S

intrinsic

Pl. t4]

extrinsic

[5]

extrinsic

[6]

midgap energy DIGS model

Effective Work Function Model

loss of 2D periodicity due V-S, to disorder of S-S, I-S, bonds at M-S interface mean hybrid orbital energy precipitation of As and P cluster at interface

pinned at Ep of metallic cluster

V-S, S-S, I-S, M-S

V-S: Vacuum-Semiconductor interface S-S: Semiconductor-Semiconductor interface I-S : Insulator-Semiconductor interface M-S: Metal-Semiconductor interface FIGURE 1. Major models for Fermi level pinning.

Based on these correlations, various models concerning the origin of Fermi level pinning have been proposed. FIGURE 1 summarizes the major models for Fermi level pinning [1-6]. Further work seems to be necessary, however, to identify the dominant mechanism for Fermi level pinning at GaAs I-S interfaces.

REFERENCES [1] [2] [3] [4] [5] [6]

W.E. Spicer, I. Lindau, P.R. Skeath, CY. Su [ J. Vac. Sci. Technol. (USA) vol. 17 (1980) p. 1019 ] M.D. Pashley, K.W. Haberern [ Phys. Rev. Lett. (USA) vol.67 (1991) p.2697 ] V. Heine [ Phys. Rev. (USA) vol. 138 (1965) p. 1689 ] J. Tersoff [ Phys. Rev. Lett. (USA) vol.52 (1984) p.465 ] H. Hasegawa, H. Ohno [ J. Vac. Sci. Technol. B (USA) vol.4 (1986) p. 1130 ] J.M. Woodall, J.F. Freeouf [ J. Vac. Sci. Technol. (USA) vol. 19 (1981) p. 794 ]

12.3 New approaches for surface passivation in GaAs H. Hasegawa May 1996

Extensive efforts in the past have shown that Fermi level pinning at GaAs I-S interfaces cannot be completely removed by the standard approaches of direct insulator formation on GaAs by traditional ways of oxidation or insulator deposition. Thus, various non-traditional approaches for surface passivation have been investigated in recent years. Representative reports are listed below. The effectiveness and practical applicability of each approach, however, seem to require further critical investigations (i)

Photochemical oxidation of GaAs in running water produces a Ga-rich oxide and removes the Fermi level pinning [1,2].

(ii)

Formation of ultrathin sulphide or selenide layers by various wet surface treatments or dry deposition processes removes the native oxide from the GaAs surface and removes the Fermi level pinning [3-7].

(iii)

Epitaxial growth of insulating fluorides such as CaxSr1^F2 on GaAs removes the Fermi level pinning [8,9].

(iv)

Insertion of an ultrathin silicon layer between the GaAs and a Si-based dielectric [10-12] realizes a pinning-free I-S interface and produces much improved GaAs MIS capacitors [13]. The pseudomorphic ultrathin Si layer, often referred to as a Si interface control layer (Si ICL), terminates the surface bonds of GaAs and makes a smooth transition to an outer silicon based dielectric layer such as SiO2 or Si3N4 [14]. This technique has successfully removed surface state effects from near surface quantum wells and wires [15,16].

(v)

Growth of an ultrathin InP or GaP epitaxial layer on GaAs increases the surface photoluminescence and removes the pinning [17].

(vi)

A hydrogen ion treatment of the native-oxide covered surface controls the surface stoichiometry and increases the quantum efficiency of photoluminescence from GaAsbased near-surface quantum wells, indicating the removal of Fermi level pinning [18].

(vii)

Vacuum deposition OfGa2O3 film from Gd3Ga5O12 produces an As-free I-S interface and removes the Fermi level pinning [19].

REFERENCES [1] [2] [3] [4] [5]

J.L. Freeouf, J.M. Woodall [ Surf. Sd. (Netherlands) vol.168 (1986) p.518 ] T. Sawada, H. Hasegawa, H. Ohno [ Jpn. J Appl Phys. (Japan) vol.26 (1987) p.L1871 ] CJ. Sandroff, R.N. Nottenberg, J.-C. Bischoff, R. Bhat [Appl Phys. Lett. (USA) vol.51 (1987) p.33] Y. Nannichi, J.-F. Fan, H. Oigawa, A. Koma [ Jpn. J Appl. Phys. (Japan) vol.27 (1988) p.L2367 ] H. Oigawa, J.-F. Fan, Y. Nannichi, K. Ando, K Saiki, A. Koma [ Jpn. J. Appl. Phys. (Japan) vol.28 (1989)p.L340]

[6] [7] [8]

[9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19]

CJ. Sandrofl; M.S. Hedge, LA. Farrow, R. Bhat, J.P. Harbison, CC. Chang [ J. Appl. Phys. (USA) vol.51 (1990) p.586] T. Kikawa, S. Takatani, Y. Tezen [ Appl. Phys. Lett. (USA) vol.60 (1992) p.2785 ] H. Ishiwara, K.H. Kim, K. Tsutsui, T. Asano, S. Furukawa [ Electrochemical Society Proc. Symp. Dielectric Films on Compound Semiconductors, Ed VJ. Kapoor, vol.88-15 (1988) p.278 (The Electrochemical Society, Pennington, NJ, USA) ] T. Waho, F. Yamanaka [IEEEElectron Device Lett. (USA) vol.EDL-9 (1988) p.548 ] H. Hasegawa, M. Akazawa, K. Matsuzaki, H. Ishii, H. Ohno [ Jpn. J. Appl. Phys. (Japan) vol.27 (1988)p.L2265] S. Tiwari, S.L. Wright, J. Batey [IEEEElectron Device Lett. (USA) vol.9 (1988) p.488 ] GG. Fountain, S.V. Hattangady, DJ. Vikavage, KA. Rudder, RJ. Markunas [ Electron. Lett. (UK) vol.24 (1988) p. 1135] Z. Wang et al [ Appl. Phys. Lett. (USA) vol.62 (1993) p.2977 ] S. Kodama, M. Akazawa, H. Hasegawa [ J. Electron. Mater. (USA) vol.22 (1993) p.289 ] S. Kodama, S. Koyanagi, T. Hashizume, H. Hasegawa [J Vac. Sci. Technol. B (USA) vol.13 (1995) p. 1794] H. Fujikura, S. Kodama, H. Hasegawa [ Proc. 8th Int. Conf. on Indium Phosphide and Related Materials, Schwabisch Gmiind, April 1996 p.323 ] Y. Wada, K. Wada [ J. Vac. Sci. Technol. B (USA) vol. 11 (1993) p. 1589 ] Y.-L. Chang et al [ Appl. Phys. Lett. (USA) vol.62 (1993) p. 1723 ] M. Passlack et al [ J. Appl. Phys. (USA) vol.77 (1995) p.686 ]

CHAPTER 13 SURFACE STRUCTURE AND OXIDATION 13.1 13.2 13.3 13.4 13.5 13.6 13.7

Surface structure of GaAs Oxide layer structures of GaAs Thermal oxides of GaAs Wet oxidation of GaAs Plasma oxidation of GaAs Laser assisted oxidation of GaAs Miscellaneous new methods of GaAs oxidation

13.1 Surface structure of GaAs Jing Zhang September 1995

A

GENERAL COMMENTS

With the exception of the non-polar (110) face, low index GaAs surfaces exhibit a large range of surface reconstruction depending on the surface orientation, stoichiometry, temperature and to a lesser extent preparation methods. The driving force for the occurrence of reconstruction is the minimization of dangling bonds with their energy at midgap. Some of the essential aspects have been discussed by Harrison [I]. Most of the studies of GaAs surface structures concentrate on the technologically important (001) polar surface and the (110) surface produced by natural cleavage. Applications of scanning tunnelling microscopy have brought many advances in the understanding of the GaAs surface structure and raised new questions. This Datareview is divided into four sections dealing with the (100), (110), (111) and high index surfaces respectively. B

THE (100) SURFACE

This surface exhibits a number of reconstructions depending on its stoichiometry. They range in order of decreasing Ga to As ratio from the Ga-rich (4><2) or c(8><2) through (4x6) or c(6x4), then (3 x 1) or (1 x6), then (2x4) or c(2x8) to the As rich c(4x4) [2-7]. Some of these

Ga droplets

facetting

F I G U R E 1. Surface structure p h a s e diagram o f GaAs(OO 1) surface misorientated 2 ° towards ( 1 1 1 ) during M B E growth. R e p r o d u c e d from D a w e r i t z e t al [ 8 ] .

reconstructions are only stable at high temperatures under constant exposure to arsenic molecules. Others can be quenched to room temperature where techniques such as STM can be applied. All these reconstructions can be studied in-situ using RHEED during MBE. A surface structure phase diagram of the 2° vicinal surface during growth has been produced by Daweritz and Hey [8] (FIGURE 1). The optical response of a range of reconstructions prepared by MBE was studied in-situ using reflectance difference spectroscopy by Kamiya et al [9]. The spectroscopic responses are very different for various reconstructions, making it possible to study the surface structures in processes such as VPE using this optical technique. Larsen et al [7,10] provided evidence, based on surface core level shifts, that the c(4><4) structure is a chemisorbed phase of trigonally bonded arsenic pairs (more frequently referred to as dimers). Their proposed model is consistent with STM measurements [11-13]. The unit cell consists of three dimers with their axis along the [110] direction and they are arranged to give a c(4><4) symmetry. Different coverages of the chemisorbed As are achieved by vacancies of As atoms in the unit cell [13]. The structure model is also consistent with grazing incidence X-ray diffraction measurements [14]. Under MOVPE conditions, the (100) surface is also reported to have a c(4><4) like reconstruction based on in-situ RDS [15] and GIXD [16] measurements. The c(2><8) surface observed in LEED has long been recognised as a special case of the (2x4) surface with anti-phase domain boundaries [17,18]. A RHEED study by Farrell and Palmstrem [19] divided this reconstruction further into a, P and y phases in ascending As coverage according to the intensity of the fractional order diffraction features obtained in [TlO] azimuth. Based on RHEED [20] and STM [21,12] studies and total energy calculations [22], it is generally accepted that the 2-fold symmetry arises from the dimerisation of surface As atoms along the [TlO] direction. The dimers are asymmetric or buckled [22]. Earlier STM results [21,12] suggested that the 4-fold periodicity is the result of a single missing dimer with a complementary argument from RHEED [23,24], a simple electron counting rule [25] and total energy calculations [23,26,27]. More recent measurements [28,13] demonstrated that many of the unit cells on the surface had two missing dimers and the P phase surface contained large ordered domains of such unit cells. This model is further supported by rocking curve measurements in RHEED [29], theoretical calculations [30] and agreement with the electron counting rule. Hashizume et al [28] also suggested a method to distinguish between two or three dimer unit cells based on examination of STM images in a kink site. The structure within the trenches produced by the missing dimer is less clear. The original model proposed by Chadi [22] consists of missing Ga atoms and dimerisation of As in the layer below. However, there is other experimental evidence suggesting the presence of Ga on the surface layers [31,32]. This evidence is subject to complications that arise from decapping procedures used in sample preparation. Both the a and y phases are reported as being disordered compared with the P phase. Based on STM results, Hashizume et al [28] believe that the different phases of the (2x4) surface arise from disorder and different structure within the trenches but Avery and co-workers [13] believe that disorder is the main cause. The band structure of this surface was investigated by angle resolved photoemission [33,34] and inverse photoemission [35]. The more Ga rich reconstructions, including the transitional phases, were studied by a number of groups using STM [12,36,37]. The (4x2) unit cell consists of two Ga dimers with two dimer vacancies with As absent in the trenches. The model is essentially a mirror image of the (2x4)

arsenic terminated surface. Xue et al [37] also proposed a unified model for the (4x2) and (4><6) Ga rich surface. However, a recent LEED study of the c(8x2) surface [38] suggested that three Ga dimers with one vacancy and anti-phase boundaries produces the best fit in I-V analysis. The phase transition between the As rich (2x4) and Ga rich (4x2) reconstructions was studied by Yamaguchi and Horikoshi [39]. The optical properties of the two surfaces have also been studied by Ren and Chang [40]. C

THE (110) SURFACE

This surface is the natural cleavage face of GaAs and one of the most understood in terms of surface crystallography. A great deal of the earlier work based on elastic LEED intensity measurements [41-45] has been reviewed by Kahn [46] and Duke [47]. The generally accepted structure consists of a buckled atomic arrangement with a constant Ga-As bond length rotated by 27°, i.e. the arsenic atom moving outwards by about 0.65 A relative to the gallium. This ultimately leads to a distortion of the next atomic layer. The arrangement is consistent with the observation that carefully prepared (cleaved) surfaces have no surface states present in the fundamental gap [48-50]. Theoretical predictions suggest that an unrelaxed surface has intrinsic surface states in the bandgap, but that these are removed by the 27° relaxation (see for example [51]). The dispersion of the surface states has been calculated by a number of groups [52,53,51]. Comparisons with photoemission [50] and inverse photoemission [54,55] measurements were also made by Zhu and co-workers [52]. Contradicting results based on high energy ion scattering experiments were reported by Gossmann and Gibson [56], who claimed a smaller 7° angular relaxation. This result also appeared to produce an adequate fit to LEED data using the X-ray R factor [57]. However, subsequent medium energy ion scattering experiments [58] and advances in the dynamic theory of LEED [59,60] re-affirmed the 27° relaxation model. Theoretical studies using total energy calculations [61,53,51] also supported this model. STM has been used to study the (110) surface from as early as 1985 [62] and has confirmed the presence of a ( l x l ) surface geometry. Using a voltage-dependent imaging technique [63,64], images of filled and empty states were obtained which were identified with images of Ga and As, respectively. Comparisons of experimentally obtained STM images were made with theoretical calculations based on the structural model. While earlier work [63] showed a small discrepancy, Wang et al [65] were able to demonstrate excellent agreement with the 27° rotation model using local density functional - pseudopotential formalism and allowing relaxation of atoms in the top three layers. Atomically resolved AFM images of the (110) surface were also observed by Ohta and co-workers [66]. There have been several applications of RHEED in determining the (110) surface structure. Jamison and co-workers [67] used this well known surface as a test for their dynamic calculations of the RHEED rocking curve. Wang [68,69] modified the existing multislice dynamic theory and demonstrated that inelastic scattering can greatly enhance the reflectivity of the surface in the case of the GaAs (110) surface.

D

THE(111)SURFACE

There are two non-equivalent GaAs (111) surfaces, namely the Ga-terminated (111)A surface and the As terminated (T T T) surface or (Hl)B surface. When prepared by MBE, both surfaces exhibit (2x2) reconstruction with the As terminated (T T T) showing additional structures [70,71]. Apart from sharing a common surface reconstruction, the two surfaces are very different in stoichiometry, surface state dispersion [72] and surface core level shifts [73]. Thornton and co-workers [74] have investigated the differences between the two surfaces produced by the desorption of an As cap using STM and photoemission measurements. The Ga rich (111) surfaces prepared by ion sputtering/annealing were studied by Tong et al [75] using LEED. They explained the 2x2 reconstruction using a vacancy-buckling model. The model is based on a Ga vacancy and buckling of surrounding Ga and As atoms (in the layer below). It was supported by a number of total energy calculations [76-78] and the Ga vacancy was confirmed by STM observation [79]. When prepared by [80], the (2x2) reconstruction exists over a large range of temperatures and As incident fluxes. The range of surface structures investigated by Kaxiras et al [77] includes an energetically favourable As adatom (in the form of trimer)induced (2x2) reconstruction. Such a surface was reported by Thornton and co-workers [74] using STM and an As decapping procedure. They demonstrated the existence of a (2x2) reconstruction induced by an As trimer. They also obtained both filled and empty state images of the Ga vacancy-buckling structure and demonstrated agreement with theoretical predictions.

JAs (molecules cm' 2 s" 1 )

The As terminated surface exhibits a range of surface reconstructions depending on the surface stoichiometry [81,3,82]. These reconstructions include ( l x l ) , (2x2) and (^19x719, rotated by 23.4°). A phase diagram of GaAs (T T T) prepared by MBE was produced by Woolf et al [80,82] (FIGURE 2). This surface was studied using STM by a number of groups [83,84,74].

FIGURE 2. Surface structure phase diagram of GaAs ( 1 1 1 ) . Reproduced from Woolf et al [82].

While original theoretical calculations suggested a vacancy [85], the STM observations favour an adsorbed As trimer structure causing the (2x2) reconstructions. This is consistent with total energy calculations [83,86]. The electronic structure of the (2><2) surface has also been studied by Cai et al [87] using normal emission AEPES. The (719x^19) reconstruction contains less surface As compared with the (2x2) surface [70]. Biegelsen et al [83] proposed a two layer hexagonal ring structure based on the STM results and total energy calculations. Kaxiras et al [86] also discussed the phase transition between the two surface symmetries. E

HIGH INDEX SURFACES

Earlier studies of high index surfaces were based on LEED with the surfaces prepared by etching/heating or ion sputtering/annealing techniques. Hren and co-workers [88] studied the (211)A and (211)B surfaces though they were not able to make positive identifications of the two. They found that both surfaces are unstable and develop large (110) facets. All three possible {110} facets were observed for one of them but only two were present for the other. These results were confirmed by Stiles and Kahn [89] who also studied the (311) surfaces. A (1x1) surface was observed to be stable up to 6000C. The (331) surfaces prepared by ion sputtering/annealing and MBE growth were studied by Horng et al [90]. They reported that the (33I)A surface exhibits (110) terraces with double layer steps which transform into a (110) and (111) faceted surface after annealing at high temperatures. The (331)B surface, on the other hand, remains flat at low temperatures. Further annealing at 5700C produces a double stepped surface with a (110) terrace on the MBE prepared sample. Annealing at temperatures above 620 0 C produces a (110) and (111) faceted surface from both sample preparation techniques. These results were further substantiated by the work of Weiss and co-workers [91]. More recently, there has been a renewed interest in high index planes due to the possibility of formation of quantum dots/wires by direct growth techniques [92] and the different doping behaviour of Si on these surfaces. Notzel and co-workers [93] made an extensive RHEED study of (331), (311), (211) and (210) surfaces. They showed that the (331) surface produced by MBE growth and annealed below 5500C had a regular step array structure with step edges running in the [llO] direction. The sides of the steps are (110) and (111) type with an overall terrace width of «1.8 nm. The surface roughens at annealing temperatures above 5500C and produces irregular macroscopic facets. These appear to be consistent with the earlier LEED work. On the (311) surface, Notzel et al [93] found a regular stepped structure along the [233] direction consistent with a two level system. The top and bottom surfaces are (311) which are linked by inclined facets of (33T) and (313). The structure is reported as being stable up to 680 0 C with a high degree of ordering. The step height was estimated as lnm with a lateral periodicity of 3.2 nm. The (211) surface is also reported as having a regular step structure but along the [TlO] direction. Observation from the (210) surface suggests that regular step structures are present in the orthogonal [001] and [120] directions which are stable below 5600C. Notzel et al proposed an asymmetric pyramid structure with (111), (101), (110) and (100) side walls. STM observations on high index surfaces have recently been made by Wassermeier et al [94] on (311)A surfaces and Yamada and co-workers [95] on (411)A surfaces. On the (311)A surface, Wassermeier et al reported the same lateral periodicity though the height of the corrugated structure was only 0.34 nm and suggested this is due to a surface reconstruction. They also

proposed a model of an (8x1) unit cell based on As dimerisation which satisfies the electron counting rule. In contrast, Yamada and co-workers observed a highly ordered structure on the (411)A surface along the [OlT] direction. They also proposed a model of the surface unit cell. The (311)B surface was reported [96] to have a (1x1) surface structure independent of the preparation techniques by ion sputtering/annealing or MBE. For an As rich surface prepared by MBE, three surface states derived from the As dangling bonds were observed. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] [II] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30]

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13.2 Oxide layer structures of GaAs H.L. Hartnagel and R. Riemenschneider July 1995

A

AMORPHISATION AND CRYSTALLISATION

Native oxide films grown near to room temperature by electrolytic or plasma anodization are usually amorphous when produced with current densities less than 2 mA/cm2. In this case, the oxide bulk is composed of a mixture OfGa2O3 and As2O3 in equal concentrations. Nevertheless, native oxides of GaAs have not been found to be homogeneous. In most cases, the native oxide can be described as a single phase, non-stoichiometric compound which alters the macroscopic picture of mixed stoichiometric binary oxides [1-3]. The initial bulk-like mixture of GaAs oxides is changed into pure Ga2O3 after exposure to temperatures above 4500C by evaporation of the volatile arsenic species [4-6]. The oxides become partly crystalline and are primarily composed OfP-Ga2O3 at temperatures above 50O0C [7-10]. At a temperature of about 50O0C, the morphology evolves as a function of time. In the early stage, well-oriented arsenic crystallites are formed [11], and as the oxidation proceeds the polycrystalline grain boundaries of the Ga oxide provide an easier diffusion path for As to escape [9,11]. In summary, polycrystalline Ga2O3 is formed at temperatures above 4500C whereas oxide films grown below 4000C almost always have an amorphous structure [12]. B

OXIDE GROWTH CONDITIONS

After the initial layer of an oxide has been formed, further oxidation cannot take place without diffusion of the reactant species, i.e. Ga and As, towards the oxide surface. On the other hand, the oxidising species must migrate towards the GaAs/oxide interface. Thus, oxidation is highly dependent on the temperature since the growth of a thick oxide is rather diffusion-controlled [8]. After anodisation under ideal growth conditions, the oxide/GaAs interface region has a width which is comparable to or less than that found in thermally oxidised Si/SiO2 systems (< 10 A) [13]. Since the surface structure and composition of GaAs evolve as a function of the cleaning procedure [14-19], such as chemical etching, sputtering, thermal desorption and plasma cleaning, it introduces variability in the oxide structure and interface composition which are dependent on surface stoichiometry, surface perfection and contamination. As a consequence, it is possible to modify the growth process, and therefore the composition of the native oxide layer, by either first depositing a thin metallic film such as Al [20-26] or Si [27-33] or by a sulphur passivation prior to oxide growth [34-37]. Experimentation with reduced-angle TEM [38], SEM [15] and other analytical techniques has shown that thin native oxides are often not uniformly deposited on etched GaAs (100) surfaces in air. In fact they are similar to the cluster formation of anodic oxides [39], with clusters of

typically 80 A wide and a few hundred A long [38]. Layer-by-layer growth of native oxides however takes place on cleaved GaAs (110) surfaces [40]. The physical adhesion of oxygen with such thin native layers is quickly affected by the high local fields of an SEM so that no measurements are normally possible in air. However, it is possible to stabilise this surface using alkaline etchants or to reduce or avoid oxidation by photochemical etching or sulphur passivation, respectively [15]. Changes in topographic STM images during oxidation in air can be seen. Enhancement of oxidation by illumination is possible. Optically-induced oxidation with deep UV light is stronger than with near-UV or visible light [41]. However, oxides formed with photoexcited oxygen appear to be closer to a Ga-enriched thermal oxide film if grown at elevated temperatures ( 2500C ) [42]. These results are altered by UV-ozone oxidation of GaAs at room temperature, which is expected to be far from thermal equilibrium, and thus results in an Asenriched oxide layer [2,43]. C

THERMODYNAMICS

Oxidation of GaAs surfaces is initiated by adsorbing oxidising species, e.g. O2, H2O, CO or N2O. As the reaction proceeds true oxidation takes place after some bonds of the GaAs substrate are broken due to a charge transfer from electrons of the valence-band bonding orbitals to the conduction band, and hence the surface composition is changed. The chemical compounds formed by this reaction are those which yield the lowest possible energy for the oxidation conditions. The equilibrium phase diagram serves as a guide to the composition of a thermal oxide and to any changes that may occur during thermal annealing [44]. Equilibrium thermodynamics predict the end products of a reaction which has overcome all kinetic barriers due to a high thermal energy such as heat. Plasma oxidation as well as UV-ozone oxidation are usually expected to yield reaction products far from thermal equilibrium. However, the elevated temperature during annealing of these kinds of oxides may provide the energy necessary to overcome the chemical and diffusion barriers and to form reaction products close to a thermal oxide [45]. Hence, the Ga-As-O phase diagram as shown in FIGURE 1 indicates that oxidation in thermal equilibrium results in the formation OfGa2O3 and elemental As [44,47]. However, the compounds As2O3, As2O5 and GaAsO4 may also form. These compounds are not stable in the proximity of the GaAs substrate and can decompose through the following reactions [2]:

D

As 2 O 3 + 2 GaAs - Ga2O3 + 4 A s

(1)

3 As2O5 + 10 GaAs - 5 Ga2O3 + 16 As

(2)

3 GaAsO4 + 5 GaAs - 4 Ga2O3 + 8 As

(3)

VAPOUR PRESSURE

The vapour pressure of GaAs and its related oxides is an important parameter in thermal oxidation, the deposition of insulators at elevated temperatures, and any kind of thermal processing such as contact alloying, diffusion, post-implantation annealing and particularly

FIGURE 1. Ga-As-O ternary phase diagram for temperatures below the melting point of As2O3 [46].

epitaxial overgrowth. It is well known that the column V elements, such as As, evaporate by forming polyatomic species like As4. The arsenic trioxides do not dissociate upon evaporation but rather form As4O6. Ga and its oxides have a low vapour pressure and dissociate into all vapour species such as Ga, GaO and Ga2O, with the Ga2O being the most abundant. E

OXIDE COMPOSITION

AES, XPS, and UPS analysis give two types of oxides, called a and P (owing to the chemical shift of electron emission and the distinct difference in the XPSAJPS spectra) [3,48]. The former one is less tightly bound and desorbs at temperatures up to 30O0C whereas the latter is more stable and desorbs at 4800C [17,48]. Thin native oxides formed by an exposure of GaAs at room temperature to various oxidising environments such as air, water, chemical solutions, any oxygencontaining gases etc. can exhibit a composition of chemically adsorbed oxygen (P- type) and physically adsorbed oxygen (a- type) whose ratio depends on the type of environment, the crystal orientation, and the perfection of the GaAs surface [49-53]. As reported by Nienhaus and Monch [52], oxide molecules are primarily physically absorbed (physisorbed) on GaAs (110) at temperatures below 160 K. These O2 molecules are converted into chemisorbed oxygen on increasing the temperature up to 170 K whereas effective oxygen chemisorption sets in above 190 K [52]. Oxides formed with excited oxygen and rather thick oxides (above 100 A) in general appear to be composed of primarily Ga oxide with elemental As at the interface (see above) which corresponds to the Ga-As-O phase diagram [47-56]. Nevertheless in most cases, kinetic rather than thermodynamic factors determine the observed phases [55] which is particularly relevant for plasma-grown oxides and UV-ozone oxides grown at room temperature. For low-temperature oxidation (350 - 4500C), accumulation of elemental As at the interface tends to increase due to a lower diffusion coefficient through the amorphous oxide, (see Section A). An oxidation-induced disordering of the GaAs surface has been generally observed for thick,

native oxides assuming a high-oxygen coverage [57-59]. Photoemission spectra in [1,60] led to the conclusion that oxidation takes place by a multi-coordinated adsorption near the semiconductor surface, i.e. it occurs in a few monolayers (ML) below the GaAs surface. F

INTERFACE STATES

Due to a low adsorption of excited oxygen (< 104 L) GaAs oxidises through one phase [1,3,5263] which involves simultaneous oxidation of both Ga and As. However, increasing the amount of oxygen chemisorbed at the GaAs surface causes local lattice rearrangement and increases the local strain and heat of adsorption [1,53,57,59] which results in strong pinning of the surface Fermi level [52,57,59,64]. Both arsenic vacancies ( as observed after plasma processing ) as well as arsenic antisite defects (in the presence of excess As) are probably responsible for the Fermi level pinning near the midgap in both p- and n-type GaAs [54, 65-71]. The thermally stimulated current (TSC) measurements in [72,73] led to the conclusion that a U-shaped continuous distribution of surface states is established at the oxide/GaAs interface. As a consequence, the CV analysis of MOS-type structures yields rather high carrier-injection type hysteresis and large frequency dispersion due to the high interface state density. Additionally, accumulation of As at the interface, predicted by the Ga-As-O phase diagram, may compensate an increase of interfacial conductivity [H]. The highly-disordered interface after oxidation, which may be far from a stoichiometric composition, also hampers the effectiveness of GaAs passivation and therefore limits device performance and reliability [74-77]. G

GaAs SURFACE PROCESSING

Surface structures formed on GaAs after applying different procedures, in particular air exposure, introduce a variability in surface composition. First, natural oxide layers grown immediately when exposed in air exhibit considerable differences in the oxide formation and composition to a degree which is not acceptable for device processing. Moreover, oxidation processes, such as UV-ozone oxidation (see Datareview 13.7), anodisation etc. may not alter the composition of the native oxide already existing on the GaAs surface. On the other hand, the concentration and composition of airborne organic molecules represents the largest variable in determining contamination sources [5,37,42,43,78-80]. In class 10 clean rooms, for example, an average concentration of 30 ^g/m3 of organic vapour has been measured [81]. Of course, the hydrocarbon concentration may differ at different locations within a CLIO room, e.g. depending on rural or urban environments. One indication of the total surface arrival rate of all organic species is given by 2 pg/cm2 which may result in an equivalent surface coverage of 1 monolayer (ML) formed in about 30 min [5]. In this context, one should be aware of the fact that even a very small amount of carbon contaminant at the oxide/GaAs interface will affect the electrical properties of epitaxial layers [82, 83], since carbon-related defects have been found to propagate into the epitaxial overlayer [5,48, 84-86]. Thermal desorption is usually applied as an appropriate pre-epitaxial surface preparation. However, complete desorption of hydrocarbons cannot be achieved because hydrogen usually gets lost and carbon bonds are formed with the GaAs surface [42, 87-89]. Existing carbon fractions have been reported to survive temperatures that may exceed the decomposition

temperature of the GaAs substrate [5,47,90]. One approach to improve this severe situation is to remove, in-situ, the native oxide layer, including surface contamination, immediately prior to a process step, e.g. epitaxial overgrowth, oxide growth, insulator deposition, metallization etc. In-situ surface pre-treatment has been successfully applied using a hydrogen plasma (see Datareview 13.5), and in electrochemically formed Schottky metallization (see Datareview 13.4). However, there are a considerable number of processes at which in-situ oxide removal may not easily be accomplished. In this case, a shortterm protection layer formed onto the GaAs surface may be used. This may be either a hydrogen passivation or a hydrocarbon-covered surface in order to saturate surface dangling bonds. Such an approach of preparing appropriate surfaces using passivating overlayers aims to control the GaAs surface structure and composition (for a review see [5]). Later on, this 'sacrificial' overlayer is readily removed just prior to any process step. Thermal desorption of this protective layer, for example, is used prior to epitaxial processes to yield stoichiometric GaAs surfaces appropriate for epitaxial overgrowth [5]. H

CONCLUSION

In the case of GaAs, thermal oxides and the formation of thick oxides (which actually is diffusioncontrolled and hence close to thermal equilibrium) usually result in non-stoichiometric oxide films of Ga2O3 and As 2O 3 as well as in pile-up of elemental As at the GaAs/oxide interface. This structural composition, which is mostly found in many oxidation processes, is clearly predicted by the Ga-O-As phase diagram, but provides poor electrical isolation and semiconductor surface protection. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] II1] [12] [13] [14] [15] [16] [17]

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13.3 Thermal oxides of GaAs H.L. Hartnagel and R. Riemenschneider July 1995

A

INTRODUCTION

The oxidation process of GaAs depends significantly on the oxidation temperature since the growth kinetics and oxide composition are completely different after formation at either low (less than 3000C) or elevated temperatures (400 - 6000C). Closely related to an oxide grown at low temperatures is the formation of a natural GaAs oxide in normal air at room temperature. For a particularly good understanding of the physico-chemical processes involved, the investigation is performed under vacuum with a controlled inlet of the oxidising species. Since the growth temperature strongly influences the oxidation process, the structural composition of the GaAs oxide and the oxide/GaAs interface is addressed in greater detail. This review therefore is divided into the following two parts. B

LOW-TEMPERATURE OXIDES

Oxidation of GaAs at temperatures less than 400 0 C including at room temperature starts with an adsorption of the oxidising species. A strong interaction between oxygen and the GaAs surface has been found in a first phase due to chemisorbed oxygen (P type). The initial phase is followed by a much slower phase which is more akin to physisorption of oxygen (a type) (see Datareview 13.2, Section E). There is considerable evidence [1-3] that a cleaved surface exhibits oxygen chemisorption which proceeds layer-by-layer, with a large kinetic barrier to oxidation beyond the first two monolayers. Extensive studies of oxidation have been carried out on cleaved GaAs with a (110) single crystal surface at room temperature [4]. The oxide of GaAs has been shown to thicken as a logarithmic function of storage time, e.g. a thickness up to 43 A has been observed after 10 min exposure in air [5,6]. It is now widely accepted that both Ga and As atoms are involved to an equal extent in the bonding with oxygen at a coverage from 0.1 to at least 1 monolayer (ML). Detailed studies have been performed using molecular oxygen or nitrous oxide in order to study the initial oxide growth [7-9]. With increasing oxide thickness, oxygen reacts to a differing degree with Ga and As, i.e. one approaches a Ga-enriched thick oxide composition in which the amount of arsenic oxide decreases [9,10]. Since the initial oxide stages depend on the physico-chemical surface reactions, crystal orientation as well as the state of surface perfection generally have a strong influence on the oxidation kinetics [I]. Interaction of oxygen with (100) and (111) GaAs surfaces depends in a more complex manner on the surface composition. The Ga-rich faces of (111) surfaces give rise to a more rapid adsorption [1,4], since the exposure at which the main uptake of oxygen sets in decreases with increasing amount of surface Ga atoms [11,12]. For (311) the main oxygen uptake was also found to start in a similar fashion (100) or (110), but after some exposure the surfaces contain 20% less oxygen as compared to the other two surface types [11]. A cylindrical GaAs sample allows the measurement of the orientation dependent deposition rate, which is strongest for (111) [13-15]. Moreover, it is suggested that oxide formation on (100) surfaces is highly

dependent on the surface cleaning and the chemical composition of the GaAs surface. During the HF treatment of GaAs (100), for example, the formation of an elemental As film has been observed [5,16] which grows linearly with the logarithm of immersion time. It is shown in [17-19] that the GaAs native oxide depends on the kind of chemical solution and on the sequence of the aqueous solutions used for surface pretreatment prior to oxidation. The reason is twofold: firstly, the stoichiometry of the GaAs surface depends on the kind of the chemical solution [5], and secondly, acid or alkali molecules left on the GaAs surface are able to enhance or diminish the oxide growth rate. The dependence of oxide formation on surface perfection suggests that defects play a key role [4]. In an early work of Barton [20], it is suggested that defect sites may serve to dissociate molecular oxygen releasing atomic oxygen which then forms strong bonds to Ga and As sites. Therefore, oxidation may be enhanced by simultaneous X-ray radiation and surface bombardment by ions or electrons as reported in [11,21-24] and in Datareview 13.7. Native oxides can play a rather crucial role in the parasitic electrical effects of the GaAs surface and may adversely affect device performance and reliability [25,26]. Some collection of carbonaceous agents is possible when oxidation is carried out in air [10,17,27,28], whereas oxidation under vacuum conditions using a defined inlet of oxidising species prevents such carbon contamination. For surface passivation to be effective, airborne oxides, and in particular surface contaminants, have to be eliminated and some other passivating films, such as encapsulating, protective layers (see Datareview 13.2), have to be deposited. Excited species of hydrogen, e.g. in a hydrogen plasma (see Datareview 13.5) are known to remove the GaAs native oxide at low temperature. On the other hand, native oxides are desorbed at temperatures above 5000C (an appropriate pre-epitaxial surface preparation), yielding a stoichiometric composition of the GaAs surface [17,27-31]. At temperatures of about 300 4500C, As oxides are desorbed while Ga oxides are removed at 450 - 600 0 C [30-36]. However, the temperature at which the thermal removal of native oxides can take place has been shown to depend on the oxidising process and oxide composition [37]. Thus, a difference of more than 50 0 C between UV/ozone-formed and airborne oxides has been observed [38]. Moreover, as found in [17], air-exposed GaAs surfaces never have a stoichiometric surface composition after thermal desorption and carbon species are not completely removed. In order to realise high performance GaAs circuits, the preparation of a good quality SiO2/GaAs interface is highly recommended. The presence of a poor quality, highly strained, oxide formed at low temperature could lead to degraded device performance [39-42]. Detailed work carried out on thermal CVD [39,40] in order to study the native oxide growth during SiO2 deposition shows evidence of an intermediate layer consisting of GaAs oxides and excess As. These results suggest that a native oxide film is formed prior to SiO2 deposition [42] whose composition is close to thermodynamic calculations which is different from bulk-like proportions. It has been pointed out in [20,43], that during the passivating processes the arsenic species tend to escape from the GaAs surface if the substrate temperature exceeds 3000C. Thus, passivation is performed below 3000C which ensures that As out-diffusion from the GaAs surface is prevented and that thermal oxidation of the GaAs surface is inhibited. Both these features are achieved when the oxidation process is carried out in an arsenic or phosphorus atmosphere [20,44-47]. Alternatively, a thin surface protective layer, such as Si or Al, may be used to control surface diffusion and thus the oxidation mechanism [36,42,45,46,48-54].

C

THERMAL OXIDES GROWN AT HIGH TEMPERATURES

The structural and chemical composition of the thermal oxide and the GaAs/oxide interface can vary significantly depending on the methods and oxidation conditions. The growth of the GaAs oxide is influenced by various factors as follows. Cl

The Temperature of the Oxidation Process

In the literature many articles are concerned with thermal oxidation of GaAs (as a review see [5557]) using a variety of analysis techniques such as TEM [58], XPS [59], AES, UPS [13-15,6062], SIMS, Raman spectroscopy [63], photoemission spectroscopy [64,65], ellipsometry [60], electron-energy loss spectroscopy (EELS, [66,67]), electron diffraction (e.g. LEED [60,68]) etc. Most of the results are consistent with the following oxidation procedure. At low temperatures, i.e. below 4000C, an amorphous oxide is formed which consists of a mixture OfGa2O3 and As2O3, and the Ga oxide is more abundant [8,48]. Oxidation temperatures between 450 and 500 0 C initially produce an epitaxial film of Y-Ga2O3. As the reaction proceeds this film becomes polycrystalline and then transforms into P-Ga2O3 (y, P represent here the crystalline structure and not the type of a thin oxide). Studies carried out using clean, polished (111) substrates reveal a parabolic growth in the low temperature range which changes into a linear growth at higher temperatures, i.e. 400- 450 0 C, and 480 - 5300C, respectively [I]. At higher temperatures, i.e. above 5000C, the polycrystalline grain boundaries provide an easier diffusion path for elemental As to escape from the interface [69,70] or for oxygen to penetrate towards the interface indicating that oxidation occurs at the interface [71,72]. Some authors [48,55,70,71,73,74] favour the formation of arsenic microcrystallites at the GaAs/oxide interface at low oxidation temperatures (400 - 450 0 C) and, as a consequence, further oxidation is hampered due to reduced diffusion through the well-oriented arsenic film. In this context, the effect of temperature is twofold; increasing the temperature accelerates oxygen diffusion processes [71] and hence enhances the oxidation rate as expected, but apparently also prevents a formation of a continuous arsenic layer which is then dispersed in the Ga oxide [55]. A corresponding effect will occur with time when oxidation is performed at lower temperature (but above 4000C) [70,71]. It should be noted that the accumulation of interfacial crystalline As gives rise to a highly conductive layer, [70,73], which of course has strong implications on the electrical stability of devices subjected to oxidation. The chemical composition of most oxides resulting from thermal oxidation at elevated temperatures is the product of predicted equilibrium thermodynamics [31,75]. This undoubtedly is due to the increased thermal energy which allows kinetic barriers to be overcome. The temperature at which the accelerated growth of GaAs oxides occurs is between 350 and 400 0 C [57]. The thermal oxide is primarily composed of Ga oxide with some As oxide, since As and arsenic oxides are more volatile as compared to the Ga species. Thus, evaporation causes loss of arsenic agents if they can reach the surface [1,69,70]. Excess As has been widely reported to pile up at the interface [48,55,76,77], since elemental As is one of the products expected from the Ga-As-O phase diagram. Some people argue that an accumulation of As is associated with the poor diffusibility of As through the Ga oxide layer.

Moreover, there are still open discussions on the morphology of the As-enriched interface, i.e. whether excess As is present in elemental, amorphous, or crystalline form. Thick native oxide films prepared by thermal oxidation in particular are utilised for masking purposes, e.g. for ion implantation or selective epitaxy [78,79]. Furthermore, they are used in connection with pre-epitaxial preparations as well as short-term, protective layers in order to avoid contamination (e.g. carbon) when exposed in air. Hence, these thermal-oxide films are supposed to be readily desorbed, and they are needed to provide an almost stoichiometric semiconductor surface appropriate for epitaxial overgrowth after being desorbed [10,17,29,34,38,80] (see also previous paragraph). C2

The Presence of Dopants / Impurities in the GaAs Substrate

Several attempts have been reported on measuring the effects of doping on the oxidation rate [55,71,73], and both enhancement and inhibition of the thermal oxidation rate has been observed depending on the type of dopant. Experimental results show a retardation effect for Si and Zn dopants on the thermal oxidation rate of GaAs [78]. In [55] it is shown that Cr-doped GaAs exhibits slower oxidation kinetics as compared to Si- and Te-doped GaAs samples. Moreover, the presence of dopants may lead to a considerable enhancement of the crystalline As built up at the GaAs/oxide interface [73] or may accelerate the escape of excess arsenic at the interface. Although the specific details are still under discussion, it is generally agreed that diffusion coefficients, and thus oxide composition and kinetics, are significantly influenced by the doping. In any case, the concentration of dopants in the GaAs bulk has been found to be significantly decreased following oxidation [55,73] indicating incorporation of dopants in the oxide or segregation at the GaAs/oxide interface. In particular, the morphology of excess As present at the interface is affected by dopants or other impurities incorporated into the oxide [73,81], e.g. crystallisation of As is retarded in the case of Si or Zn doping. Thus, it may be concluded that thermal oxidation of doped GaAs is expected to have rather strong implications on material and device properties, and thermal oxidation may be assumed to be one of the reasons for a gradual degradation in the performance of GaAs-based devices and circuits. C3

The Addition of Oxidation Promoters into the Process

The presence of impurities may diminish the oxide-induced disordering of the GaAs surface by controlling the migration of Ga-As-O species [82-85]. In addition to the effect of doping, some authors claim that the co-deposition of different compounds, like Sb2O4, Bi2S3, GeS2, Bi2O3 or PbO [86-89], may enhance oxide formation and may positively affect the interface quality. This effect appears to be mainly due to a change in the oxidation kinetics, explained by a possible oxygen transfer from oxidised compounds towards the Ga and As atoms. This effect is similar to autocatalysis. Thermodynamic calculations, e.g. for the Bi2O3-Ga system [85], confirm this understanding (AG ~ 470 kJ/mol at 8000C). Inclusion of other types of volatile components such as GaCl3, InCl3 and PbCl3 [84,85,90,91] in the oxygen atmosphere actually improves the electrophysical parameters of the oxide layer mainly by preventing the escape of As. A similar improvement has been observed after the implantation of phosphorus into GaAs which also prevents a deficiency of the group V element and promotes the dielectric stability of the thermal oxide [92].

D

CONCLUSION

During thermal oxidation, several steps are involved, i.e. absorption, dissociation, diffusion, and oxidation itself, which are all thermally activated. Hence, the reaction proceeds in a manner which is close to thermodynamic equilibrium yielding a non-stoichiometric oxide composition as described in Datareview 13.2. At elevated temperatures, the activation energy is provided by heating. In this case, other processes compete with the oxidation process since the arsenic species tend to evaporate leading to an oxide film that is strongly As deficient. In principle, this phenomenon may not be altered by dopants, surface protective layers (such as Al, Si or others), or elements which promote oxidation kinetics. Nevertheless, a profound influence of such elements on the GaAs oxide structure and composition has been reported. REFERENCES [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17]

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13.4 Wet oxidation of GaAs

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H.L. Hartnagel and R. Riemenschneider July 1995

A

INTRODUCTION

GaAs can be oxidised in most oxygen-containing liquids. The effects relate either to simple immersion into an aqueous agent, chemical oxidation, or to the simultaneous application of an anodic current, anodic oxidation. B

CHEMICAL OXIDATION

The growth of an oxide in a chemical solution or vapour is a simple way of accomplishing this process. Equilibrium thermodynamics provide a guide to the final residue products. In some cases it is desirable to etch the GaAs surface just prior to its oxidation in order to remove the air-grown oxide and to avoid surface contamination. When considering a possible aqueous solution, one has to take into account the differing chemical reactivities of the two components, As and Ga, and their oxides, i.e. arsenic oxides are highly soluble in water whereas gallium oxides are only soluble in strongly acidic etchants. As a consequence, the chemical composition of the GaAs surface after a wet chemical treatment is also dependent on the pH value of the solution used [I]. Thick native oxides can be grown in H2O2 or HNO3 solutions. The latter one is an acidic etch but is itself also an oxidiser. It will grow a Ga-As oxide and simultaneously dissolve any gallium oxide. Thus, HNO3 leaves an As-enriched oxide on the GaAs surface [2]. Infrared spectra of wet oxides chemically grown by immersion in H2O2 exhibit incorporation of hydroxyl groups and interstitial water to a greater extent than anodic oxides [3]. A wide range of oxide thickness has been reported since oxide growth depends not only on the chemical composition of the oxidising agent but also on the doping level, incident light, substrate conductivity and bath temperature [2]. Exposure to aqueous solutions of various types yields a solution-dependent oxidation rate of highly different a/p-ratios [4,5] (see Datareview 13.2, Section E). The a-type oxygen is mainly physically adsorbed whereas the p-type is chemically adsorbed. There are indications [6] that alkaline solutions produce smooth surfaces whereas acidic solutions give rough, island-type features. Other liquids, such as methanol, are also reported to result in oxide growth [7]. An in-depth analysis of the chemical composition of GaAs surfaces has been performed in [8,9]. It is concluded that the stoichiometry of the GaAs surface can be controlled by adjusting the chemical solution. The growth rate on exposure to stationary deionized water was found to be 10-20 A/h [1O]. Exposure to deionized water and O2 enhances the growth of the natural oxide (up to 1000 A) [2,11] while water rinsing in an N2 atmosphere has no noticeable oxidising effect [2,12]. Experimental results from so-called photochemical washing (PCW) gave rise to the growth of a native oxide which is highly porous and unreliable [13-17]. The adsorption rate of water molecules on cleaved GaAs (100) surfaces has been shown to be 3 orders of magnitude greater than for dry O2 [18,19]. The water molecules adsorb dissociatively [20], and OH- and H-bonds accelerate the reaction with the GaAs surface [21,22], but are also incorporated in the GaAs/oxide

interface [23,24]. Water vapour in the air is generally known to accelerate the oxide growth [2,25,26]. However, a uniform and chemically stable oxide film may be grown in an atmosphere of oxygen and water vapour at low temperature (2500C) under high pressure (-2 atm) by reducing kinetic limitations [21,22,27,28]. C

ANODIC OXIDATION

Growth kinetics of anodic oxides on III-V semiconductors are quite complex due to 'competition' between the two compound elements being either dissolved or oxidised. Hence, etching or oxidation may be selective and strongly affected by anodic current density and the electrolyte in terms of temperature, pH value and the anion species present [1,29,30]. In early studies, nucleation and island growth were reported for GaAs [31,32]. AFM surface analysis has been used to show that island growth does not occur when voltage pulses are applied [33-36]. During anodisation both in-situ removal of the air-grown oxide and subsequent metallization or oxidation are easily accomplished for applications such as MIS-Schottky diodes and nearly ideal Schottky contacts [33,34,36-39]. The chemical composition of the anodic oxide is widely accepted to contain Ga oxide and As oxide in approximately equal proportions. Recent studies using in-situ IR spectra (with multiple internal reflections) found a structural oxide composition which is not homogeneous and quite different from a simple mixture OfGa2O3 and As2O3 [3,40,41]. This may be one reason why anodic oxides are ineffective masks for Zn diffusion [42]. A transition layer with a structural composition which appears to be different from the bulk Ga-As ratio has been found [40]. The composition of the oxide/GaAs interface after anodisation is still the subject of discussion especially related to the presence of elemental As at the interface [3,29,41,43]. Metal-oxide semiconductor (MOS) structures fabricated on (T T T) or (100) oriented surfaces reveal a rather large change in the distribution of the density of states [29,44-46]. In principle, the interfacial defect structure is different in the anodic oxide on GaAs (100) and (T T T) surfaces [44,45]. However, the large density of surface states suggests an anodic oxidation which preferentially forms a Ga oxide at the interface and which gives rise to the existence of a thin Ga2O3 layer [47,48]. The interface of the oxide/GaAs of an MOS structure can be characterised by measuring the Fermi level position using CV measurements [49]. For anodically-grown oxides, CV analysis shows considerable hysteresis at high frequencies and a large frequency dispersion of the capacitance under an accumulation bias which is due to carrier injection from interface states and from traps in the oxide [29,45,50]. A minimum density of interface states has been determined to be in the order of Nss ~ 1012 cm^eV"1 [44,45]. Research efforts [44] report a conduction increase of MOS-type structures at about 5 x 104 V/cm. DLTS measurements show that anodisation of GaAs results in an increase of the EL6 type defect [51]. These results are altered in the case of anodically formed Schottky contacts which have been demonstrated to be defect free [34,35,38]. As suggested by [44,45,52-54], thermal annealing after anodic oxidation may be used to reduce the density of interface states. During thermal annealing, the arsenic oxide tends to decompose and form elemental As (as predicted by thermodynamics) at approximately the same temperature at which crystallisation of the Ga oxide is initiated (above 4500C). In this context, Weiss and Hartnagel [55] reported that the oxide layer forms P-Ga2O3 crystallites more readily with increasing current density during anodisation [55,56], and consequently, the anodic oxide films

yield a poor dielectric breakdown behaviour [57]. Thermal annealing in an arsenic atmosphere is reported to improve the electrical properties of GaAs MOS structures indicating that evaporation and escape of the arsenic oxide is inhibited [45]. Some efforts have been made to improve the electrical properties of the anodic MOS structures by introducing either a thin metallic layer or sulphur onto the GaAs surface in order to control the migration of the Ga-As-O agents. An initial deposition of a thin metal film on GaAs, e.g. aluminium, and subsequent anodic oxidation completely through and beyond the AVGaAs interface [58,59] eventually forms a new oxide, and also may completely change the interfacial structure. More recently sulphur passivation prior to the anodisation process has been found to also improve the electrical properties of an MOS structure. However, post-annealing of the anodic oxide is still required since the homogeneity and electrical stability of the anodic film degrades the characteristics of the MOS structures [52,60]. These results show that sulphur passivation stabilises the structural composition at the interface and prevents arsenic segregation and out-diffusion during the annealing process [61]. Nevertheless, it remains an open question whether sulphur affects the doping concentration and reliability of the GaAs substrate. D

CONCLUSION

Numerous investigations carried out on GaAs anodisation have proved the presence of many peculiarities in the electrical behaviour of MOS diodes. Among these, considerable frequency dispersion of capacitance and large hysteresis in both I-V and C-V characteristics have been observed. These results on anodically-formed MOS structures suggest that the electronic quality of GaAs anodic oxides is essentially hindered by a high interface-state density and a large density of traps present in the oxide, see [54,62,63]. Many people have analysed and used anodic films, especially for MOSFET applications. Although chemical anodisation is a relatively simple low temperatuare method with the advantage of introducing few defects it remains a rather complex process. It remains a challenging task to grow a homogeneous, dry and dense anodic oxide film (without post-annealing), and the deposition of insulators is considered to be a more favourable method to meet the requirements of a mature technology. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] [II]

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13.5 Plasma oxidation of GaAs H.L. Hartnagel and R. Riemenschneider July 1995

A

INTRODUCTION

The surface of a metal or semiconductor is oxidised when it is exposed to a gaseous plasma containing oxidising species, such as O2, N2O or CO2. The RF plasma is usually capacitively or inductively excited while the semiconductor is in close contact with the plasma. Recent developments of indirect plasma sources such as electron-cyclotron resonance (ECR), microwave downstream, and remote-plasma systems utilise a plasma cavity which is separated from the substrate. The plasma thus works only as a source for excited oxygen atoms which are drifting towards the GaAs surface avoiding an interaction between charged particles and the semiconductor surface [1-6]. B

PLASMA ANODISATION

If a DC bias is applied to the substrate during the plasma process, oxidation takes place in a similar way to the wet anodisation process described in Datareview 13.4. Moreover, there are many other similarities between aqueous solutions and the behaviour of a plasma. In some early publications [7-9] it was found that the oxide thickness increases linearly with oxidation time in an oxygen plasma if the anodic current is held constant. The growth rate varies from 6 - 2 0 A/s [8-10], and is strongly dependent on the substrate temperature [9,11,12]. The growth rate is relatively independent of the gas pressure and the plasma parameter indicating a diffusion-controlled process. As-grown plasma films have been found to be uniformly composed of Ga2O3 and As 2 0 3 in approximately equal concentrations. The interface appears abrupt, and may contain a low concentration of As [3,5,13-16]. However, impinging ions can sputter the surface, and thus lead to a reduced growth rate and to a modification of surface stoichiometry due to a preferential sputtering of the arsenic component [17-20] (see also Datareview 13.7). From DLTS results [9,12,21-27] it is pointed out that the conventional (not indirect) plasma technique is not a desirable process for device applications since the bombardment of highly energetic ions may adversely affect the electrical properties of the GaAs surface [11,14,20]. Some authors [26,28,29] report evidence of a significant decrease of electrical conductivity after the plasma due to a highly disturbed surface layer. The thickness of this modified layer appears to depend linearly on the substrate bias [29,30] and its composition exhibits a departure from stoichiometry due to arsenic loss [13,31], a factor which increases with plasma power [12,32,33]. Thus, plasma parameters, such as RF frequency, RF power and gaseous pressure may not affect the oxide growth, but they do affect the degree of GaAs surface degradation during the initial stage of oxide formation [24,27]. The reader may be aware that this situation is in contrast to wet anodisation which is assumed to be damage-free. Thermal annealing after plasma oxidation may be applied in order to reduce the density of interface states. As expected from the results in Datareview 13.4, annealing causes evaporation of the arsenic oxide from the surface and its decomposition into elemental As at the interface [H].

The elemental As results from the decomposition OfAs2O3 at the GaAs substrate at temperatures of more than 450 0 C, as predicted by thermodynamics. TEM was used to show that after annealing at 600 0 C the interfacial As is in the form of crystalline clusters and not uniformly distributed [34]. A thin aluminium (Al) film ( 2 - 1 0 nm) has been deposited on the GaAs surface in view of its ability to control the relative concentration of As/Ga [9,35-37]. It has been pointed out in [37,35,25,36] that oxide formation proceeds by an initial rapid grain-boundary oxidation of the Al followed by a slower oxidation towards the centres of the individual grains. Continued oxidation results in the growth of Ga-As oxide layers on both sides of the Al-oxide film analogous to wet anodization. AES and TEM [37] have shown that an amorphous Ga-As oxide layer is formed which is composed of Ga oxide and As oxide in equal concentrations all the way to the GaAs substrate. Although the structural composition appears to be improved by depositing an Al interfacial layer no marked effect on the electrical properties has been reported [9,25,35,38]. C

GaAs SURFACE STRUCTURE

The possible role of hydrogen present in an oxygen discharge has been discussed in [39,26,29] with respect to a passivation of donors in the surface region and subsequent formation of a highly resistive layer. A hydrogen plasma using H2 or NH3 is widely applied which provides in-situ removal of the native oxide and surface cleaning [1,6,13,14,39-43], e.g. prior to anodic oxide growth. Recently, low-damage, indirect plasma processes have been investigated in detail with respect to hydrogen discharges [1,44] indicating that the formation of a surface-depleted layer is closely related to the extent of plasma-damage [32]. A radiation-enhanced indiffiision of H or O may explain the large thickness (more than 20 nm and 200 nm for O and H plasmas, respectively) of these disturbed layers [17,32] that are absent in the case of damage-free plasma processes. However, indirect plasma processes show evidence of physisorbed water formed in the process and of a departure from surface stoichiometry which is due to the preferential oxidation of Ga in the proximity of water at the GaAs surface [5,33,45,46]. In contrast, the chemical structure of ECR-grown oxides appears to be similar to that of oxides formed using atomic oxygen sources (for details see Datareview 13.7) suggesting a reaction of both Ga and As towards their highest oxidation level, namely Ga2O3 and As2O5 [I]. ECR-grown oxides have been reported to be uniformly composed of As-Ga oxides in a stoichiometric ratio right from the oxide surface to the GaAs bulk [I]. Since the oxide growth rate has been found to be almost independent of substrate bias, the oxygen flux may be assumed to consist of neutral atomic oxygen in both its ground and excited states [I]. The growth of any kind of oxide by a plasma process, including the PECVD deposition OfSiO2, encounters the preferential oxidation of Ga and a deficiency of As, particularly in the presence of a hydrogen plasma [3,5,40,45]. Oxide formation in an arsenic or phosphorus atmosphere is therefore favoured [23,47-49] to preserve the stoichiometry of the GaAs surface. However, this situation appears likely to be completely different when using high-density plasma sources, such as an ECR plasma. In this case, oxide growth and structure can be very similar to that produced by atomic oxygen as discussed in Datareview 13.7.

D

CONCLUSION

In conclusion, the application of a plasma-induced oxide in the vicinity of GaAs-MOS structures is not desirable due to a departure from surface stoichiometry, which is the likely cause of a high density of interface states [22,23]. Although indirect-plasma processes may systematically avoid disturbing the GaAs surface, the growth of plasma oxides still results in an As-deficient interface. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] II1] [12] [13]

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13.6 Laser assisted oxidation of GaAs LW. Boyd September 1996

The interaction and absorption of high intensity pulsed laser radiation with compound semiconductors induces melting and subsequent segregation of the constituent atoms [I]. When GaAs is irradiated in air or O2, both hexagonal Ga2O3 and cubic As2O3 are formed, although the top 500 A is found to be Ga-rich [2]. As with Si, experiments have shown that before any significant oxygen becomes incorporated, the incident beam intensity must be nearly three times that required to melt the GaAs, E(th) [3]. Between E(th) and 3E(th) no oxides are formed. It has been shown that the native oxide must be evaporated at this higher intensity before significant new oxide formation can occur by oxygen penetration through the melted region [4]. Oxide formation on GaAs without surface melting is clearly a much more attractive proposition for device applications [5,8] as it does not lead to atomic segregation. However UHV cleaning must be performed to remove the native oxide before experimentation. Bermudez [7] has shown that irradiation with incoherent light can also produce photoenhanced solid-state oxidation of GaAs in the presence of O2, but not in CO and H2O atmospheres. Since the latter molecules do not chemisorb, this suggests that the intramolecular bonding in the oxidising species must be weakened for the reaction to occur. It was proposed that the bonding in the adsorbate is weakened by charge transfer from the bulk, and oxidation proceeds with the photoenhanced breaking of GaAs bonds. Enhancement factors up to 20 have been observed in the times required to form the first monolayer of oxide, while the rate of photoenhancement did not strongly depend upon the type of excess carriers present in the material. Petro et al [5] have found that low intensity Ar ion laser irradiation (< 3 W/cm2) of cleaved GaAs substrates increases the sticking coefficient of O2 by a factor of 1000. Heating effects and oxygen excitation were ruled out as possible causes of the enhancement in favour of electronic excitations of the GaAs surface. It has been suggested that oxidation proceeded via absorption by a loosely bound oxygen molecule, followed by break-up of the molecule into reactive species [9]. Evidence for this model indicated that in the dark the uptake of oxygen by the GaAs was strongly temperature dependent, while under laser illumination the chemisorption was totally independent of thermal activation. After experiments comparing O2 and N2O, it was concluded that the reaction step which was photoenhanced by the visible radiation was the dissociation step [10]. In related work [11] NO was found to absorb and partially react to N2O on GaAs (100) substrates at a temperature of 90 K. After laser irradiation at wavelengths between 475 and 890 run, the NO desorbed and dissociated after interaction with photogenerated carriers on the surface, leaving the surface disordered and partially oxidised. Thus visible light has been shown to produce a surface coverage of up to 2 monolayers, with no evidence of any saturation [I]. Siejka et al [13] have observed preferential oxidation of Ga, in addition to some As loss during 193 nm excimer laser (ArF) irradiation in an 16O atmosphere of GaAs 18O enriched native oxides. Additionally, very efficient exchange of oxygen was observed after 1000 laser pulses, with the total oxygen content of the films being conserved. Since thermal heating was considered

negligible, the authors appeared to have found a new mechanism of laser assisted formation of highly mobile oxygen species in the GaAs. An independent study using CW-Ar ion laser radiation has also found optically enhanced oxidation reactions for GaAs [14] and for other III-V and ternary compounds [15]. It was found that after initial rapid oxidation, film growth proceeded in proportion to the logarithm of illumination time when intensities of about 1 kW/cm2 were used. Semiconductors containing Ga oxidised easily during the initial stage of irradiation while those containing As oxidised more readily than the others only after a certain amount of film growth, or at temperatures near 300 0 C. While calculations showed that the laser increased the temperature by only around 20 0 C, the growth rate was enhanced as if the temperature had increased by 50 0 C. Since the logarithmic oxidation behaviour did not change under laser illumination, it was suggested that the enhancement was due to accelerated transport of the elements necessary for oxidation. Photo-induced reactions with heterogeneous molecules such as NO or N2O also provide routes to oxide (and even nitride) formation on GaAs. Whilst experiments with NO indicated a much smaller sticking coefficient and subsequent oxide growth than with O2, Bertrand [16] found that oxidation of GaAs irradiated by 254 and 185 A radiation from a Hg lamp occurred in the presence OfN2O. In this case the As rich surfaces oxidised much more slowly, and saturation occurred once all the absorbed oxygen became incorporated to form the oxide. The oxidation of InP has been achieved using an ArF laser (193 nm) in an atmosphere OfN2O [17]. The process is most likely initiated by the photodissociation of the N2O molecules, since it was found that for a given constant laser power the oxide thickness increased with increasing N2O pressure. When the laser beam impinged directly on the sample surface during the film growth, the oxidation rate was enhanced by approximately 20%. In line with the photon enhanced Si oxidation models and some of the related ideas connected with enhanced reaction rates also discussed throughout this Datareview, it has been suggested that the increased concentration of electron-hole pairs could assist in the weakening and possible breakage of the lattice bonding thereby increasing the availability of atoms for the reaction. Films as thick as 160 A have been grown using this method, and these appear to be amongst the thickest laser grown oxides on compound semiconductors reported so far. In general, the role of electron-hole pairs and their precise energies, is thought to be crucial in the enhancement of the oxidation reaction. Oxidation of GaAs in a UV/ozone environment using a 50 W low pressure UV lamp has also produced a mixture of stoichiometric Ga and As oxides at less than 1 nm thickness [18], quite contrary to the oxides formed thermally, or even the native oxide. This compositional difference, which exhibits subtly different thermal desorption properties, may be due to the indiscriminate nature of the reaction of the excited oxygen arising from the ozone (O3) decomposition. Below the UV/ozone formed oxides, the GaAs was found to be As-depleted [19]. However, extended oxidation (60 min) found the oxides to be less stoichiometric, with an As oxide phase being segregated closer to the surface [19]. This oxidation technique has been applied to the preparation of surfaces prior to MBE growth [20]. Laser treatment in water can enhance the incorporation of oxygen in GaAs [20]. Using 30 ns duration pulses from a ruby laser at 694 nm ensured that most of the energy was deposited in the deionised water [21]. Both As and Ga oxides were formed, although below 1 J/cm2, the growth

was very small. Oxidation was more prominent at 2 J/cm2 although for a large number of pulses, sputtering was initiated. Whilst Ga oxides dominated for pulses below 1.5 J/cm2, with only a small As2O3 presence, for 2 J/cm2 both As2O3 and As2O5 were dominant. The mechanism of light assisted oxidation appears to involve physisorption of the oxidising species on to the surface followed by dissociation which is affected by an induced charge flow to the surface. The incident light will clearly enhance this current somewhat, but apparently only when surface recombination is high [12]. Thus, rather than the individual carriers participating in the reaction, it is their excess energy released during surface recombination which enhances the oxygen uptake. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] II1] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21]

I.W. Boyd [ Laser Processing of Thin Films and Micro-Structures (Springer-Verlag, Berlin, 1987)] M. Matsuura, M. Ishida, A. Suzuki, K. Hara [ Jpn. J. Appl. Phys. (Japan) vol.20 no. 10 (1981) p.L726-8] G.G. Bentini et al [ J. Phys. Colloq. (France) vol.43 no.C-1 (1982) p.229-341 ] C. Cohen et al [in Dielectric Layers in Semiconductors Eds G.G. Bentini et al (Les Editions de Physique, Les Ulis, 1986) p. 197-207 ] W.G. Petro, I. Hino, S. Eglash, I. Lindau, CY. Su, W.E. Spicer [ J. Vac. Sd. Technol. (USA) vol.21 no.2(1982)p.405-81] F. Bartels, L. Surkamp, HJ. Clemens, W. Monch [ J. Vac. Set. Technol. B (USA) vol.1 no.3 (1983) p.756-62 ] V.M. Bermudez [ J. Appl. Phys. (USA) vol.54 no.ll (1983) p. 6795-8 ] F. Bartels, W. Monch [ Surf. Set (Netherlands) vol.143 no.2-3 (1984) p.315-41 ] K.A. Bertness, W.G. Petro, J.A. Silberman, DJ. Friedman, W.E. Spicer [J. Vac. Sci. Technol. A (USA) vol.3 no.3 pt.2 (1984) p. 1464-7 ] K.A. Bertness et al [ Surf Sci. (Netherlands) vol. 185 no.3 (1987) p.544-58 ] S.K. So, W. Ho [ Appl. Phys. A (Germany) vol.47 no.3 (1988) p. 213-17 ] K. A. Bertness et al [Appl. Phys. A (Germany) vol.47 no.3 (1988) p.219-28 ] J. Siejka, J. Perricre, R. Srinivasan [Appl. Phys. Lett. (USA) vol.46 no.8 (1985) p.773-5 ] CF. Yu, MT. Schmidt, D.V. Podlesnick, R.M. Osgood [ J. Vac. Sd. Technol. B (USA) vol.5 no.4 (1987) p. 1087-91] M. Fukuda, K. Takahei [ J. Appl. Phys. (USA) vol.57 no.l (1985) p. 129-34 ] P.A. Bertrand [ J. Electrochem. Soc. (USA) vol. 132 no.4 (1985) p. 973-6 ] M. Fathipour, P.K. Boyer, GJ. Collins, CW. Wilmsen [ J. Appl. Phys. (USA) vol.57 no. 2 (1985) p.637-42 ] S. Ingrey, W.M. Lau, N.S. Mclntyre [ J. Vac. Sd. Technol. A (USA) vol.4 (1986) p.984 ] BJ. Flinn, N.S. Mclntyre [ Surf, and Interface Anal. (UK) vol. 15 (1990) p. 19-26 ] E. Bedel, A. Munoz-Yague, C. Fontaine [Mater. Sci. Eng. B (Switzerland) vol.21 (1993) p. 157160] SS. Shushtaiian, S.M. Kanetkar, S.B. Ogale [Mater. Lett. (Netherlands) vol.13 (1992) p. 325329]

13.7 Miscellaneous new methods of GaAs oxidation H.L. Hartnagel and R. Riemenschneider July 1995

A

INTRODUCTION

There are a number of approaches for GaAs oxidation which have been of major interest in recent years. Techniques for supplying energy to the GaAs surface to be oxidised are presented, such as deep UV/ozone exposure, or electronic, ionic and neutral particle bombardment. Usually, one important point of interest is the stoichiometric composition of the oxide/GaAs interface and possible damage introduced by the oxidation process. B

UV/OZONE OXIDATION

Recently, oxides formed by UV/ozone exposure have received much attention [1-5]. Air-grown oxides of GaAs never yield stoichiometric surfaces after thermal desorption [5,6] and UV/ozone oxidation was found to be a promising approach to solving this problem. Therefore, this photo process is primarily applied to prepare sacrificial thin oxide layers which are subsequently removed by vacuum thermal desorption in order to achieve GaAs surfaces that are suitable for epitaxial overgrowth. These GaAs surfaces have to be smooth and clean on an atomic scale, and in particular they should yield a stoichiometric Ga/As atomic ratio [2,3,7-12]. In principle, higher oxidised states of As and Ga are formed under the influence of light at low substrate temperatures (below 1000C). However, oxide composition is dependent on the wavelength applied and the time of exposure. The UV wavelengths of interest are 184.9 nm and 253.7 nm. Such UV-ozone units are comprised essentially of a low pressure Hg lamp UV source. The first wavelength is absorbed by oxygen (O2) producing ozone (O3) which reacts with carbon atoms on the GaAs surface forming CO2. In such a way, carbon contamination is strongly reduced [13]. The latter wavelength (253.7 nm) is adsorbed by O3 species forming O2 and atomic oxygen which leads to oxidation of the GaAs surface [13]. As already mentioned, the oxide composition and hence the oxide desorption temperature are found to be closely linked to the time of ozone exposure [2,7]. It is reported in [7] that the As2O3 concentration follows a logarithmic growth while As2O5 and Ga2O3 formation exhibit a linear growth as a function of exposure time. Additionally, the total oxide thickness has also been found to follow a logarithmic growth characteristic. Studies of UV-ozone exposure have raised extensive controversies as to the composition of the oxide and the oxide/GaAs interface. Some investigators have found that the GaAs surface is damaged by exposure to UV/ozone, and a stoichiometric interface is not formed [2,3,6,13,14], while others claimed to achieve a GaAs oxide that is bulk-like in the Ga/As composition [5,12]. On the other hand, UV oxidation of (100) GaAs was found to produce As-rich phases whose compositions are inconsistent with the equilibrium phase diagram [3,6,13]. Hence, UV/ozone oxidation appears to grow far from thermodynamic equilibrium, suggesting that photochemical effects and kinetic factors primarily control the process [3,7]. However, it is generally accepted that oxides on (100) GaAs after short-term UV/ozone exposure ( 2 - 3 0 min) and subsequent

oxide-film desorption have a Ga/As surface atomic ratio close to unity and a low level of carbon contamination [6,12,15,16]. C

LOW-TEMPERATURE OXIDATION OF GaAs

It is widely accepted that adsorption of unexcited oxygen leads to strong Fermi-level pinning of the GaAs surface. In addition, defect formation is thought to cause a highly disordered surface layer at high oxygen coverage followed by strong pinning of the Fermi level near the midgap. Thick native oxide films, which have been formed by exposure to excited oxygen, have been shown to exhibit a similar composition to that of thermal oxides, indicating that Ga oxide is more abundant, with elemental As occurring at the interface [14,17]. Recently, much interest has emerged in N-O oxidising species, such as N2O, NO and NO2, in order to analyse initial stages of oxide formation and the formation of the first monolayer (ML) of GaAs oxide at low temperatures (25 K - 300 K). Doping the (110) GaAs surface at 300 K with nitrous oxide (N2O) usually gives very much less oxidation (less than 0.02 ML at an exposure up to 105 Langmuir) [18], since the primary dissociation process for physisorbed N2O involves the attachment of low-energy secondary electrons, e.g. generated by X-ray beam (hv ~ 1000 eV [19]). Photoelectron capture is followed by N2O dissociation which results in the formation of Ga2O3 and As2O3-like and intermediate AsOx-like bonding configurations at the surface. Photon-enhanced formation of a thick oxide at a low temperature (25 K) results in As-Ga oxides, which are almost bulk-like in composition, while at 300 K Ga2O3 growth is enhanced at the expense of arsenic oxides [19]. Additionally, photochemical oxidation of GaAs by N2O depends on the surface treatment before insertion into the vacuum chamber [20]. Moreover, it is found [21,22] that adsorption on (110) GaAs of nitric oxide (NO) predominantly reacts to form N2O as well as surface oxidation with oxygen atoms present on the GaAs surface, whereas clean surfaces yield no surface oxide formation. In contrast to O2 and NO, thermal NO2 has a high sticking probability (0.5) on (110) GaAs, which is very much greater than that of other molecules, such as N2O, NO and O2 [18]. At room temperature, a fraction of the adsorbed NO2 finds a special site or configuration and dissociates by simultaneously depositing an oxygen atom that is a highly oxidising agent. The product NO quickly desorbs. The dissociative sticking probability is of the order of 5% for exposures from 0.5 to 10 Langmuir whereas the oxygen coverage rises to approximately 0.2 of an ML. In the case of (110) GaAs, it is found [18,23] that the oxidation occurs by means OfNO2 dissociation at defects, and that the competition between desorption from the terraces and diffusion to defect sites controls the oxidation rate which saturates at approximately 1/3 of an ML coverage at exposures up to 102 Langmuir [18,24,25]. Oxidation data of (100) GaAs surfaces have been shown to be remarkably similar to the oxidation of (110) GaAs surface [18]. Nevertheless, oxidation of (100) GaAs does not appear to be defect-controlled [18]. On n-type Ga-rich (100) GaAs surfaces the oxidation saturates at an oxygen coverage of 1 ML after exposure of about 30 Langmuir by preferentially forming Ga2O [ 18,26]. Most of the chemisorbed oxygen atoms remain on the surface [18] while only a small fraction penetrate the GaAs surface (eventually at defect sites), and produce a donor level (E c - Ed - 0.4

eV) [25]. At higher oxygen exposures, NO2 oxidation results in Ga-O bond formation [27] causing Fermi level pinning in the middle of the GaAs bandgap. Additionally, the efficient oxidation of (110) and (100) GaAs by NO2 sharply contrasts with the inefficient oxidation by molecular oxygen which is known to result in a disrupted GaAs surface layer [18,25]. However, the oxidation by NO2 does not result in a complete overlayer which would be the first step in developing a passivating layer on GaAs. D

ADATOM AND OVERLAYER EFFECTS

Certain adatoms on GaAs surfaces can efficiently enhance the oxidation. Alkali-covered GaAs substrates display particularly strong enhancement in the kinetics of oxidation [28], which are much larger than those observed in electron or laser irradiated surfaces. Both Ga and As atoms are oxidised during room temperature exposure and the alkali-atoms (e.g. K) are also bonded to oxygen. Thermal processing of the sample causes oxygen to be transferred from K to the GaAs crystal [29]. However, the alkali adlayer cannot be completely removed without modifying the stoichiometry of the oxide layer. There is evidence that the alkali-metal induced enhancement of oxidation of GaAs is produced by the formation of alkali oxides [30] and not by the adatoms acting like true catalysts and dissociating oxygen molecules without being oxidised themselves in the reaction [31]. Other types of overlayer also promote oxidation of GaAs, such as Sm [32], Ag, Au, and in particular Cr [33]. E

OXIDATION BY AN ATOMIC OXYGEN BEAM

There are a number of different approaches to realise the growth of GaAs oxides at low temperatures. A non-stoichiometric oxide composition and As segregation at the interface is observed when the oxides are grown near thermal equilibrium conditions. The low temperature processes combine UV-ozone oxidation (see Section B), plasma excitation of oxygen (see Datareview 13.5), ion-beam-induced oxidation (see Section F), and simultaneous O2 and electron beam exposure [34]. Although such techniques can enhance oxide formation, they may produce an appreciable surface-disordered layer resulting in an oxide/GaAs interface that is usually deficient in either As or Ga. Furthermore, the oxidation reactions are often incomplete, in that Ga and As do not reach their highest formal oxidation state. Recent developments [35-37] consist of oxygen sources which provide a constant flux of neutral atomic oxygen with a high kinetic energy of 1 - 5 eV. Nevertheless, the atomic beam is of sufficiently low energy that displacement damage or sputtering effects are probably inhibited. Atomic oxygen sources based on a laser sustained discharge have been applied [36,38] yielding a 2.8 eV oxygen atom beam with a full width at half maximum (FWHM) of -1 eV at a flux of -10 17 atoms/cm2 (i.e. 50 -100 ML/s) [36]. Hoflund et al. [37,39-41] have realised a hyperthermal oxygen atom source using a permeable thin metallic membrane where neutral oxygen atoms (providing an ion-to-neutral ratio of about 2.5 x 10"7 [42]) are emitted towards the target by electron-stimulated desorption. In this case, the average oxygen neutral kinetic energy is about 5 eV, and a flux of 4.5 x 1012 neutrals/cm2sec with 3.6 eV FWHM is reported. These oxygen sources yield an oxygen flux of extremely high purity which is vacuum compatible to epitaxial equipment. In summary, the above-mentioned sources react with high translational

energy neutral atomic oxygen to oxidise the GaAs substrate. In principle, the formation of Ga and As oxides is achieved in their highest oxidation state which implies the formation of either Ga2O3 and As2O5 or GaAsO4. Since the oxidation temperature is very low (below 60 0 C), the oxidation products are expected to be far from thermodynamic. The oxides formed are homogeneous, and composed of As-Ga oxides in equal proportions. Raman spectroscopy [36] shows that there is no excess As (in an amorphous or crystalline phase) at the interface between the GaAs oxide and the substrate. Additionally, no noticeable difference in the oxidation mechanism has been observed for (110) and (100) GaAs surfaces indicating that the oxidation is almost independent of crystal orientation. The oxide layers are thick (200 A) without an appreciable oxidation-induced disordering of the GaAs surface as is usually observed at high oxygen exposures of GaAs (see Datareview 13.2, Section F). However, it remains to be established whether the interface state density is sufficiently low to be competitive with that attainable from low-temperature CVD passivating processes to achieve unpinning of the surface Fermi level. F

ION-BEAM INDUCED OXIDATION OF GaAs

Ion-beam induced oxidation (JBO) comprises the incidence of an oxygen beam of usually O2+ ions onto the GaAs surface with an ion energy of about 0.5-8 keV [43,44]. In principle, oxygen is athermally introduced into the GaAs substrate at low substrate temperatures (less than 1000C) which gives rise to a resulting oxide composition that is far from the thermal equilibrium condition [45]. Furthermore, the total power of the oxygen beam is considered to be below the limit for ion-beam heating, thus keeping the surface temperature low [45]. GaAs bonds are broken upon collision with the O2+ ions. As a consequence, the oxidation of the GaAs surface is enhanced by transferring energy to Ga and As atoms. Changes in the oxidation state of As and Ga as well as in the Ga/As ratio have been analysed by SIMS, AES and small spot-size XPS [43-46]. Differences in the oxide composition have been primarily explained in terms of the angle of incidence and energy of the ion beam [43,45,47-49]. At low ion energies (below 1 keV), a thin oxide film is formed which is composed OfGa2O3 and As2O3, but no elemental As is present at the interface [45]. Moreover, increased formation of arsenic oxides has been found at low ion energy and at normal angle of incidence [43], indicating a difference in the excitation energy supply between Ga and As and, at the same time, an excess supply of oxygen. In principle, a strong energy and angular dependence has been found in the formation of arsenic oxides, whereas the Ga oxidation appears to be very little influenced [43]. As the ion energy increases, preferential sputtering of As and decomposition OfAs2O3 results in an As-deficient surface layer, preventing the formation of an insulating film [45,47]. Hence, it appears that the innuence of ion energy is twofold. First, it was shown that energy is preferentially transferred to arsenic atoms that are being oxidised. Second, the incident oxygen beam also causes displacement of the As atoms which partly get lost yielding a clear angular dependence. A higher ion energy, however, increases the probability of As species being removed by sputtering effects. For the latter case, the composition of the surface layer is altered due to a preferential removal of As and preferential oxidation of Ga. The latter is of course correlated to its corresponding heat of formation [44-49].

The thickness of the modified layer has been reported to be rather large which may be explained in terms of a radiation-induced diffusion of the oxidising agents [47] suggesting that crystalline and electrical damage of the GaAs surface are likely to extend several hundreds of nm into the semiconductor surface [47,50]. However, this contrasts to the results obtained by Rutherford-backscattering spectroscopy (RBS) [45] yielding no noticeable damage after ion bombardment below 1 keV. If any form of Al is introduced into the Ga-As phase system Al2O3 is immediately formed irrespective of all experimental conditions [43,44,47]. Alternatively, electron bombardment may also be applied to initiate GaAs oxidation, leading to radiation-induced interaction between electrons, ions and atoms and also damage-enhanced diffusion of oxygen. In contrast to ion-beam induced oxidation processes, preferential sputtering of arsenic species is prevented [51]. In summary, the steady state composition of the surface after O2+ ion beam induced oxidation of GaAs (with ion energy above 1 keV) usually exhibits a disturbed surface layer that is As deficient [43,44,47]. G

CONCLUSION

The invention of ultra-pure atomic oxygen sources is expected to have great potential for GaAs surface passivation to be effective. Since the reaction products of those processes have been shown to be far from thermal equilibrium, a significant reduction of the density of states at the oxide/GaAs interface appears to be possible. As outlined in Datareviewl3.2, in-situ surface preparation is essential. Atomic oxygen sources provide these features by combining additional atomic sources for atomic hydrogen or nitrogen. However, comprehensive investigations have to be undertaken in order to study in detail the electrical interface properties relevant for device applications. UV/ozone oxidation processes have been widely established for the purpose of pre-epitaxial surface preparation. Nevertheless, a comparison of structure and composition of the native oxide/GaAs interface has led to the conclusion that UV/ozone oxidation involves a rather complex process in controlling oxide and interface properties. Further electrical measurements of MOS structures, for example, must be conducted. In the case of the other oxidation techniques mentioned above, some disadvantages can easily be identified, such as surface contamination related to adatoms remaining on the GaAs surface, changes of surface stoichiometry and charging effects due to electron or ion bombardment. Thus passivating efforts as well as material and device stability may be hampered by the surface and interface degradation involved. REFERENCES [1] [2] [3] [4] [5] [6]

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RA. Outlaw, D. Wu, M.R. Davidson, G.B. Hoflund [ J. Vac. Sci. Technol. A (USA) vol. 10 (1992) p. 1497] MR. Davidson, G.B. Hoflund, RA. Outlaw [ Surf. Sci. (Netherlands) vol.28 (1993) p. 111 ] M.R Davidson, G.B. Hoflund, RA. Outlaw [ J. Vac. Sci. Technol. A (USA) vol. 11 (1993) p.264 ] P.F.A. Meharg, E.A. Ogryzlo, I. Bello, W.M. Lau [ J. Vac. Sci. Technol. A (USA) vol. 10 no.4 pt. 1 (1992) p. 1358-64] J.-L. Alay, W. Vandervorst [ Beam Solid Interactions: Fundamentals and Applications Symp. Eds M. Nastasi, L.R Harriott, N. Herbots, RS. Averback (MRS, Pittsburgh, PA, USA, 1993) p.619-24 ] O. Vancauwenberghe, N. Herbots, H. Manoharan, M. Ahrens [ J. Vac. Sci. Technol. A (USA) vol.9 no.3pt.l (1991) p. 1035-9] JS. Solomon, J.T. Grant [ J. Vac. Sci. Technol. B(USA) vol. 12no. 1 (1994)p. 199-204] J.L. Alay, W. Vandervorst, H. Bender [ J. Appl. Phys. (USA) vol.77 no.7 (1995) p.3010-22 ] N. Herbots, O.C. Hellman, P.A. Cullen, O. Vancauwenberghe [AIP Conf. Proc. (USA) vol.167 (1988)p.259] N. Herbots, O. Hellman, P. Yeh, X. Wang, O. Vancauwenberghe [ in Low energy surface interactions, Ed J.W. Rabalais (Wiley, Cambridge, 1993) chp.VIII ] N. DasGupta, R. Riemenschneider, H.L. Hartnagel [ J. Electrochem. Soc.(USA) vol.140 (1993) p.2038-41 ] T.Wada [ Appl. Phys. Lett. (USA) vol.52 (1988) p. 1056-8 ]

CHAPTER 14 INTERFACES AND CONTACTS 14.1 14.2 14.3 14.4 14.5 14.6 14.7 14.8 14.9

The structure of the epitaxial Al/GaAs interface Structure of the Au/GaAs interface Structure of Au-Ge/GaAs interfaces Structure of silicide/GaAs interfaces Barrier height at the GaAs/Al interface Barrier height at the Ag/GaAs interface Barrier height at the GaAs/Au interface Barrier height at the Pt/GaAs and W/GaAs interfaces Conduction and valence band offsets at the AlGaAs/GaAs and (Al)GaInP/GaAs heterostructure interfaces

14.1 The structure of the epitaxial Al / GaAs interface CJ. Kiely August 1996

A

INTRODUCTION

Epitaxial Al contacts to GaAs form one of the most investigated metal-semiconductor interfaces. The structure of Al contacts formed on both (001) and (110) GaAs substrates has been studied, although the former system is by far the most technologically important. Al / GaAs (001) contacts are extensively used as Schottky diodes having a barrier height typically between 0.75 and 0.80 eV. The microstructure of the Al / GaAs (001) interface is primarily reviewed here, although some limited discussion of the Al / GaAs (110) interface is also presented. Thermodynamically AlxGa1^As alloys exhibit lower Gibbs free energies for all x than their respective two phase (Al + GaAs) counterparts. The interface is consequently reactive toward formation of the alloy, although kinetic barriers hinder the reaction under most experimental formation conditions. The metastable structures which form depend sensitively on the precise details associated with growth and primarily on the substrate temperature during or after Al deposition. B

Al / GaAs (001)

Bl

Initial Stages of Growth

Although Al absorption on GaAs has been the focus of many investigations [1] much uncertainty still exists as to the initial positions of the Al adsorbate. RHEED analysis of sub-monolayer Al coverages on clean reconstructed GaAs (001) substrates [2] have demonstrated that absorbed Al dimers are usually randomly dispersed rather than forming an ordered overlayer. With increasing coverage, Al island nucleation and growth are observed [3] suggesting that the growth mode is Stranski-Krastanow type in character. Island coalescence typically occurs after several hundred Angstroms of growth for low temperature MBE films, but islands can persist beyond 700 A in films grown by CBE [4]. B2

Orientation Relationships

Al has an f e e . structure with a lattice parameter of 4.049 A. GaAs has the cubic zincblende structure (FIGURE l(a)) with a lattice parameter of 5.654 A. The 40% difference in lattice parameters is very large for an epitaxial system. Four epitaxial orientation relationships between the Al overlayer and a GaAs (001) substrate have been reported [3-5]. These are denoted as Al(OOl)45O, Al(OOl)77, Al (011) and Al (011)R. These are considered separately in the following sections in which we summarise the current known information on a) b) c)

lattice mismatch; misfit dislocation structure; and interfacial atomic structure.

Ga at height 0 Ga at height 1/2 As at height 1/4 As at height 3/4 Al at height 0 Al at height 1/2

FIGURE 1. Schematic crystal projections of (a) GaAs (001), (b) Al (001)//f (c) Al (00I)45., (d) Al (011) and (e) Al (OH)R.

B2.1

AI (001)45o / GaAs (001)

The precise epitaxial orientation relationship is [00I]^ 1 //[00I] 01 ^ and [010]*//[11O] 01 * In this orientation the Al unit cell is rotated by 45° about the coincident [001] axis. This produces a structure, shown in FIGURE l(c), which has an isotropic misfit of-1.36% (compressive) in the interfacial plane. This is the best possible fit that can be achieved in this system, and is correspondingly the dominant orientation observed in room temperature MBE grown samples [5]. The misfit dislocation structure of the Al (001)450 / GaAs (001) interface has been thoroughly characterised by Kiely and Cherns [3]. Two distinct misfit dislocation arrays were seen to co-exist. The dominant array seen was a square array of interfacial edge-type defects with 1/4[110] and l/4[lT0] Burgers vectors (as referred to the GaAs unit cell). A mesh periodicity of 146 A for this type of interfacial defect is sufficient to completely relieve the lattice strain. The second, less common, array was found to be a square array of edge-type interfacial dislocations with 1/2[10O] and 1/2[01O] Burgers vectors. In this case a dislocation separation of 208 A is required to fully relieve the -1.36% misfit strain. The two sets of dislocation networks are rotated by 45° relative to each other and their spacing differs by a factor o f / 2 . Dislocations of the type l/2<100> are crystal defects in that the Burgers vector corresponds to a lattice translation vector in both the Al and GaAs lattices. On the other hand, the 1/4<110> defects are imperfect dislocations in that they can only exist in the interfacial plane and do not correspond to lattice translation vectors in either of the constituent 'bulk' materials comprising

the bicrystal. The l/4<110> defects can in fact be considered as partial dislocations generated from the l/2<100> perfect dislocations via the following dissociation process 1/2[10O] -

1/4[11O] + l/4[ll0]

Simple calculations involving dislocation line length and the b2 criterion suggest that the energies of arrays of 1/4<110> and l/2<100> dislocations generating the same misfit relief are in the ratio 1 : / 2 , making the 1/4<110> dislocations energetically more favourable. The critical thickness, t^, at which it becomes energetically favourable to convert from a pseudomorphic epilayer to one containing interfacial misfit dislocations has been estimated for this system [3,6]. The hc values for the 1/4<110> and l/2<100> arrays are 25 A and 40 A respectively. Al layers deposited by MBE onto room temperature substrates usually show atomically sharp interfaces [3,5]. The atomic structure of the Al (001)450 / GaAs (001) interface has been the subject of a number of experimental and theoretical investigations. Two possible models have been proposed [6]. Model A places the Al atoms directly above the atomic sites in the GaAs whilst model B introduces a rigid body shift of l/4<100>GaAs between the Al overlayer and the GaAs substrate. Since the misfit dislocations with 1/4<110> Burgers vectors are not translation vectors of either crystal, a rigid body shift of AR = l/4[100]GaAs is necessary to accommodate this dislocation array [6,7]. In contrast to this, the l/2<100> misfit dislocation array can in principle exist at either a model A or model B interface. Cherns and co-workers have used a variety of techniques to discriminate between models of the interface with and without rigid body shifts, including convergent beam electron diffraction (CBED/LACBED) [8-11], analysis of higher order laue zone (HOLZ) lines [10], atomic location by chemical microanalysis (ALCHEMI) [11] and cross-sectional high resolution electron microscopy (HREM) [12]. With the exception of HREM, all techniques showed a shift of l/4<100>GaAs; HREM appeared to show that both shifted and nonshifted variants could co-exist [12]. B2.2

Al (001)7/ / GaAs (001)

The precise orientation relationship, shown in FIGURE l(b), is; [001L // [001]GaAs and [100]^ // [100]GaAs This orientation relationship has only been reported to occur occasionally in films grown by ionised cluster beam deposition (ICB) [13] and CBE [4]. The relative rarity of articles recording this particular epitaxial orientation is unsurprising as it has an associated +40% isotropic (tensile) strain in the interface plane. HREM observations of the Al (001)// /GaAs (001) structure [7] have shown that a regular misfit dislocation array at the interface is absent. In fact the interface structure is best described as incoherent with seven {100}^ planes roughly matching to five {100}GaAs planes. B2.3

Al (011) and Al (OH)R / GaAs (001)

The precise epitaxial relationship corresponding to the Al (011) orientation, shown in FIGURE l(d), is

Al (Ol 1);

[01 l]M Il [001]GaAs and [100]* // [110]GaAs

In this case the lattice misfit is -1.36% along [110]GaAs and 40% along [TlO]GaAs The nonequivalent Al (011)R variant (FIGURE l(e)) is obtained when the Al is rotated by 45° about the [01 \]M substrate normal. Al (011)R;

[011]* // [001]GaAs and [100]* // [Tl0]GaAs

This structure has a 40% misfit along [110]GaAs and -1.36% along [Tl0]GaAs. The inequivalence of the Al (011) and Al (011)R variants derives from the fact that there is no four fold symmetry axis in the GaAs space group. The strain relief mechanisms at play in the Al (011) and Al (OH)R / GaAs (001) have been investigated by Beanland et al [5] and Bangert et al [14]. A uniaxial, but rather disordered, array of 1/4<110> misfit dislocations is usually seen to relieve the misfit strain along the direction with a -1.36% compressive strain while in the orthogonal direction the +40% mismatch is overcome by forming an incoherent ' 7 * -5 GaAs ' interface. It has also been suggested [7] that the 1/4<110> defects may be associated with interfacial steps of height l/2[001]GaAs in this particular interface. Implicit in this latter proposition is that the requirement for an interfacial rigid body shift of l/4<100>GaAs is lifted when the 1/4<110> defect is constrained at such an interfacial step. B3

Factors Affecting Orientation Distribution of Overlayer

From the standard arguments, the (001)45O orientation should be the only one observed, as a better fitting of lattice parameter is usually associated with a lower interfacial energy. This however is clearly not the case in the Al / GaAs (001) system where the Al (10O)45, Al (011) and Al (011)R orientations are frequently seen to co-exist. Although several groups of workers have observed these orientations [15-20] there is some confusion as to how the various epitaxial arrangements depend (if at all) on the state of reconstruction of the GaAs surface. There is a general consensus that the low misfit Al (001)450 orientation is favoured on the Ga-stabilised (4><6) surface. It is also agreed that the As-rich c(4x4) surface reconstruction tends to simultaneously exhibit Al (001)45o, Al (011) and Al (011)R growth although the relative proportions of each are rather variable. For the As-stabilised c(2><8) surface, some workers [15,20] report only Al (001)450 growth whereas others [16] find Al (011) and Al (011)R orientations only and a third group [17,18] find a mixture of all three types of deposit. Landgren et al [16] have proposed a chemical mechanism for the metal layer orientation in terms of the bonding configuration at the interface. This favours an Al (01 l)/Al(011)R orientation for the c(2x8) surface because of the strong Al-As tetrahedral bond, and the Al(OOl)450 orientation for the c(4x4) surface where the weaker Al-Ga metallic bond is formed. Petroff et al [17] found by TEM and Rutherford backscattering that in general, deposition on Asrich surfaces at room temperature and low deposition rates (200 - 750 A h'1) led to (011)/(011)Roriented Al films. Higher deposition rates at room temperature led to (00I)450 and (011)/(011)Roriented films or completely (00I)45 Al films at the highest deposition rates. More recently Missous et al [21] have investigated the effects of the quality of the vacuum on the Al orientations observed. It was found that leaving the ion gauge on in the MBE chamber led to mixed orientation films. If the ion gauge was turned off and a cryopump employed, the

Al(OOl)45O orientation was reproducibly formed irrespective of the type of GaAs surface reconstruction. For mixed orientation, MBE-grown Al films it has also been reported [17, 20] that the relative proportions of the various orientations depend on the film thickness. The fraction of (011 )/(011)R Al orientation is observed to progressively decrease with increasing film thickness, suggesting that during epitaxial growth re-crystallisation takes place. The thermal instability of the (011) and (011)R orientations, even in thin films when large proportions are present, has been demonstrated [20, 22]. Short rapid thermal anneals can result in the complete disappearance of the (011) and (011)R orientations in favour of Al (001)450. B4

Interfacial Reactions

Other aspects of the Al / GaAs (001) Schottky contact which have received some attention are the interface width and the thermal stability. It is important that these interfaces should remain stable at temperatures up to 5000C which are required subsequently for forming ohmic contacts to GaAs. Several workers [20,23] have found very little interdiffusion in epitaxial Al / GaAs contacts with only minor changes in electrical characteristics after heat treatments up to 500 0 C. Comparisons have been made with polycrystalline Al films deposited at room temperature in a conventional evaporator [21,24]. The polycrystalline films show diffuse interfaces which become even broader when subjected to similar heat treatments, suggesting that the enhanced diffusion that can occur at grain boundaries is highly detrimental. IfAl deposition is carried out at high temperatures, strong interactions with AlxGa1^As compound formation invariably occur [25]. Even Al layers deposited by MBE at room temperature show some evidence for localised formation of interfacial AlxGa1^As [3, 14]. The formation of regions OfAlxGa1^As at the interface has recently been correlated with a deterioration of the Schottky diode characteristics [26].

C

Al/GaAs (110)

The growth mode for Al deposition onto a room temperature GaAs (HO) substrate remains controversial in the submonolayer coverage regime [27]. Huijser et al [28] contend that for submonolayer coverage the Al forms an ordered overlayer. Other workers using LEED [29] and photoemission [30] studies suggest that the initial overlayer is dispersed. For coverages exceeding 0.5 monolayers evidence is seen for metal-like cluster formation [31]. As the coverage regime increases to 1 - 3 monolayers, evidence is seen for the formation of nucleated Al islands rather than layer-by-layer growth [32]. The epitaxial orientation relationship of the Al on GaAs (110) has been reported by Prinz et al [33] to be Al [001] // GaAs [110] and Al [110] // GaAs [001] which represents the lowest misfit epitaxy available to this system. It has also been found by Liliental-Weber [34] that the interface is atomically flat and that the Al(IlT) planes consistently form a 10° angle with the GaAs (111) planes. To date, no structural models of this particular interface have been proposed. Furthermore, studies of the interfacial misfit dislocation structure

of this interface have not yet been attempted. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] II1] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33] [34]

J.S. Bumham et al [ Phys. Rev. B (USA) vol.53 (1996) p.9901 ] S.K. Donner, K.P. Caffey, N. Winograd [ J. Vac. Sci. Technol. B (USA) vol.7 (1989) p.742 ] CJ. Kiely, D.Chems [ Philos. Mag. A (UK) vol.59 (1989) p. 1 ] D. Sun et al [ J. Cryst. Growth (Netherlands) vol. 132 (1993) p.592 ] R. Beanland, CJ. Kiely [ Interface Sd. (Netherlands) vol. 1 (1993) p.99 ] D. Chems, CJ. Kiely [Mater. ScL Eng. A (Switzerland) vol. 113 (1989) p.43 ] R. Beanland, CJ. Kiely, R C Pond [ Handbook of Semiconductors vol.3(a) Ed S. Mahajan (1994) p. 1149] DJ. Eaglesham, CJ. Kiely, D. Chems, M. Missous [ Philos. Mag. A (UK) vol.60 (1989) p. 161 ] M.A. Al-Khafaji, D. Chems, CJ. Roussouw, D.A. Woolf [ Inst. Phys. Conf. Ser. (UK) vol. 117 (1991)p.253] M.A. Al-Khafaji, D. Chems, CJ. Roussouw, D.A. Woolf [ Philos. Mag. A (UK) vol.66 (1992) p.319] M.A. Al-Khafaji, D. Chems, CJ. Roussouw, D.A. Woolf [ Philos. Mag. A (UK) vol.65 (1992) p.385] A.B. Kendrick, J.L. Hutchison, D. Cherns [ Inst. Phys. Conf. Ser. (UK) vol.98 (1989) p.383 ] I. Yamada, CJ. Palmstrom, E. Kennedy, J.W. Mayer, H. Inokawa, T. Tagaki [Mater. Res. Soc. Symp.Proc. (USA) vol.37 (1985) p.401 ] U. Bangert, B. Tang, M. Missous [ J. Cryst. Growth (Netherlands) vol. 154 (1995) p.223 ] A.Y. Cho, P.D. Dernier [ J. Appl. Phys. (USA) vol.49 (1978) p.3328 ] G. Landgren, R. Ludeke [ Solid. State Commun. (USA) vol.37 (1981) p.127 ] P.M. Petroff, L.C Feldman, A.Y. Cho, R.S. Williams [ J. Appl. Phys. (USA) vol.52 (1981) p.7317] J. Massies, N.T. Linh [ Surf Sci. (Netherlands) vol. 114 (1982) p. 147 ] W.I. Wang [ J. Vac. Sci. Technol. B (USA) vol. 1 (1983) p.574 ] M. Missous, E.H. Rhoderick, K.E. Singer [ J. Appl. Phys. (USA) vol.59 (1986) p.3189 ] M. Missous, E.H. Rhoderick, K.E. Singer [ J. Appl. Phys. (USA) vol.60 (1986) p.2439 ] G. Landgren, R. Ludeke, C Serrano [ J. Cryst. Growth (Netherlands) vol.60 (1981) p.393 ] H.B. Kim, G.G. Sweeney, T.M.S. Heng [ Inst. Phys. Conf. Ser. (UK) no.24 (1975) p.307 ] A. Christou, H.M. Day [ J. Appl. Phys. (USA) vol.47 (1976) p.4217 ] R. Ludeke [ Surf. Sci. (Netherlands) vol. 132 (1983) p. 143 ] SJ. Pilkington, M. Missous, U. Bangert [ Inst. Phys. Conf. Ser. (UK) no. 146 (1995) p.549 ] R. Ludeke [ NATO ASISeries B, Phys. (USA) vol. 163 (1987) p.319 ] A. Huijser [ Surf Sci. (Netherlands) vol. 102 (1982) p.264 ] A. Kahn, J. Carelli, D. Kanani, CB. Duke, A. Paton, L. Brillson [J. Vac. Sci. Technol. B (USA) vol.19 (1981) p.331] R.R. Daniels, A.D. Kanani, T.X. Zhao, G. Margaritondo [Phys. Rev. Lett. (USA) vol.49 (1982) p.895] K.L.I Kobayashi, N. Watanabe, H. Nakashima, M. Kubota, H. Daimon, Y. Murata [ Phys. Rev. Lett. (USA) vol.52 (1984) p. 160 ] N.G. Stofifel, M.K. Kelley, G. Margaritondo [Phys. Rev. B (USA) vol.27 (1983) p.6561 ] G.A. Prinz, J.M. Ferrari, M. Goldenberg [ Appl. Phys. Lett. (USA) vol.40 (1982) p. 155 ] Z. Liliental-Weber [ J. Vac. Sci. Technol. B (USA) vol.5 (1987) p. 1007 ]

14.2 Structure of the Au/GaAs interface C.R.M. Grovenor September 1996

A

INTRODUCTION

A complete structural model of this important interface must be able to explain the observation that the Au/GaAs Schottky barrier height decreases sharply on heating above about 4000C3 resulting in the formation of an ohmic contact. This is presumably due to some kind of metallurgical reaction at the interface. Experimental work on the Au/GaAs interface falls into two classes: detailed studies of mixing phenomena usually at low temperatures, and morphological and microstructural studies of heavily reacted interfaces to determine the nature of the intermetallic phases. The combined use of Rutherford Backscattering, X-ray diffraction and transmission electron microscopy to study these contacts is now revealing their chemistry and phase morphology in very great detail. B

OBSERVATIONS AT LOW TEMPERATURES

There is extensive experimental evidence that some interdiffusion occurs at the as-deposited Au/GaAs interface [I]. These observations have been made with XPS and UPS techniques [2,6,43], PES [7-11], AES [5,6,10,11], cathodoluminescence spectroscopy [12], and ion channelling [14]. Many of these studies have been on GaAs (110) surfaces because of the convenience of being able to cleave GaAs crystals on this plane in UHV systems [2-4,7,8,1012,43], but more recently work has begun on the technologically important (001) GaAs surface with samples grown by MBE [5,6,9]. Qualitatively the observations on these two surfaces are the same, with the gold dissociating the GaAs and reacting with Ga while the excess As escapes to the top surface or into the gas phase. At monolayer coverages of Au at room temperature, XPS studies indicate that the Fermi level position is already pinned at the same position as found in the Schottky contact between a thick Au layer and a GaAs crystal [4]. Stiles et al [15,40] have shown that the pinning of the Fermi level occurs at greater thickness of Au when the GaAs is cooled. This suggests that a thermally driven dissociation reaction at the Au/GaAs interface is necessary to pin the Fermi level. Other authors suggest that changes in the interfacial electronic state densities occur at multilayer coverages [12], so that minor modifications of the Schottky barrier height can occur as the Au layer thickens [16,17]. This two stage process of contact formation may be related to the initial formation of an intermixed Au/Ga/As layer, before nucleation and growth of fully metallic Au (or Au-Ga alloy) islands [6,8-10,40]. Care has to be taken, however, in making comparisons between Schottky barrier heights measured at Au/GaAs interfaces prepared under different conditions. A recent collation of data [56] shows that the barrier height can vary by more than 1 eV depending on the deposition conditions. Studies of Au/GaAs interfaces by ballistic electron emission microscopy have revealed nanoscale variations in the contact characteristics on a length scale much smaller than the morphological nature of the interface [55]. This kind of observation can be used very effectively to explain apparent discrepancies between macroscopic contact properties from conventional I/V and C/V measurements.

C

EFFECTS OF ANNEALING

Annealing thin Au layers on GaAs surfaces encourages the formation of thicker Au-Ga alloy islands and further evaporation of As [4,8]. Newman et al [7] have linked the formation of these Au-Ga alloys with the decrease in Schottky barrier height. Kulkarni and Lai [41] have estimated an average grain boundary diffusion coefficient of Ga in gold contacts of 3.5 x 10"9 exp(-0.21/kT) cm2 /s while Gupta et al [46] estimate a much higher activation energy for the grain boundary diffusion process of 0.9 eV/atom. This is a serious discrepancy, and calls into question the whole process of estimating diffusion coefficients from Auger depth profiles in systems where intermetallic phases are known to be formed. Weizer and Fatemi [42] have shown that the rate of dissolution of GaAs in Au contacts depends on the vacancy flux passing through the Au, (presumably a grain boundary vacancy flux). Macroscopic changes in the morphology of the contacts also occur on heating. Numerous groups have reported observation of rectangular etch pits in the GaAs surface after annealing above 3000C, see for instance [18-23], although the density of etch pits can be markedly reduced either by removing the oxide from the GaAs surface before Au deposition [23,28] or by capping the Au with a layer OfSiO2, Al2O3 or Si3N4 [24,25]. The capping layers reduce the rate of As escape from the contact, which can be very rapid at temperatures above 400 0 C [27]. Piotrowska et al [49] have shown quantitatively that the rate of As evolution when Au(Zn) contacts are annealed on (100) GaAs can be reduced by a factor of more than 10 if a 200 nm capping layer of SiO2 is applied. This work has also demonstrated that the extent of dissociation of the GaAs surface required to form contacts with low specific resistances is extremely small. The presence of an interfacial oxide layer presumably encourages localised reactions between the Au and GaAs, and removal of the oxide leads to a more homogeneous reaction. (Oxide and hydrocarbon contamination on the GaAs surface before deposition of the gold contact also reduces the interfacial adhesion. Chemical [47] and ion beam cleaning [48] processes can be used to remove these surface contaminants, significantly improving the adhesion of the metallisation.) Confusingly, Lu et al [26] have suggested that the presence of oxygen encourages a more homogeneous interfacial reaction and intermixing process. The discrepancies between these TEM and XPS results remain to be explained. However, there does not appear to be any significant difference in the Schottky properties of these two interface morphologies [23]. Williams and his group [29-32] have considered the thermodynamics of the Au/GaAs interface, including the entropy changes when As is lost to the vapour phase during annealing. They conclude that these entropy contributions dominate the reactivity of this interface, with the temperature at which reactions to form Au-Ga phases and release As occur decreasing with the background pressure [29]. As a result, experiments carried out in UHV will identify dissociation of the GaAs substance as much as 1500C below the critical temperature for reaction in poorer vacua. These authors also suggest that while Au7Ga2 is often the first intermetallic phase observed at the interface during annealing, the thermodynamically stable phase when the GaAs is present in excess is AuGa2 [30]. This phase does indeed seem to be stable on a GaAs surface at temperatures up to 5000C [32]. There is some disagreement in the literature as to the precise sequence of Au-Ga phase formation when Au/GaAs contacts are annealed. The first phase to form has been variously identified as Au7Ga2 [18,30], Au2Ga [20] and AuGa [21] in transmission electron microscopy and X-ray

diffraction experiments. An in-situ X-ray diffraction study [33,34] of Au/GaAs contacts during annealing showed the formation, and subsequent melting at 5250 C, of an Au(Ga) a solid solution, and on cooling a peritectic reaction to form Au7Ga2 at 415 0 C. They have also shown that reducing the chamber pressure during annealing encourages the loss of arsenic and lowers the melting point of the a phase by as much as 1000C in agreement with the thermodynamic predictions of Williams [30]. More recently, Barcz [52] has attempted a kinetic model of the dissolution rate at the Au/GaAs interface, which contains some interesting predictions of the development of the stoichiometry of the GaAs surface: As-rich at short annealing times switching to Ga-rich at longer times. In addition, Kaminska et al suggest that the reaction sequence between Au and GaAs is altered when Zn is included as a p-type dopant [54]. A stable Au-Zn phase replaces the Au-Ga phases normally found above 400 0 C. A simple thermodynamic model has also been used to explain why addition of Ga to the Au contact stabilises the Schottky barrier properties [44,45]. It is assumed that the driving force to form Au-Ga alloys by reaction of the metallisation with GaAs is much reduced, and interdiffusion limited at the metal/semiconductor interface. A similar reduction in the rate of dissolution of Ga into the Au contacts can be achieved by alloying the Au with Ag [57]. Microstructural observations have also helped to clarify the observations of Newman et al [7] that on annealing an ohmic contact is formed around the periphery of a mesa contact structure. It has been shown that the morphology of gold grains at the outer rim of a contact is quite different to that in the centre [35]. We might expect the loss of arsenic to be faster in these regions, and so the rate of formation of Au-Ga phases increased leading to the formation of a local ohmic contact. D

EFFECT OF ION BOMBARDMENT

The effect of ion bombardment has been studied, both of the GaAs surface before evaporation of the gold, and of the complete Au/GaAs contact. Holloway et al [36,37] used argon ions to bombard GaAs surfaces, modifying the surface composition and altering the Au/GaAs Schottky barrier height. They found that it was difficult to create ohmic contacts after ion bombardment. By contrast Jaroli et al [38] found that implanting Au/GaAs contacts with xenon ions encouraged intermixing at the interface, and gave contacts with good ohmic character after annealing above 450 0 C. Interfacial mixing and evaporation of arsenic from the sample surface has also been observed after argon ion implantation [39]. E

Au-Te CONTACTS

A considerable number of papers have been produced recently on Au-Te ohmic contacts to nGaAs. Some excellent properties have been reported after annealing above 500 0 C, but whether the mechanism of ohmic contact formation is heavy Te doping of the underlying GaAs [51] or the formation of a quasi epitaxial Ge2Te3/GaAs heterojunction [50,53] is hotly debated. The observation by cross sectional TEM that no Ge2Te3 can be found in contacts with low specific resistances does perhaps lend weight to the first mechanism [51]. REFERENCES [1] [2]

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Y.-X.Wang, P.H.Holloway [ J. Vac. Sci. Technol. B (USA) vol.2 no.4 (1984) p.613-19 ] P.H. Holloway, Y.X. Wang, X.-J. Xie, T.D. Bussing [Mater. Res. Soc. Symp. Proc. (USA) vol.47 (1985) p.235-42 ] E.Jarolietal[M/c/. Instrum. MethodsPhys. Res. B (Netherlands) vol. 19/20 pt.II (1987) p. 76772 ]; E. Jaroli et al [ Phys. Status Solidi A (East Germany) vol. 107 no.l (1988) p.K15-18 ] AJ. Barcz, M. Domanski, J. Jagielski, E. Kaminska [ Nucl. Instrum. Methods Phys. Res. B (Netherlands) vol. 19/20 pt.II (1987) p.773-6 ] K. Stiles, A. Kahn, D. Kilday, G. Margaritondo [J Vac. Sci. Technol. A (USA) vol.6 no.3 pt.II (1988)p.l511-14] A.K. Kulkarni, C. Lai [ Thin Solid Films (Switzerland) vol.164 (1988) p.435-9 ] VG. Weizer,N.S. Fatemi [ J. Appl. Phys. (USA) vol.64no.9 (1988)p.4618-24 ] D.M. Hill, F. Xu, Z. Lin, J.H. Weaver [ Phys. Rev. B (USA) vol.38 no.3 (1988) p. 1893-1900 ] S. Guha, B.M. Arora, V.P. Salvi [ Solid-State Electron. (UK) vol.20 no.5 (1977) p.431-2 ] RP. Gupta, J. Wuerfl, H.L. Hartnagel, W.S. Khokle [ IEEProc. I (UK) vol.35 no.2 (1988) p.25-8] RP. Gupta, W.S. Khokle, J. Wuerfl, H.L. Hartnagel [ Thin Solid Films (Switzerland) vol.151 no.3 (1987)p.L121-5] T.T. Bardin, J.G. Pronko, D.A. Kozak [ Appl. Phys. Lett. (USA) vol.54 no.2 (1989) p. 173-5 ] CR. Wie, J.Y. Tang, T.A. Tombrello, R.W. Grant, R.M. Housley [ Vacuum (UK) vol.38 no.3 (1988) p. 157-60] A. Piotrowska, E. Kaminska, S. Kwiatkowski, A.Turos [ J. Appl. Phys. (USA) vol.73 no.9 (1993) p.4404-7 ] J. Watte, K. Wuyts, RE. Silverans, M.Van Hove, M.Van Rossum [ J. Appl. Phys (USA) vol.75 no.4 (1994) p.2055-60 ] X.W. Lin, A. Piotrowska, E. Kaminska, Z. Liliental-Weber, J. Washburn, E.Weber [ Appl. Phys. Lett. (USA) vol.62 no.23 (1993) p.2995-7 ] AJ. Barcz [ J. Appl. Phys. (USA) vol.74 no.5 (1993) p.3172-6 ] H. Munder et al [J. Appl. Phys. (USA) vol.71 (1992) p.739 ] E. Kaminska, A. Piotrowska, E. Mizera, E. Dynowska [ Thin Solid Films (Switzerland) vol.246 (1994) p. 143-50] A.A.Talin, RS.Williams, B.A. Morgan, K.M. Ring, KX. Kavanagh [ Phys. Rev. B (USA) vol.49 no.23 (1994) p. 16474-79] A. Piotrowska, E. Kaminska [Thin Solid Films (Switzerland) vol. 193/194 (1990) p.511-27 ] S.D. Mukherjee, P. Zwicknagl, H. Lee, L. Rathbun, L.F. Eastman [ Solid State Electron. (UK) vol.29 (1986) p. 181]

14.3 Structure of Au-Ge/GaAs interfaces C.R.M. Grovenor September 1996

A

INTRODUCTION

Several review articles on practical contacts to IDL-V semiconductors have been published recently [50-52], all of them containing extensive discussion of the advantages and disadvantages of AuGe based ohmic contacts when compared to more recently developed metallisations. The eutectic Au-Ge alloy containing 12 at.% Ge is widely used to form ohmic contacts to n-GaAs [1,2]. However, understanding of this metallisation system is far from complete, although one widely held model of ohmic contact formation is thought to involve Ge incorporation onto Ga sites leading to an n+ surface region. The contact must be heated to encourage the outdiffusion of Ga and indiffiision of Ge, and this commonly results in extensive reaction of the Au with the GaAs and the formation of a range of Au-Ga intermetallic compounds. These reactions occur highly nonuniformly, probably as a result of the inhomogeneous nature of the residual oxide layers on the GaAs surfaces, and can lead to the formation of deep pits etched into the GaAs. This pitting under the Au-Ge contact renders it unsuitable for many submicron device structures. In addition, if the contact alloy is melted, 'balling up' occurs. Ni, Pt, Pd and In have all been added to the Au-Ge metallisation to improve the wetting of the GaAs by the contact melt, but also react with the GaAs to form a number of new intermetallic phases with the suppression of some phases found in annealed Au-Ge contacts. Analysis of the metallurgy of annealed Au-Ge alloy contacts to n-GaAs has been attempted by a large number of workers. Kulkarni and Lai [46] have measured diffusion rates in Au-Ge contacts on GaAs, showing that Ga diffuses much faster along grain boundaries in the Au-Ge contact than in pure Au, and that Ge diffuses more rapidly into the GaAs substrate than Au. The accuracy of these results, which are obtained by averaging chemical information over large areas, may be questionable because studies of the Au/Au-Ge/GaAs system [3] show that for sintering temperatures above 250 0 C inhomogeneous diffusion of Ge into the GaAs occurs. B

EFFECT OF ANNEALING TEMPERATURE

For sintering temperatures below the Au-Ge eutectic point (365 0 C) there is no major interaction between Au and GaAs. In situ SEM observations have identified an Au/Ge/As phase forming very rapidly at about 3650 C and forming deep pits into the GaAs surface [4]. A detailed analysis of the structure of this metastable ternary phase has been presented by Rackham and Steeds [5], and its presence confirmed by several groups [6-9]. Annealing at temperatures above 400 0 C leads to the gradual dissolution of this phase, and a significant decrease in the contact resistivity. It is tempting to link the release of Ge from the dissolving ternary phase with the formation of high quality ohmic contacts. Nebauer and Trapp have shown that there is a high concentration of Ge under AuGe contacts annealed at 420 0 C [10] and have estimated an indiffiision coefficient of 1014 cm2/ s at this temperature. SIMS analysis has also shown indiffiision of Ge at 430 0 C. Sintering at 450 0 C leads

to complete mixing of all the elements, the surface showing three distinct areas - agglomerated areas consisting of Au with some trace of Ga5 a matrix area consisting of GaAs with some small traces of Au, and rectangular particles containing about 72-85 at.% Au, 13 at.% Ga and 2 - 1 5 a t % As [3]. More detailed TEM studies have shown that Au7Ga2, Au7Ga, Au3Ga and free Ge precipitates can all be formed, usually with some free gold grains [6,7], and in some cases As2Ge and Au2Ge precipitates as well [9]. The presence of Ge particles has been confirmed by XPS experiments [H]. TEM observations show that these Ge precipitates are arranged epitaxially with the GaAs substrate [6,8,12,13]. Annealing above 5000C generally leads to extensive roughening of the contact due to excessive reaction between gold and GaAs, and degradation of the contact resistivity. C

MECHANISMS OF OHMIC CONTACT FORMATION

The model in which ohmic contacts are formed by the creation of a Ge doped n+ GaAs layer has been challenged by the observations of Kirillov and Chung [14] who present Raman scattering data indicating the presence of the free Ge particles but no Ge doped GaAs layer. Waldrop and Grant [11] and Procop et al [15] suggest that the Ge/GaAs interface has a low barrier height and forms the ohmic contact at temperatures above 400 0 C. However, Iliadis [16] has shown that a contact with a small barrier height is formed at temperatures as low as 300 0 C, well below that at which free Ge is created by the decomposition of the Au/Ge/As ternary phase. He suggests that it is the indiflusion of Au and Ge to form a disordered region which creates the ohmic contact. In addition, Wilier and Oppolzer [12] find that contacts both with and without epitaxial Ge particles can have low contact resistivities. It is thus not possible to conclude that the presence of epitaxial Ge particles is sufficient to create an ohmic contact, although they may act as sources for Ge diffusing into and doping the GaAs. Kulkarni and Lukowski [17] have proposed that the density of Ga vacancies and indiffiising Ge dopant atoms roughly balance after annealing at 4500C, but a resistive layer is formed by the presence of excess Ga vacancies in contacts fired at 500 0 C. There is currently no evidence to suggest that any of the other phases present after annealing binary Au-Ge contacts have any but deleterious effects on the electrical properties of the contacts. D

TERNARY Au-Ge CONTACTS

The most widely studied ternary contact metallisation to n-GaAs is Ni-Au-Ge [42] and the phase chemistry in annealed contacts has been extensively investigated by TEM [6,13,18-22]. The primary role of the Ni (and similarly Pt and Pd) is to react with any native oxide on the GaAs surface and improve the uniformity of the subsequent contact/semiconductor interactions [53]. There is general agreement that below 400 0 C Ni rapidly reacts with Ge in the contact to form a Ni-germanide phase, and at a later stage with As released by the decomposition of the GaAs to form NiAs near the GaAs surface. Hill et al [41] have used EBIC and Auger spectroscopy to show that the formation of Ni-As compounds results in local degradation of Schottky barrier properties. Au7Ga2 is also formed below 400 0 C [18,19]. During the middle stage of alloying (390 - 410 0 C) Ge diffuses out of the NiGe phase into the NiAs phase to form aNi2GeAs ternary phase [13,18,22,40], although there is some disagreement about the precise composition of this phase [54]. On alloying at up to 450 0 C, Au, NiGe and

Ni2GeAs grains larger than 200 nm across are observed. It has been suggested that the presence of the Ni2GeAs grains leads to contacts with the lowest resistance [22]. At the same time, Au7Ga and Au2Ga grains are observed [18]. High resolution TEM and X-ray analysis of AuGeNi contacts to GaAs have shown that the interface is locally very rough after annealing [20,21] and that the extent of this roughness is dependent on the evaporation sequence of the layered contact [20]. The Ge content in the contact metallisation has been shown to alter both the phases present after annealing and the contact morphology (on GaInAs) [49]. In particular the Ni2GeAs phase is not found in high Ge contacts, but the contact resistance is still low. The overall thickness of Au/Ge/Ni contacts and their precise composition have been shown to have a strong influence on the contact resistivity. After annealing at 4700C a minimum resistivity was achieved with 750 A of contact, and this has been correlated with a maximum interface density of a Ni+Ge rich phase which is presumably the Ni2GeAs ternary phase [23]. Patrick et al [24] have shown that with an optimum Ni content of 7% in the metallisation, low resistance contacts can be formed at temperatures as low as 3200C. More recent experiments on the contact composition indicate that the lowest contact resistances are achieved with Ni:Ge thickness ratios between 0.7 and 1 depending on the order in which the metal layers are deposited [55,56]. The importance of the layer sequence in the contacts has been further investigated with specific reference to the reproducibility of contact properties [54]. Placing a thin layer of Ni between the GaAs and the Au-Ge alloy improved the uniformity of the contact by altering the precise sequence of the metallurgical reactions, encouraging the formation of a NixGaAs phase on the GaAs surface. A comparative study on ageing of alloyed Au-Ge, Ni/(Au-Ge) and Au/Ni(Au-Ge) ohmic contacts to epitaxial n-GaAs has confirmed the importance of Ni in providing good ohmic contacts [25]. SEM observations showed the presence of nodular or cluster microstructures under both Au/Ni/(Au-Ge) and Ni/(Au-Ge) metallisations, but there are fewer of these inhomogeneities in the latter case. Capping the NiAuGe layer with 200 nm of CVD SiO2 suppresses As loss, and improves the smoothness of the contact [26]. Kim and Chung [18] have carried out a detailed study of the melting behaviour of Au-Ge and NiAu-Ge contacts, showing that the presence of Ni increases the temperature at which the first melt is formed in the contact. They have also shown that the melts formed do not have simple binary compositions, but are often ternary or quaternary mixtures containing significant levels of Ga, This implies that any melting will result in dissolution of the GaAs surface, creating a rough contact morphology. Any increase in melting temperature thus results in an improvement of the homogeneity of the contact since 'balling up' and rapid dissolution of the GaAs will be suppressed. Pt/Au-Ge contacts with excellent electrical properties have been prepared, and show very similar microstructures to Ni/Au-Ge contacts after annealing with the formation of a PtGeAs ternary phase supposed to be important in the development of low resistivity contacts [48]. Similarly, Pd/Au-Ge contacts with low resistances have been demonstrated [58]. Lin and Lee [59] have compared the performance of Pd, Pt and Ni-containing Au-Ge contacts, and have concluded that Pt gives the most uniform contact metallurgy with the most stable contact properties. It is worth noting that Au-Mn binary contacts to p-GaAs have recently been shown to have very low

resistances [60]. A recent development of the ternary contact system has been to dramatically reduce the amount of Au to limit the depth of the reaction zone with the GaAs. NiGe contacts containing very little Au [47] (or none at all) have been shown to have highly competitive electrical properties, and to offer significant improvements in the contact morphologies. The reaction chemistry of these contacts is of course completely different to that discussed above, relying on high melting point NiGe phases for the integrity at elevated temperatures.

E

NOVEL ALLOYmG TREATMENTS

While thermal annealing is the most common method of forming ohmic Au-Ge contacts to nGaAs, very brief heating treatments with laser annealing, lamp annealing and electron beam bombardment have all been used to prepare contacts with excellent properties. Several groups have shown that laser annealing contacts at energy densities around 0.5 J/cm2, and with pulse lengths around 30 ns, can give contact resistivities as low as 10"6 ohm cm2 [27-31,42]. In all cases, the contacts remain extremely smooth, showing that the limited heating period reduces the extent of reaction between Au and GaAs. Use of higher energy densities generally results in increased contact resistivities because of over-reaction and evaporation of the contact metals [28,29,42]. Kabanov et al [32] have shown that annealing with xenon discharge lamps (15-20 J/cm2 and 600 \is pulse width) gives good ohmic contacts. A scanning electron beam can also be used to anneal Au-Ge contacts for short times to give good ohmic contacts with very little reaction of Au with GaAs [33,57]. F

EFFECTS OF ION IMPLANTATION

Several groups have investigated the use of ion implantation to achieve more uniform reaction at Au-Ge/GaAs interfaces [34-37,43,44]. The general assumption is that breaking up the residual oxide layer on the GaAs surface will encourage homogeneous reaction between metal and GaAs, and that enhanced diffusion in the implanted contacts leads to a very uniform distribution of contacting elements [43,44]. After annealing the contacts are very flat and uniform with little evidence for 'balling up'. However it has been found to be difficult to produce ohmic contacts even after annealing at higher than usual temperatures. TEM analysis shows that significant lattice damage is introduced into the GaAs by the ion implantation process, and results in the high contact resistivity [35,37]. Jie and Thompson [45] have shown that below a critical ion dose of about 1015/cm2 (Se+ or Kr+) contact resistance can be slightly reduced, but this improvement is not maintained at higher doses. Callegari et al [38,39] have shown that careful sputter cleaning of the GaAs surface before evaporation of the contact metals results in an improvement in the thermal reliability and uniformity of Au-Ge contacts presumably by removing the oxide layers without introducing lattice damage. REFERENCES [1] [2] [3] [4]

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14.4 Structure of silicide/GaAs interfaces

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J. Osvald July 1995

A

INTRODUCTION

The primary aim of silicide/GaAs interface investigations is to find suitable high-temperature stable compounds for Schottky metallization in self-aligned GaAs MESFET technologies [I]. In this technology the gate serves also as a mask for source and drain implantation which means that the gate has to be deposited before the implantation and the high temperature post-implantation annealing process. In most single metals interdiffusion between GaAs and metal occurs during annealing which has a deleterious effect on the electrical parameters of Schottky diodes. From this point of view the most convenient silicides are refractory metal silicides (WSix, TaSix, MoSix, etc.). For use in GaAs technology they are prepared mainly by magnetron sputtering (DC or RF) from composite targets or from two single element targets in a multitarget system if variation of the stoichiometry of the films is needed. The silicides have an amorphous structure after deposition, with the exception of low Si/metal ratio ones [2]. In this form the silicide has good properties as a diffusion barrier. Other requirements for silicides are good adhesion properties to GaAs, low stresses in films [3] and good stopping power for ion implantation. The techniques most often used for characterizing the structure of silicide/GaAs interfaces are Xray diffraction (XRD), standard and cross-section transmission electron microscopy (TEM and XTEM), Rutherford backscattering spectroscopy (RBS), Auger electron spectroscopy (AES) and particle induced X-ray emission analysis (PIXE) as a complementary technique to RBS. This results from the poor resolution of RBS for neighbouring elements such as Ga and As. Thermoreflectance measurements have also been shown to be a sensitive tool for detecting the interface degradation during the high-temperature annealing [4]. B

WSi/GaAs INTERFACE

WSix is the most studied gate material for self-aligned MESFET technology and was first described by Yokoyama et al [5,6]. The properties of the WSixZGaAs interface depend strongly on the stoichiometry of the film. The concentration ratios which were investigated ranged from x = 0.11 [7,8] up to x = 2.3 [9]. The relatively stable compounds are those with a composition close to the stoichiometry of the crystallographic phases W5Si3 and WSi2, the former of which seems to be even better [10]. Thermodynamical considerations also anticipate improved stability OfW5Si3 because of positive heat of reaction [H]. For films with near stoichiometric composition (x = 0.6) high temperature annealing induces crystallization of the W5Si3 phase. When x differs from this value, W (or Si) atoms which are not bound in the crystallographic phase may diffuse into the GaAs. Diffusion of W and Si from nonstoichiometric films may occur through those microcrystalline silicide areas with poor crystallinity [12]. W-Ga and W-As compounds were not detected at the interface by X-ray diffraction measurements [15], as they are not stable at high temperatures. For x = 0.52, W3Si is the dominant phase crystallized in the P-W structure; the a-W structure is present only as a minor component. For higher Si content the phases were a-W, W5Si3 and WSi2 [16].

TEM showed marked differences in grain sizes in the annealed films. W-rich films had an average grain diameter of 450 nm, about ten times the average grain size of the Si-rich films [15]. TEM of cross-sections showed that the WSixZGaAs interface is no longer sharp after rapid thermal processing (RTP). According to the roughness at the interface it seems that the silicide grains penetrate into the GaAs substrate to a depth of approximately 10 nm. The interface of the Si-rich films remains abrupt and smooth after annealing. Contacts with an intermediate silicon content, W0 79Si021, W0 67Si033 and W0 57Si043 prepared by sputtering on GaAs(OOl) have been studied by Basile et al [17]. The first two compositions are characterized by an amorphous structure after annealing up to 7500C, but revert to an equilibrium structure (a-W and W5Si3 on GaAs) during annealing at higher temperatures. Corresponding electrical characteristics degrade during annealing at 7500C and above, although to a lesser extent than for other WSixZGaAs contacts. This relative stability is attributed to the amorphous structure, which precludes short-circuit diffusion along grain boundaries. Contacts consisting of W0 57Si043/GaAs are characterized by an amorphous structure after annealing up to 500 0 C, but revert to an equilibrium microstructure (W5Si3 and WSi2 on GaAs) during annealing at higher temperatures. Corresponding electrical characteristics degrade substantially during annealing at 5000C and above. Films with very low Si content have been studied [7,8]. Callegari et al have studied the electrical characteristics OfWSixZGaAs Schottky diodes with x ranging from 0.11 to 0.6. The barrier height rose and the ideality coefficient decreased as the concentration of W atoms increased [7]. The same trend in barrier height was registered by Eicher and Bruce [16]. The ideality factor remained close to unity. A higher W content lowers the resistivity of the film [8]. Ohmic behaviour or poor diode characteristics were observed for more Si-rich films. The problem of stresses is associated with stoichiometry [3,18]. It appears that the optimum Si content for Schottky diode characteristics coincides with that for minimized stress in WSix films [19]. The stress in sputtered films depends on a number of factors such as the pressure of sputtering gas, temperature of the substrate during the deposition process, etc. Many authors refer to tensile stresses in WSix films [7,8,20], The stresses decrease after annealing at 800 0 C for Wrich films and change to low compressive ones for Si-rich films [2]. Somewhat different results were published by Osvald and Dobroeka [3] for Si-rich films with x = 2. The stresses in the films were compressive after the deposition and stayed compressive after annealing. C

TaSVGaAs INTERFACE

The second most studied silicide film on GaAs is TaSix [21,22]. The stoichiometry of the film is again the most important parameter influencing the stability of the interface. Films are generally deposited by magnetron sputtering or vacuum evaporation. Tantalum silicides with x = 2.4 [23], 2 [24], 1.4 [10,25], 1.25 [23], 0.6 [26], 0.33 and 0.11 [27] have been investigated. Only two intermetallic compounds - Ta5Si3 and TaSi2, were detected by XRD, but they were found only in films with strictly corresponding compositions. No mixture of the phases was identified in the films [27]. RBS with PIXE spectroscopy [23] and with AES [27] and SIMS [23] were used for studying the TaSixZGaAs interface before and after annealing. All the techniques confirmed the presence of Ga

atoms at the surface after annealing at 7000C and higher temperatures. The dependence of Ga out-diflusion on composition has been found to be similar to that for WSix and is probably caused by the different grain sizes depending on x. Smaller grains provide more paths for out-diffusion. For silicides with x < 0.33 a reaction between Ta and GaAs after annealing was observed [27]. Rapidly annealed specimens showed no indication of interdiffiision in RBS measurements up to 975 0 C [24]. TaSi2 films crystallize with the hexagonal structure during RTP at 55O°C or below for 10 s, while Ta5Si3 films remain amorphous up to 7000C, crystallizing at 7500C in the hexagonal phase. The evaporation rates of Ga and As are higher for films undergoing crystallization and the different stability may be the consequence of the difference in the extent of crystallization [26]. Ta5Si3 layers withstand higher temperatures without damage of Schottky diode characteristics [26]. Ta5Si3 contacts remained rectifying up to at least 9000C (10 s RTP), whereas TaSi2 contacts were nonrectifying following 10 s RTP at 8000C. Morgan et al [25] found that the out-diffiision of Ga and As may be suppressed when In-doped GaAs is used. A possible explanation of this effect may be in considering the interdiffixsion between the thin metallic films and the GaAs as a crystalline defect assisted process. The low defect density in In doped GaAs then results in better interface stability. D

MoSix/GaAs INTERFACE

Stability OfMoSi2 on GaAs was studied by Truman and Holloway [28]. RBS spectra measured after 600° C and 8000C indicated out-diffusion of Ga and As into the MoSi2 and AES showed broadening of the MoSi2ZGaAs interface. The lowest temperature at which Ga diffused into MoSi2 was found by ion microprobe analysis (IMA) to be 5500C. After heating to 600 0 C, AES detected both Ga and As in the MoSi2 surface. Arsenic was found at a level of about 2 at.% at a depth of 350 nm. After 8000C annealing Ga and As atoms were found at the surface of the MoSi2. The data on Mo diffusion are not unambiguous. AES showed some diffusion of Mo and IMA profiles indicated relatively small changes after annealing at 8000C but RBS results did not confirm the results. However, the Si profiles showed apparent diffusion of Si into GaAs. MoSi2 layers seem to be unsatisfactory for self-aligned MESFET technology but MoSi2 is a stable contact material at temperatures up to 5000C. The electrical degradation above this temperature is caused by Si in-diffusion into GaAs. Chuang et al [29] have studied MoSix layers on GaAs(IOO) with compositions of x = 0.3, 0.6, 1 and 2. The MoSi03 film crystallizes into Mo3Si after annealing at 5700C for 30 min. For MoSi06, MoSi and MoSi2, no peaks were detected in the XRD measurements, indicating that these films remained amorphous. After annealing MoSi06 at 8500C the Mo5Si3 phase in tetragonal form was observed. For MoSi, the phases Mo3Siand Mo5Si3 appear. MoSi2 remains amorphous. Schottky diodes with MoSi03, MoSi, and MoSi2 films degraded only after annealing at 750 0 C. For MoSi0 6, good diode characteristics were retained after annealing at temperatures up to 8500C. E

OTHER SILICIDE/GaAs INTERFACES

Vanadium silicide, VSix, with x = 1.9 was studied by RBS [10,24]. Very little out-diffusion of Ga

and As was observed after RTP at temperatures of 9500C and 10500C. Ternary compounds TiWSix with different TiW/Si ratios have been studied (Ti03W07) [30,31]. The films analysed in [29] were deposited by a multilayer technique which enabled the stoichiometry to be changed. Films with x ranging from 0.2 to 0.83 have been investigated by electrical methods after annealing at 8000C and 8500C for 20 min in an AsH3 overpressure. Ti03W07Si033 showed reduced gallium out-diffusion and the best barrier properties after 800 0 C and 8500C annealing. SIMS measurements revealed considerable redistribution of Ti, W and Si atoms throughout the layer except for the Ti 03 W 07 Si 02 composition, for which analysis still showed evidence of discrete multilayers. Strong evidence for reduced gallium out-diffusion was also registered for this composition. From the results of [32] it seems to be promising to use a Co/Si/GaAs multilayer film structure to make a CoSi2/GaAs contact. The interfacial reaction between Co and Si is faster than that between Co and GaAs when annealed up to 600 0 C. F

CONCLUSION

The available data on the structure of silicide/GaAs interfaces have been presented. The most promising and reliable materials from the point of view of temperature stability seem to be compounds of W and Si. REFERENCES [I] [2] [3] [4] [5]

[6]

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14.5 Barrier height at the GaAs/AI interface M. Missous April 1996

A

INTRODUCTION

The Al/GaAs interface is undoubtedly one of the most studied systems of all metal/binary compound Schottky barriers. This interest stems largely from the importance of GaAs and Al from a technological point of view. Aluminium is often used to metallise GaAs structures as it serves as the gate electrode in FET devices. Another very desirable feature of this system is the fact that aluminium, having the FCC structure with a lattice parameter almost exactly l/\/(2) times that of GaAs, is easy to deposit epitaxially on the (100) surface [I]. This latter fact has tremendous consequences on both the structural and electrical properties of the contacts [2] and opens up the way for a whole class of novel devices based on the epitaxial Al/GaAs interface [3,4] including the growth of GaAs on aluminium [5,6]. The barrier height of Al/GaAs diodes depends largely on the conductivity type, ranging from 0.7 - 0.8 eV for n-type to 0.6 eV for p-type [7]. Residual semiconductor oxides and other interfacial layers (e.g. the appearance of AlGaAs at the interface as a result of annealing) are also important in controlling the value of the barrier height. B

Al/GaAs INTERFACES

Bl

Barrier Height at Etched Surfaces

There have been a large number of measurements of Schottky barrier heights of aluminium on a wide range of chemically prepared and air-cleaved n- and p-type GaAs surfaces. These have been reviewed in an excellent article by Robinson [8]. Bl. 1 Al on etched (100) GaAs surfaces Baliga et al [9] found a barrier height of 0.67 eV with an ideality factor of 1.004 for Al on LPE grown n-GaAs doped at 2 - 3 x 1015 cm"3. After annealing at 450 0 C for one hour, the barrier height increased to 0.70 eV and the ideality factor rose to 1.18. The degradation of the electrical characteristics was correlated with severe interdiffiision of Al, Ga and As at the interface as revealed by SIMS. Christou and Day [10], studying Schottky diodes formed by electron-beam evaporation of Al on 1017 cm"3 n-GaAs, obtained a barrier height of 0.70 eV which increased to 0.80 eV after annealing to 450 0 C. Using Auger electron spectroscopy (AES), they found pronounced interdiffiision at the interface with migration of Ga and As throughout the Al film as a result of heat treatment. The as-deposited diodes had rather broad Auger spectra at the interface, probably due to grain boundary diffusion across the small grain size (50 nm) of the Al film. Similar studies by Missous et al [2] on polycrystalline Al film evaporated on MBE grown n-GaAs (1016 cm"3) yielded a barrier height of 0.78 eV with an ideality factor (n) of 1.03 for as-deposited diodes. The average grain size of the Al film was around 200 nm which was substantially higher

than that reported by Christou and Day [10] and hence more abrupt interfaces were obtained. However after annealing at 5000C, the barrier height increased to 0.81 eV and n to 1.06 with a pronounced recombination current component in the I/V characteristics. The Auger depth profiles (ADP) data showed strong interdiffiision at the interface. There was however no trace of Ga in the aluminium film in contrast to Christou and Day [10] but in line with Kim et al [11]. The effects of residual oxides on the GaAs surface have been investigated in some detail by Hasegawa et al [12] who found that the barrier height on oxidised GaAs surfaces (deionised water for two hours and hydrogen peroxide for two minutes) was smaller than that on etched GaAs surfaces (H2SO4:H2O2:H2O 5:1:1 for one minute) by some 0.08 eV (the barrier heights were 0.71 eV and 0.79 eV for the two sorts of surfaces respectively). The barrier height recovered to the as-etched value when the samples were annealed at 3000C for 30 minutes, probably due to the reduction of the surface oxide by Al. The variation of the barrier height was explained by differences in the oxide composition and bonding strength of the metal-oxides (Al-O, Ga-O, AsO etc.). Bl. 2 Al on air-exposed (110) GaAs surfaces The Schottky barrier instabilities due to contamination on air-cleaved (110) Si-doped (2 - 5 x 1016 cm"3) GaAs were thoroughly investigated by Newman et al [13]. A barrier height of 0.76 eV (n = 1.07) was formed on as-deposited diodes which increased to 0.83 eV (n = 1.07) upon annealing to 370 0 C. The changes in the barrier heights were explained in terms of the reaction of the Al with the native oxide at the interface. B2

Barrier Height at Clean Surfaces

B2,1 UHV cleaved (110) surfaces 1.

Thin films

Lindau et al [14] studied the electronic structure of the (110) Al/GaAs interface and found that Fermi level pinning was completed at less than 0.2 monolayers (ML) of deposited Al. They proposed that defect states at or close to the interface were responsible for the barrier height fonnation. Photoemission core level data suggested a replacement interaction between Al and Ga in the topmost GaAs surface layer. However, recent results by Kendelewicz et al [15] on Al/ n-GaAs (110) interfaces fabricated in UHV at -800C, have shown that the Fermi level remained unpinned at least up to 3 ML coverage of Al, in contrast to room temperature deposition. The low temperature behaviour was correlated with the growth of a more uniform Al overlayer which inhibited cluster and defect formation. In a similar set of experiments, Chin et al [16] demonstrated that by controlling the surface disturbance (i.e., varying the substrate temperature and amount of Al deposited), one can modify the Schottky barrier formation process, going from the Schottky limit, which does not have pinning centres, to the Bardeen limit which does. 2.

Thick films

Thick Al contacts on (110) cleaved GaAs have been reported by several groups starting with

Spitzer and Mead [17] who used photoresponse and C/V methods to measure a range of metals on n- and p-GaAs, and more recently by Newman et al [18] who carried out similar types of experiments on n-GaAs. The results for the two groups are summarised in TABLE 1, below TABLEl Group

3.

Spitzer and Mead [ 17]

Newman et al [ 18]

Method

C/V

PR

I/V

C/V

4>b(eV)

0.78-0.92

0.80

0.80-0.85

0.84-0.93

Annealing effects on the barrier height

Newman et al [19] carried out a detailed study of the effects of annealing Ag3 Al and Au on clean, cleaved (110) n-GaAs with a doping level of 1017 cm"3. The diodes were fabricated by slow controlled evaporation of clean metal overlayers to a thickness of 100 nm, and annealed for 10 minutes in nitrogen at various temperatures.

Barrier height (eV +/-0.02)

The results shown in FIGURE 1 were obtained.

Evaporated AI/GaAs Interface I-V Method C-V Method

Anneal Temperature

( C)

FIGURE 1. Barrier height as a function of annealing temperature.

The measurements were made at room temperature after annealing. The authors used photoemission spectroscopy to correlate annealing-induced microscopic changes in the electronic and chemical structure of the interface with changes in the electrical characteristics and attributed the increase in the barrier height to the formation of an AlGaAs alloy at the interface. However in a recent publication, the same authors [20] performed annealing experiments on both n- and p-type GaAs and found that the increase in the barrier height of the diodes formed on nGaAs was equal in magnitude to the decrease in the barrier height of p-GaAs diodes and concluded that the annealing-induced changes in the electrical characteristics can be attributed to a shift in the interface Fermi level pinning position and not to the formation of an AlGaAs

interfacial layer. This last suggestion has been recently questioned by Chambers et al [21] who studied the annealing behaviour of thin epitaxial Al films grown on MBE GaAs layers at temperatures ranging from 60 to 5500C. Using X-ray photoelectron spectroscopy and low energy electron diffraction (LEED), the formation of a thin AlGaAs phase at the interface upon annealing, coupled with the lack of excess As at the interface either before or after annealing, provided strong support to the proposal that it is the presence of the larger bandgap AlGaAs that causes the larger barrier height rather than the creation and population of antisite defects as postulated by Newman et al [20]. B2. 2 Heat cleaned (100) surfaces Barrier heights at clean oxide-free Al (and 13 other metal) interfaces with (100) n-GaAs and pGaAs were measured by Waldrop [22,23]. X-ray photoemission and LEED analysis verified that the GaAs surface was clean and ordered. At least 200 nm of Al was UHV evaporated onto the surface at room temperature. The samples were LEC grown n-GaAs:Se doped to 6 x 1016 cm"3, and horizontal Bridgman grown p-GaAs:Cd doped to 8 x 1016 cm"3. The barrier height values obtained from electrical measurements are shown in TABLE 2. TABLE 2 Measurement method

I/V

C/V

Semiconductor polarity

n-type

p-type

n-type

4>b(eV)

0.85

0.61

0.84

The I/V values incorporate image force corrections of 0.04 eV (n-type) and 0.03 eV (p-type). Note that at doping levels of 6 x 1016 cm"3 there may be noticeable thermionic-field emission which would reduce the barrier height for the n-type samples. Another measurement on clean (100) MBE grown GaAs that was capped with arsenic was made by Eglash et al [24]. Up to 30 ML of Al were slowly evaporated on n-type MBE GaAs doped with Si to 5 x 1016 cm"3. There was no significant heating of the substrate during deposition. The photoemission measurements gave a barrier height value of 0.70 eV (±0.01 eV). Subsequently, 105 ML of Al were deposited on the samples which were heated to 200 0 C to melt the In used to bond the back to form the ohmic contact. A further 700 ML of Al were then evaporated in a conventional vacuum evaporator. The barrier heights obtained are shown in TABLE 3. TABLE 3 Measurement method (|)b (eV)

I/V 0.77 (±0.01)

C/V 0.99 (±0.01)

B2. 3 The effect of interfacial layers on the barrier height The influence of chalcogen atom (Se, S and Te, 0.15 to 2.2 nm thick) interlayers between Al and n-GaAs (6 x 1016 cm"3) was found by Waldrop [25] to result in a barrier height of 0.35 to 0.5 eV (with n ranging from 1.03 to 1.19) compared with 0.75 - 0.85 eV for the ideal Al/GaAs interface. Thus Se, S and Te cause a 0.40 eV decrease in the barrier height independent of the amount of chalcogen deposited. The barrier height lowering was considered to be due to the metal-chalcogen interaction, not to a chalcogen-GaAs interaction, leading to a fairly large and negative heat of formation (-173 kcal/mol for Al2S3). The chalcogenide-forming reactions associated with Ag, Au and Pd have a heat of formation greater than -20 kcal/mol and hence resulted in an increase in the barrier height. Miller and Nathan [39] have studied diode structures grown by molecular-beam epitaxy utilizing thin Si interfacial layers. These Si layers were unintentionally very heavily n-type doped with arsenic from the MBE system or from the substrate. This n-type doping is intentionally compensated with p-type (Al) doping during the growth of the Si layer in an attempt to modulate the Schottky barrier height. The resultant barrier height, as determined by I-V and C-V measurements, increased with increased acceptor doping in the Si from 0.34 eV for no Al doping to a maximum of 1.07 eV. B2. 4 MBE (100) surfaces With the advent of molecular beam epitaxy (MBE), it has become possible not only to prepare clean GaAs surfaces but to perform in-situ metallisation as well. And whereas metals deposited onto room temperature (110) III-V semiconductor surfaces are generally polycrystalline, epitaxial layers of aluminium can easily be grown on (100) surfaces by MBE. These films are generally (100) oriented with the Al unit cell rotated by 45 degrees with respect to that of the underlying GaAs unit cell [1,2,26-29]. (a)

Barrier height as a function of GaAs (100) surface reconstruction

The (100) surface of GaAs has a number of reconstructions with widely varying structure and composition. It has been reported by Cho and Dernier [26] and later by Wang [27] that the barrier height depends on the particular reconstruction adopted by the GaAs surface prior to Al deposition, the barrier height decreasing with increasing As coverage. However, detailed studies by Barret et al [28] and Svenson et al [29], and later work by Missous et al [30], find no dependence of barrier height on reconstruction. A point of interest is that the epitaxial Schottky diodes made by Missous et al [30] showed the most ideal electrical characteristics, with accurately exponential I/V characteristics over 8 decades and n values around 1.01. The barrier heights found by the different groups on the different GaAs reconstructed surfaces are shown in TABLE 4.

TABLE 4 GaAs surface 4x6 GrQU

(b)

P

c(2x8)

c(4x4)

<J>bQ/V)

4>b(C/V)

4>b(W)

^b(CZV)

4>b(I/V)

4>KC/V)

Choetal[26]

-

0.80

0.76

0.74

Svenson et al [29]

0.75

0.67

0.76

0.69

0.74

0.66

Barret et al [28]

0.76

-

0.77

-

0.80

Wang [27]

0.80

-

0.76

-

0.74

Missousetal[30]

078

077

078

0.77

077

077

High temperature stability

The good high temperature stability of epitaxial Al/GaAs, as compared to conventional polycrystalline Al films, after heat treatment to 5000C for one hour was attributed by Missous et al [2] as due to a reduction in grain boundary diffusion as attested by the abrupt Auger depth profiles both before and after annealing. The increase in barrier height (by about 0.05 eV) was explained in terms of formation of a thin AlGaAs interfacial layer. These results are at variance with those of Svenson et al [29] who found extensive deterioration of the I/V characteristics on annealing. It should be noted that their diodes had poor, non-linear C/V profiles and showed considerable variation with position across a slice. Huang et al [37] have investigated the interfacial stability, surface morphology and electrical characteristics OfMoAlx on n-GaAs by using X-ray diffraction, scanning electron microscopy, sheet resistance and current-voltage measurements. The compositions of RF-cosputtered MoAlx films were x = 0.35, 2.7 and 7.0, respectively. The contacts were annealed by rapid thermal processing in the temperature range 500 - 10000C for 20 s. The interfaces OfMoAJ0 35/ GaAs and MoAl2 7 / GaAs were stable up to 9000C, while the interfaces OfMoAl7 J GaAs were less stable and reactions occurred above 8000C. The variations of sheet resistance and the barrier heights of the Schottky diodes as a function of annealing temperature were correlated to the interfacial stability. The MoAl2 7/ n-GaAs diodes exhibited the best stability and were characterized by the highest barrier height (0.98 eV) and near unity ideality factor (1.11) after annealing at 900 0 C. (c)

Al on p-GaAs

Schottky barriers to p-GaAs have always been problematic because of the low barrier height that Al makes with p-GaAs. However for in-situ epitaxial Al and Sb films deposited by MBE, Missous et al [30,31] were able to obtain excellent diodes with barrier heights of 0.64 eV and n values of 1.01 over 4 decades in current. The C/V barrier height was also equal to 0.64 eV. The sum of the C/V barrier height on n- and p-type GaAs was equal to 0.77 + 0.64 = 1.41 eV which is virtually equal to the GaAs bandgap.These results imply that the parameters that determine the barrier height (i.e., the interfacial dipole and any interface states that may be present) are essentially the same for both n- and p-type material.

(d)

Effect of interfacial layers on the barrier height

Barrier height (eV)

Eglash et al [32] studied the effect of interfacial doped layers on the barrier height of epitaxial Al on (100) n-GaAs Si-doped to 5 x 1016 cm"3. A heavily doped interfacial layer of Be-doped pGaAs was grown on the n-type samples prior to Al deposition. The growth conditions gave abrupt, ideal interfaces. FIGURE 2 summarises the data.

Epitaxial Al/GaAs

Interfacial layer thickness (nm)

FIGURE 2.

The effect of interfacial doped layers on barrier height.

The maximum barrier height of 1.24 eV compares with a room temperature bandgap of 1.42 eV. These data show the possibility of controlling the barrier height of Schottky diodes for engineering purposes. Hirose et al [33] reported a lowering of the barrier height when rare earth metal (REM) thin films were deposited between the Al and the n-GaAs (2 x 1017 cm"3) grown by MBE. A decrease in the barrier height from 0.78 eV to 0.60 eV was achieved. Since the REM/GaAs Schottky barrier heights were found to be in the range 0.85 to 0.87 eV, the lowering was considered to be due to alloy formation and diffusion at the interface in a similar fashion to that in Al/GaAs interfaces modified by chalcogenide-interlayers [25]. Hydrogen sulphide chemisorbed onto MBE n-GaAs prior to MBE growth of an Al film considerably reduced the barrier height in an experiment by Massies et al [34]. The (100) GaAs layers were Sn-doped to 1016 to 10 19 cm "3 and of thickness 1-2 microns. The effect of 1000 Langmuirs OfH2S decreased the barrier height from 0.76 (±0.02) eV to 0.40 (±0.03) eV. This Schottky barrier lowering can be used to obtain monocrystalline non-alloyed ohmic contacts with low specific contact resistance. Svenson and Anderson [35] found Ga cluster growth on (100) MBE GaAs with c(2><8) surface reconstruction. The Ga interlayer led to an increase in barrier height and the authors concluded that this should permit barrier height tuning (0.75 to 1 eV) while still retaining Al as the contact material for device purposes. Zhang et al [38] have recently used low temperature (LT) GaAs interfacial layers to modify the

barrier height. By depositing up to 50 nm of LT GaAs, they were able to decrease the effective barrier height from 0.79 to 0.35 eV on n-type materials and to increase it from 0.55 to 0.72 eV on p-type materials. (e)

Novel device applications

Buried metal layers have attracted a great deal of attention because of their potential use in devices such as metal base transistors. The requirement of thermal stability and lattice matching between the metal and the semiconductor places stringent conditions on any metal-semiconductor pair used. Okamoto et al [5] found extensive interdiffusion between the top GaAs and the underlying Al film, resulting in complete consumption of the thin AI layer at growth temperatures greater than 5000C. Recently, Tadayon et al [6] using migration enhanced epitaxy (MEE) grew the topmost GaAs at temperatures as low as 200 0 C and thus prevented the thin Al epilayer from totally reacting. However no transistor action was reported. In related studies, Harbison et al [4] found that the intermetallic compound NiAl was not only well lattice matched to GaAs (lattice mismatch is 2 %) but also stable at high temperatures and were thus able to grow GaAs/NiAl/GaAs structures with NiAl films as thin as 3.3 nm. The films were in registry with the GaAs layers on both sides of the interface and were continuous over macroscopic dimensions. Tabatabaie et al [36] were able to fabricate double barrier resonant tunnelling structures with the central quantum well material consisting not of GaAs as in the conventional case, but of metallic NiAl. The results above attest to the exciting new avenues opened up by the repeated epitaxy of metals and semiconductors using MBE, and there is no doubt that we shall see more work devoted to this very important class of materials, opening up the way for a whole new class of devices. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] II1] [12] [13]

R. Ludeke, L.L. Chang, L. Esaki [Appl. Phys. Lett. (USA) vol.23 no.4 (1973) p.201 ] M. Missous, E.H. Rhoderick, K.E. Singer [J. Appl Phys. (USA) vol.59 no.9 (1986) p.3189 ] T. Sands et al [Appl. Phys. Lett. (USA) vol.52 (1988) p.1216 ] J.P.Harbison, T. Sands, N. Tabatabaie, W.K. Chan, L.T. Flores, V.G. Keramidas [ Appl. Phys. Lett. (USA) vol 53 no. 18 (1988) p. 1717 ] K. Okamoto, C.E.C. Wood, L. Rathburn, L.F. Eastman [ J. Appl Phys. (USA) voL53 (1982) p.4521 ] B. Tadayon et al [ Appl Phys. Lett. (USA) vol.53 no.26 (1988) p.2664 ] E.H. Rhoderick, R.H. Williams [ Metal semiconductor contacts (Clarendon, Oxford, 1988) ] G. Y. Robinson [ inPhysics and Chemistry of HI-Vcompound semiconductor interfaces Ed. CW. Wilmsen, ch.2 (Plenum Press, New York, 1985) ] BJ. Baliga, R. EhIe, A. Sears, P. Campbell, W. Garwacki, W. Katz [ IEEE Electron Device Lett. (USA) vol. EDL-3 no.7 (1982) p. 177 ] A. Christou, H.M. Day [ J. Appl. Phys. (USA) vol.47 (1976) p.4217 ] H.B. Kim, G.G. Sweeney, T.M.S. Heng [ Gallium arsenide and related compounds, Ed. J. Bok, Inst. Phys. Conf. Ser. no.24 (Institute of Physics, London, 1974) p.307 ] F. Hasegawa, M. Onomura, C. Mogi, Y. Nannichi [ Solid-State Electron. (UK) vol.31 no.2 (1988) p.223] N. Newman, Z. Liliental-Weber, E.R. Weber, J. Washburn, W.E. Spicer [ Appl Phys. Lett. (USA)

[14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33] [34] [35] [36] [37] [38]

[39]

vol.53 no.2 (1988) p. 145] I. Lindau, P.R. Skeath, CY. Su, W.E. Spicer [ Surf. Sci. (Netherlands) vol.99 no. 1 (1980) p. 192 ] T. Kendelewicz et al [Appl. Phys. Lett. (USA) vol.48 no. 14 (1986) p.919 ] K.K. Chin et al [Mater. Res. Soc. Symp. Proc. (USA) vol.77 (1987) p.297 ] W.G. Spitzer, CA. Mead [ J. Appl. Phys. (USA) vol.34 (1963) p.3061 ] N. Newman, T. Kendelewicz, L. Bowman, W.E. Spicer [ Appl. Phys. Lett. (USA) vol.46 (1985) p. 1176] N. Newman et al [ J. Vac. Sci. Technol. A (USA) vol.33 no.3 (1985) p.996 ] N. Newman, W.E Spicer, E.R. Weber [J. Vac. Sd. Technol. B (USA) vol.5 no.4 (1987) p.1020 ] S.A. Chambers [Phys. Rev. B (USA) vol.39 no.17 (1989) p.12664 ] J.R. Waldrop [ J. Vac. Sci. Technol. B (USA) vol.2 no.3 (1984) p.445 ] J.R. Waldrop [ Appl. Phys. Lett. (USA) vol.44 no. 10 (1984) p. 1002 ] SJ. Eglash, M.D. Williams, P.H. Mahowald, N. Newman, I. Lindau, W.E. Spicer [ J. Vac. Sd. Technol. B (USA) vol.2 no.3 (1984) p.481 ] J.R. Waldrop [Appl. Phys. Lett. (USA) vol.47 no.12 (1985) p.1301 ] A.Y. Cho, P.D. Dernier [ J. Appl. Phys. (USA) vol.49 no.6 (1978) p.3328 ] W.I. Wang [J. Vac. Sci. Technol. B (USA) vol. 1 no.3 (1983) p.574 ] C Barret, J. Massies [ J. Vac. Sci. Technol. B (USA) vol. 1 (1983) p.819 ] S. Svenson, G. Landgren, T.G. Anderson [J. Appl. Phys. (USA) vol.54 no.8 (1983) p.4474 ] M. Missous, E.H. Rhoderick, K.E. Singer [ J. Appl. Phys. (USA) vol.60 no.7 (1986) p.2439 ] M. Missous, E.H. Rhoderick, K.E. Singer [ Electron. Lett. (USA) vol.22 (1986) p.241 ] S.W. Eglash, S. Pan, D. Mo, W.E. Spicer, D.M. Collins [ Jpn. J. Appl. Phys. Suppl. (Japan) vol.22 suppl.22-l(1983)p.431] K. Hirose, H. Tsuda, T. Mizutani [J. Appl. Phys. (USA) vol.64 no.l 1 (1988) p.6575 ] J. Massies, J. Chaplart, M. Laviron, NT. Linh [Appl. Phys. Lett. (USA) vol.38 no.9 (1981) p.693 ] S.P.Svenson, T.G. Anderson [ J. Vac. Sci. Technol. B (USA) vol.3 no.2 (1985) p.760 ] N. Tabatabaie,T. Sands, J.P.Harbisson,H.L. Gilchrist, VG. Keramidas [Appl. Phys. Lett. (USA) vol.53 no.25 (1988) p.2528 ] T.S. Huang, J.G. Peng, C C Lin [ Advanced Metallization and Processing for Semiconductor Devices and Circuits - II. Symposium (Mater. Res. Soc, Pittsburgh, USA, 1992) p.511-6] K. Zhang, D. L. Miller [ Defect Engineering in Semiconductors Growth, Processing and Device Technology Symp. San Francisco, Ca, USA, 26 Apr.-l May 1992 (MRS, Pittsburgh, USA, 1992) p.899-904 ] T.J. Miller, M.I. Nathan [ J. Appl. Phys. (USA) vol.76 no. 1 (1994) p.371 ]

14.6 Barrier height at the Ag/GaAs interface M. Missous April 1996

A

INTRODUCTION

Silver (Ag) forms a relatively abrupt Schottky barrier with GaAs because of its inertness. The barrier height depends largely on the conductivity type of the semiconductor and to a lesser extent on the residual oxides on the semiconductor or any chemical interaction with the metal. Ag interacts very weakly with Ga or As as has been demonstrated by Ludeke et al [1] in detailed correlated chemical, structural and electronic studies of Ag on MBE-grown GaAs. Photoemission core-level studies gave clear evidence that the Ag/GaAs interface is abrupt. B

Ag/GaAs INTERFACE

Bl

Etched Surfaces

BLl

Ag on (100) and (111) GaAs surfaces

One of the first uses for silver on GaAs was as microwave mixer diodes. Genzabella and Howell [2] evaporated Ag on (111) and (100) epitaxial GaAs doped to 2 x 1016 cm"3. From I/V and C/V measurements, the barrier height was found to be 0.94 ± 0.02 eV for the (111) direction and 0.85 ± 0.02 eV for the (100) direction. B1.2

Agon air-exposed (110) surfaces

Newman et al [7] evaporated Ag on air cleaved (110) Si-doped (2-5 x 1016 cm"3) GaAs. The pressure during evaporation was approximately 10"7Torr. The contacts were then subjected to thermal annealing at 370 0 C leading to the results in TABLE 1. TABLEl 4Vv) ( e V)

Ideality factor

KC/v) (eV)

Room Temperature

0.95

1.07

1.06

370 0 C

0.79

1.06

0.85

More recently Miret et al [8] studied the effect of electrical aging of Ag/GaAs diodes formed on air-exposed (110) surfaces. The electrical aging was performed by keeping the diodes at -17 V and 2.3 x 10"3 A/cm2 from 0 to about 270 minutes, the effect of which was to reduce the I/V barrier height by as much as 75 meV from its as-deposited value of 0.96 eV. No change in the ideality factor was detected (n = 1.06). The diodes were able to recover to their original values after 5 days. The changes in electrical aging were attributed to the creation and/or annihilation of deep-level traps near the interface.

B2

Clean Surfaces

B2.1

Heat cleaned (100) surfaces

Waldrop [3,4] measured barrier heights of Ag (and 13 other metals) at clean oxide-free interfaces with (100) plane orientated GaAs samples of both conductivity type. X-ray photoemission and LEED analysis showed an ordered, contamination-free surface prior to ultra high vacuum evaporation of a 200 nm thick Ag layer at room temperature. The samples were n-type GaAs doped with Se to 6 x 1016 cm'3 (LEC grown material), and p-type GaAs Cd doped to 8 x 10 16 cm"3 (Bridgman grown material) and the barrier heights obtained are given in TABLE 2. TABLE 2 I/V method

4>b (eV)

C/V method

n-type

P-type

n-type

O90

O50

089

The current-voltage values incorporate an image force correction of 0.04 eV for the n-type samples and 0.03 eV for the p-type samples. Note that at a doping level of 6 x io 16 cm'3 there may be noticeable thermionic field emission which would reduce the barrier height for the n-type samples. Massies et al [5] measured Schottky barrier heights of Ag evaporated in an MBE-type system on bulk-grown n-type samples doped from 5 x 1016 to 1017 cm"3. The surfaces were prepared by ion etching and annealing under UHV. The barrier heights obtained depended upon the GaAs surface reconstruction and were equal to 0.67 ± 0.02 eV for the c(8><2)-Ga stable surface and 0.60 ±0.02 eV for the arsenic rich ( 1 x 1) surface. B2.2

MBE (100) surfaces

Massies et al [6] have studied in some detail the epitaxial relationship of in-situ deposited Ag on (100) GaAs as a function of surface reconstruction and the deposition temperature. The relationships shown in TABLE 3 were obtained. TABLE 3 Reconstructed surface

Growth temperature

Epitaxial relationship

(2x4) As stable or (4x2) Ga stable

150 0 C

(110)Ag//(001)GaAs with [001]AgZT[11O]GaAs

c(4x4) As rich or (2x4) As stable

> 300 0 C

(00l)Ag//(001)GaAs with [010]Ag//[l 1O]GaAs

Ludeke et al [1] have also reported on the epitaxial and electrical measurements of MBE grown Ag/GaAs structures. Ag was deposited at room temperature onto smooth, uncontaminated, undoped (100) GaAs films. The epitaxial Ag films were (110) oriented in agreement with the findings of Massies et al [6]. The barrier heights were monitored using photoemission spectra. The values obtained were:

0.083 ± 0.03 eV for the As-stabilised c(2*8) surface 0.097 ± 0.03 eV for the Ga-rich (4x6) surface. Wang et al have also reported electrical properties of molecular beam epitaxy 'in-situ' grown Ag on (001) GaAs Schottky diodes. X-ray rocking curves showed a (111) main peak for 'in-situ' Ag grown at low temperature. During annealing, the main peak of Ag rotated from (111) to (200) to closely match that of the underlying GaAs lattice. An as-deposited C/V barrier height of 0.99 eV was measured which then decreased upon annealing. The decrease in the barrier height was correlated with a dissociation of Ga from GaAs resulting in an increase in uncompensated ions at the metal/semiconductor interface therefore explaining the observation that carrier concentrations increase after annealing. B2.3

UHV cleaved (110) surfaces

(a)

Thin films

Ludeke [9] used photoemission spectroscopy to investigate the mechanism of Ag/GaAs Schottky barrier formation. Various thicknesses of Ag (up to 20 monolayers) were deposited on n-type and p-type (110) cleaved GaAs. Clustered Ag growth and undetectable chemical interactions with the substrate were observed, and the surface Fermi level changed with the thickness of the Ag layer. For the thickest layers, the barrier heights were 0.89 eV (n-type) and 0.35 eV (p-type). The slow approach to Fermi level pinning was taken as evidence that Ag forms clusters rather than a uniform layer on the GaAs surface. (b)

Thick films

Several groups have done intensive studies of thick Ag contacts on clean cleaved (110) GaAs surfaces. Their results are summarised below in TABLE 4. TABLE 4 Newman et al [10]

Ismail et al [11]

4>b(i/v) (eV)

4V/V) (eV)

4>b(w) (eV)

4>b(c/v) (eV)

0.87 - 0.90

0.95-0.99

082

L03

The C/V barrier heights are higher than those obtained by JTV by about 0.1 to 0.2 eV and both barriers are less than those on the air-exposed (110) surfaces by about 0.1 eV. The thermal annealing of Ag/GaAs (110) interfaces [12] indicates that no change in the barrier height is observed for annealing temperatures up to 400 0 C in both photoemission spectroscopy (PES) and device measurements. No detectable evidence of chemical interactions was observed in the PES spectra. Electrical aging of Ag on UHV cleaved (110) GaAs [8] showed only a very small change in the barrier height (less than 5 meV at -19 V and 1.4 A/cm2), unlike Ag contacts on air cleaved surfaces. Electrical aging seems to reduce the difference ( about 0.1 eV) of barrier height, from

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the as-deposited conditions, between the two kinds of diodes, (c)

The effects of interfacial layers on <{>b

The use of interfacial layers to modify <J)b is well established. In the case of Ag, Spaltmann et al [14] prepared Schottky diodes by depositing Mn and thereafter Ag onto ultrahigh vacuum cleaved n-type GaAs(110). The Mn interfacial thickness was varied from 0 to 15 A, while the subsequent Ag coverage was 600 A for all the samples studied. The variation of the Mn interlayer thickness allowed the barrier height to be tuned from 0.92 eV for pure Ag, with no Mn interlayer present, to 0.76 eV for 15 A of Mn (the barrier height of pure Mn on GaAs(110) was 0.74 eV). For interlayers above 1 A the barrier height was found to decrease exponentially with increasing Mn thickness. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] [II] [12] [13] [14]

R. Ludeke, T.C. Chiang, D.E. Eastman [J. Vac. Sci. Technol. (USA) vol.21 (1982) p.599 ] CF. Genzabella, CM. Howell [ Inst. Phys. Conf. Ser. (UK) vol.3 (1966) p. 131 ] J.R. Waldrop [ J. Vac. Sci. Technol. B (USA) vol.2 no.3 (1984) p.445 ] J.R. Waldrop [ Appl. Phys. Lett, vol.44 no. 10 (1984) p. 1002 ] J. Massies, P. Devoldere, N.T. Linn [ J. Vac. Sci. Technol. (USA) vol. 15 no.4 (1978) p. 1353 ] J. Massies, P. Etienne, N.T. Linh [ Surf. Sci. (Netherlands) vol.80 (1979) p.550 ] N.Newman, Z. Lilienthal-Weber, E.R. Weber, J. Washburn [ Appl. Phys. Lett. (USA) vol.53 no.2 (1988) p. 145] A. Miret, N. Newman, E.R Weber, Z. Lilienthal-Weber, J. Washburn, W.E. Spicer [ J. Appl. Phys. (USA) vol.63 no.6 (1988) p.2006 ] R. Ludeke [ J. Vac. Sci. Technol. B (USA) vol. 1 no.3 (1983) p.581 ] N. Newman, T. Kemdelewicz, O. Thompson, S.H. Pan, SJ. Eglash, W.E Spicer [ Solid-State Electron. (USA) vol.28 no.3 (1985) p.307 ] A. Ismail, J.M. Palau, L. Lassabatere [ Rev. Phys. Appl. (France) vol. 19 (1984) p.205 ] N. Newman et al [ J. Vac. Sci. Technol. A (USA) vol.3 no.3 (1985) p.996 ] Y.H Wang, M.P. Houng, F.H. Chen, P.W. Sze, M. Hong, J.P. Mannaerts [ J. Electron. Mater. (USA) vol.21 no.9 (1992) p.911 ] D. Spaltmann, J. Guerts, N. Esser, D.R.T Zahn, W. Richter, RH. Williams [ Semicond. Sci. Technol. (UK) vol.7 no.3 (1992) p.344 ]

14.7 Barrier height at the GaAs/Au interface

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M. Missous April 1996

A

INTRODUCTION

The Au/GaAs interface is one of the most intensively studied metal/GaAs contacts. This system is used both as an ohmic contact (in conjunction with Ge or Zn) and as a large Schottky barrier height in a wide variety of semiconductor devices, e.g. solar cells and field effect transistors (FETs). The interaction mechanism of Au with the GaAs surface is at the origin of the formation of the Schottky barrier as well as the preparation of ohmic contacts and has been the subject of a great deal of experimental study [1-4]. The metallurgical effects at the Au/GaAs interface show considerable chemical mixing and interdifiiision [3,4] rendering the barrier height (cj)b) extremely sensitive to the thermal processing needed for device applications. (|>b is reduced considerably by heat treatment and experimental evidence of thermally induced migration of Ga from GaAs into the Au film has been demonstrated by several workers [2-4]. The change in chemical structure at the Au/GaAs interface has been studied by a wide variety of analytical techniques including photoemission spectroscopy (PS) [5], secondary ion mass spectrometry (SIMS) [3], Auger electron spectroscopy [4] and Rutherford backscattering (RBS) [2], amongst others. B

Au/GaAs INTERFACES

Bl

Etched Surfaces

There have been numerous measurements of Schottky barrier heights on chemically etched and air-cleaved n-type GaAs, the most complete of which were those of Smith [6] who measured barrier heights on (110), (111) and (100) surfaces and obtained values of (|)bI/v of 0.93, 0.89 and 0.83 eV respectively. A dependence of (|)b on crystal orientation was also noted by other workers [7]. Ismail et al [8], Kendelewicz et al [9] and Montgomery et al [10] found virtually no influence of the oxides on c|)b since little difference in the values of (|)bI/v was measured for Au (and Al and Pd) on air-cleaved GaAs as compared to surfaces cleaved in UHV. Kowalczyk et al [11] studied the effect of oxide composition on the barrier height of several metals on (100) GaAs. Au does not reduce the oxides (as Al does) and there was considerable band bending at the free surface of the GaAs which did not change substantially when the metals were deposited. Recently, very high Schottky barrier heights were obtained by Reineke and Memming [12] using electrochemical deposition of metals on n-GaAs. For Au, $bVV of 1.19 eV was obtained with an n value of 1.22. Again the effect of the native oxide layer was found to be minimal. The thermal degradation of Au/GaAs contacts has recently been reassessed by Yeh and Holloway

[13]. Using rather heavily doped material (5.5 x 1017 cm"3 Si-doped) they found that the degradation of the contacts proceeded in two stages. In stage one, which occurred at temperatures between 200 and 2500C, the main effect of thermal annealing was a reduction in the (j)bI/v from 0.77 eV and n = 1.07 to 0.72 eV and n = 1.13. At temperatures higher than 280 0 C (stage two), <J)b decreased rapidly to attain a value of 0.50 eV with n = 1.5 at 320 0 C. An important factor in the degradation mechanism of stage two was a tremendous increase in the carrier concentration at the contact interface (from an as deposited value of 5.6 x 1017 cm'3 to 33.2 x io 17 cm"3 at 320 0 C ). B2

Effect of Hydrogenation

The effect of hydrogenation on the properties of the Schottky barrier height has been investigated by a number of workers [21-23]. Wang and Ashok have reported profound changes to (|>b of metal/GaAs interfaces by atomic hydrogen using an RF plasma in a reactive ion etching (RIE) system as well as hydrogen generated in an electron cyclotron resonance (ECR) system. I-V characteristics of Au/n-GaAs Schottky devices reveal a reduction in the barrier height following the room temperature RF plasma, and a slight increase with ECR hydrogenation at elevated temperatures. Furthermore a large increase in (J)b is seen for p-GaAs (from 0.35 to 0.84 eV for the RF plasma and 0.35 to 0.69 eV for ECR). Dopant deactivation close to the surface is observed with spreading resistance and capacitance-voltage (C-V) measurements for both conductivity types. The passivation of existing deep levels and the creation of new deep levels were found in both hydrogen RIE treated and ECR hydrogenated GaAs. The large ideality factor in I-V plots and large voltage intercept in 1/C2 plots suggest the formation of an insulator-like layer. B3

Effect of Interfacial Layers

The role of thin interfacial layers in controlling (j>bhas been studied by Fujieda [24]. They used thin (1 to 2 nm thick) nonstoichiometric GaAs layers grown by molecular beam epitaxy at 200 0 C under a wide range of arsenic pressures and placed at the interface of metal/GaAs Schottky barriers. By changing the As pressure for the interface film growth, the Schottky barrier heights obtained using Au metals varied in the range of 0.5 to 1.0 eV on n-GaAs and of 0.4 to 0.9 eV on p-GaAs. The wide variation of barrier heights, essentially independent of the metal work function, was explained by strong Fermi-level pinning controlled by defect levels in the interface layer. These defects were related to the non-stoichiometry of the low temperature GaAs layer. B4

Clean surfaces

B4.1

The (100) surface

An in-depth study of Schottky barrier height on heat cleaned, oxide-free (100) n- and p-type GaAs surfaces for Au (and 13 other metals) has been reported by Waldrop [14,15]. X-ray photoemission and low energy electron diffraction analysis verified that the GaAs surface was clean and ordered. However no information on whether the surface was reconstructed was given. At least 200 nm of Au was evaporated in UHV onto this surface at room temperature. The samples used were:

n-GaAs : Se-doped to 6 x 1016 cm'3, LEC grown material p-GaAs : Cd-doped to 8 x io 16 cm"3, horizontal Bridgman grown material. The resulting room temperature measurements gave the values in TABLE 1. TABLEl Method

I/V

C/V

Semiconductor polarity

n-type

P~tyPe

n-type

4>b(eV)

0.89

0.50

0.87

n (ideality factor)

1.03

1.15

^bW values incorporate image force corrections of 0.04 eV for n-type and 0.03 eV for p-type. The sum of the barrier heights on n- and p-type is almost equal to the GaAs bandgap. Note that for a doping level of 6 x 1016 cm"3, there may be noticeable thermionic field emission which would reduce the barrier height for the n-type samples. B4.2

The (110) surface

(a)

Ultrathin Au films

Using X-ray photoelectron spectroscopy (PES), Petro et al [16] have followed the initial stage of band bending on (110) cleaved (unpinned) GaAs as Au was deposited. Provided the evaporation was carried out slowly and without heating the GaAs substrate, (J)b of 0.90 eV was achieved after deposition of one monolayer of Au. Even at room temperature there was evidence of Au intermixing with GaAs with preferential As segregation at the surface at high Au coverages (15 ML). Heating of the substrate at temperatures greater than 1000C caused increased Au-Ga intermixing. (b)

Thick Au contacts

As far as thick metal contacts onto cleaved clean (110) GaAs surfaces are concerned, the first results are those of Spitzer and Mead [17] who used photoresponse (PR) and C/V methods to measure a range of metals on n- and p-type GaAs. Recently both Ismail et al [8] and Newman et al [18] carried out similar types of experiments on n-type GaAs. The results for the different groups are summarised in TABLE 2. Group

Spitzer et al [17]

Method

C/V

4>b(eV)

0.93-0.98

Newman et al [18] PR

W

0.90

0.92

C/V 0.99-1.05

Ismail et al [8] I/V 0.88

C/V 0.98

From the above results it can be seen that ^>baw are slightly larger than <j)bI/v. Note also that the barrier heights are similar to the ones obtained on ultrathin films.

(c)

Annealing effects on barrier heights

Newman et al [19] carried out a detailed study of the effects of annealing Ag, Al and Au on clean, cleaved (110) n-GaAs with a doping level of 1017 cm'3. The diodes were fabricated by slow controlled evaporation of clean metal overlayers to a thickness of 100 nm, and annealed for 10 minutes in nitrogen at various temperatures. The results shown in FIGURE 1 were obtained. Evaporated Au/GaAs

Barrier height (eV +/-0.02)

I/V Method C/V Method

Anneal Temperature (0C)

FIGURE 1. Effect of annealing on barrier height.

In a similar fashion the variation of ideality factor with annealing temperature was as shown in FIGURE 2.

Ideality factor

Evaporated Au/GaAs

Anneal Temperature ( C)

FIGURE 2. Variation of ideality factor with annealing temperature.

The measurements were made at room temperature after annealing. The I/V and C/V data do not include corrections for the image force potential or the Gummel-Scharfetter factor (kT/q) respectively. The authors used photoemission spectroscopy to correlate annealing-induced microscopic changes in the electronic and chemical structure of the interface with the electrical measurements and attributed the decrease in barrier heights to a change in the interface Fermi level position. An Au-Ga rich layer formed at the interface during anneals at 200 to 300 0 C, and leakage currents dominated the I/V characteristics in devices annealed above the Au-Ga eutectic temperature [20]. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] [II] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24]

A.K. Sinha,J.M. Poate [Appl. Phys. Lett. (USA) vol.23 (1973) p.666] A.K. Sinha, J.M. Poate [ in Thin films- interdiffusion and reactions, Eds Poate, Tu, Mayer (Wiley, New-York, 1978) ch.ll] H.B. Kim, G.C. Sweeney, TMS. Heng [ Inst. Phys. Conf. Ser. (UK) vol.24 (1975) p.307 ] G.Y. Robinson [ J Vac. Sci. Technol. (USA) vol. 13 no. 14 (1976) p.884 ] W.G. Petro et al [ J. Vac. Sci. Technol. (USA) vol.21 no.2 (1982) p.585 ] B.L. Smith [ Ph.D Thesis, University of Manchester (1969) ] D. Kahng [ Bell. Syst. Techn. J (USA) vol.42 (1964) p.215 ] A. Ismail, J.M. Palau, L. Lassabatere [ Rev. Phys. Appl. (France) vol. 19 (1984) p.205 ] T. Kendelewicz, N. Newman, RS. List, I. Lindau, W.E. Spicer [ J. Vac. Sci. Technol. B (USA) vol.3 (1985) p. 1206] V. Montgomery, R.H. Williams [ J. Phys. C, Solid-State Phys. (UK) vol. 15 (1982) p.5887 ] S.T. Kowalczyk, J.R. Waldrop, W.R. Grant [J Vac. Sci. Technol. (USA) vol.19 (1981) p.611 ] R. Reineke, R. Memming [ Surf. Sci. (Netherlands) vol. 192 (1987) p.66 ] L.L. Yeh, P.H. Holloway [ J. Vac. Sci. Technol. A (USA) vol.6 no.3 (1988) p. 1478 ] J.R. Waldrop [ J. Vac. Sci. Technol. B (USA) vol.2 no.3 (1984) p.445 ] J.R. Waldrop [ Appl. Phys. Lett. (USA) vol.44 no. 10 (1984) p. 1002 ] W.G. Petro, IA. Babalola, T. Kendelewicz, I. Lindau, W.E. Spicer [ J. Vac. Sci. Technol. A (USA) vol.1 no.2 (1983)p.ll81 ] W.G. Spitzer, CA. Mead [ J. Appl. Phys. (USA) vol.34 (1963) p.3061 ] N. Newman, T. Kendelewicz, L. Bowman, W.E. Spicer [Appl. Phys. Lett. (USA) vol.46 (1985) p. 1176] N. Newman et al [J. Vac. Sci. Technol. A (USA) vol.3 no.3 (1985) p.996 ] N. Newman, W.G. Petro, T. Kendelewicz, S.H. Pan, SJ. Eglash, W.E. Spicer [ J. Appl. Phys. (USA) vol.57 no.4 (1985) p. 1247] S.X. Jin eta\[Appl. Phys. Lett. (USA) vol.62 no.21 (1993) p.2719] Y.G Wang, S. Ashok [ J. Appl. Phys. (USA) vol.75 no.5 (1994) p. 2447) ] Eun Kyu Kim, Hoon Young Cho, Hyeon Soo Kim, Suk-ki Min, Taewhan Kim [ Semicond. Sci. Technol. (USA) vol.7 no.5 (1992) p. 695 ] S. Fujieda [ Appl. Phys. Lett. (USA) vol.61 no.3 (1992) p.288 ]

14.8 Barrier height at the Pt/GaAs and W/GaAs interfaces M. Missous April 1996

A

INTRODUCTION

The thermal stability of metal/semiconductor interfaces is of great practical importance in device technology because as gate metals in MESFET processing, they usually have to be subjected to thermal annealing as part of the device manufacture stage. B

THE Pt/GaAs INTERFACE

In the case of platinum (Pt), thermal annealing is known to significantly affect the barrier height as determined by the I/V and C-V methods. As far as its annealing properties are concerned, Pt has been classified by Sinha and Poate [1] along with the transition metals (Ni, Rh, Pd, Os, Ir) as belonging to a class of metals which, upon heating, form stable compounds both with the metal A(III) and the metalloid B(V) in the A(III)B(V) compound semiconductor. Bl

Etched Surfaces

The use of wet-chemical removal of native oxide in a sealed nitrogen ambient prior to deposition of metal on GaAs was shown by Ren et al [11] to be an effective method of engineering the Schottky barrier height of Pt/GaAs Schottky diodes. A barrier height of 0.98 eV for Pt on n-type GaAs was demonstrated which was considerably higher than the barrier height of conventionally processed TiPtAu contacts (0.78 eV). MESFETs fabricated using PtAu bilayer contacts showed reverse currents an order of magnitude lower than TiPtAu contacted companion devices, higher reverse breakdown voltages and much lower gate leakage. Utilizing this technology of oxide removal and the PtAu bilayer contact provided a much simpler method of enhancing the barrier height on n-type GaAs than other techniques such as counter-doping the near-surface or inserting an interfacial layer. B2

RF Sputtered Contacts

Sinha and Poate [2] have used both electrical measurements and Rutherford backscattering (RBS) analysis to investigate the effects of heat treatment on RF sputtered Pt films on 1016 cm"3 n-GaAs. The electrical properties of the contacts were highly non-ideal for both as-deposited and 500 0 C annealed samples with ideality factors (n) of 1.29 and 1.19 and (j)b of 0.84 eV and 0.88 eV respectively. The backscattering spectra from the annealed Pt films showed the appearance of a layered arrangement of Pt/PtGa, presumably due to Ga out-difiiision into Pt because of the large electronegativity difference between Ga and Pt, near the surface and a layer OfPtAs2 near the GaAs interface. This last layer was thought to be responsible for the slight increase in (J)b. Murarka [3] found a similar increase in the forward ^b0A0 from 0.92 to 0.96 eV after annealing to 350 0 C in vacuum, n remained constant at 1.09.

B3

Electron Beam Evaporated Contacts

Fontaine et al [4], using rather highly doped n-GaAs samples (1017 cm"3), obtained a cj)b(W) of 0.65 eV with n = 1.45. After annealing to 3 50°C for up to 20 hours, n improved slighty to 1.15 and c|)b increased to 0.81 eV. In contrast, the <$>HC/v) decreased from an as-deposited value of 1.1 eV to 0.98 eV after heat treatment. Photoresponse measurements on the same diodes gave values for <j)b of 0.90 eV which were almost independent of thermal annealing. To reconcile the different results obtained, the authors invoked the existence of a substantial density of interface states (Nss = 2 x 1012 cm"2) and the way in which they responded to the measurement techniques. In the I/V measurements, recombination-generation centres and enhanced thermionic field emission, due to the rather high doping used, contributed to the low barrier height as attested by the poor value of n (from 1.15 to 1.45). In the C-V measurements, partial compensation of the GaAs near the interface occurred which expanded the depletion layer width and shifted the intercept of the 1/C2 curve with the voltage-axis by an amount AV = (qNssd)/(2e) where d is the depth of the layer containing the defects and Nss their density per cm2. On the other hand, in the photoresponse measurements (|)b is less sensitive to interfacial compounds or oxides at the interface. Pt significantly dissolves gaseous hydrogen. Aspnes et al [5] found that although this reversibly lowered the work function of Pt by 1 eV, the I-V, C-V and photoresponse (|>b values were hardly affected, whereas for other Pt-group metals (Ru, Ir) (f)b was considerably changed. Onuma et al [6] have studied the effect of sintering (4000C) on the DLTS spectra OfPt(IOOO A)/ Ti(500 A)/ Au(2500 A) contacts to 1 x 1017 cm"3 n-GaAs. A deep trap at 0.46 eV below the conduction band minimum was found to be strongly dependent on annealing temperature and decreased from a concentration of 2 x 1015 cm"3 to less than 2 x 1014 cm"3 with an accompanying decrease in n from 1.15 to 1.05 after 10 minutes annealing at 4000C . The stability of the Schottky barrier was correlated with the formation of a buried PtAs2/GaAs interface. The pressure dependence of (j)b at the Pt/GaAs interface has recently been studied by Shan et al [7]. Using a 4 x io 16 cm"3 Si-doped GaAs sample and varying the pressure from 0 to 16.7 kbar, the variation of (j>b with pressure was found to follow an equation of the form : (K(P) = (J>b(0) + aP + bP2

(1)

where a = 11 meV/kbar and b = -0.26 meV/kbar2. The ideality factors of the diodes investigated were high, 1.27 at atmospheric pressure and 1.30 at 18 kbar. The linear pressure coefficient a in EQN (1) was found to be very close to that of the deep defect levels E3 and E4 which are located in the bandgap and generated by electron irradiation [8]. The E3 and E4 levels are believed to be due to the native amphoteric defects VGa (Ga vacancy) which acts as an acceptor and AsGa (As on a Ga vacancy) which acts as a donor. Based on these observations, the authors proposed that the pressure dependence of the Fermi level pinning at the

Pt/GaAs interface was consistent with the amphoteric native defect model put forward recently by Walukiewicz [9]. The above results from the various groups agree as to the extensive reactions between Pt and GaAs. The formation of PtGa near the metal surface is driven by the large electronegativity difference between Pt and Ga. The stability of the Schottky barrier as a function of annealing temperature seems to be associated with the presence OfPtAs2 at the interface. B4

Clean Surfaces

Wu et al [10] have outlined a novel method of fabricating oxide-free Pt/GaAs Schottky barriers using in-situ photopulse-assisted electrochemical processing. Nearly ideal thermionic emission characteristics with a high barrier height (j)^) of 1.07 eV and an ideality factor of n = 1. 05 were observed over a range of 7 orders of magnitude of electric current. The results of atomic force microscopy (AFM), X-ray photoemission spectroscopy (XPS) and deep level transient spectroscopy (DLTS) measurements indicated that the novel electrochemical process produces a smooth and oxide-free interface and prevents the formation of process-induced damage. C

THE W/GaAs INTERFACE

Refractory metals, especially tungsten (W), have attracted much interest because of their high thermal stability compared with noble metals [2]. These metals have fairly small electronegativities and the interfaces they make with III-V semiconductor compounds are relatively stable [I]. Contrary to earlier expectations, W and the other refractory metals (Ta, Re, Ir, Mo) are not chemically inert in contact with GaAs. According to Waldrop et al [12] their reactivity is about the same as for Fe, Cr and Ti and more than for Sn, Ag, Au and Al. Cl

Chemically Vapour Deposited (CVD) W Films

Using the reaction between tungsten hexachloride and hydrogen, Batev et al [13] produced almost ideal W on (111) n-GaAs Schottky diodes . The carrier concentration was about 1016 cm"3 and the ideality factor (n) was 1.02 (± 0.01). The room temperature (j>b was measured as 0.81 (±0.01) eV by both I-V and C-V techniques. C2

Electron Beam Evaporated Contacts

Waldrop et al [14] have studied the effect of annealing oxide-free and oxide-containing interfaces at up to 6500C. 1000 A W films were e-beam deposited on (100) n-GaAs doped to 5 x 1016 cm"3. A native oxide layer (Ga2O3 + As2O3) 10 A thick was removed from some samples by heating to 55O°C to leave an ordered contamination-free surface prior to W deposition. For both kinds of interfaces, ^>b(yW) was considerably lower than fy^c/v) (0.70 eV vs. 1.1 eV) and the ideality factor was > 1.2 in most cases, from room temperature deposition up to 650 0 C annealing for 30 minutes. The lowest ideality factor (1.1) was obtained for diodes annealed between 350 and 450 0 C . Waldrop et al [14] also determined (j>b using X-ray photoelectron spectroscopy (XPS) and found a value of 0.90 eV. The XPS barrier height is obtained by measuring the interface band bending with respect to the Fermi level and is therefore not sensitive to the interfacial compounds, oxides

or annealing factors which affect the electrical characteristics of the Schottky diodes. The difference between the I/V and XPS barrier height was attributed to a high density of interface recombination centres during current transport as supported by the high values of the ideality factor. Likewise, the high (J>b(C/v) values follow from the dielectric properties of the interface. Matsumoto et al [15] have studied the thermal stability of the interface between Se+ implanted (300 keV, 6 x 1012 cm"2) GaAs and W for 10 minute anneals at 800, 850 and 950 0 C . Tungsten patterns (1000 A thick, 130 |im in diameter) were produced by electron beam evaporation and lift off. The ideality factor stayed at 1.1 - 1.2 and (J)b(I/V) was 0.73 eV for all three anneal temperatures. He+ backscattering analysis revealed interdiffusion between the GaAs and W at 900 0 C, which nevertheless caused no deterioration in the Schottky diode characteristics. C3

Sputtered W Contacts

Yu et al [16] have investigated the Schottky barrier degradation mechanism of sputter deposited W films on 1017 cm"3 n-GaAs as a function of annealing temperature from 100 to 900 0 C . The electrical characteristics of the diodes improved after annealing between 350 and 500 0 C . The reverse leakage current decreased by an order of magnitude compared with that of as-deposited contacts, the ideality factor decreased from 1.1 to 1.05 and cf)b increased from 0.58 eV to 0.64 eV, presumably due to the annealing of the sputter damaged GaAs surface. Annealing above 600 0 C resulted in a dramatic increase in the reverse leakage current and an increse in n to 1.14. Further annealing at 7000C increased n to 1.34 and at 8000C and above, ohmic behaviour was observed. The C-V characteristics showed considerable degradation after annealing above 600 0 C, with a substantial decrease in the free electron concentration. This was attributed to in-diffusion of W atoms creating acceptor-type recombination centres and thus compensating the shallow donors. The TEM and RBS measurements showed that at 700 0 C, W was detected to a depth of 500 to 600 A below the W/GaAs interface. Above 8500C annealing, the W overlayer started to react with the GaAs and the reaction product (W2As3) tended to ball up forming islands on the GaAs surface. Somewhat contradictory results were reported by Yosefowicz and Rensch [17] who have shown that low stress, pure W films can be successfully produced using DC magnetron sputtering. These films were found to adhere strongly to the GaAs surface and were thermally stable in a furnace anneal of 8500C for 20 min, or rapid thermal anneal (RTA) at 10000C for 15 s. Auger analysis, RBS analysis and SIMS profiling showed the W films to be chemically stable during high temperature annealing. Pure W was also found to have a significantly lower resistance (1.2 ^Q.cm) compared with WSi2 and TaSi2. Investigation of (J)1^1ZV) showed that furnace annealed diodes (8000C for 15 min) had an average (|)b of 0.70 eV with an excellent uniformity (the standard deviation was 9.8 meV for 300 diodes on a 2-inch wafer). Enhancement mode MESFETs with transconductances in excess of 300 mS/mm for gate lengths of 1 ^m were produced and ring oscillators with a gate delay of 42 ps and power dissipation of 600 ^iW/gate were successfully fabricated. C4

The Use of W Contacts in GaAs Integrated Circuits

The self aligned gate (SAG) process for metal-semiconductor field effect transistor (MESFET) devices has attracted much attention as an alternative to conventional gate recess processing. The

advantage of the SAG process includes low parasitic series resistance but requires the gate metal to withstand high temperature annealing (>800°C ) following n+ implantation. Despite earlier reports of the degradation of W/GaAs Schottky diodes at annealing temperatures greater than 600 0 C 3 it has recently been shown that extremely high stability W films can be deposited on GaAs, with associated excellent MESFET devices. The advantage of using pure W includes the ease of reproducibly sputtering these films with higher purity than the silicides and hence making them serious contenders for GaAs VLSI circuits. The importance of WSiN as a Schottky contact to integrated circuits has been discussed by Zhao [18] who reported thin films on ion-implanted n-GaAs by DC magnetic controlled sputtering. AES and SIMS analysis showed that the WSiN/GaAs interface remained stable after rapid annealing at 10000C for 10 s or high-temperature furnace annealing at 8500C for 20 minutes: the Schottky barrier height and ideality factor n were 0.8 V and 1.1, respectively. WSiN self-aligned gate (SAG) E/D mode MESFETs were fabricated with a transconductance of 154 mS/mm for E-FET and 250 mS/mm for D-FET. A differential amplifier was also fabricated with a gain as high as 29.5 dB from DC to 1 GHz. Lee et al [19] have used an ion beam assisted deposition (IBAD) technique for the growth of refractory W and WNx films as gate materials for GaAs integrated circuits. W was deposited by Ar-ion beam sputtering, and simultaneously low energy nitrogen ions were introduced directly onto the GaAs substrate, using a Kaufman-type ion gun, for nitridation of the W deposits. The deposited films were characterized by electrical resistivity measurements, X-ray diffraction, and X-ray photoelectron spectroscopy analyses. The results showed that the composition and the electrical resistivity could be controlled with the nitrogen-ion dose. The thermal stability of the WNx/GaAs interfaces was also investigated by examining Schottky diode characteristics and microstructures. The WNx contacts investigated in the composition range from x = 0 to 0.47 were electrically stable up to 8500C and showed remarkable enhancements in (|)b through high temperature annealing. Such annealing-induced barrier enhancements were attributed to the dissociation of a thin native oxide layer which was observed between the as-grown film and the GaAs substrate. Comparisons of the results with previous works on conventional sputter-deposited WNx films led to the conclusion that the IBAD technique generates much less ion damage in the substrate and the film, and therefore is promising for growing refractory compounds on GaAs. REFERENCES [1] [2] [3] [4] [5] [6] [7] [8] [9] [10]

A.K. Sinha, J.M. Poate [ in Thin films-interdiffusion and reactions, Ed J.M. Poate, N.K. Tu, J. W. Mayer (Wiley, New York, 1978) ] A.K. Sinha, J.M. Poate [ Appl Phys. Lett. (USA) vol.23 no. 12 p.666 (1973) ] S.P.Murarka [ Solid State Electron. (UK) vol. 17 no.9 (1974) p.985 ] C. Fontaine, T. Okumura, K.N. Tu [ J Appl Phys. (USA) vol.54 no.3 (1984) p. 1404 ] D.E. Aspnes, A. Heller [J Vac. Sci. Technol. B (USA) vol.1 no.3 (1983) p.602 ] T. Onuma, T. Uenoyama, H. Yagita [ Appl. Phys. Lett. (USA) vol.44 no. 11 (1984) p.80 ] W. Shan, M.F. Li, P.Y.Yu, W.L. Hansen, W. Walukiewicz [ Appl. Phys. Lett. (USA) vol.53 no. 11 (1988) p.974] R.H. Wallis, A. Zylberszteijn, J.M. Besson [Appl. Phys. Lett. (USA) vol.38 (1981) p.698 ] W. Walukiewicz [J Vac. Sci. Technol. B (USA) vol.5 (1987) p.974 ] NJ. Wu, T. Hashizume, H. Hasegawa, [ Jpn.J. Appl Phys. (Japan) pt.l vol.33 no. IB (1994) p.936]

[11] [12] [13] [14] [15] [16] [17] [18] [19]

F. Ren, A.B. Emerson, SJ. Pearton, W.S. Hobson,T.R. Fullowan, J. Lothian [ J. Electron. Mater. (USA) vol.20 no.8 (1991) p.595 ] J.R. Waldrop, S.P.Kowalczyk, R.W. Grant [J. Vac. Sci. Technol. (USA) vol.21 no.2 (1982) p.607 ] P.M. Batev, M.D. Ivanovitch, E.I. Kafediiska, S.S. Simenov [ Phys. Status Solidi A (Germany) vol.45 no.2 (1978) p.671] J.R. Waldrop [ Appl. Phys. Lett. (USA) vol.42 no.4 (1982) p.350 ] K. Matsumoto,N. Hashizume, H. Tanoue, T. Kayanama [ Jpn. J. Appl. Phys. (Japan) vol.21 no.6 (1982)p.L393] K.M. Yu, S.K. Cheung, T. Sands, J.M. Jaklevic, N.W. Cheung, E.E. Haller [ J. Appl. Phys. (USA) vol.60 no.9 (1986) p.3235 ] J.Y. Yosefowicz, D. B. Rensch [ J. Vac. Sci. Technol. B (USA) vol. 15 no.6 (1987) p. 1707 ] Zhao Xinjian[ Res. Prog. SSE (China) vol. 1 no.22 (1991) p. 100 ] JS. Lee, CS. Park, J.Y. Kang, D.S. Ma, [J. Vac. Sci. Technol. B (USA) vol.8 no.5 (1990) p.1117]

14.9 Conduction and valence band offsets at the AlGaAs/GaAs and (Al)GaInP/GaAs heterostructure interfaces S. Adachi March 1996

A

INTRODUCTION

One of the most important parameters for the design and analysis of heterojunction and quantumwell electronic and optoelectronic devices is the heterojunction band offset [I]. The band offset is a consequence of the difference between the energy gaps of two semiconductors. The energygap difference is distributed between a conduction-band offset AEC and a valence-band offset AEV. In a type I (straddling lineup), we obtain AEg = AEc + AEv

(1)

while for type II (broken-gap and staggered lineups) the relation is given by AE 8 =IAE 0 -AEJ

(2)

where AEg is the bandgap difference. A large amount of experimental and theoretical work has been accumulated to address the heterojunction band offsets and related phenomena in the AlGaAs/GaAs heterojunction system [I]. The band offsets in (Al)GaInIVGaAs heterojunctions have also been studied by many groups. It has been found that these heterojunction systems belong to type I. We will discuss in Section B some results of the AlGaAs/GaAs heterojunction system. Results of the (Al)GaInIVGaAs heterojunction system are presented in Section C. The external perturbation effect is discussed in Section D. Finally, the conclusions obtained from this contribution are briefly summarized in Section E. B

AlGaAs/GaAs HETEROJUNCTION SYSTEM

Bl

Band Offsets Obtained from Optical Techniques

The measurement techniques used for determining AEC and AEV can be broadly divided into two distinct categories, namely, optical and electrical. Optical techniques are based mainly on the study of the optical properties of alternating layers of two semiconductors. In TABLE 1 we summarize the band-offset ratio AEC:AEV for AlxGa1^xAsZGaAs obtained from optical and photoelectric techniques [2-20].

TABLE 1. Band-offset ratio AEC:AEV for AlxGa^xAsZGaAs heterojunctions obtained from optical and photoelectric techniques. OA= optical absorption, XRP=X-ray photoemission, PL=photoluminescence, PC=photocurrent, LS=light scattering, ER=electroreflectance, and EL=electroluminescence. AEC:AEV

Technique

88^2

OA

1974

[2]

85^15

OA

1975

[3]

90:10

XRP

1981

[4]

70:30

XRP

1981

[4]

51:49

PL

1984

[5]

57^43

PL

1984

[6]

60:40

XRP

1985

[7]

62:38

PC

1985

[8]

(57-66):(43-34)

PC

1985

[9]

65:35

PL

1985

[10]

50:50

XRP

1986

QJJ

75:25

PL

1986

[12]

69:31

LS

1986

[13]

68:32

PL

1986

[14]

63:37

PC

1987

[15]

77:23

ER

1987

Q6]

66:34

PL

1992

[17]

62:38

PC

1993

[18]

63:37

PL

1994

[19]

69:31

I

EL

Year

|

1995

Ref

|

[20]

Optical absorption was the first technique to be applied, and the pioneering work of Dingle et al [2,3] using optical absorption measurements on multiple-quantum-well structures gave the value of AEC:AEV as 85:15. Although this value has become widely accepted, heterojunction studies by several authors have required values of AEC in the range of (0.62 - 0.66)AE g to explain their electrical and photoelectric data [21-23]. The discrepancy with Dingle's result was attributed to compositional grading or interface states. However, in 1984 Miller et al [6] reported values of 51:49 and 57:43 using photoluminescence from parabolic and rectangular wells. Duggan et al [10] have also tried essentially the same approach and arrived at 65:35. The conduction- and valence-band offset data for AlxGa1-xAs/GaAs heterojunctions published before 1992 have been compiled by Missous [24]. From the data available at that time, Missous concluded that AE 0 = (0.80 ± 0.03)x eV for 0 < x < 0.45 and

AEV= (0.51 ± 0.04)x eV for 0 < x < 1.0 (AEC:AEV= 61:39). For x > 0.45, AEC decreases steadily [24-26]. Unique optical and photoelectric approaches to the band-offset problem are the use of light scattering [13] and electroreflectance methods [16]. The basic idea for the use of light scattering is to measure the intersubband transitions in quantum wells from inelastic light-scattering spectra. The method has several appealing features such as the possibility of applying it selectively to the conduction band, where the calculation of energy levels is simpler, and the capability to determine the bandgap of the barrier material by resonant Raman scattering. Menendez et al [13] have determined the offset ratio 69:31 with inelastic light scattering in AlxGa1^AsZGaAs quantum wells. Raccah et al [16] have deduced the offset ratio 77:23 from the valence and conduction subband energies in an AlxGa1^AsZGaAs superlattice measured by electrolyte electroreflectance. Forchhammer et al [20] have also obtained the ratio 69:31 for x < 0.4 from electroluminescence measurements. Above x~0.4, Qc = AEcZAEg decreases steadily with increasing x. This corresponds to the fact that AEV increases steadily for 0 < x < 1.0 while AEC decreases with increasing x for x > 0.4. B2

Band Offsets Obtained from Electrical Techniques

We summarize in TABLE 2 the band-offset ratio AE0)AEy for the AlxGa1^AsZGaAs heterojunction obtained from electrical techniques [21,26-39]. Kroemer et al [21] proposed a new method of determining band offsets from the carrier profiles of heterojunctions by capacitance-voltage (C-V) measurements. This method is based on reverse-biased capacitance measurements for a diode structure. Since the precision of capacitance measurements is relatively high and also because no adjustable parameter is required in the determining procedure, this method is thought to be suitable for the precise determination of band offsets. As listed in TABLE 2, several authors have used the C-V method to determine the band offsets in AlxGa1^AsZGaAs heterojunctions [21,2932,34]. These works indicate that the band-offset ratio AEC:AEV is in the range (62 - 70):(38 30). Transport data also suggest the offset ratio (56 - 65):(44 - 35) (see TABLE 2). Thus, the conduction-band offsets, 56 - 70%, obtained from the electrical techniques are considerably lower than the old, widely accepted Dingle's value of 85%, but are consistent with those obtained from optical and photoelectric techniques listed in TABLE 1. B3

Type-II Alignment in Narrow Quantum Wells

In a type-I quantum well, the electrons and holes are confined spatially in the same layer. It has, however, been well established that type-II behaviour can be observed in AlAs/GaAs heterostructures [26]. The crossover between type-I and type-II behaviours occurs when the GaAs well width falls below a critical value, which is about 35 A for wide AlAs barriers. The lowest electron and highest hole states in type-I quantum-well structures with wide GaAs wells are at the GaAs F point. However, the lowest F electron shifts to higher energy with a decrease in the well thickness Lw and for Lw<35 A the lowest electron states are the X levels in the AlAs barriers. Thus, in the resulting type-II quantum-well structures the electron-hole transitions can be expected to be indirect in both momentum and real space.

TABLE 2. Band-offset ratio AEC:AEV for the AlxGa ^s/GaAs heterojunction obtained from electrical techniques. CV=C-V profiling, HE=Hall-effect measurement, IV=I-V measurement, eB=e-beam-induced current, DLTS=deep level transient spectroscopy, STM=scanning tunnelling spectroscopy, and AS=admittance spectroscopy. AEC:AEV

Technique

Year

66:34

CV

1980

[21]

62:38

HE

1984

[27]

65:35

IV

1984

[28]

67:33

CV

1985

[29]

65:35

CV, IV

1985

[30]

63:37

CV, IV

1985

[31]

62:38

CV

1985

[32]

(60-65):(40-35)

IV

1986

[33]

70^30

CV

[986

[34]

56:44

eB

1987

[35]

61:39

IV

[987

[36]

63^37

DLTS

1993

[37]

56:44

STM

1993

[38]

60:40

AS

_ J

1994

Ref

|

[39]

Studies on type-II systems have been extended to include purely binary structures of AlAs/GaAs(001) [26]. In these reports, photoluminescence (PL) and photoluminescence excitation (PLE) spectroscopy have been used extensively to identify not only the pure GaAs F transitions (type I) but also transitions related to the type-II band alignment. The main conclusion obtained by Finkman et al [40] was that the type-II exciton emission is an indirect transition involving n = 1 heavy holes in the GaAs layer and electrons in the X10^ valleys of the AlAs layer, i.e., those electrons with momenta in the layer planes, along the [100] and [010] directions. This assignment of the type-II emission as an indirect process was based on fitting the PL decay curve, which was nonexponential, to an expression for the time dependence of the decay of localized zone-boundary excitons made after allowing for scattering from a random interface potential. Support for this interpretation was provided by Ihm [41] who predicted that indeed X10^ can be the lowest minima in some circumstances. Dawson et al [42] have reported similar features in the PL and PLE spectra to those reported by Finkman et al [40] from samples consisting of 60 periods of 70 A of AlAs with either 28 or 22 A of GaAs. They, however, presented a different interpretation of the emission data. They ascribed the type-II process as zero-phonon and phonon-assisted recombination of excitons involving n = 1 heavy holes in the GaAs layer and Xlcz electrons in the AlAs layer, i.e., those electrons with momentum parallel to the growth direction, along [001]. This has the important consequence that because only the Xlcz state is mixed with F by the superlattice potential, the type-II emission process is thought to be a pseudodirect transition and not an indirect transition arising from the unmixed X10^ states.

More recently, Minami et al [43], who have studied the time dependence of the type-II emission, have questioned the analysis of Finkman et al [40]. They concluded that the nonexponential character of the decay could be explained by the recombination of excitons involving Xlcz electrons, in agreement with the conclusion of Dawson et al [42]. Moore et al [44] have also reported the results of a systematic investigation of the effects of electronic coupling of the GaAs F states on the band alignment of type-II AlAs/GaAs multiple-quantum-well structures. They studied a series of samples in which the thickness of the GaAs layers was fixed at -25 A and the AlAs thickness varied between samples from 41 to 5 A. Their results presented further evidence to support Dawson's original assignment [42] of the type-II emission as a pseudodirect process involving electrons at the Xlcz minimum. C

(Al)GaInP/GaAs HETEROJUNCTION SYSTEM

Cl

Ga052In048PZGaAs

The AlxGaxIn1^yIVGaAs heterojunction system has a broader interest since it is used for fabricating light emitting and laser diodes in the visible-wavelength region below 700 nm. The lattice-matching relation between the composition fractions x and y for AlxGaxIn1^7P latticematched to GaAs can be written as y = 0.5160 - 0.9729x (0 < x < 0.53)

(3)

EQN (3) gives an end-point alloy Ga052In048P (x = 0, y = 0.52). The band offsets in the Ga0 S2In0 48P/GaAs heterojunction have been studied both experimentally and theoretically. TABLE 3 summarizes the experimental band offsets in this heterojunction system obtained by different authors [45-53]. TABLE 3. Experimental band offsets (AE0, AEV) and offset ratio AEC: AEV in the Ga052In048PZGaAs heterojunction. CV = C-V profiling, IV = I-V measurement, DLTS = deep level transient spectroscopy, IP = internal photoemission, PL = photoluminescence and PR = photoreflectance.

AEC (me V)

AEV (meV)

AEC:AEV

Technique

Ref(Year)

220 ± 1 0

240 ± 1 0

48:52

CV

[45] (1987)

IV

[46] (1989)

-30 198

285

41:59

DLTS

[47] (1990)

108 ± 6

353 ±11

23:77

IP

[48] (1991)

60 ±20

400 ± 20

(13±4):(87±4)

PL

[49] (1991)

IV

[50] (1992)

29:71

PR

[51] (1992)

63^37

PL

[52] (1993)

210±10 130

320

I

330 ±20

1

I

PL

I

[53] (1993)

A theoretical value of the conduction-band offset AEC in the Ga 0 S2In048 P/GaAs heterojunction evaluated by Harrison was 160 meV [54], while the measured AEC value for this heterojunction

system ranges widely from -30 to 220 meV (TABLE 3). Foulon et al [55] performed a tightbinding calculation of the band lineups at the Ga052In048PZGaAs interfaces. They obtained a value ofAEc that varies between 70 and 140 meV3 depending on the detailed interface atomic structures (As-In-P-(Ga3In), Ga-As-(Ga5In), or Ga-P-(GaJn)). The C-V profiling method yielded AEC = 220 meV (AE v = 240 meV) [45], deep level transient spectroscopy on quantum wells yielded AE0= 198 meV (AEy= 285 meV) [47], and photoemission measurements resulted in AE0 =108 meV (AEy= 353 meV) [48]. On the other hand, I-V measurements on a heterojunction bipolar transistor suggested only -30 meV OfAE0 [46] and PL on a type-II multiple quantum well as a function of hydrostatic pressure indicated a value OfAE0 = 60 meV [49]. More recent study on high-pressure PL of type-II single quantum wells deduced a valence-band offset of AEy= 330 meV [53]. Assuming AEg= 460 meV, this valence-band offset gives AE 0 = 130 meV in reasonable agreement with photoemission [48] and photoreflectance measurements [51]. The conduction-band offset in the Ga0 5In0 5P/A1O^3Ga0 57As heterojunction has also been studied by Lee et al [56] using a C-V profiling method. They obtained a staggered band line-up with both bands OfGa05In05P above those OfAl043Ga057As (AEO=157 meV). The conduction-band offset in the Ga0 5In05PAn ^xGaxAs5P1-0, heterojunction system has also been determined by the same group [57] to be about 18% of the band-gap difference AEg. C2

AlxGayInlx.vP/Gao 52In048P

Studies of the band offsets in the (AlxGa1-J052In048PZGa052In048P heterojunction lattice-matched to GaAs have been studied by several groups [58-62]. We summarize these results in TABLE 4. TABLE 4. Experimental band offsets (AEC, AEV) and offset ratio AEC:AEV in the (Al x Ga 1 J 0 5 2 In 0 4 8 P / Ga 052 In 048 P heterojunction system lattice-matched to GaAs. C-V = C-V profiling, PLE = photoluminescence excitation spectroscopy, PL = photoluminescence and IP = internal photoemission.

AEC, AEV (eV) AEC: AEV:

0.27 x (0< x< 0.7) 0.35 - 0.23 x (0.7 < x < 1.0) 0.32 x ( 0 < x ± 1.0)

AEC:AEV

Technique

Ref(Year)

46:54 (0 < x ^ 0.7)

CV

[58] (1987)

65:35 (x = 0.6)

PLE

[591(1990)

67:33 (x = 0.7)

PL, PLE

[60] (1993)

AEC AEV

=0.273 ± 0 . 0 2 = 0 . 1 4 8 ± 0 . 0 2 (x = 0.58)

65:35 (x = 0.58)

PL

[61] (1995)

AEC:

0.295 x (0< x< 0.6) 0 . 1 1 5 - 0 . 0 9 x (0.6 < x < 1.0) 0.305 x (0 <; x < 1.0)

49:51 (0^x^0.6)

IP

[62] (1995)

AEV:

|

|

|

Two experiments suggest that the conduction-band offset AEC increases linearly with x for the direct gap (AlxGa1-Jj)052In048P (x < 0.7 [58] or x < 0.6 [62]) and decreases linearly for indirect gap (AlxGa1^)0 S2In048P (x > 0.7 [58] or x > 0.6 [62]), whereas the valence-band discontinuity AEV increases linearly with x for 0 < x < 1.0 [58,62]. PL and PLE measurements provide the band-

offset ratio AEC:AEV~65:35 for the direct gap (AlxGa1J052In048P [59-61]. The band offsets for (AIxGa1J0 J2In0 ^PZGayln^yP as a function of In composition y have been studied by Dawson et al using PLE [63,64]. A series of tensile and compressively strained GaxIn1^yP quantum wells have been grown on GaAs substrates by metal-organic vapour phase epitaxy. PLE spectra obtained from these samples (y = 0.4 - 0.65) indicated the conduction-band offset of about 0.7AEg. C3

Gao^Ino^gP/AlosoInosoI1

Patel et al [65] determined the band lineups at the Ga052In048PZAl05In05P interface. They performed low-temperature PL measurements at high pressure (up to 40 GPa) and obtained values ofAE c = 0.26 eV, AEV= 0.24 ± 0.05 eV, and AEC:AEV= 52:48. Determination of these values, however, required a literature value (2.45 eV) of the lowest-direct-gap energy of the barrier material. Recently, Dawson et al [66] have presented evidence that the lowest direct gap of the Al0 5Ga0 sP barrier at low-temperature (5 K) is -2.685 eV, which resulted in a revised bandoffset ratio of 67:33. D

EXTERNAL PERTURBATION EFFECT

The band offset is principally not dependent on doping concentration. Since the temperature variations of the bandgap energies are very similar among the III-V materials, the band offsets and resultant offset ratios in the III-V heterojunction system can be successfully assumed to be independent of temperature. The orientation dependence of the band offsets in AlxGa1^AsZGaAs heterojunctions were studied both theoretically and experimentally [26]. All these studies indicated that the band offsets are independent of crystal orientation. Also, there are different opinions about AlxGa1^AsZGaAs heterojunctions as to whether the band offset is commutative (i.e., independent of growth sequence) or not [26]. The possibility of band-offset control in semiconductor heterojunctions has been studied recently, both theoretically and experimentally [67]. The interest in this kind of study is well justified because of the importance of these band discontinuities in many solid-state devices for which the band offsets are the key design parameter, mainly because the efficiencies and properties in such devices depend strongly on the potential barrier at the heterojunction interface. E

CONCLUSION

We have reviewed here the band offsets in AlxGa1^AsZGaAs and (AlxGa1-X)052In048PZGaAs heterojunction systems. A flood of new results for AlxGa1^AsZGaAs heterojunctions, published since 1984, has definitely proved that the old value AEC:AEV= 85:15 is found to be no longer valid. The mass of data accumulated recently using the various experimental techniques clearly shows that the band-offset ratio converges to a (60 - 65):(40 - 35) value rather than the old one (85:15). The band offsets in the (AlxGa1^x)O52In048PZGaAs heterojunction system vary widely according to different authors. REFERENCES [1]

See, for example, F. Capasso, G. Margaritondo (Eds) [ Heterojunction Band Discontinuities: Physics and Device Applications (North-Holland, Amsterdam, 1987) ]

[2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13]

[14] [15] [16] [17] [18] [19] [20] [21]

[22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33] [34]

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[35] [36] [37] [38] [39] [40] [41] [42] [43] [44] [45] [46] [47] [48] [49] [50] [51] [52] [53] [54] [55] [56] [57] [58] [59] [60] [61] [62] [63] [64] [65] [66] [67]

A. Eisenbeiss, H. Heinrich, J. Opschoor, R.P. Tijburg, H. Preier [Appl. Phys. Lett. (USA) vol.50 (1987) p. 1583] J. Smoliner, R. Christanell, M. Hauser, E. Gornik, G. Weimann, K. Ploog [ Appl. Phys. Lett. (USA) vol.50 (1987) p. 1727] QS. Zhu, S.M. Mou, X.C. Zhou, Z.T. Zhong [ Appl. Phys. Lett. (USA) vol.62 (1993) p.2813] S. Gwo, K-J. Chao, CK. Shih, K. Sadra, B.G. Streetman [Phys. Rev. Lett. (USA) vol.71 (1993) p. 1883] S.R. Smith, F. Szmulowicz, GJ. Brown [ J. Appl. Phys. (USA) vol.75 (1994) p. 1010 ] E. Finkman, M.D. Sturge, M.C. Tamargo, [ Appl. Phys. Lett. (USA) vol.49 (1986) p. 1299 ] J. Dim [Appl. Phys. Lett. (USA) vol.50 (1987) p. 1068 ] P. Dawson, K. J. Moore, C. T. Foxon [ Proc. SPIE (USA) vol.792 (1987) p.208 ] F. Minami, K. Hirata, K. Era, T. Yao, Y. Masumoto [ Phys. Rev. B (USA) vol.36 (1987) p.2875 ] KJ. Moore, P.Dawson, CT. Foxon [ Phys. Rev. B (USA) vol.38 (1988) p.3368 ] M.A. Rao, EJ. Caine, H. Kroemer, S.I. Long, D.I. Babic [ J. Appl. Phys. (USA) vol.61 (1987) p.643 ] E.Y. Lee, LJ. Schowalter [ J. Appl. Phys. (USA) vol.65 (1989) p.4903 ] D. Biswas, N. Debbar, P. Bhattacharya, M. Razeghi, M. Defour, F. Omnes [ Appl. Phys. Lett. (USA) vol.56 (1990) p.833 ] M.A. Haase, MJ. Hafich, G.Y. Robinson [Appl. Phys. Lett. (USA) vol.58 (1991) p.616 ] J. Chen, JR. Sites,I.L. Spain, MJ. Hafich, G.Y. Robinson [ Appl. Phys. Lett. (USA) vol.58 (1991) p.744 ] T. W. Lee et al [ Appl. Phys. Lett. (USA) vol.60 (1992) p.474 ] G. Amaud, P. Boring, B. Gil, J.-C. Garcia, J.-P. Landesman, M. Leroux [ Phys. Rev. B (USA) vol.46 (1992) p. 1886] B.S. Jeong et al [ Solid State Commun (UK) vol.86 (1993) p.373 ] M. Leroux, M.L. Fille, B. Gil, J.P. Landesman, J.C. Garcia [ Phys. Rev. B (USA) vol.47 (1993) p.6465 ] W.A. Harrison [ J. Vac. Sci. Technol. (USA) vol. 14 (1977) p. 1016 ] Y. Foulon, C Priester, G. Allan, J.C. Garcia, J.P. Landesman [ J. Vac. Sci. Technol. B (USA) vol. 10 (1992) p. 1754] J.B. Lee, K.-S. Kim, B.-D. Choe [ Appl. Phys. Lett. (USA) vol.62 (1993) p.2688 ] Y-H. Cho, K.-S. Kim, S.-W. Ryu, S.-K. Kim, B.-D. Choe [ Appl. Phys. Lett. (USA) vol.66 (1995) p. 1785] K.O. Watanabe, Y. Ohba [ Appl. Phys. Lett. (USA) vol.50 (1987) p.906 ] CT.H.F. Liedenbaum, A. Valster, A.L.GJ. Severens, GW. 'tHooft [ Appl. Phys. Lett. (USA) vol.57 (1990) p.2698 ] M.D. Dawson, G. Duggan [ Phys. Rev. B (USA) vol.47 (1993) p. 12598 ] O.P. Kowalski, J.W. Cockburn, DJ. Mowbray, M.S. Skolnick, R. Teissier, M. Hopkinson [ Appl. Phys. Lett. (USA) vol.66 (1995) p.619 ] H.K. Yow, P.A. Houston, M. Hopkinson [ Appl. Phys. Lett. (USA) vol.66 (1995) p.2852 ] M.D. Dawson, G. Duggan [ Appl. Phys. Lett. (USA) vol.64 (1994) p.892 ] M.D. Dawson, G. Duggan [ Phys. Rev. B (USA) vol.51 (1995) p. 17660; see also erratum: vol.52 (1995) p. 16940] D. Patel, MJ. Hafich, G.Y. Robinson, CS. Menoni [ Phys. Rev. B (USA) vol.48 (1993) p. 18031 ] M.D. Dawson et al [ Phys. Rev. B (USA) vol.50 (1994) p. 11190 ] A. Munoz, P. Rodriguez-Hernandez [ Phys. Rev. B (USA) vol.45 (1992) p.4502, and references cited therein]

CHAPTER 15 BULK GROWTH OF GaAs 15.1 15.2 15.3 15.4 15.5 15.6 15.7

Thermodynamics of the Ga/As system LEC growth of GaAs Horizontal and vertical Bridgman growth of GaAs New types of LEC growth and vapour controlled LEC GaAs Heat treatments of GaAs ingots Heat treatments of GaAs wafers Carrier concentrations of semi-insulating GaAs

15.1 Thermodynamics of the Ga/As system D.TJ. Hurle December 1995

A

PHASE EQUILIBRIA

A major re-evaluation of the thermodynamic data for solid-liquid-vapour equilibria in the Ga/As system has recently been performed [1] following the discovery that the total arsenic pressure in equilibrium with GaAs at its congruent melting point is ~2 atm instead of ~1 atm, assumed previously. Wenzl et al [1] found that they had to apply an external pressure of inert gas of least 2 atm in order to prevent arsenic gas bubbles from forming and growing in the boric oxide encapsulant during LEC growth of GaAs. A similar result was reported in the Soviet literature many years ago [2] but was apparently ignored. Interpreting this information as implying that the equilibrium arsenic pressure over a congruent melt is 2 atm and using recently optimised thermodynamic data for arsenic [3], Wenzl et al obtained a best fit to published data on vapour pressure equilibrated to the Ga-As liquidus and liquidus composition [4]. This is summarised in TABLES 1 and 2. TABLE 1. Optimised thermodynamic data for GaAs. Quantity

Value

Reference

Formation enthalpy

AH° (298) = -87.70 kJ/mol

[22]

Formation entropy

S ° (298) = 64.18 J/mol/K

[24]

Melting point

TM = 1513 K

[1]

Entropy of fusion

ASF° = 70.0 J/mol/K

[23]

Enthalpy of fusion

AHF° = TM SF° = 105.9 kJ/mol

The liquidus is known to exhibit regular solution behaviour [5] with a temperature dependent interaction parameter. Arsenic dimers and tetramers are supposed to form ideal gaseous mixtures. Thus: «Ga.As = a - bT = -[RTln{4X As (1 - X J } + AS F (T 11 - T)] / 2(0.5 - X J 2 = [RTIn {(p^/p*^) 1 7 2 /XM} ] / (1 - X J 2

(1)

where T and T M are the actual and the congruent melting temperature, respectively, AS F is the entropy of fusion, X ^ the atom fraction of As in the melt, pM2 the arsenic dimer partial pressure in equilibrium with the liquidus and P + ^ 2 the partial pressure in equilibrium with pure liquid arsenic at temperature T. R is the gas constant. The dimer/tetramer equilibrium is given by: KA 8 = PA84-PA82"2

with the total arsenic pressure being p T = P ^ 2 + P ^ 4

(2)

TABLE 2. P-T-X data for the Ga-As liquidus [I]. T(K)

X^

P^2 (atm)

P ^ (atm)

P0* (atm)

P(atm)

800

0.0003

3.8 x 10'12

1.1 x 10"16

3.7 x 10"13

4.2 x 10'12

850

0.0007

7.1 x 10-11

5.6 x 10 1 5

4.6 x 10'12

7.5 x 10'11

900

0.0018

9.5 x 10 1 0

1.8 x 10'13

4.3 x 10"n

9.9 x 10"10

950

0.0042

9.6 x 10"9

3.9 x 10 1 2

3.2 x 10"10

10 x 10"9

1000

0.0086

7.8 x 10"8

6.4 x 10"11

1.9 x 1 0 9

8.0 x 1 0 8

1050

0.0165

5.3 xlO"7

8.IxIO 1 0

9.7 x 10"9

5.4 x 10"7

1100

0.0291

3.0 x 10"6

8.5 x 10'9

4.2 x 10'8

3.1 x 10"6

1150

0.0476

1.5 x 10"5

7.5 x 10"8

1.6 x 10"7

1.5 x 10 5

1200

0.0727

6.8 x 10"5

5.8 x 1 0 7

5.4 x 1 0 7

6.9 x 10"5

1250

0.1044

2.8 x 10'4

4.1 x 10'6

1.6 x 10'6

2.9 x 10"4

1300

0.1427

0.0010

2.8 x 10'5

4.3 x 10"6

0.0011

1350

0.1879

0.0041

1.8 x 10"4

1.0 x 10"5

0.0043

1400

0.2415

0.015

0.0013

2.2 x 10"5

0.0166

1450

0.3086

0.061

0.010

4.2 x 10"5

0.071

1500

0.4139

0.331

0.166

6.3 x 10"5

0.497

1513

0.5015

0.920

1.110

5.2 x 10"5

2.03

1500

0.5861

1.80

4.90

2.7 x 10"5

6.70

1450

0.6914

2.66

19.8

6.4 x 10"6

22.44

1400

0.7585

2.62

36.8

1.7 x 10"6

39.37

Preferred values for the thermodynamic constants of arsenic have been obtained by Gokcen [3]: In P+As2172 = 6.982 - 8489/T

(3)

Using this value, Wenzl et al obtained their best fit by adopting a value of 1513 K for the congruent melting point of GaAs which is 2 K above the usually accepted value. The optimised relation for K^ was found to be: InK^ = AH^/RT - AS^/R = 2.648 x 107 T - 17.24

(4)

This differs very slightly from Gokcen's preferred value but is in excellent agreement with Rau's vapour density measurements [6]. Since the publication of [2], Matthiesen et al [7] have reported a direct experimental determination of the total arsenic pressure over a stoichiometric melt at 1533 K, obtaining a value of 2.2 atm. Using the above mass action constants, this converts to a total pressure of 2.15 atm

Temperature CC)

at the congruent melting point which is slightly higher than the value employed by Wenzl et al. It may be therefore, that his calculated partial pressures in the vicinity of the melting point are a little low. The solid liquid phase diagram is shown in FIGURE 1.

Solution +GaAs

Solution+ Ga As

Atomic percent of arsenic

FIGURE 1. Phase diagram of GaAs. (After Koster and Thoma [14]).

B

NATIVE POINT DEFECTS AND THE PHASE EXTENT

Crystals grown under Ga-rich conditions will have a slightly different composition from those grown, at the same temperature, under As-rich conditions. This deviation from the 1:1 stoichiometric composition can be expressed in terms of concentrations of native point defects in the crystal. They are vacancies, interstitials and antisite defects on both sublattices denoted by VGa, VAS> Ga1, ASJ, As Ga , Ga^. These isolated defects can combine to form complexes as the crystal cools (e.g. to form the Schottky divacancy V ^ - VJ and they can recombine (e.g. As1 + VM = ASA8). In a pure crystal the deviation from stoichiometry (6) can be expressed as the difference between the excess numbers of atoms on each sublattice, viz: 6

=

SA8

= [AsJ - [VAS] - [ G a J + [AsGa]

SGa

= [Ga1] - [VGa] - [AsGa] + [ G a J

&As " ^Ga

(5)

where:

Several groups of workers [8,9] have made direct measurements of 5 by titrating crystals grown from melts of differing arsenic concentration X^ . These measurements show that crystals grown from melts with X^ = 0.5 are As-rich and therefore the congruent point must be on the As-rich

side. Hence a Ga-rich melt is required to grow a stoichiometric crystal. The maximum deviation from stoichiometry corresponds to point defect concentrations in excess of 1019 cm'3. The anti site concentrations are orders of magnitude lower than this: the maximum obtainable arsenic antisite concentration, [EL2], is only of the order of 1016 cm'3 (see Datareview 10.4) and most of this can be removed by annealing at high temperature. We can therefore neglect antisites in our consideration of the stoichiometry of melt-grown material. Bublik et al [10] determined the excess mass per unit cell (N) by accurately measuring crystal density (o) and lattice parameter fa) and comparing the ratio o / ao3, which is proportional to the mass per unit cell, with its value for an ideal perfect crystal with no point defects. Neglecting antisites, N is given by: N=[AsJ + [Ga1] - [ V J

- [V 0 J

(6)

Bublik et al found that N increased monotonically as X^ was increased which implies that non-stoichiometry is dominated by defects on the As-sublattice; i.e. by As1 and VM. Comparing EQNs (5) and (6) above, it is seen that 5 = N if defects on the Ga-sublattice can be neglected and that any difference between 5 and N is a measure of non-stoichiometry on the Ga-sublattice. Bublik et al interpreted their data as showing that a Ga-rich crystal resulted from growth from a melt of XM = 0.5 contrary to the titration results referred to above. However, this was because they deduced X^ from a knowledge of the temperature of an arsenic reservoir which controlled the melt composition. In doing this they used the erroneous pre-Wenzl/Gokcen data for the vapour pressure at the melting point of GaAs. When their data is corrected for this error it can be shown [11] to be in very good accord with the titration data of Oda et al [9]. The native point defects can be charged and may exist in more than one charge state. Ionisation energies of the isolated arsenic vacancy have been obtained experimentally from positron annihilation studies [12]. They show that the vacancy can act as a relatively shallow donor with an ionisation energy 0.140 eV below the conduction band but can also exist in a negatively charged state located only 0.03 eV below the conduction band edge. Using DLTS, Feng et al [13] have recently measured values of 0.13 eV and 0.037 eV, respectively, in excellent accord with the positron annihilation data. Relatively large concentrations of arsenic interstitials are known to remain in solution as melt-grown crystals are cooled to room temperature [14,15], yet electrical effects which can be ascribed to native point defects are present only in concentrations up to around 1016 cm'3 (the EL2 concentration in As-rich material). One must therefore assume that As1 is neutral in undoped material. Experimental information on charge states of native point defects on the Ga-sublattice is more sparse and comes only from self and dopant diffusion studies [16] (see Chapter 11 in this book). In undoped and n-type material Ga self diffusion is mediated by the triply negatively charged Ga vacancy VGa3" whilst in p+ material it is controlled by GaI2+. By fitting the data of Oda and Bublik to EQNs (5) and (6), Hurle [11] has estimated native point defect concentrations at the melting point. By additionally fitting electrical data obtained from crystals grown at lower temperatures, he has obtained values for the enthalpies and entropies of

the following native point defect incorporation reactions: ViAs2(R) = As1:

K ^ = [As10] PAS2 1/2

(7)

A SAS

= V^ + As1:

KFA = [As10] [V^0]

(8)

A SAS

+ VGa = As1:

K 0 ^ = [As10] / [VGa°]

(9)

where the subscripted zero indicates the neutral charge state. Values are given in TABLE 3. TABLE 3. Mass action constants for neutral vacancies and interstitials (-RT In K4 = AH1 - TAS1: 1 eV = 1.602.10"22 kJ). Arsenic interstitial (EQN (7))

-RT In K ^

= 150 - 0.040 T kJ/mol

Arsenic Frenkel pair (EQN (8))

-RT In KFA

= 332 - 0.077 T kJ/mol

Gallium vacancy (EQN (9))

-RT In K 0 ^ = 75.2 - 0.150 T kJ/mol

Note. This data is provisional only; it will be updated in a forthcoming review [26].

The phase extent of GaAs is obtained using EQN (5) and is plotted in FIGURE 2.

FIGURE 2. The GaAs solidus. The arrow marks the congruent melting point [H].

C

DOPANT SOLUBILITY

The equilibrium incorporation of a group VI donor dopant (e.g. Te) into a crystal growing from a melt or solution is given by: Te1 + VM = Te^ + + e":

KTe(As) = [Te^n 8 1 /[VM°]YT6[Te1]

(10)

where n^ is the electron concentration at the growth temperature and Yie ls the activity coefficient of Te in the melt/solution which, since the Te is present in only dilute solution, can be taken to be a constant at a given temperature. Therefore the solubility of [TeA8+] = Un, the room temperature carrier concentration, will be proportional to the melt concentration [Te1] so long as rigt is constant. However, once [Te^+] is large enough to contribute significantly to n^then the Te^ solubility will become sublinear. The 'knee' in the solubility curve will occur at about [TeA8+] = n^0 where n^0 is the carrier concentration in undoped material. Ignoring charged native point defects for the moment this will be n^0 = % the intrinsic carrier concentration. However, experimentally it is found in liquid phase epitaxial material, that n^0 ~ 100 ^. The long accepted explanation for this discrepancy is to say that growth does not occur under conditions of bulk equilibrium. Rather, it is supposed, there exists a Schottky barrier at the crystal/melt interface [17] formed by a high density of surface states which pin the Fermi level up to bulk carrier concentrations far exceeding % A second problem is that mobility measurements suggest that the donor is partially self-compensated with an acceptor to donor ratio of NA/ND -0.25, increasing at the highest doping levels. This is usually dismissed as due to a failure of mobility theory with the belief that the crystals are actually uncompensated. However, Brozel et al [18], in radio-tracer studies of Sn-doped crystals, showed that the compensation was real. The present author [19] has shown that both the problem of the position of the 'knee' and the self-compensation can be resolved (without invoking a Schottky barrier) when account is taken of the charged arsenic and gallium vacancies in determining the charge-neutrality condition at the growth temperature. In n-type material this is: H81 +

[VA8-]

+ 3[VGa3-] = n, 2Ir^ +

[VA8+]

+ n*

(11)

In liquid phase epitaxy from Ga solution this can be approximated by: H 8 1 - [ V ^ ] + n* The position of the 'knee' is then at Un ~ [TeA8+] ~ [VA8+]. This provides a measure Of[VA8+] at the growth temperature and can be used to obtain the enthalpy and entropy of the formation reaction for the arsenic vacancy (see Section B above). The constant compensation ratio is explained [19] if it is supposed that the acceptor is a complex of the donor and a gallium vacancy: [TeA8 - VGJ". (Doping with Si, Ge or C is more complex because, additionally, SIA8", GGM~ or CM' acceptors can also be formed.) Estimates of the concentration of charged gallium vacancies can be obtained by fitting to dopant

solubility curves for crystals grown under more As-rich conditions. By growing from solutions of Ga and Bi and changing the Bi/Ga ratio, one can obtain growth at a fixed temperature but with variable arsenic activity [20]. At high Bi/Ga ratio, VGa3" charged vacancies exert an influence on solubility through the electro-neutrality condition. The ionisation energies of the charge states of the gallium vacancy have not been experimentally measured. Calculations by Baraff and Schluter [25] give the following values, measured from the valence band: 0.13 Eg, 0.355 Eg and 0.49 Eg where Eg is the energy gap. Using these values, estimates of the enthalpy and entropy of reaction (9) were obtained [11] from the Bi/Ga LPE data [20] and are given in TABLE 3. The equilibrium concentration of triply-charged gallium vacancies increases as the cube of the free electron concentration at any given temperature and arsenic activity and can rise to ~1019 cm'3 in n+ material. Moreover, as such a crystal cools, the equilibrium carrier concentration, determined by a high concentration of chemical donors, will remain approximately constant. This can result in a situation where the equilibrium concentration of VGa3" actually rises slightly as the crystal cools so that, rather than becoming supersaturated with the point defect as is usually the case, the crystal actually becomes undersaturated and seeks ways of generating more VGa3" and/or of reducing the free electron concentration. This produces the complex behaviour of heat-treated n+ crystals [27] and also explains some features of n-type diffusion [21]. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] II1] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23]

H. Wenzel, A. Dahlen, A. Fattah, S. Petersen, K. Mika, D. Hennel [ J. Cryst Growth (Netherlands) vol. 109 (1991) p. 191] J P.M. Shurygin, A.L. Marbakh [ Izv. Akad. Nauk. SSSR. Neorg. Mater. (USSR) vol.6 (1970) p.1816] NA. Gokcen [ Bull. Alloy Phase Diagrams vol. 10 (1989) p. 11 ] M.B. Panish [ J. Cryst. Growth (Netherlands) vol.27 (1974) p.6 ] LJ. Vielandf ActaMetall. (USA) vol.11 (1963)p. 137] H. Rau [J Chem. Thermodyn. (UK) vol.7 (1975) p.27 ] D.H. Matthiesen [ J Cryst. Growth (Netherlands) vol. 137 (1994) p.255 ] K. Terashima [ Semi-Insulating UI-VMaterials 1988 (Adam Hilger, Bristol, 1988) p.413 ] O. Oda et al [ Semicond Sci. Technol. A (UK) vol.7 (1992) p.215 ] V.T. Bublik et al [ Sov. Phys.-Crystallogr. (USA) vol. 18 (1973) p.2183 ] D.T.J. Hurle [Mater. Sci. Forum (Switzerland) vol. 196-201 (1995) p. 179 ] K. Saarinen, P. Hautojarvi, P. Lanki, C. Corbel [Phys. Rev. B (USA) vol.44 (1991) p.10587 ] S.L. Feng, J. Zhou, L.W. Lu [ Appl. Phys. Lett. (USA) vol.66 (1995) p.2256 ] W. Koster, B . Thoma [ Z. Met. led (Germany) vol.46 (1955) p .291 ] L. Charniy, V Bublik [ J Cryst. Growth (Netherlands) vol. 135 (1994) p.302 ] T.Y. Tan, U. Gosele, S . Yu [ Crit. Rev. Solid State Mater. Sci. (USA) vol. 17 no. 1 (1991) p.473 ] K. Lehovec [ Surf. Sci. (Netherlands) vol. 1 (1964) p . 165 ] MR. Brozel, EJ. Foulkes, LR Grant, D.TJ. Hurle [J. Cryst. Growth (Netherlands) vol.80 (1987) p.323] D.TJ. Hurle [J Phys. Chem. Solids (UK) vol.40 (1979) p.613 ] NA. Yakusheva, V.G. Pogadaev [ J Cryst. Growth (Netherlands) vol. 123 (1992) p. 143 ] T.Y. Tan, H.-M. You, UM. Gosele [Appl Phys. A (USA) vol.56 (1993) p.249 ] N.N. Sirota [ Semicond. Semimet. Eds R.K. Willardson, A.C. Beer, vol.4 (Academic Press, New York and London, 1968) p.35 ] B.D. Lichter, P. Sommelet [ Trans Metall. Soc. AME (USA) vol.245 (1969) p.1021 ]

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M. Tmar, A. Gabriel, C. Chatillon, I. Ansara [ J. Cryst. Growth (Netherlands) vol.69 (1984) p.4213] G.A. Baraff, M. Schliiter [ Phys. Rev. Lett. (USA) vol.55 (1985) p. 1327 ] D.T.J. Hurle [Appl. Phys. Lett. (USA) (to be published, 1997) ] PS. Dobson et al [Inst. Phys. Conf. Ser. (UK) vol.45 (1979) p. 163 ]

15.2 LEC growth of GaAs LR. Grant August 1995

A

INTRODUCTION

GaAs compound dissociates at high temperature by arsenic evaporation loss. The total vapour pressure of arsenic, in the forms As2 and As4, over GaAs at the melting temperature of 123 80C is 0.97 atmospheres [I]. It is this characteristic which imposes the main limitation on melt growth techniques. Early melt grown single crystals were produced by solidification in quartz boats, inside sealed ampoules. A variety of such horizontal boat growth techniques exist, including Bridgman and gradient freeze, and are the subject of several review articles [2-4]. The major part of commercially produced GaAs is still made in this way. The advantages offered by these methods are:(i) (ii) (iii) (iv)

Ability to maintain precise composition control [5]. Low dislocation density [6]. Incorporation of high dopant content [7]. Large ingot growth [8].

Such material is ideally suited to optoelectronic substrates [9] but is less suitable for high purity integrated circuit applications due to residual impurities incorporated from the growth apparatus [10]. Other disadvantages include the need for growth on a <111> axial direction resulting in inefficient wafering and non-circular shapes for (100) slices, and difficulty in producing semiinsulating material without compensation doping, usually with chromium [H]. Boat growth in the vertical furnace configuration as vertical gradient freeze (VGF) or vertical Bridgman (VB) [12] is now used for commercial production, principally for doped, low dislocation density optoelectronic substrates, but increasingly for large area, undoped semiinsulating material. Czochralski growth is the most important technique for silicon single crystal manufacture. The problem of its application to dissociable compounds has been overcome by the use of liquid encapsulated Czochralski (LEC), first proposed by Metz et al [13] and applied to III-V compounds by Mullin et al [14]. B

LEC GROWTH EQUIPMENT AND METHOD

Commercial equipment for LEC growth of III-V semiconductors became available in the 1970s and has since increased in scale and complexity. Current equipment has melt capacities up to 25 kg and growth capabilities up to 150 mm (6 inches) in diameter. The elements of the system are:(i) (ii)

A stainless steel pressure vessel chamber with water cooling, A graphite resistance heater and furnace insulation assembly.

(iii) (iv) (v)

Solid shafts for crucible and seed rotation, passing through pressure seals. Viewing optics coupled to a CCTV camera to allow monitoring of the crystallisation. A load cell for continuous crystal weighing.

Two variants of the basic growth method exist, low presssure and high pressure LEC [15]. In the former the ambient pressure during growth is in the range 2 - 1 0 atmospheres and the charge is either introduced as polycrystalline compound or synthesised in-situ by bubbling arsenic from a separate source [16,17]. In the more common high pressure version, 'direct synthesis' [18] is carried out in the crucible by the reaction at high temperature (>800°C) of elemental Ga and As under the boron trioxide (B2O3) liquid encapsulant. Growth is performed at a pressure of 20 atmospheres, with a peak pressure of about 60 atmospheres during synthesis. To obtain high purity undoped semi-insulating material, Ga and As of 7N (99.99999%) purity and pyrolytic boron nitride (pBN) crucibles are employed [19]. Growth is carried out on a <100> axis under an ambient of argon or nitrogen gas [20] at a typical growth rate of about 1 cm/hr. Melt weights from 1 kg to 28 kg [21] are used for growth of crystals of diameters from 50 mm up to 150 mm. Ingot shape is maintained by automatic diameter control derived from weight monitoring [22] and can be implemented by computer [23, 24], including 'intelligent' digital control [25]. Very high single crystal yield (>90%) can be achieved in production. Some limitation is experienced in the growth of long crystals due to breakdown into polycrystalline growth at the crystal periphery. This arises from complex curvature of the growth interface near the crystal surface and can be prevented by control of heat flow distribution in the furnace [26]. C

BULK CRYSTAL PROPERTIES

Cl

Electrical Behaviour

Undoped single crystals have a high electrical resistivity, in the range 107 to 108 ohm. cm, and high electron Hall mobility, 6000 - 8000 cm2/ Vs [27]. The electrical behaviour is dominated by a native deep donor level, EL2 [28], which is usually present at a concentration of 1 - 2 x 1016 cm"3 [29]. The accepted compensation model [30] is: [EL2]>N a >N d

(1)

Na - shallow acceptor concentration Nd - shallow donor concentration The shallow acceptor is most commonly carbon (C^) at a concentration in the range 1014 to 1016 cm"3 [31]. This should be below 1015 cm"3 for the material to remain thermally stable after wafer heat treatment [32]. The starting melt composition, estimated by charge weight before and after growth, strongly influences the electrical properties [27,29] by variation of the EL2 content, leading to an association of the EL2 defect with As-rich conditions. Resistivity is strongly correlated with the CM content [33]. The residual acceptor background is important because of its influence on the doping profile and hence on the threshold voltage for FET devices made by direct ion implantation into the substrate. The substrate resistivity range

is usually precisely specified for such applications and 'undoped' semi-insulating GaAs is in effect carbon doped within controlled limits. C2

Purity

Chemical purity can be determined by mass spectroscopic analysis in various forms, e.g. secondary ion (SIMS), spark source (SSMS) [34] and glow discharge (GDMS). The principal chemical contaminant is usually boron, incorporated from the growth system [35]. Its concentration may be as high as 1018 cm"3 but it is electrically inactive [36] except for some influence on the activation efficiency of implanted silicon ions [37]. The main donor species detected are Si and S [38] in concentrations in the 1014 cm'3 range or below. C^ acceptor content is obtained from local vibrational mode IR spectrometry [39,40] down to a threshold sensitivity of about 1 x 1014 cm"3. Zinc is another common acceptor seen in mass spectroscopy [38], Raman spectroscopy [41,42] and photoluminescence. Residual moisture in the B2O3 encapsulant is an important factor in the control of impurities in the melt. Many impurities such as boron, carbon and silicon can be reduced or removed by reaction with the encapsulant [43-45]. Precise specification of the B2O3 moisture content may be achieved by IR absorption measurement of the 0-H band [46]. Gaseous oxides may be interchanged between the puller ambient, the encapsulant and the furnace graphite, influencing the melt impurity content. Deliberate addition and concentration control of carbon monoxide (CO) in the ambient can be used to achieve controlled and uniform [CM] in crystals [47]. C3

Defects

Native defects play a dominant role in the electrical behaviour of the material and various reviews exist [48,49]. The most important centre is the deep donor EL2. [EL2] is related to melt composition and thermal history and is attributed to complexes involving the antisite AsGa [50,51]. Similarly a double acceptor state is found in crystals grown from Ga-rich melts [52]. The measurement of non-stoichiometry in the solid is therefore potentially important for defect control and methods based on precise coulometric titration of Ga [53] and As [54] content have been reported. In very high purity, low acceptor content, materials other native defects such as EL3, EL5 and EL6 play an increasingly important role [55]. C4

Dislocations

LEC grown crystals have a higher density of grown-in dislocations (104 - 105 cm"2) than boat grown ingots of equivalent size and composition (103 - 104 cm"2). These arise from thermal stresses experienced during growth [56] and a number of studies have undertaken thermomechanical modelling of the growth system. Two approaches, separately or in combination, have been taken towards dislocation reduction.

(i) (ii)

Growth in reduced thermal gradients [57]. Dislocation suppression by impurity hardening.

Silicon doping is effective in 'lattice hardening' for both Bridgman [3] and LEC growth [58,59] but produces n-type conductive material. Isoelectronic doping with various group III and group V elements to achieve lattice hardened semi-insulating material was investigated by Jacob [60]. Indium was found to be the most useful dopant at a concentration of 1019 - 1020 cm"3 [61,62]. However, such a high dopant concentration leads to enhanced compositional non-uniformity and lattice parameter variations [63,64]. Isoelectronically doped GaAs has failed to achieve commercial importance. Modifications to the basic LEC process have been developed, involving a 'hot-wall' inner chamber in which a partial pressure of As can be maintained over the crystal and melt. This allows growth in low thermal gradients and hence with higher ingot surface temperatures. Such techniques are variously designated pressure controlled LEC (PC-LEC) or vapour controlled Czochralski (VCZ) and have been demonstrated to provide significantly lower dislocation density, particularly at large diameter [65]. C5

Uniformity

The distribution of dislocations within crystals is very non-uniform, exhibiting a W-shaped profile across a diameter [66] corresponding to the predicted thermal stress distribution [56]. Superimposed on this is a microstructural distribution of slip bands [70], cellular networks [67] and lineage features acting as low angle tilt boundaries [68]. Electrical properties such as resistivity [69] and carrier concentration [70] show the same average distribution as dislocation density in as-grown crystals. This results from the close association of EL2 and dislocation structures [71], down to the microscale where [EL2] distribution can be mapped by IR transmission imaging [72,73], cathodoluminescence [74] and photoluminescence [75]. The influence of substrate dislocations on ion implanted device performance is exerted by deep level and impurity gettering effects around dislocations [76,77]. [EL2] non-uniformities can be removed by high temperature heat treatment after growth [78] and this now forms a routine and important part of the processing of semi-insulating crystals. Dislocations in all LEC GaAs are decorated by microprecipitates [79]. These have been identified as hexagonal arsenic [80]. They are influential in controlling threshold voltage uniformity of FET arrays produced by direct ion implantation in substrates [81]. Some control over the distribution of these microprecipitates has been demonstrated by means of high temperature (>1100°C) treatments in combination with rapid cooling [82]. This has led to the development of multistage post growth heat treatments for both ingots [21] and wafers [83] aimed at a simultaneous controlled redistribution of micro-precipitates and optimisation of [EL2]. D

CONCLUSION

The LEC growth method is the dominant technique for the production of single crystals for integrated circuit substrates [84,85]. It is characterised by high single crystal growth yield, giving relatively lower costs, and high purity in which the electrical properties are controlled by native defects. Because of this, batch uniformity and reproducibility are achievable by means of suitable

heat treatments, making such material favoured for ion implant processing where substrate properties exert a strong influence. Increasing production of HEMT and HBT devices, involving epitaxially grown active layers, places greater emphasis on the crystal structure quality of substrates, potentially favouring growth methods such as VGF and VCZ if these can match the economics of production for LEC. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] II1] [12] [13] [14] [15] [16] [17] [18] [19] [20]

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Nantwich, England, 1982) p.2 ] S. McGuigan, B.W. Swanson, R.N. Thomas [ !Crystal Growth (Netherlands) vol.94 no. 1 (1989) p.75-84 ] H. Kimura, A.T. Hunter, E.H. Cirlin, H.M. Olsen [ J. Crystal Growth (Netherlands) vol. 85 no. 1/2 (1987) p. 116-23] H. Schink, R.D. Schnell [ Appl. Phys. Lett. (USA) vol.53 no.9 (1988) p.764-7] S. Ozawa, T. Yokotsuka, T. Fujii, T. Fukuda, S. Kojima, M. Dceda [ Mater. Lett. (Netherlands) vol.5 no.4 (1987) p. 129-33] M. Tatsumi, T. Kawase, Y. Iguchi, K.Fujita, M. Yamada [ in Proc. 8th Con/, on Semi-Insulating UI-VMaterials, Warsaw, Poland, June 6-10, 1994, Ed. M. Godlewski (World Scientific, 1994) p. 11-5] R.T. Chen, D.E. Holmes [ J. Cryst. Growth (Netherlands) vol.61 no. 1 (1983) p. 111-24 ] S. Clark, DJ. Stirland [ Inst. Phys. Con/. Ser. (UK) no.60 (1981) p.339-44 ] G.T. Brown, M.S. Skolnick, GR Jones, B.K. Tanner, S.J. Bamet [ Proc. Con/, on Semi-Insulating IU-VMaterials, Kah-Nee-Ta, OR, USA, 24-26 Apr 1984 (Shiva Publishing, Nantwich, England, 1984) p.76] I. Grant, D. Rumsby, RM.Ware, M.R. Brozel, B Tuck [Proc. Con/, on Semi-InsulatingIU-V Materials, Evian, France, 1982 (Shiva Publishing, Nantwich, England, 1982) p.98-106 ] M. Bonnet, N. Visentin, B. Gouteraux, B. Lent, J.P. Duchemin [ Tech. Digest IEEE GaAs IC Symp., New Orleans, LA, USA, 1982 (IEEE, New York, 1982) p.54-7 ] M.S. Skolnick, M.R. Brozel, LJ. Reed, I. Grant, DJ. Stirland, RM. Ware [ J. Electron. Mater. (USA) vol.13 no.l (1984) p. 107-25 ] M.R Brozel, I. Grant, RM. Ware, DJ. Stirland [Appl. Phys. Lett. (USA) vol.42 no.7 (1983) p.6102] P. Dobrilla, J.S. Blakemore. RY. Koyama [ Proc. Conf. on Semi-Insulating UI-VMaterials, KahNee-Ta, OR, USA, 24-26 Apr 1984 (Shiva Publishing, Nantwich, England, 1984) p.282 ] P.A. Leigh, I.P. HaU, CR Elliott, B. Wakefield, M.H. Lyons [ Proc. Conf. on Semi-Insulating UI-V Materials, Kah-Nee-Ta, OR, USA, 24-26 Apr 1984 (Shiva Publishing, Nantwich, England, 1984) p.214-21 ] M. Tajima [ Jpn. J. Appl. Phys. 2 (Japan) vol.21 no.4 (1982) p.L227-9 ] G.T. Brown, CA. Warwick [ J. Electrochem. Soc. (USA) vol. 133 no. 12 (1986) p.2576-9 ] S. Miyazawa, Y. Ishii, S. Ishida, Y. Nanishi [ Appl. Phys. Lett. (USA) vol.43 no.9 (1983) p.853-6 ] D. Rumsby, I. Grant, M.R Brozel, EJ. Foulkes, RM. Ware [ Proc. Conf. on Semi-Insulating IU-V Materials, Kah-Nee-Ta, OR, USA, 24-26 Apr 1984 (Shiva Publishing, Nantwich, England, 1984) p. 165-70] T. Ogawa, T. Kojima [Mater. Sci. Monograph no.44 (Elsevier, Amsterdam, 1987) p.305 ] A.G. Cullis, P.D. Augustus, DJ. Stirland [ J. Appl. Phys. (USA) vol.51 (1980) p.2256 ] H. Yamamoto, O. Oda, M. Seiwa, M. Taniguchi, H. Nakata, M. Ejima [ J. Electrochem. Soc. (USA) vol.136 no. 10 (1989) p.3098-102 ] T. Inada, Y. Otoki, K. Ohata, S. Taharasako, S. Kuma [ J. Cryst. Growth (Netherlands) vol.96 no.2 (1989) p.327-32 ] M. Mori, G. Kano, T. Inoue, H. Shimakura. H. Yamamoto, O. Oda [ in Proc. 6th Conf. on SemiInsulating UI-VMaterials, Toronto, Canada, 1990 Eds A.G. Milnes, C.I. Miner (I.O.P. Publishing, 1990) p. 155-60] RN. Thomas et al [ Semicond. Semimet. (USA) vol.20 (1984) p. 1-87 ] CG. Kirkpatrick, RT. Chen, D.E. Holmes, K.R. Elliott [ in GaAs Materials, Devices and Circuits Eds MJ. Howes, D.V. Morgan (Wiley, 1985) p.39-94 ]

15.3 Horizontal and vertical Bridgman growth of GaAs

Previous Page

K. Fujita and M. Tatsumi November 1995

A

INTRODUCTION

The horizontal Bridgman (HB) method [1] has been widely employed to produce GaAs crystals used as substrates for optical devices such as laser diodes and LEDs. The main advantage of this method compared to the liquid encapsulated Czochralski (LEC) method [2] is that it has a lower dislocation density due to a lower temperature gradient. However it has been clarified that larger than 75 mm diameter crystals are difficult to grow by the HB method because the asymmetry in the longitudinal growth direction makes it difficult to control the shape of the solid liquid interface. However, the vertical Bridgman (VB) or the vertical gradient freeze (VGF) methods [3] having symmetry in the longitudinal direction have been enthusiastically developed since the middle 1980s. On the other hand, several new HB methods to improve the original technique have been proposed and developed. This Datareview will highlight a number of recent results for horizontal and vertical Bridgmantype growth processes. B

HORIZONTAL BRIDGMAN GROWTH

Bl

Crystal Growth Methods

Three representative methods of horizontal boat growth are: (a) (b) (c)

The two temperature-zone HB (2T-HB) method [1] The three-temperature-zone HB (3 T-HB) method [4] The gradient freeze (GF) method [5].

There are several features common to all of these methods (i)

The apparatus is provided with both a melting zone and arsenic reservoir zone. It follows that the length of the apparatus is greater than that of the LEC method. This is a disadvantage that counters the advantages of accurate arsenic pressure control.

(ii)

Quartz having both high purity and transparency is mainly used for the boat and the ampoule. As the temperature of the melting zone exceeds the softening temperature of the quartz, both the boat and ampoule tend to distort during growth.

(iii)

<11 l>-seeded growth is mainly used because the <100> orientation growth tends to cause twinning. (100) wafers cut from the ingot are D-shaped, requiring considerable material loss to make a round shape.

Responding to the above disadvantages, new horizontal boat growth techniques such as modified 2T-HB [6], horizontal zone melting (HZM) [7], and <100> growth [8] have been recently

proposed and developed. Bl.1

The modified 2T-HB system

The modified 2T-HB system [6] without an arsenic reservoir zone is relatively simple compared with other horizontal growth systems. In this process, the short quartz ampoule and large charge of polycrystalline GaAs can suppress the arsenic loss and control the variation of the stoichiometry along the whole GaAs crystal within 0.05 at%. Keeping the solid/liquid interface flat with a specially designed linear construction and viewing window, dislocation densities of lower than 2,000 cm'2 have been achieved for moderately Si doped GaAs crystals of 50 mm diameter and 300 mm length [9]. Bl.2

Horizontal zone melting

As the enlargement in the diameter of GaAs crystals progresses, the increased weight of melt tends to cause considerable distortion of the quartz boats during the lengthy growth time at high temperature. A new horizontal zone melting (HZM) technique [7] has been developed not only to address the above problem, but also to minimize the longitudinal impurity distribution. Keeping the solid/liquid interface flat or slightly convex, with minimized axial temperature gradient of the melt, Cr-doped semi-insulating GaAs crystals of up to 100 mm diameter and 600 mm length have been developed. They have dislocation densities below 5,000 cm"2 and a uniform Cr distribution in the longitudinal direction [10]. B1.3

<100> growth

<100> oriented GaAs single crystals of up to 50 mm diameter and 600 mm length have been grown by the HB method with a newly designed round-shaped boat [8]. The generation of twin defects, which makes <100> oriented growth difficult, is reproducibly prevented by selecting a vertical orientation for the seed crystal together with tuning the cone angle of the boat. This is based on the observation that these defects are generated whenever (111) As facets are in direct contact with the boat surface. The grown crystals are twin and lineage free, and have a dislocation density of less than 2,000 cm"2. B1.4

Modelling

Recently, modelling efforts have been undertaken to better understand the role of heat and fluid flow on the growth of GaAs crystals. Numerical simulations of macro-segregation in the HB process with both a dilute and a concentrated low Prandtl alloy were performed by investigating the flow structures and the species distributions [11,12]. It appeared that convective motions of the melt had a great influence on macro-segregation and the convective intensity was decreased by applying a transverse magnetic field. A three-dimensional simulation of the solidification process for growing 75 mm crystals by the modified 2T-HB system was performed. The computed profiles of the solidification front and dopant concentration distributions agreed well with the experimental results [13]. The solid/liquid interface shape depends on the speed of furnace movement and the heat loss prevention arrangements. Effects due to radiative heat loss through the viewing window were analysed [14].

B2

Bulk Crystal Properties

In general the residual strain of the crystal becomes larger and more significant as the diameter increases. Yamada [15] characterized the residual strains in GaAs (100) wafers grown by the HB technique by measuring strain-induced birefringence with an infrared polariscope. It was found that the maximum strain was as low as about 5 x 10"6 and the corresponding stress value was about 3 x 105 Pa. Also there was no evident correlation between the measured residual strain and etch pit density distributions. Zakharov et al [16] have compared stress data measured by the photoelasticity technique with strain data measured by X-ray diffraction. One of the effective techniques to reduce the dislocation density is to utilize the impurity hardening effect. Moravec et al [17] investigated the influence of isovalent dopants such as In and N on crystal growth and electrical properties of GaAs crystals. It was found that an In concentration higher than 5 x 1019 cm"3 in the crystal caused the solid solution hardening effect [18]. In and N doping did not affect the electrical properties. The deviation from the stoichiometric composition is the most important factor to affect the crystalline and electrical properties of GaAs crystals. Nishizawa [19] has reviewed the importance of stoichiometry control which his group had investigated for many years. Silicon, which is a representative dopant in GaAs, behaves as an amphoteric atom, i.e. SiGa and Si^. It follows that the stoichiometry can affect the site distributions of silicon impurities. Fujii et al [20] investigated the effect of arsenic pressure on the site distribution of silicon. They found that the carrier concentration and mobility increased with an increase in arsenic pressure. Boudriot et al [21] investigated the compensation degree as a function of arsenic pressure with three different methods: local vibrational mode (LVM) infrared absorption, Hall effect and conductivity. These methods yield the same result of a minimum of compensation at a small excess of Ga in the melt. Newman [22] has reviewed the lattice locations of silicon impurities in GaAs in terms of effects due to stoichiometry, the Fermi energy, the solubility limit and DX behaviour. Defects or impurities resulting in deep levels such as EL2 and oxygen particularly influence the electrical and optical properties of semi-insulating GaAs. Wang et al [23] investigated the EL2 formation mechanism by mapping topographies and found that the EL2 distribution shows a good correlation with growth striations. However, EL2 and growth striations have only a moderate or no correlation with dislocation distribution for undoped and Si-doped GaAs regardless of growth techniques. Fujii et al [24] have made a thermodynamic analysis of the Ga-As-Si-O system in the GF method, previously performed in the HB method by Akai et al [25]. They found that concentrations of EL6 in undoped HB GaAs crystals decreased with increasing oxygen concentration. This explains why V(Ja "VA83 the origin of the EL6, might change to VGa - OM with increasing oxygen concentration.

C

VERTICAL BRIDGMAN AND VERTICAL GRADIENT FREEZE GROWTH

Cl

Crystal Growth Methods

The vertical Bridgman and vertical gradient freeze are classical growth methods which have been carried out mainly in oxide single crystals. In these methods, the crystallization proceeds from the bottom to the top in a crucible, so it is impossible to observe the state of growing in real time. This fact requires the achievement of stable and reproducible growth for application to production. For compound semiconducting materials, frequent generation of twins or polycrystals, especially in the <100> growth direction, has limited the application for practical use [26]. However, by developing the crucible or furnace structure, GaAs single crystals with large diameter have become able to be grown stably since the middle 1980s [3]. However, single crystals of other compound semiconducting materials such as InP, GaP and II-VI materials have not been successfully grown, since these crystals are liable to generate twins. There are several methods for growing crystals in the vertical direction in a crucible. CLl

Vertical Bridgman (VB) method

The crucible containing GaAs melt is moved downward relative to the heater having axial temperature gradient [27,28]. C1.2

Vertical gradient freezing (VGF) method

By adjusting the electrical power supplied to each heater, the temperature profile is shifted upward [26,29]. This method is fundamentally the same as the VB method, although generally the VGF furnace has many heater zones in order to move the temperature profile electrically. However, it is a mechanically simple structure having no moving mechanism [30]. C1.3

Vertical zone melting (VZM) method

The temperature profile with a central spike region travels to shift the GaAs molten zone by moving either the heater or crucible [31-33]. In this method, it is essential to suppress the dissociation of the GaAs melt during growth. The dissociation pressure of As is 1 atm at the melting temperature and steeply increases with the rise of temperature at the melt surface. The suppression of the dissociation is performed by the following methods. (1)

The crucible is sealed in an evacuated quartz ampoule [27]. In some cases, the arsenic pressure is controlled by placing solid arsenic at the bottom of the ampoule where the temperature is adjusted to around 617° C [29]. If different materials such as pBN are used for the ampoule, complete sealing is difficult and there is a need to continuously supply arsenic vapour during growth.

(2)

The other method is similar to the liquid encapsulated Czochralski method and the melt is covered with boric oxide encapsulant. This is the liquid encapsulated vertical Bridgman (LEVB) method reported by Hoshikawa [28]. In this case the quartz ampoule is not essential.

C1.4

Control of wetting

Wetting of the crucible surface by the GaAs melt induces the generation of polycrystals or twins. Crucibles for GaAs crystals are usually made of quartz or pBN. To reduce the extent of the wetting, the inner surface of the crucible is treated in various manners, such as surface roughening and coating. The most successful method on pBN crucibles is a thin B2O3 layer forming by oxidising the surface of the crucible. This B2O3 layer is liquefied above about 400 0 C and then prevents direct contact between the crucible and the GaAs melt. By this method reproducible growth of single crystals can be achieved [34]. C1.5

Control of thermal environment

In order to control the temperature profile in the growth furnace, various kinds of heating system and furnace structure have been tried. For example Monberg [30] reported a furnace for dynamic gradient freeze growth, which consisted of multi metal heaters with 54 elements. The temperature gradient in the furnace directly controls the shape of the solid/liquid interface and temperature gradient near the solid/liquid interface [35]. This significantly affects the perfection of the single crystal, that is, the generation of dislocations, residual strain, twins and polycrystallization. Generally it is preferable that the solid/liquid interface is not concave to the melt in GaAs crystal growth. Realizing a flat or convex interface has been tried by applying heat near the interface or by increasing the heat flow through the crucible support. However, such control has not been obtained yet. On the other hand, several analyses of the above problems have been carried out by computer simulation to derive the influence of the temperature profile, the shape of the crucible support, crucible material and growth rate on the solid/liquid interface [36,37]. The thermal stress field during VB growth was also calculated from the simulated temperature distribution [38,39]. C1.6

Application of magnetic field

Park et al investigated effects of the application of an axial magnetic field on vertical gradient freeze GaAs single crystal growth. By applying a 1200 G magnetic field, they obtained nearly striation free single crystals [40]. C2

Bulk Crystal Properties

Since the VB and VGF methods are growth in a crucible, the temperature gradient can be reduced to less than 20°C/cm without limitation of diameter controllability. So a lower dislocation density can be expected in comparison with that of LEC crystals for which the dislocation density is usually higher than 30,000/cm2 in 75 mm diameter crystals [41,42]. Results for VGF crystals have been published by researchers of AT&T [3]. They obtained an average dislocation density of 300 cm"2 for 50 mm diameter Si-doped crystals, which showed impurity hardening effects. The dopant concentration was in the range 1 ~ 3 x io 18 cm"3. An undoped single crystal of 75 mm diameter showed an average dislocation density of about 2500 cm"2 with a four-fold symmetry pattern [43]. Hoshikawa et al have reported that 75 mm diameter crystals grown by the LE-VB method exhibit a dislocation density ranging from 5000 to 40,000 cm"2. The temperature gradient is from 15 to 20°C/cm [28]. Bourret et al reported effects of total liquid encapsulation on both the crystal surface and the dislocation densities [34]. A 34 mm diameter VZM crystal was grown under a vertical temperature gradient of 9°C/cm by Henry et

al [33], and the dislocation density was a uniform (2 - 5) x 103 cm"2 throughout the crystal. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] II1] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32]

L.R. Weisberg, F.D. Rosi, P.G. Herkart [ in Properties of Elemental and Compound Semiconductors vol.5 Ed. H.C. Gatos (Interscience, New York, 1960) p.25-67 ] J.B. Mullin, RJ. Heritage, CH. Holliday, B.W. Straughan [ J. Cryst. Growth (Netherlands) vol.3-4 (1968) p.281-5] W.A. Gault, E.M. Monberg, J.E. Clemans [ J. Cryst. Growth (Netherlands) vol.74 no.3 (1986) p.491-506] S. Akai et al [ Proc. Symp. on Optoelectronics Epitaxy and Device Related Processes (Electrochemical Society, New Jersey, 1983) p.41-60 ] J.M. Woodall [ Electrochem. Technol. (USA) vol.2 (1964) p. 167-9 ] T.P. Chen, Y.D. Guo, T.S. Huang [ J Cryst. Growth (Netherlands) vol.94 (1989) p.683 ] S. Mizuniwa, M. Kashiwa, T. Kurihara, K. Nakamura, S. Okubo, K. Ikegami [J. Cryst. Growth (Netherlands) vol.99 (1990) p.676-9 ] K. Murata, T. Ishihara, M. Sato, N. Kito, M. Hirano [J. Cryst. Growth (Netherlands) vol.141 (1994) p.29-36 ] T.P. Chen, T.S. Huang, LJ. Chen, Y.D. Gou [ J. Cryst. Growth (Netherlands) vol. 106 (1990) ] S. Mizuniwa, M. Nakamori, K. Nakamura, S. Okubo [ Hitachi Rev. (Japan) no. 10 (1991) p.37 ] S. Kaddeche, H. Ben Hadid, D. Henry [ J. Cryst. Growth (Netherlands) vol. 135 (1994) p.341-53 ] S. Kaddeche, H. Ben Hadid, D. Henry [J. Cryst. Growth (Netherlands) vol.l41(1994) p.279-90 ] M.H. Hsieh, CC. Chieng, K.H. Lie, Y.D. Guo [Proc. Inst. Mech. Eng. C (UK) vol.207 no.3 (1993) p. 185-95] K.H. Lie, J.T. Hsu, Y.D. Guo, T.P. Chen [ J. Cryst. Growth (Netherlands) vol.109 (1991) p.205-11] M. Yamada [ J. Appl. Phys. (USA) vol.74 no.4 (1993) p.2436-9 ] S.N. Zakharov, S.A. Laptev, V.M. Kaganer, VT. Bublik, VL. Indendom [Phys. StatusSolidiA (Germany) vol. 131(1992) p. 143-9 ] F. Moravec, B. Stepanek, P. Doubrava [ Cryst. Res. Technol. (Germany)\ol.26 no.5 (1991) p.579-85 ] T. Inoue et al [Proc. 12th Int. Symp. GaAs and Related Compounds, Karuizawa, 1985 (Adam Hilger, Bristol, 1986) p.7-12 ] J. Nishizawa [ in Non-Stoichiometry in Semiconductors Eds KJ. Bachman, H.L. Hwang, C. Schwab (Elsevier Sciemce Publishers B.V., Netherlands, 1992) p.95-1063 ] K. Fujii, F. Orito, H. Fujita [Mater. Sci. Forum (Switzerland) vol. 117-118 (1993) p.393-8 ] H. Boudriot, W. Siegel, K. Deus, W. Buhig [ Solid State Commun. (USA) vol.89 no.10 (1994) p.889-91 ] R C Newman [ Semicond. Sci. Technol. (UK) vol.9 (1994) p. 1749-62 ] F.C Wang, M.F. Rau, J. Chars [ J. Cryst. Growth (Netherlands) vol. 103 (1990) p.311-22 ] K. Fujii, F. Orito, H. Fujita, T. Sato [ J Crys. Growth (Netherlands) vol. 121 (1992) p.255-66 ] S. Akai, K. Fujita, M. Sasaki, T. Tada [ Proc. 9th Int. Symp. GaAs and Related Compounds, Oiso, 1981 (Inst. Phys. Conf. Ser. no.63, UK, 1982) p. 13-8 ] A.G. Fischer [ in JCrystal Growth (Pergamon Press, Oxford, 1980) p357-93 ] RE. Kremer, D. Francomano, GH. Beckhart, K.M. Burke [ J. Mater. Res. (USA) vol.5 no.7 (1990) p. 1468-74] K. Hoshikawa, H. Nananishi, H. Kohda, M. Sasaura [ J. Cryst. Growth (Netherlands) vol.94 (1989)p.643-50] CR. Abernathy et al [ J. Cryst. Growth (Netherlands) vol.85 (1987) p.106-153 ] E.M. Monberg, H. Brown, CE. Bormer [J. Cryst. Growth (Netherlands) vol.94 (1989) p.109-14 ] E.M. Swiggard [J. Cryst. Growth (Netherlands) vol.84 (1989) p.S56-8 ] RS. Tang, L. Sargent, J.S. Blakemore [J.Cryst.Growth (Netherlands) vol. 103 (1990) p323-9 ]

[33] [34] [35] [36] [37] [38] [39] [40] [41] [42] [43] [44] [45] [46] [47] [48]

[49] [50]

RL. Henry, P.E.R. Nordquist, RJ. Gorman, S.B. Qadri [ J. Cryst. Growth (Netherlands) vol. 109 (1991)p228-33] E.D. Bourret, E.C. Merk [ J. Cryst. Growth (Netherlands) vol. 110 (1991) p.395-404 ] E.D. Bourret, J.B. Guitron, M.L. Galiano, E.E. Haller [ J. Cryst. Growth (Netherlands) vol.102 (1990) p.877-84 ] K. Koai, K. Sonnenberg, H. Wenzl [ J. Cryst. Growth (Netherlands) vol. 137 (1994) p.59-63 ] CW. Lan, CC. Ting [ J.Cryst. Growth (Netherlands) vol.149 (1995) p.175-86 ] W. Rosch, F. Carlson [J. Cryst. Growth (Netherlands) vol.109 (1991) p.75-81 ] A.S. Jordan, E.M. Monberg [ J. Appl. Phys. (USA) vol.73 (1993) p.4016-22 ] YJ. Park, S. Min, S. Hahn, J.K. Yoon [ J. Cryst. Growth (Netherlands) vol. 154 (1995) p. 10-8 ] G. Muller, G. Hirt, D. Hofinann [ Proc. 7th Con/, on Semi-insulating Materials, Ixtapa, Mexico, 1992 (Institute of Physics Publishing, Bristol and Philadelphia) p.73-84 ] S. Kuma, M. Shibata, T. Inada [ Proc. 20th Int. Symp. GaAs and Related Compounds, Freiburg, 1993 (Institute of Physics Publishing, Bristol and Philadelphia) p.497 ] J.E. Clemans, T.I. Ejim, W.A. Gault, E.M. Monberg [ AT&T Tech. (USA) vol.68 (1989) p.29-42 ] Z-Q. Fang, D.C. Look [ J. Appl. Phys. (USA) vol.69 (1991) p.8177-82 ] D.C. Look et al [ J. Appl. Phys. (USA) vol.66 (1989) p. 1000-2 ] M.L. Gray, L.S. Blakemore, J.M. Parsey Jr., J.E. Clemans [ J. Appl. Phys. (USA) vol.63 (1988) p.5689-93 ] RS.Tang, J.S. Blakemore, RE. Kremer, K.M. Burke [ J. Appl. Phys. (USA) vol.66 (1989) p.5428-34 ] RE. Kremer, D. Francomano, B. Freidenreich, H. Marshall, K.M. Burke [ Proc. 6th Conf. on Semi-insulating III-VMaterials, Toronto, Canada, 1990 (Adam Hilger, Bristol, Philadelphia, New York) p.205-103] B.E. Freidenreich et al [ Proc. 7th Conf on Semi-insulating Materials, Ixtapa, Mexico, 1992 (Institute of Physics Publishing, Bristol and Philadelphia) p. 105-10 ] P.E.R. Nordquist Jr., RL. Henry, J.S. Blakemore, S.B. Saban, RJ. Gorman [ J. Cryst. Growth (Netherlands) vol.141 (1994) p.343-6 ]

15.4 New types of LEC growth and vapour controlled LEC GaAs M. Tatsumi November 1995

A

BVTRODUCTION

The liquid encapsulated Czochralski (LEC) method was first applied to III-V semiconducting materials by Mullin et al in 1968 [1] and has been extensively advanced by many researchers. At present, LEC methods under a high pressure (HPLEC) are employed for the production of GaAs single crystals by using a well developed HP LEC puller. The LEC method is excellent in productivity but has some problems for the purpose of growing single crystals with higher quality, as described below. (1)

It is difficult to strictly control the stoichiometry. The boric oxide (B2O3) encapsulation is not sufficient to suppress completely the dissociation of the GaAs melt by arsenic evaporation or diffusion, throughout a growth process.

(2)

The dislocation density cannot be reduced, since the high temperature gradient is essential for suppressing the dissociation of the GaAs single crystal during LEC growth.

(3)

Convection of the high pressure inert gas in the furnace induces large temperature fluctuations at the growth interface, which generates defects and instability of the growth.

(4)

The LEC method is fundamentally normal freezing. Therefore, intentionally doped elements exhibit non-uniform distributions due to segregation phenomena. As a result, electrical properties cannot be uniform throughout the crystal.

Improvements to these drawbacks have been tried by modifying the LEC method. Several modified LEC methods have been reported, that is, (1) As injection LEC, (2) Czochralski under arsenic vapour pressure, (3) fully encapsulated Czochralski, (4) double crucible LEC and (5) magnetic field applied LEC. B

As INJECTION LEC

Synthesis of GaAs by arsenic injection was carried out by Pekarek in 1970 [2] and single crystals were grown in the same high pressure chamber. The temperature of the coldest point of the arsenic reservoir was kept between 580 and 6300C. He obtained a GaAs single crystal with a Hall mobility of 5,000 to 6,000 cm2/ Vs and with a carrier concentration of the order of 1015 cm"3. Inada et al [3] also grew GaAs crystals by using a newly developed high pressure As injection LEC growth system similar to that used by Pekarek [2]. Arsenic gas supplied from a quartz reservoir was injected into the GaAs melt which was directly synthesized in situ under a high pressure of Ar. The temperature of the reservoir was controlled between 610 and 620 0 C and the melt composition of 0.497 - 0.501 As atomic fraction was estimated from the variation in weight of the material in the crucible. They obtained 50 mm diameter <100> oriented crystals, which were grown from near stoichiometric melts. The radial distribution of the resistivity in the

As-injection-LEC crystal showed a reduced standard deviation (46%) compared to that of an LEC crystal (76%). They also fabricated 500 to 600 MESFETS on the wafer and obtained a standard deviation oVth = 28 mV which was less than half the value of 66 mV in conventional LEC GaAs. C

CZOCHRALSKI UNDER ARSENIC VAPOUR PRESSURE

Another attempt to control the stoichiometry of the GaAs melt is growth under an arsenic vapour pressure which is kept at the dissociation pressure of the melt. In this case the GaAs melt is not covered with B2O3 encapsulant layer, since the surface of the melt must be in contact with the arsenic vapour. In order to realize this it is essential to seal the growth chamber which must be equipped with movable (both translation and rotation) rods and hot walls which must be of sufficiently high temperature to prevent condensation of arsenic. The earliest technique was developed by Gremmelmaier [4] for growing GaP single crystals. The quartz chamber had a magnetically coupled pulling rod. However, this structure gave unstable movement to the seed crystal and was very complex, costly and difficult to use. Mullin et al [5] have proposed a pressure balancing technique which dynamically balances the dissociation pressure with an equivalent pressure of inert gas. This system had a simple construction, a novel liquid seal at the pull rod bearing and a liquid nanometer for sensing pressure differentials. A more simplified liquid-seal Czochralski furnace was constructed by Leung et al [6]. This furnace was also made of quartz and used molten B2O3 as a push-pull rotary seal. They grew a GaAs single crystal with a diameter of 10 - 20 mm and a length of 40 - 100 mm. The dislocation density was less than 100 cm"2. By using a similar quartz ampoule with a rotating liquid seal OfB2O3, Moulin et al [7] also grew GaAs single crystals of 20 to 30 mm diameter and 70 mm in length. The average dislocation density was about 104 cm"2. Undoped crystals grown by this method typically showed n-type with a carrier concentration of 1 x 1015 cm"3 although some crystals were semi-insulating with a carrier concentration of 2 x 108 cm"3. By sophisticating the above methods, Tomizawa et al [8] have succeeded both in sealing the hot wall inner chamber stably and in controlling the pressure of As vapour in the inner chamber by precisely controlling the lowest temperature within between 616 and 620 0 C. 50 mm diameter GaAs single crystals were grown along the <111> direction under a temperature gradient of 10°C/cm. The dislocation density was about 2 x 103 cm"2 and the radial distribution was Ushaped, while LEC crystals showed an order of magnitude higher dislocation density and a Wshaped distribution. The use of a quartz chamber made an undoped crystal conductive with a carrier concentration of 5 x 1016cm"3. In 1987 Tomizawa et al obtained undoped semi-insulating GaAs single crystals of 75 mm diameter by using a newly developed pulling apparatus [9]. The dislocation density was about 1.6 x 104 cm"2 and the resistivity was 8.5 x 107Q cm. Crystals with minimum dislocation density could be grown under As pressure with As reservoir temperatures from 610 to 624 0 C. In 1987, Tada et al reported that low dislocation density InP single crystals could be grown by the modified LEC method, where the InP melt was covered with a boric oxide layer under a phosphorus vapour pressure [10,11] The phosphorus pressure in this method plays its role, not in controlling the melt stoichiometry but in suppressing the dissociation of the single crystal under the relatively low temperature

gradient. It does not need strict control of the phosphorus vapour pressure. This method was first employed in the production of InP single crystals of high quality, and has been applied to the growth of GaAs single crystals with diameters of 100 and 150 mm under a low temperature gradient of 20°C/cm [12-14]. The dislocation density of a 100 mm diameter crystal was 4,000 to 5,000 cm"2, one order of magnitude less than that of an LEC crystal, and the residual strain was also one quarter of that of an LEC substrate. Undoped GaAs single crystals grown by this method were semi-insulating with a resistivity of 1 x 108Q cm, the concentration of carbon being 3 x 1015 cm"3.150 mm diameter single crystals have also been successfully grown by this method. Both the dislocation density and the residual strain were less than those of 100 mm LEC crystals. D

FULLY ENCAPSULATED CZOCHRALSKI (FEC)

Another method for growing single crystals with low dislocation density is the fully encapsulated Czochralski (FEC) method, in which a crystal being grown is flilly encapsulated with liquid B2O3. Since thick B2O3 encapsulant can result in a low temperature gradient of 30 - 50°C/cm, this method can reduce thermal stresses which induce dislocations and can suppress damage of the surface. Kohda et al have grown completely dislocation-free and striation fee, semi-insulating GaAs crystals with 50 mm diameter by this method [15]. In order to obtain dislocation free crystals, the impurity hardening effect was used by doping with indium at a concentration of 4.5 x 1020 atoms/cm3, and grown-in dislocations were eliminated by using both a dislocation-free seed crystal and necking. E

DOUBLE CRUCIBLE LEC

In the LEC system, intentionally or unintentionally doped impurities in the melt show segregation phenomena due to normal freezing. Therefore, the concentration of the impurity varies along the growth direction and the extent of the change depends on the segregation coefficient. (When a segregation coefficient is equal to unity, the impurity concentration is uniform.) This variation of the impurity concentration in the crystal reduces the uniformity of the crystal properties, such as resistivity, mobility, dislocation density and lattice constant. Of course, it is preferable to make crystals having a uniform concentration of impurity. The 'double crucible' method has been tried for this purpose. This method employs two crucibles, namely inner and outer crucibles. The inner crucible having several holes separates the GaAs melt into two melts with different impurity concentration. The outer melt comes into the inner crucible through the holes with the progress of crystal growth from the inner melt. This inflow of the melt keeps both the volume and the impurity concentration of the inner melt constant during the growth. This method has been tried on the growth of In-doped GaAs single crystals, where the segregation coefficient of In is 0.1. Since In atoms in GaAs crystals show an impurity hardening effect and produce dislocation free crystal, it is important to grow crystals with a uniform In concentration. Okada et al grew In-doped GaAs crystals by using an inner crucible connected to a lift system [16]. The uniformity of In concentration was improved in comparison with that in a conventional LEC crystal but it was not completely uniform because of mixing by out diffusion of the inner melt to the outer melt.

Matsumoto et al succeeded in growing 75 mm diameter GaAs single crystals having uniform In concentrations in the central 50% of the crystal (namely excluding the seed and the tail ends) [17]. They used a floating type inner crucible and optimized the size of the holes of the inner crucible by computer simulation in order to suppress the out diffusion of the melt. The concentration of [CAJ and [EL2] were in the range of 1.2-1.4 x 1015 cm"3 and 8.5-9.4 x 1015 cm"3, respectively, the same as for crystals grown by the conventional LEC method. The dislocation density of the crystal was very low for the fraction solidified from 0.0 to 0.7. Suppression of cellular growth due to excess concentration of In was elongated up to a fraction solidified of 0.7 compared to 0.5 for conventional LEC crystals. F

LEC WITH APPLIED MAGNETIC FIELD

Applying a magnetic field to the Czochralski growth of Si has resulted in useful effects on the control of oxygen concentration. In LEC GaAs systems the magnetic field makes it possible to control or suppress the melt convection. It is expected that this suppression of melt flow affects the solid-liquid interface, the temperature fluctuation of the melt, segregation of the impurity, stoichiometry of the crystal and the concentration of deep levels. There are two obvious directions of the applied magnetic field, that is, vertical (parallel to the growth direction) and horizontal (normal to the growth direction). Each direction has slightly different effects on the melt convection. The Optoelectronics Joint Research Laboratory has extensively investigated the effects of applying a magnetic field to GaAs LEC growth. A horizontal magnetic field of 0.13 T was applied to the growth system of 50 mm diameter GaAs crystals by using a superconducting magnet [18]. The temperature fluctuation of 180C in the melt was reduced continuously down to 0.1 0 C at 0.125 T. The electrical resistivity of the undoped GaAs changed from being semi-insulating (108 Q cm) to semi-conducting (10 Q cm) by increasing the magnetic field. It was concluded that this decrease in resistivity was due to a decrease in the deep level concentration, which might be caused by the temperature fluctuation in the melt. The authors also reported that a vertical magnetic field of 0.1 T could also suppress the temperature fluctuation in the melt to less than 0.3 0 C [19]. Terashima et al.[20,21] reported effects of the magnetic field on residual impurity concentrations. A magnetic field up to 0.38 T was applied horizontally to the crucible. Subsequently, the carbon and boron concentrations were found to have changed according to the applied magnetic field strength. They explained that the oxidation rate of free gallium in molten GaAs and the deoxidation rate of carbon and boron oxides were the rate limiting reactions to determine the carbon and boron incorporation rate into GaAs crystals, and that these reactions were affected by both the suppression of the melt convection and the increase of the effective melt viscosity under the magnetic field. The concentrations of the deep level EL2 were strongly affected by the melt composition of GaAs crystals grown under the magnetic field [22] and the homogeneity was better than that of conventional LEC crystals [23]. GaAs crystals of 50 mm diameter could be grown at a high pulling rate of 20 or 27 mm/hr using the magnetic field applied LEC technique. The distributions of both EL2 concentration and microscopic resistivity in these crystals were not as homogeneous as those of conventional LEC crystals grown at a pulling rate of 9 mm/ hr, though these were improved by ingot annealing at 9500C [24].

Osaka at al [25] have also reported that temperature fluctuations could be suppressed by applying a vertical magnetic field and that homogeneous crystals could be obtained by optimizing the balance between residual thermal convection under a magnetic field and forced convection generated by crystal rotation. Carlson et al [26] investigated the effects of axial magnetic fields on micro- and macrosegregation during LP-LEC growth of GaAs using IR absorption analysis with a spatial resolution of 2 nm. Kawase et al [27] grew 75 mm diameter In-doped GaAs single crystals under strong vertical magnetic fields of up to 0.6 T applied by a superconducting magnet and found that the magnetic field affected the solid-liquid interface shape, the distribution of indium concentration etc. G

CONCLUSION

The LEC growth method is the preferred technique for the production of semi-insulating GaAs single crystals. Several attempts have been carried out to improve LEC GaAs crystals. Though those methods have played an important role in both the elucidation of problems in GaAs single crystals and the improvement of the quality of GaAs crystals, they have not yet been applied to production. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9]

[10]

[II] [12] [13]

[14]

J.B. Mullin, RJ. Heritage, CH. Holliday, B.W. Straughan [ J. Cryst. Growth (Netherlands) vol.3-4 (1968)p.281-5] L. Pekarek [Czech. J. Phys. (Czechoslovakia) vol.B20 (1970) p.857-8 ] T. Inada, T. Sato, K. Ishida, T. Fukuda, S. Takahashi [J. Electron. Mater. (USA) vol.15 (1986) p.169-73 ] R. Gremmelmaier [ Z Nat. forsch A (Germany) vol.11 (1956) p.511 ] J.B. Mullin, W.R. McEwan, CH. Holliday, A.E.V. Webb [ J. Cryst. Growth (Netherlands) vol. 13-14 (1972) p.629-34] P.C. Leung, W.P. Allred [ J. Cryst. Growth (Netherlands) vol. 19 (1973) p.356-8 ] M. Moulin, M. Faure, G. Bichon [ J. Cryst.Growth (Netherlands) vol.24-25 (1974) p.376-96 ] K.Tomizawa, K.Sassa, Y.Shimanuki, J.Nishizawa [ J. Electrochem. Soc. (USA) vol.131 (1984) p2394-7 ] K. Tomizawa, K. Sassa, Y. Shimanuki, J.Nishizawa [Proc. 14th Int. Symp. GaAs and Related Compounds, Heraklion, Crete, 1987 (Institute of Physics Publishing, Bristol and Philadelphia) p.435-8 ] K. Tada, M. Tatsumi, M. Nakagawa, T. Kawase, S. Akai [Proc. 14th Int. Symp. GaAs and Related Compounds, Heraklion, Crete, 1987 (Institute of Physics Publishing, Bristol and Philadelphia) p.439-42 ] M. Tatsumi et al [ Proc. 1st. Int. Conf. on InP and Related Materials, Oklahoma, 1989 (Society of Photo-Optical Instrumentation Engineers, Washington) p. 18-28] T.Kawase et al [ Proc. 7th. Conf. on semi-insulating Materials, Ixtapa, Mexico, 1992 (Institute of Physics Publishing, Bristol and Philadelphia) p.85-90 ] T. Kawase, S. Fujiwara, K. Hashio, M. Tatsumi, T. Shirakawa, K. Tada [ Proc. 19th Int. Symp. GaAs and Related Compounds, Karuizawa 1992 (Institute of Physics Publishing Bristol and Philadelphia) p.13-8 ] M. Tatsumi, Y. Hosokawa, T. Iwasaki, N. Toyoda, K. Fujita [Mater.Sci.Eng. B (Switzerland) vol.28 (1994) p.65-71]

[15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27]

H. Kohda, K. Yamada, H. Nakanishi, T. Kobayashi, J. Osaka, K.Hoshikawa [ J. Cryst. Growth (Netherlands) vol.71 (1985) p.813-6 ] H. Okada, T. Katsumata, M. Eguchi, T. Fukuda [ J. Cryst. Growth (Netherlands) vol.82 (1987) p643-6] K. Matsumoto, M.Yamashita, R. Nakai, S. Yazu, K. Tada, S. Akai [ Proc. Sth.Conf. on Semi-insulating III-VMaterials, Malmo, 1988 (Adam Hilger, Bristol,UK) p.447-523 ] K. Terashima, T. Katsumata, F. Orito, T. Kikuta, T. Fukuda [ Jpn. J. Appl. Phys. (Japan) vol.22 (1983) p L325-7 ] K. Terashima, T. Katsumata, F. Orito, T. Fukuda [ Jpn. J. Appl. Phys. (Japan) vol.23 (1984) p.L302-4 ] K. Terashima, F. Orito, T. Katsumata, T. Fukuda [ Jpn. J Appl. Phys. (Japan) vol.23 (1984) p.L485-7] K. Terashima, J. Nshio, S. Washizuka, M. Watanabe [ J Cryst. Growth (Netherlands) vol.84 (1987)p.247-52] T. Kimura, T. Obotaka, T. Fukuda [ J. Cryst. Growth (Netherlands) vol. 84 (1987) p394-8 ] T. Kimura, T. Katsumata, M. Nakajima, T.Fukuda [ J. Cryst. Growth (Netherlands) vol.79 (1986) p.264-70 ] T. Kimura, T. Obotaka, T.Fukuda [ J. Cryst. Growth (Netherlands) vol.84 (1987) p394-8 ] J. Oska, H. Kohda, T. Kobayashi, K.Hoshikawa [ Jpn. J. Appl. Phys. (Japan) vol.23 (1984) p.L195-73 ] DJ. Carlson, A.F.Witt [ J. Cryst.Growth (Netherlands) vol. 116 (1992) p.461-72 ] T. Kawase, A. Kawasaki, K.Tada [ Proc. 13th Int. Symp.GaAs and related Compounds, Las Vegas, Nevada, 1986 (Institute of Physics Publishing, Bristol and Philadelphia) p.27-32 ]

15.5 Heat treatments of GaAs ingots O. Oda January 1996

A

INTRODUCTION

In order to improve the quality of GaAs for the purpose of applications in high-frequency devices, it was found that ingot annealing is an indispensable procedure during the preparation of bulk GaAs materials. Ingot annealing under various conditions has been performed and it is now known that annealing at moderate temperatures (800 - 9500C) for a long period is the best condition for realizing good quality wafers and high quality devices. In special cases, high temperature ingot annealing is also performed, depending on the device structures. The contribution of ingot annealing to improved device performance has been proved. B

ANNEALING EXPERIMENTS

It was first found that ingot annealing is effective in the improvement of electrical properties of GaAs in 1993 [I]. From that time, many investigations have been performed in order to optimize annealing conditions, to examine the effect of annealing on various properties and to clarify the mechanism of annealing, as seen in TABLE 1 [1-36]. In this table, studies on wafer annealing (Datareview 15.6) are also summarized for convenience. (1)

Uniformity of electrical properties and EL2 distribution is largely improved by ingot annealing at moderate temperatures [1-4,7,12,17].

(2)

Appropriate annealing improves not only the electrical properties but also the uniformity of photoluminescence and cathodoluminescence [4,7,14,32,39-44].

(3)

High temperature annealing (>1100°C) dissolves EL2 and arsenic precipitates, intrinsically existing in conventionally grown LEC-GaAs, and which deteriorate the uniformity of electrical properties [8,17,24,25,29,30,33,34,36].

(4)

Dissolved EL2 and arsenic precipitates regenerate reversibly by moderate temperature annealing [8,15,24,29,30,35].

(5)

Conversion between semi-insulating and semi-conducting properties takes place, depending on the carbon concentration in the material and the cooling rate after annealing [6,19,20,31].

Based on these results, multiple-step ingot annealing [8,17,26] has been examined. A more sophisticated ingot annealing procedure, three-step ingot annealing [25], which gives good uniformity even from the viewpoint of IR scattering, has been proposed.

TABLE 1. Summary of annealing studies. Authors

Year

Content

Ref

Rumsby et al Yokogawa et al

1983 1984

[1] [2]

Holmes et al Chin et al Osaka et al Look et al

1984 1985 1985 1986

Lohnert et al

1986

Lagowski et al Kitagawara et al Ogawa Otoki et al Osaka et al Noto et al Kang et al Kanbe et al Clark et al Kitagawara et al Reichlmaier et al Nakamura et al Asom et al Ogawa

1986 1986 1986 1986 1986 1986 1987 1987 1988 1988 1988 1988 1988 1988

Miyazawa et al Inada et al Kuma et al Lee et al Stirland et al Molva et al Otoki et al Clark et al Matsui et al Menninger et al Mori et al Oda et al Suemitsu et al Ohkubo et al

1988 1989 1989 1989 1989 1989 1990 1990 1990 1990 1990 1992 1992 1993

First attempt at ingot annealing Improvement of PL, mobility, resistivity and pinch-off voltage distributions by ingot annealing Effect of annealing on EL2 distributions Improvement of CL uniformity by annealing Ingot annealing of In-doped dislocation-free crystals Reversibility between semi-conducting and semi-insulating by annealing Improvement of IR absorption, scanning PL, CL and electrical properties Observation of inverted thermal conversion (ITC) SI conversion after annealing of low carbon content GaAs Effect of melt composition and thermal history Change of IR scattering images by annealing Mechanism of annealing effect for In-doped GaAs PL intensity after annealing of In-doped GaAs EL2 concentration increase as a function of annealing time Effect of ingot annealing on device performances Uniformity as a function of annealing temperature EL2 concentrations and arsenic precipitate densities Resistivity drop of low carbon GaAs and its mechanism Resistivity change by cooling rate after annealing Ingot annealing of Bridgman-grown GaAs Effect of arsenic fraction in the melt and thermal history on the electrical properties Model on defects around dislocations and effect of annealing Behaviour of excess arsenic as a function of annealing Three-step ingot annealing Relationship between excess As and EL2 Defect structures after annealing Mechanism of microscopic inhomogeneity IR scattering centres as a function of annealing temperature EL2 concentrations as a function of annealing temperature Effect of annealing temperature and cooling rate Effect of annealing on CL uniformity Multiple-step wafer annealing Mechanism of multiple-step wafer annealing EL2 kinetics during annealing p-type thermal conversion after high temperature annealing

C

[3] [4] [5] [6] [7] [8] [9,10] [H] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33] [34] [35] [36]

ELECTRICAL PROPERTIES

It was proved that moderate temperature annealing greatly improves the uniformity of electrical properties such as resistivity and mobility by many authors [1-4,7,12,17]. The microscopic uniformity was mainly measured by the three guard ring method [37] and it is known that the best uniformity achievable is less than 5%, which is nearly at the level of the detection limit. This uniformity is very important since it is directly correlated with the threshold voltage variation of FETs [38]. Hall mobility is also increased and homogenized by ingot annealing. The increase of mobility is believed to be due to the increase of EL2, which shifts the Fermi-level towards the conduction band and increases the free electron concentration, resulting in a lowering of mixed conduction [7].

D

OPTICAL PROPERTIES

Effects of ingot annealing on the optical properties, especially on photoluminescence and cathodoluminescence, have been extensively studied by many authors. By looking at the near band-edge emission, mainly that of the carbon-related emission line at 1.49 eV, it was found that the uniformity can be easily detected and this technique is applied often in GaAs manufacturing [7,14,39-44]. It is known that the emission intensity is increased and the uniformity is improved by moderate temperature (800 - 9500C) annealing. The increase of the emission intensity however does not literally mean that the carbon concentration is increased. The increase may be based on the decrease of nonradiative recombination centres [14], possibly that of arsenic vacancy related point defects [44]. It is also known that the emission intensity becomes very small after high temperature (> 10000 C) annealing, about two orders of magnitude smaller than that after moderate temperature (800 - 9500C) annealing [44]. The mechanism of photoluminescence homogenization by ingot annealing is not yet fully clarified even after many investigations and further systematic and comprehensive study is desired. Cathodoluminescence (CL) is also a convenient method for examining the effect of annealing. This method has a higher resolution than photoluminescence and it is convenient when examining the microscopic uniformity. In fact, Miyazawa et al [45] showed that CL nonuniformity has a correlation with the threshold voltage variation of FETs. CL uniformity measurement is therefore a good, easy technique for evaluating the effect of annealing and precise studies have been performed by various authors. It is known that the uniformity can be improved by appropriate annealing [4,7,17,32] but it was also found that the microscopic nonuniformity is not fully eliminated by ingot annealing [33]. This implies that CL may include more information on defects other than EL2 [46]. E

DEFECT CONTROL

Annealing itself naturally affects defect concentrations in GaAs. Even though not all data are completely understood, the following defect behaviour has been clarified. (1)

Moderate temperature annealing homogenizes the distribution of EL2 defects which otherwise have a large inhomogeneity along dislocation cell boundaries [23]. This EL2 homogenization is the main reason for the improvement of resistivity uniformity.

(2)

High temperature annealing has a large effect in eliminating EL2 and arsenic precipitates. This phenomenon was expected as the reversible thermodynamical reaction implied from the phase diagram [23,28,29,34] and was experimentally proved by coulometric titration analysis [34].

(3)

After dissolution by high temperature annealing, EL2 and arsenic precipitates are regenerated during annealing. The formation of these defects is also explained as a thermodynamic phenomenon. The relationship between EL2 and arsenic precipitates has also been discussed and their kinetics during annealing have been studied [18,26].

(4)

The electrical properties are converted reversibly between semi-insulating and semiconducting by annealing [6]. It is known that semi-insulating properties cannot be retained when semi-insulating GaAs with low carbon concentration is cooled slowly after crystal growth or after annealing. Low temperature annealing at around 550 - 650 0 C has an effect in increasing the concentration of midgap donors and the resistivity is converted to the level of semi-conducting (103 - 106Q cm) when it exceeds the carbon concentration [9,19,31]. This phenomenon is now clearly explained from a Shockley diagram by the formation of EL6 defects [9,10,19].

F

EFFECT ON DEVICE PERFORMANCE

It was first pointed out that dislocations affect the threshold voltages of FETs [45]. Later studies clarified that the inhomogeneity of EL2 in the vicinity of dislocations is the main reason for the effect [23]. It was also found that ingot annealing is very effective in improving the threshold voltage variation since it homogenizes EL2 concentrations. Threshold voltage variations have been improved year by year by various methods [38,49-57] as summarized in FIGURE 1 [58] and the best variation of the level of 10 meV was achieved by appropriate annealing. High temperature ingot annealing was also found to increase the microwave device fabrication yield [59]. This is due to the reduction of arsenic precipitates which largely reduce the device fabrication yield when they occur under the gate electrode [60].

As Grown Ingot Annealing Wafer Anrwafing

In Doped

PHILIPS

Cr Doped MLEC P Impurity Implantation

tfVth(mv)

SONY

TOSHIBA; HONEYWELL

HITACHI

FIGURE 1. Reduction of threshold voltage variation (oVth) as a function of year [58].

G

CONCLUSION

After much effort on ingot annealing, it has now become an indispensable process for bulk semiinsulating GaAs manufacturing. Most effects of ingot annealing have been clarified but there are

still some ambiguities in the mechanism of annealing on photoluminescence and cathodoluminescence. The behaviour of defects other than EL2, EL6 and arsenic precipitates is also under question even now. Complete clarification of these behaviours is desired. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] [II] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29]

D. Rumsby, R.M. Ware, B. Smith, M. Tyjberg, M.R. Brozel, EJ. Foulkes [ Tech. Dig. GaAs IC Symp. (New York, IEEE, 1983) p.34 ] M. Yokogawa, S. Nishine, K. Matsumoto, H. Morishita, K. Fujita, S. Akai [ Inst. Phys. Conf. Ser. (UK) vol.74 (1984) p.29] D.E. Holmes, H. Kuwamoto, CG. Kirkpatrick, R.T. Chen [ Semi-Insulating UI-VMaterials Eds D.C. Look, J.S. Blakemore, (Shiva, England, 1984) p.204 ] A.K. Chin, I. Camlibel, R. Caruso, M.S.S. Young, A.R. Von Neida [ J. Appl. Phys. (USA) vol. 57 (1985)p.2203] J. Osaka, F. Huga, K.K. Watanabe [ Appl. Phys., Lett. (USA) vol.47 (1985) p. 1307 ] D. C. Look et al [ Appl. Phys. Lett. (USA) vol.49 (1986) p. 1083 ] K. Lohnert, W. Wettling, G. Koschek [ Semi-Insulaiting UI-VMaterials Eds H. Kukimoto, S. Miyazawa, (Ohmsha, Tokyo, 1986), p. 267 ] J. Lagowski, H.C. Gatos, CH. Kang, M. Skowronski, K.Y. Ko, D.G. Lin [ Appl. Phys. Lett. (USA) vol.49 (1986) p.892] Y. Kitagawara, N. Noto, T. Takahashi, T. Takenaka [ Semi-Insulating III-VMaterials Eds H. Kukimoto, S. Miyazawa (Ohmsha, Ltd., Tokyo, 1986) p.273 ] Y. Kitagawara, N. Noto, T. Takahashi, T. Takenaka [Appl. Phys. Lett. (USA) vol.43 (1986) p.788 ] O. Ogawa [ Semi-Insulating III-V Materials Ed H. Kukimoto, S. Miyazawa (Hakone, Ohmsha, 1986)p.237] Y. Otoki, M. Nakamori, R. Nakazono, S. Kuma [ Semi-Insulating III-V Materials Eds H. Kukimoto, S. Miyazawa (Hakone, Ohmsha, 1986) p.285 ] J. Osaka, H. Okamoto, F. Huga, K. Watanabe, K. Yamada [ Semi-Insulating III-V Materials Ed H. Kukimoto, S. Miyazawa (Hakone, Ohmsha, 1986) p.279 ] N. Noto, Y. Kitagawara, T. Takahashi, T. Takenaka [ Jpn. J. Appl. Phys. (Japan) vol.25 (1986) p.L394 ] CH. Kang, J. Lagowski, H.C Gatos [ J. Appl. Phys. (USA) vol.62 (1987) p.3482 ] H. Kanber, D.C Wang [ IEEE Trans. Electron Devices (USA) vol.34 (1987) p. 1626 ] S. Clark, DJ. Stirland, M.R. Brozel, M. Smith, CA. Warwick [ Semi-Insulating III-VMaterials (Malmo, Sweden, 1988) p.31 ] Y. Kitagawara, T. Takahashi, A. Kuwabara, T. Takenaka [ Semi-Insulating III-V Materials (Malmo, Sweden, 1988) p.49] S. Reichlmaier, K. Lohnert, M. Baumgartner [ Jpn. J. Appl Phys. (Japan) vol.27 (1988) p.2329 ] Y. Nakamura, Y. Ohtsuki, T. Kikuta [ Jpn. J. Appl. Phys. (Japan) vol.27 (1988) p.Ll 148 ] M.T. Asom, J.M. Parsey Jr., L.C Kimerling, R. Sauer, F.A. Thiel [ Appl Phys. Lett. (USA) vol.52 (1988) p. 1472] O. Ogawa [ Semi-Insulating III-V Materials (Malmo, Sweden, 1988) p. 477 ] S. Miyazawa, K. Watanabe, J. Osaka, K. Ikuta [ Revue Phys. Appl. {France) vol.23 (1988) p.727 ] T. Inada, Y. Otoki, K. Ohata, S. Taharasako, S. Kuma [J Cryst. Growth (Netherlands) vol. 96 (1989)p.327] S. Kuma, Y. Otoki, T. Inada [ Oyo Buturi (Japan) vol.91 (1989) p.327 ] B. Lee, E.D. Bourret, R. Gronsky, I. Park [ J Appl. Phys. (USA) vol.65 (1989) p. 1030 ] DJ. Stirland [ Defect Control in Semiconductors Ed K. Sumino (North-Holland, Amsterdam, 1990)p.783] E. Molva, Ph. Bunod, A. Chabli, A. Lombardot, S. Dubois, F. Bertin [ J. Cryst. Growth (Netherlands) vol.103 (1990) p.91 ] Y. Otoki, M. Watanabe, T. Inada, S. Kuma [ J Cryst. Growth (Netherlands) vol. 103 (1990) p.85 ]

[30] [31] [32] [33] [34] [35] [36] [37] [38] [39] [40] [41] [42] [43] [44] [45] [46] [47] [48] [49] [50] [51] [52] [53] [54] [55] [56] [57] [58] [59] [60]

S. Clark, MR. Brozel, DJ. Stirland, D.T.J. Hurle, I. Grant [ Defect Control in Semiconductors Ed. K. Sumino (North-Holland, 1990) p.807 ] M. Matsui, T. lino, Sumitomo [MetalMining, Denzaikengiho, vol.8 (1990) p. 17 ] H. Menninger, M. Beer, R. Cleichmann, H. Raidt, B. Ulrici, G. Voigt [ Phys. Status Solidi A (Germany) vol. 121(1990) p.95 ] M. Mori, G. Kano, T. Inoue, H. Shimakura, H. Yamamoto, O. Oda [ Semi-Insulating IU-V Materials (Adam Hilger, Toronto,1990) p. 155 ] O. Oda et al [ Semicond. Sd. Technol. (UK) vol.7 (1992) p.A215 ] M. Suemitsu, K. Terada, M. Nishijima, N. Miyamoto [ Jpn. J. Appl. Phys. (Japan) vol.31 (1992) P.L1654] N. Ohkubo, M. Shishikura, S. Matsumoto [ J. Appl. Phys. (USA) vol.73 (1993) p.615 ] T. Obokata, T. Matsumura, K. Terashima, F. Orito, T. Kikuta, T. Fukuda [ Jpn. J. Appl. Phys. (Japan) vol.23 (1984) p.L602 ] S. Asai, N. Goto, T. Nozaki [ IEICE Tech. Rep. CPM87 (1987) p.55 ] M. Yokogawa, S. Nishine, M. Sasaki, K. Matsumoto, K. Fujita, S. Akai [ Jpn. J. Appl. Phys. (Japan) vol.23 (1984) p.L339 ] HJ. Hovel, M. Albert, E. Farrell, D. Guidotti, J. Becker [ Semi-Insulating IH-VMaterials (Ohmsha, 1986) p.97 ] S. Shinzawa, A. Hattori, Y. Otoki, S. Okubo [ Hitachi Cable Review (Japan) (1987) p.41 ] P.W. Yu, D.C. Look, W. Ford [ J. Appl. Phys. (USA) vol.62 (1987) p.2960 ] V. Swaminathan, R. Caruso, SJ. Pearton [ J. Appl. Phys. (USA) vol.63 (1988) p.2164 ] T. Inoue, M. Mori, G. Kano, H. Yamamoto, O. Oda [ Inst. Phys. Conf. Ser. (UK) vol. 112 (1991) p.219] S. Miyazawa, Y. Ishii, S. Ishida, Y. Nanishi [ Appl. Phys. Lett. (USA) vol.43 (1983) p.853 ] G. Koschek, H. Lakner, E. Kubalek [ Phys. Status Solidi A (Germany) vol. 108 (1988) p.683 ] I. Priecel, L. Duricek, V. Daniska, D. Korytar [ J. Cryst. Growth (Netherlands) vol.126 (1993) p. 103] Y. Ishi, S. Miyazawa, S. Ishida [ IEEE Trans. Electron. Devices (USA) vol.ED-31 (1984) p.800 ] S. Murai, K. Tada, S. Akai, T. Suzuki [ Oyo Buturi (Japan) vol.53 (1984) p. 1083 ] J. Kasahara, M. Arai, N. Watanabe [ Jpn. J. Appl. Phys. (Japan) vol.25 (1986) p.L85 ] Y. Tanaka, T. Matsumura, F. Shimura, M. Kanamori, A. Higashisaka [ Extended Abstracts of thellth Conf. Solid Devices andMaterials (1985) p.433 ] J. Maluenda, G. M. Martin, H. Schink, G. Packeiser [ Appl. Phys. Lett. (USA) vol.48 (1986) p. 715 ] I. Terashima, T. Nakanishi [ Nikkei Electron. (Japan) vol. 11 (1986) p.99 ] T. Egawa, Y. Sano, H. Nakamura, T. Ishida, K. Kaminishi [ Jpn. J. Appl. Phys. (Japan) vol.25 (1986)p.L973] K.L. Tan, H. Cung, CH. Chen [ IEEE Electron. Device Lett. (USA) vol.EDL-8 (1987) p.440 ] Y. Fujisaki, Y. Takano [ Semi-InsulatingUl-VMaterials (Malmd, Sweden, 1988) p.247 ] M. Oyake, H. Yamamoto, G. Kano, T. Inoue, O. Oda [ Inst. Phys. Conf. Ser. (UK) vol. 112 (1991) p.311] O. Oda, H. Yamamoto, K. Kainosho, T. Imaizumi, H. Okazaki [Inst. Phys. Conf Ser. (UK) vol. 135 (1993)p.285] S. Martin, P. Suchet, G.M. Martin [ Proc. 6th Conf. on Semi-Insulating HI-VMaterials (Adam Hilger, Toronto, 1990) p.l ] H. Yamamoto, O. Oda, M. Seiwa, M. Taniguchi, H. Nakata, M. Ejima [ J. Electrochem. Soc. (USA) vol.136 (1989) p.3098 ]

15,6 Heat treatments of GaAs wafers O. Oda January 1996

A

INTRODUCTION

Wafer annealing is the newest annealing procedure for improving material quality and uniformity. It is known that conventionally grown LEC-GaAs has many tiny arsenic precipitates, due to nonstoichiometry. This is inevitable because the congruent point of the GaAs phase diagram deviates toward the arsenic-rich side. Since these arsenic precipitates were found to affect the device properties and the fabrication yield, it was desired to find a new material fabrication process for reducing these arsenic precipitates. For this purpose, wafer annealing has been studied as a technology replacing ingot annealing. B

ARSENIC PRECIPITATES

Observation of precipitates has been studied by many authors [1-15] as described in Datareview 10.9. It was also found that these precipitates can greatly affect device performances [16-19]. The reduction of these precipitates therefore became indispensable for the development of GaAs electronic devices. The reason that these arsenic precipitates exist in LEC-GaAs was speculated to be the precipitation of arsenic due to the deviation of the congruent-point to the arsenic-rich side [20,21]. This speculation also came from an assumption that EL2 itself is related to arsenic precipitation [20]. In fact, EL2 and arsenic precipitates can be eliminated by high temperature annealing [22-24]. The fact that the congruent-point deviates to the arsenic-rich side was finally experimentally proved by annealing experiments and by coulometric titration analysis [25]. Because of this deviation, the appearance of arsenic precipitates is understood to be inevitable for as-grown and ingot-annealed LEC-GaAs. C

ANNEALING PROCEDURE

Wafer annealing was first applied to as-cut wafers for improving the electrical properties as in the case of ingot annealing [26-30]. This annealing was performed under similar conditions to ingot annealing at moderate temperatures (800 - 9500C), and was found to have a similar effect to ingot annealing. Since there was no specific advantage compared with ingot annealing, this moderate temperature wafer annealing was then neglected because wafer-annealed material was less uniform than ingot-annealed GaAs. Wafer annealing was later considered as a technology to overcome the disadvantages of ingot annealing. Mori et al [31] developed a multiple-step wafer annealing (MWA) where as-cut wafers are first annealed at high temperatures (>11000C) and then annealed at moderate temperatures (950 0 C). As-cut wafers are slightly etched before being annealed in quartz ampoules with an appropriate arsenic overpressure at temperatures higher than 11000C. This process was designed to control the material composition and return it to stoichiometry based on the arsenic exchange between the material and the vapour phase [25]. After this process the GaAs has poor uniformity

so that it is necessary to reanneal it at moderate temperatures, as in the case of ingot annealing, for homogenizing EL2 distributions. By this MWA process, it was found that the material was close to stoichdometry with fewer arsenic precipitates and with a good resistivity uniformity of ±5% and can be obtained reproducibly [25]. Wafer annealing was then applied to high-purity CVD epitaxial wafers in order to convert conductive epitaxial wafers to semi-insulating ones [32,33]. It was found that semi-insulating properties can be obtained for lower EL2, carbon and impurity concentrations than those in bulk material. It was also found that the concentrations of various point defects such as EL3, EL5 and EL6 can be largely reduced by this type of wafer annealing [33]. D

ELECTRICAL PROPERTIES

Mobility

Resistivity

(fi-cm)

(cm 2 /V-s)

The first step in wafer annealing was found to largely affect the electrical properties as functions of both temperature and the arsenic vapour pressure. When the temperature is increased and the arsenic pressure is decreased, the semi-insulating properties are lost and p-type conductive material is obtained [25,31] (FIGURE 1).

As Vapor Pressure (atm.)

As Vapor Pressure (atm.)

FIGURE 1. Resistivity and mobility as a function of arsenic vapour pressure during wafer annealing [25].

This p-type conversion behaviour is very similar to the case of the p-type conversion as a function of melt composition in crystal growth [34]. The p-type conversion is believed to correspond to non-stoichiometry, resulting from a decrease of EL2 concentration. It was found that by optimizing annealing conditions, semi-insulating properties can be realized with the composition very close to stoichiometry, a composition which cannot be achieved by ingot-annealing procedures [25].

Uniformity of electrical properties can be improved by moderate temperature annealing at 950 0 C as in the case of ingot annealing. This uniformity comes from the homogenization of EL2 as in the case of ingot annealing. This multiple-step wafer annealing finally gives the same uniformity of electrical properties as the best ingot annealing [25]. E

OPTICAL PROPERTIES

The annealing condition affects the uniformity of photoluminescence [35] as in the case of ingot annealing. By wafer annealing, a similar uniformity of photoluminescence can be achieved to that in ingot annealing [35]. Another feature is that CL uniformity can be largely improved in the case of MWA compared with ingot annealing [25,31]. In the case of ingot annealing, the complete homogenization of the microscopic nonuniformity of cathodoluminescence cannot be achieved, while in the case of wafer annealing, it can be completely homogenized. The reason is not yet clarified, but it seems to be due to some defect other than EL2, which is the origin of cathodoluminescence nonuniformity. F

EFFECT ON DEVICE PERFORMANCE

The effect of wafer-annealed GaAs on device performance has been examined in various cases. Oyake et al [36] proved that the threshold voltage variation can be improved by using MWA wafers compared to ingot annealing. It was also found that the device fabrication yield can be improved for integrated circuits application [37]. This is the most expected effect of wafer annealed material since in the case of large scaled ICs, a large number of inclusions, such as arsenic precipitates, will significantly deteriorate the device fabrication yield as in the case of silicon oxide in silicon ICs. This advantage of wafer-annealed GaAs becomes prominent for the case of large scale devices rather than discrete devices. Another feature of wafer-annealed material is that it was found that it improves the characteristics and the morphology of epitaxial layers, especially in the case of MBE grown AlGaAs epitaxial layers [38,39]. When wafer-annealed material is used for epitaxial growth, tiny oval defects due to arsenic precipitates along dislocations, which differ from conventional oval defects due to gallium oxides, can be significantly reduced which naturally improves the device fabrication yield. G

CONCLUSION

A wafer annealing procedure was developed which overcomes the disadvantages of ingot annealing. It has advantages in controlling the stoichiometry, improving the uniformity, increasing the IC device fabrication yield and improving the epitaxial morphology. Wafer annealing is a technology which has more controllability than ingot-annealing since it is based on the constituent exchange process between the solid and the vapour. Further scientific studies for clarifying the precise mechanism of wafer annealing are therefore desired. REFERENCES [1] [2] [3] [4]

DJ. Stirland, B.W. Straughan [ 772/« Solid Films (Switzerland) vol.31 (1976), p.139 ] A.G. Cullis, P.D. Augustus, DJ. Stirland [J. Appl Phys. (USA) vol.51 (1980) p.2556 ] P.D. Augustus, DJ. Stirland [J. Microsc. (UK) vol.118 (1980) p i l l ] DJ. Stirland, P.D. Augustus, M.R Brozel, EJ. Foulkes [ Semi-Insulating HI-VMaterials Kah-Nee-

[5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33] [34] [35] [36]

Ta (Shiva, Nantwich, UK, 1984) p.68 ] J. P. Cornier, M. Duseaux, J.P. Chevalier [ Appl. Phys. Lett. (USA) vol.45 (1984) p. 1105 ] EA. Lodge, GR. Booker,CA. Warwick, GT. Brown [Inst. Phys. Conf. Ser. (UK) vol.76 (1985) ] T. Ogawa [ Defect Recognition adn Image Processing Ed J.P. Fillard (Elsevier, 1985) p. 11 ] P. Suchet, M. Duseaux, S. Gillardin, J. Le Bris, G.M. Martin [ J. Appl. Phys. (USA) vol.62 (1987) p.3700 ] P. Suchet, M. Duseaux, [ Inst. Phys. Conf. Ser. (UK) vol.91 (1988) p.375 ] O. Ogpwa[Semi-InsulatingHI-VMaterialsMalmo, Sweden, 1-3 June 1988 (AdamHilger, Bristol, UK, 1988) p.477 ] P. Suchet, M. Duseaux, C. Schiller, G.M. Martin [ Semi-Insulating IH-VMaterials Malmo, Sweden, 1-3 June 1988 (Adam Hilger, Bristol, UK, 1988) p.483-8 ] P. Gall, J. P. Fillard, M. Castagne, J. L. Weyer, J. Bonnafe [ J. Appl. Phys. (USA) vol.64 (1988) p.5161] B-T. Lee, R. Gronsky, E.D. Bourret [ JCryst. Growth (Netherlands) vol.96 (1989) p.333 ] P. Suchet, M. Duseaux, C. Schiller, G.M. Martin [ Semi-Insulating IU-VMaterials Malmo, Sweden, 1-3 June 1988 (Adam Hilger, Bristol, UK, 1988) p.l ] Y. Otoki, M. Watanabe, T. Inada, S. Kuma [ J. Cryst. Growth (Netherlands) vol. 103 (1990) p. 85 ] S. Martin, M. Duseaux, M. Erman [ Inst. Phys. Conf. Ser (UK) vol.74 (1984) p.53 ] S. Martin, P. Suchet, G.M. Martin [ Proc. 6th Conf. on Semi-Insulating IH-VMaterials Toronto, Canada, 13-16 May 1990 (Adam Hilger, Bristol, UK, 1990) p.l ] H. Yamamoto, O. Oda, M. Seiwa, M. Taniguchi, H. Nakata, M. Ejima [ J. Electrochem. Soc. (USA) vol.136 (1989) p.3098 ] H. Yamamoto, H. Shimakura, G. Kano, M. Seiwa, H. Nakata, O. Oda [ Defect Control in Semiconductors Ed. K Sumino (North-Holland, 1990) p. 1273 ] S. Miyazawa, K. Watanabe, J. Osaka, K. Ikuta [ Rev. Phys. Appl. (France) vol.23 (1988) p.727 ] E. Molva, Ph. Bunod, A. Chabli, A. Lombardot, S. Dubois, F. Bertin [ J. Cryst. Growth (Netherlands) vol.103 (1990) p.91 ] J. Lagowski, H.C. Gatos, CH. Kang, M. Skowronski, K.Y. Ko, D.G. Lin [ Appl. Phys. Lett. (USA) vol.49 (1986) p.892 ] S. Clark, DJ. Stirland, M.R. Brozel, M. Smith, CA. Warwick [ Semi-Insulating IU-VMaterials Malmo, Sweden, 1-3 June 1988 (Adam Hilger, Bristol, UK, 1988) p.31 ] T. Inada, Y. Otoki, K. Ohata, S. Taharasako, S. Kuma [ J. Cryst. Growth (Netherlands) vol.96 (1986)p.327] O. Oda et al [ Semicond. Sci. Technol (UK) vol.7 (1992) p.A215 ] S. Miyazawa, T. Honda, Y. Ishii, S.Ishida [Appl. Phys. Lett. (USA) vol.44 (1984) p.410 ] Y. Matsumoto, H. Watanabe [ Jpn. J. Appl. Phys. (Japan) vol.21 (1982) p.L515 ] S. Asai, K. Ohata [ EICS Conference (1989) vol.ED89-109 ] K. Lohnert, W. Wettling, G. Koschek [ Semi-Insulating HI-VMaterials Eds. H. Kukimoto, S. Miyazawa (Hakone, Ohmsha,1986) p. 267 ] HJ. HoveL M. Albert, E. Farrell, D. Guidotti, J. Becker [ Semi-Insulating IH-Vmaterials (Hakone, Ohmsha, 1986) p.97] M. Mori, G. Kano, T. Inoue, H. Shimakura, H. Yamamoto, O. Oda [ Semi-Insulating Ul-V Materials Toronto, Canada, 13-16 May 1990 (Adam Hilger, Bristol, UK, 1990) p. 155 ] T. Imaizumi, H. Okazaki, O. Oda [ Appl. Phys. Lett. (USA) vol.63 (1993) p. 1390 ] T. Imaizumi, H. Okazaki, O. Oda [ J. Appl. Phys. (USA) vol.76 (1994) p.7957 ] D.E. Holmes, R.T. Chen, K.R. Elliott, CG. Kirkpatrick [ Appl. Phys. Lett. (USA) vol.57 (1990) p.2464 ] T. Inoue, M. Mori, G. Kano, H. Yamamoto, O. Oda [ Inst. Phys. Conf. Ser. (UK) vol. 112 (1991) P-219] M. Oyake, H. Yamamoto, G. Kano, T. Inoue, O. Oda [ Inst. Phys. Conf Ser. (UK) vol. 112 (1991) p.311]

[37] [38] [39]

O. Oda, H. Yamamoto, K. Kainosho, T. Imaizumi, H. Okazaki [ Semi-Insulating III-V Materials Warsaw, Poland (World Scientific, 1994) p.27 ] G. Kano, M. Takakusaki, M. Oyake, H. Yamamoto, O. Oda [ Semi-Insulating III-V Materials Toronto, Canada (Adam Hilger, 1990) p.59 ] O. Oda,H. Yamamoto,K.Kainosho,T.Imaizumi,H. Okazaki[Inst. Phys. Conf Ser. (UK)\oll35 (1993)p.285]

15.7 Carrier concentrations of semi-insulating GaAs D.C. Look December 1995

A

INTRODUCTION

By far the most dominant form of semi-insulating (SI) GaAs in use today is that grown with no doping but with an excess of As in the starting materials. Typically, undoped, SI GaAs ingots are produced by one of the following three growth techniques: (1) liquid-encapsulated Czochralski (LEC); (2) vertical gradient freeze (VGF); or (3) horizontal Bridgman (HB). All three techniques produce n-type material, as measured by the Hall effect, with electron concentrations usually ranging from 5 x 106 to 2 * 108 cm"3, and mobilities ranging from 6000 - 8000 cm2/ Vs. Interestingly, a recent study often commercial, 4-inch high-pressure LEC boules gave an average n = 2.2 x 107 cm"3, while the averages for ten low-pressure LEC boules, and ten VGF boules, were 2.0 x 107 cm"3 and 2.8 * 107 cm"3, respectively. Typical interboule standard deviations were 40 - 60% in each case. B

USEFUL FORMULA

For temperatures up to about 500 K, the carrier concentration, n, in undoped, SI GaAs is well described by n =

(^r) -

N

c ' ea/k

T3/2

CX

P ("EDO

'kT)

(D

where N^ 2 is the EL2 concentration, NA is the concentration of acceptors lying below EL2 in the bandgap, gx is the degeneracy of the singly-occupied, upper s-like EL2 state, g2 is the degeneracy of the doubly-occupied state, NC'T3/2 is the effective conduction-band density of states, and a is a temperature coefficient of the EL2 (0/+) transition energy, defined by ED = E 0 0 - ccT [I]. From temperature-dependent Hall-effect measurements, it is known that E 0 0 = 0.75 ± 0.01 eV, and near room temperature, N c ' « 8.1 x 1013 cm "3 K"3/2. Also, for an s-like state which can hold two electrons, gi = 2 and g2 = 1. Duncan and Westphal [2] have carried out a multi-boule study in which they varied NA by doping with carbon. Since the carbon concentration can be conveniently measured by local vibrational mode (LVM) absorption spectroscopy [3], the EL2 concentration by near-infrared absorption spectroscopy [4], and the carrier concentration by Hall-effect measurements [1], it is possible to determine a in EQN (1) (with the assumption NA = [C]). Duncan and Westphal [2] found that a - 3.4 x 10"4 eV K"1, giving E 0 (296 K) = 0.65 eV. However, they used a wrong degeneracy factor in EQN (1), namely ^g1 = 1/2 instead of gx/g2 = 2. Thus, the correct value of a should have been 2.2 x 10"4 eV K"1, giving En (296 K) = 0.685 eV. This latter value of a is close to that of Martin et al [5], who determined a ~ 2.4 x 10"4 eV K"1 from a deep level transient spectroscopy (DLTS) study.

Therefore, with a fair degree of confidence, we can write n = 2.0 x 1015

—— - 1

{

N

A

T 3 / 2 exp (-0.75 / kT)

(2)

J

where k = 8.617 x 10"5 eV K"1 and n is in units of cm"3. Typical values of N E L 2 and N A are 1 x 10 1 6 cm"3 and 1 x 10 15 cm"3, respectively, giving n = 1.5 x 10 7 cm"3. Practically speaking, it is relatively easy to measure n, by Hall effect, and № EL2 > by 1.1 |im wavelength infrared absorption. (Note that (NgL2/ N A - 1 ) = № E L 2 / N A , where № E L 2 is the neutral NEL 2 concentration). On the other hand, the measurement of N A , either by LVM (if it is assumed N A = [C]), or by two-wavelength absorption [6], is either time consuming or inaccurate. Thus, EQN (2) can profitably be used to determine N A by the relatively easy measurements of n and N 0 EL 2 .

As a final note, the other form of SI GaAs, namely Cr-doped GaAs, should be mentioned. In general, carrier concentrations are smaller for this material; in fact, mixed-conduction effects become important because p > n even though the Hall coefficient is still negative (i.e., n^ n 2 > pjip2). Since Cr-doped GaAs is no longer widely used, we will not discuss it further here; however, reference [7] contains detailed information as well as many references on this material. REFERENCES [1] [2] [3] [4] [5] [6] [7]

D.C. Look [ Electrical Characterization of GaAs Materials and Devices (John Wiley, New York, 1989) p. 117] W.M. Duncan, G.H. Westphal [ in GaAs and Related Compounds, Las Vegas, USA, 1986 (The Institute of Physics, Bristol, 1987) p.39 ] H. Ch. Alt, B. Dischler [Appl Phys. Lett. (USA) vol.66 (1995) p.61 ] P. Silverberg, P. Omling, L. Samuelson [ Appl Phys. Lett. (USA) vol.52 (1988) p. 1689 ] G.M. Martin, J.P. Farges, G. Jacob, J.P. Hallais [ J. Appl. Phys. (USA) vol.51 (1980) p.2840 ] S.K. Brierley, D.S. Lehr [Appl. Phys. Lett. (USA) vol.55 (1989) p.2426 ] D.C. Look [ Semicond Semimet. (USA) vol. 19 (1983) p.75 ]

CHAPTER 16 EPITAXIAL GROWTH OF GaAs 16.1 16.2 16.3 16.4 16.5 16.6 16.7 16.8 16.9

Introduction to epitaxy LPE growth of GaAs Halogen transport VPE growth of GaAs and related compounds Growth by close space vapour transport MOVPE growth of GaAs Unintentional hydrogen doping in MOCVD growth High temperature MBE growth of GaAs Metalorganic molecular beam epitaxy and atomic layer epitaxy of GaAs Epitaxial GaAs lift-off

16.1 Introduction to epitaxy R. Kaspi October 1995

A

INTRODUCTION

The process by which a layer of material, newly deposited over a substrate, forms a monocrystalline structure is referred to as epitaxy. GaAs layers that are epitaxially grown on properly prepared GaAs substrates generally exhibit a near-perfect transition from the substrate to the epilayer where the crystalline structure and orientation of the substrate is retained. Epitaxial growth of GaAs and related compounds is essential to building complex device structures because it allows the formation of high quality layers with controlled thickness and conductivity as well as abrupt interfaces. Epitaxy of GaAs and other III-V alloys is possible by several different techniques which are described in the Datareviews that follow. These techniques can be classified by the phase and chemistry of the reactants directed toward the substrate. For example, during molecular beam epitaxy (MBE) elemental evaporants are directed onto the substrate in ultra high vacuum, during metal organic chemical vapour deposition (MOCVD) gas phase organic and hydride compounds are pyrolized at the hot substrate surface, and during liquid phase epitaxy (LPE) a supersaturated solution is solidified over a cooler substrate. While each has its advantages and disadvantages, near-monolayer control of thickness and the ability to form abrupt interfaces have made MBE and MOCVD largely responsible for the production and study of complex epitaxial structures. During the last decade, there have been numerous advances in epitaxial technologies as well as an improved understanding of the surface processes leading to epitaxy. This has paved the way for the research and the synthesis of a new breed of complex GaAs-based devices which may include two-dimensional electron gas heterostructures, quantum wells or superlattices. B

EPITAXIAL GROWTH

Much of our current but limited understanding of the epitaxial process during growth of GaAs from the vapour phase comes from MBE which has allowed the in-situ probing of the interaction of simple chemical species with the film surface in an ideally suited ultrahigh vacuum environment [1,2]. Recent analyses of MOCVD grown film surfaces do show a strong similarity in the evolution of epitaxial layers when compared to MBE [3]. Much of the discussion below is limited to growth from the vapour phase (eg. MBE, MOCVD). Bl

GaAs Substrates

Substrates with low index plane surfaces such as (001), (111)B, or (110) are used to provide atomically smooth starting surfaces for epitaxy. The majority of epitaxially grown GaAs-based devices are deposited on (001) or near-(001) oriented GaAs substrates. Substrates with 50, 50.8, 76, and 100 mm diameter exhibiting n-type, p-type or semi-insulating properties are routinely available through commercial vendors at this time, while 150 mm diameter semi-insulating substrates have recently entered the market. Improvements in bulk growth techniques have led

to average etch pit density (EPD) specifications of <103/ cm2 for semi-insulating, and <20/ cm2 for Si-doped substrate wafers. Improved surface polish and cleaning procedures have enabled substrate vendors to market 'epi-ready' wafers which do not necessitate additional surface preparation for most applications. A thin oxide layer formed on GaAs substrates serves as a protection from atmospheric contamination before epitaxial growth. The oxide layer is removed immediately prior to epitaxy in order to expose a surface which is sufficiently smooth and clean for epitaxial growth. The surface of a clean substrate, as any epitaxial film surface, is characterized by atomic steps separated by smooth terraces. The nominally flat GaAs (001) surface contains two types of stable atomic steps, those that are parallel to the [TlO] direction (A-type, Ga-terminated) and those that are parallel to the [110] direction (B-type, As-terminated). Generally, a periodicity in the step structure exists due to a small deviation of the surface plane from the low index plane. This is either unintentional (< ±0.5°), or due to the intentional miscut by an angle a in the range of 1-15° toward a specific azimuth such as the [110] to expose A-type steps or the [TlO] to expose B-type steps. The average terrace width L on the GaAs surface can be adjusted by varying the miscut angle as L = 5.6533 A/ (2 tan a). B2

Stoichiometry

Stoichiometric growth is a requirement for minimizing the point defect density in epilayers. For growth of GaAs from the vapour phase, a precise control of the incident fluxes is not necessary for obtaining stoichiometric films. During MBE, for example, stoichiometric growth of GaAs is possible because when adsorbed onto the film surface at 300° to 650 0 C, gallium is essentially involatile in any form. Arsenic is vastly more volatile than gallium if unincorporated but fairly involatile once incorporated into GaAs. Therefore, under excess arsenic pressure, only as much arsenic needed to form stoichiometric GaAs is incorporated into the film while the excess is desorbed [4,5]. A similar mechanism is active in MOCVD within the transport limited regime between approximately 5500C and 8000C [6]. B3

Surface Morphology

During epitaxy from the vapour phase, adsorbates at the film surface interact with the surface potential and with each other to yield a monocrystalline arrangement, but the formation of flat, defect-free epilayers is not always assured. The surface migration of adatoms and admolecules is a key element in this process. This lateral motion is somewhat random and can only be influenced through indirect means such as substrate temperature, incident fluxes, starting surface etc. A change in any one of these parameters can substantially vary the resulting surface morphology and directly influence heterojunction interface roughness which affects device quality. Moreover, when GaAs layers are deposited under conditions which yield a smooth surface morphology, point defect densities are reduced and the electrical and optical properties are improved [7]. The GaAs surface morphology is ultimately dependent on the ability of Ga, the less mobile species, to be accommodated at existing surface steps. In the 'step-flow' growth mode, the Ga surface migration distance is well above the average terrace width, the surface step density can be indefinitely maintained and a smooth surface morphology can be achieved [8]. On the other

hand, when the Ga surface migration distance is insufficient, island formation and coalescence on existing terraces occur. This eventually leads to thickness dependent surface roughness and the formation of large regular mounds [9]. Step-flow growth can be facilitated by increasing the Ga surface migration distance by elevating the growth temperature or reducing the growth rate [10]. Reducing the incident VAII ratio in both MOCVD and MBE [11] is also known to increase the surface mobility of Ga. Flow-rate modulation epitaxy (FME) [12] during MOCVD and migration enhanced epitaxy (MEE) [13] during MBE are extensions of the latter where the alternate supply of group-Ill and group-V components is used to increase the migration distance by allowing Ga to diffuse in a temporarily As-poor environment. Additionally, intentionally miscut vicinal surfaces can be used to provide shorter terrace widths with a narrower size distribution to help promote step flow growth [14]. Finally, interrupting the growth, perhaps in combination with an annealing step, generally results in surface smoothing driven by surface diffusion [15]. The observation of rectangular two-dimensional islands nucleated on (001) singular surfaces in both MBE and MOCVD (though the direction of the elongation is opposite [16]) suggests an anisotropic diffusion coefficient and/or a large sticking probability difference between stable steps oriented along [110] and [TlO]. Moreover, the bunching of monolayer steps on vicinal surfaces has been observed [17] and modelled [18], suggesting that the migrating species incorporate more easily on descending steps rather than ascending steps. The surface reconstruction is also expected to influence the surface migration of adatoms [19]. Experiments carried out to study surface migration during epitaxial growth have generally only yielded macroscopic parameters due to the difficulty of isolating reactions at the steps from intrinsic migration. Nevertheless, the Ga surface diffusion length along the [010] direction on the As-stabilized (001) plane was measured to range between -0.2 |iim at 500 0C and ~1 ^m at 65O0C and was observed to be strongly V/III ratio dependent [20]. A variety of theoretical works have also been carried out on GaAs epitaxy using the Burton-Cabrera-Frank theory [21], molecular dynamics [22], and Monte Carlo simulations [23,24]. B4

The Formation of Abrupt Interfaces

In order to form smooth and abrupt interfaces between regions of different doping or different composition, all epitaxial techniques must address four distinct requirements. The first requirement is that the bottom layer should have a smooth surface as the interface formed between it and the overlayer will most likely conform to this morphology. For example, AlGaAs layers deposited on GaAs generally form a smoother interface (normal interface) than GaAs layers deposited on AlGaAs (inverted interface) due to the tendency of AlGaAs to form rougher surfaces. The second requirement is the ability to change the rate of arrival of the source materials onto the substrate at a rate which is relatively fast compared to the growth rate. During solid source MBE of AlGaAs, for example, group-Hi and dopant fluxes (e.g. Si and Be) can be switched on and off more rapidly than the time required to deposit a monolayer, whereas it is more time consuming to switch the group-V flux on or off or to vary the magnitude of any flux originating from a single effusion cell.

The third requirement is that the surface lifetime of all species (the time to be incorporated into the crystal) is short relative to the growth rate. Species which tend to extend their surface lifetime by segregating and remaining at the film surface can be incorporated into the film even after the external supply is terminated, causing a compositional broadening of the interface. During MBE growth of GaAs doped with Sn, for example, significant incorporation of Sn occurs even after the incident Sn flux is eliminated [25]. The final requirement is minimal interdiffusion across the interface during the remainder of the time in which the substrate is held at the growth temperature. While group-Ill interdiffusion is negligible, both Si [26] and Be [27] have been observed to diffuse across the interface during MBE and MOCVD growth of GaAs. C

RECENT ADVANCES IN EPITAXY

Cl

Advances in Epitaxial Growth

Further developments in in situ atomic scale surface analytical techniques have helped to provide a more complete understanding of the GaAs surface morphology and the epitaxial process. Scanning tunnelling microscopy (STM) has been used to observe a number of important GaAs surface features in real space. Among other things, large anisotropies in the two-dimensional island shape [28], the effect of annealing on the step edge morphology [29], the various step edge structures [30], step flow growth on vicinal surfaces [31], the structures of the As-rich (2><4) reconstructions [32], and the effect of As coverage have been studied [33]. In situ scanning electron microscopy was used to observe continuous nucleation of two-dimensional GaAs islands [34], and to estimate the migration distance of Ga atoms on a Ga-stabilized GaAs (001) surface [35]. Epitaxial growth of structures containing pseudomorphic (strained but without dislocations) layers has allowed many devices to outperform their earlier lattice-matched versions. Much effort has been dedicated to understanding the limits to pseudomorphic growth and the onset of strain relief [36,37], to optimizing growth parameters for pseudomorphic growth, and to repair the damage within the crystal after strain relaxation [38,39]. The use of surfactants which alter the kinetics of surface processes to promote two-dimensional growth of strained layers is an exciting new development [40]. Regrowth of layers following partial, ex situ, device processing has become a feasible alternative to facilitate the fabrication of complex devices. Temporary protection of the growth surface with As [41] and In [42] has been successful. Advances have also been made in combining film growth and device fabrication steps in an ultrahigh vacuum environment, paving the way for in situ processing of GaAs/AlGaAs devices [43]. Quantum wire and quantum dot structures where the charge carriers are confined in two or three dimensions, respectively, are expected to exhibit new physical properties and have received considerable attention in recent years. An improved understanding of the epitaxial process as it is affected by well defined surface features has fuelled the development of a variety of growth schemes which employ MBE or MOCVD on patterned substrates [44,45]. An example of this is given by Shimomura et al who have introduced group-Ill fluxes at a glancing angle during GaAs/AlGaAs multilayer growth in MBE so that (111)B facets are preferentially formed due to

self-shadowing on channelled substrates with mesa stripes along [011]. Subsequent growth on the (111)B facet after reorienting the substrate was shown to produce well defined quantum wire structures along the T-junction [46]. Advances have been made in selective epitaxy where preferential deposition of crystalline layers within windows that are lithographically defined on the substrate takes place [47]. For example, GaAs substrates masked by a SiO2 film with square windows was used to grow nanometre-scale GaAs quantum dot structures using MOCVD [48]. While a fundamental understanding of the selective growth process and edge effects is still lacking, monolithic integration of electronic and photonic devices using this technology is promising. Cl

Advances in Growth Hardware

Rapid flux grading in MBE has been addressed by the introduction and flow control of gas sources in gas-source molecular beam epitaxy (GSMBE) and metalorganic molecular beam epitaxy (MOMBE). These techniques have also facilitated carbon doping of GaAs [49] and the growth of phosphorus containing III-V layers [50] similar to MOCVD. A variety of in situ thin film monitoring techniques to reduce the dependency of composition and thickness control on pregrowth calibrations have been recently introduced. Most of these techniques are implemented in MBE due to its more accommodating ultra high vacuum environment, while a few have also been implemented in MOCVD. In general, these techniques measure either the optical properties of the deposited layers [51,52], the surface chemistry of the deposited layers [53], the incident fluxes [54], or the desorbed fluxes [55,56], so that thickness, composition, or temperature of the deposited layers can be inferred in real time. It is likely that a combination of techniques holds the most promise for the future. The desire to increase the throughput of epitaxial layers has led to the design and construction of scaled up MOCVD and, more recently, MBE systems that can provide multiwafer growth. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] [II] [12] [13]

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16.2 LPE growth ofGaAs M.G. Astles September 1996

A

INTRODUCTION

Although the process of liquid-phase epitaxy (LPE) has been attributed to work in the 19th century on the growth of sodium nitrate on to calcite crystals by Frankenheim [1], it was during the 1960s that the technique became popular, reflecting the increasing interest in the III-V compounds, particularly GaAs and GaP, for optoelectronic and microwave devices. While much of the early GaAs device work was carried out on bulk crystals grown by the Bridgman or liquidencapsulated Czochralski method, the potential of thin film epitaxial growth was soon appreciated, the principal advantages being: (i) (ii) (iii)

lower growth temperature, Tg, leading to reduced levels of native defects and impurities improved control and uniformity of electrical properties potential for growing p-n doped homojunctions (GaAs/GaAs) or hetero-junctions (GaAsZGaLxAlxAs).

Liquid-phase epitaxy as applied to HI-V compounds such as GaAs involves the dissolution of the material to be grown (the solute, GaAs) in a suitable solvent (normally the group III metal, Ga) at elevated temperature to form a saturated solution. Slowly cooling the solution below the liquidus (equilibrium) temperature (TL) induces supersaturation and if a suitable single crystal GaAs substrate is placed in contact with the solution, epitaxial growth occurs. At the end of the growth period, t, the solution is removed from the grown layer. The necessary liquid-solid phase diagrams for all the common III-V binary systems (GaAs, InP, GaP, InAs etc.) and ternary alloy systems (GaAlAs, GaInAs, GaInP etc.) have been extensively studied both experimentally and theoretically using regular solution models. An excellent review of the phase diagram information for EI-V binary and ternary systems was written by Panish and Ilegems in 1972 [2] and although limited regions of some of the diagrams have been researched since then, most of this earlier data still stands. The data for the Ga-As system from several authors are in excellent agreement over a wide range of temperatures [3-5]. The theoretical basis of the LPE growth process using the diffusion-limited growth model has been well established [6-8]. This explains well the observed dependences of layer thickness, d, on growth time, t, for various conditions of supersaturation. For step-cooled growth, where the solution is cooled to a fixed temperature ATS below TL before contacting the substrate, d is proportional to Vt, while for ramp-cooled growth where the solution is contacted with the substrate at the liquidus temperature and cooled at a constant rate, d is proportional to t3/2 [9]. Although LPE was initially the most widely used epitaxial technique for the growth of GaAs and the alloys of Ga1^AlxAs attention soon turned to other techniques such as the AsCl3-Ga-H2 (chloride) and AsH3-Ga (hydride) methods, molecular beam epitaxy (MBE) or metal-organic vapour phase epitaxy (MOVPE) where better control of layer thickness, improved surface morphology and better scalability to production volumes were offered. These other growth

techniques are covered elsewhere in this book. Earlier reviews of LPE growth of III-Vs include those of Dawson [10] Benz and Bauser [11], Greene [12], Brice [13] and Hsieh [14]. There was also a special edition of the Journal of Crystal Growth devoted to liquid-phase epitaxy [15]. These cover various aspects of LPE technology up to 1980. Although the volume of research on HI-V LPE has dropped there is still much work on improving the deficiencies of the process and this Datareview aims to highlight recent work and to assess the current state-of-the-art of LPE grown GaAs materials.. In Section B3 various aspects of the growth process itself are reviewed. In Section C, typical materials properties are described, and Section D briefly describes current device interest in LPE GaAs and (GaAl)As. B

GROWTH PROCESS

Bl

Introduction

There have been many different practical manifestations of LPE growth involving different ways of bringing together the solution and substrate and then separating them at the end of the growth period, such as the dipping [16], tipping [17], rotary [18], spinning substrate [19] and sliding boat [20] techniques. However, it is the latter technique which has become the most popular since it has several advantages as follows: (i) (ii) (iii)

it is economical in the use of solutions (typically 5-10 g) it gives good wipe-off of solutions at the end of growth it is well suited to the growth of complex multilayer structures.

The growth process normally takes place in a high purity silica tube, and the ambient gas is normally palladium-diffused hydrogen which contains very low concentrations of oxygen and water vapour (< 0.1 ppm). The sliding movements are usually performed with a silica pull-rod through a leak-tight sliding seal. For the growth of GaAs and (GaAl)As alloys, growth temperatures are typically around 8000C although more recent work has explored lower temperatures in order to achieve low growth rates for superlattice growth [21]. Cooling rates are typically in the range 0.1-1 K/min. Another technique which has proved useful is the semi-infinite solution dipping technique [22] where very large (> 2 kg) solutions are re-used many times with the solute concentration being topped up from time to time. This has been used primarily for the growth of solar cell structures where large-area deposition, high throughput and low cost are required. The following sections examine various aspects of the growth process in more detail. B2

Solvents

The principal requirements for a suitable LPE solvent are: (i) reasonable solubility of the solute

at moderate temperature (< 10000C), the useful range generally being 0.01 to 0.1 grammes solute per gramme of solvent, or atom fractions of As, X^, in the range 0.005 to 0.05 approximately; (ii) high surface tension to enable good wipe-off of the solution at the end of growth; (iii) low volatility at growth temperature; (iv) low melting point; (v) low incorporation of solvent species into the LPE layer (if a non-matrix solvent is used); (vi) availability in high purity form such as 99.9999% (6 9s) or 99.99999% (7 9s). TABLE 1 shows the properties of several potential solvents for GaAs LPE [23]. TABLE 1. Solvent properties for GaAs LPE. Solvent species

Melting point ( 0 C)

Surface tension at 800 0 C (dyn/cm)

Vapour pressure at800°C (atm)

Solubility at800°C (g per g solvent)

Ga

29.8

640

io- 8

3 x 10"3

Bi

271

343

io- 4

3 x 10 3

about IO"5

6 9s

In

156

514(600 0 C)

io- 6

about3 x 10 2

about2 x IO 2

7 9s

Sn

231

500

about 1 x 10 2

about 10"4

6 9s

Zn

420

about 750 0 C (420°C)

about2 x 10 2

about IO 2

6 9s

0.2

Segregation coefficient

Purity available

7 9s

Although all have reasonable solubilities, Zn and In are incorporated into the layers at a high level. The high vapour pressure of Zn is also a disadvantage. By far the most common solvent for GaAs LPE is Ga as it has near ideal properties. However much recent work has been done on the use of Bi and Bi + Ga mixed solvent for GaAs LPE [24-30]. Several advantages have been claimed for the use of Bi, namely: (i)

reduction in background carrier concentration from typical values of n - 1 x IO16 cm'3 for Ga solution growth to n ~ Ix IO14 cm"3 [24,25]

(ii)

increase in carrier mobility, |i, from typical room temperature values of 4000 with Ga solvent to 5000 cm2 /Vs using pure Bi solvent with 77 K values increasing from 5 x io 4 to 1.5 x 10 5 cm 2 /Vs [24]

(iii)

improved wipe-off of growth solution [27]

(iv)

reduced excess edge growth [29] and

(v)

increased growth rate by a factor of about 3 [27].

The purification effect involved in growth from Bi solutions is responsible for the improved background electrical properties. It appears that there is a fundamental change in the incorporation of common unintentional impurities. In layers grown from Ga solutions the main background impurities were found to be S donors and C and Si acceptors, the S and Si coming from the Ga starting material and the C from the graphite boat. These are all substitutional on the group V site as would be expected from the Ga-rich growth conditions. By contrast, the residual

impurities in Bi-grown GaAs layers are Sn and Si donors and Zn acceptors, all of which are substitutional on the group III site probably due to the higher gallium vacancy concentration in the Bi solution grown layers [26-30]. Finally, it has been found that the EL2 deep donor electron trap level commonly seen in VPE and LEC grown GaAs but never in LPE GaAs grown from Ga solutions, and thought to be due to the As anti-site defect AsGa, has been found in LPE GaAs grown from Bi solutions as might be expected from the change in the native defect concentration. By contrast, concentrations of the hole traps normally seen in Ga solution grown material and thought to be due to Ga anti-site defects, Ga^, are very much reduced in layers grown from Bi solutions [28]. B3

Boats

The predominant material for containment of LPE solutions, particularly in sliding boat systems, is graphite. This is because it has many advantages: (i) (ii) (iii) (iv) (v)

it is available in high purity form it can be readily machined into complex shapes to high tolerances it is inert to Ga solutions it is self-lubricating and it is relatively cheap.

Early problems with porosity and structural weakness have been removed with the availability of modern denser grades of graphite with lower porosity and high flexural strength. Care is needed to ensure that the boat is cleaned after machining to remove surface contamination. This can be done either by high temperature (> 10000C) baking in Cl2 gas or acid etching followed by extended soaking in boiling water, and finally an 8000C bake in hydrogen to remove residual water. Although silica [31], sapphire and boron nitride have also been used, the difficulties with machining precise shapes and the higher cost have prevented widespread use. As far as boat designs are concerned there have been many innovations, mostly around the basic sliding boat system aimed at improving wipe-off of solution [32,33], increasing the throughput of wafers per run [34,35] or growing thin layer structures [21,36,37]. Examples include: (i)

the wipingless boat of Horikoshi [32] where the solutions are pumped sequentially by pistons into a compartment over the substrate thus removing the need for slider movement

(ii)

the centrifugal design of Bauser [36] where a brief contact of solution and substrate is achieved by rapid spinning of the substrate

(iii)

the use of a continuous laminar flow of solution over a cooled substrate to achieve rapid growth rates at reduced Tg [34]

(iv)

the sliding boat of Heinen [35] whereby growth on up to 16 substrates can be performed simultaneously

(v)

the spiral boat of Kuznetsov [37] where two solutions are alternately brought into contact with the substrate by a rotating spiral tube arrangement.

B4

Furnaces

The main requirements for an LPE furnace are: (i)

a long zone of constant temperature within which the LPE boat is positioned during the growth

(ii)

good temperature stability with fluctuations of less than 0.20C

(iii)

good reproducibility of temperature (± 0.20C) from run to run.

Traditionally this has been achieved with three zone resistance wound furnaces with high thermal mass. Concentric heat-pipe liners have been used to help to provide a constant temperature zone and enhance thermal mass. However there has been an increasing interest in low thermal mass furnaces using gold-plated outer reflector tubes [38]. There are two advantages of such furnaces: (i) the outer gold coating can be made semi-transparent allowing the inside of the work tube to be viewed and (ii) rapid responses to changes in the temperature signal can be achieved. The latter is important where fast changes of temperature are required or a closed loop computer control system is being used where minimum delay is needed between the instigation of a computer-generated error correction signal and the subsequent change in temperature. B5

Computer Control

With the ready availability of cheap microcomputers and the necessary interface systems to allow them to control and monitor laboratory processes, there has been a trend towards computer controlled LPE systems. Compared to MBE or MOVPE systems the number and complexity of functions in an LPE system are low. The principal functions to be controlled or monitored are: (i) (ii) (iii) (iv) (v)

LPE boat temperature times of various stages of growth process water and oxygen concentrations in ambient gas movement of solution or substrate assembly in boat (distances, speed and acceleration) data logging.

Many common desk-top computers can be used with a suitable interface bus system such as IEEE 488 to communicate with the various components of the system such as A/D and D/A converters, stepper motor drive, relay board, printer and plotter and disk data storage. A very thorough discussion of the design considerations and the advantage of automated LPE has been given by Small et al [39]. B6

Substrates

For LPE growth of GaAs and (GaAl)As alloys, the substrate is in most cases GaAs sliced from ingots grown either by the horizontal Bridgman or Czochralski techniques. The former are sliced in a D-shape while the latter are cut in a circular form. Growth can be carried out on <100>, <111>A or <111>B crystallographic faces although <100> substrates are favoured for laser devices since they can be easily cleaved into square dice. To obtain grown surfaces free of the common terrace or 'sea shore' pattern commonly seen on LPE layers, it is necessary to orientate

and cut substrates to within 0.1 °, if possible, of the low index plane [40,41]. This requires careful Laue X-ray or laser reflection techniques. Final preparation before epitaxy normally involves a degreasing stage followed by a free chemical etch to remove residual lapping damage (see TABLE 2). The amount removed is normally only a few microns to avoid disturbing the surface flatness of the wafer. A thorough treatment of this subject has been given in a review by Kern [42]. TABLE 2. Chemical etching solutions for GaAs. Etch composition

Temperature

Etch rate

5 H2SO4-1H2O2 - 1 H2O (cone) (lOOvols)

about 40 0 C

1-2 microns/min

1NH4OH-19 H2O2 (lOOvols)

room

about 1 micron/min

(0.5-5%)Br2-CH3OH

room

1-10 microns/min (proportional to [Br2])

There has been recent interest in growing GaAs by LPE onto commercially available GaAs on Si [43-47], or growing GaAs directly onto Si by LPE [48] or onto Ge-buffered Si [49]. There are several reasons for this interest: (i) to use GaAs/Si as a cheap, large area substrate for example for LPE grown solar cells [49], (ii) to integrate optoelectronic devices such as GaAs/(GaAl)As lasers with integrated circuits in the Si substrate and (iii) to reduce the dislocation density in the original MBE grown GaAs layer by the growth of a thick LPE GaAs layer (typically from 5 x 108 cm"2 to (2 - 5) x 106 cm"2 in the LPE layer, [39]. The main difficulties to be overcome are: (i)

the dissolution of the Si substrate into the Ga solvent, which can be overcome either by using other solvents in which Si is relatively insoluble, such as Pb or Sn [46], or masking most of the GaAs-on-Si structure with SiO2 and opening up windows for the selectivearea LPE growth of GaAs [44-46]

(ii)

obtaining good nucleation of the LPE GaAs layer on the MBE grown GaAs on Si layer, probably due to the stress in the film caused by the differential contraction of the GaAs and Si on cool-down after the MBE growth. This can be overcome by using the transientmode technique of LPE where very high initial supercoolings are used to obtain a high density of nuclei.

Device results on GaAs/(GaAl)As double heterostructure high-emitting diodes [50] and solar cells [49] have been reported. Another area of LPE research involving substrates has been the work on using structured substrates. This may involve substrates with etched channels or mesastripes, or the use of windows opened in a protective film such as SiO2 or Si3N4 covering most of the substrate surface. Because the growth rate of LPE GaAs is very sensitive to the crystallographic orientation of the substrate surface, it is possible by careful choice of the solution supercooling (ATS) and the orientation on the substrate surface of the channel, mesa or window, to produce complex structures. There are three basic processes at work dependent on the type of surface structure. Firstly there is a tendency for a structured surface to become more planar during LPE growth to minimise the solid-liquid interfacial energy. Secondly, there is the orientation dependence of

growth rate mentioned above, and thirdly the presence of masked areas, where no epitaxial growth occurs adjacent to the unmasked growth areas. These can alter the diffusion flux of solute atoms and alter the growth mechanism from diffusion limited to surface kinetics limited. Examples of the applications of these techniques are: (i)

the growth of buried heterostructure lasers where two adjacent channels are etched in a standard LPE grown double heterostructure and then infilled with blocking or guiding layers in a second stage of LPE growth [51,52]

(ii)

the production of integrated optical waveguides by the selective LPE growth of GaAs on <100> GaAs substrates, where the waveguides are formed from <111> and <100> crystal facets which are the slow growing directions [53]

(iii)

lateral overgrowth of GaAs over tungsten stripes to produce higher speed GaAs-W-GaAs metal-base transistors by choosing the window and mask dimensions such as to enhance the solute diffusion parallel to the substrate surface [54].

B7

Growth Temperature

The choice of Tg for LPE is often a compromise between the requirement for a slow growth rate (lower Tg) and low background impurity concentrations (higher T8). Early work used Tg in the range 800 to 900 0 C for GaAs and (GaAl)As LPE growth with typical growth rates of 0.5 to 2 microns per 0 C cooling. Although Andre et al [55] first showed that (GaAl)As could be grown with good uniformity and control of Al content at Tg down to 700 0 C, it has been the recent interest in these GaAs/(GaAl)As structures with minimal interdiffusion for quantum well devices which has led to more research on low temperature LPE, < 450 0 C [56-62]. When using low Tg, it is important to do a pre-bake of the solution under hydrogen at a higher temperature (normally about 8000C) [60,61,63,64] and to ensure that levels of O2 and H2O in the growth system are very low (about 0.1 ppm). The LPE growth rate depends on the initial supercooling, the cooling rate and the Al concentration in the solution. Typical growth rates are as follows. TABLE 3. Temperature (° C)

Growth rate (^mZK)

800

1

GaAs

800

0.3

Ga065Al035As

630

0.2

GaAs

630

0.15

Ga065Al035As

450

0.02

GaAs

B8

Liquid Phase Electro-Epitaxy (LPEE)

The technique of electric-current induced liquid-phase epitaxy has been studied for many years and has been applied by many workers to the growth of GaAs [65-68] and (GaAl)As [69-72]. The apparatus typically consists of a modified graphite sliding boat system which allows the passage of an electric current across the substrate/solution interface, with a suitable insulator such as boron nitride between the sliding parts [74]. Although early work suggested that Peltier cooling at the liquid/solid interface was the major contribution to the growth rate [65,69,72], it was later shown that electromigration of solute species was the dominant mechanism [67,68]. The principal advantages of LPEE are the improved depth uniformity of both dopant concentration [71,72] and alloy composition [72,73] under constant current conditions. On the other hand, the segregation of the dopant can be changed by varying the current density [65,66,75]. Indeed, in the case of Si doping of GaAs or (GaAl)As layers at 900 0 C either donor (n-type) or acceptor (p-type) behaviour can be obtained depending on the current density [75,76]. Another feature of LPEE is the ability to grow very thick layers at constant temperature e.g. (GaAl)As up to 600 microns thick [73,76], and GaAs up to 4 mm thick [80], the latter having very low dislocation density and high electron mobility compared to normal bulk-grown GaAs. A very comprehensive review of the theory and application of the LPEE technique to the growth of III-V compounds and alloys has been written by Bryskiewicz [78]. B9

Doping and Gettering

One of the attractions of LPE is the wide choice of dopant species available for n- and p-type doping. The doping is normally achieved by direct addition of the elemental species to the growth solution. The main considerations in choosing a dopant species are: (i) the solubility of the dopant in the Ga-solution, (ii) the segregation coefficient k, the choice depending on the desired concentration of the dopant, (iii) the volatility of the dopant species at the growth temperature and (iv) the electrical activity of the dopant species in the GaAs lattice. Mainly the group IIA and HB elements (Mg, Zn, Cd) are used as acceptor species and the group VBB elements (Se, Te) as donors. The group IVB elements can potentially act either as donors on the group III site or as acceptors on the group V site. It has been found that while Sn and Ge are well behaved donor and acceptor species, respectively, the site distribution of Si changes such that a change-over from predominantly p-type occurs at 8900C [79]. TABLE 4 gives a list of segregation coefficients of several impurities in LPE GaAs, and in most cases agreement between workers is good [80-92]. Any discrepancies are due mainly to the use of electrical measurements to assess the dopant concentration rather than direct chemical analysis. Most dopant species show a linear relationship between log (carrier concentration) and log (dopant concentration in liquid) up to a carrier concentration of mid 1018 - 1019 atoms cm"3 when a levelling off occurs due to compensation by complex centres.

TABLE 4. Segregation coefficients of dopants in GaAs and (GaAl)As LPE. Dopants

T g (°C)

Orientation

k

Ref

Zn

800

<111>B

1.2 XlO 2

[80]

800

<111>A

2.5xlO- 2

[80]

800

<100>

1.8 XlO 2

[80]

880

<111>B

3 xlO"2

[81]

900

<100>, <111>B

4.5 x 1O"3

[82]

7 x 10-3

[83]

Ge

900

Sn

Se

Te

900

<111>B

4.8 xlO"3

[81]

900

<100>

7.5 XlO"3

[81]

700

<100>

1.1 x 10-4

[84]

700

<100>

7.4 x 10"5

[85]

700

<100>

1.8 xlO"4

[86]

850

2.8 x 10-4

[87]

860

1.2 xlO' 4

[88]

690

<100>

-16

[89]

800

<100>

-4

[89]

850

<100>

4.2

[90]

1.7

[87]

0.6

[91]

0.79

[97]

770 800

<111>B

850

O

1000

<111>B

0.37

[92]

700

<100>

4.7 x IO4

[85]

There have been few recent studies of doping in LPE GaAs or (GaAl)As, since the early work was fairly definitive. However there has been work on the addition of reactive species to the liquid phase to act as getters for residual impurities [93-98], in particular oxygen. By using in situ electrochemical monitoring of the oxygen concentration in the Ga solution, Chang et al [95] have shown that Ti or Zr rapidly deoxygenates the solution by between one and three orders of magnitude, although they concluded that the changes in background carrier concentration were due to the secondary effect of [O] on the [Si] in the solution. In some cases, rare-earth elements (Er, Yb) acting as getters have been of interest because of their potential as luminescent centres [96-98]. However, generally these very reactive species form oxides or hydrides in the solution which are then incorporated into the layers in an optically inactive form. There have been several reports on doping studies of Ga1^AlxAs alloys with Te, Sn, Si and Ge [99-105]. For donor impurities such as Te, the carrier concentration for a given dopant

concentration in solution is very similar to that found for GaAs as x increases in the range 0 to 0.2 [100,101]. As the direct-indirect gap crossover region is approached at x « 0.37, the carrier concentration falls by an order of magnitude and then levels off between x « 0.40 and x M . O [101]. By contrast, for acceptors, there is a steady reduction of acceptor concentration for a given solution dopant concentration as x increases from 0 to 0.55 [100,101]. C

LAYER PROPERTIES

Cl

Introduction

A very wide range of techniques are used in the assessment of LPE layers of GaAs and (GaAl)As. Some are routine and provide rapid feedback of results (Hall effect capacitance/voltage profiling, photoluminescence and EDAX) whereas others are more specialised and are used as part of ongoing materials research programmes. TABLE 5 lists these various techniques and the information gained. TABLE 5. Assessment techniques for LPE GaAs and Ga^xAlxAs.

Routine

Specialised

C2

Technique

Information

Hall effect (77 K - 300 K)

carrier concentration, resistivity, mobility

Electron probe micro-analysis

depth profiles or surface measurement of x

C-V profiling

carrier concentration depth profiles

Photoluminescence (77 K)

uniformity of x uniformity of luminescence efficiency minority carrier lifetime

X-ray diffraction

lattice parameter (x)

X-ray topography

crystal perfection and dislocations

Optical microscopy

layer thickness surface morphology

Etching

etch pit (dislocation) density

Photoluminescence (4 K)

impurity identification quantum well widths

Transmission electron microscopy

dislocation density layer thickness (quantum wells) interface abruptness

Secondary-ion mass spectrometry

matrix element and impuritiy depth profiles

Deep level transient spectroscopy

concentrations and energy levels of electron and hole traps

Rutherford backscattering

crystal perfection alloy composition x

Electrical Properties

When GaAs layers are grown from Ga solutions saturated with polycrystalline GaAs (n ~ 1016

cm'3) at temperatures around 8000C, the 77 K carrier concentration and mobility obtained would typically be n77K- 1016 cm"3, \i77K - 104 cm2 /Vs. The variation of background carrier density with x in Ga1^AlxAs layers has been studied by Wu and Su [106]. By prebaking of the Ga + GaAs solution at typically 700-8000C under H2 for periods up to 24 hr, these figures can be readily improved to n77K- 1014 cm'3, \x77K - 105 cm2 /Vs. Other factors to be considered are the purity of starting Ga and GaAs and the purity of the ambient gas. Best published figures are as follows: (i)

Tg = 850 0 C, silica boats [107]; n77K = 1 x io 12 cm'3 n77K = 2.5 x 105 cm2/ Vs

(ii)

Tg = 700 0 C, high purity and pre-baked graphite boats, [108]; n77K = 2 x io 14 cm'3 JI77K = 1.7 x 105 cm2/ Vs

(iii)

Tg = 8000C with controlled As vapour pressure n77K = 4 x io 12 cm'3 \i77K = 1.3 x io 5 cm2/ Vs [109] H77K = 4.5 x io 12 cm"3 n77K = 2.8 x 105 cm2/ Vs [110]

It is believed that while these results are partly due to removal of impurities from the Ga solution, there is also some compensation of Si donors and acceptors due to small changes in the site occupancy of Si atoms in the lattice. When gettering species are added, there is a tendency for the background carriers to become ptype, typically p77K~ 2 x io 14 /cm3, \i77K ~ 5 x IO3 cm2 /Vs [95], although it is not clear whether this is due to a reduction of O or Si in the layers or to the incorporation of getter atoms in the layers. C3

Luminescent and Minority Carrier Properties

Photoluminescence studies of LPE grown high purity GaAs have shown well-resolved excitonic structure at 4 K [114] and such studies can be used to identify the residual impurities in the layers [102]. Due to the well controlled stoichiometry of LPE-grown layers and the low segregation coefficients of the transition metal species which can introduce electron and hole traps into GaAs, the concentrations of deep levels are reproducibly low in LPE material. Electron traps are normally totally absent and only three dominant hole traps are generally found labelled A (Ey + 0.40 eV), B (Ev + 0.71 eV) and HB4 (Ev + 0.44 eV). The first two are believed to be related to the Ga^ antisite defect and HB4 to be possibly related to Cu [113]. The concentrations of these defects can be reduced to ~ IO12 cm"3 by decreasing Tg from 8500C to 740 0 C and using higher growth rates [114]. Positron annihilation studies on LPE-grown GaAs have shown that there are As-vacancies present at concentrations which fall from [VM]~ 2 x io 16 cm"3 at Tg = 900 0 C to -1015cm"3 at Tg = 700 0 C [115]. The internal radiative recombination efficiency in LPE grown GaAs material is generally high: hence the interest in LPE GaAs/(GaAl)As for light emitting devices. The interface recombination velocity at LPE grown GaAs/(GaAl)As interfaces is low (S < 105 cm2 /s) [116] which is important for solar cells and photocathode devices.

C4

Layer Thickness

The layer thickness of LPE grown GaAs layers as a function of growth time has been studied for various techniques of LPE (supercooling, step cooling, equilibrium cooling). The supercooling technique using an initial supercooling together with a cooling ramp has been found to give the most reproducible results in the thickness range 0.1 (im to > 20 ^m [8]. The variation of thickness is normally ± 5%, with run to run reproducibility of ±10% for both GaAs and (GaAl)As [8,117,118]. In a careful study of Ga,.xAlxAs growth (x = 0.15, x = 0.45) in the Tg range 600 to 800 0 C, Todoroki [118] has shown that for x = 0.45 and layer thickness around 2 \im the thickness uniformity improves from ± 12% at Tg = 8000C to ± 1.7% at Tg = 600 0 C. For very thin layers (< 0.1 nm) where the initial supercooling or the growth times are small, and the usual diffusion limited growth regime is not established, the thickness uniformity has been found to be independent of Tg in the range 600 to 8000C for Ga^xAlxAs (x = 0.15) growth (< ± 10%) [121]. Several workers have reported the LPE growth of layers in the thickness range 1 to 20 nm which is the region of interest for quantum well devices [21,56,58,118-121]. Generally low growth temperatures, Tg approximately 6000C, are used to achieve growth rates around 50 200 nm/min [21], as well as short growth times (< 1 sec) using automated solution movements. C5

Composition

The control of the alloy composition x in Ga1^AlxAs alloys is achieved by means of the atom fraction of aluminium X(Al) in the growth solution and the relationship between x and X(Al) at various growth temperatures is well established [122]. Run-to-run reproducibility of x is normally very good, typically ±0.01 [18,105,118]. Although published data on the uniformity of x across grown wafers is not available, it is probably ±0.005 by analogy with other LPE grown alloys.

C6

Interface Abruptness

An important consideration for many devices is the abruptness of a heterojunction or a doped junction. There are two separate effects in LPE which can lead to non-abrupt interfaces: (i) interdiffusion or impurity diffusion which can be reduced by using lower Tg and short growth times and (ii) etch back and regrowth which can be controlled by using lower Tg and increased supersaturation in the solution. The wide choice of dopant species available in LPE enables slow diffusing species such as Mg (acceptor) or Sn (donor) to be used when necessary. In work on quantum well growth, transition widths between GaAs and (GaAl)As layers of as little as 1 nm were reported [121]. C7

Structural Properties

The growth of homoepitaxial LPE GaAs layers does not normally introduce new dislocations. Threading dislocations from the substrate can be bent over at the interface or can annihilate each other if they have similar Burgers vector. Typical dislocation densities are < 104 cm'2 [122]. If, however, there is a marked difference in doping and hence lattice constant, between the GaAs layer and substrate a higher dislocation density can be found in the layer than in the substrate [123]. The growth Of(GaAl)As onto GaAs is basically a lattice matched process at a growth temperature of- 9000C. On cooling to room temperature after growth, strain is introduced due

to differential contraction. For thin layers Of(GaAl)As this can be taken up by elastic distortion of the layers [124], but for thicker layers or when slower cool-downs are used, dislocations can form by surface generation and then glide down to the (GaAl)As/GaAs interface [123]. The defect densities in LPE layers have also been found to be influenced by the vapour pressure of As during growth [125], while heavy iso-electronic doping with In in GaAs LPE has been reported to reduce dislocation density and improve luminescence efficiency [126]. D

DEVICES

Although the main interest in III-V epitaxy research is in MBE and MOVPE, particularly for complex thin-layer structures such as strained layer superlattice and multi-quantum well devices, there is still an important role for LPE in many areas of optoelectronics. Interest has been concentrated on devices where high luminescent efficiency and good minority carrier lifetimes are important. These devices include solar cells [127-131], double-heterostructure lasers [132-137] and photocathodes [138]. Although many papers have been published on the LPE growth of quantum-well devices, several mention problems with uniformity of the layer thickness and it is unlikely that there will be a role for LPE in these more demanding devices. E

UPDATE FOR THIS EDITION

Since the last edition, there have been few major breakthroughs in LPE research. Much recent work has concentrated on consolidation of earlier work on improving the purity of LPE material and on the control of native defects using solvents other than Ga. The design of LPE boats has been improved through the use of higher strength and higher density graphite and multi-wafer fixtures have been produced to increase throughput. A recent publication provides an overview of LPE growth of GaAs and other III-V compounds [139]. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] II1] [12] [13] [14] [15] [16]

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16.3 Halogen transport VPE growth of GaAs and related Previous Page compounds K. Somogyi September 1995

A

INTRODUCTION

This method is based on the early results of Ge iodide transport experiments. Deposition of III-V semiconductors (GaAs, GaP, InAs, InP) from the vapour phase by a chemical reaction was first reported by Antell and Effer using iodides and chlorides [I]. The main idea is that (a)

halogens form volatile compounds with the group III metals (and also with Ge and Si, of course) at high temperatures (the 'source' region),

(b)

these reactants can be transferred in the gas phase to the substrate at the colder end of the tube (the 'deposition' zone),

(c)

these molecules decompose at these lower temperatures and as a result

(d)

the elements of the semiconductor become adsorbed on the surface of the seed/substrate forming an epitaxial layer.

(e)

The gaseous halogen either returns to the higher temperature region (the cyclic process in the 'closed tube' system) or is exhausted from the system (continuous flow in the 'open tube' system).

We discuss the chlorine based, hot wall, open tube VPE systems with H2 carrier gas. 'Halogen transport VPE' developed in two directions: 'chloride' VPE and 'hydride' VPE. The first one uses AsCl3 (PCl3) for both chlorine and arsenic (phosphorus) sources; the second one uses HCl for chlorine and AsH3 (PH3) for arsenic (and phosphorus) sources. These are by far the most widely spread and industrialised methods. The main advantages of the classical VPE process are the high growth rate (0.3 - 0.4 (im/min and more) and the high purity of undoped epitaxial layers (< 1014 cm"3). The main drawback of these methods is the existence of a relatively wide transient layer between the substrate and the epitaxial layer. The use of direct chloride sources (GaCl3, AlCl3) has been less important till now. Most recently the use of organo-metallic chlorides (DEGaCl, DEAlCl) and HCl gas in MOVPE reactors has opened new possibilities for atomic layer and selective area epitaxy, especially that of GaAlAs compounds. B

MAIN HALOGEN TRANSPORT VPE GROWTH TECHNIQUES FOR GaAs

The main chemical and kinetic processes are similar for all group III metals (Ga, In, Al) and group V sources (As and P), so only processes related to GaAs are described.

Bl

Metal Source Chloride Transport VPE Methods

'Chloride' and 'hydride' methods are distinguished by their different arsenic source material: trichloride (AsCl3, PCl3) or hydride (AsH3, PH3). The higher purity of the chlorides compared to the hydrides at the beginning led to a differentiation in use: the chloride method is used mainly in microwave applications and the hydride method is useful in optoelectronics. These methods are reviewed in [2,3]. Bl.l

'Chloride' VPE, AsCl3ZGaZR2 system

The method using AsCl3 and GaAs sources was originally proposed by Effer in 1965 [4]. FIGURE 1 shows schematically the arrangement of such a reactor. AsCl3 as a liquid is introduced into the reactor using saturated H2 carrier gas directly at the surface of the source material. In this region and at these temperatures the AsCl3 first decomposes (EQN (I)) and then the HCl reacts with the GaAs (EQN (2)). Volatile GaCl and arsenic arrive at the substrate in the deposition zone and at a lower temperature a process opposite to EQN (2) takes place: EQN (3). AsCl3 + 3Z2 H2

-

V4 As4 + 3HCl

(1)

2HCl+ 2GaAs

-

2GaCl + ViAs4 + H2

(2)

2GaCl + 1Z2As4 + H2

-

2GaAs + 2HCl

(3)

FIGURE 1 shows typical temperatures. AsCl 3 + 3 / 2 H 2 => V 4 As 4 +3HCl

(B 1.1.1) 1

2HCl+2GaAs => 2GaCl+ / 2 As 4 +H 2

(B 1.1.2)

2GaCl+V 2 As 4 +H 2 => 2GaAs+2HCl

QUARTZTUBE

SOURCE ZONE TEMPERATURE

FURNACE

MIXING ZONE

(B 1.1.3)

GaAs SUBSTRATE

DEPOSITION ZONE

DISTANCE

FIGURE 1. Scheme of the metal source chloride reactor and of the temperature distribution along the reactor.

Pure Ga metal became widely used as a source instead of GaAs and has the advantage of higher purity. Arsenic is dissolved in the molten Ga and a GaAs crust/core is observed on the Ga melt surface [5,6]. The stability of the growth processes depends on the stability of this core. A presaturation of the Ga melt stabilizes the source [5,6]. The second AsCl3 line was proposed by Nozaki et al [8] (AsCl3 II line in FIGURE 1) for in situ etching by HCl formed in the reactor according to EQN (1) ('Effer-Nozaki'system). Furthermore, the second AsCl3 line made possible the variation of the V/III ratio in a wider range, determined otherwise by source reactions. The increase of the mole fraction of the AsCl3 also leads to a lower background carrier concentration level [9,10]. Whereas epitaxial growth from the liquid phase (see Datareview 16.2) can be considered as a process under equilibrium conditions and the pyrolitic deposition methods, e.g. MOCVD (see Datareview 16.5) or molecular beam deposition (see Datareview 16.6) are typical non equilibrium processes, halogen transport VPE is an intermediate technique. EQNs (2) and (3) together describe a dynamic equilibrium. Thermodynamic aspects of the Ga/As/Cl/H2 (and Ga/As/I/H2) system have been studied in [1116], kinetic aspects in [17-19] and chemical ones in [19-22]. The experimental growth aspects have been studied in detail, as well. These include the influence of the technological parameters (temperatures, temperature gradient, partial vapour pressures OfAsCl3, GaCl, etc., substrate surface orientation, flow rates) the surface morphology, the growth rate [23-26], and the homogeneity [17,27-31]. For the modified halide VPE technologies like the low pressure (LP) method see [32,33], for low temperature (LT) growth in the 'kinetic' region see [33-37] and for the 'flat temperature zone' arrangement, see [36,37]. The dependence of the background n-type carrier concentration on the technological parameters was established first in [9,10,27,38,39]. It was shown by indirect methods that Si plays a decisive role in the formation of the background electron concentration [38,43,44], though attempts at the direct identification of Si in layers did not lead to unambiguous results [40,41] and there were also contradicting results [42]. n-type doping can be realised by the introduction of group VI elements such as S, Se and Te [2,3,45-48]. Group IV elements do not show their amphoteric character in VPE GaAs: Ge, Si and Sn give rise to n-type conductivity material and are excellent donors [2,3,40,41,43,45,49,50]. C is a deep acceptor, which is not used widely in VPE practice [49]. Group II elements like Zn, Cd and Mg can serve as acceptor dopants [2,3]. Oxygen, iron and chromium can be used for obtaining semi-insulating (SI) layers [2,3,49]. Undoped high-resistivity buffer layers have also been grown [29,30,51,52]. Mechanisms for the incorporation of impurities have been widely investigated [37,38,45,46,53-56], but effects of simultaneous doping with different impurities were rarely tested [49,57,58]. One of the accepted disadvantages of chemical halogen transport VPE is the difficulty in growing sharp interfacial layers. The formation, the behaviour, the structure and the properties of the transition layers/regions and the influence of the substrate material on the properties of the epitaxial layers have been investigated thoroughly [29,30,51,52,59-62].

High purity, high growth rates, flexible doping, good morphology and high yield has made possible the mass production of discrete and integrated microwave [29,30,48, etc.], high voltage and high speed rectifying [63] devices. New possibilities for micro-machining [58], for shaped and conformal growth of GaAs submicron layers [64] and in selective (patterned) growth have been demonstrated. Nowadays different p-n junction devices and laser diodes are fabricated from GaAs, InP and GaP /Si heteroepitaxial structures [65-67]. There are some new investigations concerning the thermodynamic description of the method [68], the growth processes [69,70], the gas phase composition, the impurity incorporation and substrate orientation effects [71-73], and different p-type doping [74]. B1.2

'Hydride' VPE, AsH3/Ga/HCl/H2 system

The method was originally proposed by Tietjen and Amick [75]. For the chlorine and arsenic supplies HCl and AsH3 (PH3) are used. FIGURE 2 shows the scheme of a horizontal hydride reactor. The temperature distribution along the reactor can be similar to that of the chloride method, but the supersaturation of the gas phase in the growth zone can be adjusted by the choice of the flow rates [76,77].

(4) (5) (6)

(Bl.2.1) (Bl.2.2) (Bl.2.3)

OUT

Doping

LET

QUARTZ TUBE

FURNACE

GaAs SUBSTRATE

FIGURE 2. Scheme of a metal source hydride horizontal epitaxial reactor.

Here the HCl + H2 gas is in contact with the Ga metal source (inlet HCl I in FIGURE 2). Usually AsH3 + H2 is introduced directly into the transfer or deposition region, bypassing the Ga source. For in situ etching a second HCl line can be supplied (HCl II inlet).

The growth process is as follows: AsH3 decomposes into volatile As, EQN (4). HCl reacts with Ga, according to EQN (5). GaCl is developed directly from Ga; no GaAs crust formation takes place. In the deposition zone the same processes occur as in the case of the chloride method, EQN (6). The only difference between the chemical processes of the chloride and hydride methods concerns the source zone. Because of the independent III and V sources, Ga can be replaced with In to obtain InP InAs, or both Ga and In sources can easily be used simultaneously. The preparation of GaInAs, GaInP and GaInAsP is much more convenient than in chloride VPE systems. Inhibition of the decomposition OfAsH3 (and PH3) has been experimentally confirmed [76]. In this case the reaction between GaCl and AsH3 is direct at the substrate surface [77]. GaCl + AsH3 - GaAs + HCl + H2

(7)

Deposition of GaP on Si was successful at low temperatures (350 - 45O0C) by mixing the reactants (GaCl and PH3) just above the surface of the substrate: 'vapour mixing epitaxy (VME)' [78]. Probably these effects play a role in the 'vapour levitation epitaxy (VLE)5 technique, as well [79]. Red LED fabrication was established using GaAsP/GaAs structures in the early 1970s. Industrial needs for GaInAsP for long wavelength optoelectronic devices (including lasers for optical communication) can be fulfilled with high capacity hydride reactors. InP, GaAsP, GaInAsP, GaInAs and more recently AlGaAs are key materials of optoelectronics and are grown on an industrial scale. Use of hydride VPE in obtaining III-V heterostructures on Si substrates is also continuing [69,73,78,80]. Atomic layer, conformal and patterned epitaxy have been developed in hydride systems, often as a low pressure variant [66,80]. B2

Metal-Trichloride Source VPE, GaCI3/AsH3/H2 System

Monochlorides of the group III metals (GaCl, InCl, AlCl) exist only at the high temperatures of the reactor (EQNs (2) and (3)). They can also be obtained from trichlorides (GaCl3, InCl3, AlCl3) by reducing them in an H2 atmosphere. Trichlorides are solid at room temperatures. At temperatures between 50 and 100 0C they have sufficient vapour pressures to be transferable into the reactor by saturated H2 or N2 fluxes. At temperatures above 500 0C chemical reduction occurs, GaCl3 + H 2 - GaCl + 2HCl

(8)

and we obtain the halide transport agent, which can react with arsine to form GaAs (EQN (6)) [81]. In LT cases GaCl3 can react either directly with the uncracked arsine [82, 83]: GaCl3 + AsH3 - GaAs + 3HCl

(9)

AlCl3 + V2H2 + V4As4 - AlAs + 3HCl

(10)

or with As4 (EQN (4)):

This reaction is reported to be more effective for AlCl3 [80]. The main conditions correspond to the usual hydride technique. Application OfGaCl3 is possible in cold wall systems, too. The main advantage of this technique is that the reactions do not need a high temperature source zone. The method is used mainly for ALE and for the growth of ternary compounds, especially AlGaAs and for patterned/selective growth of these compounds (see Section D). B3

Metalorganic-Chloride Source VPE, DEGaCI/AsH3/H2 (or DMGaCl) Systems

MOVPE has become the most popular epitaxial technology in the last decade. This has provoked the application of organometallic-chlorides of group III sources instead of metallic ones. Initially, the carbon contamination in epitaxial layers grown by MOVPE was a major problem. Zaouk and Constant [84] reported the reduction of C contamination in GaAs using DEGaCl. Later Prakash developed the growth technique for Al and In containing alloys [85]. Application of MO-chlorides in MOVPE systems instead of TEGa or TMGa also has another advantage: that a 'back etch' by the in situ generated HCl controls the deposition of the solid phase (see also Sections C and D). DMGaCl was used to obtain GaAlAs layers by Hasegawa et al [86]. The reactor was a 'usual' hot wall chloride system and AsCl3 was used in these AlAs and AlGaAs layer growth experiments. Mori et al published results on the application of DEGaCl in 1988 [87]. In this work a MOVPE system was used under conventional MOVPE conditions including an AsH3 source. MO-chlorides of group III metals decompose easily even at relatively low temperatures (425 45O0C). Ethyl-based precursors undergo gas phase-elimination, which in turn further reduces carbon incorporation into the deposited layer (compared to methyl based ones). The decomposition is assumed to be Ga(C2H5)2Cl - GaCl + 2C2H4 + H2

(11)

and again we have GaCl which reacts with AsH3 or As4 [88]. The method is widely used in ALE, selective area and patterned growth practice, including device fabrication (see Section D). C

HALOGEN TRANSPORT ASSISTED EPITAXIAL TECHNIQUES

Cl

Introduction

The forgotten in situ etching of facility VPE has now enjoyed a great revival in MOVPE practice. This is the most recent and very important development in the practice of halide transport selective epitaxy. The aim of selective epitaxy (SE) is to grow monocrystalline GaAs (or alloys) in the openings of a dielectric mask without growth occurring on the mask. Usually, deposition of the semiconductor occurs on the surface of the mask, due to strong reactions between the Al

precursor and dielectric masks. This seriously limits the width of the passivation film. The Al content is also influenced by this effect [89]. In order to control, or to avoid, deposition on the surface of the mask, etching has been involved. First DEGaCl (DEAlCl) was used, where HCl, as a by-product, takes on the etching role of the chlorine (see Section B3) [88,90,91]. Later, the direct use of HCl was reported [92,93]. The role of halides is usually auxiliary in these cases, being used for etch back on the surface and etch away from the surface of the substrate and the masking dielectric. Some applications also in MBE practice are published. C2

Chlorine Assisted MOVPE

Both bromine and fluorine have been utilised in some experiments, but chlorine remains dominant. Chlorine can be introduced into the MOVPE system in different compositions. The use of metal chlorides, or AsCl3 (PCl3), or chlorine gas is not important. MO-chlorides, CCl4 and direct HCl are the most important Cl sources. C2.1

Metalorganic-chloride assisted MOVPE

This is a method which is based on MOVPE equipment, but uses real halide transport and HCl etch (like a modified halide VPE system): see Section B3. C2.2

HCl assisted MOVPE

In this variant, HCl is directly introduced into the MOVPE growth chamber. Selective epitaxy is possible by this technique (see Sections C and D). Thus, lasers using AlGaAs confinement layers of different wavelengths (0.78 and 0.98 mm) and with complicated structures have been fabricated [94]. One can expect that this method will be expanded. Another use of HCl is demonstrated in the case of selective area epitaxy of In compounds. In a hot wall set-up, co-injecting the conventional TMIn and HCl results in the formation of InCl [95]. In this case chlorine realises both the classical transfer of the metal (InCl) and substrate etching HCl. A 0.13 mm/min growth rate was achieved in growing InP and planar p-i-n diodes have been fabricated. C2.3

CCl4 assisted MOVPE

Heavily doped p-type layers are of interest for n-p-n heterojunction bipolar transistor structures, tunnel diodes, nonalloyed contact layers, etc. Carbon is a known acceptor with the advantage of lower diffusivity compared to other p-type dopants, e.g. Zn [96]. CCl4 is a very convenient source of C doping. HCl formed in the deposition reaction, as a byproduct, controls the deposition rate in MOVPE (AsH3ZH2) systems. It has a special importance in the case of selective area epitaxy, since this HCl inhibits polycrystalline growth on the mask [97,98]. The same effect can also be achieved by using carbon tetrabromide (CBr4) [99]. This method is preferentially used in the growth of InP and other In compounds. Recently lattice strain was controlled in AlGaAs/GaAs structures by codoping with C and In [10O].

C2.4

Group V trichloride assisted MOVPE

(MOVPEH-ASCI 3 ,

(PCI3))

Though AsH3 serves as the arsenic source in MOVPE systems, AsCl3 was used in early MOVPE experiments for in situ etching, as in chloride VPE (see Section B1.1). Recently, growth of GaAs, InP, GaP, GaInAs and GaInP has been improved (including Si substrate heteroepitaxy, selective epitaxy, C doping, etc.) using AsCl3 or PCl3. [101-104]. In these cold wall systems AsCl3 (PCl3) decomposes just above the surface of the substrate and it can be used either in atmospheric pressure or low pressure systems, and also at low temperatures (see also EQN (7)). The thermodynamics of the processes are discussed in [101]. Expansion of this approach is expected. C3

Halogen Assisted MBE/MOMBE/CBE

Success of the chlorine assisted effects in MOVPE led to the application of halogens in vacuum and very low pressure techniques. Different halogen sources have been used: DEGaCl for self limiting growth of GaAs [105], HCl or pure chlorine gas for in situ etching of GaAs substrate surfaces [106] and CBr4 for C doping [107]. CCl4 was used for the deposition of GaAlAs: C and InGaAs:C layers and also for the regrowth of GaAs:C [108]. AsCl3 has been used for etching, and also as the arsenic source in GaAs layer growth [109-112]. The decomposition processes OfAsH3 and PH3 (see Section B 1.2) have also been investigated in a CBE system [113]. D

SPECIAL APPLICATIONS OF HALOGEN TRANSPORT

There are two particular cases of peculiar interest. D.I

Selective (Area) Epitaxy and Halogen Transport (Patterned (Re)Growth, etc.)

Chloride and hydride VPE methods show potential for the patterned growth of layers, e.g. conformal growth. Modern integrated circuits need layers with different alloys and MOVPE is well suited. However, in MOVPE nucleation on the dielectric mask is practically unavoidable. The strong reactivity of Al with oxygen leads to a high nucleation density on the mask degrading the patterned growth of such layers (see Section Cl). Halogen assisted MOVPE has imparted a new impulse to selective area epitaxy, in classical halogen transport VPE, and particularly in the hydride growth of InP [114]. TABLE 1 summarises recent results of selective area growth results. D2

Atomic Layer Epitaxy and Halogen Transport

The names 'molecular layer epitaxy' (MLE) and 'digital epitaxy' [139] have been introduced for this monolayer-cycle growth. The unique characteristic of ALE is due to chemical interaction saturation between the substrate surface and the reactant gas. In the case of an AB binary compound semiconductor the simplest method of ALE is to supply A atoms and B atoms alternately to the growth surface. Molecular sources are used instead of elements in III-V ALE. The GaCl (or InCl, AlCl) source can result from the pure metal reacting with HCl (see Section B 1.2), or GaCl3 with a carrier gas (see Section B2), or an organometallic chloride cracked at

elevated temperatures (above 400 0C) producing GaCl (see Section B3). For sources of group V elements hydrides (AsH35 PH3) are always used. For ALE growth, separate inlets are usually used for the group III monochlorides and for the group V source and the substrate moves between them. Either rotating [81,129] or sliding [130] substrate holders are used for the continuous repositioning of the substrate. Alternating source supplies and a sliding graphite or quartz covering plate above the substrate can also be imagined. TABLE 1. Summary of the halogen based/assisted selective area epitaxy growth results. Compound

Method

Section

References

Remarks

GaAs

chloride

BLl

[64,115]

conformal

hydride

B1.2

[114]

MO-chloride+MOVPE

C2.1

[91,116,117]

HCHMOVPE

C2.2

[118,119]

hydride

Bl.2

[65,80]

CC14+MOVPE

C2.3

[97, 98]

hydride

B1.2

[65,95,114,120-123]

PC13+(LP)MOVPE

C2.4

[102,124]

MO-chloride+MOVPE

C2.1

[88,91,125,126]

HC1+MOVPE

C2.2

[89,90,92-94, 103, 118, 127]

AsCl3+MOVPE

C2.4

[104]

CC14+MOVPE

C2.3

[97]

JnP GaInAs GaAlAs

GaInAsP III-V

1 (LT)hydride

1 Bl.2

conformal [80]

lasers

lasers

| [114,128]

!general

TABLE 2. Summary of the halogen based atomic layer epitaxy growth results. Compound GaAs

Method metal source hydride

Section BL2

References [87, 114, 130-139]

metal-trichloride source

B2

[81]

metalorganic-chloride source

B3

[87,138,140-142]

Cl assisted MOVPE

C2J

[87,116,138,140]

Cl assisted MBE

C2

[99]

InAs

metal source hydride

BL2

[128, 121,129]

InP

metal source hydride

BL2

[114, 128, 129]

GaP

metal source hydride

BL2

[128, 129]

GaInAs

metal source hydride

B1.2

[114,130]

GaInP

metal source hydride

BL2

[128,143]

GaAsP

metal source hydride

B 1.2

[144]

InAsP

metal source hydride

Bl.2

[144]

High uniformity of the layer thickness, side-wall growth and high quality selective growth can be realised by the ALE technique [129]. Se5 Si and Zn doping have been successfully applied and FET and quantum well based devices have been fabricated [129,131-133]. The chemical and

kinetic processes, the influence of growth conditions, growth mechanisms, substrate surface orientation effects, etc. have been described [81,87,99,129,134-137]. TABLE 2 summarises recent results of ALE of different III-V semiconductors. E

CONCLUSION

Though classical halogen transport based, vapour phase epitaxial methods have been overshadowed by newer technologies, they have some unique advantages. Practically all the III-V compounds can be grown by classical VPE. Some ten years ago chloride transport based VPE was called 'old fashioned' because the presence of a large transient layer was a major drawback. However, ALE growth experiments have shown that the most effective ALE methods are based on classical VPE processes and chlorine assisted MOVPE and MBE methods. The halogen-based processes manifest their greatest advantage in selective growth processes. The most powerful types of selective epitaxy include halide transport to improve growth quality. The halogen process plays a significant role in carbon-free growth from inorganic sources. There could be another possible advantage. Adverse hydrogen incorporation in MOVPE grown layers [145] suggests that halogen VPE based devices, which are probably free of hydrogen, may show better high temperature stability and lifetime than MOVPE grown devices. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] [II] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23]

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16.4 Growth by close space vapour transport J.C. Bourgoin and H. Samic June 1996

A

INTRODUCTION

The growth of binary and ternary III-V compound semiconductors is now well mastered. Vapour phase (VPE) as well as liquid phase epitaxy (LPE) allow the production of thin layers of good quality material (free of defects) and their doping can be adjusted at will. Other techniques, namely molecular beam epitaxy (MBE) and metalorganic chemical vapour deposition (MOCVD)3 allow one to make sharp interfaces necessary for heterostructure based devices. Overall, up to date materials and heterostructures for microelectronics can be obtained with all these techniques, with the noticeable exception of semi-insulating (SI) layers. The SI substrates on which GaAs epitaxial layers must be grown are cut from ingots obtained by the liquid encapsulated Czochralski (LEC) method and the vertical gradient freeze (VGF), or Bridgman, method. However, these bulk substrates present several drawbacks: the LEC material contains a large number of dislocations and the VGF as well as the LEC materials have electrical properties which are not homogeneous over the whole surface of the wafer and vary from wafer to wafer. There is thus a need for semiconductor layers that exhibit uniform SI properties for microelectronics, but also for other specific properties corresponding to many applications, that cannot be obtained by the above-mentioned epitaxial growth techniques. For instance, one needs: cheap layers to improve the uniformity of the electrical properties of substrates, for energy conversion; SI and doped thick layers for high power electronics, for optical windows and particle detection; layers containing a controllable number of defects to adjust the lifetime for photodetection, etc. This can only be achieved with a technique allowing variation of the growth rate in a very large range. This technique must also be inexpensive, fast and should not necessitate the use of toxic gases to have a serious chance to be developed in the forthcoming years. B

PRINCIPLE

Such a technique always existed. It is usually called close space vapour transport (CSVT), a vapour phase CVD modified technique, in which the transport of matter is 100 % efficient and the rate of gas flow is not the limiting factor [I]. It consists in the chemical decomposition of a source material placed in front of the substrate, at a close distance (few mm). The temperature gradient between the source and the substrate ensures that the products of the decomposition are transported to the substrate where the reverse reaction occurs, inducing the growth of an epilayer on the substrate. In other words, it is the concentration gradient of the reacting species that drives the mass transport and hence the growth. The decomposition and recomposition reactions can take place at or near atmospheric pressure, in a gas (hydrogen) to which the chemical agent is added. Reactions with iodine [2,3], chlorine [4,5] or simply water [6-11] have been used. In the early studies, the chemical agent was not identified but was residual water contained in the hydrogen gas. Indeed, the growth rate is non negligible only in the case where water is present and is an increasing function of the water partial pressure [6,8]. The reaction of water with GaAs

is most probably to produce volatile Ga2O. The details of the transport process are not yet correctly understood. It has been suggested [7] that the reactions at the source and substrate surfaces limit the growth. But the existence of a third reaction occurring in the gas phase had to be postulated to account for the dependence of the growth rate both on the source and substrate temperatures and on the water pressure [7]. However, it has also been proposed that it is the diffusion of the reacting species in the H2 atmosphere that is the limiting factor [12,13].

GROWTH RATE (|jm mirf1)

The growth rate is thermally activated [7] with the temperature of the substrate 0, i.e. increases exponentially with 0"1 until it reaches a maximum when the temperature difference between the source (T) and substrate becomes small, then drops to zero for T = 0. As to the rate dependence with the water pressure p, it apparently (see FIGURE 1) varies with p1/2 [6,9], but the physical reason for this dependence is not yet clear.

FIGURE 1. Variation of the growth rate versus water vapour pressure for two values of substrate 0, and source (T) temperature: 0 = 760 0 C, and T = 800 0 C according to reference [8] ( • ) , and for 0 = 800 0 C and T = 900 0 C (•), (unpublished data).

The rate of growth can be varied in a wide range, from mm per hour, as in conventional techniques, to 10 |nm per minute [5,14-16] or even higher (see unpublished data in FIGURES 1 and 2). High rates are easily obtained simply because the transport is 100 % efficient. Thus layers several hundreds of microns thick can be grown in a matter of minutes, with good structural properties.

GROWTH RATE (pm mirf1) FIGURE 2. Variation of the growth rate GR versus the substrate temperature 8, for three different source temperatures: 900 0 C (A) 5 850 0 C ( • ) , and for 8 = 8000C (•).

C

RESULTS

Growth by the CSVT technique was studied as early as the classical VPE and LPE techniques that soon became the conventional techniques. It is the rapid success of the VPE technique that probably stopped the development of CSVT. Only limited attempts to grow different materials have been made: the binary and ternary III-V compounds GaAs, GaAsP3 InP and GaP as well as H-VI compounds CdTe3 ZnSe and ZnTe (for a review see [17]). Very few hetero-epitaxies have been attempted: GaAs on Ge [3,18,19], and on Al or Mo [20]. The growth rate is adjusted through the substrate temperature 0, the source temperature T and the water pressure. FIGURE 2 provides, as illustration, data in the high growth rate range (larger than 1 \im per minute) obtained for a constant pressure of water (0.2 torr). The dependence of the growth rate with the pressure is given in FIGURE 1 for two specific cases in order to illustrate its effect quantitatively. Thin GaAs layers have been obtained whose structural [21] and electronic properties can be equivalent to that of epitaxial layers grown by the conventional VPE technique when the growth conditions are properly adjusted. The surface morphology varies with the growth conditions and also with the thickness of the layer [8,13]. The electron mobility [3,9,10-12] and the luminescence spectra [22-24] can be similar in VPE and CSVT layers, which is not surprising since the growth is achieved at similar temperatures. The residual impurities and defects [12,2527] have also been characterized. The doping can be controlled [10,11,15,22,28,29] by'the impurity concentration of the source and the growth conditions. However, the question of impurity transport during growth [13,22] has not been studied systematically except for S [29]. The concentration of antisites can be adjusted via the growth rate [15,16]. When the growth rate

is large enough, the presence of defects (usually labelled EL2) associated with the As antisite apparently renders the layers semi-insulating [15]. D

CONCLUSION

The CSVT technique has never been employed for the development of devices. The success of the classical VPE and LPE techniques is probably the reason that the CSVT technique has not emerged outside laboratories. Only demonstration devices have been made such as p-n junctions for photovoltaic applications [20,28], tunnel diodes [5], and field effect transistors [30]. However, this technique possesses unique advantages which will hopefully be recognized: from the economical point of view (it is fast and the equipment is inexpensive), for safety (it is free of toxic gases) and because it can produce layers that have no equivalent (large thickness, semiinsulating and containing a controllable number of defects). REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] II1] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30]

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16.5 MOVPE growth of GaAs M.R. Leys December 1995

A

INTRODUCTION

The aim of this Datareview is to provide insight into the technology of vapour phase epitaxy using metalorganic starting compounds. The emphasis will be on the growth of the Am - B v semiconductor (most often simply called HI-V material) gallium arsenide (GaAs). This technology is mostly termed MOVPE (metalorganic vapour phase epitaxy), also termed OMVPE. The 'E' emphasizes the application to cover single crystalline substrates with single crystalline layers where the lattice planes are in perfect registry i.e. continuous throughout the whole structure. The often encountered abbreviation MOCVD (or OMCVD) is, strictly speaking, not equivalent as chemical vapour deposition (CVD) may also indicate coverage of a substrate with polycrystalline or amorphous thin films. The growth of GaAs layers using the metalorganic compound trimethylgallium ((CH3)3Ga, TMGa) as group III source combined with the hydride arsine (AsH3) as group V source, both transported to a substrate by means of a carrier gas, was first demonstrated by H.M. Manasevit. The growth of other III-V compounds was also shown to be feasible: gallium phosphide (GaP), using TMGa and phosphine (PH3); GaAsP was grown by introducing both arsine and phosphine simultaneously with TMGa. Trimethylaluminium (A12(CH3)6, TMAl) and trimethylindium ((CH3)3In, TMIn) were subsequently introduced to obtain aluminium and indium containing layers. The pioneering work is well described in [I]. The epitaxial growth of semiconductor materials is of eminent importance for the fabrication of electronic devices. The III-V materials are the candidates of choice for light emitting as well as light detecting devices. Thus, stimulated emission (laser action) from a III-V structure was first demonstrated in 1970 using a GaAs/Al,GaAs heterojunction diode [2]. Initially, such multiple layer structures were grown from the melt by liquid phase epitaxy (LPE), a technology which is described elsewhere in this book (Datareview 16.2). The success of silicon epitaxial growth from the vapour phase (for a review on silicon epitaxy and processing, see [3]) in producing high quality, uniform layers on large substrate areas stimulated research towards vapour phase growth of III-V materials. Thus, chloride and hydride VPE were developed with some degree of success for the growth of GaAs and GaAsP layers, (see Datareview 16.3 in this book and [4,5]). Additionally, during the 1970s GaAs layers were successfully grown by directly evaporating the elements in an ultra-high-vacuum chamber towards a heated substrate [6]. This technology was termed molecular beam epitaxy (MBE) (Datareview 16.6 and Chapter 17) and demonstrated as early as 1976 the ability to grow layers of GaAs with monolayer thickness control [7]. Thus, over the years, the various epitaxial techniques have stimulated each other's development. This applies to all the aspects of growth: to the topic of nucleation and layer growth mechanisms, to the comparison of structural perfection and layer quality as well as to issues concerning mass production of devices using multiple-wafer-systems. A very clear case of mutual influencing is the development of chemical beam epitaxy (CBE) which uses the metalorganic and hydride starting

materials of MOVPE and the UHV growth chamber of MBE, (see [8-10] and Datareview 16.7 in this book). The applications and the understanding of epitaxial growth are continuously increasing. A number of books [11,12] and review papers on specific MOVPE related topics [13-16] have been published. Nevertheless, many topics are still under investigation and many materials systems have not yet reached the degree of control which is required for mass production of reliable devices. The technology of MOVPE has recently received a great stimulus with the successful growth of high brightness blue light emitting devices containing gallium nitride (GaN) layers using TMGa and NH3 [17,18]. B

FLOW DYNAMICS AND REACTOR DESIGN

Studies on the behaviour of the carrier gas flow through VPE reactors already had been carried out during the development of silicon VPE. For example in [19], investigations are described using small TiO2 particles as markers of the streamline pattern through a horizontal VPE reactor. The flow through a VPE reactor is determined by the forced convection, due to the pressure difference between inlet and exhaust in combination with the free convection which arises from temperature differences inside the system. The tendency of a hot fluid to rise is termed 'buoyancy'. Both silicon and MOVPE reactors are of the cold-wall type and have only one heated area, namely the susceptor on which the substrates are placed. A continuous flow of carrier gas is passed through the reactor and the components required for growth and doping are introduced in the carrier gas via valves on a mixing manifold. Mostly hydrogen purified by a palladium diffiisor is used as the carrier gas. Initial research on MOVPE was carried out in small, laboratory scale equipment, where geometrical aspects and introduction effects are more pronounced. A typical feature for MOVPE reactor design was the requirement to grow thin (< 100 A) layers necessitating rapid changes in gas-phase composition. These two factors has led to a great variety of MOVPE reactor designs, ten of which are listed in [20]. Flow visualization studies using TiO2 particles in reactors typical for MOVPE have also been carried out [21-23]. Early studies on silicon epitaxy [19] in a horizontal reactor produced a simple but elegant model to describe the growth process: adjacent to the susceptor one has a boundary layer in which the gas velocity is low. Species are transferred by diffusion to the substrate surface. Over this boundary layer a steep thermal gradient exists. Above the boundary layer one has the free flowing layer where the gas-velocity is high and temperature differences are smaller. Variations of the growth rate in the length or width of the susceptor can thus be attributed to variations in the thickness of the boundary layer or to variations in the concentration gradient over this same layer. A thorough study of the temperature distribution inside a horizontal VPE reactor using laser interference was published in 1982 [24]. In this work strong buoyancy driven features were shown when using a carrier gas of high density and low heat conductivity, such as nitrogen or argon. Also entrance effects and geometrical considerations for reactor design are described. The inlet regions of small scale reactors were formed into special shapes [25,26] while gas introduction through a porous plug [21] was shown to lead to more homogeneous introduction. For laboratory scale applications, the horizontal and the vertical pedestal reactor are still mostly

in use. Neither of these reactors is ideal when considering the interactive effects of forced and free convection. Thus, longitudinal changes in flow [27,28] can lead to layer thickness variations while gas recirculation effects [21,29] are detrimental for abrupt changes in gas composition. The return flow in the entrance region in horizontal reactors is brought about by abrupt changes in the ratio of the horizontal and vertical pressure gradients. Therefore, the effect is still present in the inverted geometry i.e. with the wafer facing down and the gas passing upwards from below the susceptor [30]. The advantage of the inverted geometry is that defects caused by particles falling on the wafer from wall deposits cannot occur. In the dual fluid layer inverted reactor [31] very good results on homogeneity of layer thickness and composition have been achieved without resorting to substrate rotation or reduced pressure operation. Notably for the vertical pedestal reactor, substrate rotation is required to ensure layer thickness uniformity. High speed rotation (> 500 rpm) is even more beneficial as this ensures a perfect axisymmetric rotating disc flow pattern and also provides pumping action [21]. For both the horizontal and vertical reactor the buoyancy driven convection effects can be eliminated by operation at reduced pressure. This ensures a more reproducible process, albeit at the expense of a higher degree of technical complexity. For large scale production by MOVPE, multiple wafer reactors have been developed. In the barrel reactor, directly derived from silicon VPE, the substrates are placed on a polygonal rotatable drum with nearly vertical sides contained in a cylindrical reactor. These reactors operate at reduced pressure. The pedestal reactor with high speed rotation has also shown merit for multiple wafer growth. Based on this layout, reactors are available for the growth of up to 17 substrates of 2 inch diameter per run [32]. Finally, the multi-substrate planetary reactor is available for the growth of up to 0.25 square metres substrate area per run [33]. The elegance of this reactor is that the gas-flow is used to produce the rotational motion of the susceptor while each substrate, mounted on a plattern, is made to rotate individually. The rotational motion of the individual substrates is again controlled by the gas-flow enabling fine tuning of the deposition profile over each individual wafer [34]. Improvements of MOVPE reactors are not only due to empirical optimization: advances made in computer modelling have provided valuable insight into the behaviour of the carrier gas flow. A review of this aspect of MOVPE emphasising transport phenomena in vapour phase reactors is given in [13]. The importance of understanding the velocity, temperature and pressure distribution inside a VPE reactor is that this knowledge can then be applied to determine the distribution of chemical species. C

STARTING COMPOUNDS AND CHEMICAL REACTIONS

The chemical reactions taking place during MOVPE growth are quite complex. It should be emphasized here that the final step in the path of the precursor molecules is the incorporation of the relevant atom in the appropriate site of the growing crystal lattice. In this Datareview, the discussion will be the growth of gallium arsenide. This material can be grown by every technology: chloride VPE (Datareview 16.3), LPE (Datareview 16.2), MBE (Datareview 16.6), CBE (Datareview 16.8) and MOVPE.

When describing the growth process, the reaction equation of LPE is quite simple: Ga(0 + As w - GaAs(s)

(1)

Describing the growth of GaAs by MBE is only slightly more complex: Ga + V2 As4 - GaAs(s) + Vi As2(v)

(2)

This is assuming that when elemental arsenic is evaporated, the tetrameric species is formed [35] which dissociatively adsorbs on the GaAs surface with a 'sticking coefficient' which is <0.5 [36]. The halide vapour phase technique uses elemental gallium as starting compound. The gallium is transformed into a volatile molecule by passing a halide vapour (e.g. HCl) over the elemental gallium source. The volatile Ga halide molecule is transported by the carrier gas to the substrate surface where dissociation takes place: the halogen species are desorbed while the gallium atom remains. The complementary group V atom is simultaneously transported to the surface as the gaseous AsH3 or as volatile AsCl3. One assumes that halide VPE growth takes place close to equilibrium and the group V species are assumed to be tetramers. An excess of tetramers with respect to other group V species has indeed been observed by mass spectroscopic analysis [37]. In MOVPE, the gallium source is the volatile trialkyl molecule. Both trimethylgallium and triethylgallium have a suitable vapour pressure to produce partial pressures of approximately 10"4 atm in the reactor. With arsine as the gaseous arsenic source one may write the overall reaction equation as follows: (CH3)3 Ga(g) + AsH 3(gr GaAs(s) + 3CH4(g)

(3)

In order to obtain a more detailed insight into this reaction one should consider the sequence of events when performing a growth experiment. After loading and purging the reactor the substrate is heated to the required growth temperature. This is done under an AsH3 flow i.e. under arsenic pressure, to prevent decomposition of the substrate due to the loss of the more volatile As species [35]. In this pregrowth period the surface oxide is removed and an arsenic covered surface is created. The arsine at this stage is mainly decomposed by a surface reaction. It has been shown [38] that hydrogen is liberated as the molecular species (H2) and that free hydrogen radicals, i.e. 'H are not yet formed at this stage. When the TMGa is switched into the reactor the molecule undergoes homolytic fission at temperatures above about 4500C: Ga(CH3)3 - "Ga(CH3), + 'CH3

(4)

The free methyl radical "CH3 forms a CH4 molecule by collision with the surrounding H2. Subsequently, a steady state concentration of "H radicals is built up in the reactor. These enhance the decomposition rate of TMGa, according to: •H + (CH3)3 Ga - 'Ga(CH3)2 + 'CH3

(5)

The dimethylgallium swiftly loses the second methyl group to form monomethylgallium 'Ga(CH3)2 - -GaCH3 + "CH3

(6)

Early work [39,40] made note of the relative instability of the metalorganic molecules but had not anticipated the importance of this free radical mechanism. Such short lived species will not be revealed in sampling experiments: in-situ infrared absorption studies are therefore required [41]. The kinetic parameters of TMGa decomposition have been studied extensively. Thus, the chemical role of the hydrogen ambient has been evaluated using N2 as a carrier gas [42], using a deuterium (D2) ambient and also when adding the radical producer TMIn as well as the radical scavenger 1,4 cyclohexadiene [43]. Finally, in the vapour phase a direct search for atomic species when decomposing gallium and indium alkyls using atomic adsorption spectroscopy was unsuccessful [44]. Whether the tri- or the mono-methyl species arrives on the surface is dependent on growth conditions, mainly the susceptor temperature and gas temperature. Thus, at low (< 6000C) substrate temperatures such as used in atomic layer epitaxy (ALE) [45] or in the UHV conditions of CBE, the complete molecule will be adsorbed on the surface. D

SURFACE CHEMISTRY

The GaAs surface shows reconstructions which are similar to those observed in the UHV conditions of MBE [46,47] and CBE [48]. The (4x2) Ga terminated, the (2x4) As terminated and the c(4x4) and d(4x4) arsenic rich phases have been identified using reflectance difference spectroscopy [49]. The c(4x4) to (2x4) transition has been observed in temperature stimulated desorption studies [50] and the extremely long range ordering of the c(4x4) surface was deduced from surface X-ray scattering [51]. This surface configuration persists up to temperatures of 575°C and remains when growth is performed, although the domain size becomes smaller [51]. Growth by ALE requires this surface configuration [50]; the (2x4) surface has an arsenic surface coverage of only 0.75 ML. However, when comparing MBE and MOVPE surface structures, one should be aware that on MOVPE grown surfaces the arsenic species still contain bound hydrogen atoms. It has been proposed that, notably at the arsenic terminating steps, hydrogen is still present. These 'B type' steps are chemically the most reactive. At these steps the monomethylgallium can undergo a double reaction: the hydrogen atom and the methyl group desorb as CH4 while the Ga atom is incorporated in the ledge of the lattice. That the hydrogen from the AsH3 plays an important role in removing the final methyl group has been shown by growth experiments in a helium carrier gas [52] and by the result that growing GaAs by MOVPE using elemental arsenic gives rise to heavily p-doped layers [53], due to carbon. In CBE, where the hydrides are decomposed to dimeric species [54] prior to impingement on the surface, p-doped GaAs layers are best grown by introducing small amounts of TMGa together with the TEGa used for the actual layer growth [55]. The ethyl-metal bond strength is considerably weaker than the corresponding metal-methyl bond. Also at low partial pressure the triethylgallium undergoes a simple homolytic fission reaction, releasing ethyl radicals. The gas phase decomposition does not go to completion i.e. gallium atoms are not formed in the gas phase [44]. The lower thermal stability of the TEGa molecule means that the temperature at which mass-transfer controlled growth is possible is shifted to lower values while the temperature range is more narrow. This applies even more strongly when using TiBGa [56]. Low growth temperature or the absence of gas phase reactions implies that the complete trialkyl molecule arrives on the arsenic terminated surface. This will be the case in CBE. The GaAs

growth rate shows a strong temperature and flow dependent growth rate. The decomposition mechanism of TEGa and the rate constants of the intermediate steps, as well as growth rate calculations, can be found in [57]. The salient feature of this model is that the incorporation in the lattice is by the gallium atom. In MOVPE growth using TEGa one could possibly still have the double reaction with desorption OfC2H6 as the final step. Thermal decomposition in the gas phase does not lead to formation of gallium atoms [44]. E

ALTERNATIVE PRECURSORS/SAFETY

As starting compounds the hydrides have the drawback of high toxicity. Arsine has a threshold limit value (TLV, 8 hour/day exposure limit) of 0.05 ppm. At a level of 250 ppm it is lethal in a breath. Phosphine is slightly less toxic, having a TLV value of 0.1 ppm. Both gases are heavier than air. Ammonia has a TLV value of 25 ppm. Thus, use of these gases requires stringent safety measures [58,59]. Extensive research has been carried out to develop alternative, alkyl-substituted group V sources. As pointed out in Section D, removal of the final alkyl-group from the group III precursor is aided by a hydrogen atom bound on the arsenic terminating ledge. Unfortunately, the trialkyl group V precursors do not provide satisfactory layer quality, notably on carbon contamination [60,61], whereas the partially subsituted hydrides are still toxic. Research is ongoing: recent developments in this field can be found in e.g. [62,63]. At the moment of writing, only the tertiary butyl compounds are considered acceptable, albeit expensive alternatives. Another approach is to fit an arsine (phosphine) generator [64] in the MOVPE system, an approach that obviously does not eliminate the necessity of gas detection and other safety measures. Another step which has been taken to reduce potential hazards in the MOVPE process is to replace the explosive hydrogen carrier gas with inert gases such as helium [52] and argon [65]. It has been proven that using well purified nitrogen, epitaxial layer quality is at least comparable to material grown in palladium diffused hydrogen [66]. F

FINAL REMARKS

For a more detailed knowledge of the MOVPE process, the reader is referred to excellent books [11,12] and review papers [13-16]. Important topics such as doping, the growth of alloys and the potential of MOVPE for growth of other materials have not been covered here. The continuous development of precursors [63,67] remains a key research area. The challenge for MOVPE technology is towards large scale production of solar cells and (blue) light emitting diodes. Certainly a wise route for MOVPE would be to develop reactors which are much more efficient with respect to the amounts of hydride required for growth. A pyrolysis oven for the precracking OfPH3 has been proposed in the early stage of MOVPE development [68] but this has not been pursued. With the present status of reactor modelling and the knowledge of kinetic parameters, a heated porous plug design for the hydrides to pass through could be tested, the precracking of AsH3 and PH3 being a prerequisite in CBE growth, so that the V/III ratios can be drastically reduced. On the other hand, MOVPE is excellently suitable for the growth of one- and zero-dimensional structures such as quantum wires [69,70], quantum dots [71] and fractional superlattices [72]. Due to the advent of imaging techniques such as AFM and STM morphological details on the

atomic scale are further revealed. The progress which has been made on in-situ monitoring techniques [49,51,73] in combination with the work carried out in MBE and CBE is giving more and more information on the dynamics of crystal growth. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] [II] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33]

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16.6 Unintentional hydrogen doping in MOCVD growth

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N. Watanabe and K. Wada March 1996

A

INTRODUCTION

It is well known that shallow donors and acceptors in GaAs can be passivated by hydrogen (H). In particular, C acceptors are easily passivated by hydrogen atoms during growth and post-growth cool-down [1-4]. In this Datareview, we examine investigations of H incorporation into MOCVD-grown GaAs, especially C-doped GaAs. B

HYDROGEN INCORPORATION INTO UNDOPED AND n-TYPE Si-DOPED GaAs

Concentration (cm"3)

FIGURE 1 shows SIMS profiles of Si and H atoms in Si-doped GaAs epilayers.

undoped GaAs

Si-doped GaAs

sub.

Depth (|nm) FIGURE 1. SIMS profiles of Si and H atoms in GaAs epilayers. The signal intensity for H indicates the background level.

The H concentration is below the detection limit of SIMS in both epilayers, indicating that H atoms are rarely incorporated into Si-doped or undoped GaAs during MOCVD growth. C

HYDROGEN INCORPORATION INTO p-TYPE C-DOPED GaAs

Cl

Hydrogen Incorporation During Post-growth Cooling

Hydrogen incorporation into C-doped GaAs during post-growth cool-down in an ambient has been investigated by Stockman et al [2]. The C acceptors have been reported to be deactivated by compensation with hydrogen donors, H+ [5], rather than by the formation of C-H complexes [3]. Stockman et al examined three kinds of ambient gas: AsH3ZH23 PH3ZH2, and H 2 on its own. The efficiency of C passivation with these ambient gases followed the order AsH3ZH2 > PH3ZH2 > H2. Moreover, the fraction of C that was passivated increased as the C concentration and the AsH3 supply increased. The H incorporation during the post-growth cool-down can be prevented in several ways. Stockman et al showed that an n-type GaAs layer can block indiffusion of H into C-doped GaAs during the cool-down [2]. A similar blocking effect can be obtained by using an undoped AlGaAs layer. FIGURE 2 shows the SIMS profiles of C and H in C-doped layers with and without an undoped AlGaAs layer on the surface. Although the C concentration is almost the same in both samples, the H concentration in the sample with an AlGaAs layer is about half of that without AlGaAs [6]. However, the mechanism of the blocking effect of H incorporation by n-type GaAs and undoped AlGaAs layers has not yet been clarified.

i-AIGaAs

Concentration (cm"3)

C-doped GaAs i-AIGaAs

C-doped GaAs

sub

sub

Depth (jim) FIGURE 2. The SIMS profiles of C and H atoms in a C-doped GaAs layer: (a) with an undoped AlGaAs layer as a top layer and (b) without an AlGaAs layer.

C2

Hydrogen Incorporation During Growth

Hydrogen atoms are also incorporated into a C-doped GaAs layer during MOCVD growth. These H atoms are reported to form neutral C-H complexes [3]. The incorporation of H during growth into intrinsic C-doped GaAs grown from different precursors and under different conditions was investigated by Kozuch et al [4]. They reported that the fraction of C that was passivated was independent of the Ga and As source gases, but did depend on the growth temperature. The temperature range from 500 to 6000C was investigated. Below 5500C, the fraction was estimated to be roughly 50%, and above 5500C it was 20 to 30%. Hobson et al also investigated how the precursor affected the incorporation of H into C-doped GaAs grown using CCl4 as a C source [I]. The fraction of C that was passivated was independent of the Ga source, while changing the As source from AsH3 to tertiarybutylarsine (TBAs) halved the fraction of C that was passivated.

Concentration (cm"3)

FIGURE 3 shows how the H concentration in C-doped GaAs grown using TMGa-AsH3-CCl4 as a precursor depends on the C concentration [7]. In this figure, all samples had an AlGaAs top layer to prevent H incorporation during the post-growth cooling. Over the entire C concentration range, the H concentration was about 20% of the C concentration. This result shows that the fraction of C passivated by H during growth is independent of the C concentration, a result which differs from the post-growth cooling. The H atoms passivating C acceptors can be easily removed by annealing, and C acceptors reactivated [8].

hole (as-grown) H (as-grown) hole (after annealing)

Carbon Concentration (cm"3) FIGURE 3. The dependence of hole and H concentrations on the C concentration. The as-grown samples included H atoms equal to about 20% of the C concentration, and these H atoms make the C acceptors inactive. However, the inactive C acceptors were reactivated after H atoms were removed from the epilayer by annealing at 600 0 C in an N2 ambient for 10 minutes. Reproduced from [7].

D

CONCLUSION

Hydrogen atoms are unintentionally incorporated into C-doped GaAs and passivate C acceptors. Incorporation of H atoms during post-growth cooling can be prevented by growing an n-type GaAs layer or an undoped AlGaAs layer on a C-doped GaAs layer. Incorporation of H atoms during growth is strongly dependent on the growth conditions, so that H incorporation is relatively small when the growth temperature is high. Hydrogen atoms incorporated into C-doped GaAs are easily removed by annealing. REFERENCES [1] [2] [3] [4] [5] [6] [7] [8]

W.S. Hobson, J. Pearton, D.M. Kozuch, M. Stavola [Appl Phys. Lett. (USA) vol.60 (1992) p.3259 ] S.A. Stockman, A.W. Hanson, S.L. Jackson, J.E. Baker, G.E. Stillman [Appl Phys. Lett. (USA) vol.62 (1992) p.1248] M. Stavola, D.M. Kozuch, CR. Abernathy, W.S. Hobson [Mater. Res. Soc. Symp. Proc. (USA) vol.249 (1992) p.75] D.M. Kozuch, M. Stavola, SJ. Pearton, CR. Abernathy, W.S. Hobson [ J. Appl. Phys. (USA) vol.73 (1993)p.3716] H. Fushimi, K. Wada [ J. Cryst. Growth (Netherlands) vol.145 (1994) p.420 ] N. Watanabe et al [ unpublished ] N. Watanabe, T. Nittono, H. Ito [ J. Cryst. Growth (Netherlands) vol.145 (1994) p.929 ] K. Watanabe, H. Yamazaki [ Joint Procs. on 7th State-of-the-Art Program on the Compound Semiconductors and on Logic and Functional Devices for Photonics at the 182nd Meeting of the Electrochemical Society, Toronto, Canada, 1992 (Academic Press, San Diego, USA, 1993) p.86-99]

16.7 High temperature MBE growth of GaAs K.R. Evans March 1996

A

INTRODUCTION

Molecular beam epitaxy (MBE) refers to the interaction(s) between one or more thermal atomic and/or molecular beams and the surface of a heated single crystal substrate in an ultra high vacuum (UHV) environment [I]. The MBE approach provides for the highly controlled growth of non-equilibrium layered structures, comprised of very thin, high quality films, with thicknesses accurate to the atomic scale, and with heterointerfaces which are smooth at, or nearly at, the atomic scale. Wide ranges of doping and composition are available. The UHV environment provides for the use of powerful in situ real-time monitoring techniques and also provides compatibility with other high vacuum processing such as ion implantation and ion beam etching. Solid source MBE growth of GaAs utilizes atomic Ga and molecular As2 or As4 beams. This Datareview is concerned with 'high' growth temperatures, i.e., temperatures above ~ 575 K. High temperature MBE growth of GaAs is used for the general purpose of producing very low defect density, high purity GaAs films with atomically smooth surfaces, both for high quality, doped and undoped buffer layers and also for use in the active regions of electronic and optoelectronic devices. So-called 'low temperature MBE GaAs', the subject of another chapter [2] in this book, is produced at temperatures below ~ 575 K, which results in incorporation of a relatively high concentration of point defects and complexes, which in turn results in novel structural, optical and electrical properties [3]. Several books concerned solely with MBE are published. One which provides an in-depth thermodynamic description of the fundamentals of MBE growth appeared in 1993 [4]. A comprehensive bibliography of early (1958-1983) publications concerned with MBE of III-Vs has been published [5]. Several reviews of the technology and fundamentals of MBE growth are also available [6-8]. While this Datareview is concerned only with 'solid source' MBE, it is noted that several derivative techniques exist, most notable of which are gas source MBE (GSMBE) [9], chemical beam epitaxy (CBE) [10], and metal organic MBE (MOMBE) [H]. B

GROWTH SYSTEM

MBE growth systems are commercially available worldwide and generally consist of a substrate loading chamber and a growth chamber with an intermediate substrate storage/preparation/ analysis chamber. Some older or less expensive MBE systems lack the intermediate chamber, known also as the buffer chamber. Most MBE chamber components are fabricated from stainless steel. Components, such as ovens, substrate holders and heaters, which will be held at elevated temperatures, are fabricated from high purity refractory metals and ceramics, especially Ta, W, Mo and pyrolytic boron nitride (pBN). The vacuum seals are primarily metal (usually Cu) gasket type. The pumping systems

vary somewhat, with ion pumps the most common. Very high purity sources are used and extreme care is taken to minimize exposure of the vacuum chamber to air. UHV conditions result from baking the chamber at temperatures up to 475 K for up to several days. Pressures as low as 10"9 Pa are obtained before depositing arsenic; after arsenic deposition the system base pressure is typically - 1 - 2 x 10"8Pa. The growth chamber contains liquid nitrogen cooled cryogenic panels to help pump residual impurities, especially water vapour, and may contain, in place of or in addition to the ion pump, a He cryopump, turbomolecular pump, titanium sublimation pump, or an oil diffusion pump. Prior to growth, the substrates are usually cleaned via solvent degreasing and lightly etched, then loaded into the sample loading chamber. (The use of 'epi-ready' substrates is gaining prevalence, whereby commercially available substrates are used 'as received'. Resulting defect densities are improved because of the elimination of surface contamination occurring during cleaning.) After the loading chamber is evacuated, one or more of the substrates are heated to up to 475 K to remove adsorbed water and other contaminants. The substrates are then transferred to the buffer chamber for further out-gassing and storage, and are subsequently moved into the growth chamber. The substrate heater in the growth chamber may be similar in design to those in the loading and buffer chambers and typically consists of an electrically heated Ta or W filament surrounded by pBN and Ta or W regions. Heat from the filament is radiatively coupled to the backside of the substrate or the substrate holder. The substrate is either soldered (with In or Ga) to a Mo block or is mechanically fastened via clips or wires to a Mo holder. The former approach enables the use of a thermocouple (TC) for temperature sensing with the TC contacting the backside of the Mo block, but suffers from morphological degradation of the substrate backside due to its interaction with the solder. The latter approach is used mostly today and necessitates 'non-contact' thermocouple sensing, or alternative temperature sensing approaches. Such non-contact thermocouple sensing results in a large uncertainty, up to 100 K, in the knowledge of the exact substrate temperature. The Ga beam is generated by heating liquid Ga, contained in pBN or graphite crucibles, at ~ 1255 K, which produces the flux required for - 1 monolayer (ML)/sec GaAs growth rate (IML = one Ga-As bilayer = 0.283 run), based on the Ga vapour pressure curve and given the typical MBE system geometry. Other group III sources, as well as Be and Si dopants, are handled similarly. The substrate is located - 20 cm from the front of the source ovens, which are referred to as Kcells [12], after Knudsen [13], even though their flux distribution characteristics deviate significantly from Knudsen behaviour. Because sharp doping profiles are somewhat difficult to obtain using Be due to its relatively high diffusion coefficient, especially at high doping levels ([Be] > 2 x 1019 cm"3), C doping is gaining in popularity [14]. Carbon doping is accomplished in solid source MBE by either an electron beam graphite source [15] or a resistively heated graphite source [16]. Until recently, the usual practice for generating an As beam was the use of a K-cell similar to that used for Ga, which results in an As4 beam flux of ~1015 cm"2 sec"1 for an As cell temperature range of ~ 600 - 640 K. Recently, commercial As cell technology has been enhanced in three ways: 1)

larger capacity cells are used to slow As source depletion, 2) an As4 - As2 cracking zone has been added to provide the dimer species during growth, and 3) a gas flow valve has been added between the evaporation and cracking zones, providing controllable, rapid time-scale modulation of the incident As flux. These valved-arsenic cracker cells are quite popular and significantly enhance the degree of control over the growth conditions. They are especially useful when the desired As flux varies for different layers within a structure of interest. Each MBE source must be extremely carefully prepared to minimize introduction of impurities into the process. Similarly, the group III and V sources used are typically of the highest purity (99.999% to 99.999999%) commercially available. A typical III-V solid source MBE system contains primary sources of Al, Ga, In, As, and sometimes P and dopant sources of Si, Be, and sometimes C. Mechanical shutters directly in front of the source ovens provide a mechanism for turning the beams on and off. The shutters can be opened and closed on a -100 ms time-scale so that submonolayer depositions can be controllably made. Beam fluxes can be monitored by an ion gauge which can be rotated into the beam paths when a film is not being grown. The beam densities in general are not uniform across the substrate surface. To create a more uniform beam density, significant engineering of the geometry of the source ovens and their position relative to the substrate has been performed [8]. Additionally, the substrate is rotated during growth. Immediately before growth, the native oxide on the substrate surface is desorbed by heating to ~ 875 K in an As2 or As4 flux. Then the desired growth temperature is set, the time to settle to that temperature is awaited, and GaAs growth is initiated by opening the Ga beam shutter. The transition from GaAs growth to InGaAs growth, for example, can be accomplished simply by opening the In shutter once the desired GaAs thickness is produced. However, often a new substrate temperature and As flux is desired, and appropriate changes may be made during a growth interrupt, which is accomplished by closing all but the As shutter. Care must be taken to minimize growth interruption time because of the concomitant incorporation of impurities due to the imperfect background vacuum level. Compromises are often made; for example, the temperature might be ramped down during the last part of the GaAs growth, to get ready for the InGaAs growth. In fact, growth of heterostructures in general necessitates compromise in growth condition optimization, since the optimal growth conditions vary for each layer within a heterostructure. Thus, a principal duty of the MBE grower is to work closely with the growth requester to construct a combination of growth condition compromises which optimize the most important characteristics of the structure. C

GROWTH CONDITIONS

Growth conditions are nearly completely described by the incident fluxes and composition (i.e., Ga and As2 or As4) of the source beams, the substrate composition and orientation, the substrate temperature, and the background vacuum level. The growth chamber pressure must be low enough to provide for collisionless flow to the substrate, which requires in practice a level below approximately 5 x 10'2 Pa; however, considerations of impurity incorporation from background gases indicate the need for a significantly lower background pressure.

Most MBE growth of GaAs is performed on high quality GaAs substrates of (001) orientation and slight intentional mis-orientations from (001), which are readily commercially available in diameters up to 3 inches, in both semi-insulating and conductive (p- and n-type) form. A significant effort has been made to grow GaAs and related compounds on GaAs ( H l A ) and (11 IB) substrates [17], largely because of the strong electro-optic effects predicted for this orientation [18]. GaAs substrates with slight mis-orientations from (001) have been shown to give rise to improved electrical and optical properties for AlGaAs/GaAs heterostructures and AlGaAs epilayers, especially for misorientations towards (11 IA) [19]. GaAs films have also been produced on Ge substrates, which provide a relatively close (0.13%) lattice-match to GaAs [20,21]. The discussion below is concerned primarily with growth on GaAs (001) substrates. The incident beam fluxes are generally chosen such that the time required to deposit a single atomic layer, or monolayer, is on the order of 1 second. In practice the As/Ga incident flux ratio is chosen above unity, since the excess As atoms desorb as As2 or As4, resulting in a stoichiometric film. The Ga flux determines the growth rate and is 6 x io 14 cm"2 sec"1 for a growth rate of 1 ML/sec. Use of both As2 and As4 is quite common. While As2 has been shown to improve the quality of GaAs [22,23] and GaAs/AlGaAs heterostructures [24], high quality electronic and optoelectronic device structures have been produced using As4. The improvement offered by As2 over As4 is partially offset by the additional outgassing due to the hot cracking zone. Another benefit OfAs2 is a two-fold increase in As incorporation efficiency relative to As4 [25]. Substrate temperatures for GaAs growth are typically between 675 and 915 K, with 855 - 875 K most common. In this range the Ga surface diffiisivity is significantly high to enable two dimensional growth with very low (< 1014 cm"3) concentration of vacancies and antisite defects. Temperatures above 915 K result in a reduced GaAs growth rate due to significant Ga desorption [26]. Higher GaAs growth temperatures may be chosen when growing AlGaAs/GaAs heterostructures, since the optimal AlGaAs growth temperature is near 975 K. Lower growth temperatures may be used for highly doped layers of GaAs:Be and GaAs:Si, for which high growth temperatures result in significant redistribution of Si [27] and especially Be [28]. Desorptive loss of Si during AlGaAs: Si growth at high temperatures has also been reported [29]. The background vacuum must be kept very clean in order to realize high purity films of GaAs and related compounds. Since the Ga flux at the growth surface, when converted to pressure units, is approximately 5 x 10"5 Pa, a background vacuum level of approximately 5 x io~9 Pa might be expected to give an impurity level in the GaAs film of 1 part in 104, which translates to approximately 1 x 1018 cm"3 impurity level. However, many device applications require impurity levels below 1 x io 15 cm"3. Fortunately, not all background impurities incorporate readily at the growth surface. Background levels of CO, CO2, H2O and H2 are generally discernible in the growth chamber via quadrupole mass spectrometry, but they do not tend to incorporate into GaAs films during MBE growth. D

GROWTH MECHANISM

Significant research has been carried out to understand the microscopic processes giving rise to

epitaxial growth, with Si and GaAs the most studied materials systems. Many of the existing fundamental studies into GaAs MBE utilize either a kinetic or thermodynamic analysis. Dl

Kinetic Analysis of Growth

The kinetic view [30] of GaAs growth with As2 indicates that As2 is dissociatively chemisorbed on exposed Ga adatoms. Impinging As2 molecules which do not encounter an exposed Ga atom quickly desorb. Thus, the sticking coefficient OfAs2 depends on the Ga coverage and becomes unity when the Ga flux is twice the As2 flux. Stoichiometric GaAs is obtained for an incident As2/Ga flux ratio of 0.5 at relatively low temperatures. For higher growth temperatures significant As evaporation occurs, and thus the incident flux ratio must be set higher than 0.5 to maintain stoichiometry. In practice the flux ratio is set higher than the minimum necessary to maintain stoichiometry, and any excess As at the surface simply desorbs. A detailed surface phase diagram for (001) GaAs has been constructed [31] and indicates the required minimum flux ratio needed to grow stoichiometric GaAs as a function of substrate temperature. The kinetic view of GaAs growth using As4 indicates that pairs of As4 molecules interact on adjacent exposed Ga atoms and react to leave behind two As2 molecules which subsequently desorb. Thus, the maximum As4 sticking coefficient is 0.5, and the minimum As4/Ga incident flux ratio required for stoichiometric growth is 0.5. The most important surface processes occurring during growth are considered [32] to be: i)

adsorption of the constituent atoms or molecules impinging on the surface;

ii)

surface migration and dissociation of the adsorbed molecules;

iii)

incorporation of the constituent atoms;

iv)

thermal desorption of species not incorporated into the crystal lattice.

The growth temperature is usually chosen to be below - 915 K, resulting in a near unity Ga sticking coefficient (defined as the ratio of the number of Ga atoms adhering to the substrate surface to the number of Ga atoms arriving there). In this temperature range the GaAs growth rate is determined by the incident Ga flux, assuming enough As is supplied. Pooling of Ga occurs when insufficient As is supplied, and Ga desorption becomes significant when temperatures above - 9 1 5 K are used. The surfaces of growing GaAs films tend to be reconstructed as evidenced by in-situ reflection high energy electron diffraction (RHEED) analysis [33]. Additionally, periodic oscillations in the intensities of various features of the RHEED pattern, especially of the specular beam [34], during growth have been found to correspond to the monolayer deposition time. As such, RHEED oscillations are a powerful tool for monitoring growth rate. The advent of scanning tunnelling microscopy and its derivative techniques has given rise to fundamental investigations of the evolution of GaAs surface morphology during MBE growth. Issues of step bunching, step flow growth, and island shape anisotropies have been investigated [35].

D2

Thermodynamic Analysis of Growth

In the thermodynamic approach to MBE growth [36] of GaAs, the basic reaction is: Ga(g) + 1/2 As2(g) = GaAs(s). Growth takes place far from equilibrium and the incident fluxes correspond to pressures which are significantly higher than the corresponding equilibrium pressures. At a fixed As pressure, the growth rate is determined by the difference between the Ga pressure and its equilibrium value. E

IN-SITU ANALYSES

Several recent studies have focused on the use of in-situ sensors for real-time monitoring and control of growth conditions. Such work is most important for optimizing the growth conditions and improving the reproducibility of heterostructure growth. While RHEED [33,34] is a proven technique which provides much information, its value is limited to discontinuous observations and, additionally, sample rotation presents a difficulty. Non-RHEED techniques which directly probe the growing film include ellipsometry [37], pyrometric interferometry [38], photoemission oscillations [39], optical transmission spectroscopy [40], time-of-flight low energy ion scattering [41], and reflection difference spectroscopy [42]. Incident [43,44] and desorbed [45] Ga flux and incident [44] and desorbed [46,47] As flux have been determined via hollow cathode lamp light absorption spectroscopy [43], laser ionization plus time-of-flight mass spectrometry [44,47], line-of-sight mass spectrometry [45,46,48] and laserinduced fluorescence [49]. Significant advances have been made in real-time monitoring of growth parameters; however, actual demonstrations of feedback control are only beginning to be realized. Real-time integration of incident group III fluxes to give real-time monitoring of layer thickness [43,50] has been achieved, allowing real-time determination of shutter closing time to provide the desired thickness, but real-time control of incident fluxes has not been achieved. Only for the case of group III desorption has real-time feedback control been demonstrated [48]. F

FUTURE GaAs MBE

Future MBE systems for producing GaAs and related compounds will have multiple in-situ sensors for determination of incident and desorbed fluxes, surface composition, layer thickness(es), and surface temperature. These measurements will be used in conjunction with appropriate computer modelling to provide feedback to devices controlling incident fluxes, substrate temperature, and shutter event occurrences. Since composition control on the atomic distance scale is desired, which corresponds to a ~/ second growth parameter control time-scale, the sensors and flux control devices will operate on a sub-second time-scale. The valved group V sources are quite capable of making flux changes on this time-scale, as are the present nonbonded substrate heating systems and recently developed small thermal mass dopant sources. Group III sources are presently under development to enable such rapid incident flux changes [51]. REFERENCES [1]

A. Y. Cho, J.R. Arthur [Prog. Solid State Chem. (UK) vol.10 pt.3 (1975) p.157-91 ]

[2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33] [34] [35] [36] [37] [38] [39] [40] [41] [42] [43] [44] [45]

[ Datareviews in this book: Chapter 17: Low Temperature MBE GaAs ] D. Look [ Thin Solid Films (Switzerland) vol.231 (1993) p.61-73 ] J. Y. Tsao [Materials Fundamentals of Molecular Beam Epitaxy (Academic Press, 1993) ] K. Ploog, K. Graf [Molecular Beam Epitaxy of Ul-VCompounds (Springer-Verlag, 1984) ] E.H.C. Parker (Ed.) [ The Technology and Physics of Molecular Beam Epitaxy (Plenum Press, 1989)] E. Kasper,J.C. Bean (Eds.) [Silicon Molecular Beam Epitaxy (CRC Press, 1988; vols.I& II ] M.A. Herman, H. Sitter [ Molecular Beam Epitaxy: Fundamentals and Current Status (SpringerVerlag, 1989)] M.B. Panish [ J. Electrochem. Soc. (USA) vol. 127 (1980) p.2729 ] W.T. Tsang [ Appl. Phys. Lett (USA) vol.45 (1984) p. 1234 ] E. Tokumitsu, Y. Kudou, M. Kongai, KTakahashi [J. Appl. Phys. (USA) vol.55 (1984) p.3163 ] M.A. Herman [Vacuum (UK) vol.32 no.9 (1982) p.555 ] M. Knudsen [Ann. Phys. (Germany) vol.29 (1909) p. 179] T.H. Chiu, J.E. Cunningham, J.A. Ditzenberger, W.Y. Jan, S.N.G. Chu [ J. Cryst. Growth (Netherlands) vol. 111 (1991) p.274 ] J.M. Van Hove, P.P.Chow, M.F.Rosamond, G.L. Carpenter, L.A.Chow [ J. Vac. Sci. Technol. B (USA) vol. 12 (1994) p. 1200 ] RJ. Malik etal [J. Cryst. Growth (Netherlands) vol.127 (1993) p.686] K. Elcess, J.-L. Lievin, CG. Fonstad [ J. Vac. Sci. Technol. B (USA) vol.6 (1988) p.638 ] C. Mailhiot, D.L. Smith, T.C.McGill [ J. Vac. Sci. Technol. B (USA) vol. 1 (1983) p.637 ] R.K. Tsui, J.A. Curless, G.D. Kramer, M.S. Peffley, D.L. Rode [ J. Appl. Phys. (USA) vol.58 (1985)p.2570] H. Kroemer, KJ. Polasko, S.L.Wright [Appl. Phys. Lett. (USA) vol.36 (1980) p.763 ] D.L. Miller, J.S. Harris Jr. [ Appl. Phys. Lett. (USA) vol.37 (1980) p. 1104 ] H. Kuenzel, K. Ploog [ Appl. Phys. Lett. (USA) vol.37 (1980) p.416 ] L.P. Erikson et al [ J. Appl. Phys. (USA) vol.56 (1984) p.2231 ] G. Duggan, P. Dawson, CT. Foxon, G.W't. Hooft[J. Phys. Colloq. (France) vol.43 no.C-5 (1982) p. 129] CT. Foxon, B.A. Joyce [ Surf. Sci. (Netherlands) vol.77 (1975) p.434 ] J.M. Van Hove, P.I. Cohen [ Appl. Phys. Lett. (USA) vol.47 (1985) p.726 ] K. Inoue, H. Sakaki, J.Yoshino, Y.Yoshioka [ Appl. Phys. Lett. (USA) vol.46 (1985) p.973 ] M. Ilegems[J. Appl. Phys. (USA) vol.48 (1977) p. 1278] N. Chand et al [ Phys. Rev. B (USA) vol. 30 (1984) p. 4481 ] CT. Foxon, B.A. Joyce [ Surf. Sci. (Netherlands) vol.50 (1975) p.434 ] S.M. Newstead, R.A.A. Kubiak, E.H.C. Parker [J. Cryst. Growth (Netherlands) vol.81 (1987) p.49] A. Madhukar [ Surf Sci. (Netherlands) vol. 132 (1983) p.344 ] A. Y. Cho, J.R. Arthur [ Prog. Solid State Chem. (UK) vol. 10 (1975) p. 157 ] C.E.C. Wood [ Surf. Sci. (Netherlands) vol. 108 (1993) p. 1032 ] BG. Orr[J. Cryst. Growth (Netherlands) vol.127(1993) p. 1032 ] R. Heckingbottom[J. Vac. Sci. Technol. B (USA) vol.3 (1985) p.572 ] G.N.Maracasetal[J. Vac. Sci. Technol. A (USA) vol.13 (1995) p.727 ] AJ. SpringThorpe, A. Majeed [ J. Vac. Sci. Technol. B (USA) vol.8 (1990) p.266 ] J.N. Ecstein, C Webb, S.-L. Weng, K.A. Bertness [Appl. Phys. Lett. (USA) vol.61 (1992) p. 1685 ] E.S. Hellman, J.S. Harris [ J. Cryst. Growth (Netherlands) vol.81 (1987) p.38 ] K. Waters, A. Bensaoula, A. Schultz, K. Eipers-Smith, A. Freundlich [ J. Cryst. Growth (Netherlands) vol.127 (1993) p.972 ] D.E. Aspnes [ J. Vac. Sci. Technol. B (USA) vol.3 (1985) p. 1498 ] S.A. Chalmers, K.P. Killeen [ Appl. Phys. Lett. (USA) vol.63 (1993) p. 1498 ] P.G. Strupp, A.L. Alstrin, R.V. Smilgys, S.R. Leone [ Appl. Opt. (USA) vol.32 (1993) p.842 ] AJ. SpringThorpe, P. Mandeville [J. Vac. Sci. Technol. B(USA) vol. 6 (1988)p.754 ]

[46] [47] [48] [49] [50] [51]

J.Y. Tsao, T.M. Brennan, B.E. Hammons [Appl. Phys. Lett. (USA) vol.53 (1988) p.288 ] A.L. Alstrin, P.G. Strupp, S.R. Leone [Appl. Phys. Lett. (USA) vol.63 (1993) p.815 ] K.R. Evans [Mater. Res. Soc. Symp. Proc. (USA) vol.340 (1994) p. 13 ] K.L. Carleton, S.R. Leone [ J. Vac. Sci. Technol. B (USA) vol.5 (1987) p. 11441 ] FG. Cellii, Y.C. Kao, E.A. Beam, W.M. Duncan, T.S. Moise [J. Vac. Sci. Technol. B (USA) vol.11 (1993) p. 1018] CR Jones, D.L. Beasley, E.N. Taylor, K.R. Evans [ J. Vac. Sci. Technol. B (USA) vol. 13 (1995) p.739]

16.8 Metalorganic molecular beam epitaxy and atomic layer epitaxy of GaAs CR. Abernathy December 1995

A

EVTRODUCTION

One of the disadvantages of metalorganic chemical vapour deposition (MOCVD) is the tendency for the precursors to interact prior to arrival at the wafer surface. In addition to reducing the efficiency of incorporation by depleting the gas stream of reactants, this interaction also makes the attainment of good across-wafer and wafer-to-wafer uniformity difficult. Furthermore it can impose limits on the shape and magnitude of the dopant concentrations. Two schemes for eliminating gas-phase pre-reactions have been explored in recent years. One involves performing the growth in ultra high vacuum so that molecular rather than viscous flow can be achieved. This technique is commonly referred to as metalorganic molecular beam epitaxy (MOMBE) [1] or alternatively chemical beam epitaxy (CBE) [2]. The second approach involves the use of alternating flows of group III and group V reactants in order to prevent gas-phase interactions between them. The growth is typically carried out under conditions in which the growth is selflimiting, producing a growth rate of one monolayer per cycle. Hence it has been given the name atomic layer epitaxy (ALE). This Datareview will outline the relative strengths and weaknesses of these two techniques. B

METALORGANIC MOLECULAR BEAM EPITAXY

The use of ultrahigh vacuum (UHV) growth techniques for deposition of semiconductors offers several desirable features. The absence of gas phase interactions allows for precise control of both thickness and composition beyond what can be achieved by other growth methods. In addition, the use of substrate rotation produces excellent uniformity of ± 1.5% across substrates up to 4 inches in diameter. These advantages have made molecular beam epitaxy (MBE) the dominant method for growth of a number of semiconductor device structures, particularly those involving GaAs-based materials, which require abrupt interfaces and a high degree of compositional control. In spite of these advantages, MBE still retains a number of disadvantages which limit its applicability. The need to open the chamber for replenishment of the elemental sources, particularly As, can lead to substantial downtime. Additional problems associated with the use of elemental group HI sources include poor morphology due to surface defects generated from the use of group III effusion ovens. There is also limited potential for scale-up due to the large source-to-substrate distances, and hence large reactor dimensions, which are required when deposition over large areas is attempted using liquid metal effusion ovens. In response to these problems, the replacement of elemental group III sources with gaseous sources, primarily the metalorganic (MO) compounds commonly used in metalorganic chemical vapour deposition (MOCVD), was initiated, creating a new hybrid technique called metalorganic MBE (MOMBE) [1,2]. As discussed above, the use of molecular beams of gaseous precursors also eliminates many of the problems associated with MOCVD, in principle allowing MOMBE

to retain the benefits of both of the parent techniques while avoiding many of the limitations. While much of this promise has proven true, new difficulties have arisen which are unique to MOMBE. Bl

Group III Sources

One of the major problems for growth of GaAs-based materials by MOMBE has been impurity incorporation. Unlike MOCVD, where high growth temperatures and the presence of scavenging reactions in the gas phase limit the incorporation of carbon, MOMBE has no comparably effective in-situ mechanisms for impurity gettering. As a result, the surface chemistry and hence choice of precursors must be carefully tailored to the specific conditions of MOMBE. The lowest reported backgrounds for the most commonly used gallium precursors are listed in TABLE 1. As shown in FIGURE 1, the carbon background decreases with increasing growth temperature up to 873 K [3,7-9]. Above this temperature, the alkyl ligands begin to undergo significant decomposition resulting in an increase in the carbon concentration. In general, weakening the carbon-Ga bond in the precursor reduces the carbon content in the grown film. As will be discussed in Section B3, GaAs grown from trimethylgallium (TMG) contains extremely high levels of carbon [4-6] which limits its use to the growth of p-type material. Triethylgallium (TEG), tri-isopropylgallium (TIPG) or tri-isobutylgallium (TIBG) is used for n-type or undoped GaAs. TABLE 1. Minimum carbon concentrations reported for GaAs grown by MOMBE from various Ga and As sources. Concentrations were determined by SIMS except where noted. Gallium source

Arsenic source

Growth temperature (K)

TMG

AsH2

773

Carbon concentration (cm'3) 5 x IQ16 [20]

TMG TEG TIPG TIBG

DMAAs t AsH2 AsH2 AsH,

773 833 833 773

8 * IQ15 [20] 2xlQ"t[3] 1-7 x IQ14* [8] 1.5 x IQ16* [7]

I

*SIMS detection limit. ^Determined from Hall measurement. tTrisdimethylaminoarsenic.

It has been shown by a number of authors that the Ga incorporation rate in GaAs from TEG decreases markedly as the growth temperature is reduced below 723 K [10,11], as can be seen in FIGURE 2. As the temperature is reduced, pyrolysis of the Ga-ethyl bonds becomes increasingly inefficient and absorbed alkyl Ga species leave the surface without being incorporated into the growth front. Because of the weaker Ga-C bond in TIBG, the Ga incorporation rate from this source does not drop as rapidly with decreasing temperature. Above ~ 863 K, the growth rate from TEG again begins to drop [10] due to desorption of alkyl-Ga from the surface [10-12]. Because of the relative insensitivity of growth rate to growth temperature near the typical growth temperature, the thickness uniformity of GaAs grown by this method is quite good with total variations in thickness over 3 inch diameter wafers generally less than 3% (or ± 1.5%). Though TIPG and TIBG appear promising, their use is not widespread for growth of GaAs due in part to their cost, poor efficiency (in the case of TIPG), and low vapour pressure (TIBG). Consequently, trimethylgallium (TMG) and triethylgallium (TEG) are at present the most commonly used Ga precursors.

B2

Group V Sources

The standard As source in MOMBE is AsH3. Due to the thermal stability of this compound, precracking is required in order for growth to occur [13-17]. The most common cracking method is the catalytic, or low pressure, approach. The gas is introduced at the back of an injector cell and flows around a catalytic material, typically Mo or Ta. When heated to temperatures of -1073 K or higher, this catalyst allows for the efficient decomposition of the hydrides to group V dimers or monomers [15,16], even though the pressure inside the cracker cell is only ~10"3 torn These cracker cells have approximately the same dimensions as the standard MBE effusion ovens, and therefore may be used as direct replacements for other cells in the MBE system.

CARBON CONCENTRATION (cm 3)

Because of the extreme toxicity of the group V hydrides, alternative gaseous precursors have been actively sought for some time. The most commonly used replacement for AsH3 in MOCVD is tertiarybutylarsine (TBAsH2), which is simply a monosubstituted arsine. Though MOMBE growth of GaAs over a limited temperature range has been reported using this compound, the poor decomposition efficiency mandates the use of precracking [18,19] as described for AsH3. In contrast to the alkyl bonded substituted arsine compounds, trisdimethylaminoarsenic (DMAAs)

GROWTH TEMPERATURE (K)

FIGURE 1. Variation of carbon concentration (as determined by SIMS) with growth temperature for GaAs grown by MOMBE using two different alkyl-Ga precursors. The As source was AsH3. (After [7].)

GROWTH RATE (A/min)

TEG

GROWTH TEMPERATURE (K) FIGURE 2. Dependence of MOMBE GaAs growth rate from TEG on growth temperature. (After [H].)

which has the formula As{N(CH3)3}3 does appear to decompose quite efficiently over the temperature range generally used for growth in MOMBE [20,21]. GaAs layers grown from DMAAs exhibit lower carbon backgrounds than similar material grown with cracked AsH3. This effect is particularly dramatic for growth from TMG. This reduction in carbon background is believed to be due to the presence of reactive atomic hydrogen species generated at the surface by the P-hydride elimination from the DMAAs [22]. This hydrogen suppresses the uptake of carbon either by formation of volatile CH4 or by blocking adsorption sites which would otherwise be occupied by carbon. Above 723 K, no nitrogen incorporation is observed. B3

Doping

B3.1

Elemental doping sources

The incorporation of elemental dopants in MOMBE is similar to the behaviour observed in MBE, suggesting little interaction between the elemental sources and adsorbed hydrocarbon species on the growth surface. However, for some elements the charge in the effusion cell does appear to react over time with hydrocarbons desorbed from the growth surface to form carbides [23]. This results in a reduction of the dopant flux for any given cell temperature. Another potential problem with elemental doping cells is the formation of volatile dopant-carbon compounds [23]. Unintentional Sn incorporation at concentrations up to 1018 cm"3 has been observed in InP, grown in a system in which the Sn cell was heated but kept shuttered. This anomalous doping behaviour

is believed to result from the formation of alkyl Sn species on the Sn cell shutter. For these reasons, gaseous dopant sources are usually employed. B3.2

Gaseous dopant sources

B3.2.1 n-type doping Both SiH4 and Si2H6 have been explored extensively for use in the GaAs/AlGaAs system [24-27]. As expected, the maximum doping level and activation efficiency are similar to those obtained with elemental Si, though the doping efficiency is quite low, (~ 10'5), for both precursors, probably due to the stability of the adsorbed SiH3 molecule [24]. Other n-type gaseous sources such as H2S, diethyltellurium (DETe) and tetraethyltin (TESn), also suffer from poor decomposition efficiency, but appear to decompose at a rate sufficient for growth of high-quality layers. TESn is particularly attractive because of the high uncompensated doping levels (up to 1.5 x 1019 cm"3) that can be obtained with Sn [24,29,30]. Although the incorporation efficiency of this source is still somewhat low (2 x 10"3) for GaAs [24], it has been shown that the carrier concentration [24,29,30] can be reproducibly varied over a wide range by varying the TESn flow rate. The dopant concentration increases linearly with increasing TESn flux, which suggests that the incorporation efficiency is independent of Sn surface concentration. This linear variation combined with the absence of a strong dependence of carrier concentration on temperature [30] makes dopant control relatively easy using this source. Moreover, unlike elemental Sn, TESn produces very abrupt dopant transitions [30]. B3.2.2 p-type doping For GaAs and AlGaAs, carbon has proven to be an excellent p-type dopant which can be introduced quite easily in MOMBE through the use of strongly bonded metal-alkyl sources like TMG [4-6] or halocarbons such as CCl4 [31-34] or CBr4 [33-37]. Because of the low growth temperatures which can be employed in MOMBE and the low diffusivity of carbon, very high hole concentrations, up to 1021 cm"3, can be achieved using either type of carbon source [5,33]. With TMG, it has been shown that the hole concentration decreases dramatically with increasing temperature [38] or increasing V/m ratio [39]. The variation with V/III ratio is usually used to control the dopant concentration over a rather wide range. Unfortunately, while this approach works well for controlling the hole concentration, it has a dramatic effect on the growth rate as well [39], making calibration and control more difficult than when using conventional, extrinsic, p-type dopant sources, like the halogenated methyl sources. DeLyon et al [31,35] have reported that neither growth temperature nor V/IU ratio has a significant impact on the hole concentration of GaAs grown from solid Ga and CCl4 [31] or CBr4 [35]. CBr4 was shown to produce significant improvement in efficiency relative to TMG and CCl4 by factors of 750 and 150 respectively. Furthermore, since the carbon incorporation from the halogen sources has been shown to be linearly dependent on source flux [34,35], calibration of the dopant concentration is more straightforward. Their utility is compromised however by the existence of parasitic etching reactions which reduce the growth rate as the dopant flow is increased. As a result, the dependence of growth rate on hole concentration is just as problematic for CCl4 and CBr4 as for TMG, making calibration of the growth rate equally difficult for the various compounds. Other compounds like tetrazole [40] and neopentane [41] have also been explored. Tetrazole

has not yet been successful due to its low vapour pressure [42] and its tendency to explode when heated above the sublimation temperature. Neopentane has been successfully used in the MOMBE growth of GaAs, though due to its stability pre-cracking is required [41]. The hole concentration is linear with neopentane flow rate and according to the authors yields an acceptor activation of close to 100%. However, the carbon uptake rate shows a strong dependence on growth temperature such that temperatures below 673 K are required in order to achieve hole concentrations in excess of 1020 cm"3. C

ATOMIC LAYER EPITAXY

FILM THICKNESS PER CYCLE (A)

Atomic layer epitaxy (ALE) was developed as a means of overcoming some of the problems associated with the mixing of the group HI and group V flows prior to arrival at the wafer surface during growth by MOCVD. Separation of the flows was expected to provide better interfacial abruptness and better across-wafer uniformity. This effect can be accomplished either by pulsing the reactant flows and purging with hydrogen between the pulses, or by rapid rotation of the substrate through separated flow streams. Both techniques usually take advantage of the selflimiting growth behaviour obtained when using trimethylgallium (TMG) under certain growth conditions. Self-limiting behaviour further enhances the expected improvements of uniformity and thickness control. Techniques which rely on flow separation but are not self-limiting (i.e. nonsaturated) are often referred to as ALE, but in the truest sense of the term are more properly

TMG ADMITTANCE RATE (Pa 1/sec)) FIGURE 3. Atomic layer epitaxy film thickness per cycle versus TMG admittance rate for various substrate temperatures. A single monolayer of GaAs should correspond to roughly 2.8 A.

called pulsed flow or pulsed jet epitaxy. Most of this section will focus on true ALE of GaAs; however non-saturated growth will be discussed with regards to carbon doping. Cl

Growth of GaAs by ALE

Self-limiting growth occurs when the growth rate becomes independent of exposure time to the group in beam [43-45], as shown in FIGURE 3. For TMG, this situation can be achieved over a fairly wide range of conditions. For less thermally stable compounds, like TEG, such behaviour is not usually observed in the viscous flow pressure regime [46]. It is believed that the TMG, which requires more surface catalysis than TEG, is not catalyzed by the group III rich surface. Thus coverage of the growth surface with Ga or alkyl-Ga species effectively poisons the catalytic action of the surface, allowing subsequently arriving TMG molecules to leave the surface without decomposing. Hence, the maximum growth rate for any one cycle is one monolayer. If the TMG is decomposed in the gas phase, due to radiant heating for example, the self-limiting aspect of the growth is lost and the growth rate can exceed this value [47] (see FIGURE 3). For this reason, the temperatures used in most ALE applications, particularly those employing flow modulation, must be kept significantly lower than those typically used in MOCVD. The use of wafer rotation to establish the growth cycle limits the amount of time the substrate is exposed to the growth temperature and to the group III flow [48]. This allows for deposition at higher growth temperatures which is desirable due to the reduced tendency for carbon incorporation as the temperature is increased. In addition to temperature, several other conditions must be met in order for true ALE to occur. Most importantly, there must be little or no carryover of one reactant into the following portion of the cycle in order to avoid As-induced catalysis of the arriving TMG. Thus purging considerations limit the cycle time which can be employed. For this reason, pulsed flow growth is usually limited to growth rates of 0.1 - 0.2 \im/hi. Use of wafer rotation rather than flow switching can increase the cycle time and allow growth rates up to 0.4 jim/hr [48]. Even for this approach, however, the growth rate is considerably less than that typically used in other epitaxial techniques, making this a significant limitation for deposition of device structures. Purging considerations also limit the magnitude of the reactant flows. Low V/III ratios are normally required since large AsH3 flows require longer purge times. Though higher V/III ratios help to reduce the carbon contamination, the longer growth interruptions tend to increase contamination. Use of the wafer rotation scheme tends to allow the use of higher AsH3 flows. The higher V/III ratios (V/III - 130) and higher growth temperatures (853 - 873 K) which can be used in the rotation technique have produced some of the cleanest ALE-grown material to date with background electron concentrations of 1015 cm"3 and 77 K mobilities of 30000 cm2/ V sec [48]. As expected, the thickness uniformity obtained with ALE is reported to be excellent. C2

Doping

Because of the interaction of the As-hydrides and the methyl species at the growth surface, it is difficult to achieve high hole concentrations due to carbon acceptors in material grown by ALE. Chung et al [49] have shown that the maximum hole concentration can be increased from 2 x 1017 cm"3 to 1019 cm"3 if the growth is moved from saturated (low alkyl flow/pulse) to non-saturated (alkyl flow/pulse > 2 ^imol) conditions. At atmospheric pressure, growth temperatures of 723 748 K are found to produce the highest carbon content, with the level independent of temperature

and AsH3 flow in this temperature regime. D

CONCLUSION

MOMBE and ALE are promising methods for deposition of III-V materials. Both methods require optimization in order to limit impurity backgrounds to acceptable levels. These techniques offer the ability to achieve well-confined n- and p-type doping profiles. MOMBE in particular is capable of producing very high doping levels with excellent doping and thickness uniformity. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] [II] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25]

E. Veuhoff; W. Pletschen, P. Balk,, H. Luth [ J. Cryst. Growth (Netherlands) vol.55 (1981) p.3O ] W.T. Tsang [ Appl. Phys. Lett. (USA) vol.45 (1984) p. 1234 ] Y.M. Houng [ J. Cryst. Growth (Netherlands) vol. 105 (1990) p. 124 ] N. Putz, H. Heinecke, M. Heyen, P. Balk, M. Weyers, H. Luth [J. Cryst. Growth (Netherlands) vol.74 (1986) p292] M. Konagai, T. Yamada, T. Akatsuka, K. Saito, E. Tokumitsu, K. Takahashi [ J. Cryst. Growth (Netherlands) vol.105 (1990) p.359 ] CR. Abernathy, SJ. Pearton, R. Caruso, F. Ren, J. Kovalchick [Appl. Phys. Lett. (USA) vol.55 (1989) p. 1750] CR Abernathy, PW. Wisk, AC. Jones, SA. Rushworth [ Appl. Phys. Lett. (USA) vol.61 (1992) p.180] PA. Lane et al [Appl. Phys. Lett. (USA) vol.61 (1992) p.285 ] CR. Abernathy [ J. Vac. Sd., Technol. A (USA) vol. 11 (1993) p.869 ] A. Robertson Jr., T. H. Chiu, W.T. Tsang, J.E. Cunningham [ J. Appl. Phys. (USA) vol.64 (1988) p.877] N. Kobayashi, J.L. Benchimol, F. Alexandra, Y. Gao [ Appl. Phys. Lett. (USA) vol.51 (1987) p. 1907] T. Martin, C R Whitehouse [ J. Cryst. Growth (Netherlands) vol. 105 (1990) p.57 ] M.B. Panish [ J. Electrochem. Soc. (USA) vol. 127 (1980) p.2730 ] M.B. Panish, H. Temkin, S. Sumski [J. Vac. Sci. Technol. B (USA) vol.3 (1985) p.657 ] A.R. Calawa [Appl. Phys. Lett. (USA) vol.38 (1981) p.701 ] M.B. Panish, S. Sumski [ J. Appl. Phys. (USA) vol.55 (1984) p.3571 ] D. Huet, M. Lambert, D. Bonnevie, D. Defresne [ J. Vac. Sci. Technol. B (USA) vol.3 (1985) p.823 ] J. Musolf, M. Weyers, P. Balk, M. Zimmer, H. Hofinann [ J. Cryst. Growth (Netherlands) vol. 105 (1990)p.271 ] D. Ritter, M.B. Panish, RA. Hamm, D. Gershoni, I. Brener [Appl. Phys. Lett. (USA) vol.56 (1990) p. 1448] CR. Abernathy, P.W. Wisk, DA. Bohling, G.T. Muhr [ Appl. Phys. Lett. (USA) vol.60 (1992) p.2421 ] CR. Abernathy, P.W. Wisk, S.J. Pearton, F. Ren, D. A. Bohling, GT. Muhr [J. Cryst. Growth (Netherlands) vol.124 (1992) p.64 ] S. Salim, L.P. Lu, K.F. Jensen, DA. Bohling [ J. Cryst. Growth (Netherlands) vol. 124 (1992) p .16] PJ. Skevington, D.A. Andrews, GJ. Davies [ J. Cryst. Growth (Netherlands) vol.105 (1990) p.371 ] M. Weyers, J. Musolf D. Marx, A. Kohl P. Balk [ J Cryst. Growth (Netherlands) vol. 105 (1990) p.383] A. Sandhu, T. Fujii, H. Ando, T. Takahashi, H. Ishikawa, N. Yokoyama [ J. Cryst. Growth (Netherlands) vol. 111 (1991) p.559 ]

[26] [27] [28] [29] [30] [31] [32] [33] [34] [35] [36] [37] [38] [39] [40] [41] [42] [43] [44] [45] [46] [47] [48] [49]

H. Ando et al [J. Cryst. Growth (Netherlands) vol. 120 (1992) p.228 ] P.W. Wisk, CR. Abemathy, SJ. Pearton, F. Ren, T. Fullowan, J. Lothian [ Mater. Res. Soc. Symp. Proc. (USA) vol.240 (1992) p.63 ] CR. Abemathy, SJ. Pearton, N.T. Ha [ J. Cryst. Growth (Netherlands) vol. 108 (1991) p.827 ] J. Musolf, D. Marx, A. Kohl, M. Weyers, P. Balk [ J. Cryst. Growth (Netherlands) vol. 107 (1991) p. 1043] CR. Abemathy, SJ. Pearton, F. Ren, J. Song [J. Cryst. Growth (Netherlands) vol.113 (1991) P-412] TJ. deLyon et al [J. Cryst. Growth (Netherlands) vol. 111 (1991) p.564 ] T.P. Chin, P.D. Kirchner, J.M. Woodall, CW. Tu [ Appl. Phys. Lett. (USA) vol.59 (1991) p.2865 ] CR. Abemathy, SJ. Pearton, F. Ren, W.S. Hobson, P.W. Wisk [ J Vac. Sci. Technol. A (USA) vol.12 (1994) p. 1186-90] C. Palmstrom, B.P. Van der Gaag, J.I. Song, S.A. Schwarz [J. Cryst. Growth (Netherlands) vol.136 no. 1-4 (1994)] TJ. deLyon, N.I. Buchan, P.D. Kirchner, J.M. Woodall, GJ. Scilla, F. Cardone [ Appl. Phys. Lett. (USA) vol.58 (1991) p.517] Y.M. Houng, S.D. Lester, D.E. Mars, J.N. Miller [ J. Vac. Sci. Technol. B (USA) vol. 11 (1993) P-915] CW. Tu, B.W.Liang, T.P. Chin [ J. Cryst. Growth (Netherlands) vol. 136 (1994) p. 191 ] K. Saito et al [J. Appl. Phys. (USA) vol.64 (1988) p.3975 ] CR. Abemathy et al [ J. Cryst. Growth (Netherlands) vol. 105 (1990) p.375 ] D.A. Bohling, CR. Abemathy, K. Jensen [ J. Cryst. Growth (Netherlands) vol. 136 (1994) p. 118 ] M. Shirahama, K. Nagao, E. Tokumitsu, M. Konagai, K. Takahashi [ Inst. Phys. Con/. Ser. (UK) no.l29ch.8(1993)p.669] CR. Abemathy, D.A. Bohling [ unpublished ] J. Nishizawa, H. Abe, T. Kurabayashi [ J. Electrochem. Soc. (USA) vol. 132 (1985) p. 1197 ] M.A. Tischler, S.M. Bedair [Appl. Phys. Lett. (USA) vol.48 (1985) p.1681 ] CH.L. Goodman, M.V. Pessa [ J. Appl. Phys. (USA) vol.60 (1986) p.R65, references therein ] M. Ozeki, N. Ohtusuka, Y. Sakuma, K. Kodama [ J. Cryst. Growth (Netherlands) vol. 107 (1991) p. 103] J. Wisser et al [ J. Cryst. Growth (Netherlands) vol. 107 (1991) p. 111 ] J.R. Gong et al [J. Cryst. Growth (Netherlands) vol. 107 (1991) p.83 ] B.C. Chung, RT. Green, H.F. MacMillan [J. Crystal Growth (Netherlands) vol. 107 (1991) p.89 ]

16.9 Epitaxial GaAs lift-off E. Yablonovitch March 1990 (updated May 1996)

A

INTRODUCTION

Epitaxial lift-off (ELO) permits the integration of III-V films and devices onto arbitrary material substrates. In this respect, it competes with the lattice mismatched growth of GaAs directly onto silicon which will be reviewed elsewhere in this volume. The main advantages of ELO are: (1) the growth is lattice-matched, producing III-V semiconductor material of uncompromised quality and (2) the substrate choice is flexible. Historically, there have been various procedures for separating epitaxial films from their growth substrates. As an example of one such successful technology, GaAs photocathodes [1] are made by fusing glass to an epi-wafer and then etching away the entire substrate. Another approach has been to regard the substrate wafer as a kind of re-usable epitaxial growth template. For example Fan [2] has promoted a process in which an epitaxial film is cleaved away from the substrate on which it is grown. Likewise, there have been a number of attempts in the past [3,4] to use selective etching to undercut an epitaxially grown film. ELO falls into this category of thin film separation methods. Recently the problems of etch rate selectivity, bubble formation and thin film handling have been largely overcome [5], permitting rapid progress to be made. B

SELECTIVE ETCHING

ELO relies on the amazingly high etch rate selectivity OfAlxGa1^xAs alloys in hydrofluoric acid as the composition changes from x = 40% to x = 100%. Measured selectivities are larger than 100 million to 1. If an ultra-thin AlAs layer ( 2 nm to 50 nm thick) is the first layer to be grown in a multilayer epitaxial sequence, then a large-area film ( up to 2 cm x 4 cm ) can be undercut from its growth substrate. The thin film comprising the electronic device is released without the introduction of any foreign layers or substances. The GaAs substrate is left intact and can be reused if so desired, while the epitaxial thin film can be cemented or 'Van der Waals bonded' by surface tension forces to any arbitrary substrate. Etch rates are only weakly dependent upon the acid concentration. The lateral undercutting rate of pure AlAs epilayers in a narrow slot progresses at a velocity of 5 microns per minute or 0.3 mm per hour, provided there is no bubble nucleation of the gaseous etching reaction products. This speed appears to be temperature-independent. It is somewhat slower than the fastest etch rates given in TABLE 1, which were measured at a free surface. C

DEVICE EXAMPLES

When using ELO for the production of specific electronic devices the question of 'preprocessing' versus 'post-processing' arises. In 'preprocessing' the devices are fully processed on the original growth substrates and then lifted off as a complete thin film device. In this case ELO is seen as an alternative form of wafer dicing, but allowing thin film integration onto any substrate.

TABLE 1. The etch rate of various AlxGa1^As alloys in microns per minute, immersed in hydrofluoric acid of 49% aqueous concentration. The etching is thermally activated at low aluminium mole-fraction, but appears to be diffusion limited at high aluminium mole-fraction. These etch rates should be used only as a general guide since there are significant sample to sample variations. Aluminium Mole

230K

Acid temperature

Activation energy (eV)

250K

273 K

296 K

328 K

1.8 x 10'6

1.2 x 10-5

5.6 x 10'5

3.8 x 10"4

0.5

2.5

0.48

fraction 40% 50%

1.6 x 10"3

1.2 x 10"2

7.4 x 10*2

3.8 x 10'1

64%

1.0

3.2

9.8

15

0.29

80%

6.5

11

18

28

0.13

In 'post-processing', by contrast, an epitaxial film is bonded to a new substrate and only then processed into devices. This has the advantage of permitting easy alignment to existing substrate patterns using conventional lithographic tools. Depending on the specific device application, the choice of processing sequence generally falls somewhere between these two extremes. At this point a large number of devices have been demonstrated using some combination of preand post-processing: 1.

Double heterostructure GaAs/AlGaAs thin film diode lasers [6] on glass.

2.

GaAs metal-semiconductor field effect transistors (MESFETS) [7] on glass and silicon.

3.

Strained single quantum well InGaAs/GaAs high electron mobility transistors [8] (HEMTs).

4.

Regrowth of GaAs quantum wells has been demonstrated [9] on GaAs lift-off films, 'Van der Waals bonded' to silicon substrates. Regrowth after ELO is an extreme example of 'post-processing'.

5.

A high speed InP/InGaAs photodiode on a sapphire substrate [10]. This is noteworthy because it is the first application of ELO to InP growth substrates. It required the use of an ultra-thin (2 nm) pseudomorphic strained AlAs release layer.

6.

Waveguide coupling [11,12] into GaAs thin-film photodetectors lying on LiNbO3 and glass waveguide substrates.

7.

GaAs light emitting diodes on metallized silicon substrates [13]. In these post-processed devices, an alloyed electrical contact is made between the thin film GaAs structure and the substrate.

8.

GaAs MESFETs have been integrated [14] with InP waveguide structures.

9.

Double heterostructure excitonic absorption test structures [15] on glass substrates.

D

FILM CURVATURE AND GAS OUT-DIFFUSION

ELO relies upon the escape of the gaseous reaction products of etching. An important role is played by the wax providing mechanical support to the epitaxial film as it lifts off the substrate. Due to the difference in thermal contraction between the wax layer and the GaAs film upon cooling down from the wax annealing temperature, the film becomes curved [5]. This curvature opens up the escape channel for gaseous etching products. The maximum permissible [5] undercutting speed v to allow out-diffusion while preventing bubble nucleation is: v =

Dn —== m N TT sl^TTl

(i\ v J

where n is the saturation solubility of the etching reaction products (which may include H2 or possibly AsH3), N is the molar concentration of AlAs, D is the gaseous diffusion constant, m is the number of moles of reaction product per mole of AlAs, t is the thickness of the AlAs epilayer and R is the radius of curvature induced by thermal contraction. For H2 gas in aqueous media, n/N is approximately 2 x 10'5 [16], and D is approximately 5 x 10"5 cm"2 /s [17]. The fractional difference in thermal contraction, K, between the wax and GaAs produces the following radius of curvature: T3 Y R . _ ^ L _ 6K Tg2 Yg

(2)

where Tw and Tg are the thickness of the wax and GaAs films respectively, and Yw and Yg are the Young's moduli of the wax and GaAs films, respectively. EQN (2) is valid when T^T 8 » 1 and Yg/Yw » 1 . Typically TJTg is about 100, Y/Y w is about 100 and K is about 0.01 resulting in approximately a 10 cm radius of curvature. This curvature limits the permissible speed of undercutting in EQN (1). E

CONCLUSION

The ability to separate the growth substrate from the active thin film permits a considerable increase in the sophistication of materials engineering. In particular, the properties of the supporting substrate can be optimized separately from those of the epitaxially grown film. There are many attributes in which novel supporting materials would have properties superior to III-V growth substrates, among them being: substrate dielectric constant for increased speed in supercomputers; thermal conductivity for high power applications; cost for solar cell applications; weight for space applications; radiation hardness for military electronics; mechanical strength and mechanical flexibility. Most importantly ELO opens up the prospect for opto-electronic integration (i.e. GaAs films on silicon) and electro-optic integration (i.e. GaAs films on LiNbO3 and glass waveguides). The 'Van der Waals' bonding process [18] employing surface tension forces plays a role in many of these applications. ELO can be expected to impact many other research and development projects that depend on

the ability to marry thin layers of dissimilar materials. The most important applications are probably yet to be discovered.

REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] [II] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29]

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[30] [31] [32] [33] [34] [35] [36] [37] [38] [39] [40] [41] [42] [43] [44] [45] [46] [47] [48] [49] [50] [51] [52] [53] [54] [55] [56] [57] [58] [59]

[60] [61] [62] [63]

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CHAPTER 17 LOW TEMPERATURE MBE GaAs 17.1 17.2 17.3 17.4

Low temperature GaAs: growth dynamics and effects of As:Ga flux ratio Point defects in LT GaAs Uses of low temperature grown materials in MESFETs Opto-switches using low-temperature MBE GaAs

17.1 Low temperature GaAs : growth dynamics and effects of As:Ga flux ratio M. Missous April 1996

A

INTRODUCTION

The MBE growth of GaAs at low temperatures (< 2500C) has not been as well documented as that at normal temperature despite the importance of LTMBE GaAs in ever increasing applications [1-4]. In this particular system ex-situ studies (i.e. after the GaAs is removed from the MBE system) have preceded in situ studies (i.e. within the MBE environment using techniques such as reflection high energy electron diffraction (RHEED) etc.). The major advantage of GaAs growth at high temperatures is the generously wide permissible arsenic overpressure that is tolerated to achieve growth of stoichiometric GaAs. This unique feature whereby an overpressure of As is used, the 'normal growth conditions' (NGC), has been the reason for the great success of MBE in growing high quality GaAs and AlGaAs at higher temperatures. In parallel, many studies of the growth mechanism and in particular the observation of (RHEED) oscillations on (100) surfaces [5] suggesting 2 dimensional (2D), layer by layer, growth were made by a number of workers. The work on RHEED oscillations has generally been performed in the growth temperature range of -500 - 7000C [5] under NGC where sustained and strong oscillations are usually observed. Below 5000C RHEED oscillations are not observed under NGC . The growth of LT GaAs has traditionally been accomplished under conditions which are more appropriate for high temperature growth; however as the growth temperature is lowered below 3000C, NGC leads to heavily defective and strained materials with concentrations of point defects in the 1019 to 1020 cm"3 range. The exact role of the arsenic overpressure on various properties of LT-GaAs has been studied by several groups [6,7]. B

LOW TEMPERATURE MBE GROWTH

The MBE growth of LT GaAs is usually performed under an overpressure of arsenic and unlike high temperature growth no surface reconstruction or RHEED oscillations are normally observed. However when detailed investigation of the growth dynamics was undertaken by Mirin et al [8] and Missous [9], it became clear that the role of excess arsenic in determining the growth process is paramount. In particular when growing under NGC and at temperatures < 490 0 C, in the absence of an arsenic overpressure and at a background pressure of 1 * 10"7 torr, the RHEED oscillations proceeded in almost the same manner with a strong (2><4) reconstruction present at the growth temperature. In the absence of an As flux and on further lowering the temperature to 4300C, the reconstruction changed to C(4><4) and very weak and noisy RHEED oscillations are usually recorded. It is clear that at temperatures lower than ~ 450 0 C there is strong adsorption of the As species which then prohibits the 2D layer by layer growth.

Specular spot intensity (A.U.)

Time (sec) FIGURE 1. GaAs RHEED oscillations at 200 0 C and the effect of non-stoichiometiy.

However under conditions where the arsenic flux is just enough to lead to growth (stoichiometric conditions), RHEED oscillations, as shown in FIGURE 1, are sustained. This clearly indicates 2D, layer by layer, growth at these extremely low temperatures, a fact first noted by Ibbetson et al [10] and confirmed by Missous [9]. It must be noted though, despite the occurrence of strong RHEED oscillations at these extremely low temperatures, that there was still no change in the surface reconstruction which is usually a (1 x 1) surface irrespective of arsenic overpressure. The evolution of RHEED oscillations and their existence at such low temperatures where reconstruction is lost strongly points to the fact that the oscillations (i.e. the change in the step density of the growing surface) is independent of surface reconstruction. The effect of arsenic overpressure was further investigated by Missous [9] who showed that the growth rate of GaAs decreases with decreasing temperature. This small effect (-5%) was correlated with the degree of deviation from stoichiometry. Such a change in the fundamental frequency has been observed in vicinal surfaces near the transition from layer by layer to step flow growth [11] where a proportion of the atoms do not contribute to the 2D growth. RHEED oscillations are therefore an excellent tool and probe in establishing stoichiometric GaAs at low temperature. As pointed out by Ibbetson et al [10], there is no noticeable recovery in the specular spot intensity at temperatures below 4900C and as a consequence once RHEED oscillations are interrupted they decay much faster on recommencing growth. Indeed the fundamental frequency reduces again by another 5% and the intensity by a further 20% . If the arsenic flux is increased by 20 % there are no oscillations visible at all, again emphasising the importance of keeping the As flux within a tight range of that of the gallium flux. The growth mechanism and the influence of the arsenic flux as the growth temperature is lowered

Specular spot intensity (A.U.)

Time (sec) FIGURE 2. AlGaAs RHEED oscillations at 200 0 C and the effect of non-stoichiometry.

is by no means confined to GaAs. It is known that the surface adatom mobility of Al is greatly suppressed as the growth temperature is lowered, a fact that has led to most AlGaAs structures being grown at temperatures greater than 6400C. To investigate whether the Al mobility was still sufficient to lead to 2D growth of AlGaAs at low temperature, Missous repeated the GaAs experiments on AlGaAs composition of -40%, a value that encompasses the very important alloy

Normalized X-ray counts

BEP BEP BEP BEP

= = = =

Angle, arc seconds FIGURE 3. The effect of As to Ga BEP ratio on excess arsenic incorporation in LT-GaAs.

4 5 6 3

Normalised X-Ray Intensity (arb.units)

BEP-8 Undoped

Tg=250°C Undoped

Angle (arc sec) Effect of substrate temperature on excess arsenic incorporation.

FIGURE 4. The effect of reduced growth temperature on the incorporation of the excess arsenic.

composition used for most structures and devices. As for GaAs, the growth of AlGaAs can proceed via a 2D mode at temperatures as low as 200 0 C (FIGURE 2), the lowest temperature ever reported for the 2D growth of AlGaAs. Also, as for GaAs, a slight change in arsenic dependence (-20%) leads to a dramatic reduction in the RHEED oscillation frequency and intensity [9]. C

STRUCTURAL PROPERTIES

It is well known that excess arsenic incorporation in LT GaAs gives rise to extra diffraction peaks in X-ray diffraction measurements [12]. FIGURE 3 shows the double crystal X-ray diffraction (DXRD) rocking curves of various layers grown by the present author. It is quite clear that the effect of beam equivalent pressure (BEP) ratio of arsenic with respect to gallium is a very small one until one gets very close to stoichiometric conditions. In these samples, the measured concentrations of the neutral arsenic antisites [AsGa ]° using infrared absorption (see TABLE 1) are again seen to vary little with BEP and are equal to « 2 x 1019 cm"3, comparable to what is usually reported in the literature for growth at 250 0 C [13]. However, the signal from the stoichiometric layer is below the detection limit.

TABLE 1. Excess arsenic incorporated as a function of As to Ga beam equivalent pressure ratio measured using DCXR and Optical absorption.

Excess As (X-ray) (cm 3 )

[As 0 J 0 (cm 3 )

_6

1.88 x IQ20

1.8 x IQ19

_5

1.57 x 1Q;O

2.1 x IQ19

_4

1.62 x IQ20

1.5XiQ 19

BEP ratio

3

below detection limit

below detection limit

The effect of reduced growth temperature on the incorporation of the excess arsenic is shown in FIGURE 4. The inference from the data in FIGURES 3 and 4 is that at a given temperature, there is a solid solubility limit to the incorporation of excess arsenic with lower temperatures leading to higher excess arsenic incorporation. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] [II] [12] [13]

D.C. Look [ 77?/« Solid Films (Switzerland) vol. 231 (1993) p.61 ] M.R. Melloch, N. Otsuka, J.M. Woodall, A.C. Warren, J.L. Freeouf [ Appl. Phys. Lett. (USA) vol.57 no. 15 (1990) p. 1631] F.W. Smith, A.R. Calawa, CL. Chen, MJ. Manfra, LJ. Mahoney [ IEEE Electron Device Lett. (USA) vol.9 (1988) p.77 ] S. Gupta et al [Appl .Phys. Lett. (USA) vol.59 (1991) p.3276 ] J.H. Neave, B.A. Joyce, PJ. Dobson, N. Norton [ Appl. Phys. A (Germany) vol.31 (1983) p. 1 ] Z. Liliental-Weber, A. Claverie, P. Werner, W. Schaff, E.R. Weber [ Mater. Sci. Forum (Switzerland) vol.83-87 (1992) p. 1045 ] M. Missous, S.P. O'Hagan [ J. Appl. Phys. (USA) vol.75 (1994) p.3396 ] R.P. Mirin, J.P. Ibbetson, U.K. Mishra, A.C. Gossard [ Appl. Phys. Lett. (USA) vol.65 (1994) p.2335] M. Missous [ J. ApplPhys. (USA) vol.8 no.7 (1995) p.4471 ] J.P. Ibbetson, RP. Mirin, U.K. Mishra, A.C. Gossard [J. Vac. Sci. Technol. B (USA) vol. 12 (1994) p. 1050] B.A. Joyce, T. Shitara, A. Yoshinaga, D.D. Vvedensky, J.H. Neave, J. Zhang [ App. Surf. Sci. (Netherlands) vol.60-61(1992) p.200 ] M. Kaminska, Z. Liliental-Weber, E.R. Weber, T. George [ Appl. Phys. Lett. (USA) vol.54 (1989)p.l881] M.O. Manasreh, D.C. Look, K.R. Evans, CE. Stutz [ Phys. Rev.B (USA) vol.41 (1990) p. 10272 ]

17.2 Point defects in LT GaAs D.C. Look December 1995

A

INTRODUCTION

The electrical and optical properties of low-temperature-grown (LT), molecular-beam-epitaxial (MBE) GaAs are largely determined by point defects, especially the As antisite, AsGa. Most of the standard characterization techniques have been applied to LT GaAs, and many defect properties have been elucidated. In this Datareview we will be mainly concerned with the electrical, optical, and magnetic-resonance characterization of as-grown and annealed LT GaAs, all of which can be explained qualitatively and quantitatively by point defects. There is also a body of evidence that As precipitates can influence the properties of annealed (but not as-grown) LT GaAs. However, quantitative verification of this hypothesis is difficult, so we will not attempt to explore it in detail, but will refer the reader to available references. In fact, a good overall view of LT GaAs materials and devices can be found in references [1-4], which are conference proceedings or reviews, and reference [5], which is a recent work containing many references on native defects. B

GENERAL STOICHIOMETRY AND ANNEALING EFFECTS

Bl

Stoichiometry

At typical substrate temperatures of 180 - 300 0 C, and As4/Ga beam-equivalent pressure ratios (BEPs) of 10 - 40, the MBE GaAs layers are very As rich, with up to 1.5 at % of excess As [6]. Thus, native defects which might be expected include the As antisite (AsGa), the As interstitial (As1), and the Ga vacancy (VGa). The AsGa centre is known to be a double donor ((0/+) level at E c - 0.75 eV and (+/++) level at E v + 0.5 eV), and theoretical studies predict that VGa could be a triple acceptor [7], and As1 either a donor or acceptor, depending upon configuration [8]. Thus, coulombic forces could lead to all combinations of AsGa, As1, and VGa complexes, and indeed, all have been invoked in the literature. Also, a single nearest-neighbour hop can transform VGa to A5Ga" VAS [9], and this defect is also thought to exist [5]. However, of all these possible defects, onlyAsGa has been clearly and unequivocally identified to nearly everyone's satisfaction [10]. Even in that case, there is disagreement on whether the AsGa is isolated or is part of a complex. Besides isolated defects and defect-defect complexes, defect-impurity complexes have also been postulated to exist. In particular, local vibrational mode (LVM) spectroscopy has been used to identify the SiGa - VGa complex in Si-doped LT GaAs, and the AsGa - H - V^ complex in H-doped material, with the latter less definitive [5]. At lower As4/Ga BEPs, the defects mentioned above are less prevalent; in fact, at a growth temperature T = 2500C, and a BEP = 3, the layers are nearly stoichiometric [11,12]. For higher growth temperatures, TG > 4000C, the layers again are nearly stoichiometric, even at a relatively high BEP, say 20. However, in the latter case, the As011 concentration is still large enough (~ 1017 cm"3) to produce semi-insulating material [13].

B2

Annealing

Annealing LT MBE GaAs at a temperature, TA, higher than the growth temperature TG does not, in general, change the stoichiometry (i.e., amount of excess As), but does cause defect modifications. For a sample grown at 200 - 250 0 C, the optical and electrical properties will begin to be affected at annealing temperatures as low as 3500C [14-16]. The effects are not necessarily monotonic as TA is increased, suggesting complex defect reactions. As an example, for a sample having TG = 200 0 C and BEP = 20, the acceptor concentration dropped about one order of magnitude at TA = 400 0 C, and then increased back to roughly the original value at TA = 450°C[14]. For another sample, grown at 3000C, the conductivity actually changed from n-type to p-type at 450 0 C, and then back to n-type at 5000C. For TA > 5000C, most of the excess As atoms move from their AsGa and As1 positions to form large (10 - 300 A) As precipitates [17-19]. However, high enough concentrations (1017-1018 cm"3) of AsGa still exist to pin the Fermi level at midgap [14,20]. There also is evidence that As precipitates can pin the Fermi level [21], and, suffice it to say, a longstanding controversy exists on this matter [20-24]. For further information on the precipitates it is suggested that the reader consult one of the reviews, references [1-4], or a previous Datareview 10.9 of the present volume. C

DEFECT PROPERTIES

Cl

As Antisite

The As antisite As^ is unquestionably the dominant defect in MBE GaAs grown at 180 - 400 0 C, at least if the material is not annealed [2,10,20]. That is, virtually none of the major electrical and optical properties can be explained without invoking AsGa . The same may also be true for annealed material, although there is a controversy on this point, as mentioned earlier [20-24]. Although it has been known since 1978 that low growth temperatures (TG < ~ 400 0 C) will produce high-resistivity behaviour [25], and high concentrations of deep traps in MBE GaAs [26], it was not until 1989 that a positive identification of AsGa was made, by observing the well-known four-line electron paramagnetic resonance (EPR) spectrum [10]. Research on this defect in LT GaAs has greatly benefited by the existence of a wide body of literature on the EL2 defect, found in undoped, semi-insulating GaAs and known to have AsGa as one (or the only) component [27]. Some of the As^-related defects in LT GaAs appear to be identical to EL2, and some not [13,16,28]. These reports are not contradictory because, as mentioned earlier, various defect-defect complexes, such as AsGa - As1, AsGa - VGa and AsGa - VM, are expected to co-exist along with the isolated As03. Their relative concentrations can depend upon growth and annealing conditions. As an example, the dominant donor in material grown at TG = 350 450 0 C, known (from magnetic-resonance experiments [29]) to contain AsGa, has a Hall-effect Arrhenius energy of 0.65 ± 0.01 eV (the (0/+) transition), whereas the value for EL2 is 0.75 ± 0.01 eV [13]. On the other hand, the resistivity Arrhenius energy in n-i-n structures grown at 250 0 C and annealed at 6000C is about 0.77 eV, giving an energy level of 0.74 ± 0.03 eV after correcting for the expected mobility temperature dependence [28]. This latter value is quite close to that of EL2. A theoretical study [30] has predicted that an AsGa - As1 defect, in which the As1 joins a neighbouring As to form a split interstitial, has an energy of about 0.1 eV closer to the conduction

band than the isolated AsGa. Thus, the 0.65-eV defect may be AsGa- As1, although the known metastability of the 0.65-eV defect [13] is not properly predicted [30]. The near-IR (0.7 - 1.5 eV) absorption spectrum of LT GaAs looks much like that of EL2 and it seems reasonable to use the well-known photoionization cross sections of the latter [31] to determine concentrations of the former. By this procedure it has been determined by many workers that material grown at 2000C has [AsGa] - 1019 - 1020 a n 3 [32-34]. For higher growth temperatures, say 350 - 4500C, the As04 concentrations are smaller and the IR absorption is very weak. However, magnetic circular dichroism absorption (MCDA) measurements establish a strong presence of AsGa [29], and Hall-effect measurements on LT/p/LT structures show a 0.65-eV donor concentration of ~ 1017 cm"3 [35]. Thus, [AsGa] > ~ 1017 cm"3, even for growth at 4000C. Annealing at 6000C greatly reduces [As0J in material grown at 200 0 C, due to precipitate formation [17,36], although [AsGa] remains > - 1017 cm"3 [20]. For TG > - 400 0 C, annealing at 600 0 C affects the electrical properties, and thus [AsGa], very little [16]. Cl

As Interstitial

Early investigations concluded that a large fraction of the excess As was accommodated in the form of As interstitials, with [As1] - 1020 - 1021 cm"2 for material grown at 190 - 200 0 C with a BEP - 10 - 20 [37]. The reasoning here was that with about 1 - 2% excess As, and with [AsGa] < ~1020 cm"3, and [VGa] < - 1019 cm"3, there must be > 1020 cm"3 OfAs1, or perhaps small As clusters. Very recently, this hypothesis has been questioned [38], and it has been suggested that only As04 and V ^ are present in significant quantities, with all of the lattice expansion being due to the AsGa (reduced of course by the VGa). Indirect support for this point of view comes by considering the high energy of formation expected for As1, although it must be noted that MBE is not an equilibrium process in a thermodynamic sense. Another recent study [5], however, concludes that the As1 indeed do exist, and that generally [As1] > [AsGa]. Part of the problem in establishing the presence or absence OfAs1 is the lack of electrical or optical fingerprints. For example, theoretical studies find that As1 can exist in either donor or acceptor states [8], depending on configuration, and there is no known optical transition attributable to As1. Apparently, ion channelling experiments, which originally were interpreted in favour OfAs1 [37], could also be explained by the large AsGa - As^ bond [38]. C3

Ga Vacancy

Most investigators have concluded that the Ga vacancy, VGa, exists in LT GaAs and is the dominant acceptor [39]. Hard evidence, however, is lacking because there are no incontrovertible electrical, optical, or magnetic-resonance fingerprints. Positron annihilation (PA) experiments are sensitive to negatively-charged vacancies, and, indeed, such vacancies are often seen in LT GaAs [5,40]. However, it is not absolutely clear that these are VGa centres in all cases since even V^ can evidently exist in a negative charge state in n-type material [41]. The likelihood, though, is that either VGa or VGa complexes are responsible for the PA data. Further support for V03 comes from LVM experiments in Si-doped LT GaAs. Here some of the observed absorption lines are best interpreted as due to SiGa - VGa centres [5], which might be expected from autocompensation and coulombic-attraction considerations. Such considerations would suggest that V ^ would not be prevalent in Be-doped material and, indeed, no PA signature

is seen in this case [5]. Further support for the existence of V^ comes from diffusion experiments [39]. IfV08 is the dominant acceptor in LT GaAs, then the concentrations must range up to 1019 cm'3 in 200 0 C material. This conclusion comes from two-wavelength absorption experiments [33], Interestingly, the ratio NA /ND seems to hold at the approximate value 0.1 over a wide range of N0, where ND is known to be AsGa- related [35]. If VGa remains the dominant acceptor over this range, then its concentration must track that of AsGa. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] II1] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32]

G.L. Witt, R. Calawa, U. Mishra, E. Weber (Eds) [ Mater. Res. Soc. Symp.Proc. (USA) vol.241 (1992)] D.C. Look [ 77?/« Solid Films (Switzerland) vol.231 (1993) p.61 ] D.C. Look, M.R. Melloch (Eds) [ Special Issue, J. Electron. Mater. (USA) vol.22 (1993) p. 1373 ] DJ. Eaglesham [ J. Appl. Phys. (USA) vol.77 (1995) p.3597 ] R.E. Pritchard et al [ J. Appl. Phys. (USA) vol.78 (1995) p.2411 ] Z. Liliental-Weber [Mater. Res. Soc. Symp. Proc. (USA) vol.241 (1992) p. 101 ] MJ. Puska [ J. Phys., Condens. Matter (UK) vol. 1 (1989) p.7347 ] DJ. Chadi [ Phys. Rev. B (USA) vol.46 (1992) p.9400 ] G.A. Baraff, M. Schluter [ Phys. Rev. B (USA) vol.33 (1986) p.7346 ] M. Kaminska et al [Appl. Phys. Lett. (USA) vol.54 (1989) p.1881 ] M. Missous, S. O'Hagan [Mater. Res. Soc. Symp.Proc. (USA) vol.325 (1994) p.371 ] S. O'Hagan, M. Missous [ J. Appl. Phys. (USA) vol.75 (1994) p.7835 ] D.C. Look, Z-Q. Fang, J.R. Sizelove, CE. Stutz [ Phys. Rev. Lett. (USA) vol.70 (1993) p.465 ] D.C. Look, D.C. Walters, G.D. Robinson, JR. Sizelove, M.G. Mier, CE. Stutz [ J. Appl. Phys. (USA) vol.74 (1993) p.306 ] J. Darmo, F. Dubecky, P.Kordos, A. Forster, H. Luth [Mater. Sci. Eng. B (Switzerland) vol.28 (1994) p.393 ] P.Kordos, A. Forster, J. Betko, M. Morvic, J. Novak [ Appl. Phys. Lett. (USA) vol.67 (1995) p.983] M.R Melloch,N. Otsuka, J.M. Woodall, AC. Warren, JL. Freeouf [Appl. Phys. Lett. (USA) vol.57 (1990) p. 1531] Z. Liliental-Weber, A. Claverie, J. Washburn, F. Smith, A.R. Calawa [ Appl. Phys. A (Germany) vol.53 (1991) p. 141] TM. Cheng,CY. Chang, J.M. Huang [J. Appl. Phys. (USA) vol.76 (1994)p.5697] X. Liu et al [ Appl. Phys. Lett. (USA) vol.65 (1994) p.3002 ] A.C. Warren et al [Appl. Phys. Lett. (USA) vol.57 (1990) p.1331 ] D.C. Look [ J. Appl. Phys. (USA) vol.70 (1991) p. 141 ] M.R. Melloch, J.M. Woodall [ Appl. Phys. Lett. (USA) vol.67 (1995) p. 1331 ] X. Liu et al [ Appl. Phys. Lett. (USA) vol.67 (1995) p. 1333 ] T. Murotani, T. Shimanoe, S. Mitsui [J. Cryst. Growth (Netherlands) vol.45 (1978) p.302 ] RA. Stall,CEC. Wood,PD. Kirchner, LF. Eastman [Electron. Lett. (UK) vol. 16 (1980)p. 171 ] M.O. Manasreh, D.W. Fischer, W.C. Mitchel [Phys. Status Solidi B (Germany) vol. 154 (1989) p.ll] H. Fujioka, E.R. Weber, A.K. Verma [ Appl. Phys. Lett. (USA) vol.66 (1995) p.2834 ] K. Krambrock, M. Linde, J.M. Spaeth, D.C. Look, D. Bliss, W. Walukiewicz [ Semicond. Sci. Tech. (UK) vol.7 (1992) p. 1037] J.I. Landman, CG. Morgan, J.T. Schick [ Phys. Rev. Lett. (USA) vol.74 (1995) p.4007 ] P.Silverberg, P.Omling, L. Samuelson [ Appl. Phys. Lett. (USA) vol.52 (1988) p. 1689 ] M.O. Manasreh,D.C. Look, K R Evans, CE. Stutz [Phys. Rev. B (USA) vol.41 (1990)p. 10272 ]

[33] [34] [35] [36] [37] [38] [39] [40] [41]

D.C. Look, D.C. Walters, MMier, CE. Stutz, S.K. Brierley [ Appl. Phys. Lett. (USA) vol.60 (1992) p.2900 ] N. Hozhabri, S-H. Lee, K. Alavi [Appl. Phys. Lett. (USA) vol.66 (1995) p.2546 ] D.C. Look, G.D. Robinson, J.R. Sizelove, CE. Stutz [ J. Electron. Mater. (USA) vol.22 (1993) p. 1425] T.W. Fan, J.B. Liang, HJ. Deng, R.G. Li, Z.G. Wang, W.Gen [ J. Cryst. Growth (Netherlands) vol.143 (1994) p.354] K.-M. Yu, M. Kaminska, Z. Liliental-Weber [ J. Appl. Phys. (USA) vol.72 (1992) p.2850 ] X. Uu, A. Prasad, J. Nishio, ER. Weber, Z. Liliental-Weber, W. Walukiewicz [ Appl. Phys. Lett. (USA) vol.67 (1995) p.279 ] D.E. Bliss, W. Walukiewicz, J.W. Ager III, E.E. Haller, K.T. Chan, S. Tanigawa [ J. Appl. Phys. (USA) vol.71 (1992) p. 1699 ] DJ. Keeble, M.T. Umlor, P.Asoka-Kumar, K.G. Lynn, P.W.Cooke [ Appl. Phys. Lett. (USA) vol.63 (1993) p.87] K. Saarinen, P.Hautojarvi, P.Lanki, C. Corbel [ Phys. Rev. B (USA) vol.44 (1991) p. 10585 ]

17.3 Uses of low temperature grown materials in MESFETs N.X. Nguyen and U.K. Mishra March 1996

A

INTRODUCTION

Device application was the initial driving force in the research and development of low-temperature grown (LTG) GaAs; and ever since, it has always paralleled and complemented the material research of LTG GaAs and related compounds. The interest in device application of LTG material dates back to February 1988 in a paper published by Smith et al [I]. In the paper, Smith reported that backgating and light sensitivity in MESFET integrated circuits could be eliminated by employing a GaAs buffer grown at 2000C and annealed at 600 0 C. This beneficial effect of LTG-GaAs has since generated a flurry of activity in the quest for greater understanding of the material properties, and the search for further device applications of LTG materials. In this Datareview, we will review the progress in the application of LTG materials in MESFETs as a buffer layer and an epitaxial passivation layer. B

BUFFERLAYER

In a MESFET structure, the buffer is a layer that separates the channel from the semi-insulating GaAs substrate. Normally, a buffer is grown to provide a smooth and uniform starting surface for the growth of the device channel. Electrically, an ideal buffer layer would be one that is highly resistive and impervious to external stimuli such as lighting and proximity biases. In practice, this ideal buffer is often difficult to realize, a non-ideal buffer leading to a number of problems in device performance. In particular, backgating (or sidegating), the changing of device characteristics as a result of a bias applied to an adjacent, and nominally isolated device (or electrode), has been a major problem for both digital and analog circuit applications of MESFETs. Over the years, a number of different buffers have been studied and applied to devices - AlGaAs, superlattice structures (GaAs/AlGaAs), etc. but with limited success. As stated above, the advent of the LTG-GaAs buffer brought about a novel and elegant solution to the backgating problem [1,2]. As shown in FIGURE 1, LTG-GaAs buffer virtually eliminates the problem completely. Its usage as a buffer layer for MESFET circuits has gained increased acceptance in the industry and in recent years research has shed further understanding on the effect of the LTG-GaAs buffer. Although the LTG-GaAs buffer greatly enhances the DC characteristics of MESFETs, its effectiveness is reduced at microwave frequencies [3,4]. Low frequency noise of MESFETs with LTG-GaAs buffer has been studied. It was found that, although at zero backgate bias the noise level of LTG-GaAs buffer is higher, with the noise characteristics of a 1/f type noise superimposed on generation-recombination noise [5,6], the sensitivity of noise to backgating bias is significantly reduced. In addition, researchers have also found other beneficial effects that the LTG-GaAs buffer has on MESFET characteristics. Recently, it has been discovered that an LTG-GaAs buffer can significantly reduce the soft error susceptibility in GaAs MESFETs [7]. Another promising application of LTG-GaAs buffer has been found in the utilization of the layer as a 'stress reliever' in highly mismatched channel/substrate interfaces, such as those found in GaAs on silicon [8].

V 0 5 = 2.5 V SPACING = 50 /xm

V D S = 2.5 V SPACING = 50 urn

NEW BUFFER

NEW BUFFER

UNOOPED MBE 580 C GaAs

loss (0)

loss (0)

lDSS (VBG)

UNOOPEO MBE GaAs/AI 0 4 5 G a 0 55AS 7OaC SUPERLATTICE

loss (VBG)

IMPLANT

VPE EPITAXIAL LAYER IMPLANT

DARK LIGHT

DARK LIGHT

BACKGATING VOLTAGE (V)

BACKGATING VOLTAGE (V)

FIGURE 1. Reduction of backgating in MESFET with LTG-GaAs buffer (from [I]).

C

SURFACE PASSIVATION

The surface of GaAs devices has long been recognized as the source of device-to-device non-uniformity and low reliability. Efforts have been devoted to developing a reliable passivation technology for GaAs. Currently, SiNx is the standard passivant in industry. SiNx is an effective environmental passivant (preventing changes of surface properties due to environmental factors such as oxidation and humidity), but it is not an electrical passivant. Chemical treatments to electrically passivate the surface have had limited success. So far, the most successful surface treatments have included photochemical oxidation [9] and applications of sulphur-related compounds, such as Na2S and (NH4)2S [10,11]. DRAIN

SOURCE GATE

n-GaAs channel Buffer-GaAs

FIGURE 2. Gate field redistribution with LTG-GaAs epitaxial passivation; local peak field at point A is lowered.

Recently, researchers have demonstrated that LTG materials passivation can lead to an improved device performance such as higher breakdown voltage and lower 1/f noise. This novel technology is promising in that it enables researchers to engineer the electrical properties of the passivation layer in order to achieve the desired passivation characteristics. Depending upon the material process parameters, a change of resistivity of over five orders of magnitude in the LTG material can be readily achieved [12]. LTG GaAs has been demonstrated to be an effective insulator in a MfSFET structure [13-15]. Subsequently, LTG GaAs and its derivative, LTG-AlGaAs, have been successfully utilized as an epitaxial passivation layer in a MESFET structure [16-19]. It has been hypothesized that traps in the LTG passivation layer act as a charge absorber in the vicinity of the gate thereby alleviating the high field at the drain edge of the gate and increasing the breakdown voltage of the device [19] (FIGURE 2). In addition to gate field alleviation, a reduction in the frequency dispersion of transconductance and output resistance by LTG-GaAs passivation has also been reported [20]. Investigation of low frequency noise of LTG-GaAs passivated devices has revealed that the noise is significantly reduced at low offset frequencies. Ultra low phase noise has been achieved by tailoring the growth temperature of LTG-AlGaAs [21]. The reasons for this behaviour have not yet been ascertained. D

CONCLUSION

Low-temperature grown GaAs and related materials have been successfully applied to GaAs MESFETs. Improvements in device characteristics have been demonstrated with LTG materials as both buffer and passivation layers. The ability to tailor the electrical properties of the LTG materials through their growth and processing makes them very attractive to device engineers. As further understanding is gained into the material properties, the potential for useful device applications of LTG materials is just beginning. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] II1] [12]

F.W. Smith, A.B. Calawa, C-L. Chen, M. J. Manfra, L. J. Mahoney [ IEEE Electron Device Lett.fUSA) vol.9 no.2 (1988) p.77-80 ] E. Aperathids, Z. Hatzopoulos, M. Lagadas, G. Constantinidis [ Gallium Arsenide, Related Compounds 1992. Proceedings of the 19th Int. Symp. (IOP, 1993) p.767-72 ] J.-F.B. Lin, CP. Kocot, D.E. Mars, R. Jaeger [ /JEEB Trans. Electron Devices (USA) vol.37 no. 1 (1990)p.46-50] D.C Streit, M.M. Hoppe, C-H. Chen, J.K. Liu, K.-H. Yen [J Vac. Set Technol. B (USA) vol. 10 no.2 (1992) p. 819-21] A.N. Birbas, B. Brunn, A.D. van Rheenen, A. Gopinath, C-L. Chen, F. Smith [ IEEProc.-G (UK) vol.138 no.2 (1991) p. 175-8] S. Tehrani, A.D. van Rheenen, M.M. Hoogstra, J.A. Curless, M.S. Peffley [ IEEE Trans. Electron Dev. (USA) vol.39 no.5 (1992) p. 1070-4 ] T.R Weatherford, D. McMorrow, A.B. Campbell, W.R. Curhce [ Appl. Phys. Lett. (USA) vol.67 no.5 (1995) p.703-5] RA. Metzger, MJ. Delaney, L. McCray, H. Kanber, D.C Wang, T.Y. Chi [ J. Vac. Sci. Technol.B (USA) vol.8 no.2 (1990) p.297-300 ] S.D. Offsey et al [ Appl. Phys. Lett. (USA) vol.51 (1987) p.33 ] E. Yablonovich, CJ. Sarnoff, R. Bhat, T. Gmitter [Appl. Phys. Lett. (USA) vol.51 (1987) p.439 ] R.S. Besser, CE. Helms [ J. Appl. Phys. (USA) vol.65 (1989) p.4306 ] P. Kordos, A. Forster, J. Betko, M. Morvic, J. Novak [ Appl. Phys. Lett. (USA) vol.67 no.7 (1995) p.983 ]

[13] [14] [15] [16] [17] [18] [ 19] [20] [21]

L.-W. Yin, Y. Hwang, J.H. Lee, R.M. Kolbas, RJ. Trew, U.K. Mishra [ IEEE Electron Device Lett. (USA) vol.11 no.12 (1990) p.561 ] K. Lipka, B. Splingart, E. Kohn [ Electron. Lett. (UK) vol.29 no. 13 (1993) p. 1170-2 ] K. Lipka et al [ Proc. IEEE Cornell Conf. on Advanced Concepts in high SpeedSemicond.Dev. Circuits (I995) p.542 ] C-L. Chen, LJ. Mahoney, LJ. Manfra, F.W. Smith, D.H. Temme, A.R. Calawa [IEEEElectron Dev. Lett. (USA) vol. 13 no. 6 (1992) p.335 ] L.-W. Yin, N.X. Nguyen, Y. Hwang, J.P. Ibbetson, R.M. Kolbas, A.C. Gossard, U.K. Mishra [ J. Electron. Mater. (USA) vol 22 no. 12 (1993) p. 1503 ] L.-W. Yin, X.N. Nguyen, K. Kiziloglu, J.P. Ibbetson, A.C. Gossard, U.K. Mishra [ Electron. Lett. (UK) vol.29 no. 17 (1993) p. 1550 ] N.X. Nguyen, J.P. Ibbetson, W.-N. Jiang,, U.K. Mishra [ Proc. IEEE Cornell Conf. on Advanced Concepts in high Speed Semicond. Dev. Circuits (1995) p.269 ] Y. Lin, AD. van Rheenen, C-L. Chen, F.W. Smith [ J. Electron. Mater. (USA) vol.22 no.12 (1993) p. 1507] N.X. Nguyen, J.P. Ibbetson, U.K. Mishra [ to be presented at the 38th Electronic Materials Conference, Santa Barbara, 1996]

17.4 Opto-switches using low-temperature MBE GaAs J.F. Whitaker August 1995

A

INTRODUCTION

It is well-known that GaAs grown by molecular-beam epitaxy at reduced substrate temperatures around 200 - 300 0 C can exhibit a unique combination of high resistivity, ultrashort carrier trapping time, very high breakdown voltage, and moderate mobility. These properties, which until the advent of LT-GaAs were not characteristic of any one semiconductor, are also extremely desirable ones for substrates that are to be used in fast optoelectronic devices [1,2]. As a result, LT-GaAs has become a tremendously valuable material for use in optically-activated switches (opto-switches) [3]. These have been applied in many different ways: detecting short pulses of light or shaped optical signals [4,5], generating or gating very brief electrical transients [6], mixing optical signals to obtain multi-terahertz intermediate frequencies [7] and creating near-kilovolt amplitude, single-picosecond electrical pulses [8]. Furthermore, applications that require short-duration optical signals to be reproduced in the electrical domain are becoming increasingly important. These include ranging, medical imaging, high-speed ultraviolet and X-ray radiation generation, characterization of ultra-high-bandwidth electronic components [9], and, especially, communications systems that will eventually operate at speeds of > 100 gigabits per second. The future use of LT-GaAs in opto-switch devices is nearly assured. The application of LT-GaAs for opto-switches can be put into historical context by considering opto-switches which have been used to produce electrical transients of duration less than 100 ps. Since pulses of longer duration can be more effectively produced using other semiconductors with considerably higher mobility than LT-GaAs, this Datareview will consider only opto-switches with a temporal response shorter than -100 ps. A number of approaches have been investigated to achieve short fall times in the response of an opto-switch. These include the identification of photoconductive materials which have very short carrier lifetimes, the fabrication of detector patterns that rely on asymmetric excitation of long-carrier-lifetime materials [10], and the fabrication of pulse-shaping networks to change the step output of a detector into a pulse [H]. Fast-lifetime materials have included amorphous silicon [12], films of CdTe [13], O+-ion-implanted silicon-on-sapphire (SOS) [14] and proton-bombarded GaAs [15] and InP [16]. While these materials which rely on damage and deep-level trapping/recombination centres for their fast response can have very short lifetimes, typically they also have poor sensitivity to the energy in the excitation pulses due to their low mobilities. On the other hand, materials with the highest mobilities do not have fast, picosecond relaxation times. Geometrical approaches to producing short pulses have been moderately successful, although they are arguably more difficult to implement. Thus the development of LT-GaAs for use in generating picosecond and subpicosecond electrical pulses has been met with great enthusiasm. A good deal of effort has been expended in attempting to understand the physical mechanisms responsible for the behaviour of LT-GaAs and the growth conditions necessary to optimize the material for opto-switch applications. As-grown materials, for instance, are found to exhibit a

carrier relaxation time that is virtually the same as that of materials that have been annealed after growth [17]. On the other hand the highest room temperature resistivity and mobility values have been measured only for the annealed layers. The fact that as-grown materials have a high density of excess arsenic found in point defects (leading to hopping conduction through defect states) [18], along with the discovery that after annealing much of the excess arsenic forms into semi-metallic precipitates [19,20], has led to the formulation of two theories to explain the electrical properties of LT-GaAs. These are the buried-Schottky-barrier model [21] and the point-defect model [22,23]. This Datareview will explore how the behaviour of LT-GaAs opto-switches is affected by growth conditions such as temperature and annealing, and demonstrate what trade-offs are necessary to prepare optimized optoelectronic devices fabricated using LT-GaAs. B

OPTO-SWITCH OPERATION

LT-GaAs opto-switches are typically photoconductive devices that are embedded in microstrip or planar transmission lines - Auston switches [3]. When they are illuminated by an optical pulse they allow current to be switched for a period of time determined by the photoexcited carrier trapping or recombination time. The LT-GaAs acts as a resistor with a value that varies in time, so that it is essentially open-circuit before illumination (with a dark current which depends on the resistivity of the material). After carriers are excited by photons with energy greater than the GaAs bandgap, a transient will pass through the switch due to its suddenly higher conductivity, and the decay of this transient is determined by the LT-GaAs carrier relaxation and corresponding increase in resistance. This is described by the following expression (which is strictly true only in the case where the 'on-state' resistance is still much greater than Z0, the characteristic impedance of the transmission line),

WJf)

-v- I ^ T T I ;

(1

>

where Rc is the opto-switch contact resistance, Vbias is the applied DC bias, and R^t) is the time-varying resistance of the opto-switch. This resistance, which governs the behaviour of the opto-switch, can be given by q n(t) bie

+

^) w de

W

where L and w are the length and width of the opto-switch gap, respectively, q is the electronic charge, de is the excitation light absorption depth, \iG ([I11) is the electron (hole) mobility, and n(t) is the time-dependent density of photoexcited electron-hole pairs. One can see from EQN (2) that as the photoexcited carrier density relaxes, the photo-switch resistance increases (to a maximum value proportional to the substrate resistivity, p=l/[qn(O)(|ie+|ih)]) and the switch turns off. It is clear from EQNs (1) and (2) that a material with the fastest relaxation, highest resistivity, highest mobility, and highest breakdown voltage will

be able to generate the shortest pulse (which provides the highest bandwidth) with the highest amplitude, while maintaining the lowest dark current (leading to better signal-to-noise ratios). The outstanding and unique properties of LT-GaAs make it a nearly optimized material for use as a fast opto-switch. C

LT-GaAs PHOTOCONDUCTIVE SWITCHES

The key properties of LT-GaAs for opto-switch fabrication are strongly dependent on the MBE preparation of the epi-layers. For instance, a summary of some of the differences between typical as-grown and annealed LT-GaAs (see Section A) are given in TABLE 1. TABLE 1. Properties of as-grown and post-annealed LT-GaAs. As-grown (T 0 - 200 ° C)

Post-annealed (TA ~ 600 ° C for 10 min)

Excess arsenic

1019 - 1020 cm'3

As precipitates formed (1017-1018cmf3)

Crystal quality

Highly defective

Single crystal like

Dark resistivity

Low (~ 104 -10 5 Q cm)

High (~ 107 Q cm)

Mobility

Low (~ 102 cm2/ Vs)

Moderate (< 103 cm2/ Vs)

Hopping conduction at room temperature

Significant

Not significant

Photoexcited carrier lifetime

Ultrafast (-0.4 - 0.6 ps)

Ultrafast (-0.4 - 0.6 ps)

While other parameters of LT-GaAs preparation, such as beam equivalent pressure, As/Ga flux ratio, annealing temperature, time, method, epi-layer thickness, and growth rate may affect the quality and properties of the material, it is the growth temperature that has been found to give the greatest degree of control over the behaviour of LT-GaAs optoelectronic devices [24]. However, uncertainties in the measurement of the MBE substrate temperature (TG) below 400 0 C have led to a situation where different laboratories produce materials with different characteristics even when they try to use identical growth conditions. Combine the uncertainties in measured TG with the possibility of adjusting the other parameters, and one arrives at a situation where materials with different stoichiometry, defects, and optoelectronic characteristics are all being produced under the label of LT-GaAs. Ongoing research is attempting to identify a growth/annealing protocol that will allow high-quality LT-GaAs to be reproduced reliably. In the opto-switch data shown in this Datareview, only consistent samples from the same laboratory, MIT Lincoln Laboratory, have been used. Typically, Ga and As4 beam fluxes were employed at a V/in beam equivalent pressure ratio of 10, and the growth rate was 1.0 fim/hr on (OOl)-oriented semi-insulating GaAs substrates. An undoped buffer layer of GaAs was first grown at a temperature of 6000C to smooth the growth front, before the substrate temperature was lowered to the desired value to produce a l-|iim-thick layer of LT-GaAs. The annealing was performed in situ inside the growth chamber under As-overpressure by raising the substrate temperature to 6000C for a period of 10 min. All the layers studied here were of high crystalline quality, as verified by their RHEED pattern during growth.

Cl

Resistivity and Breakdown

Typically, annealed LT-GaAs samples have a resistivity of greater than 107 Q cm for TG < 300 0 C, and they have a breakdown electric field of greater than 500 kV/ cm [25]. This resistivity makes LT-GaAs ideally suited as a semiconductor substrate for opto-switches, with dark currents of ~ 10 nA for metal-semiconductor-metal (MSM) interdigitated photodetectors with 2 ^m finger width and separation demonstrated. The outstanding voltage hold-off capability of LT-GaAs also enables straight-gap opto-switches to be used for ultrafast high-voltage photoconductive switching [26], with electrical pulses of nearly a kilovolt amplitude generated in one instance. [8] C2

Time Response

Because of space constraints, opto-switch temporal characteristics which vary only as a function of growth temperature or after anneal will be considered here. For high-speed opto-switches, the growth temperature can be used, for instance, to tune the response of a device in order to obtain a certain detector bandwidth. C2.1

Ultrafast trapping

The fast photoexcited carrier lifetime in LT-GaAs, initially measured when experiments at Lincoln Laboratory failed to detect photoluminescence signals, has been verified in time-resolved measurements on many different samples. The simplest way to test for short lifetime is by a pump-probe, transient reflectivity or transmissivity technique [14,27]. A train of 100 fs-duration laser pulses is split into two beams. The first, known as the pump, excites carriers from the valence band to the conduction band, and the second, or probe beam, monitors changes in the refractive index versus a variable delay between the pump and probe pulses. The temporal resolution of this technique is better than 150 fs when 100 fs laser pulses are employed. The decay of the transient optical properties contains information about the evolution of carrier trapping and/or recombination from the conduction band. It is known that the lattice parameter expands when low-temperature growth is used for MBE-GaAs [28], and that a large density OfAs03 antisite defects is present in as-grown LT-GaAs. It is also known that this expansion and the density of excess arsenic become smaller as growth temperatures above 2000C are used. It is thus interesting to observe what happens to the carrier lifetime measured by transient transmissivity versus growth temperature. FIGURE 1 demonstrates that the carrier relaxation time gradually increases with growth temperature for both as-grown and annealed LT-GaAs layers, starting at several hundred femtoseconds for TG < 200 0 C and reaching several tens of picoseconds for TG = 3000C. Measurements of samples grown at up to TG = 4000C indicate a saturation of the relaxation time at -50 ps. This data suggests that opto-switches and detectors could be designed for particular output pulse-widths or bandwidths depending on the growth temperature of the LT-GaAs. Relaxation time measurements are also relevant to the study of the physical processes taking place in the material. As-grown LT-GaAs, with its high density of point defects, and annealed LT-GaAs, with its high density of As precipitates, have complicated but different structures. Which defect controls the fast carrier lifetime, or whether the same defect is responsible in each case, has been the subject of intense debate [29,30]. For instance, the decrease in the density of AsGa with increasing T G results in fewer trap sites being available for photoexcited carriers

1/e relaxation time (ps)

as-grown annealed

Growth Temperature (0C) FIGURE 1. Photoexcited carrier lifetime vs. growth temperature for as-grown and annealed LT-GaAs as measured by pump-probe transient absorption.

resulting in a longer relaxation time. The fact that the relaxation times for both the as-grown and annealed samples track so closely suggests that point defects may cause the fast lifetime in both materials. However, it can also be argued that point defects could control the lifetime in as-grown layers and precipitates in the post-annealed material. One group has reported pump-probe transmissivity results to support this contention [31], and they have also described how precipitates may be engineered by annealing in order to deliver the desired response [32]. Other compelling evidence for high point-defect densities, even in the annealed layers, indicates that they may be primarily responsible for the short lifetime in this material. C2.2

Time-resolved photoconductivity

Transient photoconductivity measurements on opto-switches fabricated on annealed LT-GaAs have been used to demonstrate that electrical waveforms with pulsewidths ranging from subpicosecond to tens of picoseconds can be generated [33]. In FIGURE 2, representative time-domain data of switched electrical transients for five samples with TG between 200 and 3000C, with the voltage normalized, are compared. The 1/e decay of the photogenerated pulses, x, at the different growth temperatures corresponds to the lifetimes measured in the pump-probe experiments, although the photoconductive response for the devices on the material grown below 230 0 C may have been limited by the external electro-optic sampling measurement technique. Very short pulses of -700 fs FWHM duration have been produced from these opto-switches, which were excited by -100 fs optical pulses. Other groups have reported an output pulse as short as 200-fs duration from an LT-GaAs opto-switch [34]. For the pulses shown in FIGURE 2, an electro-optic crystal placed -100 \xm from the switch gap sensed the output electric field, which was time-resolved using the ultrashort laser pulses [35]. Using the same technique, another group has reported picosecond or sub-picosecond photoconductive response measurements from LT-GaAs opto-switches for a range of growth

Normalized voltage (a.u.)

temperatures between 250 and 3070C [36]. As mentioned earlier, it is difficult to compare results across sample groups from different MBE sites due to differences in temperature calibrations.

Time (ps) FIGURE 2. Normalized photoconductive response of opto-switches fabricated on annealed LT-GaAs grown at various substrate temperatures. The voltages for each switch are the DC bias levels used.

Ultrashort electrical pulses produced by LT-GaAs opto-switches have been utilized for test and characterization of transmission lines [37], high-speed devices, and integrated circuits. The very short pulse widths that can be attained allow time-domain electrical measurements to be obtained before unwanted reflections from secondary sources can corrupt the desired signals. The pulses also provide test signals with bandwidths of hundreds of gigahertz, sometimes with usable frequencies extending out to 1 THz. In addition, the response of the LT-GaAs opto-switches can often be found to relax to a baseline of 0 V, allowing the electrical waveforms to be transformed cleanly to the frequency domain with a simple apodization function. However, some groups have reported that persistent photoconductive tails are always present, and this phenomenon is attributed to hopping conductivity [38]. Among other applications, the use of fast LT-GaAs opto-switches has also been reported in photoconductively-driven antennas in order to transmit and detect picosecond bursts of pulsed radiation [39,40]. MBE layers of LT-GaAs have also been grown on silicon and dielectric substrates, either through growth [41] or epitaxial lift-off and grafting [42], so that integration with other electronic or optoelectronic devices and circuits could be achieved. C3

Responsivity

The low carrier mobility of LT-GaAs hinders the fabrication of highly sensitive photodetectors and highlights the typical detector trade-off of speed and sensitivity. To illustrate this, the responsivities of several simple, straight-gap opto-switches are shown as a function of the growth temperature in FIGURE 3. The fastest materials suffer from the lowest mobilities due to the presence of large numbers of defects in the GaAs. However, the plot indicates that as the defect density decreases with increasing growth temperature, the mobility, and the magnitude of the switched pulses, increases. This demonstrates that if one can accept a detector with a temporal

Responsivity (A/W)

Growth Temperature (0C) FIGURE 3. Responsivity for simple straight-gap opto-switches fabricated on annealed LT-GaAs grown at different substrate temperatures.

response in the 10-ps range, LT-GaAs can be used to obtain a reasonable sensitivity. TABLE 2 summarizes the decay time, average photocurrent and responsivity of a series of opto-switches fabricated on annealed LT-GaAs grown at different temperatures. TABLE 2. Relaxation time, average photocurrent, and responsivity for simple, gap opto-switches grown at 5 different temperatures. T 0 ( 0 C)

T(ps)

Iavg(nA)

RQiA/W)

200

0.40

76

3.8

220

0.51

271

13.5

230

1.6

1570

78.5

270

5.98

5580

279

300

15.52

14100

705

While the responsivities are quite low for all the growth temperatures in FIGURE 35 an enhancement of the responsivity has been realized by the fabrication of MSM interdigitated patterns. In fact, a high responsivity of 0.1 AAV has been reported for an MSM detector grown at 1900C [4]. Due to the subpicosecond carrier response of the LT-GaAs, a bandwidth of 375 GHz was also attained. Other MSM structures on LT-GaAs have been fabricated and tested [38], although wider finger spacings and the use of LT-GaAs with different characteristics have limited the responsivity and bandwidth of these detectors as compared with [4]. D

ALTERNATIVE MATERIALS

A number of alternatives to LT-GaAs have recently been attempted to reproduce the electronic

properties of LT-GaAs without the expense and uncertainty that have been found to accompany the LT-growth. As one example, arsenic ion-implantation has been found to be an acceptable way to incorporate excess As into the GaAs lattice. When implanted samples have been annealed, arsenic precipitation has been observed [43], and opto-switches have been found to exhibit responsivities and temporal characteristics nearly as good as those using LT-GaAs [44]. Low-temperature-grown MOCVD layers of GaAs have also been grown, and these too contain metallic precipitates, and exhibit high resistivity and fast carrier lifetime when tested in an opto-switch configuration [45]. E

CONCLUSION

The short history of the application of LT-GaAs in fast, optically-activated switches has been reviewed. Key differences and similarities between as-grown and annealed materials have been presented, and the scientific and technological relevance of the ultrafast optical study of the material have been demonstrated. A representative collection of data has shown how optoelectronic properties such as carrier lifetime and responsivity vary with growth temperature. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] [II] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22]

J.F. Whitaker [Mater. Sci. Eng. B (Switzerland) vol.22 (1993) p.61-7 ] S. Gupta et al [ J. Electron. Mater. (USA) vol.22 (1993) p. 1449-55 ] D.H. Auston [ Appl. Phys. Lett. (USA) vol.26 (1975) p. 101-3 ] Y. Chen, S. Williamson, T. Brock, F.W. Smith, A.R. Calawa [ Appl. Phys. Lett. (USA) vol.59 (1991) p. 1984-6] S.Y. Chou, Y. Liu, W. Khalil, T.Y. Hsiang, S. Alexandrou [ Appl. Phys. Lett. (USA) vol.61 (1992) p.819-21] Y. Chen, S.L. Williamson, T. Brock [ Appl. Phys. Lett. (USA) vol.64 (1994) p.551-3 ] E.R. Brown, K.A. Mclntosh, K.B. Nichols, CL. Dennis [ Appl. Phys. Lett. (USA) vol.66 (1995) p.285-7 ] T. Motet, J. Nees, S. Williamson, G. Mourou [Appl. Phys. Lett. (USA) vol.59 (1991) p.1455-7 ] M.Y. Frankel, J.F. Whitaker, G.A. Mourou [ IEEE J. Quantum Electron (USA) vol.28 (1992) p.2313-24] D. Krokel, D. Grischkowsky, M.B. Ketchen [ Appl. Phys. Lett. (USA) vol.54 (1989) p. 1046-8 ] M.Y. Frankel, S. Gupta, J.A. Valdmanis, G.A. Mourou [ Electron. Lett. (UK) vol.25 (1989) p. 1363-5] W.C. Nuimally, R.B. Hammond [ in Picosecond Electronics Ed. C.H.Lee (Academic Press, 1984) p.373] M.C. Nuss, D.W. Kisker, P.R. Smith, T.E. Harvey [ Appl. Phys. Lett. (USA) vol.54 (1989) p.57 ] F.E. Doany, D. Grischkowsky, C-C. Chi [Appl. Phys. Lett. (USA) vol.50 (1987) p.460-2 ] B. Lambsdorff, J. Kuhl, J. Rosenzweig, A. Axmann, J. Schneider [ Appl. Phys. Lett. (USA) vol.58 (1991)p.l881-3] A.G. Foyt, FJ. Leonberger, R.C. Williamson [Appl. Phys. Lett. (USA) vol.40 (1982) p.448-50 ] H.H. Wang, J.F. Whitaker, A. Chin, J. Mazurowski, J.M. Ballingall [ J. Electron. Mater. (USA) vol.22 (1993) p. 1461-4] For an overview, see D.CLook [ Thin Solid Films (Switzerland) vol.231 (1993) p.61-73 ] Z. Liliental-Weber, G. Cooper, R. Mariella, C Kocot [J. Vac. Sci. Tech. B (USA), vol.9 (1991) p.2323] M.R. Melloch, N. Otsuka, J.M. Woodall, A.C. Warren, J.L. Freeouf [ Appl. Phys. Lett. (USA) vol.57 (1990) p. 1531-3] A.C. Warren et al [Appl. Phys. Lett. (USA) vol.57 (1990) p.1331-3 ] M. Kaminskaetal[^/. Phys. Lett. (USA) vol.54 (1989) p.1881-3 ]

[23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33] [34] [35] [36] [37] [38] [39] [40] [41] [42] [43] [44]

[45]

M.O. Manasreh, D.C. Look, K R Evans, CE. Stutz [ Phys. Rev. B (USA) vol.41 (1990) p. 10272 ] S. Gupta et al [ Appl. Phys. Lett. (USA) vol.59 (1991) p.3276-8 ] E.R. Brown, F.W. Smith, K.A. Mclntosh [J Appl. Phys. (USA) vol.73 (1993) p.1480 ] M.Y. Frankel, J.F. Whitaker, G.A. Mourou, F.W. Smith, A.R. Calawa [ IEEE Trans. Electron. Dev. (USA) vol.37 (1990) p.2493-8 ] J. Kuhl, E.O. Gobel, Th. Pfeiffer, A. Jonietz [ Appl. Phys. A (Germany) vol.34 (1984) p. 105 ] Z. Liliental-Weber, W. Swider, K.M. Yu, J. Kortright, F.W. Smith, A.R. Calawa [ Appl. Phys. Lett. (USA) vol.58 (1991) 2143-5 ] Z. Liliental-Weber, HJ. Cheng, S. Gupta, J. Whitaker, K. Nichols, F.W. Smith [ J. Electron. Mater. (USA) vol.22 (1993) p. 1465-9 ] A.C. Warren et al [ Phys. Rev. B (USA) vol.46 (1992) 4617-20 ] E.S. Harmon, M.R Melloch, J.M. Woodall, D.D. Nolte, N. Otsuka, CL. Chang [ Appl. Phys. Lett. (USA) vol.63 (1993) p.2248-50 ] MRMelloch et al [ in Ultra-Wideband, Short-Pulse Electromagnetics 2 Eds. L.Carin, L.B.Felsen, (Plenum Press, 1995) p.25-31 ] S. Gupta, J.F. Whitaker, G. Mourou [ IEEEJ. Quantum Electron (USA) vol.28 (1992) p.2464 ] U.D. Keil, D.R. Dykaar [ in OSA Proc. on Picosecond Electronics and Optoelectronics Eds. J. Shah, U. Mishra (Optical Society of America, 1993) p. 189-92 ] J.A. Valdmanis, G. Mourou [ IEEE J. Quantum Electron. (USA) vol.22 (1986) p.69-78 ] D.R. Dykaar et al [ Mat. Res. Soc. Symp. Proc. (USA) vol.241 (1992) p.245-50 ] H. Cheng, J.F. Whitaker, T.M. Weller, L.P.B. Katehi [ IEEE Trans. Microwave Theory Tech. (USA) vol.42 (1994) p.2399-406 ] M. Klingenstein et al [ Appl. Phys. Lett. (USA) vol.60 (1992) p.627-9 ] J.M. Chwalek, J.F. Whitaker, G.A. Mourou [ in OSA Proc. on Picosecond Electronics and Optoelectronics Eds. T.CL.G.Sollner, J.Shah, (Optical Society of America, 1991) p. 15-19 ] A.C. Warren, N. Katzenellenbogen, D.Grischkowsky, J.M. Woodall, M.R. Melloch, N. Otsuka [ Appl. Phys. Lett. (USA) vol.58 (1991) p. 1512-4 ] M.Y. Frankel, B. Tadayon, T.F. Carruthers [Appl. Phys. Lett. (USA) vol.62 (1993) p.255-7 ] H.-J. Cheng, J.F. Whitaker [ in OSA Proc. on Picosecond Electronics and Optoelectronics Eds. J. Shah, U. Mishra (Optical Society of America, 1993) p.201-4 ] A. Claverie, F. Namavar, Z. Liliental-Weber [Appl. Phys. Lett. (USA) vol.62 (1993) p.1271-3 ] H.H. Wang, J.F. Whitaker, H. Fujioka, Z. Liliental-Weber [ in Ultrafast Electronics and Optoelectronics, vol.13, 1995 OSA Technical Digest Series (Optical Society of America, 1995) p.32-4] A. Krotkus, V. Pasiskevicius, T. Lideikis, G. Treideris, D. Lescinskas, V. Jasutis [Appl. Phys. A (Germany) vol.58 (1994) p.177-81 ]

CHAPTER 18 ETCHING 18.1 18.2 18.3 18.4 18.5 18.6 18.7 18.8

Etching of GaAs: an overview Wet etching of GaAs Chemical dry etching of GaAs Plasma etching of GaAs Ion-beam milling and sputter etching of GaAs Reactive ion etching (RCE) and magnetron-enhanced reactive ion etching (MIE) of GaAs Etching of GaAs using ECR-RF reactive-ion etching Etching of GaAs using reactive-ion-beam etching (RIBE), chemicallyassisted-ion-beam etching (CAIBE), and radical-beam-ion-beam etching (RBIBE)

18.1 Etching of GaAs: an overview CLH. Ashby March 1995

This series of Datareviews covers several general types of etching processes for GaAs: wet chemical etching, dry chemical etching, plasma etching, ion beam milling or sputter etching, five types of ion-assisted chemical etching, laser- or photo-assisted etching, and electron-assisted etching. Ordinary chemical etching, wet or dry, does not use ions, electrons, or photons to promote the etching reaction, either through generation of etchants or through formation of products. Plasma etching employs a plasma to generate reactive etchants but the semiconductor is positioned within the plasma where it will not be subjected to appreciable ion bombardment at energies >50 eV. Ion beam milling, sputter etching, or ion-beam etching (EBE) relies on only physical momentum transfer processes to remove material; no chemistry is involved. Ion-assisted chemical etching remains the workhorse of GaAs device fabrication. It includes a number of techniques (RIE, ECR-RF RIE, RIBE, IBAE, CAIBE, RBIBE), all of which involve both energetic ion bombardment and chemically reactive etchants. Reactive ion etching (RIE) is a plasma-assisted etching using ions whose energy is greater than the positive space-charge potential of the plasma. ECR-RF RJE uses an electron-cyclotron-resonance plasma to generate high density, low-ionenergy plasmas remote from the sample surface and applies an RF-bias to the sample electrode to accelerate ions into the sample. Reactive ion beam etching (RIBE) involves generation of reactive chemical species such as Cl+ and Cl2+ in an ion source; a beam of these ions is used to etch GaAs. Ion-beam-assisted etching (IBAE), more commonly called chemically-assisted-ionbeam etching (CAIBE), employs an ambient reactive neutral gas, such as Cl2, and a beam of ions, usually Ar+. Radical-beam-ion-beam etching (RBIBE) is a variant of CAIBE/IBAE in which the neutral gas is predissociated to provide a beam of highly reactive radicals, such as Cl atoms from Cl2. Lasers or other high intensity photon sources produce enhanced etch rates by thermal heating, photochemical generation of reactants, photogeneration of electrons and holes that drive chemical reactions, or some combination of these processes. Electron-beam-assisted etching enhances thermal chemical reactions by bombarding the surface with an electron beam to promote product formation or product desorption. There are four important characteristics to consider in selecting an etching process. These are etch rate, profile (anisotropy and morphology), selectivity, and damage. These are discussed in some detail in each Datareview. In general, the balance between chemical and physical contributions to etching for each process largely determines these characteristics. Large chemical contributions generally produce rapid rates, compositional selectivity, and less damage; they also promote isotropic and sometimes crystallographic etching. Photons and electrons, by enhancing chemical reactions without significant physical momentum transfer, can provide increased anisotropy with minimal or no structural and electronic damage. In general, however, ion-based processes employing significant physical sputtering contributions, with their attendant damage, are used for anisotropic etching. This means that anisotropy and damage minimization are conflicting goals that require compromise in process selection.

There has been tremendous effort in GaAs etching since the 2nd edition of this book, with more than 450 new publications appearing since 1989. Future directions in GaAs etch development will include efforts to retain anisotropy while reducing damage and the development of in-situ process monitoring, mostly by optical techniques, for precise real-time control of the etching process. We may expect GaAs etching to remain an active research area for some time to come.

18.2 Wet etching of GaAs CLH. Ashby March 1995

A

INTRODUCTION

Despite the many advances in plasma-based dry etch technology, wet etching still remains important. It provides the standard when process-induced damage is of concern, and can provide highly selective etching of GaAs versus HI-V ternary compounds in heterostructures. Wet etching processes have three general applications: pattern formation, polishing, and defect or damage visualization. This Datareview is restricted to wet etches suitable for pattern fabrication. Chemomechanical methods suitable for polishing and defect visualization have been reviewed [1-3]. Some polishing etches which are 'free' etches, i.e., not chemo-mechanical, are also useful for pattern formation. There are two categories of wet etches: non-electrolytic and electrolytic. Electrolytic or electrochemical etches employ an external electrical supply to drive and control the etching reaction; non-electrolytic etches do not and are, therefore, easier to use but more difficult to precisely control. There are some common characteristics of many wet etches. First, etching reactions are generally based on oxidation of GaAs and subsequent dissolution of the oxidized products by either acids or bases. GaAs is expected to be insoluble in non-oxidizing acids over the pH range from 1 to 14 [4] and requires addition of an oxidizer, e.g., H2O2 or HNO3, for appreciable etch rates. Hydroxide-based etchants provide some degreasing of the surface by saponification. NH4OH performs better than NaOH or KOH because of its tendency to form complexes with metal ions, like Cu, Al, Cr, etc., that helps remove metal impurities from the surface [5]. Second, preferential etching based on crystallographic orientation can occur with the (111)Ga, also called (111)A, face generally etching a factor of 2 to 5 slower than the (100), (110), or (111)As, also called (111)B, faces. There is also a greater tendency for etch pits to form on (111) faces. Third, it is often possible to define different concentration ratios of the same reactants that permits either highly selective etching of GaAs relative to a related ternary compound, such as AlGaAs or InGaAs, or, conversely, equirate etching of GaAs and the ternary. Fourth, ion bombardment to amorphize the surface can greatly increase the etching rate for both electrolytic [6] and non-electrolytic [7] etching. B

NON-ELECTROLYTIC ETCHING

Bl

Rates and Profiles

Non-electrolytic etching rates (TABLE 1) can be either diffusion limited or chemical-reaction limited and the etching profile is profoundly affected by the dominant rate limitation. Whether a given etching rate is diffusion or chemical-reaction controlled is often determined by the relative proportions of the constituents of the etching solution. Dilution favours chemical-reaction control. An increase in the viscosity of the solution increases the relative importance of diffusion.

TABLE 1. Non-electrolytic wet etches for GaAs. Etchant

Etch rate (jimZmin)

T(K)

NH4OHZH2O2

0.1

NH 4 OHM 2 O 2

(Hl)GaZ(IIl)As

Ref.

Comments

298

[13]

pH = 7.04±0.02, agitated

0.017

298

[13]

pH = 7.04±0.02, stagnant

NH 4 OHM 2 O 2

4.1

298

[25]

pH = 8.4, 32:1 vs. Al016Ga084As

0.3NNH 4 OH/

0.18-0.2

RT

1/6-1/4

[47]

n-type

0.1NH 2 O 2

0.12-0.14

RT

1/6-1/4

[47]

p-type

3NNH 4 OHZlNH 2 O 2

1.4

RT

<1

[47]

little mask attack

NH4OHZH2O2Z H2O

0.5-5

293

composition dependent

[5]

very smooth

0.08MNH 4 OHZ 0.2MH 2 O 2

0.3

297

1/4

[48]

pH 10-12

INH4OHZlH2O2Z 8H 2 O

1.5

RT

[19]

250 H2O2Z INH4OH

0.1

293

[26]

NaOHZH2O2

0.2 1

303 328

[49] [49]

NH 4 S x (NH 4 S+ 4%S)

0.018

333

[50]

AlxGa1^SSeLtG up

0.1MEDTAZ0.2MH 2 O 2

0.3

297

[48]

pH 10-12

2HNO3Z3H2O2

20.5

4/5

[20]

(111)AZ(Hl)B* 1 for 1:9 mix

H2SO4ZH2O2ZH2O

3.5

273

up to 1

[15]

orientation dependence f(H2SO4:H2O2)

1H2SO4Z8H2O2Z IH 2 O

14

RT

1/4

[17]

H3PO4ZH2O2ZH2O

0.01-4

303

composition dependent

[16]

GaAsZIn0 ^ a 0 <>As> 50:1

reproducible to 10's of A for below 0.1 jimZmin

H3PO4ZH2O2Z 3CH 3 OH

2

RT

1/2

[51]

extremely smooth

•Citric acid/H 2 O 2

0.6

297

1/2

[27]

no mask attack

* 10 Citric acidZlH2O2

2

291

[28]

95:1 vs. Al03Ga07As

TABLEl. Continued. Etchant

Etch rate (umZrnin)

T(K)

•Citric acid/H2O2

0.15-0.4

*4 Citric acid/ IH 2 O 2

(Hl)GaZ (Hl)As

Ref.

Comments

300-273

[29]

selective vs. Al028Ga078As: > 80:1 to 0.7 to 1 as f(H2O2)

0.36

293

[30]

selective vs. AlxGa^xAs: x = 0.3, 0.45,1:155,260, 1450

*Citric acid/H2O2: 0.5:1 to 50:1 (5:1 max rate)

0.006-0.040 (0.31)

RT

[18]

116:1 to 1:1 vs. thick Al03Ga0 7As as f(citric:H2O2)

*4Citric/lH 2 O 2 /lH 2 O

0.35

300-273

[29]

1:1 vs. Al028Ga078As

**15 Succinic/ 1H2O2/1H2O2

0.15

RT

[31,32]

selective vs. AlxGa1^As as f(succinic:H2O2, x, pH,T)

HCl/4H 2 O 2 /40H 2 O

0.22

RT

below 1

[17]

40HCl/4H 2 O 2 /H 2 O

>5.0

300

1

[17]

isotropic

1HC1/20CH 3 OOH/3H 2 0 2

0.27

293

[24]

selective vs. GaInP as f(H2O2)

HF

0.1

353

[34]

1> 10 vs. Al fraction ;>0.4

HFZH2O2ZH2O

10-0.1

298

[21]

high(H2O2:H2O) = crystallographic low (H2O2:H2O) = isotropic

K2Cr2O7ZH2SO4ZHClZ CH 3 OOH

0.01-1

293

[14]

no mask attack or pits

K2Cr2O7ZH2SO4ZHCl

2.5 20

298 333

[14] [52]

no mask attack no pits

20HFZCrO3

0.1-0.2

RT

[53]

n = p; incubation

20HClZCrO3

0.1-0.2

RT

[54]

n = p = SI

KIZI2

1

RT

8.1K3Fe(CN)6Z 6.3K4Fe(CN)6Z 100H2O

0.023

297

5%Br 2 -CH 3 OH

5-7

RT

1HFZ3HNO3Z 3CH3COOHZ5H2O

20

297

f(H2O2)

<1

1Z5

[36] [12,26,36]

selective vs. In ol Ga O9 As>8.1;pH 6.65

[23]

lower Br2, more preferential

[8]

wafer thinning; very uniform if rotated

* Citric = 1 gm citric acid/1 g H2O (50 wt% solution); H2O2 = 30 wt% solution ** Succinic = 1 gm succinic acid/5 gm H2O; H2O2 = 30 wt% solution; NH4OH adjusts pH RT = assumed room temperature

For diffusion-limited etches, the rate is controlled either by the mass transport of reactants to the surface or of products from the surface. These etches tend to be isotropic and relatively insensitive to temperature but highly sensitive to changes in the nature and degree of agitation. Slow rotation (less than 100 rpm) permits very uniform etching for wafer thinning [8]. Fast rotation in a centrifuge decreases lateral undercutting and increases the etch rate when the centrifugal force is directed away from the surface [9]. Etching profiles at resist edges for diffusion-controlled reactions have been modelled and verified experimentally [10,11]. Local galvanic effects can lead to cathodic protection of certain facets and give crystallographic etching under normally isotropic conditions [12]. Diffusion-limited etches are also especially good for polishing wafers. For chemical-reaction-limited etches, the rate is controlled by the actual chemical reactions at the GaAs surface rather than by mass transport considerations. The oxidation rate is determined by the oxidizing strength of the solution and product dissolution is controlled largely by pH [4]. Rates increase with increasing H2O2 but surfaces roughen when H2O2 exceeds 15% [4]. Surface roughness also increases above pH = 8 [4,13]. Chemical-reaction-limited etches tend to be anisotropic with respect to certain crystallographic orientations, quite sensitive to temperature, and relatively insensitive to agitation. They are well suited for etching geometric shapes along crystallographic planes. The wall angle depends on mask orientation relative to different (110) directions [14-23]. Mask edges aligned perpendicular to the <011> direction produce sloped walls due to the slower etch rate of the (111) Ga plane. Orientation orthogonal to allows vertical profiles due to the faster etch rate of the (111)B plane [4]. Proper alignment of linear patterns at 45° to the natural cleavage planes permits vertical (100) sidewalls for (100) substrates [19]. An extensive study of rates and etch profiles produced by HF/H2O2/H2O etching of (100), (110), (111)Ga, and (111)As surfaces as a function of acid/oxidant/water ratio shows the wide variation in wall angles. A comparison is also made with H2SO4-, H3PO4-, citric-, and NH4OHperoxide and Br2-methanol etchants [21]. Etch rates can depend strongly on the history of the etching solution. Rates most frequently decrease with solution age but can vary non-monotonically as a function of time in solutions containing both HCl and H2O2 [24]. Rates have been observed to first increase with time after mixing and then decline. It is proposed that the H2O2 generates Cl2 in solution, enhancing the rate; subsequent loss OfCl2 causes the rate to decline. Trenching at the edge of etched features is also increased as the solution ages [24]. Trenching under diffusion control has been modelled [10] and experimentally verified with several common etchants [H]; it is highly dependent on exposed vs. masked areal ratio. Trenching is affected by acid:oxidant ratio in lHCl:yH2O2: IH2O solutions with negligible trenching for y < 10 and up to 10 to 20% deeper etching near mask edges than at the centre between two masked areas for y = 40 [24]. B2

Selectivity

The ability to selectively etch GaAs versus a ternary such as AlGaAs, InGaAs, or GaInP has become increasingly important due to the ubiquity of heterostructure devices, and much of the

recent wet etch work has been in this area. Selective etching of GaAs versus AlGaAs occurs at pH > 6 in NH4OHZH2O2 solutions [13,25]. Organic acid/H2O2 mixtures also selectively etch GaAs versus AlGaAs, with higher selectivity ratios than provided with NH4OHTH2O2. The relative amount of the oxidant, H2O2, and the degree of dilution at a fixed acid/oxidant ratio is often the determining factor for selectivity. When the pH of 30% H2O2 is varied between 1 and 6 by addition of either NH4OH or H3PO4, both GaAs and AlGaAs etch [13]. At pH > 6, this solution is selective for GaAs vs. AlxGa1^xAs, x > 0.1 [13]. Surfaces are smooth with 6 < pH < 7.1, but roughen at higher pH due to less efficient removal of the oxide film, which cracks into small segments and floats away during etching to permit formation of new oxide sheets [13]. Above pH = 7.5, etching at surface imperfections predominates, leading to a rough surface of overlapping etch pits [13]. GaAs etch rates increase with pH faster than AlGaAs rates [25]. At pH = 8.4, a 32:1 selectivity versus Al016Ga0 84As occurs; at pH = 8.6 and 8.8, nearly equirate etching occurs [25]. The selectivity for GaAs versus InGaAs can exceed 50 [26]. The organic acid etchants, citric acid/H2O2 [18,27-30] and succinic acid/H2O2 [31,32], have an advantage over NH4OHTH2O2 of not attacking common positive photoresists. They permit the entire range from equirate etching to > 100-fold selectivity for GaAs versus AlGaAs or 17:1 for InGaAs versus GaAs [18], with the selectivity being determined by the organic acid:oxidant ratio and the Al or In mole fraction. They produce smooth surfaces, but they also etch crystallographically [18,28]. It is important to note that selectivity ratios determined with thick layers may not be realized with very thin layers in heterostructures. Citric acid/H202has been the most extensively studied [18,27-30]. For the volume ratio 0.5 to 2.0, In 02 Ga 08 As etches faster than GaAs or Al03Ga08As. Volume ratios between 3.0 and 7.0 give approximately equirate GaAs and InGaAs etching and selectivity relative to AlGaAs. Volume ratios above 8.0 give roughly equirate etching for all three materials. Dilution with water can produce equirate GaAs/AlGaAs etching with normally selective 4:1 citric/H2O2 [29]. Succinic acid/H2O2 with pH adjustment with NH4OH [31,32] permits nearly equirate etching of GaAs and AlxGa1^As for x < 0.4 and >100:1 selectivity for x > 0.5 at pH = 4.2 [31]. Increasing the pH between 4.4 and 5.4 increases the selectivity between x = 0.2 and 0.4 to 150:1 [32]. However, surfaces roughen at pH >5.1 [32]. Decreasing the H2O2 fraction decreases AlGaAs etch rates for x = 0.1 to 0.6 with the rate of decrease greater at lower x [32]. AlxGaLxAs with x > 0.38 can be etched preferentially to GaAs with boiling HCl [33], HF [34,35], or KIfI7ZH2O [36-38]. For lower x, AlGaAs etches less than 2-fold faster than GaAs in HF [34]. The K3Fe(CN)6ZK4Fe(CN)6 system can etch either GaAs or AlGaAs selectively, depending on the ratio of the Fe(II) and Fe(III) species and the pH [36]. It also etches GaAs vs. InxGa1^As with > 8:1 selectivity for x = 0.1 and with negligible etching of InGaAs for x > 0.15. Selectivity versus Al03Ga07As was 1.4 [26]. C

ELECTROLYTIC ETCHmG

Electrolytic etches (TABLE 2) [39-45] permit accurate depth control by monitoring the amount of current passed during the etching process but require ohmic contacts to the wafer, thereby

restricting their application. They are based on anodic oxidation of GaAs followed by dissolution of products. Etching of n-type material is much slower than of p-type unless it is illuminated to photogenerate holes in the surface region or unless sufficient voltage is applied to produce breakdown of the surface barrier. Photoelectrochemical (PEC) etching is reviewed separately in this volume [46]. TABLE 2. Electrolytic wet etches for GaAs. Electrolyte

Etch rate (umZrnin)

Current density (mAZcm2)

T(K)

Ref.

Comments

3MNaOH

4

100

RT

[39]

p>n

0.5NKOH

1

20

298

[40]

p-type

INHClO 4

1

20

298

[40]

p-type

IHClO 4 / 4CH3COOH

1

20

RT

[41]

+V for n-type

O. IN HCl

0.2

10

300

[42]

surface residue if above 40 mAZcm2

0.1NHNO 3

0.12

10

RT

[42]

surface residue if above 40 mAZcm2

0.05% HF

f(scan rate)

STMInA 5 10Hz

RT

[43]

20 nm linewidth

NaOHZNa2EDTA

0.05-0.25

cycled

RT

[44]

p-type, isotropic

PtZH2SO2 (Platanex)

4 nmZpulse

25 V, 250Hz

313

[45]

smooth

ACKNOWLEDGEMENTS This work was performed at Sandia National Laboratories and supported by the U.S. Department of Energy under Contract No. DE-AC04-94AL85000.

REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] [II]

B. Tuck [ J. Mater. Sci. (UK) vol. 10 (1975) p.321 ] DJ. Stirland, BW. Straughan [ Thin Solid Films (Switzerland) vol.31 (1976) p. 139 ] W. Kern [ RCA Rev. (USA) vol.39 (1978) p.278 ] G. Franz, C. Hoyler, D. Sacher [ Jpn. J. Appl. Phys. Pt.l (Japan) vol.30. (1991) p.2693 ] L.P. Molodyakova, V.M. Andreev [ Inorg. Mater. (USA) vol. 12 (1977) p.956 ] Y. Yamamoto, S. Yano [ J. Electrochem. Soc. (USA) vol. 122 (1975) p.260 ] T. Inada, K. Kodama, M. Kitahara [ Nucl. Instrum. Methods Phys. Res. B (Netherlands) vol.7/8 (1978) p.278] J.M. O'Connor, E.F. Dvorsky, H.S. Hier, W.P. Reif [ J. Electrochem. Soc. (USA) vol. 135 (1988) p. 190] L.N. Vozmilova, N.V.Glazunova [ Inorg. Mater. (USA) vol.23 (1987) p.606 ] H.K. Kuiken, JJ. Kelly, P.H.L. Notten [ J. Electrochem. Soc. (USA) vol. 133 (1986) p. 1217 ] P.H.L. Notten, JJ. Kelly, H.K. Kuiken [J. Electrochem. Soc. (USA) vol.133 (1986) p. 1226 ]

[12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33] [34] [35] [36] [37] [38] [39] [40] [41] [42] [43] [44] [45] [46] [47] [48] [49] [50] [51] [52] [53] [54]

P.H.L. Notten, JJ. Kelly [ J. Electrochem. Soc. (USA) vol. 134 (1987) p.444 ] RA. Logan, F.K. Reinhart [ J. Appl. Phys. (USA) vol.44 (1973) p.4172-6 ] S. Adachi, K. Oe [ J. Electrochem. Soc. (USA) vol. 131 (1984) p. 126 ] S. Iida, K. Ito [ J. Electrochem. Soc. (USA) vol. 118 (1971) p.768 ] Y. Mori, N. Watanabe [J. Electrochem. Soc. (USA) vol.125 (1978) p.1510 ] D.W. Shaw [J. Electrochem. Soc. (USA) vol.128 (1981) p.874 ] G.C. DeSalvo, W.F. Tseng, J. Comas [J Electrochem. Soc. (USA) vol.139 (1992) p.831 ] S.H. Jones, D.K. Walker [ J. Electrochem. Soc. (USA) vol. 137 (1990) p. 1653 ] M.M. Berdichenko, LN. Vozmilova, V.G. Malyarova [Inorg. Mater. (USA) vol.24 (1989) 1028-9 ] T. Takebe, T. Yamamoto, M. Fijii, K. Kobayashi [ J. Electrochem. Soc. (USA) vol.140 (1993) p. 1169] C. Juang, KJ. Kuhn, R.B. Darling [ J. Vac. Sci. Technol. B (USA) vol.B8 (1990) p. 1122 ] Y.Tarui, Y.Komiya, Y.Harada [J. Electrochem. Soc.(USA) vol.118 (1971)p.l 18 ] J.R Flemish, K.A. Jones [ J. Electrochem. Soc. (USA) vol. 140 (1993) p.844 ] K. Kenefick [ Electron. Lett. (UK) Vol.21 (1985) p.558-9 ] D.G. Hill, K.L. Lear, J.S. Harris, Jr. [ J. Electrochem. Soc. (USA) vol. 137 (1990) p.2912 ] M. Otsubo, T. Oda, H. Kumabe, H. Miki [ J Electrochem. Soc. (USA) vol. 123 (1976) p.676 ] C. Juang, KJ. Kuhn, RB. Darling [ J. Vac. Sci. Technol. B (USA) vol.B8 (1990) p. 1122 ] B.-Y. Mao, J.A. Nielsen, R.A. Friedman, G.Y. Lee [ J. Electrochem. Soc. (USA) vol. 141 (1994) p. 1082] M. Tong, D.G. Ballegeer, A. Ketterson, EJ. Roan, K.Y. Cheng, I. Adesida. [ J. Electron. Mater. (USA) vol.21 (1992) p.9 ] AJ. Tang, K. Sadra, B.G. Streetman [ J. Electrochem. Soc. (USA) vol. 140 (1993). p.L82 ] S.A. Merritt, M. Dagenais [ J. Electrochem. Soc. (USA) vol. 140 (1993) p.L138 ] Zh. I. Alferov, S.A. Gurevich, M.I. Mizerov, E.L. Portnoy [Zh. Tech. Fiz. (Russia) (1975) p.2602 ] X.S. Wu, L.A. Coldren, J.L. Merz [ Electron. Lett. (UK) vol.21 (1985) p.558-9 ] J.L. Merz, RA. Logan, A.M. Sergent [ IEEE J. Quantum Electron. (USA) vol.QE-5 (1979) p.72 ] RP. Tijburg, T.van Dongen [ J. Electrochem. Soc. (USA) vol. 123 (1976) p.687 ] B.F. Levine, RA. Logan, W.T. Tsang, CG. Tethea, F.R. Merrit [ Appl. Phys. Lett. (USA) vol.42 (1983)p.339] A. Malag, J. Rataczak, J. Gazecki [Mater. Sci. Engin. B (Switzerland) vol.B20 (1993) p.332 ] CJ.Nuese, J.J.Gannon [ J. Electrochem. Soc. (USA) vol. 117 (1970) p. 1094 ] W.W. Harvey [ J. Electrochem. Soc. (USA) vol. 114 (1967) p.472 ] N.A. Pakhanov, A.S. Terekhov [ Instrum. & Expt. Tech. (USA) vol. 18 (1975) p.62 ] B. Schwartz, F. Ermanis, M.H. Brastad [ J. Electrochem. Soc. (USA) vol. 123 (1976) p. 1089 ] L.A. Nagahara, T. Thunday, S.M. Lindway [ Appl. Phys. Lett. (USA) vol.57 (1990) p.270 ] CY. Chen, W. Reichert, RM. Cohen [Mater. Lett. (Netherlands) vol. 19 (1994) p. 109 ] A. Grub, K. Frick, H.L. Hartnagel [ J. Electrochem. Soc. (USA) vol.138 (1991) p.856 ] C.I.H. Ashby [ Datareview in this book: 19.7 Laser assisted etching of GaAs ] JJ. Gannon, CJ. Nuese [ J. Electrochem. Soc. (USA) vol. 121 (1974) p. 1215 ] JJ. Kelly, A.C. Reynders [Appl. Surf Sci. (Netherlands) vol.29 (1987) p. 149 ] I. Shiota, K. Motoya, T. Ohmi, N. Miyamoto, J. Nishizawa [ J.Electrochem. Soc. (USA) vol.124 (1977) p. 155] J.-W. Seo, T. Koker, S. Agarwala, I. Adesida. [ Appl. Phys. Lett. (USA) vol.60 (1992) p. 1114 ] J.L. Mertz, RA. Logan [J. Appl. Phys. (USA) vol.47 (1976) p.3503-9 ] S. Adachi, H. Kawaguchi, G. Iwane [J. Mater. Sci. (UK) vol.16 (1981) p.2449 ] J. van de Ven, J.L. Weyher, J.E.A.M. van der Meerakker, JJ. Kelly [ J Electrochem. Soc. (USA) vol.133 (1986) p.799] J. van de Ven, A.F. Lourens, J.L. Weyher, LJ. Giling [ Chemtronics (UK) vol. 1 (1986) p. 19 ]

18.3 Chemical dry etching of GaAs CLH. Ashby March 1995

A

INTRODUCTION

Chemical dry etching (CDE)5 as defined here, employs no deliberate direct impact on the sample surface by ions, electrons, or photons. Prior to 1990, dry etching of GaAs had focused on slow etches for in-situ surface cleaning before MBE or MOVPE deposition. More recently, however, the need for virtually damage-free processing has stimulated work aimed at patterned fabrication. Some processes are outlined in TABLE 1. Most processes employ Cl species that are thermally or plasma activated. The crystallographic dependence of Cl2 CDE rates studied under high vacuum conditions depends markedly on temperature [I]. For temperatures below 450 0 C, the commonly observed dependence of (111) B > (110) = (100) > (111)A

(1)

is observed, but above 4500C all orientations etch at nearly equal rates. A strong rate dependence on temperature is observed below 1500C (Eact = 1 0 kcal/mol) and above 4500C (Eact = 1 6 kcal/mol) with only a weak dependence on substrate temperature between 150 and 4500C for ( H l ) B , (110), and (100) substrates; Cl2 transport limitation is proposed in this regime [I]. Processes based on reaction with free radical Cl atoms can be less temperature dependent than those using molecular Cl2. B

PRE-DEPOSITION ('CLEANING') ETCHES

Various processes for in-situ surface preparation for MBE, OMVPE, or MOCVD growth have been developed (TABLE 1) that result in congruent removal of Ga and As [2-4] or simply remove native oxide [5,6]. These treatments improve material and/or device quality by such effects as lowering growth defect densities [5] or reducing buried interface charge [4]. Some processes have been developed that etch GaAs and AlGaAs of varying Al fraction at the same rate [7,8]. HCl at 3050C etches InAs at 5 nm/min, 60 times faster than the negligible GaAs etch rate [9]; temperatures >400°C are required for GaAs etch rates above 1 nm/min with HCl [9]. HCl produces very smooth surfaces if used at 5000C for 5 min.; slight surface deterioration is seen at 6500C [10]. Without H2 addition to HCl, RHEED shows some As loss [10]. With H2, the product species changes from AsCl3 to AsH3 and an As- or Ga-stabilized epitaxial-quality surface is obtained, depending on temperature [10]. HCl cleaning produces an order of magnitude lower carbon at interfaces than with thermal surface cleaning [10]. High growth temperatures remove the need for thermal treatment to reactivate H-passivated dopants. With thermal Cl2, surface morphology is smooth only between 300 and 400 0 C ; degradation is seen at T < 2500C or T > 4500C [11]. Cl2 rates are nearly constant between 200 and 4500C [10]. A surface amorphous layer, probably GaClx and AsClx, is removed by heating to 4000C for 10 min to produce a surface as good as with thermal desorption of native oxide under the same conditions, i.e., atomically ordered and stoichiometric [H].

TABLE 1. Dry chemical etches for GaAs. Etchant

Etch rate (j^m/min)

Pressure (torr) (or flow in seem) ( 0 C)

Temp (0C)

Ref

Comments

MBE pre-etch

Pre-deposition etches:

HCl

native oxide

>0.1

25

[5,6]

HCl

> 0.001

2 XlO'4

>400

[9,10,23]

HC1+H2

> 0.001

2.8 xlO"4

500

[10]

no amorphous residue

Cl2

0.2

4 xlO"4

300 - 400

[2]

mirror-like

Cl2

0.2

(1)

350

[7]

MBE pre-etch equirate GaAs/AlGaAs

H 2 /AsH 3

0.0010 - 5

600 (H2) 3 (AsH3)

900

[3]

congruent, mirror-like

AsBr 3 : AsH 3 :H2

0.1

754 (H2) 5.9 (AsH3) 0.002 (AsBr3)

900

[4]

MOVPE preetch

AsCl 3

monolayer

500 - 650

[24]

layer-by-layer control

CH3IZH2

0.01

480

[8]

MOVPE preetch equirate GaAs/AlGaAs

Se+H

monolayer

450

[25]

self-limiting digital etch

(31 CH3I) (2100H2)

Patterning etches

Cl2

0.01-3

8xl0-5-103

100-700

[1]

thermal; xtallog f(T)

Cl 2 (Cl)

2-5

3 - 10 x 10"3

300

[17,18]

hot gas jet; anisotropic

Cl + Cl2

0.004

3 - 5 XlO"5

RT

[H]

ECR remote plasma smooth, stoich

CH-Cl 2

0.1

4 x 1O"4(5)

50

[13]

ECR remote plasma, f(T)

Cl2

1.5

120

[15,16]

supersonic gas expansion

Cl2

0.1-0.6

3.7 - 4.9 x IO'4

100

[21,22]

fast atom beam of neutralized (95%)1.5kV beam, vertical

CF 3 Br(Br)

0.01

3 - 10 x 10"3

RT

[17]

hot gas jet

TABLEl. Continued. Etchant

Etch rate (jim/min)

Pressure (torr) (or flow in seem) ( 0 C)

Temp ( 0 C)

Ref

Comments

I2

0.05-1.7

0.1 -1.25

270 - 330

[26]

isotropic

ICl

0.15-22

0.1 -0.8

100-300

[27]

HCl

1-3

2.5-15 x 1O"2

290

[20]

hot W filament

HCl

0.15-0.3

4 x 10-4

21 -250 (5)

[13]

ECR remote plasma ~ temp, independent

ICH4/5H2+F

0.01

195

[14]

F from SF6+Ar(ECR)

Surface cleaning with AsCl3/H2 shows a dopant-dependent micromasking effect [12]. While Be is removed without surface accumulation, Si accumulates at the surface; accumulation is more severe with delta doping than with bulk-doping. The residual Si micromasks the surface and causes roughening [12]. C

PATTERNING ETCHES

The etches developed for in-situ surface preparation tend to be quite slow since minimal material removal is desired. In contrast, faster etch rates are needed for pattern formation and much recent effort has been directed toward remote plasma processes for faster etch rates without the damage resulting from ion-based processes. The relatively high concentrations of highly reactive species present downstream from ECR plasma sources produce usable etch rates for pattern formation. Most recently developed processes referred to as 'chemical dry etching' (CDE) are such remote plasma processes. Downstream etching with 20% C1/C12 remote plasmas produces three times faster etching than with Cl2 alone [13]. The observed 45-fold rate increase between room temperature and 250 0 C is attributed to deposition of fortuitous surface contaminants, such as oxides, that prevent observation of the thermodynamically predicted temperature-independent, reactant-flux-limited rates. Remote plasma etching with HCl is essentially temperature independent, as predicted, due to the availability of H radicals to remove surface contaminants [13]. HCl does not etch significantly without plasma activation [13]. With ECR-plasma-generated Cl radical etching, an initial period of slow oxide removal is followed by constant etching at 40 A/min at room temperature with a total pressure of 4 x 10'5 torr (CHCl2) [11]. Cl2 does not etch at room temperature without prior removal of oxide; however, brief exposure to Cl radicals permits subsequent etching with Cl2 alone although at a slower rate than when radicals are present [H]. Etching produced no increase in the 50 A roughness of the original surface after etching 2640 A [ I l ] . Fluorine atoms generated in a remote ECR plasma have been used to generate methyl radicals and H atoms from CH4ZH2 in a downstream etching chamber [14]. Crystallographic profiles result,

with the (111)B plane etching faster. Strictly thermal processes that do not use plasmas also exist [1,15-19]. Profiles are isotropic with typical crystallographic dependencies [1] and undercutting observed. Between 300 and 550 K, Ga-to-As product ratios are consistent with stoichiometric etching although the major As species changes from AsCl3 below 300 K to As4 at above 550 K with a mixture of the two occurring at intermediate temperatures [19]. A fast etching process has been developed that uses a hot (>1723 K) W [17] or W-Re [18] resistively heated jet nozzle to thermally dissociate Cl2 to Cl and to direct the radical beam onto the GaAs. Higher nozzle temperatures increase the C1/C12 ratio in the beam; this increases both the etch rate and the anisotropy ratio of etched depth to undercut. Anisotropy ratios of 5 to 10 with better than 10% etch uniformity have been achieved. Hot W filaments have also been used to generate reactive Cl from HCl [20]. Supersonic beam expansions of Cl2 have been used to produce fast etching at low substrate temperatures [20,21]. Fast atom beams [21,22] have been produced by neutralizing accelerated ions to produce vertical, smooth patterned etching with relatively little damage; a damage layer thickness of only 150 A was produced with 1.5 kV atoms [21]. An aspect ratio of 10:1 with 0.6 |nm features has been achieved with a 95% neutralized beam. Some sidewall bowing can occur, as in ion processes, due to fast-atom reflection and scattering by mask edges [22]. ACKNOWLEDGEMENTS This work was performed at Sandia National Laboratories and supported by the U.S. Department of Energy under Contract No. DE-AC04-94AL85000. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] II1] [12] [13]

N. Furuhata, H. Miyamoto, A. Okamoto, K.. Ohata. [ J Electron. Mater. (USA) vol.19 (1990) p.201 ] N. Furuhata, H. Miyamoto, A. Okamoto, K Ohata [ J. Appl. Phys. (USA) vol.65 (1989) p. 168 ] A. Okubora, J. Kasahara, M. Arai, N. Watanabe [ J Appl Phys. (USA) vol.60 (1986) p. 1501 ] T.F. Kuech, E. Marshall, GJ. Scilla, R. Potemski, CM. Ransom, M.Y. Hung [J Cryst. Growth (Netherlands) vol.77 (1986) p.539 ] J.P. Contour, J. Massies, A. Saletes, M. Outrequin, F. Simondet, J.F. Rochette [ J. Vac. Sd. Technol. B (USA) vol.5 (1987) p.730 ] H. Kizuki, N. Hayafuji, N. Fujii, N. Kaneno, Y. Mihashi, T. Murotani. [ J. Cryst. Growth (Netherlands) vol.134 (1993) p.35-42 ] H.G. Lee, RJ. Fischer, GJ. Zydzik, A.Y. Cho, [J Vac. Sci. Technol. (USA) vol.Bl 1 (1993) p.989 ] CW. Krueger, CA. Wang, M. Flytzani-Stephanopoulos. [Appl. Phys. Lett. (USA) vol.60 (1992) p. 1459] Y. Kadoya, T. Yoshida, T. Someya, H. Akiyama, H. Noge, H. Sakaki. [ Jpn. J. Appl. Phys. 2 (Japan) vol.32 (1993) p.L1496 ] J. Saito, K. Kondo. [ J. Appl. Phys. (USA) vol.67 (1990) p.6274-80 ] S. Kohmoto, Y. Ide, Y. Sugimoto, K. Asakawa. [ Jpn. J. Appl. Phys. Pt. I (Japan) vol.32 (1993) p.5796 ] C A C Mendonsa, T.H. Chiu, M.D. Williams, and F.G. Storz. [ Electron. Lett. (UK) vol.30 (1994) p.1717-8] D. G. Lishan, E.L. Hu [ J Vac. Sci. Technol. B (USA) vol.8 (1990) p. 1951 ]

[14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27]

J.E. Spencer, T.R. Schimert, J.H. Dinan, D. Endres, T.R. Hayes. [J. Vac. Sci. Technol. A (USA) vol.8 (1990) p. 1690] T. Ono, H. Kashima, S. Hiraoka, K. Suzuki, A. Jahnke. [J. Vac. Sci. Technol. B (USA) vol.9 (1991) p.2798 ] L.A. DeLouise. [ J Chem. Phys. (USA) vol.94 (1991) p. 1528 ] M.W. Geis, N.N. Efremow, G.A. Lincoln [J. Vac. Sci. Technol. B (USA) vol.4 (1986) p.315 ] M.W. Geis, N.N. Efremow, S.W. Pang, A.C. Anderson [ J. Vac. Sci. Technol. B. (USA) vol.5 (1987)p.363] H-Q. Hou, Z. Zhang, S. Chen. C. Su, W. Yan, M. Vernon. [ Appl. Phys. Lett. (USA) vol.55 (1989) p.801 ] J.L. Dupuie, E. Gulari. [ J. Appl. Phys, (USA) vol.71 (1992) p.4030 ] F. Shimokawa, H. Tanaka, Y. Uenishi, R. Sawada [ J. App.Phys. (USA) vol.66 (1989) p.2613 ] F. Shimokawa. [ J Vac. Sci. Technol. (USA) vol.AlO (1992) p. 1352 ] C. Su, Z.-G. Dai, W. Luo, D.-H. Sun, M.F. Vernon, B.E. Bent. [ Surf. Sci. (Netherlands) vol.312 (1994).p.l81-97] W.T. Tsang, T.H. Chiu, R.M. Kapre [ Appl. Phys. Lett. (USA) vol.63 (1993) p.3500-2 ] S. Takatani, T. Kikawa. [ Appl. Phys. Lett. (USA) vol.65 (1994) p.2585-7 ] K-C. Wong, E.A. Ogryzlo. [ J. Vac. Sci. Technol. B (USA) vol. 10 (1992) p.668 ] L. Hahn, K.-C. Wong, E.A. Ogryzlo [ J Electrochem. Soc. (USA) vol. 140 (1993) p.226 ]

18.4 Plasma etching of GaAs C.I.H. Ashby March 1995

A

INTRODUCTION

Plasma etching of GaAs is based on the chemical reaction between the GaAs surface and plasmagenerated species such as halogen atoms or H atoms. Etch conditions are selected to minimize contributions from ion bombardment. This minimizes process-induced damage but also limits the anisotropy to such an extent that plasma etching has little utility for fabricating the small features that are characteristic of today's devices. B

RATES

Plasma etching is faster than reactive ion etching due to the higher pressure and the resultant higher densities of reactants. The dominant contribution of chemical versus ion-assisted mechanisms results in aspect ratios less than 2:1. Rates are generally increased by increasing the power density, the RF frequency, or pressure and by decreasing flow rates to increase the residence time of the reactants. The increase in rate is due to a higher density of etchant species such as atomic Cl, Br, or H. Higher etch rates tend to increase surface roughness. The etch rate often decreases with time from its initial value [I]. Cl2 provides the highest etch rates, but a factor of 2 - 3 variation in rate from run to run has been reported. The higher rates can produce surface roughness and pitting [2]. Etch rates with CCl4 are very reproducible (±10%) but polymer deposition occurs [2]. Etch rates with HCl are slower but surfaces are smooth with no polymer deposits [3,4]. Rates can be increased by adding O2 to Cl2, CCl4, or CCl3F [4,5] because the addition of O2 increases the Cl atom density. Polymer formation is also decreased. Maximum rates are achieved with O2 fractions between 40 and 60%: above this optimum concentration oxide formation reduces the etch rate. C

PROFILES

Preferential crystallographic etching is observed with (111)As > (100) > (110) > (111)Ga [I]. No difference in etch rate between n-GaAs and p-GaAs has been reported. Chlorine atoms are the most common active species but Br and H have also been used (TABLE 1). Sources of Cl include Cl2, chlorocarbons, PCl3, HCl, and COCl2 [3,5]. Chlorocarbons can deposit polymeric surface films which diminish etch rates. Cl-based etches all produce surfaces rich in Ga [6]. The etch rate with Cl2 increases with temperature with an Arrhenius-type dependence at a given power density [6]. The measured activation energy depends on power density, however, indicating that the reported activation energies do not reflect a simple thermally activated chemical process. In contrast, etch rates with chlorocarbons such as CCl4 decrease with increasing temperature. This has been attributed to a temperature-dependent morphology of the deposited chlorocarbon film, which is believed to control rates by controlling reactant and/or product diffusion [2].

TABLE 1. Plasma etching rates for GaAs. Etchant

Rate (nm/min)

Cl2 2.5 >80 0.5

Pressure (torr)

Flow (seem)

Temp. (K)

Freq. (MHz)

0.5 0.3 0.3

20 40 40

335 345 373 523

13.56 13 13 2.45GHz

0.25 0.25

Power (W/cm2)

Ref.

Comments

[3] [6] [6] [H]

pitting, 5 W to plasma

ECR plasma

5.8 37-60 20-70

0.05 0.3 0.15-0.3

10 30 30

573 373 373

0.055 14 0.1-14

0.22 0.10 0.1-0.5

[12] [1] [1]

400ZoO2ZCl2

7.3

0.05

6(O2) 10(Cl2)

573

0.055

0.22

[12]

6O0ZOCI2ZCCI4

34.5-3.3 3-0.4 1.7 0.4

0.16 0.16 0.16 0.16

10 10 10 10

573 573 440 700

0.055 0.055 0.055 0.055

0.7 0.7 0.7 1.8

[4] [4] [4] [2]

40%OJ/CC14

2.5

0.05

6(O2) 10(CCl4)

573

0.055

0.22

[12]

4O0ZoO2ZCCl4

12.3-14

0.16

10

573

0.055

0.7

[4]

slows as polymer deposits

CCl3F

2.2-0.7

0.16

150

573

0.055

0.22

[5]

f(residence time)

580ZOO2ZCCIJF

9.9-5.5

0.16

150

573

0.055

0.22

[5]

f(residence time)

CCl2F2ZHe

1.1

0.04

13.56

0.35

[13]

C2Cl3F3

0.04 0.5

0.5 0.04

HCl

0.06 0.06

0.5 0.24

0.12

0.1

Br2

H2

50 110

depends on time

slows as polymer deposits slows as polymer deposits stainless steel electrode

[14] [13]

613

13.56 13.56

335 403

13.56 13.56

6

[3] [9]

smooth, 50 W to plasma downstream from plasma

<423

30

2-10 W/cm3

[7]

tube reactor

D

SELECTIVITY

GaAs generally etches faster than its oxide. Addition of H2 to HCl or CCl2F2 decreases absolute etch rates while increasing the oxide:GaAs etch rate ratio [3]. Pure H2 plasmas etch oxide faster than GaAs and do not produce Ga-rich surfaces [7]. As a result, these plasmas are useful for insitu surface preparation prior to MBE or MOVPE growth. However, oxide growth replaces etching if base pressures exceed 10"6 torr or the H2O partial pressure exceeds 10'7 torr [7,8]. E

DAMAGE

The absence of high-energy ion bombardment results in little, if any, ion-related damage from the etching process [9]. Atomic hydrogen generated from H-containing etchants can neutralize Si donors and reduce carrier concentrations, sometimes forming a highly resistive surface layer [1O]. However, carrier concentrations can be restored by a short thermal anneal at 430 0 C [1O]. Electrode material influences film deposition. Graphite, Al, or Al2O3-coated electrodes are preferred over stainless steel. The latter produce deposits of Fe and Ni chlorides on the etched surface [4]. ACKNOWLEDGEMENTS This work was performed at Sandia National Laboratories and supported by the U.S. Department of Energy under Contract No. DE-AC04-94AL85000. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] II1] [12] [13] [14]

D.E. Ibbotson, D.L. Flamm, V.M. Donnelly [ J. Appl. Phys. (USA) vol.54 (1983) p.5974 ] R A Gottscho, G. Smolinsky, R.H. Burton [ J. Appl. Phys. (USA) vol.53 (1982) p.5908 ] G. Smolinsky, R.P. Chang, T.M. Mayer [J Vac. Sd. Technol. (USA) vol.18 (1981) p.12 ] G. Smolinsky, RA. Gottscho, SM. Abys [ J. Appl. Phys. (USA) vol.54 (1983) p.3518 ] RH. Burton, CL. Hollien, L. Marchut, S.M. Abys, G. Smolinsky, RA. Gottscho [ J. Appl. Phys. (USA) vol.54 (1983) p.6663 ] V.M. Donnelly, D.L. Flamm, CW. Tu, D.E. Ibbottson [ J. Electrochem. Soc. (USA) vol. 129 (1982) p.2533] R.P.H. Chang, C C Chang, S .Darack [ J. Vac. Sci. Technol. (USA) vol.20 (1982) p.45 ] CW. Tu, R.P.H. Chang, A.R. Schlier [ J. Vac. Sci. & Technol. A (USA) vol. 1 (1983) p.637 ] D.G. Lishan, H.F. Wong, D.L. Green, E.L. Hu, J.F. Merz, D. Kirillov [ J. Vac. Sci. Technol. B (USA) vol.7 (1989) p.556-60 ] J. Chevallier, W.C. Dautremont-Smith, CW. Tu, S.J. Pearton [ App.Phys. Lett. (USA) vol.47 (1985) p.108-10] S. Sugata, K. Asakawa [ J. Vac. Sci. Technol. B (USA) vol.5 (1987) p.894-901 ] RH. Burton, G. Smolinsky [ J. Electrochem. Soc. (USA) vol. 129 (1982) p. 1599 ] K. Hikosaka, T. Mimura, K. Joshin [ Jpn. J. Appl. Phys. (Japan) vol.20 (1981) p.L847 ] GD. Kuznetsov, E.M. Novikova, A.V. Zhuravlev [ Inorg. Mater. (USA) vol.24 (1988) p.601-4 ]

18.5 Ion-beam milling and sputter etching of GaAs CLH. Ashby March 1995

A

INTRODUCTION

Ion beam etching, sputter etching, or ion beam milling are purely physical processes based on momentum transfer from an energetic ion or neutral to surface atoms. Atom ejection can occur when the transferred momentum exceeds the atomic binding energy (10 eV). B

RATES

Sputtering rates exhibit a nonlinear dependence on ion energy [1] for energies over 100 eV. Between 100 and 40 eV (near-threshold), the log of the normalized sputtering rate is proportional to ion energy [2]. Rates are generally linearly dependent on ion current density. A pronounced dependence of etch rate on ion/neutral incidence angle is seen, with maximum etch rates near an incidence angle of 60 degrees, where surface atoms acquire a large momentum component directed away from the surface [3,4]. This angular dependence is responsible for mask and sidewall faceting, substrate trenching, and nonequal etch rates for sloped or bilevel surfaces [3]. C

PROFILES

There is little selectivity between materials, with sputter yields for a wide variety of materials ranging between 0.5 and 2.5 atoms/ion, so significant mask erosion occurs during sputter etching. Redeposition of sputtered material onto the etched sidewalls decreases the accuracy of pattern reproduction. Steeper sidewalls increase the problem of distorted pattern transfer [3]. Consequently, ion beam etching is not suitable for high aspect ratio structures. Use of thin metal masks and sample rotation decrease mask-erosion-derived problems [3]. With focused-ion-beam (FIB) sputtering, redeposition causes a V-shaped profile rather than the expected U-shaped profile that replicates the beam profile [5]. D

SELECTIVITY

Some selectivity occurs in sputter etching AlGaAs with Al mole fractions between 0.1 and 0.8 at ion energies between 200 and 1000 eV [6]. Sputtering rates depend exponentially on Al mole fraction, with higher Al fractions yielding lower rates. These can be attributed to the lower sputtering rate of the Al oxide that forms in a non-load-locked chamber. E

DAMAGE

Significant alteration of electrical and optical characteristics occurs under ion bombardment, even at low ion energies. Ar-ion sputtering produces three deep level defects with the following threshold voltages: 60 V for the 0.31 eV trap; 40 V for the 0.45 eV trap and 20V for the 0.58 eV trap [7]. Damage is therefore expected under virtually all sputter etching conditions.

TABLE 1. Ar+ ion beam etching rates for GaAs. Rate (nm/min/

Energy (eV)

Incidence angle (°)

Ref

Comments

0.10

500

0

[4]

0.15

500

30

[4]

0.11

500

60

[4]

0.04

500

75

[4]

0.01

100

0

[1]

0.08

500

0

[1]

0.10

1000

0

[1]

0.09

1500

0

[1]

0.09

500

60

[1]

0.18

400

0

[14]

0.25

500

0

[17]

0.002

40

0

[2]

diode n = 1.25

0.01

80

0

[2]

n = 1.4-1.5

0.03

100

0

[2]

n = 2.3-2.7

2

mA/cm )

Significant degradation of electrical and optical characteristics occurs under ion bombardment. Most effects are more severe at higher ion energies. Damage can be reduced by using lower ion energies and using heavier ions, which reduces the damage depth [8-10]. Ion incidence angles that favour axial channelling of ions into the bulk increase the depth of damage and the amount of damage at greater depths [10]. For n-GaAs the surface barrier, breakdown voltage, and photoluminescence intensity decrease while effective doping concentrations increase [1,11,12]. For p-type GaAs, Schottky barrier heights increase [12]. Thermal annealing can remove damage from low energy ions but damage from ions with energies of 500 eV or more is not removed [1,11,13]. Electrical properties are degraded due to amorphization of the surface [1,9,13] and due to diffusion of beam-induced damage to depths up to 10 times deeper than the amorphous region [I]. The degradation of electrical properties can generally be explained by formation of donor-like traps [7,8,11,14,15], but acceptor-like traps have also been invoked [12]. Sputtering depletes surface As, with the depletion increasing as ion energy increases [16]. Although sputtering lowers n-type Schottky barrier heights in general, the barrier height after 3 keV sputtering was higher than after 1, 4 or 5 keV sputtering due to formation of a donor-like surface damage layer, which correlates with formation of a surface layer rich in As vacancies [16]. Low-energy (< 200 eV) sputtering has been attempted to reduce damage effects. However, Ar+ ions with 50 eV energy still depress electron sheet concentrations and mobilities in GaAs/AlGaAs heterojunction structures after as little as a 15 s exposure; the lighter He+ ion produces even greater damage [15]. However, prolonged exposure to 50 eV Ar+ increases sheet carrier

Next Page

concentrations due to creation of a donor-like damage layer on the surface [15]. Schottky barrier heights decrease on n-type GaAs and increase on p-type GaAs exposed to 0 to 200 eV Ar+ etching [12]. Recovery of initial barrier height by thermal annealing is progressively less as the ion energy increases [12]. A 25% increase in diode ideality factor occurs even with 40 eV Ar+ etching without thermal annealing [2], but a 50 eV Ar+ etch followed by a 400 0 C anneal can give quasiideal diodes [12]. ACKNOWLEDGEMENTS This work was performed at Sandia National Laboratories and supported by the U. S. Department of Energy under Contract No. DE-AC04-94AL85000. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] II1] [12] [13] [14] [15] [16] [17]

M. Kawabe, N. Kanzaki, K. Masuda, S. Namba [ Appl. Opt. (USA) vol. 17 (1978) p.2556 ] J.-Z. Yu, T. Hara, M. Hamagaki, T. Yoshinaga, Y. Aoyagi, S. Namba [ J Vac. Sd. Technol. B (USA) vol.6 (1988) p. 1626-31] R.E. Lee [ J Vac. Sci. Technol. (USA) vol. 16(1979) p. 164 ] S. Somekh, H.C. Casey Jr. [ Appl. Opt. (USA) vol. 16 (1977) p. 126 ] Y. Ochiai, K.Gamo, S. Namba [J Vac. Sci. Technol B (USA) vol.3 (1985) p.67-70 ] G. Guel, CF. Schaus, K. J. Malloy [ SPIE-The International Society for Optical Engineering (USA) vol.1634 (1992) p.436-8] J.-Z. Yu et al [ Jap. J. Appl. Phys. I (Japan) vol.28 (1989) p.2391-95 ] S.W. Pang, M.W. Geis, N.N. Efremow, G.A. Lincoln [ J. Vac. Sci. Technol. B (USA) vol.3 (1985) p.398] A. Scherer, H.G. Craighead, M.L. Roukes, J.P. Harbison [ J. Vac. Sci. Technol. B (USA) vol.6 (1988)p.277-9] R Germann, A. Forchel, M. Bresch, H.P.Meier [J Vac. Sci. Technol. B (USA) vol.7 (1909) p. 14758] S.W. Pang,GA. Lincoln,R.W. McClelland,PD. DeGraff, MW. Geis, W.J.Piacentini [J. Vac. Sci. Technol. B (USA) vol.1 (1983) p. 1334 ] T. Neffati, GN. Lu, C. Barret [ Solid State Electron. (UK) vol.31 (1988) p. 1335-42 ] SJ. Pearton, U.K. Chakrabarti, AP.Perley, K.S. Jones [ J. Appl. Phys. (USA) vol.68 (1990) p.27608] RA. Powell [ Jpn. J. Appl. Phys. (Japan) vol.21 (1982) p.L170-72 ] CM. Knoedler, L. Osterling, M. Heiblum [ J. Appl. Phys. (USA) vol.65 (1989) p. 1800-02 ] Y.-X. Wang and P.H. Holloway [J. Vac. Sci. Technol. B (USA) vol.2 (1984) p.613-9 ] M.W. Geis, GA. Lincoln, N. Efremow, WJ. Piacentini [J Vac. Sci.Technol. (USA) vol.19 (1981) p. 1390-3]

18.6 Reactive ion etching (RIE) and magnetron-enhanced reactive ion etching (MIE) of GaAs Previous Page CLH. Ashby March 1995

A

INTRODUCTION

Conventional reactive ion etching (RIE) is the most common pattern delineation process for GaAs. The substrate is placed on the driven electrode of an RF parallel plate reactor and is etched by neutrals and ions whose energies are greater than the positive space-charge potential of the plasma. In general, chemical etching with its associated high rates, low anisotropy, and low damage is dominant at high pressures and low sheath voltages. Sputtering is dominant at low pressures and high sheath voltages and increases both anisotropy and damage while reducing selectivity. RIE conditions are selected to achieve an acceptable balance between etch rate, selectivity, anisotropy, and damage; some trade-off is always required. There are three common types of R E : conventional RF-only RIE, magnetron RIE (MDE), which uses a confining magnetic field to increase ionization efficiency, and ECR-RF RIE, which uses electron-cyclotron-resonance (ECR) high-density plasmas and RF bias control and is reviewed separately [I]. B

RF-ONLY RBE

Etch rates for RF-only RIE under a variety of conditions are listed in TABLE 1. Rates and optimum conditions determined in one chamber may not translate directly to another. The most commonly employed processes are based on Cl2ZBCl3 or SiCl4. Methane/H2 processes are a common choice when In-based materials are involved. Freon 12 (CCl2F2)-based processing is no longer important due to ozone depletion concerns. It is included in this Datareview for instructive purposes only. Bl

Etch Rates

The fastest etch rates are generally obtained under conditions of high pressure and low self-bias voltage where chemical contributions dominate etching. Increasing the relative importance of chemical contributions relative to ion-assisted (sputtering) contributions reduces residual damage and can sometimes reduce surface roughness [2]. However, greater anisotropy and smoother surface morphology generally result from significant sputtering contributions, which correspond to lower rates. Pronounced dependence on self-bias voltage indicates the importance of ionassisted (sputtering) mechanisms in a specific process. B1.1

Effects of pressure

As pressure increases, there is an interplay between decreasing ion energy due to decreasing selfbias voltage and increasing reactant concentrations. Consequently, at a fixed power density, rates typically increase with pressure due to increasing reactant supply, reach a maximum, and then

decrease with further increase in pressure due to decreasing ion energy [3-6]; the transitional pressure depends on the specific etchant or mixture. At constant bias voltage, etch rates increase linearly with pressure [5] for non-polymerizing etchants, although greater polymerization at higher pressures can decrease rates [7]. At low pressure, etch rates increase with increasing RF power due to the increase in self-bias voltage [2,8], but so does surface damage. At low pressure, the increase in etch rate with increasing RF power may primarily be due to substrate heating if the substrate and holder are not in good thermal contact [9]. With good thermal contact, the rate can be relatively independent of RF power at low pressure (10 mtorr) while increasing the pressure can increase the dependence of the rate on RF power [9]. B1.2

Effects of flow rates

Rates first increase due to higher etchant availability, then reach a plateau with increasing flow rate [3,10]. Further flow increases may reduce rates [2] as the residence time of reactive species decreases with increasing flow rate. Reduction of rates due to polymerization is also flow-rate dependent at fixed ion energy [11-13]. B1.3

Effects of time

Rates can decrease with time as aspect ratios increase during the etching of very small features [2] or vias [10,14-19] due to reactant/product diffusion limitations. For Cl2/BCl3/Ar or CCl2F2/BCl3/Ar, etch rates become nonlinear for etch depths greater than 20 microns [6]. The maximum achievable via depth increases with increasing Cl2 flow and higher RF power [18], and for long etch times the maximum depth is time-independent [18]. B1.4

Effects of surface area

Loading effects resulting from either the total surface area of GaAs or the areal ratios of GaAs to masking material can appreciably change rates. Increasing the total surface area can lower etch rates and alter apparent activation energies [20]. Masking effects can produce either lower (CH4ZH2 [21]) or higher (SiCl4/Cl2 [19]) rates depending on the relative reactivity of the etchant with GaAs and the masking material. B2

Profiles

More vertical walls are achieved by increasing the importance of ion-assisted processes by increasing the bias voltage or RF power density and by decreasing the pressure for a given etchant mixture. Etching profiles can be changed from overcut to vertical to undercut by chemical contributions at higher pressures [4,6,22]. Increasing the sputtering contributions by changing the gas mixtures, e.g., increasing the fraction of Ar or BCl3 [6], increases the anisotropy. Polymer deposition on sidewalls can also increase anisotropy. Either deposition of a film resulting from etchant reactions or redeposition of photoresist can improve the anisotropy and smoothness [10]. Equirate etching of GaAs and AlInGaP with smooth, >80° sidewalls is possible with SiCl4/CH4/Ar due to polymer deposition [23]. Addition of He to SiCl4/CF4/O2 reduces the anisotropy by reducing the sidewall polymer deposition [24]. Smooth via walls can be obtained with photoresist masks when they provide a protective sidewall coating [10,18] whilst vias can be rough and re-entrant with SiNx masks [10]. Wall angles of 80

to 85° are ideal for vias and such steep walls are obtained at high power and low Cl2 flow with SiCl4/Cl2 [18]. Higher power and lower pressure produce smoother walls [18]. Higher Cl2 permits deeper vias but decreases the selectivity to the photoresist [19]. Rougher walls due to crystallographic etching and re-entrant profiles can also result [19]. The source-gas ratio of ICl2ZSSiCl4 provides a good compromise between sidewall profile and etch rate [18]. Increases in RF power can slow the etch rate by sputtering photoresist into the via [19]. Alternatively, increases in the sputtering rate at higher RF power can increase the etch rate. It has been reported that etch rates are lower for vias with initial dimensions less than 100 x 100 ^m due to diffusion limitations and greater problems with PR sputtering and redeposition [19]. In contrast to the normally anisotropic profiles produced by common RIE conditions, isotropic etching of overpass microstructures at undercut rates as large as 400 nm/min have been obtained using 300 mtorr of Cl2/BCl3/Ar [25]. B3

Morphology

Although polymer deposition increases the wall smoothness, it can roughen the surface by acting as a micromask in the field, producing 'grass'. Mask sputtering and redeposition can also cause 'grass'. Higher power densities and higher pressures can increase polymerization [6], as can a higher fraction of the polymer-forming component in a gas mixture. Depending on specific conditions, a higher power density can also decrease grass by minimizing redeposition. B4

Selectivity

B4.1

Halogen-based etches

In RIE-etching of GaAs/AlGaAs structures, either equirate or highly selective etching may be desired. Equirate etching with Cl-based etchants is possible in load-locked systems that exclude H2O and O2 to prevent the oxidation of Al that inhibits etching [26-28]. Source gases that serve as oxygen (O2 or H2O) getters, such as BCl3 and SiCl4, facilitate equirate etching [6]. With SiCl4, equirate etching occurs for self-bias voltages exceeding 100 V while AlGaAs etches faster than GaAs at lower voltages [29] or higher Al fractions [30]. Selectivity for GaAs versus AlGaAs can exceed 10,000:1 at low SiCl4 flow rates and low self bias [31]. The use of F-containing etchants causes the formation of nonvolatile AlF3 and gives high GaAsto-AJGaAs selectivity [24,32-38]. By changing the SF6:SiCl4 ratio, one can vary the GaAsAlGaAs etch-rate ratio from 1:1 to 500:1 [34,38], the selectivity being greatest at lower self-bias (60 V) [38]. Adjusting parameters to increase the importance of chemical versus ion sputtering effects usually increases the selectivity and is achieved by using higher pressure, higher flow rate, or lower self-bias voltage [35,36]. The selectivity for GaAs versus AlGaAs increases with increasing Al mole fraction [36] if residual oxygen is present. AlGaAs can be used as an RIE etch-stop layer with F-containing gases, the thickness required depending on Al mole fraction [33,37]. However, ion-induced damage below the etch stop may be a problem. B4.2

Alkane-based etches

Under most alkane/hydrogen etching conditions, Al-containing materials generally etch more slowly than GaAs [39-42] and In-containing materials etch faster [42-44]. In methane-based

processes, the GaAs/AlGaAs etch-rate ratio is controlled by the CH4 concentration [39]. AlGaAs of any Al fraction can serve as an etch stop for RF powers below 0.2 W/cm2 [39]. With C2H6ZH2, AlGaAs etch rates decrease with increasing Al fraction [40]. Betwen 150 and 250 0 C, the GaAs etch rate increases 2-fold while AlGaAs etch rates remain constant [41]. Equirate etching of GaAs and In-containing materials has been achieved with SiCl4/CH4/Ar but sidewall angles differ slightly [23]. B5

Damage

B5.1

Surface damage

At ion energies less than 100 eV, damage is acceptable for most applications [45,46], although damage will occur at any energy above the 40 eV atomic displacement threshold of GaAs. Even a 2 min. exposure to 60 eV energy ions produces threshold voltage shifts and transconductance decreases in MODFETs [38]. Excellent DC and RF characteristics are obtained for 0.1 micron recessed gate GaAs MESFETs when the DC bias is kept at 80 V with CH4ZH2 etching [47]. However, undercutting can develop at these low ion energies and polymer formation may be a problem with some etchants [20]. GaAs MESFETs etched with SiCl4 using DC biases between 100 and 300 V are 'damage-free'. However, 200% over-etching with SiCl4/SF6 using an AlGaAs etch stop decreases the source saturation current due to active layer depletion [48]. More damage-sensitive HEMT devices, which operate at lower bias voltages, are completely damaged with a DC bias of 200 V and a 200% over-etch [48]. Point defect creation in the near surface leads to carrier compensation up to 0.22 \xm deep with SiCl4 at 100 V bias [49]. If the bias voltage is less than 50 V, no degradation of Schottky diode or optical properties is observed [49]. At voltages over 200 V, deep-level electron traps form, but most of these can be annealed out [50,51]. Etching with SiCl4 at 310 V primarily produces point defect-derived electron traps at 0.30, 0.42, 0.64 and 0.86 eV below the conduction band and EL2 (0.8 eV) [51]. Ion energies over 300 eV may produce permanent degradation of electrical properties [52,53]. Unfortunately, rapid thermal annealing (RTA) can cause defect propagation to greater depths [53]. Damage can be minimized by minimizing the ion penetration by reducing ion energy or using heavier ions, which reduces the ion penetration at a given energy [54,55], and by increasing the chemical contribution to etching. Higher etch rates leave less apparent damage and tend to a steady state value for etch times greater than the ion-energy-related damage depth / etch rate [51]. A prolonged etch at a lower self bias voltage can produce less damage than a faster short-time etch at higher voltage [56]. A short-time etch with low-energy ions can remove surface damage previously produced by rapid etching with high energy ions [57]. A three-step process has been proposed consisting of regular RIE, a higher RF-power step to reduce surface roughness, and a brief high pressure, low power etch to eliminate damage [58]. Lower ion energies with high ion fluxes have been achieved with magnetron-enhanced RIE (MIE) [59-61]; this yields high etch rates with minimum damage, as discussed below. High surface ion densities (101V cm3) at the GaAs surface are possible at low gas pressure (<1 mtorr) with a magnetically confined DC plasma when an AC voltage is applied to the sample to provide ion bombardment [52]. B5.2

Sidewall damage

Sidewall damage is especially important for small structures such as thin conducting channels and

quantum wires. Even though top-surface and sidewall fluxes are often similar, the sidewall damage depth is generally less due to the lower ion energy of the sidewall flux, which consists mainly of reflected primary ions, sputtered material, and chemical reaction products from the adjacent flat surface [51]. Conditions for minimizing sidewall damage include low ion energy, highly collimated beams, reactive gas etching, and efforts to minimize contamination, such as polymer coated sample holders [62]. Loss of etch directionality and slow vertical etch rates can occur in submicron trenches (< 0.25 ^m) when low energy ( 200 eV) ions are attracted to the walls by the ion-induced image potential [63]. The depth of sidewall damage increases with etch time in RIE [64,65]. Ion-neutral collisions in the sheath region increase the angular distribution of ions and result in significant ion flux on the etched walls [63]. Magnetron-enhanced RIE (MIE) generates a high ion density at low bias voltages and low presssures by confining electrons close to the powered electrode with a magnetic field. Although this produces fast rates with less damage, some undercutting can occur [59]. Adding inert gas to the Cl2 increases the sidewall damage [62]. RIE at 300 V with SiCl4 produces greater sidewall damage than surface damage due to lower sidewall etching rates [64]. However, a CH4ZH2 mixture produces less sidewall damage than SiCl4 [56]. B6

Common Process Chemistries

B6.1

Cl-based chemistries

B6.1.1 Chlorine and chlorocarbons The highest etch rates are achieved with Cl2 as the Cl source, but Cl2 produces surface roughness and exhibits low reproducibility in etch rates from run to run. Reproducibility can be improved by an initial short etch in H2 or CCl2F2 [66]. Chlorocarbons such as CCl4 and CCl2F2 produce smoother surfaces but can deposit chloropolymer films. Etching at higher power densities can lead to increased polymer deposition [2] unless offset by increased sputter desorption. The etch rate may [2,67] or may not [8] depend on impurity doping type and level. The addition of O2 to CCl4 or CCl2F2 increases the etch rates by increasing Cl atom densities [2,4]. Etching with CC12F2/O2 at greater than 380 V leads to preferential loss of Ga since thermal volatilities of reaction products are not important at high self-bias voltages [2]. The addition of Ar to Cl2, CCl4, or CCl2F2 increases anisotropy [3,66,68] by increasing sputtering. The addition OfH2 to CCl4, CCl3F, or CCl2F2 increases anisotropy, decreases apparent damage, and gives a residue-free surface [8,69]. Adding H2 to Cl2 in Ar reduces the etch rates but increases both surface smoothness and anisotropy [70]. A mixture of CH4 and Cl2 in Ar produces a rapid etch rate with very smooth surfaces and good anisotropy [70]. However, H-containing source gases raise concerns about dopant passivation and subsequent annealing requirements [68,71]. Etching with C12/BC13 is slower than with Cl2 alone but its better morphology, good anisotropy, and greater reproducibility [6,10,72] has made it a popular process chemistry. B6.1.2 Silicon halides SiCl4-based etching generally produces better surface morphologies, greater anisotropy, and less change in surface stoichiometry than the analogous Cl2-based process [49]. Etch rates are a factor of 6 slower with SiCl4/Ar than with Cl2/Ar, but SiCl4 produces little As loss whereas Cl2-etched surfaces are As deficient [49]. There is also less Cl-containing residue on the SiCl4-etched surface

[49]. However, Si incorporation has been seen in the near-surface region after 250 V bias SiCl4 etching. Secondary ion mass spectroscopy shows 1 x 1019/ cm3 of Si at the surface decreasing to below the 101V cm3 level at 30 nm below the surface [73]. More damage (a greater reduction in PL intensity) was observed for p-type than n-type GaAs [73]. Crystallographic etching occurs with SiCl4 at intermediate pressures (20 mtorr) and low power density. Addition of Ar removes the orientation dependence by increasing sputtering [22]. With SiCl4 alone, vertical, smooth etching occurs at low pressure (< 30 mtorr) and self-bias over 100 V [30]. Thus, SiCl4/Cl2 mixtures are good for etching vias [18,19]. With SiCl4 alone, AlGaAs/GaAs etching is equirate for Al fractions from 0 to 0.3, AlGaAs etching slightly faster (<2-fold) for Al fractions between 0.5 and 0.9 [30]. Very low damage, highly selective etching of GaAs versus Al03Ga07As occurs at sufficiently low flow rates that residual O can be significant [31]. Selectivity for GaAs exceeds 10,000:1 with 4 - 6 seem SiCl4 and low power ( 0.066 W/cm2) [31]. Such selectivity is lost at 15 - 20 seem and higher power densities. Emission spectra show differences in the dominant plasma species in these two selectivity regions [31]. The addition of F-containing gases like SF6 [34,35], SiF4 [38], and CF4 [24] provides high etch selectivity versus AlGaAs. These mixtures replace Freon-based chemistries, abandoned due to ozone depletion concerns. Mixtures of SiCl4 with alkanes, like CH4 [23], can etch structures comprised of both GaAs and In-containing compounds. B6.2

Alkane-based chemistries

Alkane (CH4, C2H6, or C3H8) / hydrogen etching produces smooth, anisotropic etching of most ni-Vs. Etch rates are appreciably slower than Cl-based processes. Significant polymer deposition occurs and regular O2 plasma cleaning of etching chambers is necessary even with optimized etching compositions. Etch rates initially increase as the alkane fraction is increased. Further increases in alkane fraction lead first to surface polymer films and then to gas-phase polymerization [7,11]. Polymer deposition occurs with RF powers <0.2 W/cm2 [7] and is a greater problem at higher temperatures [7]. The apparent activation energy for CH4ZH2 etching ranges from 0.7 eV in H2 to 6.5 meV at high CH4 partial pressure [74]. This near-independence of T at high CH4 partial pressure permits highly reproducible deep mesa etching [74]. Addition of noble gases alters these etch rates: etch rate (Ar) > Etch rate (H2 only) > Etch rate (Ne) > etch rate (He) [7,12]. Heavier ions (Ar) delay the onset of polymerization and improve surface morphology [7,75] but the chemical effect of H2 dominates over mass effects on polymer deposition [7]. The chemical contributions ofH2 produce less scatter in etch rates than with inert ions [12]. The threshold Ar flow for preventing polymer deposition increases with increasing CH4 flow [12]. Increasing H2 first increases etch rates by decreasing polymerization and then decreases these rates due to dilution. The optimum alkane/H2 ratios depend on the alkane with propane < ethane < methane [7,11]. In general, a CH4ZH2 ratio between 0.2 and 0.4 [76] and a C2H6ZH2 ratio < 0.4 [40] give polymer-free etching. Optimum CH4/H2 ratios are higher at lower total flow rates [13]. Increasing the number of C atoms in the precursor increases the optimal total H2 flow rate but decreases the optimal alkane flow [7,11]. Virtually all masks, including NiZCr and AuZPd, react with CH4ZH2 [21]. Plasma-photoresist

reactions significantly affect GaAs etch rate linearity above 67 0 C [21]. Hydrogen passivation of dopants necessitates post-etch thermal anneals at temperatures near 400 0 C to restore electrical activity [56,71,77]. Decreases in surface free carrier concentrations depend inversely on the initial n-type doping level. The reactivation energy for recovery of donor activity is about 1.75 eV and a 3600C anneal for 5 min gives almost complete reactivation of dopant [78]. The depth of Si dopant passivation by H increases with increasing substrate temperature during etching [41]. With C2H6ZH2 at 410 V self-bias, almost 5000C is required for restoration of initial carrier concentrations. This is due to the combined effect of H dopant passivation and the damage caused by deep penetration of the light ion [77]. B6.3

Mixed Cl-alkane chemistries

Chloromethane yields higher etch rates than methane, or alternatively, the same etch rate at lower self bias [44,79]. Etch profiles are smooth and anisotropic at lower power densities but roughen at about 1.1 W/cm2 [44]. Replacement of CH4/Y with C1CH3/Y produces the same dependence of etch rate as the ionization efficiency of Y: Ar>H 2 >O 2 >Ne»He for IClCH3: 2.5Y [44]. With C1CH3/H2, the same etch rate relationship obtained with CH4 (Al
MAGNETRON RIE (MIE)

In magnetron RIE (MIE) a magnetic field confines electrons close to the powered electrode to increase ionization efficiency. The higher electron density resulting from the reduced loss of electrons from the discharge region to the walls produces higher dissociation efficiencies. This reduces the induced bias on the sample electrode compared to conventional RIE. The higher density plasma discharge resulting from magnetic electron confinement can sustain itself at lower cathode bias voltages than in RF-only RIE. In MIE, self bias increases linearly with power density. Magnetron discharges have very thin sheaths, so there are very few collisions in the sheath. Therefore, the self-bias may be a good indication of the actual ion energy; this is not necessarily so in higher pressure plasma etching [80]. The dominant etching chemistry may actually be different in MIE than in RIE due to the different plasma species concentrations resulting from higher electron densities and lower energies in MIE. Fast rates are possible and vias have been etched using MIE with SiCl4, Cl2, and BCl3 [81]. Cl

MIE with Halides

MIE with BCl3 produces smooth morphology with no residues. Etch rates increase linearly with pressure between 2 and 6 mtorr because, in contrast to RIE, the self bias increases only 6 volts rather than decreasing [61]. MIE rates are almost an order of magnitude faster than with RIE under similar conditions due to the higher dissociation efficiency [61,82]. Dislocation loops appeared relatively benign with respect to the near-surface electrical properties that are important for Schottky diodes [61]. MCE with SiCl4 is nearly a factor of 5 faster than RIE under similar power levels and self-biases

[8O]. MIE rates first increase with power then saturate. The fast MIE rates at low DC bias produce Schottky diodes with relatively low damage after a 1 ^m/min etch [80]. MIE produces greater fragmentation of CCl2F2 and permits less C-F recombination than RIE resulting in less polymer deposits [60]. Reduced polymer deposition negates the polymer-removal effect of O2 in RIE that leads to maximum rates at 20 - 25% O2. MIE shows monotonically decreasing rates with O2 addition due to reactant dilution while the damage increases with O2 addition due to reduced etch rates. For the same gas supply and RF power, MIE produces more than 4 times the etch rate at half the self bias [60]. However, damage is appreciably lower for MIE than RIE due to the lower self bias. Larger (100 A diameter) and deeper dislocation loops form with RIE [2] than with MIE (20 - 40 A) due to both higher self bias and lower etch rates [83]. Higher MIE powers produce smaller dislocation loops at shallower depths presumably due to higher etch rates reducing damage accumulation [83]. Freon MIE produces surface As depletion, but morphologies are smoother than with RIE [83]. C2

MIE with CH4

GaAs etch rates with CH4ZH2ZAr MIE increase as CH4 is increased from 10 to 35%, then decrease with further CH4 increases due to increasingly dominant polymer deposition [84]. Similarly, AlGaAs etch rates decrease above 30% CH4. The morphology is anisotropic and smooth. Etch rates are slow without added Ar. Increasing the Ar up to 50% increases rates, while further increases decrease rates as physical sputtering becomes the dominant mechanism [84]. Extremely smooth, vertical sidewalls with slight trenching are produced with 10% Ar; surface smoothness is also enhanced by Ar [84]. Argon addition also increases the etch-induced defect density and depth [84]. MIE self-bias is up to 4 times lower than RIE under similar plasma conditions [7,84]. Schottky diode ideality factors and barrier heights are near those of unetched controls [84]. A 3000C, 30 s RTA redistributes the H-passivation, decreasing in the top 0.3 fim but increasing at greater depths. Residual H-effects remain even after an anneal at 400 0 C for 30 s [84]. ACKNOWLEDGEMENTS This work was performed at Sandia National Laboratories and supported by the U.S. Department of Energy under Contract No. DE-AC04-94AL85000. REFERENCES [1] [2] [3] [4] [5] [6] [7] [8] [9] [10]

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TABLE 1. Reactive ion etching rates for GaAs. Etchant

Rate

(nm/min)

Pressure (mtorr)

Flow (seem)

Power density (W/cm2)

0.44 0.8

Selfbias (-V)

Ref

Comments

[14] [85] [66] [62] [26] [86]

via drilling mounted with Ag grease vertical walls smooth, no sidewall damage load-lock:equirate vs AlGaAs>350V low damage, semi-isotropic

RF-onlyRIE 4.5 3.0 1.4 0.45 0.19 0.15

150 85 5 5

3-23

2

40

400 270 250 400 150-160

ICl2: 1 He

0.11

5

30

250

[59]

anisotropic

5C12:3 Ar

0.09

5

20

0.14

300

[28]

equirate vs. AlGaAs and vertical profile

4Cl2: I A r

0.80

5

200

200

[66]

severe undercut

2C12/1 Ar

0.80

5

100

[49]

As deficient

1C12:4AT

0.27

200

[66]

vertical walls

3Cl2: IAr

3.4

15

12

0.75

330

[68]

isotropic

17Cl2:83Ar

0.9

15

12

0.75

330

[68]

anisotropic

29C12:21CH4: 50Ar

0.6

0.45

140

[70]

anisotropic, smooth

17Cl2: 83BCl3

0.6

60

0.42

270

[72]

high reproducibility

17Cl2: 83BCl3

1.1

150

0.42

270

[72]

high reproducibility

1C12/1BC13/3AT

0.9

15

100

300

[10]

very smooth bias if PR mask, 100C

lCl2/4BCl3/4Ar

1.0

40

20

[6]

0.1 urn roughness

Cl2

1.05

TABLE 1. Reactive ion etching rates for GaAs (continued). Ref

Comments

0.52

[6]

>200:l vs. AL^Ga^gAs

0.20

[82]

very smooth

445

[87]

f(P,bias, BCl3ZHe)

200(385)

[88]

RIE/sputter cycling

300

[27]

equirate vs. AlGaAs profiles, resolution rates f(W,P gases)

0.16

[8]

isotropic, rough

10

0.16

[8]

isotropic, rough

10

0.16

[8]

anisotropic

200

[67]

n+ faster than SI

<100 330

[89] [36] [3]

22:1 vs. AlGaAs

[15] [3]

via drilling max at 60% Ar

Flow (seem)

(nm/min)

Pressure (mtorr)

Power density (W/cm2)

3CCl2F2/3BCl3/4Ar

3.2

80

20

BCl3

0.0017

11

20

32BCl313He

0.090

40

45

BCl3 then Ar

0.075

45(40)

9(10)

1BC13:9AT

0.28

15

0.27

CCl4

1.1

11

10

9CCl4:1O2

1.7

11

1CC14:9H2 smooth

0.45

11

2 CCl4: IAr: 2Cl2F2

0.8

15

CCl2F2

3.3 2.4 0.1210

230 30 32

0.85

0.076

250 100

70 20

ICCl2F2: IHe

1.1

38

19CCl2F2:1O2

0.07

47CC12F2:3O2

2.2

60

ICCl2F2:1O2: 10Ar

0.35

5

Etchant

Rate

2CCl2F2:3Ar

30

20

20

0.10(0.14)

0.33 0.3

Selfbias (-V)

0.6 0.35

200

[32]

1000:1 vs. AlGaAs

0.56

380

[2]

P+ and undoped n+ 25% faster; Ga loss

0.51

200

[16]

via drilling

0.5

650

[4]

smooth

TABLE 1. Reactive ion etching rates for GaAs (continued). Etchant

Rate (um/min)

Pressure (mtorr)

Flow (seem)

Power density (W/cm2)

Selfbias (-V)

Ref

Comments

SiCl4

0.129 0.067 0.4320 0.5810 0.2 0.1 0.25-0.6

10 10 3-15 3-15 14

10 10 0.10 0.10 0.2 4-6

0.81 0.27

600 275 90 120 300 40-70 20-300

[5] [5] [22] [22] [56] [31] [30]

no polymer film no polymerfilm,max at 10 mtorr (100)>(lll)>(011) non-crystallographic

2SiCl4/lAr

0.65

100

100

[49]

little As loss

SiCl4:SF6

0.14

50

60

[38]

490:1 vs. Al035Ga065As

5SiCl4/20(CF4/O2)/ 12He

0.46

10

40

[24]

420:1 Vs-Al01Ga09As

10SiCl4IlCl2

1.7

50

66

180

[17]

vias

9SiCl4/lCl2

1.1

80

40

168

[18]

vias 45 0C; statistically optimized process

25SiCl4/4Cl2

2.0

150

>150

250

[19]

100x100 ^m vias

5SiCl4/4CH4/ 9Ar

0.09

7.5

18

[23]

AlInGaP equirate

8SiCl4/lCH4/ 9Ar

0.2 7.5

18

0.56

[23]

80-90° slope

H2

0.15

5

CF4

0.01

5

2CF4:3Ar

0.02

20

10

C2F6

0.019

15

20

100

0.066 15

18.5

>104/l vs. AlGaAs anisotropic 400C, loadlock equirate vs. AlGaAs

0.12

0.56

0.4

600

[66]

500

[90]

0.55 450

[91]

PR rate 3x GaAs

[53]

damage study

TABLE 1. Reactive ion etching rates for GaAs (continued). Power density (W/cm2)

Selfbias (-V)

Ref

Comments

0.4

1000

[56]

donor passivation

24

0.75

500

[43]

f(t, P, dc bias)

30

48

0.42

580

[13]

f(flow, %CH4) 87° wall, no trench

0.040

10

60

1.1

500

[7,11, 39]

GaAs/AlGaAs as f(Al fraction)

7C2H6:93H2

0.023

100

0.6

[92]

vert, walls, smooth

IC2H6ZlOH2

0.034

44

0.56

[41]

250°C;f(T)

Etchant

Rate (jim/min)

Pressure (mtorr)

1CH4:5H2

0.02

14

1CH4:5H2

0.016

10

1CH4:3H2

>0.045

ICH4: IH2

Flow (seem)

0.036

10

43

1.1

500

[7,11]

6.7C3H8:34H2

0.038

10

43

1.1

500

[7,11]

2:1 GaAs/AlGaAs

20CH4/50Ar

0.060

15

70

1.1

[7,12]

Ar>H2>Ne>He

1C1CH3/2.5H2

0.053

15

42

0.88

480

[44]

5/1 vs. Al0 3Ga0 7As

C1CH3/O2

0.011

15

23

0.88

4000

[44]

11/1 Vs-Al03Ga07As

CCl2F2

1.2 0.188

40 4

20 20

0.53 0.52

40 145

[59] [60]

346 K low damage: n=l .09

4CC12F2/O2

0.158

4

20

0.52

[60]

n=1.29

SiCl4

1.0 0.15

2 1

15

0.24 0.76

100 184

[80,81] [93]

0.25^1 vertical, smooth

BCl3

0.085 0.50

5 2

15 15

0.16 0.8

«30 200

[61] [61]

low damage

Magnetron RIE (MIE)

Next Page

TABLE 1. Reactive ion etching rates for GaAs (continued). Etchant

Rate (pim/min)

Pressure (mtorr)

Flow (seem)

Power density (W/cm2)

Selfbias (-V)

Ref

Comments

1CH4/5H2

0.005

10

30

0.4

«95

[84]

very smooth, vertical

3CH4/6H2/lAr

0.014

10

30

0.4

«95

[84]

very smooth, vertical

lCH4/5H2/6Ar

0.043

10

30

0.4

«95

[84]

very smooth, vertical

CCl2F2

0.8

40

3000

[94]

DC plasma

lCCl2F2:9Ar

5.8

40

3000

[94]

DC plasma

lCCl4:9Ar

0.65

40

3000

[94]

DC plasma

lCF4:9Ar

0.50

40

3000

[94]

DC plasma

ICF4

0.31

40

3000

[94]

DC plasma

DC plasma

18.7 Etching of GaAs using ECR-RF reactive-ion etching

Previous Page

CLH. Ashby March 1995

A

INTRODUCTION

ECR-RF RIE (electron-cyclotron-resonance radio-frequency reactive-ion etching) is a relatively new approach to RIE in which reactive chemical species are generated in a microwave-excited high-density plasma above the sample, which is placed on an rf-driven electrode surface to provide ion bombardment for anisotropic profiles. This configuration provides independent control of reactant generation and sample bias and enables anisotropic etching at much lower ion energies than conventional RF-only RIE. In general, etching rates are higher and damage is lower when conditions favour chemical contributions to etching. In contrast, anisotropy is greater and surfaces tend to be smoother with increased sputtering contributions to etching. As with RF-only RDE, the acceptable balance between etch rate, anisotropy, and residual damage for a particular application determines the best etching conditions. ECR-RF RIE has significant advantages over conventional RF-only RIE for fabrication of devices requiring low-damage processing for optimal performance [I]. In a microwave-excited (2.45 GHz) ECR-RF RIE, an ECR plasma is used to generate reactive chemical species in a region not in direct contact with the sample. The high electron densities in ECR discharges (up to 1012/cm3 [2]) produce higher ionization efficiencies than in simple RFdischarges, which permits lower pressure operation. There is a high density of neutral radicals and ions with energies typically 15 V. It is important to note that the higher ionization efficiency also means that the relative concentration of various chemical species in the plasma and, consequently, the etching behaviour, may be quite different for nominally the same gas conditions in ECR-RF and RF-only RIE. B

RATES

Etch rates increase with increasing electron densities due to increased generation of highly reactive species in the plasma. In both CC12F2/O2 and CH4/H2/Ar plasmas, electron densities and rates have been shown to increase with increasing pressure and RF power [2]. Electron densities also increase with increasing O2 or Ar fraction. However, the concomitant decrease in reactant source gas by dilution counteracts the effect of increasing electron densities [2]. Application of an RF bias at 13.56 MHz to the sample permits selection of the DC self-bias, independent of chemical reactant generation; this permits much greater process control than in RF-only RIE. Etch rates increase linearly with increases in DC self-bias voltage, i.e., RF power [3]. Lower pressures in ECR-RF RIE reduce collisions and permit lower accelerating voltages for anisotropic etching, which results in less residual damage. Comparable etch rates can be achieved with ECR-RF RIE using bias voltages more than a factor of two lower than those employed in RF-only RIE.

C

PROFILES AND UNIFORMITY

Anisotropy is supplied by the RF-driven ion bombardment. As in all ion-assisted processing, there is a trade-off between higher ion energies for greater anisotropy and higher rates, and lower ion energies for minimum damage. ECR-RF RIE is desirable for low-damage etching because it is anisotropic at lower bias voltages (typically < 150 V) than RF-only RIE (typically >250 V). Reduced temperature (-300C) etching provides an alternative way to increase anisotropy. Vertical sidewalk have been obtained with only 50 V bias. A Cl-containing residue, which is removed with a 5 min H2 plasma, is probably playing the protective role often provided by polymer deposition [3-5]. Etch uniformity is best when the magnetic field lines are perpendicular to the surface. This can be achieved without magnetic confinement far from the ECR source or with the sample located within a magnetic bottle that establishes parallel field lines. Application of a magnetic field downstream from the ECR plasma source permits collimation or even focusing of electrons [6]. This can increase ion densities and, concomitantly, increase etch rates. Smooth morphologies and vertical walls have been etched with a self-bias of only -10 V with a downstream electronconfining magnetic field [6]. Increasing the separation of the sample from the ECR source to improve uniformity also increases the sample self bias [7]. This increased separation can also decrease the concentration of reactive species at the surface [7]. The relative importance of ionassisted processes or chemical reactions as the rate limiting step for a given reactant gas mixture determines whether the etch rate increases or decreases as source-sample separation is changed [7]. Different source hardware configurations have been examined for etching characteristics. Comparison of multipolar and magnetic mirror ECR sources with CH4ZH2 etching chemistry has shown them to be similar in terms of uniformity, rate, and anisotropy [8]. D

DAMAGE

ECR-RF RIE is promoted as a low damage process because of the lower ion energies employed to achieve anisotropy than in other ion-bombardment-based processes. Nevertheless, three types of damage must still be considered that can alter electrical and optical properties of devices: 1) damage resulting from ion bombardment, 2) passivation of donor or acceptor atoms by hydrogen, and 3) changes in surface stoichiometry due to either chemical or ion-induced effects. Dl

Damage from Ion Bombardment

Ion-induced damage results from the rf-derived self-bias needed for anisotropy. Ion-related damage will generally be less with greater chemical contribution to etching, lower ion energy, higher ion mass, and shorter exposure time to the plasma, i.e., higher etch rates. Point defects formed in the near-surface region can migrate to depths greatly in excess of projected ion ranges, even including channelling effects. These deep level defects serve as carrier traps and degrade device operation. The lattice damage threshold for GaAs is on the order of 40 eV [3], and it has been proposed that further ion energy reduction below this level may reduce anisotropy with minimal return in improved electrical properties due to effects of electron bombardment, UV illumination, or loss of surface stoichiometry [9]. Even with no applied substrate RF bias,

cathodoluminescence decreases with both H and Ar plasmas [9]. Damage studies using Ar or H2 to separate ion and chemical effects have shown damage to depths of 2 nm with 20 V bias with Ar [10], and up to 100 nm with 300 V bias with H2 or Ar [11]. Damage is deeper but less severe with the lighter ion. Channelling can increase penetration of damage by as much as a factor of 8 times the calculated ion range. Migration of point defects is the probable source of damage at 100-nm depths [H]. Such lattice damage is particularly of concern for Schottky gate-controlled devices. Examples include MESFETs, and HEMTs (high electron mobility transistors) where deep-level traps can trap charge carriers needed for optimum current conduction. Transmission line measurements following CC12F2/He etching show reduced conduction due to increased depletion layer thicknesses [12]. With CH4/H2/Ar, etching with 0 to 40 V bias produces no significant increase relative to wet etching in source-drain resistance at 300 K for GaAs/AlGaAs HEMTs. At 1.2 K, however, the two dimensional electron gas (2DEG) carrier density, n, was decreased by 20% and mobility by 60%, which is nevertheless much better than the decrease in n of up to three orders of magnitude seen with RF-only RIE [13]. At 50 V bias, surface barrier height and diode ideality factors are degraded by only 3% relative to wet etching [3]. For 1017 cm"3 n-GaAs, photoluminescence decreases by only 25% after CCl2F2 or CHCl2F etching at 150 V bias [14] compared to a decrease of more than a factor of 2 with 300 - 400 V Cl2-based RF-only RBE etching [3]. D2

Passivation Effects

Whenever H is a constituent plasma species, passivation of dopants will occur and restoration of dopant electrical activity requires an anneal in excess of 3000C. The H passivates nonradiative traps, increasing both photoluminescence (PL) and cathodoluminescence (CL) intensity [9]. Hpassivation of shallow donors to a depth of 300 nm occurs during a 60 min exposure to a CH4ZH2ZAr discharge with 100-V bias; similar passivation is seen with RF-only RIE [3]. Migration of point defects rather than direct ion damage permits the deep H penetration. [3,9]. D3

Surface Stoichiometry Degradation

Surface stoichiometry changes affect surface electronic properties by changing dominant surface states. Either Ga or As enrichment is possible. Stoichiometric changes usually occur when surfaces are roughened, especially at higher microwave powers [2]. Above 200 W microwave power, both CCl2F2ZO2 and CHCl2FZO2 produce As-rich surfaces over the entire range of pressure, DC bias, and gas flow rates studied [14]. Above 200 W microwave power, these etchants produce an As-rich surface. For 5:17:5 CH4ZH2ZAr with 250 W microwave power and 100 V DC bias, there is no As loss; RF-only RIE with these gases at >300 V bias does cause loss of As, however, [3]. WithHBrZH2ZAr, surfaces are Ga- and AlZGa-rich above 200 0 C [15]. Morphology is smooth with HIZH2ZAr etching from 50 to 250 0 C, but loss of some As occurs in the top 50 A at higher temperatures [15]. Photoresist (PR) degradation can occur at high microwave powers [16]. High H2 partial pressures increase positive PR etching. Without RF-biasing (5 - 10 V self-bias), PR etch rates can exceed GaAs etch rates by more than two orders of magnitude [16]. However, since GaAs rates increase linearly in bias voltage, good selectivity with respect to the photoresist is achieved under typical ECR-RF RIE etch conditions [3].

E

PROCESS CHEMISTRIES

Two main chemistries have been employed to date: halogen-based (mostly chlorine) and alkanebased (mostly methane). In general Cl-based processes are faster than alkane processes, but the alkane processes produce smoother surfaces, especially for In-containing materials [17]. The main chemistries employed are Cl2ZAr and CH4ZH2ZAr. Ozone depletion concerns have made the CCl2F2 chemistries obsolete; they are included in this Datareview for illustrative purposes only. El

Halogen-Based Chemistries

Etching with Cl2ZAr can give smooth, vertical features with 10:1 aspect ratios using low microwave power, low Cl2 concentration, low pressure, and relatively high RF power (60 - 100 W) [7]. At fixed microwave power, higher RF power increases DC self-bias and initially produces faster rates until a plateau value is reached [7]. Increasing microwave power increases dissociation efficiency but decreases DC self-bias [7]. Rates increase with microwave power if etching is limited by reactant supply but decrease if limited by ion-enhanced reactions or product desorption. Higher microwave powers decrease surface quality and decrease anisotropy by increasing chemical contributions while increasing RF power or adding Ar improves morphology and anisotropy by increasing sputtering contributions [7]. Increasing the pressure increases the DC bias slightly, in contrast to RF-only RIE where increasing pressure decreases DC bias [7]. Self bias increases from 155 to 200 V as the pressure increases from 0.5 to 20 mtorr. The opposite occurs in RF-only RIE where self bias decreases as the pressure rises due to the increase in ion collisions [7]. Surfaces roughen at higher Cl2 pressures (rough at 2 mtorr but smooth at 0.5 mtorr) [7]. Increasing the temperature to 380 0 C increases etching rates but also roughens surfaces and increases undercut as chemical contributions to etching become dominant [7]. The dopant type has been observed to influence etch rate, with Cl2ZHe mixtures etching n+-GaAs 50% faster than P+-GaAs [18]. Surface damage, as measured by C-V and I-V characteristics of Schottky diodes, is reduced in Cl2ZN2 ECR-RF etching by using low pressure, high microwave power, short source to sample distance, low RF power, and high concentrations of reactive species [19]. Addition of F to the plasma with CCl2F2 or CHCl2F allows the formation of low-volatility AlF3, which permits very high selectivity for GaAs vs. AlGaAs at low bias voltages (virtually total for 50 to 75 V bias) [14]. Selectivity is higher for higher Al mole fraction and lower at higher bias voltages [14]. Addition of F sources to other Cl-source gases should provide similar selectivity [20]. With BCl3, etch rates decrease with increasing RF power without added Ar; the formation of Bbased compounds was proposed as the possible explanation. With Ar addition, rates increase with increasing RF power. Higher RF power also increases anisotropy and produces smoother surfaces [7]. Equirate etching of GaAs and AlxGa^xAs (x = 0 to 1) occurs with BCl3 and BCl3ZAr mixtures; dilution OfBCl3 with SF6 gives > 600 fold selectivity at 150 V bias with the selectivity decreasing as the bias increases [20]. Etching with PCl3 gives somewhat rougher surfaces and less vertical sidewalls than CCl I^ etching, but they are still smoother than those etched in pure Cl2 discharges [14].

2

Bromine- and iodine-based etching chemistries have also been explored. Whereas Cl-based etchants will generally etch with even 10 - 20 V DC biases, higher biases (>100 V) are required to get appreciable etching with HBr-based plasmas [21]. With HBr/H2/Ar, surfaces are Ga- and Al/Ga-rich above 200 0 C [15]. With BJM2ZAr etching, morphology is smooth at temperatures between 50 and 250 0 C but loss of some As occurs in the top 50 A at high temperatures. Hydrogen passivation to a depth of 3500 A is observed after 10 minutes at elevated temperature. Carrier profiles are restored to their original condition by a 5 min anneal in N2 at 400 0 C [15]. Etching with CH3I, C2H5I, or C3H7I is slower than with HI but up to two times faster than with the corresponding alkane for bias voltages of 150 V [22]. Polymer deposition is minimized at low pressure (10 mtorr) and with H2 dilution [22]. In general, surface stoichiometry is retained and morphology is both smooth and anisotropic over a significantly wider process window than with the corresponding alkane [22]. E2

Alkane-Based Chemistries

Alkane-based processes give lower rates, in general, than Cl-based processes, but produce smoother surfaces. The presence of H in the plasma ensures that lightly doped ( I x 1017/ cm3) ntype materials will require thermal annealing between 350 and 400 0 C to release hydrogen from the Si donors. Acceptor atoms such as Zn can also be passivated. Mixtures of CH4 or C2H6 with H2 have been shown to etch the Ga-containing ternaries with In and Al. For alkane-based etching (CH4 or C2H6) rates generally decrease in the order In-based > Ga-based >Al-based materials [3,16]. Ethane produces up to 50% faster rates than methane under the same conditions [H]. Rates increase linearly with both RF and microwave power [11]. Smooth surfaces result with microwave powers of 200 W or less; higher power produces rough surfaces [H]. Some surface contamination with carbon occurs under most conditions [H]. Severe polymer deposition can occur if H2/CH4 ratios are less than 3:1 [3]. ACKNOWLEDGEMENTS This work was performed at Sandia National Laboratories and supported by the U.S. Department of Energy under Contract No. DE-AC04-94AL85000. REFERENCES [1] [2] [3] [4] [5] [6] [7] [8] [9]

SJ. Pearton. [Int. J Mod. Phy B (Singapore) vol. 8 (1994) p.1781-1876 ] SJ. Pearton, T. Nakano, R.A Gottscho. [ J Appl. Phys. (USA) vol 69 (1991) p.4206 ] C. Constantine et al [ J Vac. ScL Technol. B (USA) vol.8 (1990) p.596 ] SJ. Pearton, F. Ren, CR Abernathy [Appl. Phys. Lett. (USA) vol.64 (1994) p.1673 ] SJ. Pearton, CR. Abernathy, RF. Kopf, F. Ren [J Electrochem. Soc. (USA) vol.141 (1994) p.2250 ] K.D. Choquette, RC Wetzd, RS. Freund, RF. Kopf. [J Vac. Sci. Technol. B (USA) vol. 10 (1992) p.2725 ] S.W. Pang, K.K. Ko [ J Vac. Sci. Technol. B (USA) vol. 10 (1992) p.2703 ] SJ. Pearton, CR. Abernathy, RF. Kopf, F. Ren, W.S. Hobson. [J Vac. Sci. Technol. B (USA) vol.12 (1994) p. 1333] V. Swaminathan, M.T. Asom, U.K. Chakrabarti, SJ. Pearton. [Appl. Phys. Lett. (USA) vol.58 (1991)p.l256]

[10] [11] [12] [13] [14] [15] [16] [17]

[18] [19] [20] [21] [22] [23] [24]

[25] [26] [27] [28] [29]

T. Bickl, B. Jacobs, J. Straka, A. Forchel [ Appl. Phys. Lett. (USA) vol.62 (1993) p. 1137 ] SJ. Pearton, U.K. Chakrabarti, A.P. Perley, W.S. Hobson, M. Geva. [ J. Electrochem. Soc. (USA) vol.138 (1991) p.1432] M.A. Foad, S. Thorns, CD.W. Wilkinson. [ J. Vac. Sci. Technol. B (USA) vol. 11 (1993) p.20 ] VJ. Law, SG. Ingram, G.A.C. Jones. [ Semicond. Sd. Technol. (UK) vol. 6 (1992) p.945 ] SJ. Pearton, U.K. Chakrabarti, A. Katz, A.P. Perley, W.S. Hobson, M. Geva. [ Plasma Chem. and Plasma Process. (USA) vol.11 (1991) p.405-22 ] U.K. Chakrabarti, F. Ren, SJ. Pearton, CR. Abernathy [ J. Vac. Sci. Technol. A (USA) vol.12 (1994) p. 1129] VJ. Law, M. Tewordt, S.G. Ingram, G.A.C Jones [J. Vac. Sci. Technol. B. (USA) vol.9 (1991) p. 1449] SJ. Pearton. [ Proc. Long-Wavelength Semiconductor Devices, Materials, and Processes Symp., Boston, MA, 26-29 Nov. 1990, Eds. A. Katz, R.M. Biefeld, R.L. Gunshor, and RJ. Malik (Materials Research Society, Pittsburgh, USA, 1991) p.277-90 ] S. Miyakuni et al [ J. Vac. Sci. Technol. B (USA) vol. 12 (1994) p.530 ] K.K. Ko, S.W. Pang [ J. Electrochem Soc. (USA) vol. 141 (1994) p.255 ] SJ. Pearton et al [ Plasma Chem. Plasma Process. (USA) vol. 13 (1993) p.311 ] SJ. Pearton, U.K. Chakrabarti, E. Lane, A.P. Perley, CR. Abernathy, W.S. Hobson. [ J. Electrochem. Soc. (USA) vol.139 (1992) p.856 ] U.K. Chakrabarti, SJ. Pearton, A. Katz, W.S. Hobson, CR. Abernathy. [ J. Vac. Sci. Technol. B (USA) vol.10 (1992) p.2378 ] SJ. Pearton, F. Ren, A. Katz, J.R. Lothian, T.R Fullowan, B. Tseng [ J. Vac. Sci. Technol. B (USA) vol.11 (1993) p. 152] SJ. Pearton, F. Ren, A. Katz, B. Tseng, J.R. Lothian, T.R. Fullowan [ Proc. Symp.IU-VElectronic and Photonic Device Fabrication and Performance, 12-15 April, 1993. Eds. K.S. Jones, SJ. Pearton, H. Kanber. (Mater. Res. Soc., Pittsburgh, USA) p. 153 ] S.W. Pang, Y. Liu, K.T. Sung [ J. Vac. Sci. Technol. B (USA) vol.9 (1991) p.3530 ] R. Cheung, Y.H. Lee, CM. Knoedler, K.Y. Lee, T.P. Smith III, D.P. Kern [ Appl. Phys. Lett. (USA) vol.54 (1989) p.2130] R Cheung, Y.H. Lee, KY. Lee, T.P. Smith III, D.P. Kern, SP. Beaumont, CD. W. Wilkinson [ J. Vac. Sci. Technol. B (USA) vol.7 (1989) p. 1462 ] SJ. Pearton, W.S. Hobson [ Semicond. Sci. Technol. (UK) vol 6 (1991) p.948 ] S.K. Noh, K. Ishibashi, Y. Aoyagi, S. Namba, Y. Yoshizako [ J. Appl. Phys. (USA) vol.67 (1990) p.2591 ]

TABLE 1. ECR-RF reactive ion etching rates for GaAs. Etchant

Rate

T ( 0 C)

(nm/min)

Pressure mtorr

Flow (seem)

|i-wave power (W)

RF power (W)

Selfbias (-V)

| Ref

Comments

Halogen-based: Cl-based Cl2

0.23

R.T.

2

10

50

0

lCl2/7He

0.2-0.13

RT.

0.4

20

200

20

8Cl2/2Ar

0.9

R.T.

10

250

10

Cl2/2Ar

0.037

25

70

5Cl2/15Ar

0.50 0.16

BCl3

0.5

[19]

damage removal

[18]

p+=2/3 n+

[7]

f(^wave and RF power, T, P, source sample separation)

80

[7]

smooth, vertical

20

1 1

20 20

500 500

50 50

[4] [4,5]

GaAs=AlGaAs no lateral etch

0.080

1

30

200

150

[20]

GaAs=AlGaAs

65BCl3/35Ar

0.032

1

30

200

150

[20]

GaAs=AlGaAs

65BC13/35SF6

0.071

1

30

200

150

[20]

GaAs/AlGaAs>600/l

65BC13/15C12

0.8-1.1

10-20

80

300

150

[23,24]

30-^im vias

10BCl3/15Ar

0.5

25

500

50

[5]

GaAs=AlGaAs

1CC12F2/2C12

3.4

20

6

200

[25]

nonresonant; rough

CCl2F2ZHe

0.2

5

10

200

1:1

0.09 0.008

5 0.35

10 5

200

0 to +70 -30

20-60

400

20-30

200

[26,27]

80 70

[26] [12]

10±nm damage

TABLE 1. ECR-RF reactive ion etching rates for GaAs (continued). Etchant

Rate (jim/min)

T( 0 C)

CC12F2/O2 28:2

0.0060.086

>35

CC12F2/O2 38:2

0.0060.086

<55

10

CC12F2/O2 38:2

0.0040.086

<55

1

<55

1

<55

1

CC12F2/O2

0.043

Pressure mtorr

Self bias (-V)

Ref

Comments

88

[2]

rough if jiwave power >200 W

10-50

127350

[H]

GaAsZAl03Ga07As 10(30 W RF); (50 W RF)

150

3-8

50110

[14]

GaAs/AlGaAs: oo (50 V); 40 (75 V)

40

250

30

[14]

<20 A halocarbon & hydrocarbon residue

40

250

30

[14]

more damage than with CCl2F2

Flow (seem)

H-wave power (W)

RF power (W)

30

50 - 300

10

40

150

40

CHC12F/O2 38:2

0.041

CHCl2F

0.043

<55

1

40

250

30

[14]

CCl2F2

0.050

<55

1

40

250

30

[14]

CCl2F2ZPCl3 20:7

0.0130.245

1

27

100

100-300

[14]

PCl3

0.0310.076 0.045-0.40

1

7

50-300

100

[14]

1

7

100

100-300

[14]

7PCl3ZlOAr

0.05

1

17

200

80

[28]

SiCl4

0.0440.090 0.0260.068

0.1

7

400

0-100

[6]

confining B field

0.1

7

400

0-100

[6]

no confining field

equirate vs. AlGaAs

TABLE 1. ECR-RF reactive ion etching rates for GaAs (continued). Etchant

Rate (fim/min)

T( 0 C)

Pressure mtorr

Flow (seem)

H-wave power (W)

RF power (W)

Selfbias (-V)

Ref

Comments

f(p, ^iwave W, V)

Br-based: 3HBr/lAr

0.001 0.011

1

20

150

200-400

[21]

3HBr/lH2

0.0020.021

1

20

150

150-400

[21]

3HBr/lCH4

0.004 0.023

1

20

150

150-400

[21]

5HBr/5H2/lAr

0.0100.029

25

250

200

[15]

non-Arrhenius, smooth to 200 ° C

non-Arrhenius, As loss at high T, smooth

50250

I-based: 5HI/5H2/lAr

0.093-0.81

50250

1

25

250

200

[15]

10CH3/

0.023

<80

1

25

250

200

[22]

10C2H5I/ 10H25Ar

0.017

<80

1

25

250

200

[22]

smooth, anisotropic

50

[3]

GaAs»AlGaAs. InGaAs/GaAs 2:1. Stoichiometric surf.

0-40

[14]

-no damage if <40 V

Alkane-based: CH4ZH2ZAr 5:17:5

0.008

CH4ZH2Z5Ar

<0.025

1CH4Z5H2

0.01

27

40500

75

36

250

400

0-60 0-100

[16]

TABLE 1. ECR-RF reactive ion etching rates for GaAs (continued). Etchant

Rate

T( 0 C)

(nm/min) IC2H6ZlOH2

0.015

CH4/H2/Ar 5:17:8 C2H6ZH2ZAr 5:17:8

40500

Pressure mtorr

Flow (seem)

ji-wave

power (W)

75

600

0.0070.016

10

50-300

0.0090.017

10

50-300 33

H 2 OrAr

none

Ar

0.00520.0004

Ar

0.00750.010

RF power (W)

Selfbias (-V)

Ref

Comments

[16] 30

[H]

30

[H]

Other: 125

0-150

[9]

0.4-2.3

400

0,20 250

[10]

damagefromVUV at 0 V bias

0.5 0.5

150 300

[29] [29]

damage study

5-28

4 4

18.8 Etching of GaAs using reactive-ion-beam etching (RIBE)9 chemically-assisted-ion-beam etching (CAIBE)9 and radicalbeam-ion-beam etching (RBIBE) CLH. Ashby March 1995

A

INTRODUCTION

Several approaches to GaAs etching combine a chemically reactive gas with a collimated or focused ion beam to provide highly anisotropic etching. All provide the independent control of ion energy and of the fluxes of chemically reactive neutrals and ions that is lacking in RF-only RIE since ions in these techniques are generated in an ion source rather than in a plasma in direct contact with the sample to be etched. The use of an ion source also reduces the ion-energy spread to a few eV compared to the energy spread of up to several hundred eV in RIE. Reactive-ion-beam etching (RIBE) involves the generation of reactive chemical species such as Cl+ and Cl2+ in an ion source using source gases such as Cl2, CCl4, and BCl3; a beam of these ions is used to etch GaAs. Chemically-assisted-ion-beam etching (CAIBE)5 also called ion beam assisted etching (IBAE), uses a beam of ions, usually Ar+, Xe+, or Ga+, that is directed at a substrate immersed in an ambient of neutral gas, such as Cl2. Focused ion beam (FIB) etching is a variant of CAIBE using a scanned, tightly focused ion beam capable of direct-write etching of small features. Radical-beam-ion-beam etching (RBIBE) is a process in which the neutral gas is predissociated to provide a beam of highly reactive radicals, such as Cl atoms from Cl2, which strikes the surface in concert with an ion beam. In all these processes, GaAs is etched through the combined effect of chemical reaction and physical sputtering. Rates are intermediate between those of physical sputtering and plasma etching. Higher rates usually occur when parameters are selected to increase the importance of chemical etching versus ion sputtering. Appropriate reaction conditions can give both highly anisotropic etching and reasonable etching rates. Etching rates are listed in TABLE 1 for REBE, in TABLE 2 for CAIBE, and in TABLE 3 for RBIBE. Process considerations specific to each process type are discussed in subsequent sections. For all three processes, rates and profiles are controlled in a synergistic fashion by four major parameters: reactive gas pressure, substrate temperature, ion beam current density, and incident ion energy and angle. In general, increasing the chemical contribution to etching enhances etch rates for all three techniques. Increasing the ion energy also increases etch rates. Rates can also be increased by increasing the proportion of ion-derived collision events near the surface to produce a greater number of excited surface atoms to undergo chemical reaction. For example, etch rates with Ar are nearly an order of magnitude faster than with He due to the greater energy deposition in the near surface by the heavier ion [I]. Similarly, maximum rates occur for angles between 60 and 70 degrees where the maximum amount of energy is depositied in the nearsurface region [2,3]. However, non-vertical etch profiles result in oblique incidence angles. Normalizing etch rates measured at lower ion current densities to 1 mA/cm2, as is frequently done, can be misleading. In some cases, the increased sample heating resulting from a higher ion flux can increase thermally-activated chemical reactions. Consequently, normalized rates derived at

lower ion fluxes may be significantly lower than actual rates at higher ion fluxes. Thermallyactivated chemical reactions can also reduce the degree of anisotropy at high ion flux [1,4]. Alternatively, for processes such as FIB-CAIBE [2] and RBIBE [5], where an optimum ratio between molecular or radical flux and ion flux exists, etch rates can decrease at higher fluxes of either type. Etch profiles for all ion-beam-based etching processes have certain characteristics in common due to their beam-dependent character. Several factors affect the edge profile: 1) ion incidence angles, including beam divergence, 2) sidewall etching by reflected primary ions or energetic secondary ion species, 3) deposition of nonvolatile materials on the side wall, 4) chemically derived crystallographic etching, 5) etch-mask sidewall profile, and 6) etch mask degradation. Ion incidence angles may have the most significant effect under common etching conditions. Etch profiles can be changed by changing beam incidence angle [2] and multifaceted or controlledcurvature structures can be made by sequential etch steps or by continuous alteration of the beam angle during etching [6]. Normal incidence can produce overcutting when sputtering is a dominant mechanism, but increasing the incidence angle slightly can produce vertical walls [1,7]. Sample tilt and rotation also increases vertically [7]. Very small structures (9 nm quantum wires) have been etched using angled CAIBE on thin larger structures (80 nm quantum wires) previously etched with vertical CAIBE [8]. The feature size in FIB-CAIBE is limited primarily by the focused ion-beam diameter. Increasing the neutral flux, which increases the chemical contributions to etching, can both improve electronic properties [1,4] and decrease overcutting [9]. Undercutting can be minimized by improving beam collimation [4], which also decreases sidewall damage [10]. Beam divergence is more pronounced at low ion energies, while higher ion energies increase mask erosion [11] and substrate surface damage. Some beam spread due to scattering from neutrals can occur for ion energies below 500 eV [4]. Higher temperatures can favour undercutting and crystallographic etching due to enhanced chemical effects [H]. As in any etching process, etch mask degradation can often dominate the final etch profile since changes in the mask during etching will be transferred into the final pattern on the sample. Higher ion energies increase the potential severity of this complication. B

RIBE

RIBE processes using ECR-plasma ion sources have been the most commonly reported, although RF-plasma ion sources are becoming popular. An alternative plasma source based on electronbeam-excited plasmas (EBEP) produces higher ion densities, thereby permitting lower acceleration voltages for a given etch rate [12,13]. EBEP sources are reported to provide up to several hundred mA/cm2 ion-current densities with 5 - 130 V acceleration over a 30 cm diameter with ±1.2% beam uniformity [12]. Bl

Etch Rates

With an ECR plasma source and Cl2 gas, the etch rates that track the ion current density are proportional to (extraction voltage,V)3/2 [H]. This is seen at 1000C between 200 and 500 V acceleration [11], where etching is 'RIE-like'. When sputtering is a dominant mechanism, such

TABLE 1. Reactive-ion-beam etching (RIBE) rates for GaAs. Normalized rate (micron/min/ 1 mA/cm2)

Pressure (mtorr)

Ion energy (eV)

Current density (mA/cm2)

Beam angle off normal (°)

Ref

Cl2

0.6

0.25

550

0.5

0

[14]

Cl2

8.9 2.2 0.33 0.19 0.067 0.230 0.405

1 1 3.3 3.3 0.2 0.07 0.3

200 200 100 200 500 500

0.09 0.09 6.4 6.4 0.05 0.22 0.21

0 0 0 0 30 30 30

[H] [H] [12] [12] [41] [41] [41]

473 K 373 K e-beam plasma e-beam plasma no loss of As no loss of As no loss of As

0.78 0.06

0.01 0.23

750 500

0.4

0 0

[7] [42]

rate depends on sample source distance

CCl2F2

0.065

0.23

500

55

0

[42]

BCl3

0.12*

1.5

600

0

[17]

2%O2/Ar

0.02

0.14

100

0

[43]

Etchant

Comments

RIBE

CCl4

* rate not current-normalized

0.15

p damaged more than n

as occurs at low temperature, rates depend on both the number and energy of the ions and exhibit a V5/2 dependence [H]: etching without heating at 300 and 400 V approximates this behaviour [H]. When etching is predominantly due to chemical reaction with radicals or molecules, etching becomes increasingly independent of V [H]. At 200 0 C and 200 V, GaAs etch rates are four times higher than at 1000C due to enhanced chemical contributions and are sublinearly dependent OnV[Il]. With a high-density EBEP plasma source, etch rates of 1.2 ^m/min have been reported for GaAs with 5 V acceleration voltage [12]. B2

Profiles

Very smooth surfaces and sidewalls are possible with RIBE, as are near-vertical sidewalls. Lower extraction voltages permit greater beam divergence and increase undercutting. However, higher voltages can cause greater mask degradation [11] and consequent degradation of etched profiles. Etching by scattered primary ions or energetic secondary species can also be greater at higher voltages, increasing undercutting. Higher sample temperatures contribute to crystallographic etching and undercutting; some is seen at 1000C with 300 V beam energies [H]. B3

Selectivity

Equirate etching of GaAs and AlGaAs is possible over the typical operating range of Cl2 pressures and ion energies with both ECR [14,15] and EBEP sources [13] when the H2O partial pressure is below 4 x 10"8 torr [15]. Equirate etching even at ion acceleration voltages of 5 V is attributed to the high etch rate under the associated high ion-current density [13]. B4

Damage

Although the beam energies sometimes used for RIBE can cause significant surface damage, the highly collimated nature of the RIBE beam generally makes RIBE a 'low-damage' process for sidewalls [16]. This makes RIBE well suited for etching very small features where sidewall, and not surface, characteristics are of primary importance. RIBE-induced surface damage is higher at higher ion energies. Although etching with BCl3 and 500 V Ar+ ions produced little change in the Schottky diode parameters for n-type material, it drastically increased the specific contact resistance by decreasing the effective carrier concentration in the sub 20 nm surface damage layer [17]. Damage of p-type material is significantly greater than for n-type GaAs. With 600 V RIBE in BCl3, damage of p-type GaAs with a carrier concentration of 2.2 x io 19 cm"3 extended 20 nm beneath the surface and resulted in a carrier loss of the order of 1019 cm"3 [17]. The use of a high-rate, moderate energy (>200 eV) etch with Cl2 followed by a short-time, lower energy (<100 eV) etch to remove damage yields diodes with ideality factors the same as those fabricated by wet etching [18]. With Cl2 RIBE between 30 and 200 V, diode behaviour is comparable to wet-etched controls [19]. Degradation of both ideality factor and barrier height occurs with extraction voltages above 300 V [19]. With EBEP-RIBE, no degradation of ideality factor was observed up to 120 V ion

energies [12]. EBEP-RIBE with Cl2 or Ar produced 3 deep level defects with the following threshold voltages: 60 V for the 0.31 eV trap, 40 V for the 0.45 eV trap, and 20V for the 0.58 eV trap [20]. Thermal annealing of RIBE-induced deep levels at 400 0 C for 10 min essentially removes them [19]. Caution is required in interpreting low extraction voltage data. For extraction voltages less than 100 V, the extraction voltage and incident ion energy are not the same because the ECR plasma potential is of the order of several tens of eV. For example, extraction at nominally 10 V produced a peak in Cl+ ion energy at 17 eV and 30 V extraction produced a peak at 32 eV with 2*10- 3 torrCl 2 [19]. C

CAIBE

CAIBE employs two general approaches to providing reactive gas. In the most common, a reactive gas, such as Cl2, is locally introduced near the sample surface through a tube or nozzle, forming a jet [1,4,8,10,20-28]. In the second configuration, the gas is introduced nonlocally into the reaction chamber [2,3,9,29]. Similarly, there are two general ion-bombardment configurations. The first uses a broad-area ion gun, usually a Kaufman-type, to provide a uniform beam current density (ion flux) over the entire sample. The second type, called focused ion beam (FIB) etching, uses a focused ion beam, usually Ga+, to promote the etch reaction locally on the surface in a scanned or rastered mode. Cl

Rates

For broad-area CAIBE, etch rates initially increase linearly with Cl2 flow [21,29] due to increasing Cl coverage on the surface. The coverage saturates at relatively low flows and a plateau in etch rate versus flow results [29]. The Cl2 flow rate at which etch-rate saturation occurs increases with increasing beam-current density or ion energy [28]. Over some ranges of flow rate and ion current density, the near-saturation etch rate depends approximately linearly on ion energy in contrast to the E1/2 dependence expected from nuclear stopping considerations [28]. The etch-rate dependence on ion energy is much less at lower Cl2 flow rates; this is attributed to reaction limitation by surface chloride supply [28]. At very high Cl2 fluxes, excessive surface coverage with Cl2 can hinder desorption of etch products [22] and reduce rates. The transition point will depend ontheCl 2 /Ar + ratio[28]. For a given ion current in FIB CAIBE, there is a maximum in rate versus Cl2 gas pressure or flow rate. The highest rates occur when there is an optimal balance between the Cl2 reaction to form a thin chloride surface layer and the ion-induced desorption of products [2]. Coverage of the surface with excessive Cl2 adsorption retards etching [2]. Pulsed FIB etch rates can be more than an order of magnitude higher than continuous FIB rates [30-32]; pulsed yields are greater with shorter ion pulse widths down to 0.02 ^s and for longer off-time between pulses [30]. The FIB rate dependence on ion energy is close to E1/2 under most conditions studied [30]. C2

Profiles

Smooth surfaces are possible with CAIBE etching, but greater initial surface cleanliness is required than RIE. Solvent cleaning of samples just before loading gives smooth surfaces [27].

TABLE 2. Chemically-assisted-ion-beam etching (CAIBE) rates for GaAs. Rate (jim/min/ 1 mA/cm2)

Pressure (mtorr)

Ion energy (eV)

Current density (mA/cm2)

5 0.085 5 3 1.32 0.50 1.36 0.65 1.43

10 10 0.08 0.08 0.25 0.25 0.25 0.25 10 seem

1000 500 500 500 200 200 400 400 500

0.02 0.25 0.02 0.025 0.100 0.025 0.100 0.23

I2+Ar+

1.4

0.14

3000

1.0

Cl2 + (Ar + CCl4)+

0.43

0.20

400

20 20 29 29 29

35 35 0.2

Etchant

Angle off normal (°)

Ref

Comments

0 0 0 0

[4] [4] [1] [1] [5] [5] [5] [5] [21]

±5% uniformity

75-78

[25]

rough if >0.05 mtorr

0.65

0

[9]

RIBE/CAIBE combined

80 80 7.1 7.1 7.1

0 70 0 0 0

[2] [2] [32] [32] [32]

CW CW CW: 16x faster if pulsed

Broad Area: Cl2+Ar+

heating effects

equirate vs. AlGaAs 0-5° laser facet angles

Focused ion beam (FIB): Cl2+Ga+

1.4 3.6 1.6 2.4 4.3

15

CAIBE etch profiles depend on the incident angles of ion-beams [4,6,9,29], the gas molecules [1,2,21,26,33], and on temperature [1,22]. Beam collimation affects the profiles in broad-area CAIBE [4]. The feature size in FIB-CAIBE is limited by the focused beam diameter. FIB-CAIBE has the additional capability of controllably varying the depth of etched features across a sample by varying the beam dwell time. Beam collimation, which affects undercutting, is influenced by several factors. A narrower beam divergence angle (smaller angular spread of ions) results from a greater ion source-to-sample distance, since the collimation is determined to first order by the ratio of beam diameter at source to the source-sample separation for >500 eV ions [4]. At lower ion energies, scattering of ions by neutrals can reduce collimation [4]. As Cl2 flow increases, the mean-free-path OfAr+ ions can be reduced by scattering from Cl2 molecules, decreasing collimation and increasing undercutting [29]. Trenching can result at feature edges when poorly collimated ions are deflected from sidewalls onto the surface; this is more serious at higher energies. When trenching was observed at high ion energy (800 eV), it was greatly reduced by decreasing the ion energy to 500 eV [9]. Sidewall angles depend on the current density for fixed Cl2 flow and Ar+ ion energy. At low current densities with a substantial chemical etch component, mask undercutting and sloped facets result [21]. Facet angles between 0 and 5 degrees have been achieved with beam densities between 0.20 and 0.25 mA/cm2, which also produces equirate etching of GaAs and AlGaAs [21]. The actual angle depends somewhat on the relative position of the facet and the Cl2 jet nozzle [21]. In general, CAIBE reaction rates depend on the Cl2 flux at the surface. For nozzle-based processes, high surface concentrations OfCl2 are possible with relatively low background chamber pressure, which protects the ion source from corrosion and facilitates beam collimation. However, angular spreading of the gas as it exits the nozzle can lead to concentration gradients and resultant etch-rate nonuniformity across the sample. Nozzle configuration including angle and separation from the sample surface is, therefore, important [2,26]. The reactive flux from a nozzle can vary by more than an order of magnitude across a wafer. This can produce non-uniformities across the wafer in both etch rate and damage as measured by the diode ideality factor [I]. Nozzle design and placement are, therefore, critical to uniformity. At room temperature, the beam-nozzle angle can affect surface profiles, producing angled surfaces as features deepen due to the shadowing of surface regions from simultaneous reactive gas and ion flux. This reduces etch rates in shadowed areas toward the limit of the sputter-only rate [I]. Higher chemical reactivity at 1000C removes this effect [I]. At room temperature, reactive-gas scattering from vertical walls can also produce trenching with a nonvertical gas flux [I]. When gas and ion flux are both vertical, aspect ratios greater than 50:1 have been achieved [I]. With vertical-beam FIB-CAIBE, U-shaped profiles occur due to the absence of redeposition, which is responsible for the V-shaped profiles characteristic of FIB sputtering without chemical assistance [2], Increasing the ion-beam angle from 0 to 80° increases the etch rate but produces angled etch surfaces [2]. Independent adjustment of etch rate and sidewall profiles has been achieved by controlling the sample temperature [22] to shift the balance between physical and chemical etching contributions.

High beam currents (1 mA/cm2) produce high etch rates but can degrade anisotropy due to sample heating that enhances chemical contributions [4]. Beam-current heating effects on rate have been seen with current densities as low as 0.25 mA/cm2 [I]. C3

Selectivity

Selectivity is determined by the balance between chemical and physical contributions to CAIBE etching, with greater physical contributions facilitating equirate GaAs/AlGaAs etching. In the absence of a load lock to prevent residual water-induced oxidation of AlGaAs, CAIBE etching of GaAs can be several times faster than etching of AlGaAs up to beam current densities of at least 0.6 mA/cm2 [34]. Both materials have etch rates that are linearly dependent on beam current density. However, while GaAs etch rates monotonically increased with Cl2 flow rates with 500 eV Ar+ at 0.2 mA/cm2, AlGaAs etch rates were relatively independent of Cl2 flow without load locking [34]. When the water partial pressure is low, selectivities up to 2:1 for GaAs versus AlGaAs occur at low current densities (< 0.15 mA/cm2), but etching is equirate and increases linearly with higher current densities as physical sputtering plays an increasingly important role [21]. C4

Damage

Surface electronic damage in broad-area CAIBE depends on the degree of chemical contribution to etching. Whereas CAIBE with a Cl2 flux in excess of 1 x 1018/ cm2 sec with 500 eV Ar+ ions produced diode ideality factors (n) equal to the control, an order of magnitude lower Cl2 flux several mm away on the same wafer yielded significant increases in n [I]. The use of typically higher energy ions in FIB-CAIBE than in broad-area CAIBE leads to photoluminescence (PL) degradation that required 6000C annealing to restore partial PL intensity following 1 and 10 keV etching; the difference between 1 and 10 keV etching was slight [35]. For FIB-CAIBE, the damage depth is shallower for a given ion energy when etching yield is increased by increasing Cl2 flux [36]. Increasing the substrate temperature also decreases the damage depth; for example, etching at 1500C reduced damage below the PL detection limit of 0.5 ^m for 4 x 1017/ cm3 n-type GaAs [36]. Sidewall damage can be negligible in CAIBE with well collimated ion beams. Although the depth of sidewall damage increases significantly with etch time in RIE, relatively little etch-time dependence has been observed using CAIBE-etched quantum wires [10]. Comparison of 500 eV CAIBE with 250 V RIE showed similar reductions of the saturation current due to damagereduced channel widths in narrow wires [27]. The stoichiometry of etched surfaces depends on the balance between sputtering and chemical reaction mechanisms as determined by the ArVCl2 flux ratio [23]. For GaAs(110) surfaces, Garich surfaces form under high Ar+ fluxes and As-rich surfaces under higher Cl2 fluxes for Ar + energies from 500 eV to 3 keV at 300 K [23]. However, if etching occurs at >400 K, a Ga/As ratio of unity is obtained regardless of the Cl2/Ar+ flux ratio. This may be due to surface segregation of bulk As to replace the depleted surface As [23]. In FIB-CAIBE, the surface stoichiometry depends on the etched depth and the beam energy, with surfaces being initially Ga-rich but becoming As-rich as etching deepens [32]. The amount of As enrichment decreases as the beam energy increases from 200 eV to 15 keV [32]. Surface

enrichment in As is also greater under higher Cl2 flux [32], indicating that the physical/chemical sputtering balance plays an important role in determining surface stoichiometry. D

RBIBE

Radical-beam-ion-beam etching (RBIBE) is a process in which the neutral gas is predissociated to provide a beam of highly reactive radicals, such as Cl atoms from Cl2, which strikes the surface in concert with an ion beam. In RBIBE, the primary role of the ions is merely to assist the desorption of products already formed by the reaction with the radical beam. This distinction from RIBE and CAIBE leads to some of the unique characteristics of RBEBE, such as temperature insensitivity and non-monotonic dependence on ion current, discussed below. Dl

Rates

Comparisons of Cl RBIBE with CAIBE (no predissociation to radicals) and RBE (radical beam etching without ion bombardment) at room temperature show an 8-fold increase in rate versus CAIBE due to the higher reactivity of Cl versus Cl2 and an 18-fold increase versus RBE due to enhanced product desorption rates [37]. The relative increase in rate with RBIBE versus CAIBE was greater at lower beam current densities [5]. The RBIBE etch rate with Cl was essentially temperature independent between 5 and 900C [37], as is common for free-radical reactions. Etch rates track roughly with excited-state Cl* emission intensity [5]. Optimum rates occur when the flow of radicals to the surface and the ion current density are selected to balance product formation and desorption. Increases in ion current density can actually decrease both the actual etch rate and the normalized etch rate per ion [5]. Addition of H as HCl or (H2+C12) can increase the GaAs etch rates 3-fold [38], although the mechanism is unclear. FIB etching with Cl instead of Cl2 gives the same maximum sputter yield of 20 atoms per ion, but the higher reactivity of Cl provides this maximum rate over a wider range of scanning rates [39]. A digital RBIBE variation exists where pulses of Cl2 are injected into a cw Ar ECR discharge to produce Cl radical pulses to the surface. These are alternated with Ar+ beam pulses to cyclically form and desorb monolayers of chloride etch products [40]. D2

Profiles

Although the bottom surface is rough for beam current densities below 0.02 mA/cm2 with 200 eV Ar+, surfaces are smooth for current densities between 0.05 and 0.2 mA/ cm2 [37,38]. The degree of anisotropy also increases as current density increases over this range [37]. Facets etched at 0.2 mA/ cm2 were straight and smooth with facet angle determined by ion-beam angle and not Cl beam nozzle angle [37]. Lower current densities are required for smooth surfaces when etching at higher temperature [38]. A higher current density is required for smooth surfaces with HCl (0.2 mA/ cm2) than with Cl2 source gas (0.05 mA/ cm2) while surfaces are smoothest under conditions that produce equirate etching of GaAs and AlGaAs [38]. D3

Selectivity

Equirate etching of GaAs and Al06Ga04As requires base pressures below 3 x 10"7 torr [5]. At low beam current density, etch rates of AlGaAs versus GaAs decrease with the addition of H radicals, either as HCl or (H2-I-Cl2) source gases [38]; the onset of slower AlGaAs etching occurs at lower

TABLE 3. Radical-beam-ion-beam etching (RBIBE) rates for GaAs. Etchant

Rate(jim/min /lmA/cm2)

Flow (seem)

Pressure (mtorr)

Ion energy (eV)

Current density (mA/cm2)

Angle off normal (°)

Ref

Comments

Cl2+Ar+

2.8 6.4 1.9 8.0 1.6 0.75

9 5 5 5 5 5

0.8 0.25 0.25 0.25 0.25 0.25

200 200 200 400 400 200

0.2 0.025 0.100 0.025 0.100 0.200

10

[37] [5] [5] [5] [5] [38]

straight, smooth facets low damage, 300C equirate vs. Al06Ga04As

HCH-Ar+

1.0

3

0.25

200

0.200

[38]

3.0

5+3

0.25

200

0.200

[38]

3 seem H2

(Cl2+H2+N2)+Ar+

0.7

5+3

0.25

200

0.200

[38]

3 seem N2

(Cl2+Ar)+Ar+

1.3

5+3

0.25

200

0.200

[38]

3 seem Ar

Cl then Ar+

monolayer per pulse

5(CW) 0.5 (pulse)

0.5

25

1.9

[44]

sequential Cl2 pulses into CW Ar flow for digital etch

(Cl2+H2)+Ar

+

Al mole fractions with lower current densities [38]. At higher current densities, etching can be equirate. An ion current density of 0.01 mA/ cm2 is sufficient for equirate etching with Cl2 up to Al = 0.5, but equirate etching with HCl requires 0.20 mA/ cm2 at 3O0C [38]. As the temperature increases, etch rates increase faster for Cl2 than for HCl and show less dependence on Al mole fraction [38], A higher Al mole fraction also correlates with increased surface roughness with H addition [38]. D4

Damage

RBIBE with a low beam current density at 200 V can produce diode ideality factors (n) comparable to wet-etched controls, but higher current densities increase n. RBIBE at 200 V also produces high breakdown voltages [5]. Cathodoluminescence (CL) shows damage to a depth greater than 100 nm with 200 V RBIBE, but the CL intensity reduction is greater for 200 V, 0.025 mA/cm2 CACBE than RBIBE [5]. Schottky diode ideality factors from RBIBE are also lower than from comparable CAIBE-etched surfaces [5]. ACKNOWLEDGEMENTS This work was performed at SandiaNational Laboratories and supported by the U.S. Department of Energy under Contract No. DE-AC04-94AL85000. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] [II] [12] [13] [14] [15] [16] [17] [18] [19]

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[20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31]

[32] [33] [34] [35] [36] [37] [38] [39] [40] [41] [42] [43] [44]

J.-Z. Yu et al [ Jpn. J. Appl. Phys. 1 (Japan) vol. 28 (1989) p.2391-95 ] P. Buchmann, H.P. Dietrich, G. Sasso, P. Vettiger {Microelectron. Engin. (Netherlands) vol.9 (1989) p.485-9 ] WJ. Grande, J.E. Johnson, CL. Tang [ J. Vac. Sd. Technol. B (USA) vol.8 (1990) p. 1075-79 ] L.A. DeLouise [J. Appl. Phys. (USA) vol.70 (1991) p.1718-29 ] Z. Xiao, B. Nilsson [ J. Electrochem. Soc. (USA) vol. 138 (1991) p.3086-9 ] L.M. Bharadwaj, P. Bonhomme, J. Faure, G. Balossier, R.P. Bajpai [ Proc. SPIE-The International Society for Optical Engineering (USA) vol. 1593 (1992) p. 186-92 ] RJ. Young, J.R.A. Cleaver, H. Ahmed [ J. Vac. Sci. Technol. B (USA) vol. 11 no.2 (1993) p.234] SW. Pang, WD. Goodhue, TM. Lyszczara, DJ. Ehrlich, RB. Goodman, G.D. Johnson [J. Vac. Sci. Technol. B (USA) vol.6 (1988) p. 1916-20 ] RJ. Davis, E.D. Wolf [ J. Vac. Sci. Technol. (USA) vol.8 (1990) p. 1798-803 ] M. Hagberg, B. Jonsson, A.G. Larsson [ J. Vac. Sci. Technol. B (USA) vol. 12 (1994) p.555-66 ] T. Kosugi, H. Iwase, K. Gamo [ Jpn.J. Appl. Phys. I (Japan) vol.32 no.6 (1993) p.3051-2 ] K. Gamo. [Proc. Symp.Beam Solid Interactions: Fundamentals and Applications, Boston, MA, 30Nov-4Dec. 1992, Ed. M. Nastasi, L.R. Harriott, N. Herbots, RS. Averback (Mater. Res. Soc, Pittsburgh, PA, 1993) p.577-86 ] T. Kosugi, H. Iwase, K. Gamo [ J. Vac. Sci Technol. B (USA) vol. 11(1993) p.2214-18 ] A. Gandhi, J. Orloff [ J Vac. Sci. Technol. B (USA) vol.8 (1990) p.1814-19 ] Z. Xiao, B. Nilsson [ J. Electrochem. Soc. (USA) vol. 138 (1991) p.3086-9 ] T. Kosugi, K. Gamo, S. Namba, R. Aihara [ J. Vac. Sci. Technol. B (USA) vol. 9 (1991) p.2660 ] Y. Sugimoto, M. Taneya, H. Hidaka, K. Akita [ J. Appl. Phys. (USA) vol.68 (1990) p.2392-99 ] J.A. Skidmore, L.A. Coldren, E.L. Hu, J.L. Merz, K. Asakawa [ J. Vac. Sci. Technol. B (USA) vol.6 (1988) p. 1885-8] J.A. Skidmore, D.G. Lishan, D.B. Young, E.L. Hu, L.A. Coldren [ J. Vac. Sci. Technol. B (USA) vol.10 (1992) p.2720-4] N. Takado et al [Mater. Res. Soc. Proc. (USA) vol.75 (1987) p. 107-14 ] Y. Aoyagi, K. Shinmura, K. Kawasaki, T. Tanaka, K. Gamo. S. Namba, I. Nakamoto [Appl. Phys. Lett. (USA) vol.60 (1992) p.968-70 ] R.A.Barker, T.M.Mayer, RH.Burton [ Appl. Phys. Lett. (USA) vol.40 (1982) p.583-86 ] W.-X. Chen, L.M. Walpita, CC. Sun, W.S.C Chang [ J. Vac. Sci. Technol B (USA) vol.4 (1986) p.701-5 ] H. Kinoshita, T. Ishida, K. Kaminishi [ Appl. Phys. Lett. (USA) vol.49 (1986) p.204-6 ] I. Ishii, T. Meguro, H. Kodama, Y. Yamamoto, Y. Aoyagi. [Jpn. J. Appl. Phys. 1 (Japan) vol.31 (1992)p.2212-15]

CHAPTER 19 ION IMPLANTATION AND RAPID THERMAL ANNEALING 19.1 19.2 19.3 19.4 19.5 19.6

Ion implantation in GaAs: an overview Ion ranges in GaAs: discussion Rapid thermal annealing of GaAs: overview Maximum concentrations and activation efficiencies for each ion species Controlled atmosphere annealing of GaAs Compensation mechanisms in GaAs at high dopant fluences

19.1 Ion implantation in GaAs: an overview BJ. Sealy August 1995

A

INTRODUCTION

Ion implantation is widely used to make devices and integrated circuits in GaAs and related materials. For example, the channel, source and drain regions of field effect transistors (MESFETs) are normally fabricated using silicon ion implantation, with applications both to monolithic microwave and digital integrated circuits [I]. The technology is therefore well established, but there is still little understanding of the processes taking place during the electrical activation of implanted dopants. It is therefore difficult to model the process as a whole, although there are a few empirical models which are helpful to the device engineer [2]. This Datareview aims to summarise and provide an overview of achievements, our understanding and problems. B

ACTIVATION AND ANNEALING

All incident ions create disorder in the GaAs lattice during the implantation process. However, a factor that has become apparent in recent years is the need to have very careful control over the implantation conditions if reproducible results are to be achieved. For example, two parameters which affect the degree of electrical activation are the dose rate and the substrate temperature during implantation. It is well known that increasing the substrate temperature to about 200 0 C will prevent an amorphous layer forming for heavy mass or high dose implants, but such implants are rarely utilised for device production. As there is an annealing stage near room temperature, even a small rise in temperature during implantation will reduce significantly the amount of damage introduced into the material for any implant, compared with nominal room temperature implants. At the same time variations in dose rate (beam current density) also influence the amount of damage retained following implantation [3,4]. It is probably the latter parameter that has been the cause of the often large variations in published data. Many ion species have been studied after implantation into GaAs, but the most important ion, the one that has received by far the most attention, is the donor, silicon. However, the complication with silicon is that it is an amphoteric dopant in GaAs, this behaviour being prevalent only for high concentrations, say >5 * 1018 cm"3, and temperatures above about 9500C. It is therefore possible to avoid the problem by choice of dose (peak atomic concentration) and annealing temperature. There are two types of annealing process commonly used by industry, (i) conventional furnace annealing, and (ii) rapid thermal annealing (RTA). Furnace anneals are usually performed at 850 - 900 0 C for times of at least fifteen minutes. In the case of RTA, temperatures are normally higher in the range 900 - 9500C, but only for times of, say, ten seconds. In general, the electrical properties of ion implanted GaAs depend on both annealing time and temperature [5,6]. The sheet carrier concentration increases with increasing time at a given temperature until a saturation value is reached, which is dependent on temperature alone. A characteristic activation energy can be deduced for each of these two regions (see TABLES 1 and 2) [7]. A value of about 2.5 eV, the diffusion energy, E^ has been found for many ion species in the time dependent region. However, some rather higher values of around 4.5 eV have also been recorded. In

contrast, in the time independent region, saturation energies, Ea, are found to lie in two bands of values, ranging from 0.3 - 0.5 eV and 1.0 - 1.8 eV. TABLE 1. Activation energies for the electrical activation of various ions implanted into GaAs. Ion

Implant temperature CQ

Dose (cm"2)

Saturation energy (eV)

Diffusion energy (eV)

Ref

Be, P

RT

1 x IQ14 Be+ 7.7 x IQ13P

0.4, 1.3

-

[81

14

Be

RT RT RT

5 x 10 1 x 1015 1 x IQ15

0.32 +/- 0.05 0.70 035

2.3 +/- 0.1 -

[9] [10] [U]

BeF

RT

2 x IQ14

1.1-1.2

-

[42]

Mg

RT RT

I x 1014 5 x IQ14

0.98 0.98+/-0.10

2.45+/-0,1

[9] [9]

Zn

RT

1 x IQ15

0.35 +/- 0.05

U

[12]

Cd

LiqN RT 400

5 x 1013 5 x 1013 5 x IQ13

0.8 0.6/0.7 04

-

[13] [13] [13]

Hg

200

1 x IQ14

0.4 +/- 0.05

-

[14]

S

RT RT 400

I x 1014 2 xlO 14 2 x IQ14

1.1 +/-Ol 1.2 L2

-

[9] [15] [15]

Se

RT RT

2.3

5.4,5.8 -

[16] [17]

RT RT RT RT RT 400

2 x 1013 3 x 1013 4 x 1014 1 x 1014 1 x 1014 1 x 1014 2 x 1014 1 x 1015 2 x IQ14

1.2 1.2 1.2+/-0.1 2.2 1.3 L2

4.5 2.5+/-0.1 -

[18] [19] [9] [15] [20] [20]

Te

LiqN RT 200 400

5 x 1013, 1015 5 x 1013, 1015 5 x 1013, 1015 5 x IQ13, IQ15

1.8 0.9 0.7 0.6/0.7

-

[13] [13] [13] [13]

Sn

LiqN RT RT RT 200

5 x 1013, 1015 1 x 1014 1 x 1015 5 x 1013, 1015 5 x 1013, 1015

1.8 1.2+/-Ol 1.2+/-01 1.6 08

2.5+/-01 2.5+/-01 -

[13] [21] [21] [13] [13]

I

400

1

5 x IQ 13 , IQ 15

1

0.7

I

-

I [131

TABLE 2. Activation energies for the electrical activation of silicon ion implanted GaAs. Ion

Energy (keV)

Dose (cm'2)

Saturation (eV)

Diffusion (eV}

Ref

Si Si Si Si Si Si Si Si Si Si Si Si Si, P Si Si Si Si5P Si, P Si, B Si5B Si, Ga, As

30 60 100 100 100

4.5 xlO 13 I x 1015 1 x 1013 1 x 1014 3 x 1012 5 x 1013 2 x 1013 1 x 1013 5 xlO 13 6 x 1013 4 xlO 12 3 xlO 13 1 x 1014 6 x 1012 5 xlO 13 1 x 1014 5 xlO 13 1 x 1014 6 xlO 12 12 6 x 10 Si+ 3 x 1013B 2 x IQ13 Si+1 x IQ13 Ga/As

0.48 0.51 0.43 0.53

-

[22] [10] [11] [11]

0.55 0.79 0.76 0.5-0.7 0.54 1.0 1.0 1.45 1.45 0.63,0.5 1.0,0.5 -

4.75 3.0 2.3 1.7 2.5

[23] [16] [24] [25] [26] [27] [27] [28] [29] [30] [30] [30] [30] [30] [30] [16]

120 150 150 150 150 150 200 200 200 200 200 200 200 1 120,270,290

|

[

|

|

It is difficult to draw any definitive conclusions from the data in TABLES 1 and 2. However, it has been suggested that the low value saturation energies (about 0.5 eV) are related to the energy required to place an interstitial atom on the correct lattice site, e.g. Be or Zn interstitials onto a gallium vacancy [5]. The larger values of about 1.0 - 1.5 eV may correspond to the energy required to break up a complex defect of the dopant with a neighbouring vacancy [5]. The diffusion energies are mostly in the range 2 - 3 eV which are very close to values recorded for the diffusion of gallium and arsenic in GaAs [31,32] and also close to values for the diffusion of some of the impurity atoms [33]. C

PROFILES

It is relatively straightforward to dope GaAs either p-type or n-type by ion implantation and the technique is used widely in industry. However, a problem that needed to be overcome in order to get good uniformity of device characteristics across a wafer was the channelling tail, which is always present to a greater or lesser extent when silicon is used to create n-type regions for MESFETs. In essence the solution is simple, that is, to implant an electrically compensating atom into the tail of the silicon profile. Various ion species have been used successfully, such as Mg, Be, B, C, O and P. The result is invariably a sharpening of the profile. However, the best results have been obtained when carbon is used, because this atom is a very effective compensating species in the presence of silicon, but beyond the channelling tail its effect is as a weak acceptor with negligible electrical activity [34,35]. Thus the tail is removed without the formation of a buried p-type layer, which is the ideal situation and an important requirement. If other ion species are used, it is necessary to control their degree of electrical activity very carefully so that the layer is fiilly depleted, rather than becoming p-type. This can be done and excellent devices have been made using this process [36].

Electrical/atomic profiles of high dose implanted acceptor atoms tend to suffer from diffusional broadening during subsequent heat treatment. Several methods can be used to control this phenomenon, ranging from low temperature and/or short time annealing, to dual implants which are believed to introduce defects which inhibit diffusion. For example, implanting Be and P into GaAs produces a higher peak acceptor concentration with little diffusion, even under the same annealing conditions which would produce significant profile broadening for a single implant of Be [8]. The reason for this result can be understood by consideration of the probable relative concentrations of gallium and arsenic vacancies. When P is implanted with the Be, arsenic vacancies tend to be occupied by the additional P atoms. In turn, assuming that thermodynamic equilibrium is approached during annealing, the concentration of gallium vacancies should increase in proportion to the number of P atoms occupying arsenic sites. The larger the concentration of gallium vacancies, the more likely it is that Be atoms will be able to locate and occupy such vacancies. This reduces the likelihood of fast interstitial diffusion, which would otherwise take place. D

APPLICATIONS OF ION IMPLANTATION TO GaAs

For GaAs, ion implantation is used most frequently in the fabrication of the channel (active) region and the more highly doped source and drain regions of MESFETs. It is used also for introducing deep levels in the bandgap which trap mobile carriers and therefore create semi-insulating regions which separate one circuit component from another. Some modern MESFET devices also utilise large angle implants to generate lightly doped drain (LDD) structures, commonly used in advanced silicon devices. In order to get this technology to work for submicron device structures, it is necessary to produce highly resistive layers beneath the active channel of the MESFET. Beryllium ions have been used for this purpose because they produce little damage due to a combination of their small mass and the very low dose required [36]. Analogous to this, a fiilly implanted JFET device has been reported recently which involves using a carbon implant to sharpen the tail of the silicon doped n-channel and to produce a fiilly depleted region (high resistivity) beneath [37]. This work implies the fabrication of p+n junctions by ion implantation which has been discussed by the same authors elsewhere, who compared Mg and Zn gate implants to produce very shallow and abrupt junctions for JFETs [38]. E

CONCLUSION

In this Datareview, no distinction has been made between the use of long time furnace and rapid thermal annealing, since they are both used widely. The section on activation and annealing refers to both types of annealing process by virtue of the time spent at the maximum temperature attained. The reader is referred to several recently published review articles on ion implantation of GaAs [39,40] and for a more detailed discussion of RTA, see [41] and references cited therein. REFERENCES [1] [2] [3] [4] [5]

J. Mun (Ed) [ GaAs Integrated Circuits - Design and Technology (BSP, Oxford, 1988) ] S.E Hansen, M.D. Deal [ SUPREM3.5 - A Program for Process and Device Simulation of GaAs (Stanford University, January 1990) ] T.E. Haynes, R. Morton, S.S. Lau [Appl Phys. Lett. (USA) vol.64 (1994) p.991-3 ] E. Wendler, W. Wesch, G. Gotz [Nucl Instrum. Methods Phys. Res. B (Netherlands) vol.52 (1990) p.57-62 ] BJ. Sealy [ Semicond. Sci. Technol. (UK) vol.3 (1988) p.448-51 ]

[6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33] [34] [35] [36] [37] [38] [39] [40] [41] [42]

N. Morris, BJ. Sealy [ Nucl. Instrum. Methods Phys. Res. (Netherlands) vol.39 (1989) p.453-6 ] B.J. Sealy [NATOASISer. E, Appl Sd. (Netherlands) vol.144 (1988)p.215-38 ] A.C.T. Tang, B.J. Sealy, A.A. Rezazadeh [ Electron. Lett. (UK) vol.25 (1989) p. 1077-9 ] N. Morris, B.J. Sealy [ Inst. Phys. Conf. Ser. (UK) no.91 (1988) p. 145-8 ] K.D. Cummings, S.J. Pearton, G.P. Vella-Coleira [Mater. Res. Soc. Symp. Proc. (USA) vol.52 (1986) p.375-81 and J. Appl. Phys. (USA) vol.60 (1986) p.163-8 ] S.J. Pearton et al [J Appl. Phys. (USA) vol.67 (1990) p.2396-409 ] R. Bensalem, B.J. Sealy [ Appl. Phys. Lett. (USA) vol.50 (1987) p. 1382-3 ] S.J. Pearton et al [ J. Appl. Phys. (USA) vol.65 (1989) p. 1089-98 ] A.C.T. Tang, S.R. Gardener, B.J. Sealy, W.P. Gillin [ Electron. Lett. (UK) vol.25 (1989) p. 161820] W.H. van Berlo, T. Pihl [ Nucl. Instrum. Methods Phys. Res. B (Netherlands) vol.55 (1991) p.785-8 ] T. Kanayama, H. Tanoue [ Mater. Res. Soc. Symp. Proc. (USA) vol.279 (1992) p. 159-64 ] B.D. Choe, S.I. Kwun, Z.G. Khim, J.S. Chang [NewPhys. (Korean Phys. Soc.) (South Korea) vol.21 (1981) p.219-25 and J. Korean Phys. Soc. (South Korea) vol. 16 (1983) p.72-7 ] B.J. Sealy, N.J. Barrett, R. Bensalem [ J. Phys. D (UK) vol. 19 (1986) p.2147-55 ] N.J. Barrett, J.D. Grange, B.J. Sealy, K.G. Stephens [J. Appl. Phys. (USA) vol.56 (1984) p.3503-7 ] R.L. Chapman, J.C.C. Fan, J.P. Donnelly, B.Y. Tsaur [ Appl. Phys. Lett. (USA) vol.40 (1982) p.805 ] R. Bensalem, B.J. Sealy [ Vacuum (UK) vol.36 (1986) p.921-3 ] J.P. de Souza, D.K. Sadana, HJ. Hovel [Mater. Res. Soc. Symp. Proc. (USA) vol.144 (1989) p.495-9 ] Y. Suzuki, S. Komatsuzaki, J. Kasahara [ Inst. Phys. Conf. Ser. (UK) no. 106 (1989) p.539-44 ] T. Hiramoto, Y. Mochizuki, T. Saito, T. Ikoma [ Jpn. J. Appl. Phys. (Japan) vol.24 (1985) p.L921-4] S. Adachi [ J. Appl. Phys. (USA) vol.63 (1988) p.64-7 ] S.K. Tiku, W.M. Duncan [ J. Electrochem. Soc. (USA) vol. 132 (1985) p.2237 ] W.M. Paulson, RN. Legge, CE. Weitzel [ J. Electron. Mater. (USA) vol. 16 (1987) p. 187-93 ] CW. Farley, T.S. Kim, B.G. Streetman [ J. Elec. Mater. (USA) vol. 16 (1987) p.79-85 ] RM. Gwilliam, RS. Deol, R Blunt, B.J. Sealy [Mater. Res. Soc. Symp. Proc. (USA) vol.92 (1987) p.437-44 ] RM. Gwilliam [ Ph.D. Thesis, University of Surrey, UK (1991) ] H.D. Palfrey, M. Brown, A.F.W. Willoughby [J. Electrochem. Soc. (USA) vol.128 (1981) p.2224 ] H.D. Palfrey, M. Brown, A.F.W. Willoughby [ J. Elec. Mater. (USA) vol. 12 (1983) p.863-77 ] I. Harrison, B. Tuck [ inProperties of Gallium Arsenide (INSPEC, IEE, London, UK, 1990) ch. 14 p.341-59] B.P. Davies, P. Davies, D.M. Brookbanks, DJ. Warner, RH. Wallis [ Inst. Phys. Conf. Ser. (UK) no.l 12 (1990) p.275-80 ]; see also [ Datareviews 13.2 and 13.3 in this book ] RM. Gwilliam, RJ. Wilson, T.D. Hunt, BJ. Sealy [ Nucl. Instrum. Methods Phys. Res. B (Netherlands) vol.74 (1993) p.94-7 ] K. Onodera, K. Nishimura, K. Asai, S. Sugitani [ IEEE Trans. Electron Devices (USA) vol.40 (1993) p. 18-24] J.C Zolper, M.E. Sherwin, A.G. Baca, RJ. Schul, J.F. Klem, V.M. Hietala [ IEEE Electron. Devices Lett. (USA) vol.15 (1994)p.493-5 ] M.E. Sherwin et al [ J. Electron. Mater. (USA) vol.23 (1994) p. 809-18 ] S.J. Pearton [Mater. Sci. Rep. (Netherlands) vol.4 (1990) p.313-67 ] S.J. Pearton [ Int. J. Mod. Phys. B (Singapore) vol.7 (1993) p.4687-761 ] BJ. Sealy [ Datareview in this book 19.3 Rapid thermal annealing of GaAs: overview ] CE. Wu, T. Keller, K. Evans [ J. Electrochem. Soc. (USA) vol. 140 (1993) p.3230-2 ]

19.2 Ion ranges in GaAs: discussion BJ. Sealy August 1995

A

INTRODUCTION

When an energetic ion penetrates a solid target, it will eventually come to rest at a depth characteristic of the ion energy and the mass and atomic number of both the ion and target atoms. The ions are slowed down by losing energy in two ways: (i) by excitation and ionisation of electrons, and (ii) by elastic collisions with target nuclei. These two processes are called (i) electronic, and (ii) nuclear stopping, respectively. Initially, the ions are slowed down by electronic stopping, but as their velocity decreases, nuclear stopping becomes predominant until the ions come to rest at a depth called the projected range. The contribution of electronic stopping to lattice damage is very small, but nuclear stopping can cause severe damage to crystalline targets and, in the extreme, the surface layer can be rendered amorphous. The projected range corresponds to the effective distance travelled perpendicular to the sample surface, which, due to the random nature of the stopping process, varies from ion to ion. Thus for a large number of incident ions, there is a distribution of projected ranges which forms an approximately Gaussian shape. This distribution can, therefore, be characterised by a mean projected range, Rp, and a standard deviation of the range, ARp. Experimentally determined atomic profiles with depth are often found to deviate from a Gaussian distribution. For example, a tail on the deep side of the atomic distribution invariably occurs when implanting into crystalline targets, to produce a skewed profile. In order to model this profile distortion, it is necessary to introduce two additional parameters which represent the skewness (y) and curtosis (P) [I]. Data has been published of the values of all four moments Rp, ARp, y and P for a variety of ions and ion energies in GaAs, using SIMS to determine the atomic profiles. There is, however, some disagreement in the literature due to the inherent problems in accurately determining atomic profiles by SIMS. The following section contains some of the published data and references to other results. Theoretically determined values of profile parameters for many ion species implanted into GaAs over the energy range 50 keV to 2 MeV are tabulated in [I]. B

PUBLISHED DATA

Most data on range statistics make use of the SIMS technique to determine the atomic profile following ion implantation. This profile is then modelled using an appropriate mathematical technique. The simplest of these is to fit a Gaussian distribution to the experimental data to quantify the mean projected range and standard deviation. However, this process, although often adequate, can give poor agreement between the theoretical model and measurement. It is therefore necessary to use a more complicated mathematical model to get a better fit to experimental results. Thus profiles have been interpreted as 'modified Gaussian', 'hyperbolic Gaussian' or 'Pearson type' distributions [2]. Hence the introduction of additional moments which characterise the skewness and curtosis of profiles.

Anholt [3] studied profiles of Si, Se and Be in the range 20 - 400 keV and noted that it is necessary to use a Pearson IV distribution to get exact agreement between theory and experiment, i.e. it is necessary to use four moments. These authors found that the mean projected range was proportional to some power of the ion energy. For example, in the case of silicon ions the mean projected range in GaAs was determined to be 2.5E0853 nm, where E is the ion energy in keV. In contrast, Favennec [4] found that the mean projected range is a linear function of ion energy for a significant range of values and that, for silicon ions, Rp (in microns) = 11.2 x 10"4E, where E is measured in units of keV and is valid up to 600 keV. Values of the proportionality constant were determined for Be, B, C, O, Mg, S, Zn, As, Se, Cd, and Te (see TABLE 1). Thompson has determined the range statistics and profile shape factors for 1, 2, 4 and 6 MeV silicon ions in GaAs from SIMS measurements of atomic profiles [5]. Additionally, Wilson [6] has carried out a very detailed study of profiles of H, He, rare earths, Be, Mg, Zn, C, Si, Ge, S, Se and Te in GaP, GaAs and InP. Ion energies were varied between 100 keV and 6 MeV. Again, raw data generated from the SIMS technique was fitted to a Pearson type IV distribution and the results compared with TRIM and LSS calculations (see [I]). Wilson found that there is often a significant difference between calculated ranges and measured values, sometimes by up to 50%, although this is not typical. An important point to appreciate is that profiles can be calculated by a number of different procedures (see [I]) but can often be inaccurate. Because the differences between calculated and measured values are often greater than device designs can tolerate, it is important to experimentally determine profiles using techniques such as SIMS or Rutherford backscattering. Wilson [7] has also investigated profiles of hydrogen in GaAs at ion energies of 50, 100 and 200 keV in random and <100> and <110> channelling directions. The <100> channelling ranges are slightly deeper than the (100) random ranges. However, the <110> channelling ranges are 1.4 to 1.7 times deeper than the random range. Range and shape factors were determined for both 1 H and 2H ions. Similar data has also been obtained for 200 keV He ions [8]. The ranges of 25 - 50 meV oxygen ions in GaAs and InP have been measured recently by Pearton [9]. Using SIMS, the projected ranges were determined to be 14.0 and 28.8 ^m for 25 MeV and 50 MeV ions in GaAs, respectively. TABLE 1. The variation of K with ion mass for which the relationship Rp = KE is approximately true [4]. Ion

Mass

K (10'4microns/keV)

Valid up to energy E(keV)

Be B C O Mg Si S Zn As Se Cd

9 11 12 16 24 28 32 64 75 80 114

40 28 26 20 13.4 11.2 9.7 4.8 4.2 3.6 2.8

100 200 200 400 500 600 700 1000 1000 1000 1000

Te

I

130

1

2J63

|

1000

Simonton [10] has studied the effect of the incident angle of the ion beam with 150 mm diameter

GaAs wafers to gauge the effect of channelling. They conclude that a 10° tilt is not generally adequate to avoid the effects of planar channelling across the full diameter of bare GaAs wafers greater than about 100 mm diameter. The conclusion is that for large wafers, tilt angles in the range 12 - 15° may be required to minimise channelling. Device fabrication often entails implanting through a deposited film, perhaps photoresist or a dielectric layer such as silicon oxide or silicon nitride. Whilst it is often convenient to assume that the stopping powers of ions in the layer and in the semiconductor are equal, this can give rise to errors, although it provides a useful guide to designing an experiment. Since this is an important practical topic associated with the application of ion implantation to device fabrication, it is necessary to be able to simulate the process of implanting through layers. Thus several authors have developed and discussed analytical models to do this [2,11-14] with satisfactory results. SUPREM 3.5 [15] and SUPREM - IV.GS [16] are probably the most advanced simulators which include the concept of implantation through layers. However, it is still not possible to predict from first principles an accurate value of the entire atomic (electrical) profile for implants into GaAs, which could then be used in a device simulator. Reasons for this are associated with the variability of the annealing process, ion channelling and control of the wafer temperature during implantation (see [17]). C

CONCLUSION

It is possible to calculate the mean projected range of any ion species implanted into any target. However, depending on the details, this calculation can have a large error. In practice, it is important to measure the atomic distribution in order to have confidence in the experiment being performed. Theory should merely be used as a guide to get started. The reader is referred to the reference list for further information on profile measurement and simulation [I]. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] [II] [12] [13] [14] [15]

R.P. Webb, B.J. Sealy [ Datareview in this book: 19.2 Ion ranges in GaAs: discussion ] K. Tabatabaei-Alavi, LW. Smith [ IEEE Trans. Electron Devices (USA) vol.37 (1990) p.96-106 ] R. Anholt, P. Balasingham, S.Y. Chou, T.W. Sigmon, M. Deal [ J. Appl. Phys. (USA) vol.64 (1988) p.3429-38 ] P.N. Favennec, M. Gauneau, M. Salvi [ Solid State Phenom. B vol. 1-2 (1988) p.377-416 ] P.E. Thompson, R.G. Wilson, D.C. Ingram, P.P. Pronko [J Appl. Phys. (USA) vol.65 (1989) p.2986-90 ] R.G. Wilson [J Electrochem. Soc. (USA) vol.138 (1991) p.718-22 ] R.G. Wilson [ J Appl. Phys. (USA) vol.61 (1987) p.2826-35 ] R.G. Wilson [ J. Appl. Phys. (USA) vol.61 (1987) p.2489-91 ] SJ. Pearton et al [ J. Appl. Phys. (USA) vol.71 (1992) p.2663-8 ] R.B. Simonton, D.H. Rosenblatt, E. Corcoran, D. Kamanitsa [ J Appl. Phys. (USA) vol.71 (1992) p.2441-8] H. Ryssel, W. Kruger, J. Lorenz [ Nucl. Instrum. Methods Phys. Res. B (Netherlands) vol. 19/20 (1987)p.40-4] L. Duricek [ Cryst. Prop. Prep. (Switzerland) vol. 19/20 (1989) p. 173-8 ] RP. Webb [ in Practical Surface Analysis (2nd Edition) Eds D. Briggs, M.P. Seah, vol.2 Ion and Neutral Spectroscopy (Wiley, 1992) p.657-704 ] M. Deal et al [ Tech. Dig. Int. Electr. Dev. Mtg. (1987) (IEEE, New York, USA, 1987) cat.no.87CH2515-5p.256-9] S.E. Hansen, M.D. Deal [ Suprem 3.5 -A Program for Process and Device Simulation of GaAs

[ 16] [17]

(Stanford University, USA, January 1990) ] S.E. Hansen, M.D. Deal (Eds) [ Suprem-IV. GS - Two Dimensional Process Simulation for Silicon and Gallium Arsenide (Stanford University, USA, 1993) ] BJ. Sealy [ Datareview in this book: 19.1 Ion implantation of GaAs: overview ]

19.3 Rapid thermal annealing of GaAs: overview BJ. Sealy August 1995

A

INTRODUCTION

The pioneering work on rapid thermal annealing (RTA) of ion implanted GaAs was published in 1976 [1,2]. It was demonstrated for the first time that a saturation of the electrical activity occurred at each annealing temperature as the annealing time was increased, where the electrical activity is defined as the ratio of the sheet carrier concentration to the retained ion dose. The time required to reach the saturation level decreased with increasing temperature, as one would expect from thermodynamic considerations. The early work employed a graphite strip heater, but later studies introduced the idea of using quartz-halogen lamps [3] and most, if not all, commercial RTA systems make use of this method of heating [4,5].

B

ADVANTAGES OF RTA OVER FURNACE ANNEALING

Employing the rapid thermal annealing technique means that the effects of thermally activated processes, such as diffusion and decomposition, are reduced, especially if the temperature is low as well as the time being short. Thus for compound semiconductors like GaAs, there is a distinct advantage in using RTA compared with conventional furnace annealing, because during annealing the surface is less likely to decompose and lose arsenic. Although this makes the reliability of the encapsulant less problematic and the encapsulant process step less critical, some way of protecting the surface is still required. This can be carried out in one of several ways, with the most usual being a thin deposited layer of PECVD silicon nitride. Other techniques include proximity or capless methods which are normally limited to a maximum temperature of about 900 0 C, even if the time is limited to a few seconds [6]. The proximity technique normally involves resting a sacrificial wafer of GaAs or silicon on top of the sample. Assuming that the two surfaces are similarly flat, the volume available to evaporating arsenic is small and it is as if the sample is being annealed in an atmosphere of arsenic. The capless technique usually refers to the use of a controlled atmosphere of arsenic which engulfs the sample and minimises the loss of arsenic from the surface [6]. C

PROBLEMS WITH THE USE OF RTA

Despite the obvious advantages of RTA over furnace annealing, there are various problems which make the technique difficult to employ on a commercial basis. The measurement of sample temperature during RTA is notoriously difficult. However, for a given piece of equipment, the reproducibility of the sample temperature can be extremely good (within a few degrees) and acceptable for device and circuit processing. The major issues associated with the application of the technique are concerned with (i) the uniformity of temperature/processing of individual wafers, (ii) the repeatability of the entire thermal cycle, (iii) the problem of introducing stress into the wafer as a result of the rapid treatment, and (iv) throughput, since currently the technique is applied to a single wafer, rather than a batch of wafers (as in the furnace annealing process). Thus, to change technology from a furnace (batch) to an RTA process (single wafer) requires some trade-offs of the above parameters.

For processing advanced silicon integrated circuits, the RTA (RTP) technique is essential because of the small dimensions over which very fine control is required. Such precise control of the overall process may not be so important for annealing ion implanted GaAs (III-V materials), since the degree of integration is much lower and the design rules much less stringent than for silicon integrated circuits. A particular problem for GaAs is the likely decomposition during capless or proximity annealing. However, this can be avoided by using a deposited encapsulating layer of, for example, Si3N4, SiO2, AlN, SiON, or WN, although none of these is perfect. Silicon nitride is the preferred and most widely used encapsulant, but it is necessary to control deposition conditions very carefully in order to avoid variations in composition which can significantly affect the outcome of an annealing cycle. D

CONCLUSION

There have been many publications on the topic of RTA of GaAs and III-V compound semiconductors [7-15] and the reader is referred to these for a much more detailed description of work done and achievements to date. A book entitled 'Rapid Thermal Processing' [5] has recently been published, and, even though it is aimed entirely at applications to silicon, it is a useful and informative reference text. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] [II] [12] [13] [14] [15]

R.K. Surridge, BJ. Sealy, A.E.D. D'Cruz, K.G. Stephens [ Inst. Phys. Con/. Proc. A (UK) no.33 (1977)p.l61-7] BJ. Sealy, S.S. GiU [Les Editions de Physique, Proc. MRS-Europe 1985, USA (1985) p.361-72 ] M. Arai, K. Nishiyama, N. Watanabe [ Jpn. J. Appl Phys. (Japan) vol.20 (1981) p.L124-6 ] S.S. Gill [ Phys. Technol. (UK) vol. 17 (1986) p.245-53 ] RB. Fair (Ed.) [ Rapid Thermal Processing (Academic Press, 1993) ] BJ. Sealy [ Datareviews in this book: 19.1 Ion implantation in GaAs: overview, and 19.5 Controlled atmosphere annealing ] M. Kuzuhara, H. Kohzu, Y. Takayama [ Mater. Res. Soc. Symp. Proc. (USA) vol.23 (1984) p.651-62] D. Eirug Davies [ Mater. Res. Soc. Symp. Proc. (USA) vol.45 (1985) p.261-71 ] D. Eirug Davies [ Nucl Instrum. Methods Phys. Res. B (Netherlands) vol.7/8 (1985) p.387-94 ] S.S. Gill, BJ. Sealy [ J. Electrochem. Soc. (USA) vol. 133 (1986) p.2590-6 ] W. Wesch, G Gotz [ Phys. Status Solidi A (Germany) vol.94 (1986) p.745-66 ] R Singh [J Appl. Phys. (USA) vol.63 (1988)p.R56-R114 ] S.S. Gill [SolidStatePhenom. vol. 1/2 (1988)p.281-342 ] BJ. Sealy [ Semicond. Sci. Technol. (UK) vol.3 (1988) p.448-51 ] SJ. Pearton [ Int. J Mod Phys. B (Singapore) vol.7 (1993) p.4687-761 ]

19.4 Maximum concentrations and activation efficiencies for each ion species BJ. Sealy August 1995

A

INTRODUCTION

One major problem in the application of ion implantation to GaAs is that, in general, only a fraction of the dose of the implanted ion species becomes electrically activated after annealing. The actual percentage varies between ion species and also with dose, low doses having the highest percentage activity and high doses having increasingly lower values. It has also been established that n-type dopants do not activate so readily as p-type dopants. Thus for doses in the range 1014 to 1015 ions cm"2, it is possible to obtain approximately 100% activity for p-type dopants such as Be, Mg, Zn and Cd, but far less, say 1 - 10%, for n-type dopants. For the silicon ion doses (1 x io 12 to 5 x io 12 cm"2) required for producing the active channel of MESFET devices, activations of 50 - 100% are normally achieved, with mobilities of around 4000 Cm2V"1 s"1, but for the more highly doped regions needed for the source and drain, the activities are less. This Datareview lists the maximum values of peak carrier concentration achieved for a range of ion species and doses, sometimes under quite extreme annealing conditions. The encapsulants used to obtain these results are various and are not mentioned. For further details the reader is referred to the list of references and the many review papers published (see [18] for a recent review). B

PEAK CARRIER CONCENTRATIONS

Typical peak carrier concentrations are about 2 x io 19 cm"3 and 2 x 1018 cm"3 following furnace anneals at about 750 0 C and 8500C, for p-type and n-type material respectively. However, TABLE 1 shows that values almost an order of magnitude higher than these can be achieved, but often at annealing temperatures which are far too high for device fabrication. In brief, SILICON is the most used ion species for device fabrication and good activations can be achieved, but it is amphoteric. GERMANIUM is also amphoteric producing, in general, low levels of activation with low mobilities. SELENIUM is also activated readily for doses below about IO13 ions cm"2, but can produce electron concentrations of around IO19 cm"3 only after very high temperature anneals. SULPHUR diffuses rapidly and is not suitable for device fabrication, whilst TELLURIUM, being a heavy element, creates much lattice damage unless implanted into a hot substrate, and even then the peak electron concentrations are still not particularly high. TIN is similar in mass to tellurium, but seems to activate more readily to give electron concentrations in the low 1018 cm"3, with maximum values close to IO19 cm"3 being recorded. All of the acceptor atoms diffuse fast and so profiles can broaden significantly during annealing, even if RTA is used. The implantation of As or P together with the dopant has been shown not only to reduce diffusion, but also to increase the peak hole concentration. CARBON has been studied as a possible p-type dopant, but the percentage activity is very low, unless it is co-implanted with gallium which can produce peak concentrations around IO19 cm"3 [18].

TABLE 1. Maximum peak carrier concentrations for various ion species in GaAs.

Ion species

Dose (cm'2)

Peak carrier concentration (cm'3)

Anneal temperature ( 0 C)

Time (sec)

Si Si S S Se Se Se Se Sn Sn Te Be Mg Mg 5 As Mg 9 As Mg, P Mg, As Zn, As Zn Zn Cd

4 x 1014 5 xlO 1 4 3 x 1013 1 x 1014 1 x 1014 1 x 1014 1 x 1015 1 x 1015 1 x 1015 1 x 1015 1 x 1015 1 x 1015 1 x 1015 1 x 1O15 1 x 1015 5 xlO 1 4 5 x 1O15 1 x 1O15 1 x io 1 5 1 x IO15 1 x IQ15

6 x 1018 8 xio18 2 x 1018 5 - 6 xlO 1 8 4 x 1018 2-5 x 1018 3 x 1019 1.5 xlO 1 9 8 x 1018 8 x 1018 I x 1018 I x 1019 6 x 1018 1-3 x 1019 1-3 x 1019 2 xlO 1 9 4 x 1020 8 xlO 1 9 8 xlO 1 9 2 x 1019 2 x IQ19

>1000 1140 1000 1000 1000 1000 1050 1050 1090 850 850 900 800 900 900-1000 1050 750 1020 1120 800-900 850

1 5 2 2 8 2 10 10 5 1200 1200 5 5 30 2-20 5 300 1 3 300 1200

I

I

|

1

Ref

[1] [2] [3] [3] [4] [5] [6] [6] [4,7] [8] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] I [8]

C

CONCLUSION

1.

Ion doses below about 1013 cm"2 produce activities close to 100% for both n-type and p-type dopants.

2.

Doses in the range 1014 cm"2 to 1015 cm"2 are activated at the 100% level for p-type dopants (Be, Mg, Zn, Cd) but activities are only of the order of 1% to 10% for n-type dopants (Si, Se, Sn).

3.

Maximum carrier concentrations achieved are around 1019 cm"3 and approaching 1020 cm"3 for n-type and p-type layers respectively, but these are achieved mostly after very high temperature anneals.

4.

More typical and easily attainable values are 2 x 1018 cm"3 and 2 x 1019 cm"3 for n-type and p-type material, respectively.

REFERENCES [1] [2] [3] [4] [5] [6]

D.E. Davies, PJ. McNaIIy5 PJ. Lorenzo, M. Julian [ IEEE Electron. Device Lett. (USA) vol.3 (1982) p.25] M. Kuzuhara, T. Nazaki, H. Kohzu [ J. Appl. Phys. (USA) vol.58 (1985) p. 1204 ] H. Kohzu, M. Kuzuhara, Y. Takayama [ J. Appl Phys. (USA) vol.54 (1983) p.4498-5003 ] K.K. Patel, R. Bensalem, M.A. Shahid, BJ. Sealy [ Nucl. Instrum. Methods Phys. Res. B (Netherlands) vol.7/8 (1985) p.418-22 ] NJ. Barrett, J.D. Grange, BJ. Sealy, K.G. Stephens [J. Appl. Phys. (USA) vol.56 (1984) p.3503-7 ] BJ. Sealy, R. Bensalem, K.K. Patel [ Nucl. Instrum. Methods Phys. Res. B (Netherlands) vol.6

[7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18]

(1985) p.325-9 ] M.A. Shahid, R. Bensalem, BJ. Sealy, P.N. Favennec, M. Gauneau [ Nucl. Instrum. Methods Phys. Res. B (Netherlands) vol.30 (1988) p.531-9 ] SJ. Pearton et al [ J. Appl. Phys. (USA) vol.65 (1989) p. 1089-98 ] NJ. Barrett [ Mater. Res. Soc. Symp. Proc. (USA) vol.35 (1985) p.451-6 ] R. Szweda, M.S.M. Lamb, R. Blunt [ Chemtronics (UK) vol.2 (1987) p.89-92 ] K.K. Patel, BJ. Sealy [Appl. Phys. Lett. (USA) vol.48 (1986) p.1467-9 ] A N M . MasumChoudhury, CA. Armiento [Appl. Phys. Lett. (USA) vol.50 (1987)p.448-50 ] H. Shen, Z. Zhou, H. Xu, G. Xia, S. Zou [Appl. Phys. Lett. (USA) vol.61 (1992) p.2093-5 ] G. Landgren, W.H. van Berlo [ J. Appl. Phys. (USA) vol.63 (1988) p.2783-6 ] D.E. Davies, PJ. McNaIIy [Appl. Phys. Lett. (USA) vol.44 (1984) p.304 ] NJ. Barrett, J.D. Grange, BJ. Sealy, K.G. Stephens [ J. Appl. Phys. (USA) vol.57 (1985) p.5470-6] S.S. Kular, BJ. Sealy, Y. Ono, K.G. Stephens [ Solid State Electron. (UK) vol.27 (1984) p.83-8 ] S.J. Pearton [ Int. J. Mod. Phys. B (Singapore) vol.7 (1993) p.4687-761 ]

19.5 Controlled atmosphere annealing of GaAs BJ. Sealy August 1995

A

INTRODUCTION

A problem with using a thin film encapsulating layer to protect the surface of GaAs during annealing is that there is, with few exceptions, a difference in expansion coefficient between the two materials. This causes significant strain near the surface and can, therefore, affect the result of an annealing cycle. For example, it can cause an increase in diffusivity of some impurities and can also change the degree of electrical activity for silicon implants [I]. A method of overcoming this problem is not to use a capping layer, but to heat in an atmosphere which will limit the loss of arsenic. This Datareview discusses some of the advantages and disadvantages related to this method of annealing. B

DISCUSSION

From the thermodynamic viewpoint, the ideal way of annealing GaAs is in an ambient containing arsenic, such as AsH3, or a mixture of arsenic vapour and hydrogen (As-H2). However, this method is hazardous and problems with gas purity have also been noted. An additional problem associated with annealing in an overpressure of arsine is the large thermal mass of the required furnace which means that short time anneals are not easy to perform, and hence diffusion of acceptors will occur [2]. In general, for any capless annealing technique loss of acceptor atoms to the surface is likely above about 900 0 C. See, for example, the results of Barrett et al [3] who annealed zinc implanted GaAs in arsine at 8500C and 8800C for twenty minutes and found that at least 50% of the dose (1015 cm"2) was lost through the surface. There are some commercially available RTA systems [4,5] which are designed to allow treatment in AsH3, but most capless techniques tend to be of the proximity type, in which two GaAs wafers are placed face to face. The little free space between the wafers acts to contain evaporating arsenic which creates an effective overpressure to limit decomposition. The technique does not produce perfect surfaces and often the periphery of wafers is degraded visibly after annealing. Proximity annealing also can create microscratches and baked on contamination of the wafer surface. In addition, loss of arsenic at the higher temperatures coats both the quartzware and lamps and can adversely affect the repeatability and uniformity of the process [6]. One of the earliest papers to compare capless (proximity) with capped annealing was by Kuzuhara et al [7], who identified annealing conditions for which decomposition was a minimum. They concluded, for example, that a GaAs capping wafer gave superior results to the use of a silicon capping wafer. They found, however, that silicon wafers allowed anneals up to 950 0 C for two seconds. Their results for silicon implants in the dose range 3 x 1012 cm"2 to 3 x 1013 cm"2 showed evidence of the amphoteric behaviour of silicon above about 9000C. There have been many other publications which demonstrate the efficacy of capless annealing for silicon implants in GaAs with or without an arsenic containing overpressure (see [2,5-12]). For example, Jackson et al [9] demonstrated a multi-wafer arsine ambient RTA system for MESFET fabrication, whilst Ehrenheim and Vidimari [11] showed that too low an arsenic overpressure produced p-type layers

from silicon implants. It was necessary to use high pressures (2 torr) of arsenic to completely control surface decomposition and to generate a non-diffused, n-type silicon profile. C

CONCLUSION

Capless annealing is a commonly used technique for processing ion implanted GaAs. It does, however, have a number of drawbacks. The ultimate way forward is the development of an ideal encapsulant which can protect wafers throughout the processing schedule. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] II1] [12]

Y. Saito [ J. Appl. Phys. (USA) vol.71 (1992) p.3544-61 ] SJ. Pearton [ Solid State Phenom. vol. 1/2 (1988) p.247-80 ] NJ. Barrett, J.D. Grange, BJ. Sealy, K.G. Stephens [J. Appl. Phys. (USA) vol.57 (1985) p.5470-6 ] S.S. Gill [ Phys. Technol. (UK) vol. 17 (1986) p.245-53 ] SJ. Pearton [ Int. J. Mod. Phys. B (Singapore) vol.7 (1993) p.4687-761 ] R. Singh [ J. Appl. Phys. (USA) vol.63 (1988) p.R59-Rl 14 ] M. Kuzuhara, H. Kohzu, Y. Takayama [ Mater. Res. Soc. Symp. Proc. (USA) vol.23 (1984) p.651-61] SJ. Pearton, J.M. Poate, F. Sette, J.M. Gibson, D.C. Jacobson, J.S. Williams [ Nucl. Instrum. Methods Phys. Res. B (Netherlands) vol. 19/20 (1987) p.369-80 ] T.N. Jackson, J.F. Degelormo, G. Pepper [ Mater. Res. Soc. Symp. Proc. (USA) vol.144 (1989) p.403-8 ] O. Paz [Mater. Res. Soc. Symp. Proc. (USA) vol.144 (1989) p.385-90 ] A. Ehrenheim, F. Vidimari [ Inst. Phys. Conf. Ser. (UK) no. 106 (1989) p.557-61 ] T.E. Kazior, SK. Brierley, FJ. Piekarski [ IEEE Trans. Semicond. Manuf. (USA) vol.4 (1991) p.21-5]

19.6 Compensation mechanisms in GaAs at high dopant fluences BJ. Sealy August 1995

A

INTRODUCTION

This Datareview will summarise briefly the explanations for the observed electrical compensation in ion implanted GaAs. B

DISCUSSION

Problems with electrical compensation occur only for n-type dopants, since for most p-type implants all atoms are electrically activated and the concentration of acceptor atoms equals the concentration of holes. For n-type dopants (Si, S, Se, Sn, Te) implanted to doses equal to or greater than 1013 cm"2, not all atoms appear to be electrically activated, even though the majority occupy lattice sites. The explanation for such behaviour is unclear despite the fact that this problem was first encountered in the 1970s. Thus, it has often been observed that 1020 atoms cm"3 occupy lattice sites and yet the electron concentration is only about 1018 cm"3 [I]. Thus, substitutionality is not the only prerequisite for electrical activation. There have been many studies of n-type implants in GaAs over the past twenty years which have considered this problem. Perhaps the favoured explanation is that complex defects are formed between the dopant atoms and neighbouring vacancies, such defects being deep acceptors. There are many reports based on photoluminescence and, more recently, several papers on extended X-ray absorption fine structure (EXAFS) measurements which indicate that complex defects with vacancies and donor atoms form in implanted GaAs. These observations are also consistent with the results of annealing kinetics studies [2]. The previous comments apply to all the above mentioned donor atoms. The situation with silicon is complicated because it is amphoteric, i.e. it is a simple donor when occupying a gallium lattice site and an acceptor when occupying an arsenic site. Above a certain atomic concentration (dose), the percentage activation saturates. It is argued that when the silicon concentration is raised above about 1018 cm"3, some of the silicon atoms start to occupy arsenic vacancies (Si x J rather than gallium vacancies (SiGa). This causes self-compensation, but additionally, other defects are likely to form such as neutral pairs (Si^-Si^) and complex defects with vacancies, (SiGa-VGa) and (Si^-VGa). Thus the explanation for electrical compensation in silicon implanted GaAs is complicated, and, as yet, the role each of the above- mentioned defects plays in this process is not clear. C

CONCLUSION

In general, electrical compensation does not occur for p-type implants in GaAs, since close to 100% electrical activity is obtained. However, the electrical activation of n-type dopants saturates at a carrier concentration above about 1018 cm"3 due to electrical compensation. The explanation is complicated involving complex defect formation between the implanted atoms and vacancies and, in the case of silicon, occupancy of both lattice sites (amphoteric behaviour) causing

self-compensation.

REFERENCES [1] [2]

SJ. Pearton [ Int. J. Mod. Phys. B (Singapore) vol.7 (1993) p.4687-761 ] BJ. Sealy [ Datareview in this book: 19.1 Ion implantation in GaAs: Overview ]

CHAPTER 20 EXPLOITATION OF GaAs IN MICROWAVE AND HIGH SPEED DIGITAL CIRCUITS 20.1 20.2 20.3 20.4 20.5 20.6 20.7 20.8

GaAs MESFET: discrete, power and MMIC microwave devices GaAs MESFET: digital and opto-ICs High electron mobility transistors GaAs in bipolar integrated circuits GaAsMMWICs The transferred electron effect Back-gating in GaAs MESFET integrated circuits GaAs IMPATT diodes

20.1 GaAs MESFET: discrete, power and MMIC microwave devices M. Feng, PJ. Apostolakis and W-H. Chang December 1995

A

INTRODUCTION

The first transistor, made of Ge3 was the point contact bipolar transistor discovered on 16 December 1947, and reported in 1948 by J. Bardeen and W. H. Brattain [1,2]. The identification of minority carrier injection was first understood by Bardeen [4]. Subsequently, the p-n junction transistor was reported by Shockley in 1949 [3]. It is worth mentioning that Bardeen held the basic patent (filed before the transistor patent on 26 February 1948) on the use of the inversion layer in the field effect transistor (MOSFET) [2]. The GaAs MESFET (metal semiconductor field effect transistor) concept was first introduced by C. Mead in 1966 [5]. The first microwave GaAs MESFET, fabricated by Hooper in 1967, achieved a unity current gain cutoff frequency, ft = 3 GHz [6]. Although there are many different methods used to prepare the active channel of GaAs MESFETs, such as LPE, VPE, MBE, MOCVD and ion implantation, it is most commonly prepared by the ion implantation technique [7], followed in popularity by epitaxial methods such as metal organic chemical vapour deposition (MOCVD) [8] and molecular beam epitaxy (MBE) [9]. There are significant cost savings with ion implantation compared with other techniques. The first ion implanted GaAs MESFET was reported in 1973 by Hunsperger [10], and subsequently by Ohata et al [11] and Higgins et al [12]. However, ion implanted GaAs MESFETs were not widely used at this time because of the lack of reproducible semi-insulating GaAs substrates, and appropriate annealing methods. In 1978, Swiggard et al [13] reported the use of substrates grown by the high pressure liquid-encapsulated Czochralski (HPLEC) technique. Feng et al [14] reported the electrical and chemical profiles of Si implanted in semi-insulating GaAs substrates annealed using both a SiO2 cap and without a cap using AsH3 over pressure capless anneal methods. The combination of the understanding of annealing methods, and HPLEC substrate growth was a major breakthrough in developing direct ion implanted GaAs MESFETs for microwave and digital IC applications. This work has stimulated much research in the search for higher speed devices [15-18] for both low noise and power applications. Since 1990, GaAs MESFETs have achieved ft >100 GHz [19-21], and have become the basic building blocks used in commercial GaAs foundries [23] for monolithic microwave integrated circuits (MMICs) [22] and digital integrated circuits (DICs), with an integration level of over a million transistors [24,25]. Since 1994, GaAs MESFETs have enjoyed an increase of 40% growth per year, due to expanding applications. This includes use in cellular phones, pagers, other wireless communication products, direct broadcasting systems (DBS), collision avoidance radar, global positioning system (GPS) and phased array radar, as well as fibre optic drivers and receivers. As wireless products continue to enter the marketplace, the low-cost, manufacturable ion implanted MESFET will be the technology of choice.

B

GaAs MESFET MMIC AND OEIC PROCESS SEQUENCE

The typical MESFET-based MMIC or OEIC is fabricated with a nine mask-level process [26-29]. A cross-section of the process is shown in FIGURE 1. The transistor regions are defined by a mesa etching isolation process. Layered metal (gold-germanium, nickel, gold) ohmic contacts are deposited and alloyed in a hydrogen rich atmosphere. The typical 0.5 ^m gate is defined by optical photolithography.

Via Metal 2 Metal 1 Ohmir

Gate

Ohmic

N-channei

GaAs Semi-insulating Substrate

FIGURE 1. A cross section of GaAs MESFET-based MMIC and OEIC process.

Gates smaller than 0.25 ^m are defined by direct E-beam lithography. A typical gate cross section of an E-beam defined gate of Lg = 0.12 \im is shown in FIGURE 2. The detailed e-beam processing steps for fabricating 0.1 ^m gate GaAs MESFETs are described in detail in [30,31]. The titanium-platinum-gold Schottky gates are deposited after a gate recess step which sets the Isat and V,

FIGURE 2. Typical 0.1 ^m gate in cross-section.

The circuit elements are then interconnected with a metal layer similar in composition to that of the gate layer. The interconnection layer is followed by a low stress silicon nitride film, used for surface passivation, anti-reflection coating, capacitors and electrical isolation. The final metal interconnect layer follows a two-step via and bonding pad etch. The complete process sequence is shown in FIGURE 3. Electrical qualification steps taken periodically are an integral part of the process. Process lots which do not meet specification at particular test points during the fabrication process are scrapped. Ion Implantation and CAT Anneal

Mesa Isolation

Ohrnlc Metal and Alloy

Gate Metal

Metal 1

Silicon Nitride Deposition

AIr Bridge

Final Passivation and Scribe Street

Via and Pad Etch

Metal 2

FIGURE 3. Process flow chart for MMIC and OEIC MESFET process.

C

GaAs MESFET TOPOLOGY

Many high-frequency performance characteristics are significantly affected by the gate of the MESFET. Simple changes in device topology can enhance the FET speed and noise performance by reducing the parasitic elements. There are two core types of FET layouts, namely, 'T' and 'inter-digitated' gate [32,33]. SEM micrographs of a coplanar 'T' gate and 'inter-digitated' gate configuration are shown in FIGURE 4. The standard CT' gate of 0.25 x 200 (2 x 100) ^m

FIGURE 4.. Coplanar T (left), 'inter-digitated' (right) gate low noise MESFET.

(FIGURE 4 (a)) employs a single input with two opposing gate 'fingers'. The typical gate finger widths range between 10 and 100 |im. The maximum gate width is limited to 200 \im using this method, due to yield and gate resistance considerations. The 'inter-digitated' gate is more complex to fabricate, requiring air bridge technology, but overall is better. It is commonly used in industry because of its ability to scale total gate width and reduce source resistance and input gate resistance using the parallel gate effect. A typical low noise 0.25 x 300 (4 x 75) |um MESFET, as fabricated using an E-beam direct write technique, is shown in FIGURE 4 (b). The drain to source spacing is 3 |um, and the source is air-bridge connected. Both devices are fully passivated by silicon nitride. The topology of a typical power FET is dictated by its functional requirements: specifically, by increasing the gate width using parallel groups of unit cells [34]. FIGURE 5 shows a 10 GHz power MESFET with a unit cell of 600 |um (4 x 150 \xm) gate. A total gate width of 2400 jim is achieved using 4 unit cells in parallel. This device employs backside vias, not air bridges, to connect the source to the backside ground in an effort to reduce the source inductance. A power combiner can be used to both increase the power level, and provide the input and output match.

FIGURE 5. X-Band Power MESFET for 2400 mm gate width. Each unit cell has 4 x 150 ^m gate width.

D

INTRINSIC SPEED MODEL OF InGaAs MESFETS AND p-HEMTS

A simple model to determine the speed limitations of sub-micron gate length Ga 10 Jn x As channel MESFETs and HEMTs based upon fundamental materials properties, rather than the layer structure (2-DEG effect), was proposed by Feng et al [35] in 1990. In this case, we refer to speed as the microwave ft performance.

(1)

(2)

(3)

In these equations, Leff is the effective gate length of the FET; Ctot is the total capacitance, including parasitic effects due to pads, layout geometry and the source drain and gate metallization. The energy difference between the F-L conduction band minimum is denoted by AE1x. Based on EQNs (1) - (3), we conclude that ft is directly proportional to the square root of AE1x. Hence, we can estimate that the maximum speed of the GaAs MESFET should be about 10% faster than for the GaAs HEMT [36] because the conduction band energy separation between the P and L minima, AE1x, for the GaAs MESFET is 0.33 eV, while for the GaAs HEMT is only 0.29 eV. Based on the experimental extrinsic ft = 110 GHz for 0.15 |im gate GaAs MESFETs, the value of AE1x and the electron effective mass OfGa1^xInxAs (TABLE 1), we can predict the extrinsic ft limitation for a typical 0.1 ^m gate MESFET or p-HEMT as a function of the indium mole fraction, FIGURE 6. Also shown are state-of-the-art published extrinsic ft results for Ga1-JnxAs channel MESFETs and HEMTs from TRW, Hughes, GE, NTT, the University of Illinois and the University of Michigan [37,42]. The gate lengths of these devices vary from 0.1 \im to 0.2 urn. As can be seen in FIGURE 6, the experimental ^ data agree very well with this simple prediction. The intrinsic limit, based upon parasitic de-embedding [36], is also shown in FIGURE 6. TABLE 1. Material properties OfInxGa^xAs. GaAs

In053Ga047As

InAs

E g (eV)

1.42

0.76

0.36

AE rL (eV)

0.33

0.55

0.9

Hi6Vm0

0.067

0.051

0.021

Ft (GHz)

MESFET HEWT Extrinsic Limit Intrinsic Limit

Indium Concentration, x FIGURE 6. Extrinsic and intrinsic ft OfInxGa^xAs MESFETs and HEMTs versus indium concentration.

E

LOW NOISE GaAs MESFET AND MMICS

El

Ion Implanted GaAs MESFET Structure

Gallium arsenide MESFETs are fabricated by the direct ion implantation of silicon and beryllium into semi-insulating (100) GaAs substrates, the former to form the active channel and the latter penetrating to the rear of the channel to ensure better pinchoff voltage characteristics. The low-noise channel is formed by three distinct implants. The Si channel implant, with a penetration depth of 1500 A, achieves a peak n-type carrier concentration of 1 x 1018 cm"3; a Si implant with a penetration of 600 A achieves a peak n+-type carrier concentration of 2 x 1018 cm"3 and the deep p-type implant of 50 keV is used to improve the pinchoff characteristics. The controlled atmosphere (CAT) anneal technique is used to activate the implant [44,45]. This is a capless furnace anneal performed at 8500C under an arsine over-pressure, as shown in FIGURE 7.

FIGURE 7. CAT anneal furnace tube at the University of Illinois.

For a typical wafer lot, the average sheet resistance is 220 ohms/square and the percent standard deviation across a 3-inch wafer is 1.4%. The channel shape, as formed by ion implantation and anneal, is critical to the high performance of these devices. A typical electron concentration-depth profile is shown in FIGURE 8.

Concentration (cm"3)

Conventional Implant UIUC Low Noise Implant

Depth (jum)

FIGURE 8. Si implant low noise profile.

The drain current, for a typical 0.25 \im gate length MESFET at 300 K, as a fixnction of the drain-to-source voltage, with gate voltage steps of -0.2 V starting from +0.4V, is shown in FIGURE 9 (a). The typical drain current and transconductance at 300 K as a function of the gate voltage is shown in FIGURE 9 (b). The peak transconductance for the GaAs MESFET is 105 mS (350 mS/mm) at I^ = 167 mA/mm and the pinchoff voltage is Vp = -0.4 V.

NIESFET

ld, (mA) AG-1(InS)

MESFET

FIGURE 9. Current versus voltage profile (a) for 0.25 ^m gate x 300 ^m wide GaAs MESFET and (b) transconductance.

The current gain |H21| is calculated from the measured S-parameters, and the ft is determined by extrapolating |H211, using a slope of-6 dB/octave, to the unity gain point. The measured ft as a function of I^ is shown in FIGURE 10 (a). For a typical 0.25 ^m gate GaAs MESFET3 ft = 61 GHz3 nearly the same as for a p-HEMT (62 GHz) for a 0.25 |iim gate, as shown in FIGURE 10. A comparison between state-of-the-art GaAs MESFETs and GaAs HEMTs (not p-HEMT) from several sources as a function of gate length is shown in FIGURE 10(b).

P-HEMT MESFET

MESFET HEMT

Gate Length (jim)

Ids (mA)

FIGURE 10. (a) Unity current gain cutoff frequency, ft, for a MESFET and a p-HEMT. (b) State-of the-art ft comparison for MESFETs and HEMTs as a function of gate length.

Minimum Noise Figure (dB)

The noise figure and associated gain at 10 GHz as a fixnction of drain current at 300K of a typical ion implanted GaAs MESFET and GaAs p-HEMT are shown in FIGURE 11 (a). These measurements show a typical minimum noise figure of 0.7 dB with 12 dB associated gain occurring at 13.5 mA (or 45 mA/mm) for both the MESFET and p-HEMT. The noise figures are practically identical over the same I68 (bias) region. The noise figure as a function of frequency is shown in FIGURE 1 l(b). The noise figure of the MESFET is comparable to the p-HEMT [46-51].

Fmin p-HEMT Fmin MESFET

Frequency (GHz)

FIGURE 11. (a) Noise and associated gain for MESFET and p-HEMT at 18 GHz. (b) Minimum noise figure versus frequency.

Molecular beam epitaxy (MBE) has been used for GaAs FET channel formation since 1976 [52]. The typical MBE low noise structure is composed of a 1.5 ^m undoped p-type buffer layer with an impurity concentration < 1014 cm"3 and a 0.3 jim Si doped n-type active layer with a carrier concentration of 3 x 1017 cm"3. The transition distance for the electron concentration to drop from 1017 cm"3 to 1016 cm"3 is less than 500 A. In 1982, an MBE grown MESFET achieved a 1.5 dB noise figure with 10 dB associated gain at 10 GHz using a 0.6 x 300 ^m gate technology [53], comparable to the best reported ion implanted GaAs MESFETs. E2

Low Noise GaAs MESFET

Millimetre-wave device modelling begins with the extraction of an accurate small signal equivalent circuit model using least-square fits to measured s-parameter data from 45 MHz to 110 GHz. FIGURE 12 is an example of a small signal 0.25 ^im MESFET model. Noise parameters are then measured from 2 to 18 GHz using a Cascade Microtech noise parameter measurement system. An equivalent noise representation based on the Statz model with parasitic noise sources, in

addition to correlated gate and drain noise, is then developed. The coefficients for the model are extracted from the measured data using an algorithm developed at the University of Illinois. This noise model can then be extrapolated to millimetre-wave frequencies in order to accurately design matching networks for low noise amplifiers. The 33-75 GHz noise figure and gain measurement system will allow for the verification and fine tuning of the equivalent noise model at millimetre-wave frequencies.

.25 x 200 Mm MESFET Model 50% ldss a 40 mA

FIGURE 12. Example of 0.25 ^m x 200 ^m MESFET small signal circuit model.

An amplifier was designed, based on the high performance quarter micron gate ion implanted GaAs MESFETs developed at UIUC. The millimetre-wave circuits were fabricated on the pilot line at Raytheon ADC using a standard 0.25 |Lim gate MMIC process. A SEM picture of the five stage Ka-band low noise amplifier is shown in FIGURE 13(a).

Frequency [QHz)

FIGURE 13. SEM photograph offive-stageion implanted Ka-band MESFET LNA (a). On-wafer noise associated gain data for 10 samples of MESFET LNA (b).

The chip size in FIGURE 13(a) is 3.53 mm x 1.815 mm. The on-wafer millimetre wave performance for 10 samples of the MESFET-based LNA, biased with a single power supply of 3 V at I68 = 50% I ^ (84 mA) is shown in FIGURE 13(b). The power consumption for this circuit is 252 mW. The average on-wafer probe gain is 30 dB with a standard deviation of 0.6 dB, and the on-wafer probe noise figure is between 2 and 3 dB over the frequency range 2 7 - 3 3 GHz.

The input/output reflection coefficients are less than -10 dB at 33 GHz. These results compare well with Ka-band LNA results achieved using GaAs p-HEMT technology, which exhibit a 2.5 dB noise figure with 13 dB associated gain [54]. TABLE 2. Low noise amplifier technology characteristics. Company/Technology

Noise

Associated Gain

UIUC 5-stage LNA [54] 0.25 ^m Gate GaAs MESFET

2-3 dB

30 dB

TRW 1-stage LNA [55] 0.2 ^m GaAs p-HEMT

3.5-4.5 dB

10 dB

Hughes 2-stage LNA [56]

3.5 dB

17 dB

0.25 ^m GaAs p-HEMT

F

POWER GaAs MESFETS AND AMPLIFIERS

Fl

The Epitaxial Power GaAs MESFET and MMICs

The epitaxial MESFET structure has three major advantages over the ion implanted MESFET. (1) The VPE [57-60], MOCVD [61] and MBE [62] growth techniques can produce N+ concentrations up to 6 x 1018 cm"3, while the N+ layer by ion implantation is limited to 2 x 1018 cm"3. (2)

The epitaxial MESFET has a sharp transition between the channel layer and the semiinsulating substrate. Therefore, it can provide better power performance than the MESFET, provided material uniformity and reproducibility are available. Additional channel epitaxial layers with 'graded' or 'delta' doping profiles can improve both the pinch-off and the linearity of the current and voltage characteristics [60].

(3)

The epitaxial layer growth can support heterostructure FETs (HFETs) such as AlGaAs/GaAs [63]; InGaAs/GaAs [64,65]; and InGaP/GaAs [66]. These can exhibit improved device performance over conventional ion implanted MESFETs.

The typical structure of a GaAs power MESFET, grown by vapour phase epitaxy, consists of a 0.4 |nm thick active channel layer with a carrier concentration of 1.6 x 1017 cm"3 and a 4 \im undoped buffer layer grown on a Si-doped GaAs substrate. The \\im * 2400 \xm gate power FET, as shown in FIGURE 5, has demonstrated an output power of 2 watts at 10 GHz, with an associated gain of 5 dB and a power added efficiency of 31%. The output power and power added efficiency of a single (600 \im) and a multi-cell power FET (2400 fim) plotted versus input power, at a Vds = 9 V, is shown in FIGURE 14 [58] . There are two issues concerning epitaxial wafers for power applications. First, material availability is a problem for both end-users and suppliers. Currently, there exist two major MBE epitaxial suppliers (Quantum Epitaxial, PicoGiga) and four MOCVD epitaxial suppliers (Sumitomo Electronics, Kopin, Sumitomo Chemical and Epitronics). Securing large quantities (100-500 wafers per week) of qualified epitaxial wafers for a given device structure is difficult. Second, the

Pout (dB)

PAE (%)

Frequency (GHz) FIGURE 14. Output power and power added efficiency (PAE) of single and multi-cell FETs.

cost of each epitaxial wafer is still very high ($500 to $1,000 above the substrate cost). Hence, a manufacturable, low cost epitaxial technology needs to be developed to realize the advantages of the epitaxial structure for power MMIC applications. F2

Ion Implanted Power GaAs MESFETs

The channel layer for ion implanted power MESFETs is formed by a deep implant at an energy of 250 - 300 keV and a dose level of 8 x 1012 cm"2. The typical anneal temperature is 890 0 C. The peak carrier concentration is approximately 2 x 1017 cm"3, and the typical Hall electron mobility is 3800 cm2/ Vs. GaAs power MESFETs have been made using ion implanted channel layers since 1981 [67,68]. The layout geometry consists of 4 parallel cells of FETs, FIGURE 5. A typical 1 |im gate length by 2400 \xm gate width device exhibits a transconductance of 280 mS (117 mS/mm). The microwave power performance at 10 GHz, consisting of output power and power-added efficiency (PAE), for a typical single cell (1 x 600 ^m) and four cell (1 x 2400 ^m) power MESFET, is plotted in FIGURE 15 as a function of input power at a drain voltage of 10 V, [67]. The single-cell device generates nearly 26 dBm (0.4 W or 0.66 W/mm) output power with 7.6 dB associated gain and 35% power added efficiency. The linear gain is 11.4 dB and the small-signal gain is 13.65 dB. The four cell device delivers 32 dBm (1.63 W or 0.68 W/mm) of output power with 6.9 dB associated gain, 35% power added efficiency and 9.7 dB linear gain. In FIGURE 16, the power added efficiency for a 4 cell FET is above 40% for a drain voltage in the range 5.5 V - 8.5 V. The peak power added efficiency for this device is 42.5% at 7.5 V, with an output of 31.5 dBm and 7.3 dB associated gain.

Pout (dBm) single cell

PAE (%) single cell PAE (%) four cell

PAE (%)

Pout (dBm)

Pout (dBm) four cell

Input Power (dBm)

FIGURE 15. Output power and power added efficiency for a 600 and 2400 jum gate MESFET.

PAE (%)

Pout (dBm)

Over the past 10 years, considerable improvements have been made in PAE for ion implanted GaAs power MESFETs. A manufacturable X-band power MMIC fabricated by direct ion implantation has been developed by Hughes. This chip delivers 4 W with 35% power added efficiency at 10 GHz. Ka-band power MMICs fabricated by direct ion implantation have also been developed by Hughes. This chip delivers 250 mW with a power added efficiency exceeding 20% at 38 GHz. In the commercial area, ITT has developed a 0.6 ^m gate self-aligned gate power MMIC process for low voltage RF power amplifier applications. The ITT power MMICs3 operating at 900 MHz, deliver 1 W with 35% PAE at 3V and 58% PAE at 6 V.

FIGURE 16. Output power, power added efficiency, and gain for a 2400 ^m gate GaAs MESFET as a function of drain to source voltage.

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20.2 GaAs MESFET: digital and opto-ICs M. Feng, PJ. Apostolakis and W-H. Chang January 1996

A

INTRODUCTION

In order to accurately simulate and predict digital circuit performance, large-signal models which can reliably predict MESFET behaviour are essential. Large-signal models are not used very often in MMIC design, since they operate in the linear region. A typical low noise amplifier usually contains only a few transistors and is modelled using different small-signal models at their corresponding bias points. However, small-signal models are not sufficient for circuits containing ten or more transistors, where each may be biased differently. For digital circuits as well as for more 'traditional' circuits like operational amplifiers, transimpedance amplifiers and limiting amplifiers, large signal models are necessary for the nonlinear simulations which these designs require. B

LARGE-SIGNAL MODELLING TABLE 1. List of large-signal MESFET models supported by selected commercial simulators. Simulator/Extractor

Model Category

XTRACT 3.6/ mwSPICE 3.51

IC-CAP 4.32

MDS 6.202

EEsofIV comms 5.O1

Eldo3

HSPICE4

Curtice

Curticet

CGaaslt CGaas2J CGaashfJ

Curtice Quadratic^ Curtice CubicJ Advanced Curtice

CURTICE2T CURTICE3J CURTICE20t CURTICE30J

Level 6f

Level 3 SAT=Of SAT=IJ

Statz

Statz (Raytheon)

UCBGaas UGaashf

STATZ

STATZ STATZO

Level 8

Level 3 SAT=2

TOM*

level 9 update* level 9 update 2*

Level 3 SAT=3±

Materka

MATERKA

Tajima

TAJIMA HPRootFet

Root

HPROOT

Triquint

EEFET

EEFETl EEFET2

Others

User Defined

EEFET3

Symbolic Model (user defined) ± t Curtice Quadratic, J Curtice Cubic, Triquint's Own Model (TOMl), * TOM-2 Model. 1 from HP-EEsof,2 from HP, 3 from ANACAD,4 from Meta-Software.

Empirical equation-based large-signal models are popular in circuit simulation due to their efficiency and accuracy. Physical equation-based models are usually only suitable for single device simulation, due to their slow simulation speed. Several empirical large-signal GaAs MESFET models have been developed and are used in commercial software simulator packages. TABLE 1 lists several large-signal GaAs MESFET models available from some commercial simulators and model parameter extractors. Most of the programs support the Curtice models. The original Curtice model (Curtice Quadratic) [1] is based on the Schichman-Hodges IGFET (insulated-gate field-effect transistor) model [2], but later adopted a hyperbolic tangent function of Vds to better model the I^ versus V^ relation below the saturation region. Another major improvement is the inclusion of transit-time effects to account for the delay between changes of V88 and Ifc The Curtice Cubic model was proposed in 1985 [3], and uses a cubic approximation for I^ to simulate the non-square-law relationship between the saturation current and V88, as in EQN(I). I63 = (A0 + A1V1 + A2V12 + A3V13) tanh (Y V J

(1)

where I^ is the saturation current, V^ is the drain to source voltage, A 0 -A 3 and y are the fitting constants. V1 is the effective gate-source voltage, which includes the observed shift of the pinch-off voltage with VJ8, as defined in EQN (2). V1 = V88(I-T)[I + P ( V 4 8 0 - V J ]

(2)

P is the pinch-off voltage coefficient, V480 is the drain to source voltage at which A 0 -A 3 are evaluated, V88 is the gate to source voltage, T is the internal time delay of FET. The model proposed by Statz et al [4] improves upon the Curtice Quadratic model by introducing a factor, [ 1 + b (V83 - V1)], in the drain-source current equation to account for the non-quadratic behaviour of I^ with respect to (V88-V1), when V88 is both larger and not equal to Vt. The hyperbolic-tangent function is replaced by a simple series expansion to reduce the computation time. The modified equations for I48 are given in EQNs (3) and (4).

I =

s* " Vt)2 I l - (1 - ^Lf 1(1 3 1 + KV88 - V1) I J

I

8s " ^ (i + XV.) for V, * 1 + KVg8 - Vt) * * a

=

P(V

+

XVJ for 0 < V^ < 1 a

(3)

(4) K }

a is a factor in determining 'knee' voltages of I^ -V^ curves, b is an empirical coefficient to fit 1,J8 for Vg8- V, > 0.3, and X is the channel length modulation coefficient. McCamant et al proposed a new model [5] to address the insufficiencies in the Statz or Curtice

Cubic models. The fitting of the I-V curves near pinch-off and the output conductance variation with respect to V88 are improved by (1) making the pinch-off voltage, V0 a function of V^, as in EQN (5) to address the poor fit of I^ near pinch-off, and (2) providing a feedback factor in the expression for I&, as in EQN (6). V, ^V 4 0 -YV*

Ids =

i • «Vj*.

(5)

(6)

The term in EQN (6) is meant to address the decrease of the slope of I^ at higher values of current and voltage for some MESFETs. This model refinement was reported [5] to provide an improved fit for GaAs MESFETs compared to previous models. The GaAs FET Root model [6] is essentially a look-up table, large-signal model in which the measured and calculated nonlinear model functions are stored in tabular form as functions of two independent controlling terminal voltages (V88 and V418). Device data between measured points is interpolated by a two dimensional spline function. The simulation time can be reduced, and high accuracy can be achieved as long as the measured data is accurate and well behaved. FIGURE 1 shows the general flow-chart of a large-signal parameter extraction procedure. Device Data Collection (I-V, S-parameters, etc.)

Small-Signal Parameters and Parasitic Extraction (Cgs, Cgd, Cds, gm, Rd, Rg, etc.)

Large-Signal Parameter Extraction

Modify Parameter Values

Model Fit?

END FIGURE 1. MESFET model parameter extraction procedure.

Device data, including I^ - Y& curves, I^ - V^ curves, and s-parameters versus frequency across the bias range of interest need to be measured and recorded in the correct format for subsequent processing. The first data processing step entails removing parasitic parameters, usually assumed

to be constant over the range of bias conditions. A set of small-signal parameter values is then extracted at every measured bias point (V^ and V88). A 10 x 10 array of V^ and V88 points is usually chosen to represent the device under all possible operating conditions. Small signal parameters are then exported into a large-signal, model parameter, extraction program. Here, large-signal model parameters are optimized to give the best fit. The fit between the measured and simulated data can be determined by inspection or by using pre-defined error coefficients. Further optimization may be necessary if the desired fitting accuracy is not reached. TABLE 2 shows the Curtice Cubic model parameters extracted from a D-mode 0.25 ^m * 150 jum MESFET fabricated at the University of Illinois. The measured and modelled I-V curves are shown in FIGURE 2, where it is clear that the saturation region of this device (0.4 V - V88 < OV) is well modelled. TABLE 2. Curtice/Ettenberg large-signal model parameters extracted from a UIUC D-mode MESFET with gate length = 0.25 jim and 2 x 7 5 jim gate width. (Resistances are normalized with respect to gate width). Parameter

Unit

Value

Lg

H

6.23e-ll

I8

A

1.993e-13

Ld

H

6.99e-ll

1.349375

L8

H

4.56e-14

N V bi

V

0.593936

V8^

V

-0.4

Rg

Q.fim

411.76

V ^

V

2

Rd

Q. jim

101.9

R1n

Q. pun

le-6

R8

O.jim

240.91

Tau

s

8.45e-13

Cd8

F

4.38e-16

Beta

0.159407

0.5

Gamma

2.181209

Fc C 880

F

9.4e-16

Rd80

Q.^im

395.1564

C gd0

F

4.32e-16

V^0

V

0.3579

A0

1.8712e-4

V^ 0

V

4.224324

A1

3.0694e-4

Vt

V

-0.8081

A2

1.189e-4

C rf

F

1.0e-8

A3

3.33e-14

Rc

Q.\un

320.2217

The near sub-threshold region (Vg8 < -0.6V) is also simulated by the Curtice Cubic model. Most analytical equation-based models, such as the Curtice and the Statz models are not able to predict device behaviour in highly nonlinear regions such as the sub-threshold and the near-breakdown regions due to the inherent limits of empirical equations. However, such modelling is sufficient for analogue circuits in which transistors are operated in the saturation region. In oscillator circuits and digital circuits where transistors can swing across most of the measured bias area, and can be turned off during operation, accurate modelling in the subthreshold and near break-down regions is necessary in order to get precise simulations. FIGURE 3(a) and 3(b) show modelling, performed at UIUC, for Vitesse 8 x 40 \im HGaAsIII

Ids (mA)

squares: measured solid-lines: modeled

Vds (V)

FIGURE 2. Measured (squares) and modelled (solid-lines) drain currents for a UIUC 0.25 urn MESFET using the Curtice Cubic model.

depletion and enhancement type MESFETs. Rectangles and solid lines represent measured and simulated data, respectively. Both devices are well modelled by the Root model over the entire bias range. Vitesse 8x40um D-MESFET

Vitesse 8x40um E-MESFET

FIGURE 3. HP RootFET models for a 0.611 m 8 x 40 um Vitesse (a) D-mode and (b) E-mode MESFET.

C

OPTICAL RECEIVER DESIGN

In optical fibre communications, optical receivers are critical for converting the light signals into electrical signals for further processing (amplification, regeneration, and signal processing) in the electrical domain. The basic building blocks of an optical receiver are photodiodes and current-to-voltage converters. The current-to-voltage converters are needed as an interface between the photo-diodes and subsequent amplifying stages. Because of the good combination

of speed, dynamic range and noise performance, transimpedance amplifiers are the most widely used converters for this purpose. FIGURE 4 shows a block diagram of an optical receiver with extra amplification stages and an output buffer to drive an external load.

post

OUT

FIGURE 4. Simplified block-diagram of an optical receiver.

In the design of optical receiver circuits, the bias conditions for the active devices need to be optimized. FIGURE 5 (a)-(f) shows 3-D plots of parameter values in V88 -V^ space for a UIUC 0.25 |um 2 x 75 |iim D-mode MESFET. Such plots help to visualize the performance trends along two-dimensional bias conditions and are used as a reference in selecting bias conditions. In FIGURE 5(a), the maximum transconductance is seen for Vg8 = 0 V and values of V^ between 1 - 2 V. As expected, the transconductance depends more strongly on V88 than V^. Usually ion-implanted GaAs MESFETs exhibit a ^ peak immediately before the gate-drain Schottky-diode turn-on (Vg8 - 0.5 V). The V88 should be set below the turn-on of the gate Schottky diode, to prevent a large gate leakage current. The output impedance shown in FIGURE 5(b) increases sharply when the Vg8 is reduced towards the pinch-off voltage. This occurs as the conducting channel reduces in size, causing the resistance between the drain and source to increase. FIGURE 5(c) is the ideal single stage common-source transistor voltage gain under the condition of infinite load resistance, computed as the product of gm and R^8. The maximum ideal gain occurs at Vg8 = 0 and high V^. Conversely, the maximum current-gain cutoff frequency, as shown in FIGURE 5(d), occurs at V88 = 0 V and V68 = 0.5 V. The trade-off between maximum gain (FIGURE 5(c)) and cut-off frequency (FIGURE 5(d)) must be carefully considered with respect to the circuit performance requirements. The gain-bandwidth product, as shown in FIGURE 5(e), represents the product of the bandwidth and the voltage gain assuming single transistor operation. The gate-source capacitance (C88), as shown in FIGURE 5(f), and the transconductance (gjj, as shown in FIGURE 5(a), are usually the dominant components in defining the ft of the transistor. The C88 value is also important in the multiple-stage condition, where it represents the output loading capacitance of the previous stage. Reduction of this value is extremely important when the dominant pole of the previous stage lies at the output node. D

OEIC RECEIVER PERFORMANCE

A particular optical receiver array consists of four channels. Each channel has an MSM (metalsemiconductor-metal) photo-detector, a transimpedance amplifier, and three stages of post amplification with an output buffer for 2.5 Gb/s performance. MSM photo-diodes were used because of their compatibility with the standard MESFET process sequence. FIGURE 6 shows

(Slemens'ohm) Gm'Rds

FIGURE 5. Parametric plots of (a) transconductance; (b) output drain-source resistance; (c) ideal single-stage common-source gain; (d) current gain cutoff frequency; (e) gain-bandwidth product; (f) gate-source capacitance, as functions of both V^ and Vgg for a UIUC 0.25 ^m gate length by 2x75 ^m gate width GaAs MESFET.

a cross-section of an MSM monolithically integrated with a MESFET. Gate metal is used for both the gate of the MESFET and the fingers of the MSM. The latter are usually placed directly on the semi-insulating substrate to achieve a low dark current and a low breakdown voltage.

MESFET

S.I. GaAs Substrate

implanted active layer gate metal ohmic metal

FIGURE 6. Cross-section of monolithically integrated MSM and MESFET.

The chip size is 2.4 mm x 3.1 mm. All of the DC pads which provide power to the chip are located together along the edge of the chip.

FIGURE 7. SEM micrograph of a four-channel OEIC receiver fabricated at UIUC.

Complementary outputs are provided for each channel. This chip is compatible with both on-wafer high speed probing, and wire-bonding. MSM photodetectors are located on the left side of the chip to convert modulated optical signals into photocurrent, which is processed by the transimpedance amplifiers. FIGURE 8 shows the circuit diagram of a complete channel in the receiver array. Both eye-diagram and bit error rate (BER) measurements on the OEIC receiver circuit are

FIGURE 8. Circuit schematic of one channel in the receiver array.

Both eye-diagram and bit error rate (BER) measurements on the OEIC receiver circuit are performed by using a bit error rate tester (BERT) and a digital sampling scope (FIGURE 9(a)). The pattern generator creates a pseudo-random binary signal (PRBS) and modulates a high speed laser diode. The laser diode DC bias and temperature are maintained by a laser diode controller, capable of 3 GHz operation. A Cascade light wave probe (LWP), with an optical fibre core size of 50 jim, is used to couple the optical signal into the on-chip MSM detectors. Each MSM detector is 75 |im by 75 urn, and represents a trade-off between speed (capacitive loss of a large area MSM) and the coupling efficiency of a small area detector. During eye diagram measurement, the electrical signal at the output of the receiver is collected by the high speed microwave probe and sent to the digital sampling scope. The electrical signal is then fed to the error detector of the BERT tester and compared with the original signal from the pattern generator to determine the bit error rate in the optical link. FIGURE 9(b) shows the eye diagram for an optical signal measured on a UIUC four channel receiver circuit. Good eye opening and a BER of < 10"9 has been observed at 1 GHz indicating that the circuit can be used in high speed optical data links at or above 1 Gb/s. Enhancement-mode (E-mode) MESFETs can be used to achieve lower power consumption for the same high-speed performance as depletion-mode MESFETs. This is because a higher gain can be achieved at a lower I^ for E-mode devices. Enhancement-mode GaAs MESFETs can be fabricated by reducing the ion-implantation into the GaAs channel; shallow, low-dose implants raise the pinch-off voltage to above zero gate-source voltage.

BERT Error Detector Oscilloscope Clock Pattern Generator Source

LD

LD Controller

OEIC Rx

mw Probe Probe Card Power Supply

FIGURE 9. (a) Testing setup for optical receivers and (b) measured eye-diagram of an OEIC four-channel optical receiver with -16 dBm NRZ 27-l PRBS optical input at 1 Gb/s.

In FIGURE 10(a), a ^ of 42 mS is obtained at an I^ - 43 mA for a 150 |um depletion mode MESFET. In FIGURE 10(b), the same transconductance is obtained by a 150 jim E-mode device at an 1^-8 mA.

FIGURE 10. Gm and I* versus V88 for 150 |um (a) D-mode and (b) E-mode MESFETs.

However, the E-mode MESFET allows less of a V88 swing between the pinch-off voltage and the diode turn-on voltage. The process tolerance for the pinch-off voltage shift is consequently much smaller for E-mode devices and if such devices are to be used, circuits with a high tolerance of

pinch-off voltage variation must be carefully designed. Front end transimpedance amplifiers using E/D GaAs MESFET technology have been designed, fabricated and tested. A photograph of a fabricated E/D transimpedence amplifier (MOSIS/Vitesse foundry service) and the resulting eye diagram measured for this circuit at 1 Gb/s is shown in FIGURE 11.

FIGURE 11. (a) a transimpedance amplifier using E/D technology, (b) an eye-diagram of the same amplifier for 1 Gb/s operation.

A photograph of a fabricated optical receiver using an E/D MESFET process (MOSIS/Vitesse foundry service) is shown in FIGURE 12. Simulation and preliminary data indicate that the circuit should operate at 2 Gb/s with a power consumption of about 120 mW per channel.

FIGURE 12. Fabricated E/D optical receiver.

Next Page

REFERENCES [1] [2] [3] [4] [5] [6] [7] [8]

W.R. Curtice [ IEEE Trans. Microw, Theory Tech. (USA) vol.28 no.5 (1980) p.448 ] H. Shichman, D.A. Hodges [ IEEE J. Solid State Circuits (USA) vol.SC-3 no.3 (1968) p.285 ] W.R. Curtice, M. Ettenberg [ IEEE Trans. Microw. Theory Tech. (USA) vol.33 no. 12 (1985) p. 1383] H. Statz, P. Newman, LW. Smith, R.A. Pucel, H. A. Haus [ IEEE Trans. Electron Devices (USA) vol.ED-34 no.2 (1987) p. 160 ] A.J. McCamant, GD. McCormack, D.H. Smith [ IEEE Trans. Microw. Theory Tech. (USA) vol.38 no.6(1990)p.822] D.E. Root [Microwave J. (USA) (1991) p. 126 ] Ravender Goyal (Ed) [ High frequency analog integrated circuit design (Wiley-Interscience, 1995)] K. Lee, M. Shur, T.A. Fjeldly, T. Ytterdal [ Semiconductor device modelling for VLSI (Prentice Hall, Series in Electronics and VLSI, 1993) ]

20.3 High electron mobility transistors

Previous Page

R.H. Wallis October 1996

A

EVTRODUCTION

The use of modulation doping to achieve enhanced electron mobilities in AlGaAs/GaAs multilayer structures (reviewed in Datareviews 2.9 and 2.10) led to the proposal that improved device performance could be achieved if modulation doped material replaced the simple doped GaAs layer in a MESFET. The modulation-doped AlGaAs/GaAs multi-quantum well structure in which mobility enhancement was originally observed [1] was not well suited for inclusion in a MESFET5 because charge control could not be achieved throughout the whole layer structure. The subsequent demonstration of enhanced mobility at a single selectively doped AlGaAs/GaAs heterojunction gave a structure more amenable to device fabrication [2]. The first demonstrations that the electron density at the heterojunction could be controlled by a Schottky barrier, so forming a field-effect transistor, were reported in 1980 by Mimura et al [3] and D. Delagedelagebeaudeuf et al [4]. Mimura et al called their device a high electron mobility transistor (HEMT), a name which has been widely adopted, although the alternative name MODFET (from modulation-doped field effect transistor) introduced by Morkoc [5] is also common. The same basic device was also given other names by early workers such as TEGFET (two-dimension electron gas field effect transistor) and SDHT (selectively doped heterojunction transistor) although these are now rarely, if ever, used. The term HEMT will be used exclusively in this Datareview. Since this first demonstration of a transistor based on a selectively-doped heterojunction, the HEMT has been the subject of intense development. These efforts have led to the HEMT achieving the highest operating frequency of any type of transistor, extending the operating frequencies of three-terminal devices and transistor-based circuits up to 100 GHz and beyond. In addition, the superior noise performance of HEMTs compared to MESFETs has resulted in them being used for performance-critical applications at lower microwave frequencies. A particularly striking example is their widespread use in low-noise receivers for direct broadcast by satellite (DBS) television at around 12 GHz. More recently their high gain at microwave frequencies has led to the development of power HEMTs with high power-added efficiency, leading to their widespread use in mobile and satellite communications systems. The intense effort expended on the development of HEMTs since the early 1980s has resulted in several thousand publications describing their operation, fabrication and applications. Much of this material is already the subject of numerous reviews [6-15] giving more detail than space allows here. B

TYPES OF HEMT AND MATERIAL SYSTEMS

The earliest HEMTs were based on a single modulation-doped AlxGa1^As /GaAs heterojunction grown on a semi-insulating GaAs substrate. A cross-section through the basic structure of the

device is shown in FIGURE 1. The device channel is created by electrons which transfer from the donors in the wide-gap AlxGa1^As layer to form a two-dimensional electron gas (2DEG) on the GaAs side of the heterojunction. To reduce Coulomb scattering and so enhance electron mobility, the first few atomic layers OfAlxGa1^As immediately adjacent to the heterojunction are generally left undoped. In a HEMT such a spacer layer is much thinner (typically 2 - 5 nm) than in modulation-doped structures designed for very high mobilities, since thicker spacer layers reduce the density of electrons which transfer from the AlxGa1^As into the 2DEG. The rest of the AlxGa^xAs layer (known as the supply layer), typically 30 to 60 nm thick, is uniformly doped at 1 - 2.5 x 1018 cm"3. The heavily doped (typically 2 - 4 * 1018 cm "3) GaAs cap layer above the AlxGa^xAs supply layer serves to facilitate the fabrication of good ohmic contacts for the source and drain and to reduce the series resistance between these contacts and the gated part of the channel. The gate is generally placed on the AlxGa1^As supply layer by etching a recess through the GaAs cap. To avoid deleterious effects due to trapping associated with DX centres in the AlxGa1^As layers (see Datareview 7.5 by Maude), the Al mole fraction x does not generally exceed 23%.

SOURCE

DRAIN

n-type GaAs Cap layer n-type AIGaAs supply layer Undoped AIGaAs spacer layer Undoped GaAs Two-dimensional electron gas

Semi-insulating GaAs substrate

FIGURE 1. Cross-section of a conventional AlxGa^xAsZGaAs HEMT.

The major weakness of this basic structure is that the conduction band discontinuity AEC at the AlxGa^xAs /GaAs heterojunction is not large enough to produce an efficient transfer of electrons from the AlxGa^xAs supply layer into the GaAs3 resulting in the maximum electron density in the 2DEG channel being not much greater than about 1 x 1012 cm"2. Excess charge remaining in the AlxGa1^As supply layer screens the 2DEG from the gate and forms a parasitic AlxGa1^As MESFET in parallel. Because the mobility and saturated drift velocities of electrons in heavily doped AlxGa^xAs are low, the effect is to reduce the device transconductance until the gate bias is sufficient to deplete the residual charge in the AlxGa^xAs. A larger value of the conduction band discontinuity AEC, and hence more efficient charge transfer, can be achieved by using an A^Ga^xAs/InyGa^yAs heterojunction in place of the AlxGa^xAs /GaAs

heterojunction. Unlike AlxGa^xAs, IriyGa^yAs is not lattice-matched to GaAs. However the work of Matthews and Blakeslee [16] has shown that a lattice-mismatched layer can elastically accommodate the strain without forming misfit dislocations as long as its thickness remains below a critical value. The use of an elastically strained InxGa1 _yAs layer to form the channel of a HEMT was first suggested and demonstrated by Rosenberg et al [17], who described such a layer as 'pseudomorphic'. Since then the 'pseudomorphic HEMT' structure (commonly abbreviated to PHEMT)5 as shown in FIGURE 2, has largely replaced the original structure shown in FIGURE 1. The earliest work used 15 nm thick pseudomorphic layers with compositions of In0^5Ga0 85As. The trend has been towards thinner (10-12 nm) layers with higher (20 - 30%) In mole fractions.

GaAs Cap

40 nm

4 x 1018 cm 3

AI 023 Ga 077 As

25 nm

2 x 1018 cm 3

Al0 23 Ga 0 77As

4 nm

Undoped

In0 2 Ga 0 8As Channel

12 nm

Undoped

GaAs Buffer layer

0.5 jim

Undoped

Semi-insulating GaAs Substrate FIGURE 2. Typical layer structure for an A^Ga^As/IriyGa^yAs/GaAs pseudomorphic HEMT.

Several other refinements to the structures shown in FIGURES 1 and 2 are commonly used. One example is the use of delta-doping instead of uniformly doping the AlxGa1^xAs supply layer. Here all the donors are introduced into a single plane close to the AlxGa1^AsZGaAs or AlxGa1^AsZ ItIyGa1^yAs heterojunction. This results in better transfer of electrons into the channel and also has the advantage of placing the Schottky barrier gate on undoped material, giving better breakdown characteristics. Another example is the introduction OfAlxGa^xAsZGaAs superlattices in the GaAs buffer layer, which is claimed to give improved material quality in the active device layers and to improve carrier confinement by reducing injection into the substrate. Although AlxGa1^As is most commonly used as the wide gap semiconductor, the use of GaInP (with a composition lattice-matched to GaAs) as an alternative has also been reported [18]. The principle of the HEMT has also been transferred to other materials systems grown on substrates other than GaAs. In particular, the highest frequency devices have been achieved in the AlInAsZInGaAs system grown either lattice-matched or pseudomorphically on InP substrates [19]. Only GaAs-based devices are considered in this Datareview.

C

HEMT OPERATION

The operation of a HEMT is basically similar to that of a MESFET5 the major difference being that the presence of the heterojunction modifies the distribution of the electrons in the structure and so changes the charge control by the gate electrode. The charge control behaviour of a HEMT was first derived by Delagebeaudeuf and Linh [20]. Since the heterojunction holds the electron gas in the channel at approximately constant distance from the gate electrode, the charge density in the channel varies approximately linearly with gate voltage, similar to a Si MOSFET [21]. The physical reasons for the improved performance of HEMTs compared to MESFETs are essentially two-fold. Firstly, by separating the electrons from the ionised donors, modulation doping gives improved electron transport in the channel, with a higher effective saturation velocity. Secondly, because the electrons are confined by heterojunction barriers it is possible to make very short gate length devices while still maintaining a high aspect ratio of the gate length, Lg, to the separation between the gate and the channel. This ratio must be at least five (and ideally more) to avoid 'short-channel effects' such as increased output conductance, which reduces device gain. By making the AlGaAs layer beneath the gate very thin ( - 2 0 nm), a good aspect ratio can be maintained even for gate lengths as short as 0.1 ^m. This cannot be achieved in MESFETs, where the channel is defined only by the potential arising from the charged donors. For any field-effect transistor the intrinsic current-gain cut-off frequency, fT, is given (approximately) by g^TtC^, where g^ is the transconductance and C88 is the gate to source capacitance. In a HEMT, if only the electrons in the channel were modulated by the gate potential and all electrons moved at their saturated velocity vsat, then fT would be given by: fT

=

g m / 2TiC68* vsat/2TTLg

(1)

In reality this is too simplistic: as the gate to source bias V88 approaches pinch-off, an increasing fraction of the modulated electrons in the channel are not at their saturated velocity, while at low values OfVg8 there is significant modulation of the charge in the AlGaAs supply layer. This led Foisy et al [22] to introduce the concept of a 'modulation efficiency', r|, defined as the ratio of the modulated velocity-saturated charge in the channel to the total modulated charge. The value of fT is then reduced from vsat / 27iLg to vsat / 27uLg r\. Since the two mechanisms which reduce r\ below 1 have opposite dependences on V88, fT passes through a maximum as a function OfV88. As brought out clearly in [22], the better performance of PHEMTs compared with conventional AlGaAs/GaAs HEMTs is a result of the better confinement in the InGaAs channel separating these two mechanisms which reduce gm , giving a higher and broader maximum in the curve of Sm vs - Vgs- T h i s i s ft% described in [13,15]. D

HEMT FABRICATION

The fabrication of HEMTs is in general very similar to the processes used to fabricate MESFETs, although the thinner layer structure and smaller feature sizes impose some additional requirements. In particular HEMTs are almost universally made with gate lengths of 0.25 urn or shorter. The growth of HEMT structures is most commonly carried out by molecular beam epitaxy (MBE) to achieve the necessary control of layer thicknesses, compositions and doping. However the use

of metallo-organic vapour phase epitaxy (MOVPE) has been reported for both conventional [23] and pseudomorphic devices [24]. The most common fabrication sequence is first to define the area of the device by mesa etching through the active layers to the buffer layer and then to define the source and drain ohmic contacts by standard lithographic and deposition techniques. Typically the spacing between the source and drain contacts is 3 ^m or less. The most common ohmic contact metallisation is evaporated AuGeNi, annealed at around 400 0 C. However many other metallisation schemes have been reported. The next stage is the etching of the gate recess and the deposition of the gate metallisation. These are normally performed in a single lithographic stage. To achieve good high frequency performance very short gate lengths are used, typically from 0.25 ^m down to below 0.1 ^m, requiring the use of electron beam (e-beam) lithography. For good noise performance the parasitic source and gate resistances R8 and Rg must be kept small. To minimise R8 the gate is placed as close to the source ohmic contact as possible (typically - 0.6 jum), since the access resistance between the source contact and the start of the gated channel contributes a significant part OfR8. Typical values OfR8 are 0.5 Q.mm (of source metallization width) or less. To minimise Rg requires the gate metallisation to have a large crosssection, which conflicts with the need for short gate lengths for high fT. To resolve this conflict it is common to use a gate metallisation with a 'mushroom' cross-section, so that a short 'footprint' can still be achieved with a large cross-section of metal. Mushroom gates typically have Rg below ~ 200 Q/mm (of gate width), even for gate lengths of 0.1 |um. To achieve such a mushroom cross-secion requires the use of multi-layer e-beam resists with an appropriate exposure strategy to produce the mushroom profile in the resists [25]. After exposure and development, the e-beam resist layers first serve as a mask for etching to produce the gate recess. Wet chemical etching is frequently used but, for better control, low damage dry etching processes are preferred [26]. In particular, selective dry etching combined with the correct layer design can be used to etch the gate recess through the GaAs cap layer and stop at the AlGaAs layer below. Deposition of the gate metal then follows, the most common metallisation being TiPtAu (although again there are many other metallisation schemes in common use). Finally the device is completed with appropriate metal overlays and passivation layers. E

DEVICE RESULTS

El

Low-noise Microwave and Millimetre-wave Devices

Although the HEMT was originally conceived as a high speed device for digital applications, its potential for low-noise microwave use was recognised early in its development [H]. Its performance for this application made rapid progress throughout the 1980s, resulting in the HEMT being the preferred device for very low-noise applications (such as DBS television receivers) below 20 GHz, while extending the range of operation of transistor-based circuits up to 100 GHz and beyond. The requirements for a high performance microwave device are: (i) a high value of the DC transconductance, gm (ii) a short gate length, L8, to give a low gate to source capacitance, C&, and hence, with the high gjn, a high current-gain cut-off frequency, fT

(iii)

(iv)

low values of other parasitic elements, especially the output conductance and the gate to drain capacitance, Cgd, to give a high value of the unity power gain frequency, fmax, and hence a high gain low values of the parasitic source and drain resistance, since these contribute to the noise figure.

The earliest HEMTs had ^ values of about 200 mS/mm: material and process improvements raised this value to 570 mS/mm for 0.33 ^m gate length devices by 1984 [28]. The improved carrier confinement of PHEMTs gave even higher values, early 1 |im gate length devices having gn values of 270 mS/mm [29]. A value of 1380 mS/mm was achieved as early as 1987 [30], with the highest value to date for a PHEMT being 1510 mS/mm [31]. Values of fT for PHEMTs have progressed from 24.5 GHz for early 1 ^m gate length devices [32] to 120 GHz for 0.2 \im gate length devices [33], 152 GHz for 0.15 |nm gate length devices [34] to 220 GHz for 0.1 |um gate length devices [30]. These results are all roughly consistent with a 1988 study by Hikosaka et al giving the fT * Lg product for PHEMTs to be approximately 21 GHz. jim [35]. This consistency between different establishments is a consequence of fT being predominately determined only by the transport properties of the material system and the gate length. Gain, noise figure and ^ 3x a r e m o r e variable between different organisations since these quantities depend critically on the ability of the process technology to minimise parasitic elements. Nevertheless by about 1990 noise figures below 0.7 dB and associated gains above 12 dB had become benchmark figures for 0.25 jum gate length devices for volume applications such as DBS receivers. Shorter gate lengths have been used to achieve low-noise performance at higher frequencies. PHEMTs with 28% In mole fraction in the pseudomorphic channel and with 0.15 |um gates have demonstrated noise figures of 1.5 dB and associated gains of 6.1 dB at 60 GHz [36]. Devices with 0.10 |nm gates have achieved 2.1 dB noise figure and 6.3 dB associated gain at 94 GHz, with an extrapolated ^113x of 290 GHz [37]. A compilation of published results up to 1995 for noise performance as a function of frequency is given in figure 3 of [14]. The trade-offs between optimising a device for a high value of fT or for a high value of fmax have been discussed by Lester et al [38], who achieved an extrapolated value of fmax of 350 GHz by introducing a double gate recess (involving an additional lithography step) to reduce Cgd. In the early stages of HEMT development, an impressive demonstration of their noise performance was provided by their use in low-noise receivers for signals from the Voyager spaceprobe (launched several years before the invention of HEMTs!) during its fly-by of Neptune in 1988, at a distance of more than 4 billion km from Earth [39]. E2

Power Devices

Although originally developed as a device for low-noise operation at microwavefrequencies,the HEMT has more recently become an important microwave and millimetre-wave power device. The first measurements of the power performance of HEMTs were reported by Smith et al [40] in 1985. They showed that even a conventional AlGaAs/GaAs HEMT designed for small signal operation could achieve a power density of 0.34 W per mm of gate width and a power-added efficiency (PAE) of 37% at 15GHz. Subsequently there have been extensive reports of devices specifically optimised for their power performance. An essential requirement for a power device

is to increase the current per mm of gate width. The PHEMT, with its higher carrier density in the channel, has been widely adopted in place of the AlGaAs/GaAs HEMT. In addition, more complex doping strategies have been investigated. Lester et al [41] compared the power density (in watts per mm of gate width) and power-added efficiency at 35 GHz of 0.25 \xm gate length devices fabricated on the basic low-noise structure, on a double heterojunction structure with delta-doped AlGaAs both above and below the channel, and on a structure where the channel itself was doped, and with a structure combining both of these features. They achieved the highest power density (0.97 W/mm) with the double doped heterojunction structure, a layer design which has been adopted by several organisations [42-46]. The use of delta-doping in the AlGaAs supply layer has the double advantage both of increasing the efficiency of charge transfer into the channel and of placing the gate metal on undoped AlGaAs, thereby increasing the gate to channel breakdown voltage. In some power HEMT processes, a double gate recess [38] is used to increase the gate to drain breakdown voltage [42,45,47]. Other organisations [48-51] have used a different layer structure in which the channel is doped and the AlGaAs left undoped. Such a device is more properly called a heterojunction FET (HFET) rather than a HEMT, since the principle of modulation-doping is not employed. Again the advantages are high channel currents and increased gate breakdown voltage due to the Schottky gate being on undoped material. Power HEMTs and HFETs are attracting considerable attention for two reasons. Firstly, short gate length power HEMTs are unchallenged as power transistors for mm-wave applications. Results include: 0.2 ^m gate length devices giving 500 mW output power with 35% PAE at 32 GHz and almost the same gain with 30% PAE at 44 GHz [47], 184 mW with 25% PAE at 55 GHz [44] and 63 mW with 13.2% PAE at 94 GHz [52]. Compilations of results for output power and PAE as functions of frequency are given in Figures 8 and 9 of [14]. Secondly, HEMTs are now finding widespread application at lower frequencies where power MESFETs or Si bipolar transistors have until now been the dominant power devices. The advantage of HEMTs stems from their higher PAE, which results from their very high gain and low knee voltage (the drain voltage at which the drain current saturates). Consequently HEMTs are becoming widely adopted when efficient conversion of DC to microwave power is required. This is particularly the case for battery-powered applications such as mobile communications [53]. F

HEMTS IN INTEGRATED CIRCUITS

Much of the original motivation for the development of HEMTs was aimed at very high speed digital circuits. While many impressive results were achieved, reviewed in [6,8,9], the difficulties in fabricating large integrated circuits with the necessary control (particularly of threshold voltage) have led to interest in HEMTs for high speed digital circuits declining considerably in recent years. In contrast to digital circuits, interest in the use of HEMTs in monolithic microwave and millimetre-wave integrated circuits for analogue applications continues to expand rapidly. A vast number of HEMT-based circuits have now been reported, covering a range of functions including low-noise amplifiers, broad-band amplifiers, power amplifiers, up-converters, down-converters, oscillators and multi-function circuits. References to many advances in the design and fabrication of such circuits may be found in the Annual IEEE MTT-S Int. Microwave Symp. Digest and the Annual Technical Digest of the IEEE GaAs Integrated Circuits Symp. Many such HEMT-based monolithic microwave and millimetre-wave integrated circuits are now available commercially.

REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] [II] [12] [13] [14] [15] [16] [ 17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31]

R. Dingle, H.L.Stormer, A.C. Gossard W. Wiegmann [ Appl. Phys. Lett. (USA) vol.33 (1978) p.665-7 ] H. Stormer, R. Dingle, A.C. Gossard, W. Wiegmann M.D. Sturge [ Solid-State Commun. (USA) vol.29 (1979) p.705-9] T. Mimura, S. Hiyamizu, T. Fujii K. Nanbu [ Jpn. J. Appl. Phys. (Japan) vol. 19(1980) p.L225-7 ] D. Delagebeaudeuf, P. Delescluse, P. Etienne, M. Laviron, J. Chaplart N.T. Linh [ Electron. Lett. (UK) vol.16 (1980) p.667-8] H. Morkoc [ IEEE Electron Device Lett. (USA) vol.EDL-2 (1981) p.260-2 ] H. Morkoc H. UnIu [ Semicond. Semimet. Ed.R. Dingle vol.24 (Academic Press Inc, San Diego, 1987) p. 135-201] N.T. Linh [ Semicond. Semimet. Ed. R. Dingle vol.24 (Academic Press Inc, San Diego, 1987) p.203-47 ] M. Abe, T. Mimura, K. Nishiuchi, A. Shibatomi, M. Kobayashi T. Misugi [ Semicond. Semimet. Ed. R. Dingle vol.24 (Academic Press Inc, San Diego, 1987) p.249-78 ] T. Mimura [ Semicond. Semimet. Ed.T. Ikoma vol.30 (Academic Press Inc, San Diego, 1990) p. 157-93] SJ. Pearton NJ. Shah [in High Speed Semiconductor Devices Ed. S.M. Sze (J. Wiley sons Inc, New York, 1990) p.283-334 ] H. Morkoc, H. UnIu G. Ji [ in Principles and Technology of MODFETs vols. 1&2 (J. Wiley Sons Inc, Chichester, 1991)] P.C. Chao, A.Swanson, A. Brown, U. Mishra, F. AU, C. Yuen [ in HEMTs andHBTs Eds. F. AIi A. Gupta (Artech House Inc, Boston 1991) p.77-190 ] L.D. Nguyen, L.E. Larson U.K. Mishra [ Proc. IEEE (USA) vol.80 (1992) p.494-518 ] S. Takamiya, N. Yoshida, N. Hayafuji, T. Sonoda S. Mitsui [ Solid-State Electron. (UK) vol.38 (1995)p.l581-88] PJ. Tasker [ in Physics and Technology of Heterojunction Devices Eds. D.V. Morgan RH. Williams (Peter Peregrinus Ltd, London, 1991) p. 146-200 ] J.W. Matthews A.E. Blakeslee [ J. Cryst. Growth (Netherlands) vol.27 (1974) p. 118-25 ] JJ. Rosenberg, M. Benlamri, P.D. Kirchner, J.M. Woodall G.D. Pettit [ IEEE Electron Device Lett. (USA) vol.EDL-6 (1985) p.491-3 ] H. Suehiro, T Miyata, S.Kuroda, N. Hara M. Takikawa [ IEEE Trans. Electron Devices (USA) vol.41 (1994) p. 1742-6] L.D. Nguyen, A.S. Brown, M.A. Thompson L.M. Jelloian [ IEEE Trans. Electron Devices (USA) vol.39 (1992) p.2007-14] D. Delagebeaudeuf N.T. Linh [ IEEE Trans. Electron Devices (USA) vol.ED-28, (1981) p.790 ] RF. Pierret M.S. Lundstrom [ IEEE Trans. Electron Devices (USA) vol.ED-31 (1984) p.383-5 ] M.C. Foisy, PJ. Tasker, B. Hughes L.F. Eastman [ IEEE Trans. Electron Devices (USA) vol.ED35(1988)p.871-8] H. Takakuwa, Y. Kato, S. Watanabe Y. Mori [ Electron. Lett. (UK) vol.21 (1985) p. 125-6 ] A.G. Thompson, B-Y Mao G.Y. Lee [Appl. Phys. Lett. (USA) vol.55 (1989) p.2208-10 ] P.C. Chao, P.M. Smith, S.C. Palmateer J.C.M. Hwang [ IEEE Trans. Electron Devices (USA) vol.ED-32 (1985) p. 1042-6 ] F. Ray et al [ IEEE Trans. Electron Devices (USA) vol.ED-39 (1992) p.2701-6 ] M. Laviron, D. Delagebeaudeuf, P.Delescluse, J. Chaplart N.T. Linh [Electron. Lett.fUK) vol.7 (1981)p.536-7] L.H. Camnitz, PJ. Tasker, H. Lee, D. van der Merwe L.F. Eastman [ IEDM Tech. Dig. (1984) p.360-3 ] A.A. Ketterson et al [ IEEE Trans. Electron Devices (USA) vol.ED-33 (1986) p.564-71 ] P.C. Chao et al [ Proc. IEDM. (1987) p.410-3 ] F. Diette, D. Langrez, JL. Codron, E. Delos,D. Theron G. Salmer [Electron. Lett. (UK) vol.32

[32] [33] [34] [35] [36] [37] [38] [39] [40] [41] [42] [43] [44] [45] [46] [47] [48] [49] [50] [51] [52] [53]

(1996), p.849-50 ] A.A. Ketterson et al [ IEEE Trans. Electron Devices (USA) vol.ED-33 (1986) p.564-71 ] L.D. Nguyen, D.C. Radulescu, PJ. Tasker, WJ. Schaff LF. Eastman [ IEEE Electron Device Lett. (USA) vol.9 (1988) p.374-6 ] L.D. Nguyen, PJ. Tasker, D.C. Radulescu L.F. Eastman [ IEEE Trans. Electron Devices (USA) vol.ED-36(1989)p.2243-8] S. Hikosaka, S. Sada, N. Harada S. Kuroda [IEEEElectron Device (USA) Lett, vol.9 (1988) p.2413] K.L. Tan et al [ IEEE Electron Device Lett. (USA) vol.EDL-12 (1991) p.23-6 ] K.L. Tan et al [ IEEE Electron Device Lett. (USA) vol.EDL-11 (1990) p.585-7 ] L.F. Lester et al [ Proc. IEDM. (1988) p. 172-5 ] K.H.G. Duh, B.C. Adams [ Microwave J. (USA) (1990) p. 122-6 ] P.M. Smith, U.K. Mishra, P.C. Chao, S.C. Palmateer J.C.M. Hwang [ IEEE Electron Device Lett. (USA) vol.EDL-6 (1985) p.86-7 ] L.F. Lester et al [ IEEE Trans. Electron Devices (USA) vol.ED-36 (1989) p.2616-7 ] M-Y. Kao, S-T. Fu, P.M. Smith, P.C. Chao, KJ. Nordheden S. Wang [ Proc. IEDM (1992) p.31921] S.T. Fu et al [ Proc. IEEEMTT-SDigest (USA) (1993) p. 1469-72 ] K.L. Tan et al [IEEEElectron Device Lett. (USA) vol.EDL-12 (1991) p.213-4 ] J.C. Huang et al [ IEEEMTT-S Digest (USA) (1991) p.713-6 ] CS. Wu, F. Ren, SJ. Pearton, M. Hu, CK. Pao RF. Wang [ IEEE Trans. Electron Devices (USA) vol.ED-42(1995)p.l419-23] J.C. Huang et al [IEEEElectron Device Lett. (USA) vol. EDL-14, (1993) pp. 456-8 ] B. Kim, H.Q. Tserng, J.W. Lee [IEEEElectron Device Lett. (USA) vol. EDL-8 (1987) pp. 638-9 ] H.Q. Tserng, P. Saunier, Y-C Kao [ Electron. Lett. (USA) vol. 29, (1993) p.304-6 ] T. Sonoda et al [ Electron. Lett. (USA) vol. 27 (1991) p. 1303-5 ] J. Udomoto et al [ IEEEMTT-SDigest (USA) (1995) p.339-42 ] K.L. Tan et al [IEEEElectron Device Lett. (USA) vol.EDL-12 (1991) p.213-4 ] Y-L. Lai et al [IEEEElectron Device Lett. (USA) vol.EDL-17 (1996) p.229-31 ]

20.4 GaAs in bipolar integrated circuits J.A. Turner December 1995

A

INTRODUCTION

Over the past few years the GaAs based heterojunction bipolar transistor (HBT) has emerged as a very versatile component in a number of circuit applications [1-3]. Its material structure gives it many beneficial properties over both silicon bipolar transistors and the current families of MESFETs and heterostructure FETs. In comparison to silicon bipolar devices the HBT has a higher cut-off frequency, ft, reduced base resistance, lower base emitter capacitance and higher substrate resistivity. In comparison to the GaAs MESFET, the HBT shows higher transconductance, higher current and power density, better threshold voltage matching, and lower 1/f noise. Considerable effort is now being made to demonstrate a cost effective manufacturing process; the material structure is quite complex and requires a high degree of control of the epitaxial growth process. The process technology is undergoing 'streamlining' to reduce the number of masking operations. Early technologies used a non self-aligned fabrication approach in order to access the various contact areas of the devices. Generally, this has now been replaced by a self aligned process in order to place the base contact as close to the emitter as possible and thereby reduce the base resistance of the structure. The lithographic techniques used to produce the device, even those operating at high microwave frequencies (10 GHz), do not need to be sophisticated. Emitter metallisation widths of 2 microns are commonplace. The high frequency performance is determined mainly by the sub-micron base width which is controlled by epitaxial, not lithographic processes. Many laboratories across the world have published exciting results with the device being used in both analogue and digital applications. The design techniques used in these circuits closely follow that for the silicon bipolar transistor, for which highly sophisticated CAD packages are available. High volume applications of the device are emerging. For instance, they are now being used as transmitter components in hand held cellular telephones. B

PERFORMANCE

Bl

Microwave Power Gain

One figure of merit for the HBT is the maximum frequency of oscillation, fmax,

max

" ( s*R b cJ

0)

where R1, is the base resistance and C1x. is the base-collector capacitance. ft is the cut-off frequency given by DnAV where Dn is the diffusion coefficient of electrons in the p-type base and W is the

base width. To optimise fmaK requires the reduction OfR1, and Cbc which requires using relatively large values of collector-base depletion width to reduce C1x.. This will increase the overall emitter to collector transit time and reduce ft from its optimal value. Best values of fmax reported for GaAs/GaAlAs devices are 218 GHz with a corresponding ft of 160 GHz [I]. B2

Microwave Output Power

One of the key applications of the HBT is for high transmitted power. GaAs HBTs have the same attributes as the silicon device (high current density and high voltage capability) but with an improved frequency response. CW power densities of over 6W per millimetre of emitter length have been reported at 9 GHz [4] with power added efficiencies of 62 %. Millimetre wave operation has also been demonstrated at 59 GHz [5] where 1.7W per millimetre of emitter length has been reported. 8.5dB of gain at 5OW (2.5W/mm) has also been achieved at 35 GHz [6]. The pulsed performance of GaAs based HBTs is also very impressive. The highest power density of 18.7 W/mm has been achieved at 10 GHz from a device giving 5 dB of gain with an output powerof560mW[7]. B3

Low Frequency Noise Performance

Current GaAs HBTs do not compare favourably with GaAs MESFETs and HEMTs for microwave noise performance. Johnson noise in the base resistance and shot noise are the major sources. However 1/f noise in HBTs has been shown to be considerably reduced over GaAs MESFETs due mainly to a reduction in surface effects [8,9]. This has led to their usage in oscillator applications where low close to carrier noise is important. Phase noise values some 10 dB lower than for GaAs MESFETs are obtained [3]. B4

Digital Switching Speed

HBTs can switch at very high speeds due to their high fj and transconductance though in common with silicon bipolar circuits their power dissipation limits the level of integration that can be achieved. For this reason, the highest speed circuitry (operating at around fmax/2) will be of relatively low complexity. C

GaAs HBTs IN CIRCUIT APPLICATIONS

Cl

Analogue Circuits

The fundamental device performance of the GaAs HBT has been used to advantage in a number of analogue MMIC applications. These include: (a) (b) (c) (d) (e)

Wide band amplifiers for communications and instrumentation [8,10] Low noise amplifiers [11] - 2.OdB noise figure, 2.1 mW power dissipation Voltage comparators [12] Highly linear amplifiers [3,13,14] Logarithmic amplifiers [15]

(f) (g) (h) (i)

Mixers [1] Oscillators [5,16-19] - up to 62 GHz Power amplifiers [4,20-24] 12 W, 12 GHz5 PAE 48% Mixed signal circuits [25] - InP based devices

C2

Digital Applications

Over the past few years there has been a dramatic increase in the level of integration of HBT digital circuits. High speed operation records are being broken by relatively simple divider circuits, the latest being 50 GHz [26]. However, advanced work particularly at Texas and IBM/Rockwell in the US, has been aimed at MSI and LSI levels of integration. Texas Instruments [27] have produced a 32 bit microprocessor containing 12,900 equivalent NOR gates implemented in I2L logic with a clock frequency of 100 MHz. IBM/Rockwell [28] have used an ECL approach to produce a gate array with 1200 equivalent NOR gates. Flip flops in this circuit toggle at 7.7 GHz. C3

Multifunction Circuits

The microwave and digital capability of the HBT has been put to use in a number of mixed function circuits. Single chip A/D converters have been fabricated with sample rates around 4 Gbit/sec. With the availability of dividers operating above 10 GHz and DC - 10 GHz wide band amplifiers, this frequency will soon be extended. The integration of lasers with HBT driver circuitry has also been reported [29]. C4

Reliability

The reliability of HBTs has shown a marked improvement during the last two years and they have been shown by some companies and some process technologies to have lifetimes comparable to the GaAs MESFET. Temperature stress testing of GaAs/GaAlAs HBTs at three different temperatures has given a median time to failure (MTTF) at 1250 C of greater than 108 hours and MMICs utilising these devices have operated successfully at temperatures of greater than 300 0 C [30]. GalnAs/AlInAs/InP devices have been reported [30] to have MTTFs of 1.27 x io 7 hours at 125 0 C. In these devices the failure site is the base collector junction. GaInAs/AlInAs/InP HBTs do not show hydrogen related degradation effects that have been reported for MESFETs and GaAs/GaAlAs based HBTs [31,32]. C4

Radiation Hardness

The search for increased radiation hardness in semiconductor devices continues, and already GaAs HBTs have been subjected to harsh space-like environments. Initial tests on a limited number of GaAs/GaAlAs devices [33] have shown that their radiation hardness is considerably better than silicon transistors and compares favourably with GaAs MESFETs. D

FUTURE HBT IMPROVEMENTS

New material structures (other than GaAs / AlGaAs) have recently improved the performance of the HBT. Although in the late 1980s good results were also achieved for InP/InGaAs HBTs, [20-22, 27-29, 34-36] work has now concentrated on GaInP / GaAs and GaAlAs / GaAs devices

on GaAs substrates [37] and Al InAs/GalnAs on InP substrates [25]. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] [II] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33] [34] [35] [36]

M.E. Kim et al [ IEEE Trans. MTT (USA) vol.37 no. 9 (1989) p.1286 ] P.M. Asbeck et al [ IEEE GaAs IC Symp. (USA) (1991) p.7-10 ] B. Bayraktaroglu et al [ Proc. IEEE (USA) vol.81 no. 12 (1993) ] B. Bayraktaroglu et al [ IEEEMTT-S Int. Microwave Symp. Digest (USA) (1989) p. 1057 ] J.A. Higgins [ Tech. Digest GaAs IC Symp. (USA) (1988) p.23 ] NL. Wang, WJ. Ho, J. Higgins [ IEEE Microwave and Guided Lett. (USA) vol.2 (1992) p.397 ] M. Adlerstein et al [ Electronic Lett. (UK) vol.27 (1991) p.148-149 ] M.E. Kim et al [ Tech. Digest GaAs IC Symp. (UK) (1988) p.117 ] W. Fonger [ Transistors 1 (USA) Princeton, NJ, RCA Labs. (1956) p.239 ] K.W. Kobayashi et al [ Tech. Digest GaAs IC Symp. (USA) (1989) p.87 ] K. Kobayashi et al [ IEEE Microwave and Millimetre Wave Monolithic Circuit Symp. (USA) (1995) p.73] P.M. Asbeck et al [ IEEE Trans. Electron Dev. (USA) vol.36 no. 10 (1989) p.2032 ] B.L. Nelson et al [ Tech. Digest GaAs IC Symp (USA). (1989) p.79 ] P. Walters et al [ GaAs IC Symp. Digest (USA) (1995) p.67-70 ] G.M. Gorman et al [ IEEEMTT-SInt. Microwave Symp. Digest (USA) (1989) p.537 ] K. W. Lee et al [ Tech. Digest GaAs IC Symp. (USA) (1987) p.227 ] M. Madihianetal [Tech. Digest GaAs IC Symp. (USA) (1988) p. 113] K. Matsvnagaeta\[ IEEE Microwave and Guided Wave Lett. (USA) vol.5 no. 11 (1995)p.4O2] H. Blanck et al [ IEEE Microwave and Millimetre Wave Monolithic Circuit Symp. Dig. (USA) (1994)] H. Wang et al [ IEEE Microwave and Guided Wave Lett. (USA) vol.5 no. 11 (1995) p.388-90 ] K. Sakumo et al [IEEE Microwave andMillimetre Wave Monolithic Circuit Symp. Digest (USA) (1994)p.63-6] A. Khatibzadeh et al [ IEEE Microwave and Millimetre Wave Monolithic Circuit Symp. Digest (USA) (1994) p. 117-20] F. AIi et al [ IEEE Microwave and Millimetre Wave Monolithic Circuit Symp. Digest (USA) (1994) p. 113-16] JJ. Komiak et al [ IEEE Microwave andMillimetre Wave Monolithic Circuit Symp. Digest (USA) (1995) p. 17-20] W. Stanchina et al [ GaAs IC Symp. Digest (USA) (1995) p.31-4 ] Y. Yamanchi et al [ Tech. Digest GaAs IC Symp. (USA) (1989) p. 121 ] DA. Whitmire et al [ Tech. DigestISSCC (USA) (1988) p.34 ] J. George et al [ Tech. Digest ISSCC (USA) (1988) p.30 ] J. Shibataetal[vl/>/>/. Phys. Lett. (USA)vo\A5 (1984) p. 191] M. Hafizi et al [ IEEE Microwave andMillimetre Wave Monolithic Circuit Symp. Digest (USA) (1995)p.lll-14] F. Yamada et al [ GaAs Reliability Workshop Proc (USA). (1995) p.2-10 ] WJ. Roesch [ GaAs Reliability Workshop Digest (USA) (1994) ] Y.D. Song et al [ Tech. Digest GaAs IC Symp. (USA) (1989) p. 115 ] R.Nottenbvirg[Tech. Digest GaAs IC Symp. (USA) (1989) p. 135] Y.K. Chen et al [ IEDMTech. Digest (USA) (1988) p.876 ] U.K. Mishra et al [ IEDMTech. Digest (USA) (1988) p.873 ]

20.5 GaAsMMWICs MJ. Howes August 1995

A

INTRODUCTION

GaAs has proved to be an excellent material for semiconductor devices operating in the microwave frequency band and beyond. It has essentially replaced other technologies and materials as a base for low power devices operating in the 2 - 4 0 GHz frequency band and GaAs devices are also mounting a significant challenge to vacuum tube devices at higher power levels. High performance and reliable two, three and four terminal discrete devices have found applications in virtually all aspects of microwave engineering. Moreover, monolithic GaAs integrated circuit technology is now in a sufficiently mature state that analogue and digital circuits are commercially available. Although the development of mature devices based on GaAs has been relatively slow as compared with those based on silicon due to the more difficult technological aspects, they have emerged much faster than device modelling and circuit design techniques and the exploitation of the maximum performance of such devices is still a subject for intense research and development. B

ANALOGUE ICs

There is no doubt that monolithic mm-wave integrated circuit (MMWIC) technology is the key to producing the high frequency, miniature low-cost electronic circuits necessary for future systems [I]. Recent developments have indicated that GaAs based MMWIC technology will form the basis of many applications up to 100 GHz and beyond. These new developments have largely centred around the application of nanoelectronic technology to planar devices such as the GaAs/AlGaAs high electron mobility transistor (HEMT) and the InGaAs/AlGaAs pseudomorphic HEMT. The result has been the demonstration of transistors with gate lengths of 50 nm and fT of >300 GHz. Consequently, the building blocks of the analogue aspects of electronic systems (amplifiers, mixers, oscillators) can be designed with sufficient performance such that mm-wave applications such as radar, phased arrays, radiometers and communication systems become a commercial reality. Bl

Substrates

For MMW7ICs the substrate is normally semi-insulating GaAs and the resulting circuit is ideally planar in form. Consequently the electrical modelling of the substrate is as critical as the electrical modelling of the circuit components themselves. The microwave permittivity of semi-insulating GaAs is a fundamental parameter in the design of GaAs active and passive circuit components and may be written 6 =

r

^ "K

(D

The imaginary part of the permittivity is a measure of the dielectric loss of the material and a figure of merit of the material, termed the loss tangent, can be defined which is given by

tan8 = -^-

(2)

The parameters er' and er" are both Sanctions of frequency and hence for accurate MMWIC design it is essential that the complex dielectric constant of GaAs is known over the frequency range of interest. Bl. 1

Frequency and temperature dependence of permittivity

Theoretically, for a low loss material such as semi-insulating GaAs, the permittivity should not vary significantly from DC up to millimetre wave frequencies. This value is referred to as the low frequency or static dielectric constant and has been measured over a wide range of frequencies. GaAs is a polar material and it therefore displays a characteristic resonance in the far infrared region of the electromagnetic spectrum (at a wavelength of about 35 \xm in GaAs). At frequencies above this resonance the permittivity drops to a somewhat lower value referred to as the infrared or optical permittivity. Room temperature measurements of the real relative permittivity and loss tangent of semiinsulating GaAs as reported in the literature are given in TABLE 1. The large range of measured values of er' is quite perplexing and cannot be totally explained by measurement error. It is quite clear that an investigation into the value of the complex permittivity of GaAs and its dependence on frequency, crystal imperfections, defect densities and fabrication processing is long overdue. The temperature dependence of the permittivity shows that er increases almost linearly with temperature and is of the form e/T) = €^0) [1 + aT]

(3)

The values of er'(0), a and room temperature values of e r ', er'(295), are shown in TABLE 2 for various frequencies. Although there is some spread in the results, a appears to decrease with frequency. It is no surprise that measurements of the loss tangent of GaAs indicate that €r" increases with frequency (at a rate of 2 x W6ZGHz) and with temperature. B2

MMWIC circuits

The availability of energy sources is the key to the success of a circuit technology [I]. If high performance oscillators can be realised then other subsystems such as amplifiers and mixers will follow. Due to its very low phase noise and high DC to RF conversion efficiency, the bipolar junction transistor is normally preferred up to frequencies of around 10 GHz. The phase noise

TABLE 1. Summary of measured values of er' and tanS of GaAs at RT. Frequency (GHz)

er'

tanS

Date

2-10

11

0.001

1966

9

12.35

0.00056

1967

9-11

13.3

0.0016

1968

33

13.25

2-36

12.95

4-18

12.9

8-71

12.7

71

12.95

1967

71

13.18

1968

71

13.17

1968

60

1968 0.0002-0.0006

1977 1980

13.3

1967

0.0005

1973

0.0009-0.0014

1983

100-400

12.87-12.92

IR

13.13

1964

IR

13.1

1969

TABLE 2. Published results on the temperature dependence of er' of GaAs. Frequency (GHz)

Temperature range (K)

er'(0)

104Ct(K-1)

er'(295)

4.5

300-400

12.36

12.94

1.60

71

100-300

12.73

13.18

1.20

71

100-600

12.79

13.17

1.00

Far infrared

8 and 300

12.80

12.80

0.56

performance associated with MESFETs is very poor, but oscillators based on this device and its variants such as the HEMT can deliver useful amounts of power well into the mm-wave range of the frequency spectrum. High Q dielectric resonators are often used to reduce the phase noise of oscillators based on FETs to an acceptable level and varactors based on GaAs technology are used in narrow and broad band voltage controlled oscillators. For multi-octave tuning and when slew rate and power consumption is of minor concern, yttrium iron garnet (YIG) tuning elements can be used. Transistor oscillators are normally based on a synthesised two terminal negative conductance. The synthesis is achieved by using a positive feedback amplifier topology. The transistor may be in the common gate, common source or reverse channel common drain configuration. Common gate is preferred for realising low power broadband voltage controlled oscillators since it is

relatively easy to obtain broadband small signal negative conductance in this way. For higher power, fixed frequency or narrow band tunable oscillators the common source configuration is to be preferred with the possibility of using a dielectric resonator to improve the phase noise performance. The basic circuit configurations for the synthesis of negative conductances are shown in FIGURE 1.

Common Source

Common Gate

Common Drain

Reverse Channel CD

FIGURE 1. The basic circuit configurations for the synthesis of negative conductances.

Dielectric resonators can be used in a variety of ways to stabilise and/or improve the noise performance of FET oscillators. i.

DSOs - Passive stabilisation of free running oscillators for which the oscillation frequency is sensitive to load variations.

ii.

DROs - The frequency determining circuit is replaced or augmented by the dielectric resonator and forms an integral part of the design. Two of the many topologies possible are shown in FIGURE 2.

Common Source Resonator in Feedback Ckciant

Transmission Resonator to Provide Drain to Gate Feedback

FIGURE 2. Typical DRO-FET oscillator topologies.

The use of dielectric resonators is widespread in microwave and mm-wave oscillators particularly in circuits realised in MIC form. They are also used, off chip, to stabilise MMIC oscillators. At frequencies in excess of 100 GHz DR technology becomes difficult (TABLE 3) and more traditional circuit techniques have to be employed. TABLE 3. Dielectric resonator materials

Frequency range GHz

Manufacturer

Ba2Ti9O20

1 to 100

Bell Labs

(Zr-Sn)TiO4

1 to 100

Thomson; Transtech; Murata

Ba (Zn173 Ta273) O2

4 to 100

Murata

Ba (Mg173 Ta273) O2

4 to 100

Sumimoto

BaO-PbO-Nd2O3-TiO2

<4

Murata; Transtech

As an indication of the state of the art technology, FIGURE 3 shows a recent example of a Dband MMIC oscillator based on a pseudomorphic double heterojunction InAlAsAn0^Ga0 3 As on an InP substrate which delivers -8 dBm at 131 GHz [5]. The oscillator was designed using a methodology based on a synthesised negative conductance realised by using a dual feedback topology: series feedback from source to ground and parallel feedback from drain to gate. At these frequencies and with the primitive active physical devices available the negative conductance is extremely small (a single ended amplifier using such a device is unlikely to have a gain greater than 2 dB) and circuit design becomes absolutely critical. Matching Stubs

2X45jim HEMT

FIGURE 3. D-band (1319 Hz) synthesised negative conductance p-HEMT oscillator using the selective dual feedback technique.

Electronic tuning of high frequency oscillators is best achieved using the variable capacitance properties of Schottky barrier diode structures. Important aspects of design for broadband operation using MIC technology are large capacitance ratios (and hence hyperabrupt diodes), minimal parasitics (and hence chip devices) and tight coupling between active device and tuning diode [6]. If careful attention is paid to all these points octave tuning can be achieved whilst retaining the high slew rate which characterises varactor tuned oscillators. Due to the difficulty

of fabricating high performance varactor diodes in monolithic form an FET is often turned into a planar varactor by suitably configuring its structure. C

CONCLUSION

Monolithic mm-wave integrated circuit (MMIC) products have more than trebled since 1987. Monolithic oscillators [7, 8], amplifiers [10], mixers [9,11] and switches [12] are readily available at lower frequencies and multi-function chips are coming on stream [13]. Operating frequencies now extend beyond 100 GHz and are likely to reach 200 GHz by the end of the decade. GaAs devices dominate the low to medium power applications at frequencies in the band 5 - 4 0 GHz. This band is continuing to expand as MESFETs challenge the BJT at L-band and HEMTs show their full capabilities above Q-band. Hybrid circuits based on GaAs and InP transferred electron devices will continue to have a substantial role to play as the basis of low noise, low power (10 - 1000 mW) oscillators at frequencies up to 110 GHz until the noise problem with MESFETs under large signal conditions is solved. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] [II] [12] [13]

P.H. Ladbrooke [MMICDesign: GaAs FETs andHEMTs (Artech House, 1989) ] F.B. Fank [ Proc. SPIE - Int. Soc. Opt. Eng. (USA) vol.337 (1982) p. 12 ] R.K. Mains [ IEEE Trans. Microw. Theory Tech. (USA) vol. 32 no.2 (1984) p.208 ] D.W. Mooney, FJ. Bayuk [ IEEE Trans. Microw. Theory Tech. (USA) vol.31 no.2 (1983) p. 171] Y. Kwon et al [ IEEE Trans. Microw. Theory Tech. (USA) vol.41 no. 12 (1993) p.2335 ] D. Large [The Microwave Journal (1970) p.49 ] H. Wang et al [ IEEE Trans. Microw. Theory Tech. (USA) vol.43 no.5 (1995) p. 1010 ] S.E. Rosenbaum et al [ IEEE Trans. Microw. Theory Tech. (USA) vol.43 no.4 (1995) p.927 ] T.H. Chen et al [ IEEE Trans. Microw. Theory Tech. (USA) vol.43 no.3 (1995) p.477 ] S. Banba, H. Ogawa [ IEEE Trans. Microw. Theory Tech. (USA) vol.43 no.3, (1995) p.485 ] A. Paolella et al [ IEEE Trans. Microw. Theory Tech. (USA) vol.43 no.3 (1995) p.518 ] J.L. Lee et al [ IEEE Trans. Microw. Theory Tech. (USA) vol.43 no.2 (1995) p.250 ] CS. Wu et al [ IEEE Trans. Microw. Theory Tech. (USA) vol.43 no.2 (1995) p.257 ]

20.6 The transferred electron effect CG. Diskus, K. Luebke, A.L. Springer and H.W. Thim April 1996

A

INTRODUCTION

In the early 1960s Ridley and Watkins [1] and Hilsum [2] showed theoretically that semiconductors with two ohmic contacts on opposite surfaces may exhibit a negative differential resistance if the charge carriers occupy high-mobility (low effective mass) states at low electric fields and are transferred into low-mobility (large effect mass) states above some critical field. Such a device can thus be used for amplifying or generating microwave signals. Particularly well suited are certain III-V compound materials, whose conduction band consists of a high mobility central valley and several energetically higher-lying satellite valleys. This so called 'transferred electron effect1 was observed in 1963 by Gunn [3] in gallium arsenide and indium phosphide and is now frequently associated with his name. Since their discovery transferred electron devices (TEDs) or Gunn diodes have been developed for frequencies up to 170 GHz. However field effect transistors and, more recently, heterobipolar transistors are becoming strong competitors below 80 GHz not only because of their good performance as far as output power, efficiency and noise are concerned. The fact that they are three terminal planar devices and therefore MMIC compatible makes them very attractive to the circuit designer especially for constructing stable amplifiers. However, two terminal devices such as TEDs are preferred for building millimetre-wave oscillators if high spectral purity and power levels of the order of hundreds of milliwatts are needed. Monolithic integration remains difficult unless properly designed planar structures are used. This Datareview discusses the relevant results obtained at millimetre-wave frequencies with both sandwich and planar TEDs including those having modified cathode and anode contacts. B

DEVICE STRUCTURES

Bl

N+-N-N+ Structures

The most widely used transferred electron device structure is the N+-N-N+ sandwich structure. These conventional TEDs are usually operated at the transit-time frequency which is proportional to the velocity (~107 cm/s) of the cathode injected accumulation layer divided by the thickness of the active N-layer. Unfortunately, a lossy region (the so-called dead zone) exists near the cathode where cold electrons are injected from the N+ contact due to the low electric field. The electrons may travel a distance up to half a micron in order to gain the energy necessary for the transfer into upper valleys. The influence of this lossy region will be strong in devices whose active region is approaching 1 \im, i.e., at frequencies near and above 100 GHz. A way around this problem is to operate TEDs in the second or third harmonic mode or to use an injection limiting cathode contact as will be described in Sections B2 and B3.

Energy relaxation (heating and cooling) and intervalley transfer times responsible for the formation of the dead zone also determine the ultimate upper frequency limit of the transferred electron effect which occurs around 130 GHz in GaAs and perhaps around 250 GHz in InP [4,5]. The best experimental results obtained with N+-N-N+ devices are summarized in TABLE 1. Highlights in this table are 96 mW at 94 GHz (GaAs) in second harmonic operation [6], 60 mW at 151 GHz (InP) and 130 mW at 131 GHz (InP) in fundamental mode operation [7]. Devices with doping gradients seem to perform better than uniformly doped devices due to the electric field enhancement at the cathode contact leading to a reduction of the dead zone. Additional data can be found in references [8-12]. TABLE 1. Output power and efficiency obtained with N+ N N+ TEDs. Material

Doping

Length

Frequency

Output power

Efficiency

Mode

GaAs

1016cm"3

2 ^m

70 GHz

110mW[6]

2.8%

fundamental

GaAs

2 x 1016 cm'3

1 ^m

94 GHz

96 mW [6]

2.7%

2nd harmonic

GaAs

1O16Cm"3

2 ^m

98 GHz

6OmW [6]

3.4%

2nd harmonic

InP

1016 cm' 3 *

1 urn

131 GHz

130 mW [7]

2.5%

fundamental

InP

10 16 cm" 3

1.8 urn

138GHz

65 mW [8]

2.6%

2nd harmonic

InP

1O 16 Cm 3

1 urn

151 GHz

6OmW [7]

2.5%

fundamental

InP

1016 o n 3

1 |nm

279 GHz

0.2 m W [9]

3rd harmonic

* graded doping profile (0.75-2 x 1016 cm"3), increasing from the cathode towards the anode

B2

Sandwich Structures with Hot Electron Injection

In order to fiilly exploit the transferred electron effect at millimetre-wave frequencies, provision has to be taken to avoid the dead zone. The simplest idea is to inject hot electrons. This can be achieved either by replacing the ohmic cathode contact by a reverse-bias current limiting contact where electrons are injected through a high field region [13,14] or by using a forward-biased graded heterojunction injector [15] which injects accumulation layers consisting of hot electrons. If the number of electrons injected from a reverse-biased junction is sufficiently small the formation of accumulation layers may be prevented yielding a stationary field distribution throughout the drift region. This stable field distribution can be uniform or slightly falling depending on the level of electron injection [14,16,17]. Computer simulations [14,17] have revealed that such a device exhibits a frequency-independent negative resistance shunted by its parallelplate capacitance. Thus the operating frequency is determined by the load impedance rather than by the transit-time and, as a consequence, the power-impedance product does not fall off with 1/f2 because longer drift zones can be used at higher frequencies. Another advantage is that the efficiency can approach the maximum theoretical value attainable in a uniform-field mode (32% for GaAs and 40% for InP), but delays due to slow energy relaxation and intervalley scattering will reduce it above 50 GHz, say. Different cathode contacts applied to sandwich structures have indeed led to improved efficiencies but the frequency of operation always corresponds to the transit-time frequency or, in one case

[15], to its second harmonic. The three most successful approaches are the shallow Schottkybarrier [18-20], the two-zone cathode contact [21] and the graded-gap AlGaAs/GaAs launcher [15,22]. The first two types of contacts suffer from problems related to metal-semiconducor interfaces and their associated processing difficulties and are not very reproducible. But efficiencies are very high (24% at X-band, 10% at 55 GHz and 4.6% at 94 GHz). The efficiencies obtained with AlGaAs/GaAs launchers are lower, but reproducibility is very good. The best results are summarized in TABLE 2. TABLE 2. Output power and efficiency obtained with injection controlled TEDs. Material

Doping

Length

Frequency

Output power

Efficiency

Mode

InP

2 x 1015 cm'3

8 ^m

12 GHz

110 mW [21]

24

fundamental

InP

1O16Cm3

1 urn

55 GHz

30OmW [4]

10

fundamental

InP

1O16Cm3

\\un

94 GHz

175 mW [4]

4.6

fundamental

InP

7xl0 1 5 cm" 3

2 \im

90 GHz

126mW[19]

4.7

fundamental

AlGaAs/ GaAs

2 x 1016 cm"3

1 ^m

80 GHz

60 mW [15]

1.5

2nd harmonic

AlGaAs/ GaAs

2 x 1016 cm"3

1 urn

90 GHz

80 mW [15]

2.1

2nd harmonic

AlGaAs/ GaAs

2 x 1016 cm 3

1 jim

94 GHz

58 mW [15]

2.4

2nd harmonic

B3

Planar (MMIC Compatible) Transferred Electron Devices

The sandwich structures described above (Sections Bl and B2) are capable of producing high output power levels due to the large cross-sectional area. However, they are not well suited for monolithic integration because heat extraction through the relatively thick substrate is difficult. Since the use of monolithic integration (MMIC) technology provides significant improvements such as size, weight and cost reduction, high reliability and greater uniformity [23,24], MMICs are being developed all over the world, with the key active element being the planar GaAs field effect transistor (MESFETs and HEMTs). Planar transferred electron devices have been successfully developed using both GaAs [25] and InP [26]. A special feature of the GaAs devices is a MESFET cathode contact with an integrated capacitor providing an RF short between gate and source [16,25], At sufficiently large negative gate voltage the formation of accumulation or dipole layers is suppressed and, instead, a stationary high field region develops in the gate-drain (drift) region. This exhibits negative differential resistance over a wide band of frequencies extending from about half the transit-time frequency up to well above the transit-time frequency. Using this device, which has been called the field effect controlled transferred electron device or FECTED, MMIC-oscillators have been built and operated at 35 GHz (14 mW, 3.8%) [27] and 60 GHz (8 mW, 1.6%) [28]. InP devices have an ion-implanted doping notch next to the ohmic cathode contact which creates a local high field region. However, accumulation layer injection is probably not prevented by the doping notch allowing only transit-time operation. 29.1 mW with 6.7% efficiency at 29.9 GHz

has been obtained with such an oscillator mounted in a waveguide cavity. At 72 GHz only 0.9 mW with much lower efficiency was obtained. A fully monolithic version of this oscillator produced 0.1 mW at 75 GHz3 but the circuit was not optimized [26]. Another promising scheme is a monolithic transferred electron oscillator utilizing a unique flip chip design, whereby a sandwich device sits directly on the heatsink. The device is connected to the RF matching circuit on the back of the semi-insulating GaAs substrate through an integrated via hole. With this approach 7.7 dBm, 11.3 dBm and 13 dBm of output power have been achieved at 96 GHz, 72 GHz and 63 GHz, respectively [29,24]. C

TUNING CAPABILITIES

The frequency of operation of transferred electron effect devices has been found to be sensitive to bias voltage. This can be attributed to both the bias dependent transit-time of travelling accumulation layers and the bias dependent device impedance. The tuning range can be as large as one octave, but the useful tuning range is typically only several hundred megahertz if constant output power is required. For phase locking purposes this is normally sufficient. If larger tuning ranges with good spectral purity are required, both varactor and YIG tuning can be employed. Great frequency variation can be obtained if the appropriate circuit element has a large storage capability for microwave energy. The big advantage of varactor tuning is its fast response (~GHz/|ns), but the frequency-voltage relationship is nonlinear. YIG tuning is linear with current, but the sweep rate is restricted to 1 MHz/^s. With GaAs varactor diodes, tuning ranges of several GHz have been achieved [30-33]. In one case [31] ultra-wideband tuning (69 - 91 GHz) was obtained by tuning the voltage controlled oscillator (VCO) at the fundamental frequency from 23 to 30.33 GHz and using the in-situ generated Gunn-diode 3rd harmonic for output. The achieved output power (-1 dBm to -13 dBm) is sufficient for driving a biased or low barrier height mixer. The phase noise of the 3rd harmonic VCO is extrapolated to be -90 dBc/Hz at a 100 kHz offset frequency. Much higher output power (16.3 dBm ±0.45 dBm) although at much lower frequency (X-band) has been obtained with a hybrid circuit design which integrates a Gunn and a varactor diode with a coplanar waveguide and a slotline to create a planar VCO. The tuning bandwidth was 300 MHz at 10.4 GHz. The achieved spectral purity and stability are comparable to waveguide and dielectric resonator stabilized microstrip oscillators [33]. D

PHASE NOISE

Due to hot electron diffusion, transferred electron devices cannot be used as low noise amplifiers showing noise figures of several dB. However, when operated as oscillators they exhibit very good noise properties and are superior to MESFET and HEMT oscillators, which suffer from large FM noise. A typical noise level produced by a 93 GHz oscillator with an output power of 57 mW is -87 dBc/Hz at 100 kHz off carrier [15]. Of course, the Q-factor of the resonator plays an important role, and 10-20 dB lower noise levels can be achieved if resonators with large RF energy storage such as dielectric resonators are used [34,35]. In any case it is important that the impedance locus of the active device intersects the load impedance line at a large angle [36].

E

CONCLUSION

Transferred electron oscillators are very attractive millimetre-wave sources capable of producing hundreds of milliwatts at efficiencies of several percent with high spectral purity. GaAs devices should soon reach the 100 mW level at 100 GHz and InP devices perhaps 200 mW at 100 GHz and 100 mW at 150 GHz. However, monolithic integration is more difficult in comparison with field effect transistors and hetero-bipolar transistors. A more comprehensive review on TEDs can be found in [37]. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] [II] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33] [34]

B.K. Ridley, T.B. Watkins [ Proc. Phys. Soc. vol.78 (1961) p.293 ] C. Hilsum [ Proc. IRE vol.50 (1962) p. 185 ] JB. Gunn [ Solid State Commun. (USA) vol.1 (1963) p.88 ] PA. Rolland, M.R. Friscourt, C. Dalle, D. Lippens, N. Haese [ Digest 20th EuMC (Microwave Exhibitions and Publishers Ltd.) Budapest (1990) p.34 ] KF. Wu, J. Czekaj,M.P. Shaw[ J. Appl. Phys. (USA) vol.74 (1993)p.315 ] S.J.J. Teng, R.E. Goldwasser [IEEEElectron Device Lett. (USA) vol.10 (1989) p.412 ] H. Eisele, G.I. Haddad [ IEEEMicrow. Guid. Wave Lett. (USA) vol.5 (1995) p.385 ] JD. Crowleyetal [Electron. Lett. (UK) vol.30(1994)p.499] A. Rydberg [ IEEE Electron Device Lett. (USA) vol. 11 (1990) p.439 ] J.M. Szubert, J. Barstow [ Solid-State Electron. (UK) vol.33 (1990) p. 1035 ] R. Kamoua, H. Eisele, G.I. Haddad [ Solid-State Electron. (UK) vol.36 (1993) p. 1547 ] H. Eisele, G.I. Haddad [ Electron. Lett. (UK) vol.30 (1994) p. 1950 ] H. Kroemer [ IEEE Trans. Electron Dev. (USA) vol. 15(1968) p. 819 ] H.W. Thim [ US Patent No. 3 537 021 (1970) ] N.R. Couch et al [ Solid-State Electron. (UK) vol.31 (1988) p.613 ] H.W. Thim [ US Patent No. 3 740 666 (1973) ] G. Rieder,H.W. Thim, R. Kuch,K. Lubke [Arch. Elektron. Ubertrag. Tech. (Germany) vol.37 (1983)p.217] DJ. Colliver, L.D. Irving, JE. Pattison H.D. Rees [ Electron. Lett. (UK). vol. 10 (1974) p.221 ] J.D. Crowley, JJ. Sowers, BA. Janis, F.B. Fank [ Electron. Lett. (UK) vol. 16 (1980) p.705 ] MA. DiForte-Poisson et al [ Proc. 1st Int. Conf. on InP & ReI. Mat. Norman, Oklahoma, USA, 2023 March 1989, p.551-60] K.W. Gray et al [ Electron. Lett. (UK) vol. 11 (1975) p.402 ] S. Neylon et al [ IEEEMMT-SInt. Microw. Symp. Dig. (USA) (1989) p.519 ] B.E. Spielmann [IEEEMMT-SInt. Microw. Symp. Dig. (USA) (1988) p.405 ] M.I. Herman, G.-L. Lan, J.C. Chen, C-K. Pao [ Proc. IEEE (USA) vol.79 (1991) p.342 ] Kuch et al [ Inst. Phys. Conf. Ser. (UK) no.65 (1983) p.439 ] S.C. Binari, R.E. Neidert, K.E. Meissner [IEEEMMT-SInt. Microw. Symp. Dig. (USA) (1988) p.683 ] CG. Diskus et al [ IEEEMicow. Guid. Wave Lett. (USA) vol.3 (1993) p. 180 ] AL. Springer et al [IEEEMicow. Guid. Wave Lett. (USA) vol.5 (1995) p. 114 ] J.C. Chen, C-K. Pao, D.W. Wang [ IEEEMicrow. Millimeter-wave Monolithic Circuits Symp. (1987) p.89] RA. Gough, B.H. Newton [ IEEE Trans. Electron Dev. (USA) vol.20 (1973) p.863 ] L.D. Cohen [IEEEMMT-S Int. Microw. Symp. Dig. (USA) (1991) p.937 ] C Kim, C Dunnrowicz, J. Crawly, B. Fank, L. Wandinger [Microw. J. (USA) vol.32 (1989) p.91 ] JA. Navarro, Yong-Hui Shu, K. Chang [ IEEE MMT-S Int. Microw. Symp. Dig. (USA) (1991) p.1187] H. Barth [ Digest 21th EuMC (Microwave Exhibitions and Publishers Ltd.) Stuttgart (1991) p.364 ]

Next Page

[35] [36] [37]

D. Cros,C. Tronche,P. Guillon,B. 1hsron[IEEEMTT-SInt. Microw. Symp. Dig. (USA) (1991) p.929 ] K. Kurokawa [ Proc. IEEE (USA) vol.61 (1973) p. 1386 ] H. W. Thim [ Handbook on Semiconductors (North Holland, Amsterdam, 1993) vol.4 p.475 ]

20.7 Backgating in GaAs MESFET integrated circuits

Previous Page

L.G. Salmon September 1996

A

INTRODUCTION

Backgating (sometimes called sidegating) is an interaction phenomenon between devices in a GaAs MESFET integrated circuit and is of great concern to integrated circuit designers because it can complicate integrated circuit design and simulation to the point of near impossibility. Backgating is the name given to a reduction in the current output of a MESFET device that occurs when a negative voltage is applied to a neighbouring contact caused by a partial depletion of the MESFET channel by fields induced by the neighbouring contact. The current characteristics of a transistor in an integrated circuit become a function of the voltage of surrounding contacts and unless ameliorated, prediction of gate performance becomes extremely complicated. When backgating characteristics are uniform, the problem is difficult, and when they are nonuniform, the problem can be intractable. Great efforts have been made to understand and minimize backgating, resulting in considerable success. VLSI GaAs integrated circuits with complexity of more than 100 K gates have been reported, and limited numbers have been utilized in military and commercial computer systems [1,2]. Efforts to minimize backgating can be separated into three categories: adapting substrate growth parameters to minimize intrinsic backgating, developing processes to improve device isolation beyond what is possible with intrinsic material, and defining circuit design and layout concepts to minimize the effects of backgating on fabricated integrated circuits. Current theoretical understanding of backgating is first reviewed and then important developments in each of the three categories listed above are summarized. Testing and evaluation of backgating also is reviewed. B

THEORY

FIGURE 1 illustrates the effects of a backgating potential on the I-V curve of an unprotected GaAs MESFET. FIGURE 2 shows the layout of the test structure measured to obtain the data displayed in FIGURE 1. The upper structure is the MESFET device and the bottom metal contact is called the backgating electrode. In an integrated circuit the backgating electrode would be the ohmic contact of an adjacent MESFET device. The I-V curves in FIGURE 1 show that the net current output of the device at a given voltage is reduced as a negative voltage, Vbg, is applied to the backgating electrode. FIGURE 3 shows a profile of the fields in the test structure and illustrates that the backgating electrode is connected to the backside of the gate depletion region. Backgating data from many GaAs integrated circuit fabricators exhibit a threshold phenomenon in the backgating current (the current between the backgating electrode and the source contact

Drain Current, /ds (mA)

Source-drain voltage, Vds (V)

FIGURE 1. Drain current, I68, as a function of drain voltage with the gate shorted to the source. Each curve is the I-V curve at a different backgate bias, Vbg. The data were taken without illumination and a delay of 1 second separated each scan (after [26]). of the MESFET). FIGURE 4 shows such an I-V relationship for the backgating current measured in the structure outlined in FIGURE 2. Two regions are evident in this figure. The first is a linear

Channel implant n+ Implant Gate metal Ohmic metal Metal 1 Via Metal 2

FIGURE 2. Layout schematic of the test structure used to obtain the data. Distances are measured from implant edge to implant edge (after [26]).

region at low backgating bias and the second is a super-linear region at voltages higher than a critical value. The critical voltage is frequently called the backgating threshold voltage, Vtbg. Wafers from many, but not all, GaAs integrated circuit fabricators have a clear backgating threshold similar to that shown in this figure.

Backgate electrode

Source

Gate

Drain

n + GaAs

High-resistivity substrate (p-type)

FIGURE 3. Profile of backgate structure illustrating the fields in the device. Note that the backgate electrode is connected to the depletion region of the MESFET through the substrate (after [26]).

Backgate current /bg (A)

Threshold backgating behaviour and the value of Vtbg were first explained by Lee et al [3] using the trap-filled-limited (TFL) model of Lampert and Mark [4]. The TFL model predicts that backgate conduction is dominated by the filling of deep-level traps in the GaAs substrate. Carriers injected into the substrate from a backgating contact fill ionized, deep-level traps until a critical voltage, Ym. Above Vm all deep-level trap sites are filled, injected carriers are no longer trapped, and the backgating current rises sharply. Below Vm, conduction is dominated by thermionic emission, the I-V curve is principally ohmic in character, and the backgating voltage causes a

Backgate voltage, l/bg (V) FIGURE 4. Backgating current as a function of the backgate voltage. Note the position of Vtbg.

small reduction in I4188. Above Ym, injected carriers fill available deep-level traps, resulting in rapidly increasing conduction based on a power-law dependence. Lee et al associated the Vm of the TFL model of Lampert and Mark with the V ^ shown in FIGURE 4 although measured values of the Vtbg are much lower than the value of the Ym predicted by the TFL model. Not all GaAs MESFET devices show threshold backgating behaviour. FIGURE 5 shows the backgating characteristics of MESFET structures fabricated at three different GaAs integrated circuit fabrication facilities. Devices fabricated at one facility show a sharp threshold in backgating behaviour, devices fabricated at another facility show a much smoother threshold, while devices from a third facility do not show any threshold behaviour.

Backgate bias, Kbg (V)

FIGURE 5. AIj38 as a function of backgate voltage for representative structures fabricated at three different facilities. Note that the backgating data from two facilities show threshold behaviour and backgating data from the other fabricator does not (after [26]).

Because not all GaAs integrated circuits show a clear backgating threshold, and because Ibg is difficult to measure, another quantity, AI^8, is usually used to measure backgating. AI^8 is the percent reduction in the saturation current of the MESFET caused by a backgating voltage. AI^8 is defined as:

Mdss =

xT

dss

b

g/T o ' -Mss

where I ^ is the current through the MESFET with the gate and source shorted to ground (for a depletion-mode device), I4880 is I488 when the backgate contact is grounded and I^ is I^ when the backgate contact is biased. FIGURE 6 shows the variation in AI488 with backgating voltage and shows the same characteristics observed in FIGURE 4: a linear region at low backgating bias and a super-linear region at more negative bias. The critical voltage in this curve is also Vtbg. In the past few years backgating phenomena have been observed that cannot be explained by the TFL model. Liu et al [6] observed that backgating current was increased by the presence of a Schottky contact between the backgating electrode and the MESFET channel. They proposed that the explanation for this effect is hole injection at the Schottky contact as proposed by Goto et al [7,8]. This result was further supported by the modelling efforts reported of Chang and Lee [9] where two-dimensional simulations showed the importance of hole traps in the substrate and

Backgate voltage, Vbg (V)

FIGURE 6. AI^5 as a function of backgate voltage. Note that the position of Vtbg is the same for both curves (after [26]). the influence of bias on Schottky contacts to shield or enhance backgating. Harrison [10] also observed the correlation of increased backgating at high drain biases with short-channel effects in sub-micron GaAs MESFET devices. He attributed the enhanced backgating to hole injection from the MESFET channel into the GaAs substrate. Harrison's conclusions were supported by simulations performed by Horio and Usami [11] where they were able to predict the observed enhanced backgating with hole generation in the substrate. These recent results demonstrate the great complexity of backgating phenomena. A single unified theoretical description of backgating is still not available, but it is clear that substrate traps and hole injection are important to backgating in GaAs integrated circuits. The participation of traps in backgating is indicated by two commonly observed characteristics of backgating. First, experimental data show a correlation between backgating and the type and density of traps in the GaAs substrate [12,13]. Second, backgating increases with decreasing substrate temperature as shown in FIGURE 7. FIGURE 7 shows the temperature dependence of backgating for a 10 ^m wide MESFET with a 4 jim backgating gap [14]. The data show that AI^8 at a given backgating voltage is much lower at -50 0 C than it is at +1000C (54% as opposed to 91% at Vbg= -7 volts). The accepted explanation for this behaviour is that backgating conduction occurs after traps are filled and therefore at low temperatures backgating increases because the number of ionized traps decreases [15]. The backgating phenomena discussed in the above theoretical summary are due to currents that

Backgate bias, l/ bg (V)

FIGURE 7. Variation of AI^8 with backgate bias. Each curve represents a different ambient temperature for the measurement: (•) -500C5 0 ) O0C, (D) 500C and (x) 1000C. (after [14] and [26]). flow between devices through the substrate and the substrate/channel interface. A separate type of backgating phenomenon is called surface backgating and is due to enhanced conduction along the surface area between devices. The two types of backgating are frequently indistinguishable in experimental data and can be coupled through a common variable, such as substrate resistivity [16]. It is important, however, to distinguish between the two types because each must be minimized in a different way. Bulk backgating is most easily reduced by improving substrate material, while surface backgating is most easily reduced by improving process techniques. Section D discusses processing approaches that are employed to reducing backgating. C

MATERIALS

Bulk backgating is intimately related to materials and can be greatly reduced by improving substrate material. Backgating characteristics are most commonly described in terms of substrate traps. As a result, efforts to reduce backgating focus on controlling the type and density of deep-level traps. This task is made difficult by four complications. The first complication is the difficulty of measuring low concentrations of traps in GaAs. The most commonly used method, deep-level transient spectroscopy (DLTS), requires special test structures, and the data it yields is often ambiguous. The test is time consuming and is not suitable for large-scale implementation in production quantities. The second complication is that the correlation between measured trap densities and backgating measurements is often not clear. There are many types of traps in GaAs and still more

types of substrate conduction. Identification of all of these conduction sources is very difficult but this is necessary before an unambiguous correlation can be made. The third complication is that backgating is a localized phenomenon. FIGURE 8 shows the distribution of AI^8 across a three-inch wafer. The distribution shows clear non-uniformities which may be caused by variations in substrate material. Thus, without extensive localized measurements, it is difficult to predict precisely the backgating characteristics of a wafer or a boule. The fourth complication is that processing can change the density and distribution of substrate traps. High-temperature processes used during integrated circuit fabrication (e.g. implant annealing and ohmic contact alloying) are known to cause changes in the distribution of substrate traps [17]. Moreover, because every facility processes the substrate differently, it is possible that substrate material which provides excellent backgating characteristics when processed at one facility, provides poor backgating characteristics when processed at another.

Percentage of sites

Due to the difficulty in obtaining a comprehensive theoretical prediction of backgating, most GaAs MESFET foundries use an empirical approach to predict and monitor backgating. They focus not on understanding backgating, but simply on screening out material that has poor backgating characteristics in their process. Many foundries regularly reject bulk substrate material because backgating structures from processed trial wafers fail a set of backgating specifications. The failure to meet the test specification is unambiguous, but the cause of the failure is frequently unclear, clouded by incomplete data and conjecture. Some wafers fail the backgating evaluation criteria while other wafers, fabricated using the same process and from substrate boules having identical evaluation results (resistivity, DLTS data, surface polish, etc.) pass. It is clear that since backgating results are different on material for which all known boule material parameters are the same, critical parameters are unknown and inadequately controlled at either the wafer production or in the integrated circuit fabrication process.

A/dss (%) FIGURE 8. Histogram of distribution in AI1188 across a wafer with poor backgating characteristics (after [26]).

D

PROCESSING

GaAs integrated circuit engineers have two backgating goals when developing processes. The first is to minimize backgating response and the second is to ensure that the remaining level of backgating is uniform across the wafer and reproducible from wafer to wafer. Numerous processing techniques have been developed to realize these goals. Most of these process approaches can be grouped into three categories: device surface and interface passivation, implant isolation, and field termination barriers. Dl

Surface and Interface Passivation

Surface backgating is the type that is most amenable to solution using processing techniques. This is because the source of surface backgating, the GaAs surface, is localized to the area of the device and it can be modified by changing the processing. All major GaAs integrated circuit fabrication facilities have developed processes that reduce surface backgating to a negligible level by either minimizing surface depletion of compensating traps or by minimizing the concentration of ionized impurities at the surface. The depletion of traps at the surface is corrected by the use of proper surface preparation and annealing conditions and by the use of passivation layers to limit out-diffusion of traps from the semiconductor surface. Each fabrication facility uses its own set of surface preparation and annealing processes to ensure that trap depletion does not occur. Passivation layers of vapour or sputter-deposited SiO2 and Si3N4 are often used to ensure confinement of traps in the substrate during processing. Reduction of ionized impurities at the surface of the device is controlled by careful selection and careful use of process chemicals and procedures. The Si integrated circuit industry has pioneered the development of such procedures and their implementation in GaAs fabrication facilities is common. D2

Implant Isolation

Implant isolation processes locally decrease conduction through the substrate by introducing damage to the crystal lattice. FIGURE 9 shows a schematic cross-section between two devices. The active device region is masked with photoresist and the region between devices is damaged using a high-damage implant. Implant damage locally increases the trap density and thus decreases backgating between implant isolated devices. The first implants used for this purpose in GaAs were protons [18]. Other implant species such as oxygen and boron are used with similar results [19,20]. Nelson et al [14] demonstrated the relationship between backgating and the energy and dose of an isolation implant.

Backgate electrode

Isolation implant

Source

Gate

Drain

n + GaAs

High-resistivity substrate (p-type)

FIGURE 9. Profile of backgating structure illustrating the function of damage implant isolation (after [26]).

Implant isolation serves both goals of the process engineer. The damage implant reduces backgating response as shown in FIGURE 10 . It also provides greater uniformity across a wafer and from wafer to wafer. FIGURE 11 shows backgating data from two wafers, one with five times the isolation implant dose of the other. The structures with a higher dose show a more uniform result than the devices fabricated with a lower dose. The costs for this improvement are

Backgate bias, V^g (V) FIGURE 10. AIj88 as a function of backgate bias for (•) implant isolated and (A) non-isolated structures (after [14] and [26]).

Percentage of sites

an extra masking and implant step (increasing device cost and lengthening process turn-around time) and a decrease in device packing density necessitated by the increased device spacing caused by the implant masking step. Most foundries find that the cost incurred to implement implant isolation is more than compensated by the improved backgating characteristics resulting from isolation. Some foundries producing integrated circuits for applications less limited by backgating (notably MMICs) do not utilize the implant isolation process.

FIGURE 11. Histogram of distribution in AI653 across two wafers processed with (•) a low isolation implant dose and (•) 5 x that isolation implant dose. All other process variables were the same (after [26]). D3

Field Termination Barrier

The other process approach to minimizing backgating, the field termination barrier approach, uses structures that terminate the fields between devices. FIGURE 12 shows the cross-section schematic of one such structure, a p-well barrier, reported by Inokuchi et al [21]. Trap-filledlimited conduction remains unchanged by the barrier, but the fields are terminated by the potential applied at the barrier electrode. Current flows from the backgating contact to the barrier electrode, but the I-V characteristics of the subject device remain the same. There are reports of other barrier-type structures based on the field termination barrier approach being used to limit backgating [22,23]. The major disadvantage to the barrier approach is that it sharply increases device-to-device spacing, thus decreasing device packing density. FIGURE 13 illustrates the comparative FET-to-FET distance achievable with no device isolation, with implant isolation, and with barrier isolation. Typical 1.0 |im gate layout design rules are used to show the relative spacing. E

DESIGN

Some backgating remains even after all possible improvements are achieved in materials and processing. Circuits must be designed to function within these limitations. The first challenge is to document backgating sensitivity through layout design rules that help the designer avoid

Backgate electrode

p-Well barrier electrode

Source

Gate

Drain

n+ GaAs

High-resistivity substrate (p-type)

FIGURE 12. Profile of p-well barrier backgating structure illustrating the function of the p-well barrier (after [26]).

backgating-related failures. The second challenge is to design circuits that are tolerant of backgating phenomena El

Layout Design Rules

The basis of backgating design rules is the scaling of backgating effects with distance and voltage. FIGURE 14 shows an example of how AI^8 increases as the distance between the backgating contact and the device under test increases. Backgating design rules consist of restrictions on the proximity of devices with differing potentials. Some facilities use a design rule that defines a

FIGURE 13. Comparison of maximum MESFET packing distances for three different isolation methods: (a) intrinsic material isolation, (b) damage implant isolation, and (c) p-well barrier isolation. The same set of nominal spacing design rules was used for each case (after [26]).

minimum device spacing independent of the potential difference between the devices. A more sophisticated method uses data such as those shown in FIGURE 14 to scale the minimum devicedevice distance with the potential difference between devices. FIGURE 14 shows that the minimum distance must also be increased for smaller width devices. All backgating design rules must be based on worst-case data to assure proper performance of all designs.

Backgate distance (jam)

FIGURE 14. Variation OfAI^8 with backgate distance for three different MESFET widths: (•) 10 ^m5 (X) 25 urn, and (•) 50 ^m (after [14] and [26]). Another parameter that must be considered in developing backgating design rules is the operating temperature range of the integrated circuit. FIGURE 7 showed that backgating is more severe at low temperatures. As a result, backgating design rules are more stringent for integrated circuits that must operate over the full military temperature range of-55 0 C to +1250C than for circuits designed to operate over a commercial operating range of 0 ° C to +100 ° C. Strict implementation of backgating layout design rules such as those described above permit high yield for VLSI circuits operating over the full military temperature range. Backgating design rules used at Rockwell have permitted the production of circuits such as the high-yield 7 K gate array reported by Kezer et al [24] that are free of backgating effects down to temperatures as low as -55 0 C. E2

Circuit Design

Circuit design techniques have been developed to reduce the impact of backgating on integrated circuit performance and density. Two of these techniques are reducing the power supply voltages in the circuit and using a circuit layout that minimizes close spacing of power buses and devices. The backgating-limited maximum device packing density is partially determined by the magnitude of negative voltages in the circuit. Circuit technologies can achieve higher density by using low power supply voltages. One of the advantages of low-power enhancement/depletion (E/D) logic families is that the power supply voltages are lower, the backgating determined device spacings are decreased, and the device packing density is increased. Vitesse Semiconductor has utilized this approach to achieve circuit densities greater than 100 K gates using a low-voltage DCFL

logic family based on E/D technology [25]. Backgating is reduced in a circuit with a given power supply voltage by laying out the circuit to maximize spacing between device nodes at different voltages. This approach is used as a constraint on high-level layout of circuit functions and as a consideration for layout at the device level. Use of this method to reduce backgating is limited due to the multitude of other constraints driving integrated circuit layout, but designers that consider backgating reduction in integrated circuit partitioning can achieve better performance and smaller circuits. F

TESTING

Testing of backgating is a complicated subject, partially because of the complicated nature of the backgating phenomenon, but also partially because of the variety of test methods used by different groups. Each company defines backgating in terms of its backgating test structures and uses its own unique test method. In general, research facilities utilize extensive backgating tests, while production facilities utilize less extensive backgating tests that can be made and analyzed quickly. In this section, the principal methods used to measure backgating are discussed and some of the difficulties in comparing data taken using the different methods are indicated. A relatively complete description of backgating characteristics for a given wafer should include Aldss as a function of distance, V bg time, device bias, temperature, backgating voltage, and backgating voltage rise time taken at multiple points on the wafer. Acquisition of all these data is too time consuming and expensive for all or even most of the integrated circuit wafers fabricated in manufacturing facilities. Each fabrication group uses a selected subset of this data to characterize backgating to its satisfaction. Many groups use a parameter called the critical backgating voltage (Vcbg), defined as the voltage at which a backgating electrode at a distance d from an FET either reduces I ^ of an FET by a given percentage x or reduces the threshold voltage of an FET by a given percentage, x. Other groups measure AI^8 at a given d and a single backgating bias. Measuring AI4188 at a single bias point is less time-consuming than measuring Vcbg. However, measuring AI488 at a single bias point must be indirectly correlated with the parameter utilized in most backgating design rules: the voltage/distance permitted before backgating becomes extreme. Most production facilities measure AI488 at a single backgating bias and correlate these data with overall backgating performance of the wafer. There is not a generally accepted value for either x or d. The backgating distance, d, is sometimes defined as the nearest spacing distance between the doping implants of two devices. Other definitions of d are: the spacing from metal contact to metal contact, the width of the isolation implant between two devices, or the spacing between the backgating contact doping implant and the device channel implant. The percentage reduction, x, is generally defined to be between 10% and 20%. There are two subtle effects that also must be controlled in order to measure backgating reproducibly: voltage sweep times and illumination. Since backgating is related to trapping phenomena it is strongly time dependent and is sensitive to illumination. FIGURE 15 illustrates the variation of AI488 with sweep time. The value of AI488 changes with sweep rate and is asymptotic at long sweep times. The typical method of normalizing these parameters is to set a

sweep time long enough to be in the asymptotic range and to measure the data with the device in the dark.

(a) top curve

(b) bottom curve

Backgate voltage, Vbg (V)

FIGURE 15. Variation of backdating with sweep time. The data for the two curves were obtained at two different sweep rates for the backgating bias: 72 millisecond/volt (top curve) and 5.3 second/volt (bottom curve) (after [26]). The geometry of backgating test structures is also a critical variable in development of backgating test procedures. Liu et al [6] determined that the placement of Schottky metal on the GaAs substrate can greatly enhance backgating currents in test structures. Position and orientation of backgating contacts is also an important factor in determination of backgating data. It is important to verify that test structures are similar or, better, identical in geometry before data is compared. G

SUMMARY

Changes in materials, processes, and design techniques have reduced backgating to the point that VLSI circuits have been fabricated operating over the military and commercial temperature ranges. However, as VLSI circuit complexity increases and gate lengths decrease, further reductions in backgating will be necessary. New methods to reduce backgating must be accompanied by standard definitions and tests for backgating and an improved understanding of the origins of backgating phenomena. REFERENCES [1] [2] [3] [4] [5]

D. Kiefer, J. Heightley [ IEEE GaAs IC Symp. Digest (1987) p.3-6 ] F. Davis [PC Week (UK) vol.8 no. 19 (1991) p. 150 ] CP. Lee, SJ. Lee, B.M. Welch [IEEE Electron Device Lett (USA) vol.EDL-3 (1982) p.97-8 ] M.A. Lampert, P. Mark [ Current Injection in Solids (Academic Press, New York, 1970 ] ZM. Li, DJ. Day, S.P. McAlister, CM. Hurd [ IEEE Electron Device Lett. (USA) vol.EDL-11 (1990)p.342-5]

[6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26]

Y. Liu, R.W. Dutton, M.D. Deal [ IEEE Electron Device Lett. (USA) vol.EDL-13 (1992) p. 149 ] Y. Ohno, N. Goto [J. Appl. Phys. (USA) vol.66 (1989) p.1217-21 ] N. Goto, Y. Ohno, H. Yano [IEEETrans. Electron. Devices (USA) vol.ED-37 (1990) p.1821-7 ] S.-J. Chang, C-P. Lee [ IEEE Trans. Electron Devices (USA) vol.ED-40 (1993) p.698-704 ] A. Harrison [IEEEElectron Device Lett. (USA) vol.EDL-13 (1992) p.381-3 ] K. Horio, K. Usami [ IEEE Electron Device Lett. (USA) vol.EDL-16 (1995) p.277-9 ] H. Goronkin [Proc. 1983IEEEConf. on High-Speed Semiconductor Devices, Ithaca, USA (1983) p.26-37 ] C. Kocot, CA. Stolte [ IEEE Trans. Electron. Devices (USA) vol.ED-29 (1982) p. 1059-63 ] D.A. Nelson, Y.D. Shen, B.M. Welch [ J. Electrochem. Soc. (USA) vol. 134 (1987) p.2549-52 ] CP. Lee, M.F. Chang [IEEEElectron Device Lett. (USA), vol.EDL-6 (1985) p.428-30 ] [ Private communication with Alan Harrison of Bell-Northern Research ] S. Makram-Ebeid, D. Gautard, P. Devillard, G.M. Martin [Appl. Phys. Lett. (USA) vol.40 (1982) p.161-3] D.C D'Avanzo [ IEEE Trans. Electron Devices (USA) vol.ED-29 (1982) p. 1051-9 ] W.M. Paulson, M.S. Birrittella, T.H. Meirs, K.L. McLaughlin [ IEEE GaAs IC Symp. Digest (1982) p. 166-8] S. Makram-Ebeid, P. Minondo [ IEEE GaAs IC Symp. Digest (1983) p. 142-3 ] K. Inokuchi, M. Tsunotani, T. Ichioka, Y. Sano, K. Kaminishi [ IEEE GaAs IC Symp. Digest (1987) p. 117-20] CP. Lee, M.F. Chang [IEEEElectron Device Lett. (USA) vol.EDL-6 (1985) p.169-71 ] A.S. Blum, L.D. Flesner [IEEEElectron Device Lett. (USA) vol.EDL-6 (1985) p.97-9 ] R.C. Kezer et al [ IEEE GaAs IC Symp. Digest (USA) (1987) p. 193-6 ] R. Nunn [ Des. News (USA) (1991) p.79-80 ] L.G. Salmon [J Mater. ScL Mater. Electron. (USA) vol.3 (1992) p.263-71 ]

20.8 GaAs IMPATT diodes J. Freyer and M. Claassen October 1996

A

INTRODUCTION

IMPATT (impact avalanche transit-time) diodes, first proposed by Read [1] and Tager [2] in 1958, are active two-terminal semiconductor devices for RF-power generation in the frequency range from about 8 GHz up to more than 200 GHz. Since the growing progress of three-terminal devices (FET, HEMT, HBT, etc.) at microwave and lower millimetre-wave frequencies in the last years, the interest in IMPATT diodes below about 60 GHz has been somewhat reduced and their application is now more directed to higher frequencies where the power capability of threeterminal devices is relatively poor. B

PRINCIPLE OF OPERATION

In IMPATT diodes, carriers (electrons and holes) are generated in a high-field region of a reverse biased diode structure due to avalanche multiplication and then drift through an adjacent low-field space-charge region of the device. Due to the inductive behaviour of the avalanche multiplication process and the transit-time delay of the drifting carriers, a phase difference between diode RF voltage and RF current of more than 90 degrees can be realised, i.e. a negative RF resistance is produced which is utilised for RF power generation and amplification. Principally, the active region of an IMPATT device can be separated into two parts: 1. 2.

High-field avalanche region (width wa) for the generation of the carriers. Low-field drift region (width wd) for the drift of the carriers.

For optimum power generation and efficiency the following principles can be stated: The avalanche region wa should be as short as possible to reduce DC-losses and the length of the drift region wd must be adapted to the desired oscillation frequency via the transit-time relation of the drifting carriers. The ratio wa/w (where w = wa + wd) determines severely the power capability of IMPATT diodes. C

BASIC EQUATIONS

Four basic material parameters are important in determining the RF performance of the diodes: 1.

The avalanche multiplication rate a as a function of the electric field, in particular its derivative with respect to the electric field a' = da/dE.

2.

The tunnel generation rate gt as a function of the electric field, especially its derivative with respect to the electric field g/ = dg/dE.

3.

The saturated drift velocity vs which determines the delay of the carriers.

4.

The thermal conductivity which limits the maximum DC-power capability and, accordingly, the DC-current density J0.

For power generation, the impedance of an IMPATT diode must be matched to the load of the resonator circuit. The impedance of a single-drift IMPATT diode in a quasi-static small-signal approximation is given [3] Z = 1/IJcoC, {1- K/G))2}] + l/(G)Cd) [g(0)/{ 1- (G)/G)a)2} - j]

(1)

Ca and Cd are the capacitances of the avalanche and drift region, respectively, g(0) is the transittime relation; g(6) = {l-cos6+jsin6}/6

(2)

0 = cowd/vs

(3)

where 0 is the 'transit angle',

Also, a) is the angular frequency and G)a is the angular avalanche frequency [4] given by, G)a = (3cc'v s yef 2

(4)

In EQN (4), e is the dielectric permittivity. For a given device structure, the negative RF resistance of the IMPATT diode is achieved only at angular frequencies G) > G)a. Above G)a the negative RF-resistance decreases strongly. It follows that the avalanche frequency is also a measure for the maximum oscillation frequency. A rough approximation of the drift region length wd for a desired frequency f, obtained from an optimisation of the transit-time relation (see EQNS (2) and (3)), leads to the technological device design rule for the drift length [3] wd = 3vs/8f D

RELATED DEVICES

Dl

MITATT Diode

(5)

For applications at high frequencies (f > 90 GHz), short avalanche regions (wa < 50 nm) with corresponding high electric field (E > 1 MV/cm) have to be applied. In these cases the carrier generation due to the tunnel effect must be taken into consideration. Devices with mixed avalanche and tunnel generation mechanisms are known as MITATT (mixed tunnel avalanche transit-time) diodes [5-7]. The additional carrier generation by tunnelling compensates to a certain degree the decrease of a' at high electric fields, a' can approximately be replaced by a'eff « / e f f =« / + qgt7J0 where q is the electronic charge.

(6)

The design of a MITATT diode is equivalent to that of an IMPATT diode if a'eff is taken into consideration instead of a . D2

TUNNETT Diode

Generation of the carriers predominantly by tunnelling leads to the TUNNETT (tunnel transittime) diode [8-10] with an in-phase component of carrier injection. A reduced power capability results as compared to the inductive phase of IMPATT and MITATT diodes. The design of a TUNNETT diode differs from that of an IMPATT diode because a wider drift region must be employed due to the missing inductive phase lag. Output power and efficiency of TUNNETT diodes are essentially lower than for IMPATT and MITATT diodes. However lower noise may be expected. E

BASIC MATERIAL PARAMETERS

In this section, the average ionisation rate a, the Zener-tunnelling rate g t , and the carrier drift velocities Vn and vp of GaAs are given as functions of electric field E. These are pure material parameters determined from Monte-Carlo simulations for single carriers under idealised homogeneous field conditions [H]. The scattering mechanisms in the Monte-Carlo algorithm have been adjusted to fit experimental data for different device structures. In actual devices, non-local effects such as 'dead spaces' for impact ionisation, tunnelling channel lengths and velocity overshoot make these material parameters functions not only of electric field but also of device dimensions. They are significant for short lengths below 100 nm and lead to lower overall carrier generation by avalanching and Zener tunnelling (the latter at higher breakdown voltage) as well as higher drift velocity in the generation region. For accurate simulation, Monte-Carlo calculations are appropriate because they may be adjusted by varying the input parameters. El

Impact Ionisation Rate

FIGURE 1 shows the averaged ionisation rate of electrons and holes a = (an-ap)1/2 versus inverse electric field, a can be used as electrons and holes contribute equally in the breakdown regime [H]. In the typical operating range a can be approximated by a = 1.1 x io 6 cm"1 exp(-1.5 x 106 V/cm) E2

(7)

Zener Tunnelling Rate

In homogeneous GaAs5 Zener tunnelling is well described by the Kane formula [12] gt = AE 2 exp(-B/E)

(8)

if an appropriate reduced mass parameter is applied [13]. For 300 K this leads to A =1.42 x 1020V2Cm-1S"1 B =1.70 x 107V/cm

(9) (10)

alpha [1/cm]

1/E[cm/MV] FIGURE 1 Averaged impact ionisation rate of electrons and holes for homogeneous bulk GaAs at 300 K for low doping concentration [H]. The full line corresponds to the experimentally verified range, the broken line to Monte-Carlo extrapolation.

To apply EQN (8) to inhomogeneous electric fields like those in pn-junction devices, gt may be transformed to a tunnelling rate per electron energy interval, and the electric field be expressed by the barrier width between the valence and conduction band at constant electron energy [11], which is well defined wherever tunnelling can occur. E3

Drift Velocities

In FIGURE 2, the drift velocities of electrons and holes are depicted versus electric field [14]. The increase of velocity beyond 300 kV/cm is due to velocity overshoot after ionising collisions. In an actual device it depends strongly on the number of ionising processes. E4

Thermal Conductivity

Besides a' and vs, the thermal conductivity of GaAs [15] is one of the important parameters for GaAs IMPATT diodes. The thermal conductivity of the semiconductor material, together with the heat-sink of the device encapsulation, determine for a given device diameter the thermal resistance of the diode. To avoid thermal degradation, the maximum diode temperature should not exceed ~ 250 0 C. This demand limits the maximum DC power density and thus also the maximum DC current density J0, which in turn limits the avalanche frequency o)a. For optimum heat dissipation, the devices must be thermocompression bonded onto diamond heat-sinks.

drift velocity [cm/s]

E [V/cm] FIGURE 2 : Drift velocities Vn and vp of electrons and holes, respectively, for homogeneous bulk GaAs at 300 K for low doping concentration [14].

F

DEVICE STRUCTURES

Fl

Single-drift Device

The simplest form of IMPATT diode is the so-called single-drift flat-profile structure consisting of a one-sided abrupt p+n-junction, an n-doped layer of definite length and doping concentration and an nn+-junction (p+nn+-structure) [16-18]. Instead of the abrupt p+n-junction a Schottky contact can be used [19]. For these device structures the maximum electric field and therefore the maximum value of a' are determined by the length and doping concentration of the n-layer which represents the generation and drift region. The high-field generation region can be approximated to 0.3 - 0.5 times the total device length depending on the doping concentration [20,21]. With increasing frequencies the total length of the device must be reduced according to the transit-time relation. However, the relative amount of the avalanche region increases which leads to a decrease in output power and efficiency. So, for optimum device behaviour, the ratio wa/w must be reduced by more sophisticated device structures. F2

Read-type Device

A limitation of the high-field generation region is obtained by the application of Read-type device structures, where high-field and low-field regions are separated by different doping concentrations in the respective regions [1,2]. These structures are mainly employed for low-frequency highpower devices [22,23] and for frequencies above 100 GHz [24]. For these more complicated device structures the proper design as well as high-precision fabrication of the doping profiles, e.g. by MBE and mature device technology, are absolutely necessary.

F3

Double-drift Device

Output power and efficiency can be increased if, instead of only one type of carrier (electrons in the case of p+nn+-IMPATT diodes) both electrons and holes, are used in the drift process. This leads to the double-drift structure, where the carrier generating high-field region is in the middle of the device [25,26]. An efficient increase in output power at high frequencies can be achieved if both principles, double-drift and Read-type device, are combined [27,28]. F4

Examples of Device Structures

In the following, examples of the dimensions of IMPATT diodes for use at low and high frequencies, respectively, are given. 1.

An 8 GHz single-drift flat-profile (p+nn+) IMPATT diode. This has a length of the active region (avalanche- and drift-region) of about 5 |um and a diameter of about 150 jum.

2.

An optimised, more complex double-drift Read-type (p+p"pp+n"n+nn"n+) device structure for 170 GHz. This has nine differently doped layers and a total active layer thickness of only 230 nm as well as a diameter of about 20 |um [27].

G

OUTPUT POWER AND EFFICIENCY

cw-output power [W]

The output power of IMPATT diodes is limited by the maximum DC current density and the maximum device area. The relatively poor heat conductivity of GaAs limits the DC current density.

frequency [GHz] FIGURE 3. Rf-output power versusfrequencyfor GaAs IMPATT diodes (A: [28,29], ±: [29],0: [18], o: [30], • : [31], • :[32],x:[33],D:[34]).

The maximum device area is restricted by the minimum matchable impedance as well as by RF losses in the active device and the resonator that is used. FIGURE 3 shows a summary of the best values for output power of cw GaAs IMPATT diodes versus frequency. The best values for the efficiency are relatively high at low frequencies and reach 37 % at X-band frequencies [22,23]. These reduce, however, to about 10% at 90 GHz [29], and reach only about 1% at 170 GHz [28]. H

CONCLUSION

For IMPATT diodes the high-field transport parameters the avalanche multiplication rate, the tunnel generation rate and saturated drift velocities are of major importance. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] II1] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33] [34]

W.T. Read [ Bell Syst. Tech. J. (USA) vol.37 (1958) p.401-46 ] A.S. Tager [ Sov. Phys. Usp. (USA) vol.9 (1958) p.892-912 ] M. Gilden, M.E. Hines [ IEEE Trans. (USA) vol. 13 (1966) p. 169-75 ] H.K. Gummel, J.L. Blue [ IEEE Trans. (USA) vol. 14 (1967) p.563-80 ] M.E. Elta, G.I. Haddad [ IEEE ED (USA) vol.25 (1978) p.694-702 ] ME. Eltaet al [IEEEEDL (USA) vol.1 (1980)p.l 15-6 ] M. Poebl, W. Bogner, L. Gaul [ Electron. Lett. (UK) vol.30 (1994) p. 1316-7 ] J. Nishizawa, Y. Watanabe [ Sci. Rep.Res. Inst. Tohoku Univ. (Japan) vol. 10 (1958) p.91-108 ] M. Poebl, C. Dalle, J. Freyer, W. Harth [ Electron. Lett. (UK) vol.26 (1990) p. 1540-2 ] J. Freyer et al [ Proc. 20th EuMC, Budapest, Hungary, 10-13 Sept 1990 p.599-604 ] C. Benz, M. Claassen, D. Liebig [ J. Appl. Phys. (USA) (1996) ] E.O. Kane [J Phys. Chem. Solids (UK) vol.12 (1959) p.181-8 ] J.B. Krieger [Ann. Phys. (France) vol.36 (1966) p. 1-60 ] D. Liebig [ Proc. 6th Int. Conf. on SISDEP, Erlangen, Germany, 1995 p.74-7 ] J.C. Brice [^Properties of Gallium Arsenide, ^Edition, EMIS Datareviews Series no.2 (INSPEC, IEE, 1990), ch. lp.21-3] R.L. Johnston, B.C. De Loach, B.G. Cohen [ Bell Syst. Tech. J. (USA) vol.44 (1965) p.369-72 ] H. Eisele, J. Freyer [ Electron. Lett. (UK) vol.22 (1986) p.224-5 ] H. Eisele [Microwave J. (USA) (1991) p.275-81 ] S.M. Sze, M.P. Lepselter, R.W. MacDonald [ Solid-State Electron. (UK) vol. 12 (1969) p. 107 ] T. Misawa [ Solid-State Electron. (UK) vol. 15 (1972) p.457-65 ] W.E. Schroeder, G.I. Haddad [ Proc. IEEE (USA) vol.59 (1971) p. 1245-8 ] R.E. Goldwasser, F.E. Rosztoczy [Appl. Phys. Lett. (USA) vol.25 (1974) p.92-4 ] CO. Bozler et al [ Appl. Phys. Lett. (USA) vol.29 (1976) p. 123-5 ] M. Tschernitz, J. Freyer, H. Grothe [ Electron. Lett. (UK) vol.30 (1994) p. 170-1 ] D.L. Scharfetter, WJ. Evans, R.L. Johnston [ Proc. IEEE (USA) vol.58 (1970) p. 1131-3 ] T.E. Seidel, R.E. Davis, D.E. Iglesias [ Proc. IEEE (USA) vol.59 (1971) p. 1222-8 ] M. Tschemitz, J. Freyer [ Electron. Lett. (UK) vol.31 (1995) p.582-3 ] J. Freyer et al [ Proc. 3rd Int. Workshop Terahertz Electronics, Switzerland, 1995, p.22-3 ] M. Tschernitz [ PhD-Thesis at Technical University Munich (1995) ] X. Zhang [ PhD-Thesis at Technical University Munich (1985) ] M.G. Adlerstein, S.L.G. Chu [IEEEEDL Trans. (USA) vol.5 (1984) p.97-8 ] W.R. Wissemann et al [ IEEEED Trans. (USA) vol.21 (1974) p.317-23 ] J.R. Grierson et al [ Electron. Lett. (UK) vol. 15 (1979) p. 13-15 ] Y. Hirachietal [IEEEED Trans. (USA) vol.25 (1978) p.666-74 ]

CHAPTER 21 GaAs IN OPTOELECTRONICS 21.1 21.2 21.3 21.4 21.5 21.6

GaAs based heterostructures for optoelectronic devices GaAs light emitting diodes Quantum well heterostructure lasers GaAs mid- and far- infrared detectors The GaAs solar cell GaAs optoelectronic integrated circuits and future applications

21.1 GaAs-based heterostructures for optoelectronic devices P.K. Bhattacharya May 1996

A

INTRODUCTION

As more stringent performance requirements are being demanded from optoelectronic devices, epitaxial growth techniques need to be modified and perfected to meet these challenges. It is now well-recognized that the use of heterostructures, quantum wells, and superlattices not only vastly improves the performance of more conventional devices, but has also helped in the conception and realization of many new ones. Liquid phase epitaxy (LPE) [1] and vapour phase epitaxy [2] were the mainstay for the growth of optoelectronic device structures for several years and these techniques produced very high purity, defect-free materials. As device design necessitated the incorporation of ultrathin layers, sharp doping profiles, quantum wells and superlattices, it became clear that these growth techniques were inadequate. Fortuitously, the development of epitaxial techniques such as molecular beam epitaxy (MBE) [3] and metal-organic vapour phase epitaxy (MOVPE) [4] have taken place almost concurrently, and these epitaxial techniques can deliver epitaxial layers with unprecedented control of layer composition and thickness. Heterostructure technology is based on chemical modulation in real space using two or more atomic species. Thus a critical ingredient of a heterostructure is the interface between the two materials. As the modulation distance becomes smaller, the role of the interface region becomes more critical. In addition to the interface, the quality of the bulk layers is equally important. Poor quality of the interface and/or bulk regions can cause serious degradation of the potential properties of the heterostructure devices. It is thus extremely important that growth be conducted under conditions such that the heterostructure quality is as good as possible. Strained layers and their heterostructures have emerged as important materials for application in high speed electronic and optoelectronic devices. They provide an additional degree of freedom in bandgap engineering. More recently, it has become clear that in the pseudomorphic regime, large changes can be made in the valence band structure of III-V compounds, which result in dramatic changes in their electronic and optical properties. The growth of pseudomorphic materials by techniques such as MBE and the associated growth modes are also areas where great understanding has recently developed. It is clear, for example, that for large misfits (lattice mismatch > 1.5%) pseudomorphic layers prefer to grow in well ordered three-dimensional islands instead of two-dimensional layers, as is the case for lattice-matched layers [5,6]. B

MATERIALS CHARACTERIZATION

Materials characterization forms an integral part of epitaxial growth and device fabrication. More conventional characterization techniques, that are generally employed to measure the properties of single-layer materials, are not always applicable to multilayered heterostructures. Techniques that are capable of providing microscopic information and that can be used in a feedback loop to improve growth conditions need to be identified. Many measurements have been used for such

purposes including in-situ techniques such as reflection high energy electron diffraction (RHEED) [7-9], and ex-situ techniques such as excitonic photoluminescence (PL) and absorption [10-13], low temperature mobility [14], capacitance-voltage studies [15,16], X-ray studies [17], transmission electron microscopy (TEM) studies [18], etc. Although all these techniques can provide valuable information on heterostructure quality, in situ RHEED and ex-situ photoluminescence have been proven to be extremely useful. Both these techniques are capable of providing microscopic information on the heterostructure quality; RHEED studies provide information on growth modes as well as step density, while excitonic line widths provide information on interface quality as well as on the randomness in alloys grown by MBE. C

PROPERTIES OF QUANTUM WELLS GROWN BY DIFFERENT EPITAXIAL TECHNIQUES

FIGURE 1 shows the low-temperature photoluminescence spectrum of a GaAs/AlGaAs single quantum well [19] grown by MBE. The measured linewidth of 0.1 meV of the excitonic transition is the narrowest ever achieved and agrees with the theoretical limit. Similar results have been obtained with pseudomorphic InGaAs/GaAs single quantum wells (SQW) and multiple quantum wells (MQW) [20]. Examples of photoluminescence and absorption spectra using these materials are shown in FIGURE 2. In addition to the narrow linewidths, the small Stokes shifts also confirm the high quality of the heterointerfaces.

GaAs-SUPERLATTlCE SQW (L z =120A) T= 4.2 K RESOLUTION=0.5A

PL

INTENSITY (ARB. UNITS)

Impressive results have also been obtained with the growth of GaAs and InP-based quantum wells

FIGURE 1. Photoluminescence in a 120 A GaAs single quantum well with superlattice barriers. The superlattice is composed of 30 A GaAs and 43 A Al0 3Ga07As wells and barriers, respectively. The main exciton-related transitions are seen in the spectrum. The peak at 1.5356 eV is composed of light and heavy hole donor-bound excitons. The inset shows luminescence from the superlattice and a weak transition at 1.5996 eV, thought to arise from higher order excitonic transitions.

by MOVPE [21-23]. In a typical system for the growth of GaAs-based materials the epitaxial layers are grown in a low-pressure, RF-heated reactor having a circular or rectangular cross section. The gas handling system is designed for minimum dead space and is carefully pressurebalanced between the reactor run and vent lines. The sources are typically triethylgallium (TEGa), triethylaluminium (TEAl), 10% AsH3 in H2, and alkyl dopant sources. The characteristics of quantum wells grown by MOVPE are comparable to those grown by MBE. Chemical beam epitaxy, CBE [24], is a technique similar to MBE except that all the sources are either alkyl or hydride. Growth is performed under high vacuum in a conventional MBE growth system. Several variations of this technique are being used. For example, the use of solid sources for group III elements and hydrides for the group V elements is termed gas-source MBE (GSMBE); the use of alkyl group III sources and elemental group V sources is termed metalorganic MBE (MOMBE); and the use of alkyl group III sources and hydride group V sources is called chemical beam epitaxy (CBE). All these techniques combine the advantages of the MBE technique with semi-infinite sources. The use of gas sources, particularly for the group III elements, results in low defect densities on the surface of the growing crystal. The alkyl species are usually introduced into the system by injectors, while the hydride species are either injected at room temperature, or cracked in a high temperature cell and effused. Abrupt interfaces [34] are obtained in this technique by switching mass-flow controllers and the vacuum in the growth ambient.

Transmission Ca.uJ

PL Intensity la.u.l

Energy IeVI

Wavelength IA)

FIGURE 2. Low temperature photoluminescence and absorption spectra of 2 ^m In007Ga093As (100 A) GaAs (100 A) MQW structures grown directly on GaAs. The different identified excitonic transitions in the absorption spectra are labelled.

D

LASER HETEROSTRUCTURES

Double heterostructure GaAs/AlGaAs lasers, such as the one shown in FIGURE 3(a), were

Top ohmic contact P+GaAs contact layer P+AlGaAjS p GaA-S active layer H+AlGaAs n + GaAs Bottom ohmic contact

Ohmic contact Oxide Regrown H-AlGaAs

FIGURE 3. Schematics of single mode double heterostructure GaAs/AlGaAs lasers: (a) mesa-etched ridge waveguide, and (b) buried heterostructure.

Inner cladding

Outer cladding

Active region


initially grown by LPE [25] but are now produced by MBE and MOVPE [26,27]. In a semiconductor laser, index guiding normal to the junction plane is provided by cladding layers that

Critical Thickness (A) Wavelength (jam)

In Composition, x FIGURE 5. Variation of critical thickness, laser emission wavelength and threshold current with In composition x in pseudomorphic InxGa1^AsZGaAs quantum well lasers (courtesy: JJ. Coleman).

have a lower index relative to that of the active region. For high performance lasers with low threshold current, it is also important to have index guiding along the junction plane. This is achieved by burying the active region in a material of lower refractive index. Fabrication of these buried heterostructures involves growth over patterned substrates or multiple growths. A buried heterostructure laser is schematically shown in FIGURE 3(b). It is perhaps fair to state that quantum well lasers have demonstrated the highest levels of performance amongst different active region materials. This is due to the step-like density of states. However, at the same time the optical confinement factor is reduced. Although very low transparency and threshold currents have been reported for single quantum well (SQW) GaAs/AlGaAs lasers [28], multiquantum well (MQW) lasers eliminate the effects of gain saturation and exhibit higher gains and differential gains. Therefore, higher modulation bandwidths can be achieved. Typical QW laser heterostructures are shown in FIGURE 4. Separate confinement heterostructures (SCH) and graded refractive index-SCH structures (GRIN-SCH) have now become a standard in high-performance lasers. The well and barrier compositions in the MQW have to be carefully chosen to balance the effects of uniform injection and sufficient confinement. It has been shown that Al019Ga081As barriers provide the lowest threshold currents. Strained quantum well lasers can improve the performance characteristics further through a modification of the valence band structure and density of states [29-41]. Strained quantum well lasers have been fabricated using InxGa1^As quantum well and GaAs barrier layers for x < 0.2 on a GaAs substrate using MOVPE or MBE growth techniques. These lasers exhibit threshold current densities as low as 50 A/cm2 [42] and high CW and pulsed output powers [43] have been reported. These lasers emit in the wavelength range 0.85 to 1 jam. The reduction OfJ111 with In

composition x in an Ii^Ga^As/GaAs/AlGaAs strained MQW laser is depicted in FIGURE 5. Recently the tunnelling injection mechanism has been demonstrated in QW lasers to reduce hot carrier effects [44]. The band diagram and the very high modulation bandwidth achieved in such a laser grown by MOVPE are shown in FIGURES 6(a) and 6(b). Almost all laser heterostructures are grown on (001) GaAs substrates. Due to the slight modification of the band structure in (111) oriented crystals, GaAs/AlGaAs lasers grown on Ill-oriented GaAs have demonstrated lower threshold currents [45]. Lower-dimensional quantum confined structures such as quantum wires and boxes promise lower threshold currents and higher differential gains due to the singular density of states function in these structures. Processing-induced damage and disorder can eliminate most of the expected advantages. The three-dimensional island growth modes mentioned in Section A allow the formation of self-organized quantum boxes, which can then be buried to form the gain region in an SCH laser heterostructure. Room temperature lasing from such a quantum box laser, observed by the author and co-workers, is shown in FIGURE 7. E

HETEROSTRUCTURE AND QUANTUM WELL PHOTODIODES

Energy («V)

Modulation Response (dB)

The specific application and the associated noise and bandwidth considerations determine the choice of photodetection devices. The PIN photodiode based on GaAs/AlGaAs heterostructures is a simple device and has demonstrated low noise and large bandwidth operation [46]. GaAs based devices are not suitable for fibre-optical communication at 1.3 or 1.55 ^m. To achieve photodetection with associated gain, the photoconductor or the avalanche photodiode (APD) is used. The photoconductor is not generally favoured because of its large dark current and associated thermal noise [47]. The APD is a versatile device which exhibits gain resulting from

Distance (nm)

Frequency (GHz)

FIGURE 6. (a) Band diagram and (b) small-signal modulation response of multiquantum well GaAs/InGaAs/AlGaAs tunnelling injection laser emitting at 0.98 ^m.

Intensity (arb.unit)

quantum dots

Substrate n GaA;

Wavelength (A)

FIGURE 7. Room temperature lasing from an In04Ga06AsZGaAs quantum box laser. The quantum box is formed by self organizing growth. The laser heterostructure is shown in (a) and the lasing characteristics of a broad area 90 ^m x 1 mm laser are shown in (b).

the impact ionization process. Avalanche gains of 100 can easily be achieved. A device structure that exhibits high performance and is very versatile is the separate absorption and multiplication (SAM) APD [48] or the separate absorption and graded multiplication (SAGM)APD [49]. These structures combine low leakage, due to the junction being placed in the high bandgap material, with absorption in a lower bandgap material, e.g., GaAs. Admittedly, however, these structures are used more commonly in InP-based devices for long-wavelength applications. The best performance of an avalanche photodiode is obtained when the ionization rates of electrons and holes, ae and och, are very different, i.e. ajah - 0 or °°. Unfortunately, in most compound semiconductor materials ae I <xh « 1. Different ways and means to enhance or reduce the value of a e / ah in artificially structured materials have therefore been explored. A viable multilayered structure, in which enhancement of ae / ah can be obtained, is shown in FIGURE 8(a). It is called a staircase superlattice and in a practical photodiode this material would form the avalanching region. It should be mentioned that the staircase superlattice is difficult to realize, since it involves the growth of precisely controlled quaternary graded layers in each period. In order to understand the process of impact ionization in a staircase superlattice, it is necessary to look at the band profile under an applied bias, as shown in FIGURE 8(b).

FIGURE 8. A staircase superlattice (a) under zero bias, and (b) with an applied transverse bias.

The electrons, which drift from left to right, gain in energy at each step and undergo impact ionization. The process is repeated at each step, resulting in a large multiplication of the number of electrons. The holes, on the other hand, do not encounter such large potential steps and their multiplication remains as in bulk semiconductors. The structure therefore behaves as a solid state photomultiplier of electrons and a large value of ae / ah can be obtained [50]. Similar enhancement of ae / ah (~ 8 - 10) can be obtained in a multiquantum well structure with large well and barrier thicknesses (200 - 500 A) and/or deep wells in the conduction band [51]. Similarly, ac I ah will be enhanced if the band offset is higher in the valence band, resulting in deep-hole quantum wells. The process of preferential multiplication of one type of carrier is similar to that in the staircase superlattice. For example, for the GaAs/Al0 3Ga0 7As MQW, where Aec: A ^ is 57:43, enhancement of ae / ah can still be obtained. In these structures, the electrons are not truly confined in the wells because of their small effective mass. The electron behaviour is similar to that in a bulk semiconductor. The holes, on the other hand, are more confined and experience more scattering in the wells, and their ionization rate can be reduced. The overall effect is an enhancement of ae / ah . Superlattice avalanche photodiodes [52,53], in addition to promising low noise performance and gain, can also provide tunability of the spectral response. Quantum-well based photodiodes have also been developed for long wavelength (> 3 ^m) photodetection. These devices, based on intersubband transitions in a doped well, multiquantum well [54-56], are called quantum well infrared photodetectors (QWIPs). On absorbing a photon, carriers in the wells (conduction or valence band) are raised to a higher confined state, from which they are emitted and contribute to the photocurrent.

F

MODULATORS

For electric fields perpendicular to the quantum well layers, the optical absorption near the bandgap energy can be shifted to lower photon energies without destroying the strong excitonic features in the absorption spectrum. The effect, known as the quantum-confined Stark effect (QCSE) [57], is of great practical interest for absorption modulators and switches. The QCSE in quantum wells is much larger than the Stark effect or Franz-Keldysh effect in bulk semiconductors. A large refractive index variation, based on the QCSE, is predicted in quantum well materials [58], and this makes them extremely attractive for electro-optic devices such as directional couplers and Mach-Zehnder interferometers. It is this possibility that has spurred theoretical and experimental work in this field. The conventional technique of applying a transfer bias across an MQW is to apply a reverse bias across a p-i(MQW)-n diode as shown in FIGURE 9(a). Bias-dependent absorption measured in such a device having 40 periods of GaAs(90 AyAl03GdL07As(SO A) at room temperature is shown in FIGURE 9(b) [59]. The red shift in the absorption spectrum is clearly evident. It may also be

4

-1

Absorption coefficient (10 cm" )

Wavelength (nm)

substrate

Photon energy (eV)

FIGURE 9. (a) Top-illuminated p-i(MWQ)-n modulator and (b) bias-dependent absorption spectra illustrating the quantum confined Stark effect.

noticed that the shift is accompanied by a quenching of the peak height of the excitonic resonances. This is because of the transformation of the electron and hole bound states in the quantum well into quasi-bound states and the corresponding decrease in the overlap of the electron and hole wavefimctions. With proper tailoring of well and barrier composition and thicknesses and device size, amplitude modulation bandwidths of -40 GHz have been achieved [60]. Low-bias modulation can be achieved [61-63] in these devices by using coupled or asymmetric quantum wells Due to the presence in the absorption spectra of the HH and LH excitonic resonances of quantum

well materials, a negative resistance region is produced in the photocurrent-voltage characteristics when the incident photon energy is slightly smaller than the HH peak energy. The potential of this negative resistance region has been exploited to develop a number of photonic switching and logic devices. The first and most important of these is called the self electro-optic effect device (SEED) [64]. This device exhibits photonic switching, bistability, and optically induced oscillations through the negative differential resistance in the photocurrent. By incorporating the MQW in the collector region of a heterojunction bipolar transistor (HBT)5 an integrated controllermodulator device has been developed, as shown in FIGURE 10. Several memory, switching, and logic functions, for very low input levels of light, have been demonstrated with this circuit [65,66]. Due to the lack of inversion symmetry in III-V compounds, a linear electro-optic effect is observed. However, this effect is much weaker than that in lithium niobate. In a quantum well heterostructure, the electro-optic effect is different from that in bulk semiconductors [67-69]; in a quantum well, there will be a strong interaction of the electric field with the optical wave. Because QCSE is a quadratic effect with respect to electric field, the quadratic electro-optic effect is strongly exhibited in quantum wells. It is important to remember that the QCSE is an excitonic effect, and therefore the energy of the light to be modulated should be close to and smaller than the (Q1 - Mi1) transition energy. A quantum well electro-optic phase modulator is typically a guided wave device with 2 - 1 0 periods of quantum wells in the centre of the guiding region. The transverse electric field is applied with a Schottky barrier or junction diode incorporated on the top surface. Phase modulation in GaAs/AlGaAs in quantum well electro-optic modulators has been demonstrated [68-70].

Modulator Controller

n + substrate

FIGURE 10. Integrated controller-modulator switching circuit utilizing QCSE. The MQW is in the collector region of the heterojunction bipolar transistor (HBT). The HBT controller and the p-i-n modulators are grown by on-step epitaxy.

G

CONCLUSION

Heterostructures and quantum wells are used for the realization of high-performance GaAs-based optoelectronic devices. The epitaxial growth and photoluminescence characteristics of these heterostructures have been described, together with optoelectronic device designs incorporating these heterostructures. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] II1] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28]

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21.2 GaAs light emitting diodes E.F. Schubert May 1996

A

INTRODUCTION

Light-emitting diodes (LEDs) are pn-junction devices emitting spontaneous radiation. To date LEDs are the most common compound semiconductor device and the number of manufactured LEDs exceeds the number of all other compound semiconductor devices. This Datareview discusses LEDs with GaAs epitaxial layers and LEDs of other active materials but grown on GaAs substrates. Depending on the epitaxial layer material and composition, the emission wavelengths of GaAs based LEDs include the visible colours, green (500 - 580 nm), yellow (580 600 nm), orange (600 - 620 nm), and red (620 - 680 nm) and infrared (680 - 1000 nm). The GaAs emission wavelength at room temperature is 870 nm. Shorter wavelengths can be attained by epitaxial materials lattice matched to GaAs substrates, such as AlxGa1^As and Ga05In05P. Lattice mismatched materials, for example GaAs1^Px, are used as well, but at the expense of reduced radiative efficiency. The number of applications for LEDs is growing rapidly at the present time. These applications include optical fibre communication, free space communication, indicator lamps, infrared (IR) illumination, visible illumination and automotive applications. More recently, LEDs have been used in flat panel displays, printers, and medical imaging applications This Datareview concerns the fundamentals of LEDs including internal and external efficiency (Section B), LED structures and fabrication processes (Section C), LED materials and wavelength ranges (Section D), and novel device concepts (Section E). B

EFFICIENCY AND BRIGHTNESS OF LEDs

There are several figures of merit for LEDs including the internal efficiency, external efficiency, power efficiency, brightness and luminosity. Depending on the application, different figures of merit are relevant. For example, the brightness is important for communication LEDs, whilst the luminosity is relevant for visible LEDs for observation by the human eye. The figures of merit will be discussed below. The internal quantum efficiency, X]^x, of an LED is defined as the ratio of the number of light quanta (photons) emitted by the active region to the number of charge quanta (electrons) injected into the active region mt

number of internally emitted photons number of injected electrons

A perfect direct-gap semiconductor would have an internal quantum efficiency of unity (100 %). However, native defects, dislocations, and deep level impurities will reduce ri^ to lower values due to Shockley-Read recombination processes. In the 1960s and 1970s, typical internal

efficiencies of GaAs were a few percent (see, for example, [ I ] ) . This value has increased significantly over the last decade and values exceeding 99 % have been reported for O M V P E grown GaAs [2]. Clearly all design and growth parameters must b e optimized t o achieve such high efficiencies. The external efficiency of semiconductor L E D s is defined as the ratio of the number of photons emitted into free space t o the number of electrons injected into the active region ext

number of externally emitted photons number of injected electrons

The highest external efficiency of commercial GaAs L E D s exceeds 50 % (see, for example, [3]). The internal and the external efficiency can be very different, an important reason being that photons emitted by the active region may not escape into free space due t o reabsorption or total internal reflection at the semiconductor surface. The probability that a photon emitted by the active region can leave the L E D is called the escape probability or extraction efficiency. It is the ratio of the external t o the internal quantum efficiency, i.e. extraction efficiency = r|ext / r ^

(3)

As will be discussed below, the extraction efficiency depends strongly on the device geometry. For fibre optics applications, the brightness, i.e. the optical power emitted per unit solid angle, is more important than the total power because the light needs to be coupled into the small core of an optical fibre. The brightness efficiency is defined as _ ^brightness

P

oPt ' j

Q

(4)

where P^is the optical power emitted into the solid angle Q and I is the injection current of the device. The brightness efficiency can be used to calculate the power that can be coupled into a fibre [4]. The power efficiency, also called the 'wallplug efficiency', is the ratio of the optical LED output power to the electrical input power. It is related to the external quantum efficiency by

where V and I are the diode voltage and current, respectively, h is Planck's constant, c the velocity of light in vacuum, and e is the electron's charge. The power efficiency allows one to calculate the heat dissipated in the LED. The luminosity is used for LEDs emitting in the visible wavelength range and it is a measure of the subjective intensity as experienced by the human eye. The sensitivity of the human eye peaks

in the green at X = 555 nm. The relative eye sensitivity, V(A)5 is defined as unity at that particular wavelength, i.e., V(A = 555 nm) = 1 ImAV. The sensitivity decreases on both sides of the peak wavelength as shown in FIGURE 1. The luminosity and the relative eye sensitivity are related by L = 680 V(A) Popt

(6)

(LUMENS/WATT) LUMINOSITY RED

ORANGE

YELLOW

GREEN

BLUE

VIOLET

RELATIVE EYE SENSITIVITY V(X)

Thus an LED emitting 1 mW at 555 nm has a luminosity of 0.680 Im. The luminous efficiency is obtained by dividing the luminosity by the electrical input power to the LED. The unit of the luminous efficiency is ImAV. The luminosity is used only for LEDs emitting in the visible, whereas radiometric units (mW) are used for IR LEDS.

WAVELENGTH X (nm)

FIGURE 1. Relative human eye sensitivity and luminosity as defined by the CIE (Commission Internationale de l'Eclairage) for normal photopic vision [5].

C

LED STRUCTURES AND FABRICATION

The device structure and fabrication process of GaAs-based LEDs strongly influence the escape probability of photons. To illustrate this fact, consider a point-like active region located below a planar semiconductor/air interface. Consider further that the active region emits photons in an isotropic pattern. Such a configuration is shown in FIGURE 2. Light emitted along the normal direction with respect to the semiconductor/air interface can escape from the semiconductor. The transmission probability is given by the Fresnel coefficients. If the angle of the light incident on the semiconductor surface exceeds the angle for total internal reflection, then the light is reflected back into the semiconductor. The solid angle of the light-escape cone is given by Q = 2 TT(I-cos 0 C )

(7)

where 0 C = arc sin (n2 / Xi1) is the critical angle of total internal reflection, and nx and n2 are the refractive indices of the semiconductor and air, respectively. For GaAs, EQN (7) yields 0C = 16.6°. The total fraction of light that can escape from the semiconductor can be calculated by dividing EQN (7) by the solid angle of the unit sphere (4n). One obtains

(8)

If the semiconductor is not coated with an optical anti-reflection layer, Fresnel reflection losses need to be taken into account. For large refractive index differences (Ti1 » n2), the angle of incidence at the semiconductor/air interface is near the surface normal. Including the Fresnel coefficients for normal incidence yields

Using Xi1» n2 and neglecting Fresnel losses, EQN (8) simplifies to

Thus a large index difference is an impediment for the realization of high-efficiency LEDs. As an example, we consider a point-like light source located below a planar GaAs surface and calculate the fraction of light that can escape from the semiconductor. Using nx = 3.5 and n 2 = l , EQN (8) yields Q / 4 n = 2 %. This result illustrates a fundamental problem of high-index LEDs, namely the difficulty of extracting the light from the semiconductor. Photons reflected back into the semiconductor are likely to be lost for external emission due to reabsorption processes, e.g. in the substrate. The inherently high defect density of substrates makes non-radiative recombination the

Air (D= 1) Semiconductor (n = 3.5)

Lightemitting region FIGURE 2. Light escape cone defined by the angle of total internal reflection, 0 C , of a semiconductor/air interface.

dominant recombination process in substrate materials. That is, reabsorption of photons in the substrate decreases the external quantum efficiency. The ideal LED structure would consist of a point-like active region in the centre of a spherically shaped semiconductor coated with an antireflection coating. This geometry would ensure that light emitted from the active region would be incident on the semiconductor surface at an angle of 90°, resulting in the absence of total internal reflection losses [6]. However, such spherical LEDs are incompatible with today's planar semiconductor technology. There are a number of methods to improve the escape probability beyond the 2 % calculated above. These methods include (1) multiple escape cones, (2) mesa structures, (3) transparent epitaxial layers and substrates , (4) lensed structures, (5) reflective contacts, (6) transparent contacts, (7) anti-reflection coatings, and (8) epoxy domes. The structures and fabrication methods will be discussed below. (1)

Multiple escape cones allow for more efficient extraction of light from a high refractive index semiconductor. The principle of multiple escape cones is illustrated in FIGURE 3 (a) which shows a cubic LED with a point-like, light-emitting region in the centre of the epitaxial layer. Light escapes through cones directed towards the top, four sides, and the substrate. Three of the six cones are illustrated in FIGURE 3(a). The escape probability is increased by a factor of six compared to a planar LED.

(2)

Mesa-etched structures further increase the light escape probability (FIGURE 3(b)). It is evident from geometric optics that a cylindrical mesa has the optimum shape. For such a circular mesa, the four discrete escape cones are merged into a single escape ring as shown in FIGURE 3(b). Further improvement results if a hemispherical structure could be used. However, the fabrication of hemispherical structures is difficult and involves, for example, photochemical etching processes [7].

Top escape cone Epitaxial layer

Substrate

Top escape cone Side escape ring

Side escape cone

Substrate

FIGURE 3. (a) Light escape cones of a cubic LED. (b) Top light escape cone and side escape ring of a cylindrical LED.

(3)

Transparent epitaxial layers and transparent substrates reduce reabsorption losses in the epitaxial layers and the substrate. For emission wavelengths X < 870 nm (e.g. LEDs with AlxGa^xAs, (AlGa)0 5In0 5P5 or GaAsP active regions), the emitted light will be reabsorbed in GaAs buffer layers and in the GaAs substrate. Due to the low internal quantum efficiency of substrates, absorbed photons will not be re-emitted. Transparency of the epitaxial layers is achieved by avoiding the growth of semiconductors with a lower bandgap, or by keeping the layer thickness of absorbing layers very thin. Transparent substrates can also be obtained by removal of the absorbing GaAs substrates. In this process, which is also called epitaxial lift-off, the substrate is removed, and the fragile epitaxial film is subsequently bonded to another substrate by means of Van der Waals bonds [8,9]. GaP, glass, and sapphire are suitable transparent substrate materials. Another alternative to achieve transparent substrates is the growth of very thick layers of AlxGa1^As on the GaAs substrate and the subsequent removal of the original GaAs substrate [10]. If the AlxGa^xAs is sufficiently thick (> 100 |im), it has enough mechanical stability to serve as a transparent substrate. The capability of high growth rates make vapour-phase epitaxy (VPE) a suitable technique for the growth of the thick AlxGa1^xAs.

(4)

Lenses have been employed to increase the brightness of communication LEDs [7]. The lenses were made by illuminating the sample with a suitable intensity pattern and by simultaneous photosensitive etching. Lenses on GaAs LEDs have also been used to generate collimated beams for free-space optical interconnects [H].

(5)

Reflective contacts have been used to reflect the light emanating from the active region to the direction of interest. For a 100% reflective mirror, a doubling of the LED brightness can be achieved. Two types of reflectors have been employed, namely metallic reflectors [12] and distributed Bragg reflectors [13].

(6)

Transparent ohmic contacts can be made from semiconducting oxides, in particular InSnO and CdSnO [14]. Absorption occurring in metallic contacts is thereby avoided.

(7)

Anti-reflection (AR) coatings are used on communication LEDs in order to reduce the Fresnel losses. For normal incidence, the reflectivity of an uncoated GaAs/air interface is approximately 30 %. Antireflection coatings with thickness X 14 and refractive index n AR ~ (11GaAs)12 reduce the reflectivity to zero at the wavelength of interest.

(8)

Epoxy domes as shown in FIGURE 4 are used on all commercial LEDS. The epoxy refractive index (n = 1.5) reduces the index contrast and, according to EQN (10), increases the size of the light escape cone. Total internal reflection at the epoxy/air interface does not occur due to the hemispherical shape of the epoxy dome and the resulting perpendicular angle of incidence at the epoxy/air interface.

For communication applications, surface-emitting as well as edge-emitting LEDs are employed. Surface emitting LEDs [12] have diameters of the light-emitting area ranging from 20 \im to 50 Hm. These diameters are suited to efficiently couple light into the core of multimode optical fibres with typical core diameters of 62.5 ^m. Single mode fibres with core diameters of 5 \im require

Epoxy Dome LED Chip

Leads FIGURE 4. Typical LED structure with a hemispherical epoxy dome.

much smaller light-emitting regions. For single-mode fibres, the use of edge-emitting LEDs is advantageous. The structure of edge-emitting LEDs [12] is similar to that of edge-emitting semiconductor lasers including the double heterostructure active region which serves as a waveguide. However, edgeemitting LEDs lack at least one of the two reflectors of a Fabry-Perot cavity laser. The lack of a cavity prevents edge-emitting LEDs from lasing. When the devices are driven to transparency, the heterostructures guide the light emitted into waveguiding modes all the way to the edge of the LED. Thus the light intensity is proportional to the length of the device. The size of the lightemitting area at the edge of the LED is given by the thickness and width of the waveguide. These dimensions can be very small (e.g. 0.5 ^m x 5 \im). The high intensity and the small emission area of edge-emitting LEDs make them well suited for efficiently coupling the output light into singlemode fibres. Superluminescent LEDs [12], also called super-radiant LEDS, are pumped beyond transparency and use stimulated emission and optical gain to further increase the light intensity. The light intensity increases superlinearly with increasing injection current due to stimulated emission processes. Only the absence of reflectors prevents superluminescent LEDs from lasing. The advantages of superluminescent LEDs are the simplicity of fabrication and the higher reliability as compared to lasers. However, like lasers, super-luminescent LEDs are generally temperature sensitive. D

MATERIALS AND WAVELENGTH RANGES

The wavelengths accessible to GaAs LEDs range from the green at A = 560 nm to IR wavelengths of < 1000 nm. These wavelengths are covered by epitaxial layers that are lattice matched to GaAs substrates. TABLE 1 lists the active region materials and their properties for GaAs-based LEDs. Both lattice-matched and lattice-mismatched semiconductors are included in TABLE 1.

TABLE 1. Epitaxial materials for GaAs based LEDs. Active region

Gap

Substrate

Lattice match

Wavelength (nm)

Colour

(AlxGa1J05In05P

D

GaAs

Yes

560-640

Green, yellow, orange, red

AlxGa,_xAs-GaAs MQW

D (x < 0.45)

GaAs

Yes

630-870

Red to IR

AlxGa1^As

D

GaAs

Yes

630-870

Red to IR

(x < 0.45) GaAs06P04

D

GaAs

No

650

Red

GaAs

D

GaAs

Yes

870

IR

GaAs:Zn

D

GaAs

Yes

870-940

IR

GaAs doping SL

I-RS

GaAs

Yes

870-1000

IR

Ga1JnxAs

D

GaAs

Yes

>870

IR

MQW - Multiple quantum well SL - superlattice D - direct gap in k space I-RS - indirect in real space IR - infrared

In the visible wavelength range, (AlxGa1^)0 5In05P, AlxGa1^xAs and GaAs0 6 P 04 are currently used in commercial LEDs [10]. The (AlxGa1^)0 5In0 5P material system covers the colours green, yellow, orange and red. (AlxGa1-J05In05P has been demonstrated to be a highly efficient LED material [15]. AlxGa^xAsZGaAs multiple quantum wells (MQW) and AlxGa1^xAs bulk active regions are suitable for wavelengths X > 630 nm. At X = 630 nm, the Al mole fraction OfAlxGa1^As is x = 45 % and the alloy becomes indirect in k space. As a result, the radiative efficiency drops rapidly for x > 0.45. The active regions of LEDs are typically bulk layers or multiple quantum wells. Single quantum well structures are rarely used for LED applications due to the lower total recombination rates. GaAs0 6P04 is mismatched to GaAs substrates and consequently the epitaxial layer has misfit dislocations which are associated with a large concentration of deep levels [16,17]. Even though the efficiency OfGaAs06P04 diodes is low, the material is used for low-cost LEDS. In the infrared wavelength range (X > 870 nm), GaAs LEDs based on highly Zn-doped GaAs, GaAs doping superlattices, and Ga1^InxAs have been demonstrated. The primary use of GaAs LEDs (X = 870 nm) is for infrared free-space communication (e.g. remote controls). The emission can be extended to longer wavelength by heavy Zn doping of the active region which leads to bandgap narrowing (see, for example, [18]). Doping superlattice LEDs have been demonstrated in the wavelength range 870 nm -1000 nm [19]. The addition of In to GaAs active regions is another possibility to extend the wavelength range of GaAs [20]. However, Ga1^InxAs

is not lattice matched to GaAs. As a result, the efficiency of strain-relaxed diodes decreases with increasing In mole fraction and active layer thickness due to the formation of misfit dislocations and the concomitant generation of deep levels. The doping concentration in the active region determines the minority carrier lifetime and thus the modulation speed in LEDs. Be3 Zn and C are popular p-type dopants in III-V semiconductors. 3 dB frequencies exceeding 1 GHz have been demonstrated for heavily doped GalnP/GaAs LEDs [21]. However, high doping concentrations can also result in high concentrations of native defects which act as nonradiative recombination centres [18]. The internal quantum efficiency of C-doped GaAs active regions was found to be a factor of 20 lower than that of Be-doped GaAs active layers [22]. GaAs LEDs have also been grown on Si substrates. Due to the large lattice mismatch between the Si substrates and the epitaxial GaAs layers, a high concentration of dislocations exist in the GaAs epilayer which act as luminescence killers. Yazawa et al [23] and Sakai et al [24] demonstrated a hybrid growth technology with improved crystallinity and reduced aging rates compared to the direct growth of GaAs on Si. Recently Egawa et al [25] reported 2000 hours of stable operation of InGaP/GaAs/Si LEDs. Rare earth doping in semiconductors is motivated by the possibility of atomic-like optical transitions. Coupling of free electron-hole pairs to the optically active electronic states of rare earth atoms is required for such transitions. GaAs LEDs implanted with neodymium (Nd) have been reported by Chang [26]. Atomic optical transitions near 1.3 |im with an external quantum efficiency of 5 x 10"7 have been reported. E

NOVEL DEVICE CONCEPTS

During recent years, several new LED device concepts have been demonstrated or proposed. These have the potential to significantly increase LED performance. Three new concepts are discussed here: the resonant-cavity LED, the photon-recycling LED, and the single-mode LED. El

The Resonant Cavity LED

Resonant-cavity light-emitting diodes (RCLEDs) are capable of more than 10 times the brightness achievable with conventional LEDs [27,4,9]. RCLEDs employ the enhancement of spontaneous emission achievable by placing the active region into a resonant microcavity. The structure of a GalnAs/GaAs/AlGaAs RCLED is shown in FIGURE 5. It consists of a GalnAs/GaAs strained pseudomorphic quantum well active region located between an Ag-reflector and an AlAs/GaAs distributed Bragg reflector (DBR). These two reflectors define the optical cavity of the device. The fundamental optical mode of the cavity is in resonance with the emission wavelength of the active region. In the case of resonance, the spontaneous emission along the axis of the cavity is strongly enhanced [4]. Resonant cavity enhanced photonic devices were recently reviewed by UnIu and Strite [28]. RCLEDs have been realized in the wavelength range 850 - 950 nm [2931,27,4,32-35] as well as in the visible wavelength range [36,37]. The light power versus current (LI) characteristic of an RCLED is shown in FIGURE 6(a) [4]. Also shown are the typical range of LI curves for the best conventional LEDs (shaded area) and the LI curve of an idealized theoretical device denoted as the ideal isotropic emitter. This planar

Au Ti SiO2 p+-type GaAs p-type AIGaAs p-type AIc2Ga0-8As 4 GalnAs/GaAs QWs n-type AI0.2Ga0.8As

Reflector Confinement Active Region Confinement

12 Period AlAs/GaAs DBR

Reflector

AuGe

FIGURE 5. Structure of a GaAs-based resonant-cavity light-emitting diode (RCLED).

INTENSITY / ftiW/Steradian)

device is assumed to have an isotropically emitting, 100 % internal quantum efficiency active region. Comparison reveals that the intensity of the RCLED exceeds the intensity of the ideal isotropic emitter and of the best conventional LEDs (the ODL50 is a commercial AT&T data link product). The enhancement of the spontaneous emission along the cavity axis makes the RCLED well suited for communications applications [33]. In display applications, quantum efficiencies as high as 16 % have been achieved with RCLEDs [30].

RCLED Contact diam.: 21 jum ?i = 910nm T= 200C

Ideal isotropic emitter

CURRENT / (mA) FIGURE 6. Light intensity vs. current of an RCLED, of the ideal isotropic emitter and of the best conventional LEDs.

E2

The Photon Recycling LED

Photon recycling LEDs employ the recycling of photons in order to increase their efficiency. The physical principle of the device is shown in FIGURE 7. Photons emitted into the escape cone can

be emitted into free space. All other photons are reflected back into the semiconductor and, if the semiconductor does not have any optical losses, will be reabsorbed by the active region. If the internal efficiency of the active material is 100 %, all reabsorbed photons will eventually escape from the semiconductor. FIGURE 7 schematically shows the principle of photon recycling. Photons either escape after the first emission process or after one or several total internal reflection absorption re-emission ('recycling') processes. The physics of photon recycling in optoelectronic devices was first discussed by Stern and Woodall [38]. Schnitzer et al [2,39] showed that high external quantum efficiencies can be achieved by photon recycling in optically pumped structures. Numai et al [40] used Au-coated semiconductor micropillars to enhance photon recycling in vertical-cavity structures.

Active Region Substrate Reflector

FIGURE 7. Illustration of photon recycling processes consisting of emission, total internal reflection, re-absorption and re-emission.

E3

The Single Mode LED

Single mode LEDs are based on an optical resonator structure which has just one single optical mode resonant with the optically active medium. In this case, the entire radiation couples to that particular optical mode. If such a device could have 100% internal quantum efficiency, the L vs. I characteristic of the device would be indistinguishable from that of a laser with a zero threshold current [41]. Therefore, the device has also been termed the 'zero-threshold laser' [40]. The simplest structure of a single-mode LED would be a semiconductor micropillar, whose cavity quality factor Q could be enhanced by high-reflectivity coatings. Even though the zero-threshold laser would have a laser-like L vs. I characteristic, the noise properties would be similar to those of an LED. The single-mode LED is inherently a low current device. A low optical mode density implies that the total volume of the device is very small (i.e. the device volume * A3). Since the current density in a CW operated device is limited to the kA/cm2 range (e.g. 20 kA/cm2), the maximum current injected into the device must be limited as well. Due to the small size of the device, problems associated with the high surface recombination velocity of GaAs need to be addressed.

REFERENCES [I] [2]

K. Seeger [ Semiconductor Physics, 2nd Edition (Springer, Berlin, 1982), p. 404 ] I. Schnitzer, E. Yablonovitch, C. Caneau, TJ. Gmitter [ Appl. Phys. Lett. (USA) vol.62 (1993) P-131]

[3] [4] [5] [6] [7] [8] [9] [ 10] [II] [12] [13] [14] [15] [16] [17] [ 18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33] [34] [35] [36] [37]

Hewlett Packard [ Optoelectronic Designer's Catalog (1996)] E.F. Schubert et al [ Science (USA) vol.256 (1994) p.943 ] S.M. Sze [ Physics of Semiconductor Devices (Wiley, New York, 1981) p. 689 ] W.N. Carr, G.E. Pittman [Appl. Phys. Lett. (USA) vol.3 (1963) p. 173 ] F.W. Ostermeyer Jr., P.A. Kohl, R.H. Burton [ Appl. Phys. Lett. (USA) vol.43 (1983) p.642 ] E. Yablonovitch, D.M. Huang, TJ. Gmitter [ Appl. Phys. Lett. (USA) vol.56 (1990) p.2419 ] F.A. Kish [ Compound Semiconductors 1994 Conference, San Diego, CA (invited talk) ] F.A. Kish [ in Kirk-Othmer Encyclopedia of Chemical Technology, Fourth Edition vol. 15 (Wiley, New York, 1995) ] B. Dhoedt, P. De Dobbelaere, J. Blondelle, P. Van Daele, P. Demeester, R. Baets [ IEEEJ. Lightwave Technology (USA) vol. 13 (1995) p. 1065 ] R A Saul, T.P. Lee, CA. Burrus [in Semiconductor and Semimetals, Eds. W.T. Tsang, vol.22 pt.C eds. RK. Willardson, A.C. Beer (Academic Press, Orlando) (1985) ] T. Kato et al [ J. Cryst. Growth (USA) vol. 107 (1991) p.832 ] L-W. Tu et al [ Appl. Phys. Lett. (USA) vol.58 (1991) p.790 ] H. Sugawara, M. Ishikawa, G. Hatakoshi [Appl. Phys. Lett. (USA) vol. 58, (1991) p.1010 ] M.S. Abrahams, JJ. Tietjen [ J. Phys. Chem. Solids (USA) vol.30 (1969) p.2491 ] GH. Olsen, M.S. Abrahams, CJ. Buiocchi, TJ. Zamerowski [ J. Appl. Phys. (UK) vol.46 (1975) p. 1643] E.F. Schubert [ Doping in III- VSemiconductors (Cambridge University Press, Cambridge, 1993)] E.F. Schubert, A. Fischer, K. Ploog [ Electron. Lett. (USA) vol.21 (1985) p.411 ] W.T. Tsang [Appl. Phys. Lett. (USA)vol.3% (1981)p.661] TJ. de Lyon et al [Appl. Phys. Lett. (USA) vol.60 (1992) p.353 ] TJ. de Lyon et al [ Appl. Phys. Lett. (USA) vol.59 (1991) p.402 ] Y. Yazawa, T. Minemura, J. Asano, T. Unno [ Appl. Phys. Lett. (USA) vol.58 (1991) p. 1292 ] S. Sakai et al [Appl. Phys. Lett. (USA) vol.53 (1988) p.1201 ] T. Egawa, T. Jimbo, M. Umeno [ Appl. Phys. Lett. (USA) vol.67,(1995) p.3605 ] SJ. Chang [J. Appl. Phys. (USA) vol.78 (1995) p.4279 ] E.F. Schubert, Y-H. Wang, A.Y. Cho, L-W. Tu, GJ. Zydzik [ Appl. Phys. Lett. (USA) vol.60 (1992)p.921] M.S. Urdu, S. Strite [ J. Appl. Phys. (USA) vol.78, (1995) p.607 ] J. Blondelle, H. De Neve, P. Demeester, P. Van Daele, G. Borghs, R. Baets [ Electron. Lett. (UK) vol.30 (1994) p. 1787] J. Blondelle, H. De Neve, P. Demeester, P. Van Daele, G. Borghs, R. Baets [ Electron. Lett. (UK) vol.31 (1995) p. 1286] H. De Neve, J. Blondelle, R. Baets, P. Demeester, P. Van Daele, G. Borghs [ IEEE Photonics Technol. Lett. (USA) vol.7 (1995) p.287] E.F. Schubert, N.E.J. Hunt, RJ. Malik, M. Micovic, D.L. Miller [ J. Lightwave Technol., submitted for publication (1996) ] NE.J. Hunt, E.F. Schubert, RF. Kopf, D.L. Sivco, A.Y. Cho, GJ. Zydzik [ Appl. Phys. Lett. (USA) vol.63 (1993) p.2600] S.T. Wilkinson, N.M. Jokerst, RP. Leavitt [ Appl. Optics (USA) vol.34 (1995) p.8298 ] D.L. Huffaker, CC. Lin, D.G. Deppe [Appl. Phys. Lett. (USA) vol.66 no.23 (1995) p.3096 ] J.A. Lott, RP. Schneider Jr., J.C Zolper, KJ. Malloy [ IEEEPhotnics Technol. Lett. (USA) vol.5 (1993)p.631] J.A. Lott, RP. Schneider Jr., G.A. Vawter, J.C. Zolper, KJ. Malloy [ Electron. Lett. (USA) vol.29 (1993)p.328]

Next Page

[38] [39] [40] [41] [42] [43]

F. Stern, J.M. Woodall [ J. Appl. Phys. (USA) vol.45 (1974) p.3904 ] I. Schnitzer, E. Yablonovitch, C. Caneau, TJ. Gmitter, A. Scherer [ Appl. Phys. Lett. (USA) vol.63 (1993) p.2174] T. Numai, H. Kosaka, I. Ogura, K. Kurihara, M. Sugimoto, K. Kasahara [ IEEE J. Quantum Electron. (USA) vol.29 (1993) p.403 ] H. Yokoyama [ Science (USA) vol.256 (1992) p.66 ] D. Faklis [ Opt. Photonics News (USA) vol.6 no. 10 (1995) p.28 ] J. Jahns [ in Optical Computing Hardware, Eds. J. Jahns, S.H. Lee (Academic Press, San Diego, 1994) p. 137]

21.3 Quantum well heterostructure lasers

Previous Page

R.M. Kolbas August 1996

A

INTRODUCTION

Quantum well (QW) laser diodes moved very rapidly from their first demonstration (grown by liquid phase epitaxy) in 1977 [1] into production in the mid 1980s. The speed at which this field developed scientifically, technologically and commercially stems from the unique characteristics and high performance of semiconductor quantum well heterostructures (QWHs). This Datareview was constructed to provide the detailed information needed by the readers to establish their own modelling capability for AlGaAs/GaAs quantum well lasers. A physical model and mathematical technique for calculating the emission wavelengths, confined carrier energies and wave functions are provided. The range of applicability of the model and its success in fitting experimental data are described. Also, design rules are presented with emphasis on the differences between quantum well lasers and conventional heterostructure lasers. Finally, some of the unique phenomena observed in quantum well lasers but not in conventional semiconductor lasers are briefly described. Al

Quantum Size Effects

Quantum size effects occur when the de Broglie wavelength of the confined particle is larger than the spatial extent of the confining potential. To observe quantum size effects the following conditions should be satisfied [2]: kBT < En+1 - E n

(a)

™ < En+1 " En

(b)

ALZ < < AFermi

(C)

(D

where the definitions of the variables can be found in the Appendix. Condition l(a) requires that the difference in adjacent energy levels be larger than the thermal energy. Condition l(b) requires that the mean free path of the particle exceeds the spatial extent, Lz, of the confining potential. Finally, condition l(c) requires that the variation in the thickness of the layer in which the particle is confined should be much smaller than the wavelength of the particle (with an energy equal to the Fermi energy). It is quite easy to meet all of these requirements in a semiconductor quantum well (in contrast to a metal layer that must be much thinner). B

ENERGY LEVELS EV SEMICONDUCTOR QUANTUM WELLS

To determine the energy-momentum or band structure of a semiconductor quantum well (or superlattice) there are two basic approaches:

1)

first principles approach (solve Schrodinger's equation with the full Hamiltonian for the periodic and nonperiodic atomic potentials);

2)

approximate problem by breaking the Hamiltonian down into parts that have already been solved or are readily solved (e.g. separate the periodic and nonperiodic parts).

The second approach is much easier and provides solutions that are in close agreement with experiments. The trick is to start with the band structure of the bulk semiconductor which has continuous energy-momentum relations in the x, y, and z directions and restrict (or quantize) the momentum in one of the directions, say the z direction, due to the potential well. In this approach the energy gaps, electron effective masses and hole effective masses are already available and the problem reduces to finding the bound state of a one dimensional square well. The justification, implementation and success of this approach are explained in Sections B1-B4. Bl

Motivation for the Square Well Model

The Anderson model [3] for semiconductor heterojunctions provides guidance for determining the alignment of the conduction and valence bands in two dissimilar semiconductors in intimate contact. In order to align the Fermi levels in both materials both band bending near the interface and band discontinuities at the interface occur. The spatial extent of the band bending is proportional to the Debye length, L 0 , which is typically several hundred to a few thousand Angstroms. ek F172(T]) D

~ \

e2N

F_ 1/2 (TI)

nondegenerate

6

^T

\ ~^N~

(2)

The extent of the quantum well, however, is typically 200 A or less. Hence, a good approximation is that the quantum well width, Lz, is much less than the Debye length.

FIGURE 1. Schematic representation of an AIxGa^xAsZGaAsZAlxGa1 _xAs quantum well heterostructure. (a) An ultrathin narrow bandgap layer of GaAs is sandwiched between two thick wider-gap Al x Ga 1 ^As layers, (b) Electrons (and holes) are trapped in the thin GaAs layer by a finite square well potential.

Under these conditions the band bending across the quantum well or near the quantum well is so slight that a square well model is justified. It should be kept in mind that band bending will become important if the material is heavily doped (L 0 is small) or the spatial extent of a large multi-quantum well array or a superlattice is comparable to the Debye length. With these approximations a thin layer of a narrow bandgap material such as GaAs sandwiched between wider bandgap layers of AlGaAs, FIGURE l(a), results in finite square well potentials for electrons and holes as shown in FIGURE l(b). These attractive potential wells produce a series of bound states for the electrons and holes as shown in FIGURE 2. The justification for this model, a detailed approach to finding the allowed energies and wave functions, and the success/limitations of the model are presented in Sections B2 - B5.

FIGURE 2. Square well potential that is characteristic of an AlGaAs/GaAs/AlGaAs quantum-well heterostructure. For well thicknesses less than the electron (hole) de Broglie wavelength, size quantization occurs and results in a series of discrete energy levels given by the bound state energies of a finite square well. A potential well exists in both the conduction band and the valence band giving rise to a series of bound states En* for electrons, En1* for heavy holes and Enm for light holes.

B2

Approximating a 3D System as a 2D Electron Gas plus a One Dimensional Square Well

A simple but usefixl approximation to the semiconductor quantum well problem is to start with the expression for the band structure of the semiconductor using the parabolic band (E a k2) approximation: *2k2

*2k2

*2k2

m

E(k x , ky, k z )

which is applicable at the bottom of the conduction band and the top of the heavy and light hole bands. In bulk material the wave vectors Icx, ky and kz are all continuous having values from zero out to the edge of the Brillouin zone. The introduction of a quantum well (thin layer of GaAs between wider bandgap AlGaAs) with thickness in the z direction restricts the allowed values of

kz. For an infinite square well this restriction in kz arises because the wave functions must have wavelengths that exactly satisfy nA/2 = L2, that is, the oscillatory wave functions must go to zero where the potential goes to infinity (like the fixed ends of a vibrating string). The finite well case requires a similar restriction on the wavelengths but the wave function can penetrate into the barriers giving it a slightly longer wavelength and slightly lower energy than the infinite well case. With this restriction on the allowed kz values the energy momentum relation can be written as: ^k x 2 tfk2 E(kx, ky, kz) = — 1 + _ * • + E n 2mx 2my where kx, ky are continuous and where kz is discrete ^2K

_ p

—

«

,

n/hprp

^

[I TC n

E =

—

z

2m z

so

2n

b-

2mzLz2

x.

n,

•

A

n7r

2TT —

—

2L2Zn

Lz

(4)

11X

(infinite well)

The same approach is applicable to the finite square well problem but the energies, En, are not given by a simple closed form solution but must be computed as described in Sections B3 and B4. Using this approach the problem of modelling a semiconductor square well is reduced to a two dimensional gas of electrons (or holes) that is free to move in the x and y directions with effective masses Inx and m^ but is restricted in the z direction by a finite one dimensional square well. The beauty of this approach is that the overall problem is now tractable (without being an expert) and also provides remarkably good agreement with experiment. For the curious it is an interesting exercise to take the limit of this approximation to see if the full three dimensional problem is recovered in the limit of a very thick quantum well [4]. B3

Approaches to Finding the Energy Levels in Square Well Potentials

Discussions of the finite square well problem can be found in text books on Quantum Mechanics [5] and review articles [6]. Numerous graphical and numerical techniques have been used to solve the finite square well problem. (See references 1-23 in [6] of this Datareview.) The principal challenge in solving the one-dimensional-symmetric finite square well potential is that the governing equations are transcendental in the variable of interest, the energy E. ( 2mE)

m

.

( 2mE]

m

L z

{ 2In(V-E)) m ( symmetric states'!

{—) "U-J T H ^ H U= .,3,5.. 0ddJ I 2mE Im

M

0

J f 2mE Im

L z

^

f 2m(V-E) | m ( antisymmetric states^

I M Tj='I—*T-J I n - 2,4,6... even j

Armed with only a simple calculator it is straightforward (but tedious) to solve the single finite square well problem by guessing E values that satisfy the equations to obtain all the allowed

bound state energies En. Finding the bound state energies of multiple square wells requires a more sophisticated approach especially if the potential lacks symmetry. The following technique is both general and easy to implement on a computer. Both the energies En and the wave functions T n are computed. The premise of the technique is that all bound states must have well-behaved wave functions of finite extent (must decay exponentially in the semi-infinite barrier layers). Each region of the potential is assigned a wave function of the form: Y

i

= A ekjZ+B e

i

i "kjZ

(6)

where j is an integer corresponding to each region, Aj and Bj are constants, m is the mass of the particle, and the wave vector k is given by: k- = y^m/Vj - E) / * 2

(7)

The numbering scheme for the index j is shown in FIGURE 3.

FIGURE 3. Example of a complex multiple square well potential that can be solved by the iterative technique presented in the text. The numbering scheme (consistent with EQNS 6-11) for each region and interface is shown.

In regions where the particle's energy is greater than the potential, as in the well where E > Vj and kj is imaginary, the wave functions are oscillatory. In regions where the particle's energy is less than the potential, as in the barriers where E < Vj and kj is real, the wave functions must decay exponentially. The assignment of the wave functions and the coefficients for a single finite square well is shown in FIGURE 4.

FIGURE 4. Analytical expressions for the wave functions in the well region and the confining layers of a single quantum well heterostructure. Since the wave function must decay exponentially in the confining layers the coefficients B0 and A2 must be zero.

In the region j = 0 (z < 0) the B 0 coefficient must be zero (or the wave function would increase exponentially without bound) and the A0 coefficient is arbitrarily set to one. The coefficients in the adjacent regions are found by matching boundary conditions (continuity of the wave function 1 F. = 1P+ and the first derivative *F_/m_ = *P+/m+ of the wave function) at each interface. The introduction of the masses in the second boundary condition is to conserve probability current across the interface. The resulting recursion relations are:

(8)

(9)

for the three regions in a single quantum well, or in general

(10)

(H) The coefficients Aj+1 and Bj+1 can be computed, with equations (10) and (11), by knowing only A1 and Bj and the physical parameters of the quantum well (see TABLE 1). The mass ratio mj+1/mj reduces to 1 if the mass is constant in adjacent layers. The difference in the bound state energies with/without this mass ratio is only a few percent for the AlGaAs-GaAs material system. TABLE 1. Required square well parameters. Square well

Semiconductor square well

potential = Vj

potential for electrons, heavy and light holes bandgaps and electron affinities or bandgaps and % band discontinuities

mass = mj

effective m a s s e s = m e , Hi11115 h^

spatial dimensions = Lj = Zj+1 - Zj

well width(s) = Lj = Z 1+1 -Zj

The bound state energies are found using the following procedure: 1)

Set B 0 = 0 and A 0 = 1 in the semi-infinite barrier (z<0).

2)

Guess an energy E and compute the kj using EQN (7).

3)

Successively compute the coefficients Aj+1 and Bj+1 using the recursion relations in EQNS (10) and (11).

4)

Examine the coefficient Aj in the last region. IfAj is zero (or sufficiently small) then the guess for E is one of the allowed bound states. If not guess again.

A systematic search for the zero values of A^ can be done by bisection or other more advanced root finding techniques. A plot of Aj as a function of energy for single, double, triple and quadruple multi-quantum well systems can be found in Figure 10 of [6]. B4

Calculating Energy Levels in Single and Multiple Semiconductor Quantum Wells

The parameters needed to solve the semiconductor quantum well problem are shown in TABLE 1. The relevant material parameters for the AlxGa1^AsZGaAs system are the band gaps (in units of eV) for the F, L and X conduction bands: Er(x,T) = 1.519+ 1.247x - 5.405 E - 4 T2 Z (T + 204) = 1.519 + 1.247x + 1.147(1 - x)2 - 5.40 E - 4 T2 Z (T + 204)

for x < 0.45 x > 0.45

(12)

(13)

and the electron, heavy hole and light hole effective masses: me(x) HIu1(X) In111(X)

= 0.0665 + 0.0835x = 0.45 + 0.302x = 0.08 + 0.057x

(14)

The discontinuity in the conduction band at the interface between two layers is given by the difference in the respective electron affinities, AEC = Xj+i - Xy The valence band discontinuity can be found from AEg = AEC + AEV. A more common way to present the same information is to specify the % discontinuity in the conduction band. This value is generally accepted to be 60%

FIGURE 5. The electron (or hole) energies in a quantum well heterostructure are described by one discrete quantum number n (kz discrete) and two continuous quantum numbers Icx and ky. Allowed energies are given by the intersection of planes of constant Ic2 with the paraboloid of revolution (bulk E-k relation) as shown in the top diagram. Viewed along the Ic2 axis the same information is plotted as E versus k xor k ybelow. Each parabola represents a subband (constant Ic2) displaced from E = O (the band edge) by En, where En are the bound state energies of a finite square well. Note that there are no allowed energy states at E = 0 in contrast to the bulk E-k relation (narrow curve). Since the two dimensional density of states is a constant, the density of states g(E) increases only when E reaches another subband as shown in the lower right. The resulting stepped or staircased cumulative density of states is less than the bulk density of states. The two dimensional density should be divided by a length (L2) to have the same units to compare with the three dimensional density of states.

in the conduction band and 40% in the valence band for the AlGaAs material system. This leads to two finite square wells, one for electrons in the conduction band and one for holes in the valence band as shown in FIGURE 2. The quantization of the allowed kz wave vectors has a profound effect on the density of states in a quantum well. The discrete nature OfIc2 has the effect of slicing the E-k relation as shown in the top portion of FIGURE 5. Looking down the kz axis the cross sections of these slices appear as a series of stacked parabolas displaced along the energy axis by an amount equal to the bound state energies En as shown in the bottom left of FIGURE 5. Associated with each parabola (arising from the free movement in Icx and ky) is a two dimensional density of states arising from the two dimensional electron (hole) gas. Note that the two dimensional density of states is independent of energy and has units of l/(energy x area). The density of states is zero for energies below the first bound state, a constant value of g0 between E1 and E2, a value of g0 + g0 = 2go between E2 and E3, etc., leading to a stepped or staircased density of states as shown in the lower right of FIGURE 5 [7]. This modified density of states leads to unique and useful optical properties [8]. A density of state diagram for both electrons and holes is shown in FIGURE 6. Optical transitions from like quantum states (An = 0) occur when bound electron states recombine with bound heavy- or light-hole states as shown in FIGURE 6. The energies of these transitions are: T-, e - * h h

T^

T^ e

T^ hh

EUjL = 1hv = E + E n + E n photon n n g

i

T^ e~*lh

i

-,-.

^e

T-, Ih

and E 1 . = hv = E0 + Enn + Enn photon s

/1C. (15)

where the primed transitions correspond to electron to light-hole recombination and the unprimed

An = 0 transitions

Heavy Holes

Light Holes

FIGURE 6. Plot of increasing electron energy and increasing hole energy (downward) as a function of the density of states for a quantum well heterostructure. The smooth parabolas correspond to the three dimensional bulk density of states. Interband optical transitions (n = 0 selection rules assumed) occur from a bound state in the conduction band to a bound state in the valence band. Electron transitions can occur to heavy hole states or light hole states (primed).

transitions correspond to electron to heavy-hole transitions. Broken transitions (An * 0) are much less likely because electron and hole wave functions with different quantum numbers are nearly orthogonal.

Electron Energy (meV)

The electron bound state energies (n = 1, 2, 3, 4) for an Al035Ga0 65As/GaAs/Al0.35Ga0 65As QWH as a function of well thickness are shown in FIGURE 7. Note that the energies increase rapidly as the well width is reduced. Also note that none of the bound states exceed the depth of the conduction band quantum well which is 262 meV for this Al0 35G%65 As/GaAs/Al0 ^5Ga0 65 As quantum well. There is always at least one bound state in a symmetric square well. However, there is no guarantee of higher order bound states (e.g. the n = 2 state does not exist for thickness less than 40 A in FIGURE 7).

Well Width (A) FIGURE 7. Calculated electron energies as a function of layer thickness Lz for a thin GaAs layer sandwiched between two wider bandgap Al035Ga0 65As confining or barrier layers (T = 300 K). The depth of the potential well is 262 me V. Each point on the graph corresponds to a difference of one monolayer of GaAs (half unit cell = 2.8267 A).

The heavy and light hole bound state energies (n = 1, 2, 3, 4) for an Al035Ga065As/ GaAsZAl035Ga065As quantum well as a function of thickness are shown in FIGURE 8. The depth of the valence band potential is 175 meV. Note that there can be more heavy hole than light hole bound states because of the difference in effective mass. The results shown in FIGURES 7 and 8 do not change significantly as a function of temperature since the effective mass has very little temperature dependence and the temperature dependences of GaAs and AlGaAs are nearly the same (Eg (T) « constant). The wavelengths for the electron to heavy hole and electron to light hole transitions at 300 K are shown in FIGURE 9 for an Al0J5Ga0 65As/GaAs/Al0J5Ga0 65As quantum well. This wavelength tunability as a function of layer thickness has greatly increased the range of colours available from semiconductor laser diodes. Unfortunately, the gain spectra of all semiconductor lasers shift with temperature (approximately 2.5-3.0A/°Cfor GaAs). To calculate the emission wavelengths at other temperatures simply shift the energies in FIGURE 9 with the same bandgap dependence as GaAs.

Hole Energies (meV)

Well Width (A) FIGURE 8. Calculated heavy- (solid squares) and light-hole (open circles) energies as a function of layer thickness Lz for a thin GaAs layer sandwiched between two wider bandgap Al035Ga065As barrier layers (T=300 K). The depth of the potential well is 175 meV. Each point on the graph corresponds to a difference of one monolayer of GaAs (2.8267A). T=300K

Emission Wavelength (A)

AIGaAs-GaAs (x=0.35)

Well Width (A) FIGURE 9. Calculated emission wavelength for an Al035Ga0 65As/GaAs/Al0 35Ga065As quantum well as a function of well width at T = 300 K. The transition wavelengths are computed by first using EQN (15) and then using X = 1.239852/E. Note the large shift in emission wavelength as a function of well width especially for ultrathin wells.

B5

Success and Limitations of the Square Well Model

The square well model provides solutions that are in close agreement with experiment from large values of Lz down to a few monolayers. There are many numerical techniques available to compute the bound state energies. The technique presented in Section B3 has the added advantage that the wave functions are also found. These wave functions are actually the envelope functions that modulate the more rapidly varying atomic wave functions associated with the atomic potential used to calculate the band structure. These envelope wave functions can be used to compute selection rules , transition probabilities , and probability densities <*Pn | Yn>. In Section C the spatial extent of the probability densities will be used to

explain the 'unusual' carrier collection characteristics of ultrathin quantum wells. The semiconductor square well model has limitations associated with the physics that was iost' in order to make the problem more tractable. For example, note that some of the light hole curves cross the heavy hole curves in FIGURE 8. These crossings do not actually occur because interaction effects between the heavy and light hole bands prevent the crossings. This is not reflected in the square well model because sufficient physics has not been included. A more fundamental approach is required to obtain detailed band structure and second order effects in the valence band. The earliest work on AlGaAs/GaAs quantum wells was based on an 85%/15% split. Changing the % discontinuity has an impact on the magnitude of the electron, heavy hole and light hole bound state energies (En6, En1*, En111) especially for higher energy states. It does not, however, have much impact on the electron to heavy hole, or the electron to light hole optical transition energies. Hence photoluminescence measurements of these transitions are not an accurate method of determining the band discontinuities. Transitions that involve bound electrons and unbound holes are much more sensitive to differences in the band discontinuities. Some rather subtle strengths of the iterative technique (EQNS (10) and (H)) include: 1)

Simplicity of including an energy dependent effective mass (when calculating the kj) to account for nonparabolicity in the conduction or valence bands.

2)

Simplicity of calculating transmission coefficients (or probabilities) in complex resonant tunnelling structures by using alternative boundary conditions for A0, B0, A1, B1. Likewise long lived resonant states can be found.

3)

Ease of modelling pseudomorphic strained layer quantum well heterostructures (e.g. InGaAs/GaAs). Strain corrections are introduced by modifying the depth of the square well potentials.

Finally, all of the discussions in this review assume that the heterojunction interface is a Type I interface resulting in an attractive potential for both electrons and holes in the same layer. The square well model can be used for Type II misaligned and Type II staggered interfaces but the determination of the potential depths and physical location of the confined carriers is different from that presented in this Datareview [9], C

DESIGN RULES FOR QUANTUM WELL LASERS

Since the first demonstration of infrared and visible semiconductor lasers, numerous advances have been made in their design and performance. Three fundamental design rules were recognized for building low current threshold diode lasers: 1)

Current confinement is required to define a stripe geometry device and to limit the area of heat dissipation.

2)

Carrier confinement is required to hold a sufficiently high density of electron-hole pairs in a small volume to achieve stimulated emission.

3)

Optical confinement is required to guide the optical field (beam quality) and ensure good overlap between the optical field and the region of high electron-hole pair density (optical field aligned with the gain region).

These fundamental design rules are well documented and discussed in numerous books [10] and review articles [H]. Quantum well lasers require additional design considerations. The importance of layer thickness (L2 < de Broglie wavelength) on the emission wavelength has already been described in Section B4. Small dimensions also impact the scattering and subsequent capture of hot electrons and holes as they pass by or through a quantum well. The qualitative aspects of carrier collection and thermalization are described in the remainder of this Section. Cl

Importance of Carrier Collection in Quantum Well Lasers

With the development of quantum well lasers the issue of carrier collection and thermalization within the quantum well emerged as an important design consideration. The initial work on quantum well lasers began, quite fortuitously, with well widths in the order of 200-250 A. These operated quite successfully as photopumped lasers operating on quantum state transitions as high as n = 5 [12]. As the well width was reduced (to shorten the emission wavelength) it became difficult to obtain low threshold single quantum well laser operation for L2 < 80 A. The cause of this difficulty is that the scattering path length for an electron in GaAs (dominated by longitudinal optical (LO) phonon scattering) is approximately 63 A. For conventional lasers with thick active layers of the order of 1000 A the injected electrons and holes are certain to be scattered and

(a)

U >

l e P

Lz > ' p h

(b)

L7, < / p e Lz > /pfc

FIGURE 10. Schematic diagram of the inelastic scattering and subsequent capture of an electron or hole by a quantum well, (a) If the well width is larger than the scattering path length then scattering and capture occur before the carrier traverses the well, (b) If the scattering path length is larger than the well width then the carrier can pass through (over) the well without being captured.

thermalize to the band edge of the GaAs active layer. For a thin quantum well, however, some electrons can pass through (over) the well without being collected if L2 is on the order of the

scattering path length. A schematic representation of this process is shown in FIGURE 10. The solution to this problem was to fabricate multiple-coupled quantum wells (with Lz = 50 A) to ensure that hot carriers would be captured in one of several wells. This approach was a resounding success and multi-quantum well lasers became synonymous with high performance and low threshold. Several years later it was observed that some ultra-thin quantum wells had exceptionally bright spontaneous emission and the question of ultrathin single quantum well lasers was revisited. It was subsequently demonstrated that a quantum well as thin as a single monolayer [13] could be operated as a laser with the optical field travelling parallel (edge emitter) or perpendicular [14] (surface emitter) to the plane of the quantum well. The surprising performance of ultrathin quantum wells as low threshold lasers can be understood if the spatial extent of the wave function is considered rather than the well width Lz. The probability densities <*F | Y> for the bound states of an AlGaAs quantum well of various well widths is shown in FIGURE 11. The full width at half maximum (FWHM) of the probability density at first decreases as the well width decreases but then begins to increase for very thin well widths. This is a very well known effect and it occurs if a particle (wave) is confined in a region of space that is much smaller than the characteristic wavelength. For an electron (hole) in the AlGaAs layer to be captured (bound) by the quantum well the carrier must scatter and lose energy. For a thick well the carrier must scatter within the well where the bound state wave function is located. However, for ultrathin wells the wave function extends considerably outside the well and the carrier can scatter and be captured by the well outside the physical dimensions of the quantum well. Hence, the spatial extent of the wave function, not the physical dimensions

L 2 = 201 A, FWHM = 125A

L2 = 99A, FWHM = 74^

L2 = 51 A, FWHM = 51A

L2 = 8.5A, FWHM = 77 A

FIGURE 11. Probability densities for the bound electron states of quantum wells of various widths (L2). Note that the full width at half maximum (FWHM) of the probability density at first decreases with decreasing well width but then increases for very thin wells. Also, the FWHM of the wave function (not shown) is much wider than the probability density.

of the well, should be considered when evaluating the capture of the carriers.

The calculated full width at half maximum of the probability density for the n = 1 electron state of an AlxGa1^As-GaAs quantum well as a function of well width is shown in FIGURE 12 for various aluminium compositions. For high aluminium compositions (x = 0.45, deep wells) the probability density is well confined and does not begin to spread until the well width is very small (3 monolayers, 8.5 A). For lower aluminum compositions (x = 0.1) the smallest spatial extent of the probability density occurs at 12 monolayers or 34 A and grows rapidly for smaller well widths. This thickness dependence effect is even more dramatic for the wave function (compared to the probability density < T | Y>.

Full Width at Half Maximum Electron Probability Density (A)

n=1 Electron Probability Densities

Well Width (A) FIGURE 12. Calculated full width at half maximum of the probability density of the n = 1 bound electron as a function of well width for five different aluminium compositions. The discrete points are calculated at increments of one monolayer (half the unit cell of GaAs; 2.8267 A). Note that for small well widths the spatial extent of the probability density increases rapidly in large steps.

The calculated full width at half maximum of the probability density for a n = 1 electron, heavy hole and light hole in an Al015Ga085As-GaAs quantum well is shown in FIGURE 13. Note that the shape of the heavy hole curve differs from the electron and light hole curves. This implies that the wave function overlap between an electron and a heavy hole
Probability Density Full Width at Half Maximum (A)

n=1 Probability Densities Electrons Heavy Holes Light Holes

Well Width (A)

FIGURE 13. Calculated Ml width at half maximum of the n = 1 bound electron, heavy-hole and-light hole probability densities as a function of well width (in increments of one monolayer). The shapes of the curves for the electrons and heavy holes are different because of the differences in the effective masses and the relative splitting of the conduction and valence band discontinuities.

rapid thermalization of hot carriers. However, it is also possible to use the carrier collection 'problem' to demonstrate new device functionality [16]. A three terminal dual colour (AA > 500 A) optical emitter was demonstrated based on the selective collection of electrons and holes in two different quantum wells. The optical intensity is varied with a current applied to one terminal while the colour is selected with an applied voltage to the other terminal [17]. The carrier collection process can be both a problem and an asset in the design of light emitting diodes and lasers. D

UNIQUE PHENOMENA IN QW HETEROSTRUCTURE LASERS

The physical dimensions and modified density of states of a QW can have a profound effect on the emission wavelength, optical gain, carrier lifetime, mode confinement factor and optical scattering loss which are all important laser diode design parameters. This section focuses on band filling and phonon assisted stimulated emission in QWHs. Band filling can be observed in thick semiconductor layers but the degree of band filling is relatively small (a few meV). In QWHs band filling can be so extreme as to fill the F conduction band all the way up to the L indirect minima (EL - Ex « 290 meV) or support laser operation more than 250 meV above the GaAs band edge [18]. Prior to these photoluminescence and laser experiments on QWs, it was unprecedented to examine a semiconductor band so far above the band edge with a recombination process. This had previously been the exclusive domain of absorption and reflectance spectroscopy. Simple charge analysis shows that the excess carrier density in the conduction band of a QWH is given by:

e Lz

O6)

where J is the current density or effective current density (watts/photon energy in eV for

photopumped samples) and x is the excess carrier lifetime (« 10"9 sec). Using, for convenience, the usual parabolic density of states the energy to which the band will fill above the band edge is given by: SE - 3.65 x l(T 1 5 — (8n)2/3 eV me

nn, vu)

Thus it is possible to band fill a quantum well hundreds of meV with a reasonable current density so long as the excess carrier lifetime does not become too short due to stimulated emission. Phonon assisted stimulated emission was first observed in AlGaAs/GaAs QWHs [19]. This was quite astounding since phonon assisted transitions do not appear in the spontaneous or stimulated emission spectra of thick GaAs layers. Note, however, that in polar II-VI semiconductor compounds phonon replicas appear in (and sometimes dominate) the spontaneous and stimulated emission spectra. Phonon assisted stimulated emission manifests itself as a strong laser emission peak shifted downward by one longitudinal optical phonon energy (36 meV) from one of the quantum states. The process is most often observed 36 meV down from the n = 1 state but has been observed downshifted relative to higher quantum states or the photon pump energy in photoexcitation experiments [20]. The initial reports of phonon assisted stimulated emission were viewed sceptically by some researchers who had difficulty reproducing the results. However, during the past several years there have been confirming observations in other material systems (InGaAs/GaAs), materials grown by alternative growth techniques and time resolved measurements that distinguish the process from other radiative recombination processes [21]. E

CONCLUSION

The finite square well model provides accurate quantitative results for several key characteristics of quantum well heterostructure lasers. The model is also useful in the qualitative analysis of more complex design variables such as carrier collection. The model and numerical solution of the multiple-finite-square-well problem (Sections B3 and B4) can be implemented by a novice but can be expanded with minor modifications to include sophisticated refinements (Section B5) to satisfy experts in the field. To adequately design a quantum well laser carrier collection must be added to the traditional list of current confinement, carrier confinement and optical confinement. REFERENCES [1] [2] [3] [4] [5] [6] [7] [8] [9] [10]

E.A. Rezek et al [Appl Phys. Lett. (USA) vol.31 (1977) p.288-90 ] B.A. Tavger, V.Ya. Demikhovskii [ Sov. Phys. (USA) vol. 11 (1969) p.644 ] R.L. Anderson [ Solid-State Electron. (UK) vol.5 (1962) p.341 ] M.W. Prairie, R.M. Kolbas [ Superlattices Microstruct. (UK) vol.7 no.4 (1990) p.269-77 ] L.I. Shiff [ Quantum Mechanics (McGraw Hill, 1968) ] R.M. Kolbas, N. Holonyak Jr. [Am. J. Phys. (USA) vol.52 no.5 (1984) p.431-7 ] R. Dingle [ Festk. Prob. XV, Advances in Solid State Physics (Pergamon) p.21-48 ] R.D. Dupuis, P.D. Dapkus, R.M. Kolbas, N. Holonyak Jr. [ IEEE J. Quantum Electron. (USA) vol.15 (1979) p.756-11] L. Esaki [Highlights in Condensed Matter Physics and Future Prospects (Plenum, 1991) p.55-83 ] H. Kressel, K. Butler [ Semiconductor Lasers andHeterojunction LEDs (Associated Press, 1977) ]; H.C. Casey Jr., M.B. Panish [ Heterostructure Lasers pt.A & B (Academic Press, 1978) ]; G.H.B. Thompson [Physics of Semiconductor Laser Devices (Wiley and Sons, 1980) ]

[11]

[12] [13] [14] [15] [16] [17] [ 18] [19] [20] [21]

Special Issue on Semiconductor Lasers [IEEEJ. Quantum Electron. (USA) vol.23 (1987)p.6491089]; Special Issue on Quantum Well Heterostructures and Superlattices [ IEEE J. Quantum Electron. (USA) vol.24 (1988) p. 1579-1791 ] R.M. Kolbas, N. Holonyak, R.D. Dupuis, P.D. Dapkus [ Sov. Tech. Phys. Lett. (USA) vol.4 (1978)p.28-30] J.H. Lee, K.Y. Hsieh, RM. Kolbas [Phys. Rev. B (USA) vol.41 (1990) p.7684 ] S.D. Benjamin, T. Zhang, Y.L. Hwang, M.S. Mytych, RM. Kolbas [Appl. Phys. Lett. (USA) vol.60 (1992) p. 1800-2] H. Shichijo, R.M. Kolbas, N. Holonyak Jr., R.D. Dupuis, P.D. Dapkus [ Solid State Commun. (USA) vol.27 (1978) p. 1029-32 ] D. Zhang, F.E. Reed, T. Zhang, N.V.Edwards, R.M. Kolbas [Appl. Phys. Lett. (USA) vol.63 (1993)p.3367-9] F.E. Reed, D. Zhang, T. Zhang, R.M. Kolbas [Appl. Phys. Lett. (USA) vol.65 (1994) p.570-2 ] RM. Kolbas [ Luminescence Characteristics of Single and Multiple AlGaAs-GaAs Quantum Well Heterostructure Lasers (Ph.D. thesis, U. of Illinois, 1979) p.50-85 ] RM. Kolbas etal[ Solid State Commun. (USA) vol.31 (1979) p. 1033-7] D. Zhang, RM. Kolbas [ Solid State Commun. (USA) vol.98 (1996) p.645-9 ] S.D. Benjamin, J.H. Lee, Y.L. Hwang, T. Zhang, RM. Kolbas [Appl. Phys. Lett. (USA) vol.59 (1991)p.351-3 ]

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APPENDIX: LIST OF S Y M B O L S e E Er EL Ex En FiZ2Cn) g(E) >i kj kx, ky, k z kB LD Lj Lz m x , niy, m z nij m j J1n nij m n N P T Vj x Zj A ^Fermi 1 P x xt X

electron charge energy bandgap for the F conduction band bandgap of the L conduction band bandgap of the X conduction band bound state energy associated with the n* allowed level in a square well potential for electrons, E n e , for heavy holes, En1*, and for light holes, E n m Fermi Dirac Integral of order Vi density of states (as a function of energy referenced from the band edge) Planck's constant/2TT wave vector in the j * region wave vector in the x, y and z direction; k 2 = kx2 + ky2 + k z 2 Boltzmann constant Debye length thickness of each semiconductor layer in the heterostructure thickness of the quantum well effective mass in the x, y and z directions effective mass of a particle in the j * region heavy hole effective mass in the f1 region light hole effective mass in the j m region quantum number for the bound states of a quantum well electron concentration hole concentration temperature in Kelvin potential in the j t h region aluminium mole fraction in Al x Ga^ x As spatial location of the semiconductor heterojunction interfaces wavelength A = 27i/k de Broglie wavelength for a particle at the Fermi energy wave function lifetime of excess carrier relaxation time associated with transport electron affinity

< T n I Ym>

overlap integral f Y * T m dxdydz

21.4 GaAs mid- and far-infrared detectors

Previous Page

S.D. Gunapala and S.V. Bandara July 1996

A

INTRODUCTION

It is customary to make infrared (IR) detectors in the long wavelength range (8 - 20 |nm) by utilizing the interband transition which promotes an electron across the bandgap (Eg) from the valence band to the conduction band. These photo-electrons can be collected efficiently, thereby producing a photocurrent in the external circuit. Since the incoming photon has to promote an electron from the valence band to the conduction band, the energy of the photon (hv) must be higher than the Eg of the photosensitive material. Therefore, the spectral response of the detectors can be controlled by controlling the Eg of the photosensitive material. Detection of very long wavelength IR radiation up to 20 \xm requires small bandgaps down to 62 meV. Examples of such materials meeting these requirements are Hg1^CdxTe and Pb1^xSnxTe in which the energy gap can be controlled by varying x. It is well known that these low bandgap materials are more difficult to grow and process than large bandgap semiconductors such as GaAs. These difficulties motivate the exploration of utilizing the intersubband transitions in multi quantum well (MQW) structures made of large bandgap semiconductors [1-8] (see FIGURE 1).

continuum

energy

conduction band

cross section TEM

bound state photocurrent GaAs

"dark current" mechanisms

position FIGURE 1. Schematic diagram of the conduction band in a bound-to-quasibound quantum well infrared photodetector (QWIP) in an externally applied electric field. Absorption of IR photons can photoexcite electrons from the ground state of the quantum well into the continuum, causing a photocurrent. Three dark current mechanisms are also shown: ground state tunnelling (1); thermally assisted tunnelling (2); and thermionic emission (3). The inset shows a cross-section transmission electron micrograph of a QWIP sample.

B

QUANTUM WELL INFRARED PHOTODETECTORS

An MQW structure designed to detect infrared (IR) light is called a quantum well infrared photodetector (QWIP). An elegant candidate for a QWIP is the square quantum well of basic quantum mechanics [9]. When the quantum well is sufficiently deep and narrow, its energy states are quantized (discrete). The potential depth and width of the well can be adjusted so that it holds only two energy states: a ground state near the well bottom, and a first excited state near the well top. A photon striking the well will excite an electron from the ground state to the first excited state. An externally-applied voltage can sweep it out producing a photocurrent (FIGURE 1). Only photons having energies corresponding to the energy separation between the two states are absorbed in this way, resulting in a detector with a sharp absorption spectrum. Designing a quantum well to detect light of a particular wavelength becomes a simple matter of tailoring the potential depth and width of the well to produce two states separated by the desired photon energy. The GaAsZAlxGa1-XAs material system allows the quantum well shape to be varied over a range wide enough to enable light detection at wavelengths longer than ~ 6 jim [10,11]. Fabricated entirely from large bandgap materials which are easy to grow and process, it is now possible to obtain large uniform focal plane arrays (FPA) of QWIPs tuned to detect light at wavelengths from 6 to 25 jum in the GaAs/AlxGa^xAs material system [12-16]. Currently, there is a great interest in the GaAsZAlxGa^xAs based QWIP due to its high sensitivity, high uniformity, high yield and low cost. Therefore, in this Datareview, we are focusing most of our attention on GaAsZAlxGa1-XAs based QWIPS. Typically each period of the multi-quantum well (MQW) structure consists of a 30 - 100 A well of GaAs (doped n = 1 x 1018 cm'3) and a 500 A barrier OfAlxGa1-XAs [10,11]. Stacking many identical quantum wells (typically 50) together increases photon absorption. Ground state electrons are provided in the detector by doping the GaAs well layers with Si. This photosensitive MQW structure is sandwiched between 0.5 \im GaAs top and bottom contact layers doped at n = 1 x 1018 cm"3, grown on a semi-insulating GaAs substrate by molecular beam epitaxy (MBE). Then a GaAs cap layer about 1 |um thick is grown in situ on top of the photosensitive MQW layers to fabricate the light coupling optical cavity. C

PROPERTIES OF QWIP

Cl

Dark Current

Improving QWIP performance depends largely on minimizing the dark current, the current that flows through a biased detector in the dark, i.e., with no photons impinging on it, and is the parasitic current that plagues all light detectors. In QWIPs, the dark current originates from three different mechanisms [17-21]. As shown in FIGURE 1, the dark current arising from the first process is due to quantum mechanical tunnelling from well to well through the AlxGa^xAs barriers (sequential tunnelling). This process is independent of temperature. Sequential tunnelling dominates the dark current at very low temperatures (<30 K). The second mechanism is thermally assisted tunnelling which involves a thermal excitation and tunnelling through the tip of the barrier into the continuum energy levels. This process governs the dark current at medium temperatures. The third mechanism is classical thermionic emission and it dominates the dark current at higher temperatures (>55 K for 9 |um cutoff QWIPs). The thermal generation rate associated with this current depends on the well doping density and the lifetime of the carriers which will be determined by the thickness of the AlxGa^xAs barriers [22]. Consequently, for QWIPs operating

at higher temperatures the last mechanism is the major source of dark current [16,17]. The best previous QWIPs (pioneered by Levine et al [6] at AT&T Bell Labs) were of the boundto-continuum variety, so-called because the first excited state was a continuum energy band above the well top (typically 10 meV). Later, Gunapala et al designed the bound-to-quasibound quantum well by placing the first excited state exactly at the well top as shown in FIGURE 2. Dropping the first excited state to the well top causes the barrier to thermionic emission (roughly the energy height from the ground state to the well top) to be ~ 10 meV more in the bound-toquasibound QWBP than in the bound-to-continuum one, theoretically causing the dark current to drop by a factor of - 6 at a temperature of 70 K [16]. The dark current-voltage curve of the 8.5 |um peaked bound-to-quasibound QWIP is shown in FIGURE 2. This theoretical reduction in dark current compares well with the experimentally observed factor of - 4 drop for devices having the same peak wavelength. CONTINUUM STATES

VIRTUAL STATE

BOUND-TO-CONTINUUM QWIP

DARK CURRENT (A)

GROUND STATE

DARK CURRENT REDUCTION >x10

AREA = 3.14 x10- 4 cm- 2

BIAS VOLTAGE (V)

CONTINUUM STATES

QUASI BOUND STATE BOUND-TO-QUASIBOUND QWIP GROUND STATE

FIGURE 2. Comparison of dark currents of bound-to-continuum and bound-to-quasibound QWIPs as a function of bias voltage at a temperature T = 55 K.

C2

Responsivity

Typical responsivity spectra of bound-to-bound, bound-to-continuum, and bound-to-quasibound QWIPs are shown in FIGURE 3. Unlike the responsivity spectra of intrinsic infrared detectors, the responsivity spectra of QWIPs are much narrower and sharper due to their resonance intersubband absorption. The normalized responsivity spectra R(A) are given in FIGURE 3 for samples A-F.

RESPONSIVITY R

WAVELENGTH X ftim)

FIGURE 3. Normalized responsivity spectra versus wavelength measured at T = 20 K for samples A-F.

The bound and quasibound excited state QWIPs (samples E and F) exhibit a much narrower AA/A = 10% - 11% than the continuum QWIPs with MIX = 19% - 28% (samples A-D). TABLE 1 gives the responsivity peak Ap and cutoff wavelengths Ac as well as the responsivity spectral width AA. TABLE 1. Responsivity and spectral parameters for samples A-F, including peak responsivity wavelength Ap, long wavelength cut-off X0, spectral width AA, and fractional spectral width AXIX [30]. Sample

A.p (jim)

Xc (urn)

AX (um)

AX/X (%)

A

8.95

9.8

2.25

25

B

9.8

10.7

2.0

20

C

13.2

14.0

2.5

19

D

16.6

19

4.6

28

E

8.1

8.5

0.8

10

F

8.4

8.9

0.9

11

The absolute peak responsivity, Rp, can be written in terms of quantum efficiency r| and photoconductive gain, g, as R15 = CeZhV)TIg

(1)

The bias dependence of Rp is shown in FIGURE 4. Note that at low bias the responsivity is nearly linearly dependent on bias and it saturates at high bias. This saturation occurs due to the saturation of carrier drift velocity. For the longest wavelength sample D, where A c = 19 |im, the dark current becomes too large at high bias to observe the saturation in Rp. The responsivity of a quasibound QWIP (sample F) behaves quite similarly to the continuum QWIPs of FIGURE 4.

RESPONSIVITY R p (AAA/)

BIAS VOLTAGE V b (V)

FIGURE 4. Bias dependent peak responsivity Rp0 (X = Xp) measured at T = 20 K for samples A-D. The insert shows the conduction band diagram.

The fully bound sample E has a significantly different shape: the responsivity does not start out linearly with bias but is in fact zero for finite bias. That is, there is a zero bias offset of more than 1 V, due to the necessity of field assisted tunnelling for the photoexcited carrier to escape from the well [5, 23-25]. C3

Detector Noise

The dark current noise (shot noise) for sample B at temperature T = 77 K is shown in FIGURE 5. The solid circles were measured for negative bias (mesa top negative) while the open circles were measured for positive bias. Smooth curves have been drawn through the experimental data. Note that the current shot noise for positive bias is much larger than that for negative bias (e.g., at Vb = 3.5 V it is 4 times larger), and also that near Vb = 4 V there is a sudden increase in the noise due to a different mechanism (possibly due to the avalanche gain process) [25]. This asymmetry in the dark current noise is due to asymmetry in the dark current for positive and negative bias voltages [10]. The photoconductive gain g can now be obtained using the current shot noise expression [17,19,27] In = / I i I ^ A f

(2)

where Af is the bandwidth (taken as Af = 1 Hz). This expression is valid for small quantum well capture probabilities (i.e., p c « 1). QWIPs satisfy this condition at usual operating bias (i.e., 2 3 V). There have been many studies [28-30] on the optical and electrical gain and the relationship between gain and noise properties of QWIPs. A more general formula [29] which can apply even under low bias conditions where well capture probabilities are high is given by in2 = 4eIdgAf(l - pc/2)

CURRENT NOISE in (pA/VRz)

SAMPLE B

POSITIVE BIAS

NEGATIVE BIAS

BIAS VOLTAGE V b (V)

FIGURE 5. Dark current noise in (at T = 77 K) versus bias voltage Vb for sample B. Both positive (open circles) and negative (solid circles) bias is shown. Smooth curves are drawn through the measured data. The insert shows the conduction band diagram.

C4

Detectivity

Detectivity D* is a commonly used figure of merit and it is defined as [17,19]

D-= R JM

(3)

n

where A is the detector area and Af = 1 Hz. This is plotted as a function of bias for a continuum (A), a bound (E), and a quasibound (F) QWIP in FIGURE 6. (The dashed lines near the origin are extrapolations.) For all three samples D* has a maximum value at a bias between Vb = -2 and -3 V. Since these QWIPs all have different cutoff wavelengths, these maximum D* values cannot be simply compared. The primary noise source in QWIPs is the shot noise produced by the dark current. Therefore, unlike the narrow bandgap detectors in which the noise is dominated by temperature independent processes at low temperatures, QWIP performance can be further improved by cooling to cryogenic temperatures [10].

DETECTIVITY D* (1010 cm VHz/W)

BIAS VOLTAGE V b (V)

FIGURE 6. Detectivity D* (at T = 77 K) versus bias voltage Vb for samples A, E and F. The inserts show the conduction band diagram.

D

LIGHT COUPLING

QWIPs do not absorb radiation incident normal to the surface since the light polarization must have an electric field component normal to the superlattice (growth direction) to be absorbed by the confined carriers. When the incoming light contains no polarization component along the growth direction the matrix element of the interaction vanishes (i.e. e • p z = 0 where e is the polarization and pz is the momentum along the z direction). As a consequence, these detectors have to be illuminated through a 45 ° polished facet [10]. Clearly, this illumination scheme limits the configuration of detectors to linear arrays and single elements. For imaging, it is necessary to be able to couple light uniformly to two-dimensional arrays of these detectors. Several monolithic grating structures [31-35] have been demonstrated for efficient light coupling to a QWIP, and this technology advance has made two-dimensional QWIP imaging arrays feasible. It has been shown that many more passes of IR light, and significantly higher absorption, can be achieved with a randomly roughened reflecting surface [36]. By careful design of surface texture randomization, a factor of eight enhancement in responsivity compared to 45° illumination has been demonstrated experimentally [36]. Naturally, thinning down the substrate enables more bounces of light and therefore higher responsivity. One of the main differences between the effect of the cross grating and the random reflector is the shape of the responsivity curve. Unlike the cross grating, the random reflector has little impact on the bandwidth of the response curve since the scattering efficiency of the random reflector is significantly less wavelength dependent than for the regular grating. Therefore, for QWIPs with random reflectors the integrated responsivity is enhanced by nearly the same amount as the peak responsivity. E

IMAGING ARRAYS

After the grating or random reflector array is defined by lithography and dry etching, the photoconductive QWIP pixels of a large focal plane array (FPA) (i.e., 256 x 256 or larger) can

be fabricated by wet or dry etching through the photosensitive GaAs/AlxGa^xAs MQW layers into the doped GaAs bottom contact layer. Typical pixel-to-pixel distance (i.e., pitch) of these large area FPAs varies from 30 to 50 um. The random reflectors or gratings on top of the detectors are then covered with Au/Ge and Au for ohmic contact and reflection. FIGURE 7 shows twenty five processed QWIP FPAs on a 3 inch GaAs wafer. After indium bumps are evaporated on top of the detectors for Si readout circuit hybridization, these QWIP FPAs are hybridized (via the indium bumpbonding process) to silicon CMOS readout multiplexers to yield the final imaging FPA units. At temperatures below 72 K, the signal to noise ratio of the QWIP FPAs is limited by array nonuniformity, multiplexer readout noise, and photo current (photon flux) noise. At temperatures above 72 K, temporal noise due to the QWIP's higher dark current becomes the limitation. As mentioned earlier this higher dark current is due to thermionic emission and thus causes the charge storage capacitors of the readout circuitry to saturate. Since the QWIP is a high impedance device, it should yield a very high charge injection coupling efficiency into the integration capacitor of the multiplexer. In fact charge injection efficiencies approaching 90% have been demonstrated [12,16]. These FPAs have been back-illuminated through the flat thinned substrate membrane (thickness = 1300 A). These QWIP FPAs gave excellent images with 99.98% of the pixels working (number of dead pixels « 10) , demonstrating the high yield of this GaAs technology [15,16]. The operability was defined as the percentage of pixels having a noise equivalent differential temperature of less than 100 mK at 300 K background and the usual operability of GaAs based QWIP FPAs closely agrees with the pixel yield. FIGURE 8 shows a typical noise equivalent differential temperature (NEAT) histogram of a QWIP FPA at an operating temperature of T = 70 K at 300 K background with a mean value of 26 mK [16]. A typical uncorrected photocurrent

FIGURE 7. Twenty five 256 x 256 QWIP focal plane arrays on a 3 inch GaAs wafer.

NUMBER OF PIXELS

non-uniformity of a large area QWIP FPA is about 5% (= sigma/mean) [16]. The non-uniformity after two-point (17° and 27°C) correction improves the uniformity to an impressive 0.05%. As mentioned earlier, this high yield is due to the excellent GaAs growth uniformity and the mature GaAs processing technology.

NEAT (mKELVIN)

FIGURE 8. Photosignal histogram of the 65,536 pixels of the 256 x 256 array showing the high uniformity of the FPA. The uncorrected non-uniformity (= standard deviation/mean) of the FPA is only 6.8% including 1% non-uniformity of ROC and 1.4% non-uniformity due to the cold-stop not being able to give the same field view to all the pixels in the FPA.

F

OTHER TYPES OF MQW IR DETECTORS

Fl

Miniband Transport Multi Quantum Well Infrared Detector

The miniband transport (MBT) IR detector, which is functionally equivalent to the bound-tominiband design [37], uses doped GaAs quantum wells containing two bound states separated by short period AlxGa^xAs superlattice barrier layers. The quantum wells are designed such that the higher energy level is resonant with the ground state miniband in the superlattice barrier. The infrared radiation is absorbed in the doped quantum wells, exciting an electron which is transported in the miniband and generates photocurrent. The structure parameters of the quantum well (well width and barrier height) have some flexibility because it is possible to obtain the same operating wavelength with a continuous range of well widths and barrier heights [14,38]. The performances of 128 x 128 and 256 x 256 FPAs based on GaAsZAlxGa1^As MBT MQWs demonstrate that these are very competitive with the other LWIR technologies. The detectors typically have a peak wavelength of 9.1 ^m and spectral bandwidth of 1.2 ^m and values of peak D* of 2 x io10 cm Hz172AV at 77 K. The temporal and fixed pattern NEAT of 15 mK and 30 mK, respectively, have been reported on these FPAs operating at 60 K. F2

Infrared Hot Electron Transistor

The infrared hot electron transistor (IHET) [39] is an infrared sensitive three terminal device which utilizes an IR sensitive GaAs/AlGaAs multi quantum well structure as emitter, a wide (1500 A) GaAs layer as base, and a thick quantum barrier placed in front of the collector as an electron energy high pass filter. The energy filter is designed to selectively filter higher energy electrons

(mostly photocurrent) to the collector and to reject lower energy electrons (mostly dark current) which are drained through the base. In general, an MET can be considered as a two-stage device, in which each stage has a unique function. The emitter stage is designed to give desirable optical properties, and the collector stage is designed to improve its electrical properties. By using an InGaAs layer instead of a GaAs layer as the base, a value of D* at 77 K of 1.4 x io 10 cm Hz172AV for a device with cut-off wavelength 9.5 ^m has been reported [40]. In addition to a reduction of dark current, these devices are expected to reduce readout noise due to their large output impedance. Although substantial research has been carried out on these IHET devices, an FPA or an infrared imager based on such a device is yet to be demonstrated. F3

p-Doped QWIPs

In all the work discussed in this Datareview the quantum wells were doped n-type. For these ndoped GaAs quantum wells the quantum mechanical selection rule forbids normal incident absorption (Section D). For p-doped QWIPs, however, the strong mixing between the light and heavy holes in the valence band permits normal-incidence absorption. In order to take maximum advantage of this, the GaAs quantum wells have to be doped heavily (typically 4 x io 18 cm"3 with Be). A normal incidence quantum efficiency of 28% and detectivity of D* = 3 x io 10 cm Hz172AV at T = 77 K, for a cutoff wavelength of 7.9 (im, has been achieved for bound-to-continuum p-type QWIPs [41]. It is worth noting that the optical gain of p-type QWIPs is over one order of magnitude smaller than the corresponding value for n-type QWIPs due to the lower carrier velocity associated with higher hole effective mass. G

CONCLUSION

The basic advantages of the GaAs based LWIR detectors, namely the highly mature GaAs growth and processing technologies, become more important at longer wavelengths where narrow bandgap materials become more difficult to work with. Exceptionally rapid progress has been made in the performance (i.e., detectivity, NEAT, minimum resolvable temperature difference, uniformity, etc.) of long wavelength QWIPs (i.e., 6 - 2 5 |nm), starting with bound-to-bound QWIPs which had relatively poor sensitivity, and culminating in high performance bound-toquasibound QWIPs. Extremely good progress has been reported in light coupling schemes starting from the 45° polished face to efficient cross gratings and random reflectors. Detectivities higher than Ix 1011 cm Hz1/2/W have been achieved with 8 - 10 jxm QWIPs at 77 K and 13 - 15 |um QWIPs at 50 K. These operating temperatures can be easily achieved by single stage Stirling coolers. Due to the high performance and excellent uniformity of GaAs based QWIPs, several groups [1216] have demonstrated IR imaging cameras based on large (128 x 128, 256 x 256, and 640 x 480 pixels) QWIP arrays up to a cut-off wavelength of 15 |im. Due to the availability of large area FPAS, easy manufacturability, high internal impedance, low 1/f noise, high radiation hardness, low cost (i.e., high yield), easy hybridization to readout electronics, and the high uniformity of QWIPs, they are potential candidates for a large variety of ground-based and space-based IR applications.

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L. Esaki, H. Sakaki[/BM 7ec/2. Disci. Bull. (USA) vol.20 (1977) p.2456] J.S. Smith, L.C. Chiu, S. Margalit, A. Yariv, A.Y. Cho [ J Vac. Sci. Technol. B (USA) vol.1 (1983)p.376] D.D. Coon,R.P.G. Karunasiri [ y^p/. Phys. Lett. (USA) vol.45 (1984) p.649 ] L.C. West, S.J. Eglash [/*/>/>/. Phys. Lett. (USA) vol.46 (1985) p. 1156] B F . Levine, KK. Choi, CG. Bethea, J. Walker, RJ. MaMk [Appl. Phys. Lett. (USA) vol.50 (1987) p. 1092] B.F. Levine, CG. Bethea, G. Hasnain, J. Walker, R.J. Malik [ Appl. Phys. Lett. (USA) vol.53 (1988)p.296] D.D. Coon, K.M.S.V. Bandara [ in Physics of Thin Films, Eds. M.H. Francombe, J.L. Vossen, (Academic, New York, 1991) vol. 15 ] K.K. Choi [J. Appl. Phys. (USA) vol.73 (1993) p.5230 ] C. Weisbuch [ in Semiconductors and Semimetals, vol.24, Ed. R. Dingle (Academic Press, NY, 1987)p.l-133] S.D. Gunapala, K.M.S.V. Bandara [ inPhysics of Thin Films, Eds. M.H. Francombe, J. L. Vossen vol.21 (Academic Press, New York, 1995) p. 113-237] B.F. Levine [J. Appl. Phys. (USA) vol.74 (1993) p.Rl ] CG. Bethea et al [ IEEE Trans. Electron Devices (USA) vol.40 (1993) ] LJ. Kozlowski et al [ IEEE Trans. Electron Devices (USA) vol.ED-38 (1991) p.l 124 ] W. A. Beck et al [ Second International Symposium on 2-20 fxm Wavelength Infrared Detectors and Arrays: Physics and Applications, October 10-12, 1994, Miami Beach, Florida ] S.D. Gunapala et al [ IEEE Trans. Electron Devices (USA) (to be published) ] S.D. Gunapala et al [ IEEE Trans. Electron Devices (USA) (to be published) ] B.F. Levine et al [ Appl. Phys. Lett. (USA) vol.56 (1990) p.851 ] S.D. Gunapala, B.F. Levine, L. Pfeiffer, K. West [ J Appl. Phys. (USA) vol.69 (1990) p.6517 ] A. Zussman, B.F. Levine, J.M. Kuo, J. de Jong [ J. Appl. Phys. (USA) vol.70 (1991) p.5101 ] E. Pelve et al [ J Appl. Phys. (USA) vol.66 (1989) p.5656 ] M.A. Kinch, A. Yariv [Appl. Phys. Lett. (USA) vol.55 (1989) p.2093 ] E. Rosencher, F. Lue, P. Bois, J. Nagle, Y. Cordier [ Appl. Phys. Lett. (USA) vol.63 (1993) p.3312] B.F. Levine, CG. Bethea, K.K. Choi, J. Walker, RJ. Malik [ Appl. Phys. Lett. (USA) vol.53 (1987)p.231] K.K. Choi, B.F. Levine, CG. Bethea, J. Walker, RJ. Malik [ Phys. Rev. Lett. (USA) vol.59 (1987) p. 245 9 ] N. Vodjdani, B. Vinter, V. Berger, E. Bockenhoff, E. Costard [ Appl. Phys. Lett. (USA) vol.59 (1991)p.555] B.F. Levine, K.K. Choi, CG. Bethea, J. Walker, R. J. Malik [Appl. Phys. Lett. (USA) vol.51 (1987)p.934] G. Hasnain, B.F. Levine, S. Gunapala, N. Chand [ Appl. Phys. Lett. (USA) vol.57 (1990) p.608] H.C Liu [Appl. Phys. Lett. (USA) vol.61 (1992) p.2703 ] W.A. Beck [.4/7/7/. Phys. Lett. (USA) vol.63 (1993) p.3589] K.K. Choi [Appl. Phys. Lett. (USA) vol.65 (1994) p. 1266 ] K.W. Goosen, S.A. Lyon [ Appl. Phys. Lett. (USA) vol.47 (1985) p. 1257 ] G. Hasnain, B.F. Levine, CG. Bethea,RA. Logan, J. Walker,RJ. Malik [Appl. Phys. Lett. (USA) vol.54 (1989) p.2515] J.Y. Andersson, L. Lundqvist, Z. F. Paska [Appl. Phys. Lett. (USA) vol.58 (1991) p.2264 ] J.Y. Andersson, L. Lundqvist [Appl. Phys. Lett. (USA) vol.59 (1991) p.857] J.Y. Andersson, L. Lundqvist, Z. F. Paska [ J Appl. Phys. (USA) vol.71 (1991) p.3600 ] G. Sarusi, B.F. Levine, S. J. Pearton, K.M.S.V. Bandara, R.E. Leibenguth [Appl. Phys. Lett. (USA) vol.64 (1994) p.960]

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2L5 The GaAs solar cell J.E. Parrott August 1995

A

INTRODUCTION

At a very early stage in the development of solar cells it was realised that, in terms of efficiency, there is an optimum energy gap [1] and that the gap of gallium arsenide was closer to this optimum than that of silicon. It was therefore disappointing when the first GaAs cells to be made showed no advantages [2] compared with contemporary silicon devices. The cause of this was quickly understood to be excessive recombination at the illuminated surface. Two approaches were adopted to overcome this. Firstly, improvements were made in the design of the cells [3,4] and, secondly, materials studies led to some degree of passivation of the front surface. As a result, there were improvements in efficiency which now considerably exceeds that found in silicon. The overriding condition for the widespread terrestrial use of photovoltaic energy conversion is the reduction of its cost. This requirement is particularly severe as far as GaAs is concerned because of the high cost of the basic material. The situation is different where space applications are being considered; here specific performance (W/kg) and radiation hardness are more important and, recently, GaAs devices have been showing themselves to be increasingly competitive with silicon solar cells [5]. Despite its problems, GaAs has not been ignored for terrestrial photovoltaic conversion. One possible method of reducing the importance of material costs is to use concentrated sunlight [6] so reducing the cell area needed to produce a given quantity of power. Another method is to make thin film cells [4]; the high optical absorption coefficient close to the edge facilitates this approach. The fragility of GaAs has meant that thin film cells are often deposited on foreign substrates. Finally, GaAs and related materials are often used as a section of a multigap (tandem) cell structure; such structures may be expected to achieve efficiencies higher than those obtainable from single unit solar cells. All of the topics listed in the above two paragraphs are discussed at greater length in what follows. Throughout there will be a large number of references to the solar energy conversion efficiencies reported by different laboratories. The efficiency found for a particular cell depends considerably on the experimental conditions under which it was measured. The principal relevant factors are: (a)

The illuminating spectrum: this is often specified in terms of Air Mass (AM). AMO is the spectrum outside the earth's atmosphere while AMI.5 is used as a standard for terrestrial work. It corresponds to the spectrum found when the sun is at an elevation of about 45 °. A further distinction is made between global (AMI .5G) and direct (AMI .5D) illumination

Vl (b)

The intensity of illumination. The standard spectra have particular intensities associated with them but, in addition, the cell may be designed for, and therefore tested in, concentrated sunlight.

(c)

The temperature of the cell when the measurements are made: the standard is 25 0 C. To enable proper comparisons to be made all of the above should be specified when reporting new results. It is also desirable that cells should be independently tested at a centre recognised for this purpose.

One completely different application which is worth a mention concerns the conversion of light generated by lasers or some other bright source into electric power. The distinguishing feature of this type of conversion is that the light is essentially monochromatic. Spitzer et al [8] describe a suitable system for this purpose making use of gallium arsenide. B

THEORETICAL LIMITS FOR GaAs SOLAR CELLS

The fundamental theoretical limit for photovoltaic energy conversion was first investigated in a classical paper by Shockley and Queisser [I]. This was based on the fact that the minimum dark current must be that due to radiative recombination which cannot be eliminated since it arises from the same physical process which makes possible the absorption of radiation and the generation of electron hole pairs. The authors called this the 'detailed balance' limit. They showed that, for a given illumination, the maximum efficiency was a function of the semiconductor energy gap. There was an optimum value of this parameter close to that for GaAs when this illumination was due to blackbody radiation characterised by the temperature of the sun's surface. The forecast upper limit is 30% for one sun illumination and 34% at 200 suns. Although the calculations in [1] have been corrected and refined, this has not led to any substantial change in these figures. However there is another intrinsic recombination process, ignored by Shockley and Queisser, which has been shown to be of considerable importance. This is the Auger mechanism which, in the case of silicon, is believed to be decisive in setting the efficiency limit [9,10] rather than radiative recombination. One effect of this is the efficiency which no longer increases with increasing concentration but levels off at about 36%. Because of its direct gap, radiative recombination is generally much more significant in GaAs but there are circumstances in which Auger processes dominate in this material also [H]. If limitations of acceptance angle can be combined with a back surface reflector, GaAs shows distinct advantages over silicon at rather high concentrations (>1000), the efficiency levelling off at 39%. The superiority of GaAs is even more marked if temperature effects are taken into account. As yet, it still remains to be seen whether these advantages can be realised in practice. Because of its direct gap, GaAs has a very high optical absorptivity close to the absorption edge. 97% of incident AMI photons whose energy exceeds the gap are absorbed within 2 microns of the surface making it possible to use much less material than is required for silicon solar cells. C

MODELLING GaAs SOLAR CELLS

The considerations outlined in the previous section provide reasons for choosing GaAs as a photovoltaic material but do not tell us how to design an efficient cell. For this we require accurate simulation of actual device structures by computer. The earliest attempt to model a GaAs window cell was made by Hutchby and Fudurich [12]. A more complete system is that described by Demoulin and Lundstrom [13] and we outline their method and some of their results below. Using a two-dimensional model suitable for heterostructure simulation, they solve Poisson's

equation and the electron and hole diffusion equations modified to allow for the presence of pseudofields and changes of parameters like permittivity which arise from the spatial dependence of the band structure. Its two dimensional character means that it can deal with the transverse currents which flow near the surface in a self consistent fashion. The model also includes the effects of bandgap narrowing and carrier degeneracy which occur in heavily doped regions. In p-type material these effects can be allowed for by means of a single (nonphysical) bandgap shrinkage parameter. In n-type GaAs where degeneracy effects are very strong, it is assumed that it completely offsets any true bandgap narrowing, enabling one to ignore this effect. Both p-n and n-p heteroface cells with homojunction backsurface fields were studied at 500 suns (AMI .5).Their model was validated by comparing with changes of illuminated and dark currents in cells whose active regions were thinned by repeated etching. The data used for lifetimes and surface recombination velocity were chosen to represent those realised in practice. The optimised design gave efficiencies of 30.8% for p-n and 29.8% for n-p devices. In both cells the collection efficiency was 95% but the open-circuit voltage was 39 mV larger in the p-n cell. Bandgap narrowing reduced the efficiency of both cells; it reduced Voc in both cases but in addition it also reduced the short-circuit current for the n-p cell because of bandgap shrinkage in the backsurface layer. The situation was improved by the use of heteroface backsurface layers. One effect whose significance has been recently appreciated is the phenomenon of photon recycling [14]. This occurs as a result of the absorption of photons emitted in radiative recombination. If this process is ignored in modelling calculations, then the effective recombination rate is exaggerated leading to an underestimate of the open-circuit voltage and, hence, of the efficiency. Taking photon recycling into account resulted in a simulated efficiency increasing from 22.5% to 23.5%. D

HIGH PERFORMANCE ONE-SUN CELLS FOR TERRESTRIAL USE

The first method used for improving the efficiency of GaAs solar cells involved the use of heteroface windows; this produced an impressive improvement. Hovel and Woodall [15] obtained an efficiency of 22% at AMI in 1977. For a time thereafter, improvement was slow but recently there have been important gains. The current record efficiency was reported by Kurtz et al [16]. As well as a window layer, their cell also has a heteroface back surface layer. Instead of the AlGaAs conventionally used for these layers, they employed GaInP2 which is less sensitive to water and oxygen contamination. The n-p cell, on a p+ GaAs substrate, had a thickness of 3.7 \im. A double layer antireflection coating (MgF2ZZnS) was applied. In the case of their best cell, the short-circuit current density was 28.54 mA/cm2, the open-circuit voltage was 1.039 V, and the fill factor was 86.84% leading to an efficiency of 25.7%. The illumination was AMI.5 global. The authors attribute the high performance to exceedingly low recombination velocity at the GaAs/GaInP2 interface. The level of performance described above will not be realised in modules made up of many solar cells. To get some idea of what is achievable in this situation, reference can be made to the work of McClelland et al [17]. They used monolithically connected thin film AlGaAs/GaAs cells made by the CLEFT process (see Section F). Under AMI .5 global illumination they were able to reach efficiencies of 21 %.

E

GaAs CONCENTRATOR SOLAR CELLS

At an early stage in their development it was realised that there were economic advantages to be gained by the use of GaAs cells under concentrated illumination since this would reduce the amount of expensive semiconductor needed to convert a given amount of solar energy. Operation under intense illumination intensifies two problems, however. Firstly, the requirement for a low series resistance becomes much more stringent and secondly, effects of the increase of temperature which is likely to occur must be allowed for. If these are strongly adverse, then the heat dissipation must be enhanced. Neither of these has turned out to be a particularly severe problem for GaAs. In fact the temperature effect is much less pronounced than in the case of silicon. The first results obtained were rather impressive. In 1978 Sahai et al [6] described p-n window cells with an efficiency of 24.7% under 180 suns (AM2). The key to this high performance was the use of an AlGaAs layer of only 50 nm thickness leading to a much improved short circuit current. It was several years before this efficiency was significantly improved. MacMillan et al [18] have recently described high performance GaAs concentrator cells. Internally these are of relatively familiar design, AlGaAs/GaAs window cells, but the cell itself is circular with a circular grid pattern to match the solar image. Both p-n and n-p cells were fabricated. Cells of the former type reached 27.5% efficiency at around 1000 suns whilst the latter attained 28.1% at 400 suns. The improvements were primarily due to advances in cell processing technology. The use of prismatic cover glass [19] to reduce the grid line shadowing resulted in a cell with an efficiency of 29.2% at about 200 AMI.5 suns. As in the case of cells without concentration, the efficiency of the multicell module is distinctly lower. Using cells of the same type as those described above in a system with a Fresnel lens concentrator with passive cooling gave an efficiency of 22.7% under a 942 times concentration of AM 1.5 direct sunlight [20]. The temperature near the cells was 490 C. F

THEV FILM GaAs SOLAR CELLS

The possibility of making high performance GaAs thin film cells is an attractive one. The very rapid increase of optical absorption for photon energies only just greater than the energy gap provides the physical basis of this possibility which would lead to very large cost reductions. In space use the reduction in weight would also be an advantage in competition with silicon which has a lower density than GaAs. Unfortunately the poor mechanical strength of GaAs has made the realisation of these possibilities exceedingly difficult. For this reason, the deposition of GaAs on foreign substrates such as silicon has attracted widespread interest. In addition, these heteroepitaxial structures should form the basis for the fabrication of monolithic multijunction tandem solar cells of very high efficiency. The starting point of most of the thin film work was the fabrication by epitaxy on GaAs substrates of n+-p-p+ cells with very shallow junctions only 45 nm deep showing an AMI efficiency of 20% without concentration [4]. Note that because the junction was so shallow it became advantageous to use a p-type base layer. In order to reduce the amount of GaAs used, the same group developed the CLEFT process. In this a thin film GaAs solar cell is peeled off a reusable substrate [21]. Since cells only 5 jim thick have been prepared in this way, large material savings should,

in principle, be possible. GaAs is closely lattice matched to germanium but shows a 4% mismatch with silicon. The thermal expansion coefficients of GaAs and Si also differ considerably. It is therefore not surprising that good results were obtained at an early stage for GaAs/Ge cells. lies et al [22] were able to routinely produce cells on inactive Ge whose AMO efficiency exceeded 20%. The use of Si substrates has the advantage of low cost and low density but the mismatches mentioned earlier lead to high dislocation densities, antiphase domains and cracking. To avoid this it was necessary to use complex structures and processing. For example, strained layer superlattice buffers were inserted between the Si and the GaAs to reduce the dislocation density. Only partial success was achieved; using MOCVD fabrication, Kadota et al [23] reported 18.3% AMO efficiency. G

THE EFFECTS OF IRRADIATION ON GaAs SPACE CELLS

Work on thin film cells was largely motivated by space applications and it is necessary to demonstrate an advantage over silicon cells. A basis of comparison has been proposed by Castaner and Calderer [24]. Stella et al [5] have recently reviewed the present status of GaAs space cells. For a recent account of the fabrication of ultrathin GaAs space cells, see Hardingham et al [25]. One most important question concerns the effect of radiation on cell performance. Both proton and electron damage are involved and Anspaugh [26] has recently measured the damage coefficients for GaAs/Ge solar cells. The results were similar to those found earlier for LPE GaAs cells. The shielding effects of cover glasses were also measured. These results, and others for Si and InP, were used to calculate expected end-of-life performance by Summers et al [27] for different solar cells. The predictions suggested that GaAs/Ge cells would be unusable after 10 years in orbits between 1700 and 7000 nautical miles. InP/Si cells maintained useful performance in all orbits. H

STACKED CELLS INVOLVING GaAs

To break out of the constraints assumed by Shockley and Queisser [1], it is necessary to go to systems of two or more solar cells of different energy gap [28]. The total solar spectrum can then be divided so that each cell is better matched to the radiation falling on it. A number of different ways of realising this proposal have been identified. There are 'spectrum splitting' systems in which an external optical system of selectively reflecting mirrors directs different parts of the solar spectrum onto distinct solar cells of differing gaps. Alternatively, there are stacked or tandem cells where the separate cells are one behind the other. These may be either mechanically stacked or parts of a monolithic structure. The latter arrangement, however, generally imposes the requirement that the same current must flow through each cell. This can be arranged for a particular illuminating spectrum but should this change then there would be a decline in efficiency. Theoretical studies have shown how the efficiency may be expected to increase with two, three or more stages [29]. In each case there will be an optimum set of energy gaps.. For two junction structures, the larger gap in the top cell should be about 1.7 eV and the smaller about 1.0 eV. Although these figures depend on Air Mass, concentration ratio and whether the cells are

electrically in series, the changes produced by these factors are relatively small. In any case, the consequent reductions in performance are small for energy gap variations of, say, 0.1 eV particularly when the component cells are electrically independent. A number of such systems involve a GaAs cell as one of the components. An interesting example was recently described by Marti et al [30]. This uses spectrum splitting combined with a light confining cavity. GaAs and Si cells are used for the photovoltaic conversion. The illumination was direct AMI. 5 concentrated by a factor of 180. The system efficiency was 29.4% of which 22% was contributed by the GaAs cell. However the highest efficiency involving a GaAs cell was reported by Fraas [31]. This used mechanically stacked GaAs and GaSb cells. The latter is better matched to GaAs than is silicon. The illumination was 100 times concentration and the efficiency was 37% of which the GaAs provided 28.9%. In the case of monolithic structures it is necessary to provide a very low resistance, tunnel junction between the component cells. Considerable progress has been made in this direction [32]; the tunnel junction could be used at concentrations up to 500 times. However good results have been obtained with unconcentrated sunlight and monolithic structures. Bertness et al [33] have achieved 29.5% at AMI.5 global and 25.7% at AMO with a GalnP/GaAs tandem cell. They attribute this performance to careful optimisation of the grid pattern, the use of an AlInP window layer and a GaInP back surface field layer and changes in the tunnel junction fabrication. I

CONCLUSION

Despite its superior performance, the one-sun terrestrial GaAs solar cell has not established itself except in niche markets. Under concentrated sunlight, the situation is more promising though suitable conditions for its application are limited geographically. GaAs thin film cells have made considerable progress and may expect to find expanding use as space power sources. The most rapid development has occurred in stacked cell systems of all kinds and it is to be expected that this will continue in the future. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] II1] [12] [13] [14] [15] [16] [17] [18]

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21.6 GaAs optoelectronic integrated circuits and future applications S.L. Chuang August 1996

A

INTRODUCTION

As early as 1982, it was demonstrated [1] that a buried-heterostructure GaAs/AlGaAs laser diode could be fabricated monolithically on a GaAs substrate with a field-effect transistor (FET). In this rather basic arrangement, the drain current of the FET supplies the injection current into the laser diode. The field of integrated optoelectronics has emerged rapidly since it provides a new technology for applications such as optical interconnects. For short-haul communication applications such as local-area networks (LANs), it is desirable to employ GaAs-based circuits for their ability to work at high temperature. Optoelectronic devices, such as laser diodes, photodetectors, and electroabsorption modulators, have been monolithically fabricated with electronic components such as MESFETs, HEMTs, HBTs, and Si CMOS components. The physics of semiconductor optical and electronic devices can be found in [2-4]. Monolithic integration of optoelectronic integrated circuits (OEICs) has been realized using GaAs, InP and Si substrates. In this Datareview, we will focus mostly on GaAs OEICs. It should be noted that OEICs using InP substrates have also been under intensive research because of their applications in optical communication systems at 1.3 and 1.55 jum wavelengths. B

OPTOELECTRONIC INTEGRATED CIRCUITS

Bl

Integration of Laser Diodes (LDs) with Microwave Components

Laser-diode transmitter modules require LD drivers using FETs, HEMTs, HBTs, or silicon-bipolar devices. Both hybrid- and monolithic-integrated circuit LD modules have been investigated. The major issues for the design of LD transmitter modules include [5]: (i)

low threshold current and high differential gain of the laser diodes

(ii)

high breakdown voltage of the driver integrated circuit

(iii)

small waveform distortion due to impedance mismatching between the LD and the LD driver

(iv)

superimposing LD bias current on modulation current without degrading the signal waveform

(v)

easy packaging of the collimating lens systems

(vi)

good thermal stability of the LD considering the heat emitted by the LD driver.

Bl.l

Integration of LDs with MESFETs

Usually, the transistors employed in monolithic integrated circuits are MESFETs, although HBTs or HEMTs have also been used. Impedance matching plays an important role in the interconnection between the LD and the LD driver IC. A strained InGaAsP multi-quantum-well (MQW) distributed-feedback (DFB) laser diode chip and an LD driver using a GaAs-MESFET IC has been connected [5] by a microstrip line with a characteristic impedance of 25 Q. The LD is biased at 1.1 times the threshold current (I± = 10 mA) and the modulation current is 80 mA peak-peak modulation following a 10-Gbit/s nonreturn-to-zero (NRZ) pattern. The extinction ratio is 11 dB. The combination of high quantum efficiency of 0.19 W/A and low optical isolator loss of 4.2 dB achieves the high output of+4.6 dBm at the fibre pig tail with a clear eye opening. Monolithic integrated circuits of a GaAs/AlGaAs/InGaAs laser and GaAs MESFETs in a differential pair configuration have been demonstrated [6] using a layered structure grown by metalorganic chemical vapour deposition (MOCVD). The FETs were found to have a cutoff frequency of 6.3 GHz and ^ x = 8.5 GHz. The circuits exhibited a 3 dB bandwidth as high as 3.4 GHz. Monolithic integration of an AlGaAs/GaAs multi-quantum-well LD with a GaAs MESFET grown on a Si substrate using selective regrowth has also been reported [7]. The laser diode has a room-temperature threshold current of 38 mA (current density is 1.81 kA/cm2), which is small enough to be controlled by varying the gate voltage of the MESFET in a series connection. The GaAs MESFET exhibits a maximum transconductance of 88 mS/mm and a pinch-off voltage of -2.25 V for a gate length of 2.5 jim and a gate width of 400 \im. Using GaAs-on-InP heteroepitaxial technology, a 1.3 jum OEIC transmitter integrating a GaInAsP laser with two GaAs MESFETs has been fabricated on an InP substrate [8]. The laser has a low threshold current (20 mA) and a high output power (15 mW at 100 mA), while the MESFETs have a transconductance of 60 mS/mm for a 1 |nm long, 250 |um wide gate. The transmitter has an 8 Gbit/s NRZ modulation bit rate. Monolithic integration of GaAs MESFETs with vertical-cavity surface-emitting lasers (VCSELs) has also been realized [9]. The entire epitaxial layer structure for both GaAs VCSELs and MESFETs was grown at the same time by MBE. The VCSEL structure was grown on an n-type GaAs substrate, consisting of 27 and a half pairs of n-Al As/Al01Ga09 As layers as a distributed Bragg reflector (DBR) for the bottom mirror, an active region with a GaAs single quantum well sandwiched between p- and n-Al0J5Ga065As cladding layers, an Al01Ga09As current conduction layer above the active region, and 21 and a half pairs of undoped AlAsZAl01Ga09As DBR as the top mirror. Then a 1500 A thick n-GaAs channel for the MESFETs was grown on the top of the undoped top mirror, which serves as a highly resistive layer to isolate the VCSELs from the MESFETs and reduce leakage in the MESFETs. The VCSELs with a 10 jum diameter in the active region have an average threshold current of 6 mA and a CW power of 1.1 mW. The MESFETs with a 3 |im gate length have a transconductance of 50 mS/mm. The laser output is modulated by the gate voltage of the MESFETs and exhibits an optical-to-electrical conversion factorof0.5mW/V. The VCSEL has also been monolithically integrated with MESFETs and MSM photodiodes for switching devices to perform NOR- and OR-types of operations with optical gain [10].

B1.2

Integration of LDs with HEMTs

An In015Ga085As MODFET (or HEMT) and a graded-index separate-confinement heterostructure (GRINSCH) single quantum-well laser have been vertically integrated [11] by molecular beam epitaxy. The MODFET survived the long high temperature growth of the laser without noticeable degradation of the operation characteristics. The MODFETs with 2 ^m x 500 \im gates had a maximum source-to-drain current of 250 mA/mm, a typical unity current-gain frequency ft of 5 GHz, and a unity power-gain frequency ^ x of 4 GHz. The integrated LD and MODFET devices had a maximum microwave modulation 3-dB bandwidth of 3.5 GHz. B1.3

Integration of LDs with HBTs

Monolithic integration of a GaAs/AlGaAs HBT with a GaAs vertical-cavity surface-emitting laser (VCSEL) for an optoelectronic switch has also been demonstrated [12]. The VCSEL structure was grown by MBE on an n-type GaAs substrate. The HBT structure was epitaxially regrown on the underlying VCSEL and isolation layers by MBE5 and it contained a graded heterojunction interface with a thicker and lower-doped base layer to relax the processing tolerances. High-performance switching has been achieved, including high current gain (700), low switching current (7 ^A at threshold), and high electrical-to-optical power conversion efficiency (150 W/A) with a DC power dissipation of 27 and 76 mW at an output power of 0.4 and 1.4 mW, respectively. An AlGaAs/GaAs Npn HBT laser driver and a pseudomorphic InGaAs/GaAs/AlGaAs GRINSCH single quantum-well laser have been laterally integrated using selective organometallic vapour phase epitaxy (OMVPE) regrowth [13]. A quasi-self-aligned GalnP/GaAs HBT laser driver circuit has been designed [14] to drive a DFB laser with a 14 Gbit/s digital optical transmission rate. The laser diode is connected directly to the driver IC to avoid waveform distortion. B2

Integration of Photodetectors with Microwave Components

Metal-semiconductor-metal (MSM) photodiodes and avalanche photodiodes (APDs) integrated with microwave transistors for photoreceivers have been of great interest. The bandwidth and the noise are the key design issues. For the APD receiver, the issues include, (i) the broad APD bandwidth which is limited by carrier transit time in the low multiplication gain region, (ii) the gain-bandwidth product in the high gain region, and (iii) the large transimpedance bandwidth of the preamplifier [5]. The noise issues include the low excess multiplication noise and low dark current of the APD, as well as the low input equivalent noise and high transimpedance gain of the preamplifier. Superlattice APDs have been used for their large gain-bandwidth product and low excess multiplication noise. InGaAs/InAlAs superlattice APDs suffer from large dark currents in the high gain region. InGaAsP/InALAs superlattice APDs have both low dark current and high gain-bandwidth products of over 110 GHz. AlO Gbit/s-APD optical receiver using an InGaAsP/InAlAs superlattice APD and a GaAs-MESFET IC preamplifier with flip-chip bonding has a measured sensitivity better than -27 dBm.

B2.1

Integration of Photodetectors with MESFETs

Monolithic integration of a pin photodiode and a MESFET for photoreceivers was realized using a single MBE growth step [15]. The structure was grown on a semi-insulating GaAs substrate and consisted of a 1000 A layer of unintentionally doped GaAs, 1000 A of Al0 4Ga06As, an n-type channel layer (2300 A GaAs) with an electron concentration of 4 x 1017 cm"3, and a 3000 A GaAs n-contact layer doped to 2 x 1018 cm"3. This contact layer served as the ohmic contact for both the MESFET and the pin diode. These layers were then followed by the pin structure, which consisted of 2000 A OfAl04Ga06As (doped n-type to 2 x 1018 cm"3), a 6000 A undoped GaAs absorption layer, 2000 A OfP-Al04Ga06As (doped to 2 x 1018 cm"3), and the GaAs p-contact layer doped to 5 x 1018 cm"3. The pin layers were then selectively removed by wet chemical etching to define the MESFET areas, and isolation mesas were formed around the pins and MESFETs by etching down to the semi-insulating substrate. Photolithographic techniques were used to define the ohmic contacts, gates, Ti resistors, and interconnect metal. The photoreceiver exhibited flat band gains as high as 17 dB and a bandwidth of 2 GHz. Monolithic integration of ion-implanted GaAs MESFETs and metal-semiconductor-metal (MSM) photodetectors has also been demonstrated with a bandwidth exceeding 10 GHz for an OEIC photoreceiver [16]. Cascadeable smart pixels using monolithic integrated circuits of AlGaAs/GaAs LEDs, MESFETs, and photodiodes have also been demonstrated [17]. A 1.3 ^m Fabry-Perot LD transmitter module with a GaAs MESFET LD driver and pin photodiode receiver has been developed for use at 5 Gb/s operation [18]. B2.2

Integration of Photodetectors with HEMTs

A 0.85 |im wavelength photoreceiver using a 0.25 |im InGaAs/GaAs MODFET and GaAs MSM photodiode was demonstrated with an overall bandwidth of 11 GHz [19]. Another 0.85 |im wavelength photoreceiver using a multistage AlGaAs/GaAs HEMT amplifier has a bandwidth of 13 GHz, a sensitivity that is better than -14.7 dBm at 12.5 Gbit/s, and a bit error rate BER = 10"9 [20]. Monolithic integration of an InGaAs MSM photodiode and AlGaAs/GaAs HEMTs grown on a GaAs substrate for 1.3 nm-wavelength operation has also been achieved with a bandwidth of 7.1 GHz [21,22]. B2.3

Integration of Photodetectors with HBTs

A 2.5 Gbit/s four-channel In053Ga047As pin photodetector and an AlGaAs/GaAs HBT photoreceiver array has been demonstrated with hybrid integration. It has a uniform responsivity of 1.6 V/mW per channel at a wavelength of 1.55 \im. The sensitivity was -26.5 dBm (measured for a BER of 10"9 and a wavelength of 0.98 \xm) at 1 Gbit/s and the interchannel crosstalk is below -40 dB [23]. Design considerations using a comprehensive circuit and device model can be found in [24]. B3

Integration of Modulators with MESFETs9 HBTs and CMOS

GaAs quantum-well electroabsorption modulators are based on the quantum-confined Stark effect. They are low-energy, electrically-driven devices, compatible with electronics in materials and voltages and with laser diodes in wavelength and power levels. Combining photodetectors, modulators, and other circuitry can make a self-electro-optic effect device (SEED). A field-effect

transistor SEED is realized by integrating FETs, photodiodes, and modulators [25,26]. Monolithic integration of FET-SEED for both optical input and output is a desirable function for making a so-called 'smart pixel' [27]. Optical receivers are one of the major components of smart pixels, the other being the transmitter. Several FET-SEED optical receiver arrays have been tested. They displayed a mean response of 0.7 mV/^W and were capable of larger than 100 Mb/s per channel operation [28]. A sensitivity of-25 dBm for a 40 Mb/s rate was measured for a BER of less than 10~9 These two-dimensional arrays of GaAs FET-SEED differential transimpedance receivers have applications in massively parallel optical data link, board-to-board interconnections. Integration of a GaAs/AJGaAs surface-normal modulator with HBTs or with silicon CMOS components has also been demonstrated [29-31]. The hybrid GaAs modulators bonded to a 0.8 jim silicon CMOS chip function as smart-pixel optical receivers at 1 Gbit/s operation with a BER ofl x IO"9. C

POTENTIAL FUTURE SYSTEM APPLICATIONS

Photonic techniques for steering a microwave phased-array radar have been of great interest recently. A guided lightwave using optical waveguides or fibres is used as a carrier for distributing and delaying the microwave signals that drive and control the phase of the antenna radiating elements [32]. GaAs monolithic microwave integrated circuits (MMICs) are used as array output modules in radiating phased array antennas for space communication applications. An optical beam-forming network using the interconnect of MMIC modules with fibre-optical links appears attractive for future radar systems with optical fibres being used to carry the control signals to the variable phase shifters and amplifiers. A high-speed GaAs MESFET integrated-circuit optical receiver/demultiplexer has been applied [33] to the optical control of a Ka-band GaAs monolithic phase shifter. Monolithic integration of GaAs rib waveguides has been achieved and InGaAs waveguide coupled MSM photodetectors have been fabricated for an optical time-delay network [34] used in the optical control of phased arrays. Other applications of optical and microwave integrated circuits include a fibre optical link for cable television, antenna measurement instrumentation, and signal processing in radar systems [32]. There is also a growing demand for high-speed medium-span optical links in applications including inter-LAN communications, gigabit per second computer network communications, and trunk line systems of telecommunication networks [5]. D

CONCLUSION

GaAs OEICs have been reviewed. The monolithic integration of optical components such as laser diodes, photodetectors and modulators with electronic components such as MESFETs, HEMTs, and HBTs, provides a new technology compatible with MMICs. Future applications such as photonics and microwave phase array radar and optical interconnects have been briefly discussed. REFERENCES [1] [2] [3] [4]

I. Ury, K.Y. Lau, N. Bar-Chaim, A. Yariv [AppL Phys. Lett (USA) vol.41 (1982) p. 126-8 ] S. Sze [ Physics of Semiconductor Devices (Wiley-Interscience, 1981) ] A. Yariv [ Optical Electronics (Holt, Rinehart, and Winston, 1985) ] SX. Chuang [ Physics of Optoelectronic Devices (Wiley-Interscience, 1995) ]

[5] [6] [7] [8] [ 9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33] [34]

Y. Miyamoto et al [ J. Lightwave Technol. (USA) vol. 12 (1996) p.332-42 ] D.T. Nichols et al [ Electron. Lett. (UK) vol.30 (1994) p.490-1 ] T. Egawa, T. Jimbo, M. Umeno [ IEEE Photonics. Technol. Lett. (USA) vol.4 (1992) p.612-4 ] O. Calliger et al [ IEEProc. Optoelectron. (UK) vol. 142 (1995) p. 13-6 ] YJ. Yang et al [ Appl. Phys. Lett. (USA) vol.62 (1993) p.600-2 ] S. Matsuo et al [ IEEEPhotonicsTechnol. Lett. (USA) vol. 1 (1995) p. 1165-7 ] S.D. Offsey et al [ Electron. Lett. (UK) vol.26 (1990) p.350-2 ] P. Zhou et al [ IEEE Photonics Technol. Lett. (USA) vol.5 (1993) p. 1035-8 ] D.B. Slater Jr. et al [IEEEPhotonics Technol. Lett. (USA) vol.5 (1993) p.791-4 ] M. Menouni et al [ Electron. Lett. (UK) vol.32 (1996) p.231-3 ] D. Nichols et al [ Electron. Lett. (UK) vol.29 (1993) p. 1133-4 ] J.S. Wang et al [ IEEE Photonics Technol. Lett. (USA) vol.5 (1993) p.316-8 ] U. Kehrli et al [ Electron. Lett. (UK) vol.30 (1994) p.2069-70 ] Y. Hakamata, K. Takemoto, T. Nakanishi, J. Nakano [ IEEE Photonics Technol. Lett. (USA) vol.5 (1993) p.251-3] A. Ketterson et al [ IEEE Trans. Electron Devices (USA) vol.40 (1993) p. 1406-16 ] V. Hurm et al [ Electron. Lett. (UK) vol.32 (1996) p.683-5 ] V. Hurm et al [ Electron. Lett. (UK) vol.31 (1995) p.67-8 ] V. Hurm et al [ Electron. Lett. (UK) vol.32 (1996) p.391-2 ] M. Govindarajan, S. Siala, R.N. Nottenburg [ IEEE Photonics Technol. Lett. (USA) vol.5 (1993) p.1397-1400] M. Govindarajan, S. Forrest [ J. Lightwave Technol. (USA) vol. 11 (1993) p.367-78 ] D.A.B. Miller et al [ IEEE Photonics Technol. Lett. (USA) vol. 1 (1989) p.62-4 ] T.K. Woodward et al [ IEEE Photonics Technol. Lett. (USA) vol.4 (1992) pp.614-7 ] T.K. Woodward, A.L. Lentine, L.M.F. Chirovsky [ IEEE J. Quantum Electron. (USA) vol.30 (1994) p.2319-24] R. Novotny et al [ J. Lightwave Technol. (USA) vol. 13 (1995) p.606-14 ] K. W. Goossen, J.E. Cunningham, W. Y. Jan [ IEEE Photonics Technol. Lett. (USA) vol.4 (1992) p.393-5] K. W. Goossen et al [ IEEE Photonics Technol. Lett. (USA) vol.7 (1995) p.360-2 ] T.K Woodward et al [ IEEE Photonics Technol. Lett. (USA) vol.8 (1996) p.422-4 ] R. Simons [ Optical Control of Microwave Devices (Artech House 1990) ] K.B. Bhasin et al [ IEEE Trans. Microw. Theory Techn. (USA) vol.38 (1990) p.686-8 ] W. Ng, D. Yap, A. Narayanan, A. Walston [ IEEE Photonics Technol. Lett. (USA) vol.6 (1994) p.231-4]

CHAPTER 22 GaAs IN OTHER APPLICATIONS 22.1 22.2 22.3

Sensors GaAs nuclear particle detectors High temperature applications for GaAs

22.1 Sensors Y. Sugiyama August 1996

A

INTRODUCTION

This Datareview considers various sensors made from GaAs and related materials. Sensors based on GaAs are based on the transfer of a physical quantity from one energy into another using the optical, electrical, magnetic, thermal, mechanical or chemical properties of the crystal. Optical sensors, magnetic sensors, pressure sensors, mass flow sensors, piezoelectric sensors, power sensors, gas sensors and biosensors are summarized in this Datareview. B

OPTICAL SENSORS

GaAs-based optical sensors such as photodiodes are widely used in various applications such as photosensors, displacement sensors, pressure sensors, and so on. A new material technology for optical sensors has recently been presented that has exploited low dimensional heterostructures [I]. This has resulted in an infrared detector using inter-subband absorption in AlGaAs/GaAs multi-quantum-wells (MQWs) [2]. In general, a GaAs photodetector is composed of a metal-semiconductor-metal (MSM) Schottkybarrier structure while metal-insulator- semiconductor (MIS) structures are typically used for Si [3]. An interdigitated MSM Schottky-barrier photodetector with direct coupled FET logic (DCFL) has been demonstrated for use in multiplexers and demultiplexers in optical communications over the frequency range 0.9 -1.3 GHz. This technology offers low power and high frequency operation [4]. The photosensor shows a photoresponsivity of 1.2 AAV and a dark current of 4 nA at a wavelength of 820 nm and at a bias voltage of 2 V. A theoretical bandwidth of 10 GHz has been estimated. An optical displacement sensor with a Michelson interferometer composed of a DFB laser, a photodetector and transparent waveguides integrated on a GaAs substrate has also been developed [5]. Interference fringes were seen at a distance of up to 45 cm with a spatial resolution of approximately 100 nm. Fibre-optic interferometric systems with a high resolution using low-coherence light sources have been reviewed in [6]. C

MAGNETIC SENSORS

Cl

Hall Devices

Semiconductor magnetic sensors can work on the basis of the magnetoresistance effect or the Hall effect [7]. Magnetic sensors such as Hall devices and magnetoresistors totalling a billion devices per year are widely used for magnetic field probes, displacement sensors in brushless motors of videotape recorders and disk players, bill sensors of vending machines, potentiometers, and contactless switches in factory and office automation and in domestic electronics [8]. Most of these Hall devices are made of III-V compound semiconductors with high electron mobility such as InSb, InAs, and GaAs.

The output of a Hall device, that is the Hall voltage, VH, is expressed as follows: VH = KHI/B

(1)

where KH is called the product sensitivity which is expressed as the Hall coefficient divided by the thickness of the device, I is the drive current in amps, and B is the magnetic flux density in tesla. The product sensitivity, KH, the magnetic sensitivity per unit current, is expressed as 'V/AT'. On the other hand, we can use another expression for the magnetic sensitivity, the magnetic flux density sensitivity, KB = VH/B = KHI. KB is expressed as 'V/T' and is used to describe the sensitivity of the device independent of its size. GaAs is a better material than Si for a Hall sensor because of its high sensitivity. This is due to the high electron mobility, good thermal stability due to its large energy gap, and excellent electrical isolation resulting from the use of a semi-insulating substrate. GaAs Hall sensors fabricated by Si+ ion implantation into semi-insulating GaAs substrates are most popular for applications requiring good thermal stability [9]. It has recently been possible to obtain high quality material with high electron mobility using new technologies such as molecular beam epitaxy (MBE) and metal organic chemical vapour deposition (MOCVD). GaAs Hall sensors made by MOCVD have been shown to have a sensitivity of 2.2 V/T with an offset of 3.4 mT. Constant sensitivity can be obtained in the wide temperature range from -1900C to 20 0 C [1O]. Quantum-well Hall sensors, taking advantage of the properties of quantum-well heterostructures with a two-dimensional electron gas, have been reported. High product sensitivities of 1000 V/AT but with high impedance and a poor temperature coefficient of the sensitivity of -0.7%/° C (due to the presence of the deep DX trap) were achieved for AlGaAs/GaAs single heterostructure Hall sensors [H]. This drawback was overcome in AJAs/GaAs superlattice Hall sensors with a Hall mobility of 5000 cm2/ Vs [12]. A 200 ^m Greek-cross Hall sensor with high sheet resistance of 2.2 kQ showed a high product sensitivity of 1200 V/AT with a moderate temperature coefficient of -0.1%/°C. Higher electron mobility can be achieved in the ternary alloys of InGaAs grown on either GaAs or InP substrates. Hall sensors have been made of pseudomorphic In0 52A1O^8AsZIn0^Ga0 2As heterostructures grown on InP substrates by MBE [13]. They have excellent mobilities of 16000 cm2/ Vs and 160000 cm2/ Vs at 300 K and 10 K, respectively. They exhibit a product sensitivity of 320 V/AT, a maximum magnetic flux density sensitivity of 15 V/T, a temperature coefficient of -0.035 %/°C, and a signal-to-noise ratio corresponding to a resolution of 60 nT/Hz1/2 at 1 kHz with a Hooge's noise parameter of 9 x 10'4. The temperature dependence of the sensitivity in the range from low temperatures to room temperature was measured [14] and maximum magnetic flux density sensitivities of 33.5 and 22.5 V/T at 11 and 296 K, respectively, were obtained. In practice, MOCVD may be available for factory production with the result that GaAs-based devices will offer good cost performance. Good device properties were achieved in In0 52A1O^8AsAn0^3Ga0 47As heterostructures made by MOCVD lattice-matched to InP substrates [15] and in pseudomorphic AlGaAs/InGaAs structures grown on GaAs substrates by MBE. They exhibited a low thermal drift of-160 ppm/°C and a product sensitivity of 900 V/AT [16].

C2

Micro-Hall Devices

Semiconductor Hall devices with high spatial resolution are useful for various applications. Conventional Hall devices normally have an active region of some hundred microns square and one micron thickness. However, Greek-cross micro-Hall devices made from InSb bulk material and GaAs epitaxial films have been reported [17]. The spatial resolution of the device is defined by the central square area of a cross-shaped Hall sensor. A GaAs Hall sensor with a spatial resolution of 4 |im square has been demonstrated with a product sensitivity of 500 V/AT and an impedance of 5 kQ. This should be compared with an InSb Hall sensor with a 4 |im square resolution, a product sensitivity of 140 V/AT, an impedance of 130 Q, and a signal-to noise ratio of 105 dB at 0. IT. A quarter-micron thick GaAs Hall device has been fabricated by focused ion implantation (FIB) [18]. The implant dose into the central part of the active region was only 5 x 108 ions/cm2. The electrical properties of a typical Greek-cross shaped specimen with 0.3 ^m width are a product sensitivity of 85 V/AT, a maximum magnetic flux density sensitivity of 48 mV/T, an impedance of 14 kQ, and a Hall mobility of 3700 cm2/ Vs. The maximum sensitivity is limited by the saturation of the electron velocity. A Greek-cross Hall device made of a 5-doped InAlAs/InGaAs pseudomorphic quantum-well heterostructure with a spatial resolution of 0.4 |im square has been reported [19]. The sheet resistance was 170 Q/square and the product sensitivity was measured to be 230 V/AT. The minimum detectable magnetic field is 7 mT at an operating frequency of 1 Hz and 230 mT at 1 kHz. The device works in the wide temperature range from 10 K to more than 300 K, the temperature coefficient being -0.08 %/K at room temperature. C3

Magnetotransistors

Semiconductor magnetic sensors are normally based on the carrier deflection that results from the Lorentz force. Most of them can be realized with standard IC processes. In order to take advantage of higher electron mobility, magnetotransistors using the possibilities of GaAs and related materials have been investigated, and they have been shown to exhibit very good sensitivities. Split-contact magnetotransistors made of AlAs/GaAs superlattice structures [20, 21] with a Hall mobility of 6400 cm2/ Vs and a resistivity of 1.4 mQ have shown high sensitivities of 46 %/T in the drive current. This is three times as large as that of Si-CMOS devices. Multidrain GaAs magnetotransistors can reach high sensitivities of 70%/T as well as having low noise [22]. Noise and resolution of magnetotransistors such as GaAs and Si have been compared in [23,24] where a resolution of 10 - 40 |iT/Hz1/2 for GaAs devices was obtained. D

PRESSURE SENSORS

For the measurement of uniform pressure, the pressure must be converted to a uniaxial stress which can then be coupled to a crystal with the aid of a mechanical system, such as a diaphragm. However, the inertia of the mechanical system increases the response time of the device far beyond that of the semiconductor crystal alone. Semiconductor alloy crystals such as AlGaAs [25, 26] and GaAsP [27] are used for pressure sensors. In these crystals uniaxial stress distorts the conduction band so that an n-type Te-doped

AlxGa^xAs pressure sensor produced by growing graded-composition alloys on a semi-insulating GaAs substrate exhibits a high sensitivity of the resistivity to stress. One part of the crystal will be low resistance (where carriers occupy the F-region) and the other will be of high resistance (the X-region). The total resistance of the crystal in the direction perpendicular to the composition gradient will be the sum of the resistances of the T and X regions. The sensor has a high and linear sensitivity of 5 -10 %/100 MPa with a small temperature coefficient of 0.02 %/K. It also exhibits a fast response time of less than 1 \is, and works up to 1 GPa in the temperature range 120 - 400 K. The housing for the pressure sensor which has high natural frequency response up to 300 kHz is useful as protection against aggressive media [26]. The energy shifts to higher energies of the narrow excitonic line in photoluminescence (PL) under hydrostatic pressures represent optical properties that can be used for high pressure sensors. Optical pressure sensors with semiconductor multi-quantum-wells (MQWs) such as AlGaAs/GaAs [28], InGaAs/GaAs [29,30] and InGaAs/InAJAs [31] have been reported. Quantum-well optical spectra can be used for the calibration of very high pressures up to 10 GPa with good accuracy (0.1%). Pressure sensors with quantum-wells of AlGaAs/GaAs and InGaAs/GaAs are useful up to 4 GPa and 5 GPa, respectively. Other materials like InGaAsAnAlAs on InP should extend the high pressure range to 10 GPa. The use of an AlGaAs/GaAs pressure sensor with 10 multi-quantum-wells, a well width of 30 nm, a barrier width of 10 nm, and Al content of 0.3 has been demonstrated up to 40 kbar with an accuracy of 0.02 kbar. The sensor works in a wide temperature range from cryogenic temperatures to room temperature [28]. Free-standing microstructures fabricated by micro-machining can be used for sensors and actuators. A free-standing GaAs cantilever of 600 nm thickness, 10 |im width and 50 |nm length for use as a single chip capacitive pressure sensor has been made [32]. Another device with chevron-shaped free-standing GaAs wires about 67 mm in length terminating on mesa structures with Au pads for use as a gas-pressure sensor [33] has been demonstrated. The operation of this pressure sensor is as follows: the chevron-shaped free-standing wires are heated to a high temperature (2500C) by the passage of a current whilst consuming little power (25 mW) due to their thermal isolation from the substrate. The resistance depends on their temperature, which in turn depends on the variation of heat loss with the pressure of the surrounding gas. Thus they act as gas pressure microsensors, FIGURE 1. E

MASS FLOW SENSORS

An integrated mass-flow sensor fabricated on GaAs is based on the principle of convective heat transfer from a heater kept at a constant temperature [34]. Two cascaded thermopiles, a thermally-isolated heater and a differential MESFET amplifier have been integrated on a GaAs wafer with an AlGaAs membrane using micromachining technologies. The features of such a GaAs mass-flow sensor are (i) high temperature operation, because of the large band gap of GaAs, (ii) a large thermal resistivity of (2.27 + 28.83x - 3Ox2) K cm /W for Al1^GaxAs, and (iii) the possibility of creating a microstructure by micromachining. The sensor works up to an air velocity of 20 m/s and up to a temperature of 4000C. Unfortunately, the sensitivity decreases for temperatures over 240 0 C.

Change in current (jiA)

Pressure (mbar) FIGURE 1. Dependence of the change in the current through the chevron wires at different pressures for applied potentials of 12 (•), 13 (o) and 14 V (*). (Taken from [33]).

F

PIEZOELECTRIC SENSORS

GaAs is a non-centrosymmetric crystal with non-zero components of the piezoelectric tensors. The piezoelectric effect in GaAs has been exploited by fabricating the GaAs into a cantilever [35]. A typical piezoelectric pressure sensor consists of a GaAs MESFET in which the transverse piezoelectric effect is used to change the potential on the gate electrode [37]. The MESFET itself acts as an integrated charge amplifier. The sensitivity is 200 mV/bar in the range 0 - 2 0 bar. Piezoelectrically excited flexural vibrations in GaAs can be used for resonant temperature sensors with a temperature coefficient of-52 ppm/K [39]. Such a sensor composed of micromachined tuning forks of 4.25 mm length, 0.85 mm width and 0.25 mm thickness in semi-insulating GaAs has a very high Q-value of 27000 at 33 kHz at atmospheric pressure (FIGURE 2).

FIGURE 2. Micromachined tuning forks in semi-insulating GaAs, with electrodes. Dimensions: 4250 x 850 * 250 ^m3 (L x w x t). Three bond pads are located at the bottom part of the base (taken from [39]).

There are applications for piezoelectric sensors in the fields of pressure [36,37], stress [38] and temperature measurement [39]. In addition, piezoelectrically activated flexural, torsional and longitudinal, vibrating GaAs structures have good potential in other applications in sensors. The piezoelectric semiconductor GaAs is an interesting material for acoustic sensor applications. It offers the possibility of monolithic integration of fast active electric components as well as highly sensitive surface acoustic wave (SAW) modules on the same substrate. SAW sensors on GaAs offer the possibility of a temperature-stabilized mass-sensitivity of-0.55 ppm assuming a mass loading of 1 ng/cm2 [40]. G

POWER SENSORS

Power sensors can be used for power monitoring, gain control, or protection in electronic circuits such as power converters, inverter power sources, and microwave power devices. A monolithic GaAs MESFET power sensor microsystem with a three-terminal thermoconverter (TTTC) is based on a thermally isolated heater with dynamically controlled impedance [41]. The circuit diagram is composed of two thermally-symmetrical FETs for power-control and a diode as a temperature sensor. This enables the creation of a self-balancing thermal bridge which eliminates sensitivity to fluctuations in the ambient temperature. A large thermal resistance for GaAs is expected because the thermal conductivity of GaAs is three times lower than that of Si. The sensor has a three-dimensional configuration with two independent very thin GaAs cantilever beams. The GaAs power sensor has a thermal resistance of 5200 KAV and a working point of 100 K above ambient temperature, and can be excited by <20 mW of dissipated power. An ambient integrated microwave power sensor with a 50 ohm coplanar wave guide (CPW) to work at microwave frequencies has been fabricated. It is based on the conversion of electrical power into heat resulting in a local temperature increase which is measured [42]. The central conductor of a 50 ohm CPW is evaporated onto a 1 |nm thin AlGaAs membrane of high thermal resistance while the earth conductors cross the membrane in a bridge configuration with an air gap of 3 mm to prevent undesirable heat loss. The temperature difference between the central conductor and the rim of the chip is detected by 24 series connected AlGaAs/Cr-Au thermocouples. This structure has a sensitivity of 1.6 mV/K. Such an integrated MMIC power sensor has been shown to have an RF power sensitivity of 1.1 VAV with a flat responsivity up to 15 GHz (FIGURE 3).

thermopile

thermopile

spray etching membrane pre-etching FIGURE 3. Central conductor has contact to membrane so that its power-loss induced temperature increase can be detected by thermopile underneath mass conductor (taken from [42]).

H

GAS SENSORS

Catalytic metals can be used in hydrogen and ammonia gas sensors. In a catalytic metal-based Schottky diode, the barrier height decreases when the device is exposed to a hydrogen-containing atmosphere. The reduction of the Schottky barrier can be explained by the atomic hydrogeninduced dipole model [43,44]. Effects of different catalytic metals such as platinum, palladium and iridium on the sensitivity of GaAs Schottky diode gas sensors have been reported [45]. Pt and Pd metal Schottky diodes formed on GaAs are sensitive to H2 gas concentrations ranging from 6 to 100 ppm. Pd devices are appropriate for detecting H2 gas at room temperature. Porous films of Pt and Ir are sensitive to NH3 gas at concentrations up to 10 ppm [46]. A single chip hydrogen detector monolithically integrated on GaAs using a Pd Schottky diode gas sensor, a resistance heater, temperature detector and MESFET has been reported [47]. The sensitivity of this device is 26 mV/ppm at 520 K for a H2 concentration ranging from 0.1-50 ppm. I

BIO SENSORS

Techniques favoured at present usually involve measuring changes in absorption or refractive index of a sensitive element produced by the gas molecules. However, a sensor using evanescent wave interaction (interferometric, plasmon resonance, absorption, etc.) has been proposed [48]. An optical biosensor with monolithically integrated infrared (0.86 - 1.1 |nm) LED, waveguide and photodetector with AlGaAs/GaAs heterostructures on a GaAs substrate for use as a near-infrared biosensor has been developed [49]. A waveguide with a sensing area of 750 mm length and 200 \im width serves as an intermediary between an LED and a photodetector. A GaAs-based MachZehnder interferometer with a sensing area which consists of an SiO2 buffer, Ta2O5 waveguide and sensitive membrane has been designed. It is expected to be useful for chemical, biological, immunological or environmental sensors [50] (FIGURE 4).

Optical Mode In

Cladding

Sensing Membrane

Cladding

Cladding

Optical Mode Out

Cladding

GaAs Substrate

FIGURE 4. Schematic cross section of the sensor pad, filling a hole etched into the semiconductor waveguide. Not to scale. (Taken from [50]).

New materials and transducers for chemical sensors are widely reviewed elsewhere [51]. REFERENCES [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33] [34] [35] [36]

B. Hamilton [ Sens. Actuators A (Switzerland) vol.47-51 (1992) p.47-51 ] J. S. Smith et al [ J Vac. Sci. Technol B (USA) vol. 1 (1983) p.376 ] A. Y. VuI, A.T. Dideikin [ Sens. Actuators A (Switzerland) vol.39 (1993) p.7-18 ] H. Pohjonen, M. Andersson [ Sens. Actuators A (Switzerland) vol.21-23 (1990) p. 1124-7 ] D. Hofstetter, H.P. Zappe, R. Dandliker [ Electron. Lett. (UK) vol.31 no. 24 (1995) p.2121-2 ] Y.N. Ning et al [ Sens. Actuators A (Switzerland) vol.30 (1992) p.181-92 ] H.P. Baltes, R.S. Popovic [ Proc. IEEE (USA) vol.74 no. 8 (1986) p. 1107-32 ] Y. Sugiyama [ J. Vac. ScL Technol. B (USA) vol. 13 no. 3 (1995) p. 1075-83 ] H. Tanoue, T. Tsurushima, S. Kataoka [ IEEE Trans. Electron Devices (USA) vol.27 no. 6 (1980) p. 1188-92] R. Campesato, C. Flores, A. Passaseo, S. Verni [ Sens. Actuators A (Switzerland) vol.32 (1992) p.651-5] Y. Sugiyama, T. Taguchi, M. Tacano [ Proc. 6th Sensor Symp.(Japan) (1986) p.55-8 ] Y. Sugiyama, H. Soga, M. Tacano [J. Cryst Growth (Netherlands) vol.95 (1989) p.394-7 ] Y. Sugiyama, Y. Takeuchi, M. Tacano [ Sens. Actuators A (Switzerland) vol.34 (1992) p. 131 ] Y. Sugiyama [ Sens. Actuators A (Switzerland) vol.40 (1994) p. 135-40 ] R. Kyburz, J. Schmid, R. S. Popovic, M. Melchior [ IEEE Trans. Electron Devices (USA) vol.41 no. 3 (1994) p.315-20] V. Mosser et al [ Sens. Actuators A (Switzerland) vol.43 (1994) p. 135-40 ] Y. Sugiyama, S. Kataoka [ Sens. Actuators (Switzerland) vol.8 (1985) p.29-38 ] T. Kanayama, H. Hiroshima, M. Komuro [J. Vac. Sci. Technol. B (USA) vol.6 no. 3 (1988) p. 1010 ] Y. Sugiyama [ Digest. 8th Int. Conf. on Solid-State Sensors and Actuators, Stockholm, Sweden, 25-29 June 1995 vol.2 ses.A7-D13 p.225-8 ] Y. Sugiyama, H. Soga, M. Tacano, H. Baltes [ IEEE Trans. Electron Devices (USA) vol.36 no.9 (1989) p. 1639-43] A. Nathan, H.P. Baltes, R. Castagnetti, Y. Sugiyama, D. R. Briglio [ Sens. Actuators A (Switzerland) vol.21-23 (1990) p.776-9 ] N. Mathieu, A. Chovet, R. Frauquembergue, P. Descherdeer, A. Leroy [ Sens. Actuators A (Switzerland) vol.25-27 (19891) p.741-5 ] N. Mathieu, A. Chovet, R. Frauquembergue, P.Descherdeer, A. Leroy, P.Giordano [ Sens. Actuators A (Switzerland) vol.33 (1992) p.57-61 ] A. Chovet, N. Mathieu [ Sens. Actuators A (Switzerland) vol.32 (1992) p.682-7 ] S. Zilionis, V. Stankevic [ Sens. Actuators A (Switzerland) vol.25-27 (1991) p.295-9 ] S. Zilionis, K. Pyragas, G. Tautvaisaa [ Sens. Actuators A (Switzerland) vol.32 (1992) p.622-7 ] V.A. Gridchin, Y. A. Pucklyakov, L.B. Shurman [ Sens. Actuators A (Switzerland) vol.30 (1992) p. 139-42] W. Trzeciakowski et al [ Sens. Actuators A (Switzerland) vol.32 (1992) p.632-8 ] W. Trzeciakowski [ Sens. Actuators A (Switzerland) vol.41 -42 (1994) p.247-50 ] T.P. Sosin, P. Perin, W. Trzeciakowski, R. Tober, R. Zarecka [ Sens. Actuators A (Switzerland) vol.41-42(1994)p.654-7] A.R. Goni, K. Syassen, Y. Zhang, K. Ploog, A. Cantarero, A. Cros [ Phys. Rev. B (USA) vol.45 (1992)p.6809-18] J. Miao, H. L. Hartnagel, D. Ruch, K. Fricke [ Sens. Actuators A (Switzerland) vol.46-47 (1995) p.30-4 ] P.C. Hpyle, J. R. A. Cleaver, H. Ahmed [ Sens. Actuators A (Switzerland) vol.50 (1995) p.31-7 ] K. Flicke [ Sens. Actuators A (Switzerland) vol.45 (1994) p.91-4 ] Q. Huang, Q. Tong, S. Lu [ Sens. Actuators A (Switzerland) vol.35 (1993) p.247-54 ] K. Flicke, H. Hartnagel [ Electron. Lett. (UK) vol.26 no. 11 (1990) p.693-4 ]

[37] [38] [39] [40] [41] [42] [43] [44] [45] [46] [47] [48] [49] [50] [51]

G. Schweeger, C. Lang, K. Flicke, H. L. Hartnagel, R. Dolt, G. Hohenberg [ Electron. Lett. (UK) vol.30 no. 16 (1994) p. 1355-6] K. Flicke [ J. Appl. Phys. (USA) vol.70 no.2 (1991) p.914-8 ] J. Soderkvist, K. Hjort [ Sens. Actuators A (Switzerland) vol.39 (1993) p. 133-9 ] J. Enderlein, S. Makarov, E. Chilla, H. J. Frohlich [ Sens. Actuators B (Switzerland) vol.24-25 (1995)p.65-8] T. Lalinsky, J. Kuzmik, M. Porges, S. Hascik, Z. Mozolova, L. Grno [ Electron. Lett. (UK) vol.31 no.22 (1995) p. 1914-5] A. Dehe, V. Krozer, K. Fricke, H. Klingbeil, K. Beilenhoff, H.L. Hartnagel [ Electron. Lett. (UK) vol.31 no.25 (1995) p.2187-8] P.F. Ruths. A. Ashok. S. J. Fonash. J. M. Ruths [ IEEE Trans. Electron Devices (USA) vol.ED-28 (1981) p. 1003-9] L.M. Lechuga, A. Calle, D. Golmayo, P.Tejedor, F. Briones [ J. Electrochem. Soc. (USA) vol. 138 (1991)p.l59-62] L.M. Lechuga, A. Calle, D. Golmayo, F. Briones [ Sens. Actuators B (Switzerland) vol.7 (1992) p.614-8 ] L.M. Lechuga, A. Calle, D. Golmayo, F. Briones, J. De Abajo, J. G. De La Campa [ Sens. Actuators B (Switzerland) vol.8 (1992) p.249-52 ] V.A. Berezkin, V.P.Grabchak, V.N. Inkin, G.S. Rychkov [ Sens. Actuators A (Switzerland) vol.28 (1991)p.l91-5] Y. Zhou et al [ Biosensors Bioelectron. (UK) vol.6 (1991) p.595-607 ] P.Huertas, G. Mier, M.L. Dotor, J.V. Anguita, D. Golmayo, F. Briones [ Sens. Actuators A (Switzerland) vol.37-38 (1993) p.512-5 ] H. Zappe, H.E.G. Arnot, R.E. Kunz [ Sens. Actuators A (Switzerland) vol.41-43 (1994) p. 141 ] W. Gopel [ Sens. Actuators B (Switzerland) vol.18-19 (1994) p.1-21 ]

22.2 GaAs nuclear particle detectors C. Buttar July 1996

A

INTRODUCTION

Radiation detectors based on gallium arsenide are being developed for several applications. For many years X-ray and y-ray detectors, based on a variety of semiconductors (GaAs, CdTe, HgI2), have been under development [I]. These materials have a large enough bandgap that the detectors have, in theory, only a small leakage current and so can operate at room temperature providing resolutions similar to those achieved with Ge detectors. X-ray and y-ray detectors based on GaAs are being developed for X-ray astronomy [2,3], medical imaging [4] and synchrotron radiation imaging applications. Pixel and microstrip detectors based on GaAs for high precision spatial measurements have been developed for use in high energy physics experiments [5]. Some of the important material parameters for use of GaAs as a detector are given in TABLE 1. The photoelectric-Compton cross-over is the energy above which the photoelectric interaction is no longer the dominant photon interaction. The mean free drift lengths are a measure of how far carriers will drift before trapping, when the device is under operating voltage. dE/dx for minimum ionising particles is the most probable energy loss of high energy charged particles. Radiation length is the length of material that will reduce the energy of a high energy electron by 1/e. TABLE 1. GaAs material properties relevant for radiation detection. Parameter Photoelectric-Compton crossover Linear attenuation 60keV Linear attenuation 511 keV dE/dx for minimum ionising particles Radiation length e-h pair creation energy Fano factor Electron mean free drift length (bulk GaAs) Electron mean free drift length (LPE GaAs) Hole mean free drift length (bulk GaAs) Hole mean free drift length (LPE GaAs)

B

EPITAXIAL DETECTORS

Bl

Material

Value 120 ke V 10.8 cm'1 0.436 cm"1 5.6 MeV/cm'1 2.3 cm 4.3 eV 0.18 65-200 \xm > 1 mm 200-800 ym > 1 mm

Reference [6] [6,7]

see Section B [8]

Detectors have been fabricated using material grown by liquid phase epitaxy (LPE). Early work used material grown by the tilt-tube furnace method, giving layers up to 100 \xm thick with carrier concentrations in the range 5 x 1012- 1014cm'3 and mobilities > 8000 cm2/Vs [9]. Other workers used a vertical dipping method which gave 'undoped' carrier concentrations of around 7 - 15 x 1014cm"3. By careful addition of Fe the material was compensated and carrier concentrations in

the range 2 x 1012 - 1014 cm"3 were achieved for layer thicknesses between 40 and 100 |um. Mobilities of 7000 - 7500 cm2/ Vs at 300 K were reported [10]. Detectors were fabricated by evaporating Au layers onto the epi-layer to form a rectifying contact. The Au layer was usually a pad with a diameter of a few millimetres. The epi-layer was grown on an n+-substrate and this had an ohmic contact formed on it using, for example, silver paste. The detectors are operated as reverse biased diodes. B2

Results

The good quality of LPE material has resulted in detectors with good energy resolution. Some results of the energy resolution measured for y-rays are given in TABLE 2 below. TABLE 2. Summary of results of y-ray resolution for devices based on LPE material. Energy (ke V) 59.5 140 122 59.5 28.5 122 59.5 59.5 59.5 59.5 122 _8&

59.5 B3

Resolution 1.0% 2% 2№> 4% 9% 2.4% 6A% 4.9% 3_9% ^0% 33 32

I 3.9

Comments 130 K, best result Measured at 300 K Using low leakage Current sample 60-70 pun thick Fe compensation

Measured at RT Measured at 313 K

I

Ref [8] [8] [8] [8] [8] QO] [10] [11] [12] [12] [13] [13]

I [13]

Electron-Hole Pair Creation Energy

The energy required to generate an electron-hole pair, e, in GaAs was measured using LPE detectors [8,10] irradiated by alphas and betas. The results are presented in TABLE 3 below. The detectors were calibrated using a precision capacitor and comparing to the response of a Si detector for which the electron-hole energy is assumed. Eberhardt [8] states that the detectors were not operated at saturated voltage which may lead to a systematically high result for €. The data for measurements with 241Am alpha particles of Kobyashi and Eberhardt disagree at the level of 2%. There is a systematic shift between the data across the entire temperature range over which e was measured. This suggests that there are differences in the calibrations for the two results. The exact situation is unresolved. Although the 2% disagreement is small it introduces an error in the energy calibration of around 2%, e.g, 100 keV on the 5.5 MeV 241Am alpha peak.

TABLE 3. Electron-hole pair creation energy. Radiation 244Cm, 241Am, 239Pu Beta 241 Am

e at 300 K 4.35 +/- 0.02 eV 4.57 +/- 0.05 eV I 4.27 +/- 0.05 eV

€ as a function Eg 2.35 Eg+ 0.74 eV [ 2.7 Eg + 0.41 eV

Ref [10] [10] | [8]

In addition to the discrepancy between the two measurements for alpha particles, Kobyashi reported a different value for electrons [10] at the level of 5%. This is a small effect and in general does not affect the operation of a device except where the ultimate calibration is required. Knowing e allows the expected number of electron-hole pairs, N, generated when a particle loses energy E to be calculated from:

€

B4

Conclusion

The excellent quality of LPE GaAs has resulted in detectors with good energy resolution but the difficulty of growing layers of sufficient thickness and low enough impurity concentration has meant that the use of this material has been limited. C

BULK GROWN GaAs

Cl

Material

Early work on bulk grown GaAs found that the devices had poor charge collection and hence poor energy resolution as well as being plagued by low frequency current oscillations. Consequently early workers concentrated on good quality epitaxially grown GaAs (see above). More recently, with improved material, there has been a resurgence in the interest in GaAs detectors fabricated from bulk grown GaAs for future high energy physics experiments and as room temperature y-ray spectrometers. In general the detectors are fabricated on standard commercial SI wafers grown by either the LEC or VGF technique, although results have also been obtained with 'low' resistivity material: p - 106 - 107 Q cm. The detectors are usually fabricated as a surface barrier device with a Schottky contact e.g. Ti/Au, on one face and an 'ohmic' contact, e.g. Ni/AuGe/Ni/Au, on the other. For further details see, for example, [14]. The devices are then operated as reverse biased diodes and have typical leakage current densities of 10 - 30 nA mm"2 at room temperature. C2

Studies of charge collection

Early detectors were very inefficient with little of the charge being collected. Initially it was thought that bulk grown GaAs detectors had a uniform electric field across their bulk and the loss of charge was due to the effect of traps. This is quantified as the charge collection efficiency (CCE), defined as: p p p __ ^measured ^ionised

where Qmeasured is the charge measured using a calibrated amplifier and Qionised the charge expected from calculating the number of electron-hole pairs generated using e. However McGregor [15] showed that the presence of deep levels, specifically EL2, could modify the electric field. This leads to a high field region (with high charge collection efficiency) which penetrates into the GaAs below the blocking contact, below which there is a low field region where the charge is lost. The penetration rate has been measured in a number of ways: i) ii) iii)

SEM [16] OBICS [17] optical probe [18].

The field structure measured by probing the surface voltage across a cleaved surface is shown in FIGURE 1 [18].

Electric field (volts per micron)

Applied voltage

Distance from p

contact (microns)

FIGUREl.

All these techniques show that the high field region typically penetrates the bulk at a rate of 0.7 1.0 jim/V. It is not clear if the variation in the penetration rate is a systematic difference between the different measuring techniques or due to the measurements being carried out on a wide variety of samples. The electric field has been calculated numerically [19,20]. The resulting field profile does not agree with the field map measured by the surface probe method. Agreement can be achieved by modifying the capture cross-section of EL2 so that it increases above a critical field [21].

This improved understanding of the behaviour of the devices has indicated that devices that have a high field region through the entire bulk would have a high CCE. This has been difficult to achieve due to the devices breaking down near the voltage at which they are fully active. However by optimising the 'ohmic' contact 100 |im thick devices with breakdown voltages exceeding 500 V have been fabricated and they have shown 100% CCE [22]. Although fixll penetration of the electric field has led to some devices with 100% CCE others have lower CCEs due to trapping in the high field region. The method to optimise the CCE is not clear. Studies have shown that there is a variation in the mean free drift path for different substrates, depending on various parameters such as manufacturer and wafer resistivity. A direct correlation between wafer resistivity, [EL2+] measured using MCDM and the charge collection efficiency has been observed. This indicates that electron trapping in these devices is primarily due to EL2+. For EL2+ to have a significant effect compared to other ionised deep levels requires EL2 to have a field enhanced capture cross-section consistent with the capture cross section behaviour required to produce the observed electric field profiles [21]. Another important effect is wafer preparation. By careful polishing before metallisation, detectors with high CCE have been fabricated regardless of substrate. C3

Results

The Y-ray energy resolution and the charge collection efficiency of various detectors are given in TABLE 4. TABLE 4. Results of radiation detectors based on bulk SI GaAs. Energy 122 _88 59.5 _22 59.5 59.5 59.5 122 59.5 122

CCE 87.6 +/- 3.5 87.0 +/- 3.9 81.9+/-6.1 88.0 +/- 7.0 43.7 +/- 8.8 63.1 +/- 14.7 87.6 +/- 4.8 82 71 I 52

Resolution 4.0 +/- 0.2% 4.5 +/- 0.2% 7.5+/-0.6 8.7 +/- 7.0 20.2 +/- 4.1 23.4 +/- 5.4 5.5 +/- 0.3 9/7% 11.3% 1 12.8%

Comments 80 [im thick, room temperature, V = 80 V Room temperature Limited by electronic noise 350 ym thick, V = 350 V 200 nmthick,V = 160 V 70 ^m thick, V = 80 V 200 nm thick, V = 500 V 200 inn thick V = 400 V 1 400 jim thick V = 500 V

Ref [23] [23] [23] [23] [4] [4] [4] [24] [24] | [24]

Energy resolutions have improved considerably, especially for thin detectors where the resolution is starting to approach that of LPE devices. This is due to the high CCEs that are being achieved for thin detectors. However, achieving similar resolutions in thicker devices remains problematical because of the lower CCEs. In general the performance at room temperature seems to be ultimately limited by leakage current. D

CONCLUSION

Improved understanding of the operation of detectors based on bulk grown GaAs has led to a much improved performance in terms of X-ray and y-ray resolution and microstrip devices with sufficient signal-to-noise to be used as position sensitive detectors. However, there are still uncertainties as to how to optimise the fabrication of these devices as the relative effects on CCE due to substrate choice, contact fabrication and wafer preparation are not yet fUlly understood.

Radiation detectors based on GaAs grown by LPE, LEC and VGF have been developed and used as X-ray and Y"raY detectors, and position sensitive detectors. The LPE devices have demonstrated good energy resolution but material of sufficient thickness and low doping concentration is difficult to produce. Devices based on bulk grown GaAs are subject to charge trapping effects and require voltages >1 V/|am to ensure that they are fully active. However progress has been made in understanding their operation leading to an improvement in performance. The energy resolutions of thin (~100 (im) devices are approaching those of LPE devices but further work is required to produce thicker devices with good energy resolution. REFERENCES [I]

[2] [3] [4]

[5] [6] [7] [8] [9] [10] II1] [12] [13] [14] [15] [16] [17] [18]

[19] [20] [21] [22] [23] [24]

D.S. McGregor, J.E. Kammerand [ in Semiconductors for Room Temperature Nuclear Detector Applications, Ed. T.E. Schlesinger, RB. James, Semiconductors and Semimetals, vol.43 (Academic Press, 1995)] T J . Sumner, S.M. Grant, D. Alexiev, K.S.A. Butcher [Nucl. InstrumMethods Phys. Res. A (Netherlands) vol. 348 (1994) p.518-21 ] A.D. Holland, A.D.T. Short, T. Cross [ Nucl InstrumMethods Phys. Res. A (Netherlands) vol.346 (1994) p.346-71] M.E. Fantacci [Proc. 3rdlnt. Workshop on GaAs and Related Compounds, San Miniato, Italy, 2124 Mar 1995, Eds. P.G. Pelfer, J. Ludwig, K. Runge, H.S. Rupprecht (World Scientific, Singapore, 1996)p.225-31] K.M. Smith [Nucl. InstrumMethods Phys. Res. A (Netherlands) vol.368 (1995) p.220-3 ] G.F. Knoll [ Radiation detection and measurement (John Wiley & Sons 1989) ] W. Bencivelli et al [ Nucl. InstrumMethods Phys. Res. A (Netherlands) vol.310 (1991) p.210 ] J.E. Eberhardt, RD. Ryan, AJ. Tavendale [ Nucl. InstrumMethods Phys. Res. A (Netherlands) vol.94 (1971) p.463-76] H.G.B. Hicks, D.F. Manley [ Solid State commun. (USA) vol. 7 (1969) p. 1463-5 ] T. Kobyashi et al [ IEEE Trans. Nucl. Sci. (USA) vol.23 (1976) p.97-101 ] P.E. Gibbons, J.H. Howes [ IEEE Trans. Nucl.Sci. (USA) vol. 19 (1972) p.353-7 ] D. Alexiev, K.S.A. Butcher [Nucl. InstrumMethods Phys. Res. A (Netherlands) vol.317 (1992) p.111-5] W. Bencivelli et al [ Nucl. InstrumMethods Phys. Res. A (Netherlands) vol.346 (1994) p.372 ] S.P. Beaumont et al [Nucl. InstrumMethods Phys. Res. A (Netherlands) vol.321 (1992) p. 172 ] D.S. McGregor, G.F. Knoll, Y. Eisen, R Brake [ IEEE Trans. Nucl. Sci. (USA) vol.39 (1992) p. 1226-36] S.B. Beaumont et al [ Nucl. Phys. B (Netherlands) vol.32 (1993) p.296-9 ] M. Alietti et al [Nucl. InstrumMethods Phys. Res. A (Netherlands) vol.355 (1995) p.420-4 ] K. Berwick, M.R. Brozel, CM. Buttar, M. Cowperthwaite, Y.Hou [ Proc. Semiconductors for room-temperature radiation detector applications, San Francisco, California, USA, 12-16 April 1993 (MRS Symp. Proc vol. 302) ] T. Kubicki et al [Nucl. InstrumMethods Phys. Res. A (Netherlands) vol.345 (1994) p.468-73 ] J.W. Chenet al [Nucl InstrumMethods Phys. Res. A (Netherlands) vol.365 (1995) p.273-84 ] D.S. McGregor et al [ J. Appl. Phys. (USA) vol.75 (1994) p.7910-15 ] M. Aletti et al [Nucl InstrumMethods Phys. Res. A (Netherlands) vol.362 (1995) p.344-8 ] W. Bencivelli et al [Nucl. InstrumMethods Phys. Res. A (Netherlands) vol.355 (1995) p.425-7 ] M. Dogru et al [Nucl. InstrumMethods Phys. Res. A (Netherlands) vol.348 (1994) p.510-3 ]

22.3 High temperature applications for GaAs L.P. Sadwick and RJ. Hwu August 1996

A

INTRODUCTION

This Datareview outlines the data on high temperature applications for GaAs. GaAs has a larger bandgap than silicon and has the most mature processing and fabrication technologies of the group III-V and wide bandgap semiconductors. High temperature applications and potential limitations of GaAs-based electronics are given. B

HIGH TEMPERATURE APPLICATIONS FOR GaAs

High temperature electronics for automobile, aeronautical, geothermal, nuclear, and space applications have received increased attention in the last few years [1-14] and the use of high temperature electronics is expected to increase substantially in the near future. Commercial and military applications for high temperature electronics include oil well logging and other geothermal applications, automobile electronics, aerospace electronics and optoelectronics, ground and space communications, industrial controllers, distributed control systems [15,16], nuclear reactors, high power solid state microwave circuits, high temperature stable sensors [16], elevated environment systems, and accelerated reliability testing of electronics including GaAsbased monolithic microwave integrated circuits (MMICs) [17-19]. GaAs-based devices and circuits have been investigated for these high temperature applications. The economic benefits of using distributed control systems is well understood [15]. However to realize distributed controls for applications including commercial and military automobile and aerospace engine systems will require high temperature analog, digital and sensor technology. Instrumentation for high temperature and harsh environments needs to localize 'intelligence' to obtain increased accuracy and resolution, reduce weight and system response time, and to increase reliability and efficiency [20]. The distributed control building blocks are i) ii) iii)

input sensing signal amplification, conditioning and processing actuator activation and control [21].

The major high temperature components needed to realize high temperature distributed control systems are i) ii) iii) iv)

microprocessors and microcontrollers pressure, temperature, vibration, speed, and light sensors relay, solenoid, motor drivers, etc. data transmission links [15].

Over 50% of electronic failures are due to high temperature [21]. High temperature electronics

will result in more reliable systems and will also eliminate or reduce the need for active or passive cooling of these systems which will further reduce the system weight and complexity. The cost of electronic components in an average automobile has increased from $20 in 1955 to $782 in 1993 and is projected to be over $1000 by the turn of the century [22]. Examples of projected operating temperature ranges for commercial automobiles are passenger compartment: -40 to 85°C; engine compartment: -40 to 165°C; and wheel-mounted systems: -40 to 250 0 C [21,22]. Oil and gas wells will require electronics that operate at 225 0 C and above. Geothermal applications will have a need for electronics that can operate at 300 0 C and above. Aerospace will need electronics that can operate to i) 250 0 C for braking systems and mounted avionics electronic controls; ii) 3000C for engine control and monitoring; and iii) 350 0 C for smart composite skin applications [21]. With the exception of microwave ovens (which may require 500 0 C operation) most consumer applications will require electronics that can operate in the range of 200 0 C. A very promising candidate for high temperature applications is the GaAs metal-semiconductor field effect transistor (MESFET) and related heteroFET devices. Some reports on the high temperature characteristics of GaAs MESFETs [7-14] have indicated problems caused by leakage currents between the MESFET terminals and the semi-insulating GaAs substrate at high temperatures. A recent comparison of GaAs junction field effect transistors (JFETs) with MESFETs for high temperature operation indicated that both devices behave similarly and show a notable degradation at temperatures primarily due to an increase in substrate conduction resulting from thermally generated carriers [14]. One potential advantage that JFETs may have over MESFETs is an effective higher p-gate/n-channel junction barrier height as compared to the Schottky/n-channel barrier height. The exact temperature at which thermal leakage currents affect device/circuit performance depends to a large extent on the device/circuit structure/architecture and materials properties.The leakage current due to thermally generated electrons being injected from the substrate into the channel can, at high temperatures, initiate an avalanche multiplication that results in substantially larger source and drain currents compared to the zero or low substrate injection condition [23]. These substantially larger drain and source currents cannot be controlled by gate bias. At temperatures above 2000C, the GaAs substrate should not be viewed as being semi-insulating but, instead, as being semiconducting. From room temperature to over 300 0 C, the thermal leakage current for proton bombarded or semi-insulating GaAs follows EQN (1) below [24]: I(T) = I0 exp(-Ea/kT)

(1)

where I(T) is the temperature dependent thermal leakage current, I0 is a constant dependent on the geometry and type of conduction mechanism, Ea is the thermal activation energy of the conduction current process, k is the Boltzmann constant, and T is the temperature in Kelvin. TABLE 1. Thermal activation energies observed in GaAs. Sample

Substrate type

Ea(eV)

Identification

MEHl

LEC

0.74

EL2

TM-I

LEC

0.76

EL2

14B

HB

0.46

Cr3+

MOD 12-2

LEC

0.75

EL2

C5-1

LEC

0.67

NPl

MOD 5-2

LEC

0.74

EL2

Ln (I) Amps

TABLE 1 lists typical activation energies of the dominant thermal conduction currents observed in semi-insulating and proton bombarded GaAs and the deep levels that correspond to the respective activation energies. Proton bombardment is ineffectual in isolating GaAs devices at elevated temperatures due to the thermal release of electrons from the traps associated with proton bombardment. These electrons become free carriers and contribute to the current conduction process. Provided that the temperature does not exceed 4000C5 at which point the traps associated with the proton bombardment are annihilated [25], reversible retrapping takes place as the temperature decreases toward room temperature [24]. FIGURE 1 shows plots of thermal leakage current in proton bombarded material versus reciprocal temperature (K"1). The Arrhenius plots in FIGURE 1 represent four different LEC substrate samples: C5-1, TM-I, Mod 5-2, and Mod 12-2 which had activation energies of 0.67, 0.76, 0.74, and 0.75 eV, respectively. In general, for a given sample, all Arrhenius plots obtained from data analyzed from that sample in the voltage range of 5 to 25 V (i.e., a set of five Arrhenius plots at 5, 10, 15, 20, and 25 volts, respectively) resulted in similar activation energy values usually within ±15 meV of each other. As the thermally activated conduction currents are not influenced or controlled by the gate bias they contribute a highly undesirable leakage current to MESFET and HBT device operation at

FIGURE 1. Arrhenius plots of samples C5-1, TM-I, Mod 5-2, and Mod 12-2. The curve fits are as follows: for sample TM-I at 5 V; In(I) = 10.406 - 0.75878 (q/kT), R2 = 0.999; for Mod 5-2 at 20 V, In(I) = 8.1629 - 0.73668(q/kT), R2 = 0.998; for Mod 12-2 at 20 V, In(I) = 8.5514 - 0.75433 (q/kT), R2 = 1.000, and for C5-1 at 10V, In(I) = 4.3145 0.67070 (q/kT), R2 = 0.999. Note that the slope values are equal to the effective activation energy, Ea.

high temperatures. FIGURE 2 shows a typical family of current vs. voltage (I-V) curves measured at temperatures ranging from room temperature to 2500C of a proton isolated GaAs structure. The data shown in FIGURE 2 were reversible and reproducible with no permanent changes observed for samples heated to the range of 250 to 3000C. After exposure to temperatures as high as 300 0 C and subsequent cooling to room temperature, no measurable changes in the room temperature I-V characteristics were detected nor was there any evidence or indication of a change in the original room temperature proton isolation characteristics. The same situation holds true for samples that underwent multiple high/room temperature cycling or were held at high temperatures for long durations (i.e., from 1 hour to several days) [24]. There are several important observations to make with regards to FIGURES 1 and 2: i) over a large range of applied voltage (i.e., except at very low and very high applied voltages), the current-voltage relationship is linear at a given, fixed temperature; ii) the current is relatively small at or near room temperature (i.e., 2.7 x 1O"8 amps at 5 volts for sample 14B; 6.6 x 10'10 amps at 5 V and 2.15 x iO"9at 10 V, respectively, for sample MEHl; and 6.3 * 10"9 amps at 10 V for sample TM-I) indicative of the good isolation between adjacent MESFET structures at room temperature; and iii) the current increases in a superUnear (exponential) fashion with temperature.

I (Amps)

There have been several reports on low leakage GaAs MESFETs [14] and GaAs-based heterojunction FET devices [26-29] for high temperature applications. The techniques include biasing of the substrate at elevated temperatures, etching the substrate and/or multilayer epitaxial

V (Volts) FIGURE 2. Family of I-V curves for a proton isolated GaAs sample as a function of temperature. Note the values of current for temperatures less than 1000C are below the resolution of the linear 1 mA full scale plot and thus appear to overlap each other at zero current on the voltage axis.

3

CO

O

O

O

O

O

<

I (Ohmic Contact to Substrate) (A)

structures. The essence of the thermal current reduction technique that employs substrate biasing is the high temperature electronic technique (HTET) [23,31-33]. The HTET technique involves using a substrate that exhibits rectifying behaviour to the source and drain ohmic contacts as shown in FIGURE 3.

FIGURE 3. Linear plot of the substrate to drain ohmic contact leakage current, I, versus bias voltage, V, as a function of temperature. Voltage applied to source or drain with substrate at common potential. I-V curves range from 75 to 3000C in 25 0 C increments.

A single isolated MESFET with the source at ground potential will be used to illustrate and explain the substrate bias effect. When the substrate is left floating or grounded and a positive bias is applied to the drain, the drain-to-substrate 'diode' becomes forward-biased with a resultant thermal 'leakage' current that is exponentially dependent on temperature. FIGURE 4 is a plot of the drain contact to substrate leakage current for VDS = 5 V, Vsub = 0 V, as a function of temperature. The key to successful high temperature operation using HTET is to eliminate or substantially reduce the drain-to-substrate (and source to substrate) leakage current. By treating the substrate as an active element (essentially, the fourth terminal) of the MESFET, it is then possible to turn off the source-to-substrate and drain-to-substrate 'diodes' by applying a reverse bias voltage. This can be accomplished by applying a voltage to the substrate that is more positive by typically 1 V (or less) than the largest positive voltage applied to the drain. Turning off these 'parasitic' diodes thus produces a remarkable reduction in the drain leakage current.

Isub-to-drain (mA)

(A) Isub-to-drain

Temperature

(C)

FIGURE 4. Substrate to drain contact current vs. temperature for VDS = 5 V and V8111, = 0 V.

The effects of HTET on MESFET I-V characteristics are illustrated in FIGURES 5 and 6. FIGURE 5 shows the large, uncontrolled increase in the saturation region (UnZdV08 at 25O0C resulting from thermally-generated leakage current for a substrate bias of 0 V. The effects of transistor breakdown are also evident in FIGURE 5. FIGURE 6 shows the ID-VDS curves of the

Vds (V) FIGURE 5. Drain current, ID, for a 20 x 1 \im depletion MESFET at 2500C. VGS is from 0.4 V to -1.8 V in -0.2 Vsteps; Vmh is 0 V. Note: the current scale is 2.5 times as large as that of FIGURE 6.

same MESFET at 2500C with an applied substrate voltage of 6 V. The ID-VDS curves in FIGURE 6 clearly show that MESFET operation is well-controlled and that breakdown of the MESFET is avoided.

Vds (V) FIGURE 6. Drain current, ID, for a 20 x 1 ^m depletion MESFET at 25O0C. VGS is from 0.4 V to -1.8 V in -0.2 Vsteps; V ^ is 6 V.

HTET has been shown to increase the high frequency gain (S21) of discrete MESFETs and high electron mobility transistors (HEMTs) [34]. Substrate biasing also proved effective in enhancing the gain and frequency response of GaAs MESFET-based operational amplifiers and inverters [33] at temperatures at or above 175°C compared to the 0 V or floating substrate cases. Transconductance (g^ = diydV^) was favourably controlled by the application of the HTET. Using the HTET also increased the breakdown voltages of GaAs MESFET devices at high temperatures. FIGURES 7-10 show the effect of HTET on controlling the thermally generated leakage current component in the drain current at elevated temperatures. FIGURE 11 shows the output conductance, G0, as a function of temperature with (Vsub = 6 V) and without (Vsub = 0 V) substrate bias. FIGURES 12 and 13 show the effect of HTET on G0 at elevated temperatures. It can be seen that a substantially smaller G0 value is obtained for substrate voltages nearly equal to or greater than the drain voltage. Another promising approach is to grow epitaxial wide bandgap layers (e.g., AlAs) between the substrate and the active GaAs channel. By including these higher bandgap epilayers in GaAs FET structures, it is possible to suppress the substrate thermal leakage current component. It has been demonstrated that an increase in the AlAs buffer layer thickness for epitaxially grown GaAs MESFETs having NiGeInW ohmics and TiPtAu Schottky gates resulted in a substantial reduction in the substrate thermal current and the output conductance [26]. FIGURE 14 shows the effect of the AlAs buffer layer on the 350 0C ID-VDS characteristics.

< > O

Il

CO

O) > <§> U)

T?

FIGURE 8. Total drain current, ID, versus VDS for VGS = 0 V at selected temperatures. Note that large leakage currents are present at 175 and 250 0 C.

Absolute Current (A)

FIGURE 7. Typical 30O0C plots of ID and I s vs. VDS for V№b = 0 and 6 V for a 20 x 1 ^m depletion MESFET.

Temperature (C) FIGURE 9. Comparison of ID versus VDS for VGS = 0 V at Vmh = 0 and 6 V. Note the thermal leakage current resulting in breakdown conditions for Vmh = 0. For Vmh = 6 V, the curve is well behaved showing no signs of breakdown.

FIGURE 10. Typical absolute values of ID, I s , I G and I rob vs. temperature for V8111, = 0 and 6 V for a 20 x 1 ^m depletion MESFET. VGS = 0VandV D S = 5 V.

Temperature

(C)

Vsub (V)

FIGURE 11. Output conductance, G0, vs. temperature for a 20 x l jLim depletion MESFET for V8110 = 0 and 6 V. VGS = 0 and VDS = 2.75 V.

FIGURE 12. Output conductance, G0, vs. V8110 for a 20 x 1 urn depletion MESFET at VGS = 0.4,0, and -0.4 V andV DS = 3.75V.

Vsub (V) FIGURE 13. Output conductance, G0, vs. V8110 for a 20 x 1 u m depletion MESFET at VGS = 0.4, 0, and -0.4 V and VDS = 4.75V.

It is well known that, in the absence of significant thermal leakage currents and before the onset of thermal breakdown, the FET drain saturation current decreases with increasing temperature due to a decrease in mobility and saturation velocity at higher temperatures. This decrease in the FET current can be balanced, under certain conditions, by a shift in the negative direction of threshold voltage, V1, which produces a stable, zero temperature coefficient (ZTC) bias point (at only one drain current (IJ, drain-to-voltage ( V J pair) for the FET I4-V^ characteristics [11,13].

MESFET I-V Characteristics with 1500A AlAs Buffer at 350 0 C

MESFET I-V Characteristics with 2500A AlAs Buffer at 3500C

FIGURE 14. The effects of AlAs buffer thickness on the 350 0C MESFET I-V characteristics.

Traditional Schottky diode level shifters do not work well at elevated temperatures due to the decrease in barrier height with increasing temperature. Nichrome (NiCr) resistors coupled with a ZTC-biased MESFET acting as a constant current source have been shown to have a very low temperature output level sensitivity [35] as shown in FIGURE 15. Recently a comparison between GaAs MESFET, AlGaAs/GaAs HEMT5 and AlGaAs/GaAs HBT devices was performed [36]. Due to a larger Schottky barrier height to AlGaAs as compared to GaAs, the HEMT devices had gate leakage currents that were typically three to four times less than the MESFET devices. It was also experimentally observed that MESFETs and HEMTs had superior high temperature RF performance compared to HBTs. Conventional heterojunction bipolar transistors (HBTs) do not perform well above 1750 C due to a number of factors including emitter-base interdiffusion and leakage currents which translate to a lack of Vbe control. FIGURES 16 and 17 show typical DC H ^ (current gain, P) characteristics of lower and higher Al-content AlGaAs/GaAs HBTs, respectively, as a function of temperature. As can be seen from these two figures, the gain decreases as the temperature increases. The higher Al-content AlGaAs/GaAs HBT has a higher room temperature gain but undergoes permanent degradation during operation at temperatures of approximately 1000C or higher. This degradation is presumably due to interdiffusion in the emitter base junction region. For both types of HBTs, the leakage current increased dramatically at higher temperatures.

Levelshift with diodes

Temperature (K) Levelshift circuits.

Simulation of different levelshifters with a shift-voltage of about 6 V.

FIGURE 15. Simulation of the level-shift output for a shift-voltage of approximately 6 V for a level shifter consisting of nichrome (NiCr) resistors coupled with a ZTC-biased MESFET acting as a constant current source as a function of temperature.

Detailed simulations of AlGaAs/GaAs HBTs have shown that above 200 0 C, thermal generation currents dominate at low base currents thus making it difficult to turn off the HBT [37]. AlGaAs/GaAs/AlGaAs DHBT circuits have been shown to work up to 200 0 C and individual DHBTs were functional at 300°C with considerably lower P and Early voltage values compared to room temperature operation [38]. In addition to thermally generated currents, there are other technological considerations and limitations when operating GaAs-based devices at high temperatures. The main concern is the HBT Gain Curves

HBT Gain Curves

FIGURE 16. Current gain vs. collector current as a function of temperature for a low Al-content AlGaAs HBT.

FIGURE 17. Current gain vs. collector current as a function of temperature for an AlGaAs HBT with a higher content than in FIGURE 16.

limitations when operating GaAs-based devices at high temperatures. The main concern is the ohmic and Schottky contacts. Conventional ohmic and Schottky contacts are problematic at high temperatures. TiPtAu Schottky contacts on GaAs are not suitable for use above 200 0 C. The widely used alloy tunnel ohmic contact, AuGeNi, is acceptable for use up to 250 0 C but is extremely marginal and not reliable at higher temperatures and unacceptable for long term operation at temperatures approaching 3000C. AuGeNi contacts are formed by a rapid alloying cycle limiting the kinetics of a liquid phase reaction that occurs at a relatively low temperature. Since the rapid alloying cycle does not allow the reactions to go to completion, these contacts are highly metastable and further reactions can take place particularly for prolonged operation at high temperature. Schottky contacts that can withstand annealing temperatures of 8500C and higher that are used in self-aligned gate (SAG) fabrication processes generally work quite well for high temperature GaAs-based FET applications. The Schottky contacts that have been shown to work at high temperatures include TiW [39], TiWSi [27], TiWN [10,41] and LaB6 [42,43]. All Schottky contacts to semiconductors show a reduced barrier height at high temperatures. The main effects of a reduced barrier height are an increased reverse bias leakage current and a smaller forward bias voltage swing. The traditional and highly utilized AuGeNi ohmic contact attains a low contact resistivity value due to alloy formation producing a P-AuGa phase. AuGeNi is considered to be the best ohmic contact in terms of electrical properties. However, the AuGeNi/GaAs system is thermally unstable for reasons primarily related to the P-AuGa phase [44-51]. A number of potential high temperature-stable ohmic contacts can be found in the literature [4673]. From a device standpoint the primary considerations are contact resistivity and contact resistance to GaAs and related compounds (i.e., InGaAs). Higher performance ohmic contacts lead directly to higher frequency devices. Pdln-based ohmic contacts have been investigated as

an alternative to AuGeNi. Indium agglomeration was observed on the PdIn contact due to the low melting point of In and poor In flux control [52]. This is not desirable for device applications. The major impetus for devising thermally stable, low resistance ohmic contact schemes is greater processing freedom and uniform dimensional downsize scaling for VLSI GaAs technology. Replacing Au with a refractory ohmic contact metal that forms high melting point intermetallic compounds when annealed at high temperatures appears to be a very promising approach [53-63]. However the contact resistance may be significantly higher for these refractory ohmic contacts compared to that of AuGeNi. A potential drawback to some of these refractory ohmic contacts is the need for a high temperature anneal to form the proper low resistance intermetallic compounds. W-based ohmic contacts are currently employed in HBT technology [38,71,72]. It has also been experimentally observed that In-based ohmic contacts that react with GaAs to form InGaAs layers at elevated temperatures produce ohmic contacts that are significantly more stable at high temperatures than AuGeNi [50,64,67,69]. Solid phase epitaxy and limited reaction ohmic contact formation has also been extensively studied [50,64-70]. Again the driving force for studying these contact technologies has been enhanced process integration. Towards this end, high temperature contact stability and resistivity data is typically acquired after 100 hours at temperatures in the range of 350 to 400 0 C. A complete and detailed discussion of alternative ohmic contacts is beyond the scope of this Datareview. A partial list of alternative ohmic contacts studied to date includes NiInW [56,57], Ni/In/Ge/Mo [40], GeMoW [51], NiInWNx [55], MoGeInW [63], GeInW, NiInW [60], NiGeW with low Au content [48], NiGe [46], solid-phase regrowth Si/Ni [64], Pd-In [50], Ge/Pd [66,68,69], Si/Pd [66] and Pd-In-Ge [67,69] contacts. Ge/Pd contacts appear to be thermally stable at 300 0 C but degrade at higher temperatures. Further work is still needed on high temperature stable contacts to GaAs. REFERENCES [I] [2] [3] [4] [5] [6] [7] [8] [9] [10] II1] [12] [13] [14] [15] [16]

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H.R Kawata, T. Oku, A. Otsuki, M. Murakami [J. Appl. Phys. (USA) vol.75 no.5 (1994) p.25306] N. Lustig, M. Murakami, M. Norcott, K. McGann [ Appl. Phys. Lett. (USA) vol.58 no. 19 (1991) p.2093-5 ] M. Masanori, HJ. Kim, Yih-Cheng Shih, W.H. Price, CC. Parks [ Appl. Surf. Sd. (Netherlands) vol.41/42 (1989) p. 195-200] H.G. Fu, T.S. Huang [ Solid-State Electron. (UK) vol.38 no. 1 (1995) p.89-94 ] K.G. Merkel, V.M. Bright, S.N. Schauer, C.I. Huang, G.D. Robinson [ Electron. Lett. (UK) vol.29 no.5 (1993) p.480-1 ] L.C. Wang, XZ. Wang, S.S. Lau, T. Sands, W.K. Chan, T.F. Kuech [ Appl. Phys. Lett. (USA) vol.56 (1990) p.2129-31] K. Mitani, Y. Imamura [ Electron. Lett. (UK) vol.29 no.7 (1993) p.589-90 ] K.G. Merkel, V.M. Bright, S.N. Schauer, J. Barrette [Mater. ScI Eng. B (Switzerland) vol.25 (1994) p. 175-8] M. Murakami,N. Lustig, W.H. Price, A. Fleischman [Appl. Phys. Lett. (USA) vol.59 no.19 (1991) p.2409-11] M.C Hugon, B. Agius, F. Varniere, C Dubon-Chevallier, J.F. Bresse, M. Froment [Appl. Phys. Lett. (USA) vol.58 no.24 (1991) p.2773-5 ] M.C. Hugon, B. Agius, F. Varniere, M. Froment, F. Pillier [J. Appl Phys. (USA) vol.72 no.8 (1992) p.3570-7 ] M. Ketata, K. Ketata, C. Dubon-Chevallier, R. Debrie [ J. Phys. D, Appl. Phys. (UK) vol.20 (1993) p.2055-9 ] K. Shenoy, V. Fonstad, G. Clifton Jr., J.M. Mikkelson [ IEEE Electron Device Lett. (USA) vol. 15 no.3 (1994) p. 106-8] HJ. Kim, M. Murakami, W.H. Price, M. Norcott [J. Appl. Phys. (USA) vol.67 no.9 (1990) p.4183 ] M. Murakami, Y-C. Shih, W.H. Price, E.L. Wilkie, K.D. Childs, CC. Parks [ J. Appl. Phys. (USA) vol.64 no.4 (1988) p. 1974-81 ] H.-J. Kim, M. Murakami, S.L. Wright, M. Norcott, W.H. Price, D. La Tulipe [ J. Appl. Phys. (USA) vol.68 no.5 (1990) p.2475-81 ] M. Murakami, W.H. Price, Y.-C Shih, N. Braslau, K.D. Childs, C C Parks [ J. Appl. Phys. (USA) vol.62 no. 15 (1987) p.3295-303 ] T. Sands, E.D. Marshall, L.C. Wang [J. Mater. Res. (USA) vol.3 no.5 (1988) p.914-21 ] E.D. Marshall et al [ J. Appl. Phys. (USA) vol.62 no.3 (1987) p.942-7 ] L.S. Yu, LC. Wang, E.D. Marshall, S.S. Lau, T.F. Kuech [ J. Appl. Phys. (USA) vol.65 no.4 (1989) p. 1621-5] L.C. Wang et al [ J. Appl. Phys. (USA) vol.69 no.8 (1991) p.4364-72 ] J. Tsuchimoto, S. Shikata, H. Hayashi [ J. Appl. Phys. (USA) vol.69 no.9 (1991) p.6556-63 ] L.C. Wang [ J. Appl. Phys. (USA) vol.77 no.4 (1995) p. 1607-10 ] CY. Nee, CY. Chang, T.F. Cheng, T.S. Huang [ J. Mater. Sci. Lett. (UK) vol. 7 (1988) p. 1187 ] J. Wurfl, A.G. Nassibian, H.L. Hartnagel, R. Langfeld, C Maurer [ Int. J. Electron. (UK) vol.66 no.2(1989)p.213-5] M. Missous, T. Taskin [ Semicond. Sci. Technol. (UK) vol.8 (1993) p. 1848-53 ] T. Nozu,N. Iizuka, Y. Kuriyama, S. Kongo [Electron. Lett. (UK) vol.29 no.23 (1993) p.2069-70 ]

Index

Index terms

Links

A Acoustic mode deformation potential

132

acoustic sensors

938

A/D converters

822

aerospace electronics

949

Aharanov-Bohm effect

178

179

68

Al, lattice parameter

499

AlAs, deformation potentials

172

AlGaAs carbon doping

102

DX centres

251

253

etch rates

672

673

etching

711 759

714

low temperature MBE growth

681

682

LPE growth

614

619

photoluminescence intensity

323

324

segregation coefficients of dopants

616

selective area epitaxy

633

thermal resistivity

936

727

728

758

AlGaAs/GaAs bandgaps

893

894

conduction band discontinuity

894

895

electron effective mass

894

electron mobility

52

hole effective mass

894

valence band discontinuity

894

895

AlGaAs/GaAs 2DEGs, see also quantum wells low temperature mobility

61

AlGaAs/GaAs heterostructures electron mobility enhancement

61

inverted interface

71

modulation doping

70

normal interface

71 This page has been reformatted by Knovel to provide easier navigation.

71

963

964

Index terms

Links

AlGaAs/GaAs quantum well lasers

887

alpha particle detection

943

Amorphous GaAs

155

crystallization temperature

156

dielectric constant

155

flash evaporation

155

growth

155

hydrogenation effects

157

optical absorption coefficient

156

159

160

optical gap

156

157

159

optical properties

155

phonon frequencies

158

photoconductivity

159

Raman spectra

159

refractive index

156

157

Urbach energy

157

158

amphoteric impurities

57 781

108

228

302

776

amphoteric native defect model

541

Amplifiers

793 954

799

821

822

829

low noise

793

794

operational

954

power

794

transimpedance

804

809

analogue circuits

821

822

anodic oxidation

478

479

anti-phase domain boundaries

456

457

antisite defects

338 685

344 686

346

559

560

anti-Stokes emission

293 102

630

632

647

Atomic layer epitaxy (ALE)

54

156

159

668 doping

69

GaAsP

633

GaInP

633

GaP

633

growth rate

668

InAs

633

InAsP

633

InGaAs

633

160

670

669

This page has been reformatted by Knovel to provide easier navigation.

159

824

965

Index terms

Links

Atomic layer epitaxy (ALE) (Continued) InP

633

atomic oxygen beam oxidation

491

492

71

73

Auger recombination

137

138

Auston switches

694

autocompensation

435

436

automobile electronics

948

949

Avalanche photodiodes, APD

866

atomic planar doping see also delta doping

SAGM APD

867

SAM APD

867

B Backgating

689

690

circuit design

847

848

field termination barriers

845

846

implant isolation

843

layout design

846

passivation

843

substrate material

841

surface

841

temperature dependence

840

841

testing

848

849

threshold voltage

838

trap-filled-limited model

838

Ballistic transport

66

horizontal

67

vertical

67

band edge

286

band filling

902

836

847 842

839 68

903

Band offset ratio AlGaAs/GaAs

545

AlGaInP/GaInP

550

GaInP/AlInP

551

GaInP/GaAs

549

Band structure

145

551

151

symmetry point energies

146

temperature dependence

147

148

Bandgap

151

186

direct

151 This page has been reformatted by Knovel to provide easier navigation.

186

789

966

Index terms

Links

Bandgap (Continued) effective bandgap narrowing

186

InAs

789

indirect

154

InGaAs

789

pressure dependence

153

154

temperature dependence

151

152

beta particle detection

943

bio sensors

939

bipolar integrated circuits

820

breakdown voltage

195

Bridgman growth

572

204

205

422

423

196

see also HBGaAs, VBGaAs Bulk GaAs, see Bridgman growth, Czochralski growth, LEG, etc. Bulk GaAs, electronic properties electron mean free drift length electron mobility

942 46

hole mean free drift length

47

942

Bulk modulus

14

15

adiabatic

14

15

isothermal

14

15

pressure dependence

14

15

temperature dependence

14

15

Burgers vector

418

420

buried heterostructure lasers

864

865

buried metal layers

528

Burstein-Moss effect

321

Burstein shift

139

С Capless annealing

779

Carbon doping

101

AlGaAs

101

102

device applications

111

112

GSMBE GaAs

103

lattice site location

108

109

MBE GaAs

102

10

MOCVD GaAs

101

102

MOMBE GaAs

103

thermal stability

109

111

359

368

carrier capture cross sections

780

This page has been reformatted by Knovel to provide easier navigation.

369

500

967

Index terms

Links

Carrier concentration

596

LPE GaAs

597

618

Carrier ionization coefficients

190

doping dependence

196

197

electric field dependence

191

orientation dependence

194

195

temperature dependence

195

196

carrier lifetime, see electron lifetime, hole lifetime carrier mobility, see electron mobility, hole mobility cathode sputtering of amorphous GaAs

155

cathodoluminescence efficiency

335

cathodoluminescence nonuniformity

587

593

Caughey-Thomas parameters

42

43

CBE GaAs

54

LVM frequencies

231

CBE growth

647

655

halogen assisted

632

cellular telephones

785

820

Chemical dry etching

416

714

AlGaAs

714

etch rates

714

etchants

715

InAs

714

716

chemical oxidation

477

478

chemical sensors

940

Chemically-assisted-ion-beam etching (CAIBE)

751

755

damage

758

759

etchants

756

profiles

755

757

rates

755

756

selectivity

758

trenching

757

758

close space vapour transport (CSVT)

639

CMOS integrated circuits

929

compensation

227

228

compensation mechanisms

781

782

compensation ratio

42 238

78

79

Compliance

17

21

22

adiabatic

21

isothermal

21 This page has been reformatted by Knovel to provide easier navigation.

22

438 117

132

968

Index terms

Links

Compliance (Continued) pressure dependence

22

temperature dependence

22

compliance tensor

162

compressibility

23

computer network communications

27

929

Conduction band discontinuity AlGaAs/GaAs

894

895

conduction band dispersion

148

149

Conduction band offset

545

898

see also conduction band offset

AlGaAs/GaAs

546

547

AlGaInP/GaInP

550

551

GaInP/AlGaAs

550

GaInP/AlInP

551

GaInP/GaAs

549

GaInP/InGaAsP

550

congruent melting temperature

557

consumer applications

949

contact resistivity

112

controlled atmosphere annealing

779

convergent beam electron diffraction

419

Cotterell atmospheres

401

CSVT growth

639

growth rate

639

Czochralski-growth

565

lattice parameter

9

Czochralski growth under As vapour pressure

580

550 558

780

579 581

D Data transmission links

948

DBS television receivers

815

Debye function

29

Debye temperature

29

temperature dependence

29

X-ray value

31

deep defect state characterization

358

deep level passivation

414

Deep level transient spectroscopy (DLTS)

264

constant capacitance

366

double

367

816

367

This page has been reformatted by Knovel to provide easier navigation.

561

591

969

Index terms

Links

Deep level transient spectroscopy (DLTS) (Continued) Laplace transform Deep levels

367

368

358

376

380

see also EL2 capture cross sections

377

charge exchange mechanisms

359

concentration

377

electron

376

emission energy

377

hole

376

occupancy

359

thermal emission

360

360

Defects (general) densities

371

380

381

383

685

line defects

371

372

point defects

373

685

volume defects

372

373

energy levels

244

346

347

380

394

EL2

341

electron damage

394

electron traps

380

383

394

395

397

hole traps

380

384

390

395

396

ion implantation effects

396

397

LPE GaAs

380

MBE GaAs

383

MOVPE GaAs

387

neutron irradiation effects

397

proton irradiation effects

397

VPE GaAs

385

see also dislocations

Defects (point)

defect structure

335

Deformation potentials, electrons

162

acoustic mode

132

AlAs

172

Ge

172

InP

172

interband

166

intervalley

162

intraband

167

intravalley

162

Si

172

341 178

169

This page has been reformatted by Knovel to provide easier navigation.

179

970

Index terms

Links

Deformation potentials, holes intravalley

177 177

degenerate conduction

105

Delta doping

71

carrier density

74

free carrier saturation

74

temperature dependence

74

demultiplexers

73 75

933

Density

3

dislocation effects

4

ideal

6

impurity effects

3

neutron irradiation effects

6

stoichiometry

5

temperature dependence

5

X-ray estimate

3

Dielectric constant

6

203

amorphous GaAs

155

infrared

203

temperature dependence

203

Dielectric function

201

far infrared

203

reststrahlen region

203

static

203

sum rules

202

204

204

dielectric resonators

827

828

Diffusion of impurities

431

438

Cr

442

443

noble metals

443

shallow acceptors

432

shallow donors

438

transition metals

443

442

443

435

436

Diffusion of impurities (coefficients) Be

434

C

109

110

Cr

442

443

D (deuterium)

414

Mn

433

S

439

Se

440

Si

438

This page has been reformatted by Knovel to provide easier navigation.

971

Index terms

Links

Diffusion of impurities (coefficients) (Continued) Sn

439

Te

440

Zn

432

433

diffusion of minority carriers

363

385

387

diffusion profiles

431

438

442

54

632

digital integrated circuits

785

802

dimers

456

dislocation arrays

371

372

dislocation climb

372

420

422

dislocation counting

327

dislocation density

572

579

619

620

dislocation glide

420

dislocation loops

420

dislocation mobility

421

422

Dislocations

335 567

371 568

372

399

edge

420

glide

420

partial

420

screw

420

shuffle

420

421

slip

371

372

digital epitaxy

distributed Bragg reflector

112

distributed control systems

948

440

822

421

420

Dopant-hydrogen complexes dissociation energies

411

412

vibrational modes

412

413

dopant solubility

562

563

dopant uniformity imaging

403

404

doping striations

328

403

dopant identification: see also photoluminescence spectra etc.

doping superlattice laser

404

73

double barrier resonant tunneling structures (DBRTS)

528

double crucible LEC

581

582

double heterostructure GaAs/AlGaAs lasers

620

673

DX centres

250

AlGaAs

251

see also Czochralski, LEC

253

This page has been reformatted by Knovel to provide easier navigation.

863

864

419

972

Index terms

Links

DX centres (Continued) capture barrier

251

chemical shift

253

emission barrier

251

ionization energy

251

photoionization threshold

251

pressure coefficient

253

253 253

254

E Effective bandgap narrowing

186

effective intrinsic carrier concentration

187

188

Einstein coefficient

136

139

EL2

259 368 946

260 376

273 385

336 560

343

348

activation energy

341

annealing effects

587

electronic structure

346

emission cross section

341

identification

344

Laplace DLTS spectra

368

metastability

341

348

photoquenching

337

338

pressure effects

351

stress dependence

353

354

symmetry

344

353

354

thermal ionization energy

341

thermal stability

348 21

22

EL2 imaging

347

399

elastic properties

16

elasto-optic coefficient

217

electrical compensation

781

782

Electrical resistivity

103

566

annealing effects

592

LEC GaAs

566

LT GaAs

696

MOCVD GaAs

103

MOMBE GaAs

103

Electroabsorption

220

electric field dependence

221

polarization dependence

221

221

This page has been reformatted by Knovel to provide easier navigation.

592

696

341 685

973

Index terms

Links

Electron concentration LEC GaAs

596

SI GaAs

596

VGF GaAs

596

Electron-cyclotron-resonance RF (ECR-RF) reactive ion etching

741

damage

742

etchants

747

process chemistries

744

profiles

742

rates

741

uniformity

742

Electron diffusion length

106

743 745 747 107

MOMBE GaAs

106

electron drift velocity

854

855

Electron effective mass

148

149

band edge

148

149

InAs

789

InGaAs

789

LPE GaAs

184

pressure dependence

184

temperature dependence

149

184

789

electron-hole pair creation energy

942

electron ionization rate

853

854

50

227

231

394

88

106

107

57

Electron-irradiation effects annealing

396

Electron lifetime

87

LPE GaAs

87

MBE GaAs

87

MOCVD GaAs

107

MOMBE GaAs

107

MOVPE GaAs

87

electron mean free drift length Electron mobility

942 41

54

55

70

71

78

bulk GaAs

46

47

carrier density dependence

42

44

Caughey-Thomas parameters

42

43

compensation ratio dependence

78

doping effects

41

42

55

81

electron irradiation effects

50

This page has been reformatted by Knovel to provide easier navigation.

50

78

47

50

59

54

974

Index terms

Links

Electron mobility (Continued) GaAs heterostructures

61

HB GaAs

46

higher energy conduction band states

43

ion-implanted GaAs

57

laser irradiation effects

50

LEG GaAs

46

47

566

LPE GaAs

48

49

59

MBE GaAs

54

55

minority carrier

44

45

modulation-doped GaAs

70

71

MOVPE GaAs

50

pressure dependence

59

60

substrate effects

52

55

substrate orientation effects

55

temperature dependence

41

43

two-dimensional electron gas effects

44

55

VPE GaAs

50

Electron mobility enhancement

44

61

73

delta doping

73

74

electron density dependence

63

64

GaAs heterostructures

61

illumination effects

64

65

impurity effects

63

64

interface roughness effects

63

spacer thickness dependence

62

temperature effects

64

electron traps

380

electron-two phonon interactions

172

electron velocity

244

electronic stopping

770

Electro-optic effect

870

linear

870

quadratic

870

81

50

128

74

63 383

66

electronic absorption bands

60

electro-optic integration

674

Electro-optical coefficients

217

218

clamped

217

218

linear

217

218

quadratic

217

temperature dependence

218

This page has been reformatted by Knovel to provide easier navigation.

394

395

618

975

Index terms

Links

Electro-optical coefficients (Continued) unclamped

217

wavelength dependence

218

electrorefraction

220

energy dispersive X-ray analysis

419

enthalpy

361

enthalpy of fusion

557

entropy

361

entropy of fusion

557

Epitaxial growth

601

218 221 362 362

see also specific growth technique abrupt interfaces

603

hardware

605

stoichiometry

602

substrates

601

602

surface morphology

602

603

Epitaxial lift-off

604

672

film curvature

674

gas out-diffusion

674

post-processing

672

673

pre-processing

672

673

selective etching

672

etch rate selectivity

672

Etch rates AlGaAs

672

673

CAIBE

755

756

chemical dry etching

416

613

crystallographic orientation dependence

707

714

ECR-RF RIE

741

747

FIB

756

ion beam etching

722

723

MIE

731

732

plasma etching

719

720

pressure dependence

715 736

716 747

RBIBE

759

760

RF-only RIE

725

726

736

RIBE

752

temperature dependence

708

709

712

747 wet etching

707

This page has been reformatted by Knovel to provide easier navigation.

714

739

740

720 753

725 756

726 760

714

720

976

Index terms

Links

eutectic point, Au-Ge

510

extended defects

335

Extinction coefficient

201

wavelength dependence

202

207

449

207

F Facets

459

Fano factor

942

faulted loops

423

Fermi level pinning

252 523

447

models

449

450

pressure dependence

540

FET-SEED optical receiver arrays

929

FET threshold voltage variation

588

fibre optical links

785

field effect controlled transferred electron device

832

flash evaporation of amorphous GaAs

155

flow-rate modulation epitaxy

603

focal plane arrays

907

912

Focused ion beam etching

751

752

damage

758

profiles

757

rates

756

929 156

focused ion beam sputtering

722

formation enthalpy

557

560

formation entropy

557

560

fractional quantum Hall effect

71

fractional superlattices

648

Franz-Keldysh absorption imaging

404

Franz-Keldysh effect

220

Free-carrier absorption

235

237

absorption mechanisms

237

238

n-GaAs

237

p-GaAs

240

Free-carrier absorption coefficient

405

238

acoustic phonon scattering

239

impurity scattering

239

magnetic field effects

239

optical phonon scattering

239

This page has been reformatted by Knovel to provide easier navigation.

755

466

522

977

Index terms

Links

Free carrier saturation

74

electronic

75

structural

75

76

free electron concentration

240

241

free exciton binding energy

152

fully encapsulated Czochralski (FEC) growth

581

furnace annealing

774

G GaAs/AlGaAs electron mobility

43

44

two-dimensional electron density

43

44

GaAs/AlGaAs QWIP

906

GaAs/AlGaAs SQW

862

54

GaAs on Si microstructure

424

GaAsP atomic layer epitaxy

633

GaInAs/AlInAs/InP HBTs

822

GaInAsP selective area epitaxy

633

GaInP atomic layer epitaxy

633

gamma ray detection

943

GaP atomic layer epitaxy

633

gas pressure microsensors

936

gas sensors

939

Ge deformation potentials

172

geothermal applications

948

gettering

616

618

Gibbs free energy

361

362

graded refractive index SCH (GRINSCH) structures

864

865

grain boundary diffusion

506

510

GRINSCH SQW laser

927

944

946

947

937

499 521

526

614

619

growth processes: see specific growth techniques Growth rate: see also MBE growth rate etc. ALE

668

669

CSVT

639

low temperature MBE

680

LPE

610

613

MBE

656

658

MOMBE

664

666

MOVPE

648

This page has been reformatted by Knovel to provide easier navigation.

667

978

Index terms

Links

Growth rate: see also MBE growth rate etc. (Continued) VPE

625

Grüneisen parameter

23

GSMBE GaAs

103

GSMBE growth

103

carbon doping

655

103

Gunn diodes

830

H Hall factor

99

117

131

Hall mobility

42 59 121

43 91 131

46 95 566

50 98

51 117

Hall sensors

933

934

QW

934

sensitivity

934

temperature coefficient

934

820

950

958

halogen transport assisted epitaxy

630

HB GaAs EL2 concentration

341

electron mobility

46

growth

565

horizontal zone melting

573

microprecipitates

335

modeling

573

modified 2T-HB system

573

residual strain

574

stoichiometry control

574

HBT

342 572

111 959

112

C-doped

111

112

circuit applications

821

822

collector current

187

digital switching speed

821

high temperature operation

950

integrated circuits

927

leakage current

950

low frequency noise performance

821

microwave output power

821

microwave power gain

820

radiation hardness

822

958

821

This page has been reformatted by Knovel to provide easier navigation.

959

979

Index terms

Links

HBT (Continued) reliability

822

HEMT

70

71

673

788

824

829

954

957

fabrication

814

815

high temperature operation

957

958

InGaAs

789

790

812

813

integrated circuits

817

927

928

low noise

792

794

815

operation

814

pseudomorphic

813

structure

811

unity current gain cutoff frequency

791

Heterostructures

811

816

816 792

861

characterization

861

high resolution electron microscopy

418

high temperature applications

948

high temperature electronic techniques

952

high temperature MBE growth

655

862

see also MBE growth hole drift mobility

99

hole drift velocity

854

855

Hole effective mass

149

184

heavy hole

149

light hole

149

pressure dependence

184

185

hole ionization rate

853

854

Hole lifetime

135

LPE GaAs

138

MBE GaAs

141

MOCVD GaAs

140

hole mean free drift length

942

Hole mobility

84 117

91 131

98

91

117

126

127

133

134

92

93

118

126

119

C-doped GaAs

118

compensation ratio dependence doping effects Ge-doped GaAs ion-implanted GaAs

139

85 123

Be-doped GaAs carrier concentration dependence

185

120 94 This page has been reformatted by Knovel to provide easier navigation.

95

103

980

Index terms

Links

Hole mobility (Continued) LEG GaAs

95

LPE GaAs

93

100

MBE GaAs

91

118

120

Mg-doped GaAs

120

minority carrier

99

123

MOCVD GaAs

92

93

104

MOMBE GaAs

92

104

118

MOVPE GaAs

93

118

pressure dependence

100

temperature dependence

84 104

Zn-doped GaAs

119

hole traps

380

85 127

93

98

99

384

390

395

396

415

416

651

homojunction transistors collector current

187

horizontal boat (HB) growth

565

see also Bridgman growth, HB GaAs horizontal zone melting

573

see also HB GaAs hot carrier luminescence

299

300

409

410

see also photoluminescence Hydrogen incorporation CVD

416

epitaxial growth

415

etching

416

implantation

410

MOCVD

651

plasma

409

410

112

409

Hydrogen passivation deep levels

414

defects

409

shallow dopants

410

416

Hydrogen-dopant complexes dissociation energies

411

412

vibrational modes

412

413

167

177

hydrostatic deformation potential

I IC device fabrication yield

593

impact ionization rates

853

854

This page has been reformatted by Knovel to provide easier navigation.

178

981

Index terms

Links

IMPATT diodes

851

avalanche frequency

852

avalanche multiplication rate

851

double-drift

856

drift length

852

efficiency

857

output power

856

857

Read-type

855

856

saturated drift velocity

851

single-drift

855

tunnel generation rate

851

impurity concentrations

373

856

see also defects Impurity energy levels

244

248

pressure coefficient

248

249

transition metal

248

249

574

576

249

see also defects, deep levels etc.

impurity hardening effect impurity identification: see IR absorption photoluminescence spectra etc. InAs atomic layer epitaxy

633

bandgap

789

electron effective mass

789

etching

714

InAsP, atomic layer epitaxy

633

InGaAs atomic layer epitaxy

633

bandgap

789

electron effective mass

789

etching

711

MBE growth

657

selective area epitaxy

633

thermal conductivity

35

InGaAs/GaAs MQW

863

InGaAs/GaAs quantum well lasers

865

866

InGaAs/GaAs quantum well piezoelectric constant

224

InGaAs/GaAs SQW

862

Ingot annealing

585

defect control

587

588

This page has been reformatted by Knovel to provide easier navigation.

581

982

Index terms

Links

Ingot annealing (Continued) device performance

588

electrical properties

585

multiple-step

585

optical properties

585

three-step

585

586

588

587

InP atomic layer epitaxy

633

chlorine assisted MOVPE growth

631

deformation potentials

172

hydride growth

632

LEG growth

580

selective area epitaxy

633

InP/InGaAs photodiode

673

inter-conduction band absorption

241

581

see also optical absorption interface properties semiconductor/air

876

877

interface recombination velocity GaAs/AlGaAs

618

interface state density

447

449

478

484

540

507

510

511

517

473

478

Interface structure Ag/GaAs

531

annealing effects

506 518

Au/GaAs

505

Au-Ge/GaAs

510

Au-Te/GaAs

507

CoSi2/GaAs

519

epitaxial Al/GaAs

499

GaAs/AlGaAs

425

GaAs/InGaAs

425

ion bombardment

507

MoSix/GaAs

518

orientation relationships

499

oxide/GaAs

466

silicide/GaAs

516

SiO2/GaAs

471

substrate/LPE interface degradation

279

280

TaSix/GaAs

517

518

thermal stability

503

516

TiWSix/GaAs

519

525

472

This page has been reformatted by Knovel to provide easier navigation.

983

Index terms

Links

Interface structure (Continued) VSix/GaAs

518

519

WSix/GaAs

516

517

interstitial loops

420

interstitials

559

560

561

686

765

776

see also defects Inter-valence band absorption

241

see also optical absorption n-type GaAs

241

p-type GaAs

241

intervalley deformation potentials electrons intracentre absorption of localized centres

162

169

245

246

see also optical absorption,EL2 etc. intracentre transitions

244

see also above Intravalley deformation potentials electrons

162

holes

177

ion beam assisted etching

751

Ion beam etching

722

damage

722

profiles

722

rates

722

723

ion-beam induced oxidation

492

493

ion beam milling

722

Ion implantation

423

424

activation energy

765

annealing effects

765

applications

768

electrical activation

774

peak carrier concentrations

776

777

profiles

767

768

777

Ion-implanted GaAs defect energy levels

397

electrical compensation

781

782

electron mobility

57

hole mobility

94

95

peak carrier concentration

776

777

rapid thermal annealing

774

775

This page has been reformatted by Knovel to provide easier navigation.

777

984

Index terms

Links

Ion ranges in GaAs

770

channeling

772

energy dependence

771

ionization coefficient ratio

191

ionized impurity concentration

195

235

342

399

231

394

78

IR absorption

227

free carrier

235

impurity

227

quenching

350

IR absorption coefficients

237

IR detectors

906

detectivity

192

351

915

IR hot electron transistor

914

IR imaging

399

IR imaging cameras

915

915

Irradiation effects electron

50

laser

51

neutron

6

proton

227 397

397

isoconcentration technique

433

J JFET

768

949

see also MESFET, etc. high temperature operation

949

K Kerr effect

217

Kramers-Kronig relationship

202

L LANs

925

laser assisted oxidation

486

laser diode drivers

925

laser diode transmitter modules

925

Laser heterostructures

863

quantum well

864

laser irradiation effects

929

865

51

laser scanning microscopy

403

This page has been reformatted by Knovel to provide easier navigation.

985

Index terms

Links

Lasers AlGaAs/GaAs QW

887

buried heterostructure

864

doping superlattice

865

73

double heterostructure GaAs/AlGaAs

620

673

GRINSCH SQW

927

InGaAs/GaAs QW

865

MQW

864

quantum box

866

867

quantum well heterostructure

864

887

SQW

864

865

tunneling injection

866

VCSEL

112

zero-threshold

884

lattice frequency

227

lattice hardening

371

568

lattice imaging

418

419

Al/GaAs

500

502

GaAs/Ge

658

922

GaAs/Si

922

NiAl/GaAs

528

863

864

567

866

Lattice mismatch

Lattice parameter

3

Al

8

499

dislocation effects

4

11

doping effects

9

10

facet effect

8

homogeneity effects

8

ideal

12

impurity effects

5

stoichiometry

11

surface damage effects

8

temperature dependence

9

X-ray value

9

lattice relaxation LDD structures

12

250 768

LEG GaAs defect density

371

dislocations

371

372

420

EL2 concentration

341

342

566

electrical resistivity

566

This page has been reformatted by Knovel to provide easier navigation.

568

986

Index terms

Links

LEG GaAs (Continued) electron Hall mobility

566

electron mobility

46

hole mobility

95

lattice parameter

47

9

LVM frequencies

229

native defect structure

335

near-IR mapping

399

photoluminescence mapping

328

photoluminescence spectra

273

LEG growth

567

565

579

As injection

579

580

Czochralski under As vapour pressure

580

581

dislocation reduction

567

568

double crucible

581

582

equipment

565

566

FEC

581

high pressure

566

579

InP

580

581

low pressure

566

583

magnetic field effects

582

583

purity

567

uniformity

568

LEVB growth

575

level shifters

957

Light emitting diodes

673

874

efficiency

874

875

fabrication

878

Fresnel losses

877

879

luminosity

875

876

materials

880

photon recycling

883

884

resonant cavity

882

883

single mode

884

structures

878

super-luminescent

880

wavelength ranges

880

light scattering tomography

403

line defects

371

576

372

see also dislocations, etc.

This page has been reformatted by Knovel to provide easier navigation.

987

Index terms

Links

lineage

335

372

401

48

49

615 232

see also dislocations, etc. liquid phase electro-epitaxy Localized vibrational mode (LVM) frequencies

227

calibration

227

228

hydrogen-related

227

231

stretch mode

231

wag mode

231

Lorentz number

33

loss tangent

825

826

Low noise amplifiers

793

794

gain

794

low noise microwave devices

815

816

annealing effects

685

695

breakdown electric field

696

carrier lifetime

696

EL2 concentration

342

electrical resistivity

696

Low temperature MBE (LTMBE) GaAs

697

electron mobility

55

growth dynamics

679

MESFETs

689

microstructure

424

opto-switches

693

point defects

684

stoichiometry

684

ultrafast trapping

696

697

689

690

295

296

low temperature MBE (LTMBE) GaAs buffer low temperature MBE (LTMBE) GaAs photoluminescence spectra Low temperature MBE growth

679

AlGaAs

681

As:Ga flux ratio effects

679

growth rate

680

temperature dependence

680

low temperature MOCVD GaAs

700

low temperature oxidation

490

682

491

LPE GaAs carrier concentration

618

defect energy levels

380

electron effective mass

184

This page has been reformatted by Knovel to provide easier navigation.

988

Index terms

Links

LPE GaAs (Continued) electron lifetime

87

electron mean free drift length electron mobility

942 48

49

hole lifetime

138

139

hole mean free drift length

942

hole mobility

93

100

hole traps

618

photoconductivity spectra

264

photoluminescence spectra

278

shallow donors

264

solution bakeout

280

281

substrate/interface degradation

279

280

608

646

AlGaAs

614

619

boats

611

chemical reaction

646

computer control

612

doping

615

furnaces

612

gettering

616

growth rate

610

growth temperature

614

interface abruptness

619

layer thickness

619

sliding boat technique

609

solvents

609

substrates

612

LPE growth

LPE layer assessment

613

611

617

LTMBE, see low temperature MBE luminescence efficiency

295

Luttinger parameters

149

Lydane-Sachs-Teller relation

203

308

M Mach-Zehnder interferometer

939

Magnetic sensors

933

Hall devices

933

magnetotransistors

935

micro-Hall devices

935

uses

933

934

This page has been reformatted by Knovel to provide easier navigation.

59

60

120

618

614

619

618

989

Index terms

Links

Magnetoluminescence VPE GaAs Magnetotransistors

284

285

284

285

935

sensitivity

935

Magnetron-enhanced reactive ion etching

725

731

732

739

740

damage

732

etchants

739

740

rates

731

732

739 105

mass flow sensors

936

Matthiessen’s rule

91

93

carbon doping

102

103

defect energy levels

383

740

MBE GaAs

delta doping

73

EL2 concentration

342

electron lifetime

87

electron mobility

54

hole lifetime

141

hole mobility

91

hydrogenation

295

photoconductivity spectra

266

photoluminescence spectra

291

piezoelectric constant

55 118 267

54

shallow donors

266

267

surface structure

455

456

458

102 663

103 679

632

carbon doping

102

103

chemical reaction

646

growth conditions

657

658

growth rate

656

658

growth system

655

halogen assisted

632

high temperature

655

InGaAs

657

in-situ analysis

660

kinetic analysis

659

low temperature

679

solid source

655

substrate temperature

658

thermodynamic analysis

660

thermal expansion coefficient MBE growth

26

This page has been reformatted by Knovel to provide easier navigation.

646

655

990

Index terms

Links

MBE growth (Continued) medical imaging

942

melting point

28

pressure dependence

36

MESFET

689

36

557

565

785

799

836

949

836 953

active channel preparation

785

backgating

689

690

buffer layers

689

690

depletion mode

543

803

808

809

enhancement mode

542

543

803

807

epitaxial

796

795

epitaxial lift-off

673

fabrication

786

high temperature characteristics

949

InGaAs

788

intrinsic speed

788

789

ion implantation

768

790

795

796

large-signal models

799

leakage currents

949

low noise

790

low temperature grown materials

689

power

788

precipitates

373

surface passivation

690

691

topology

787

788

transconductance

804

805

unity current gain cutoff frequency

785

789

804

805

542 926

543 928

673 929

804

836

389

686

335

373

402

403

MESFET integrated circuits backgating

836

large signal modeling

799

optical receivers

804

sidegating

448

metastability

348

microcontrollers

948

microdefects

335

Micro-Hall devices

935

sensitivity

935

temperature coefficient

935

micro-machining

628

microprecipitates

330 568

787

794

This page has been reformatted by Knovel to provide easier navigation.

991

Index terms

Links

microprocessors

948

microsensors

936

937

microtwins

420

422

microwave devices

816

microwave ovens

949

microwave permittivity

824

microwave phased-array radar

929

migration enhanced epitaxy

603

millimetre wave devices

816

Miniband transport MQWIR detectors

914

detectivity

423

825

914

Minority carrier lifetime electrons holes

363

385

87

88

387

135

Minority carrier mobility doping dependence

83

84

126

electron

44

45

81

hole

99

123

temperature dependence

84

85

minority carrier transient spectroscopy

368

MISFET

447

misfit dislocation structure

500

MITATT diode

852

853

mixers

822

829

MMIC

785 845

790 929

fabrication

786

787

low noise

790

power

795

MMWIC

824

substrates

99

127

822 938

828 948

691

796

824

825

carbon doping

101

102

conductivity-type conversion

287

electrical resistivity

103

electron lifetime

107

hole concentration

103

hole lifetime

140

hole mobility

92

93

photoconductivity spectra

265

266

photoluminescence spectra

286

MOCVD GaAs, see also MOVPE GaAs

This page has been reformatted by Knovel to provide easier navigation.

103

832

992

Index terms

Links

MOCVD GaAs (Continued) shallow donors

265

thermal stability

109

MOCVD growth

266

101

102

C doping

101

102

unintentional H doping

651

651

see also MOVPE growth

MODFET

70

811

927

928

70

71

modulator integrated circuits

928

929

modulators

869

870

molecular layer epitaxy

632

see also HEMT integrated circuits modulation doping

MOMBE GaAs

54

carbon concentration

664

carbon doping

103

compensation

227

electrical resistivity

103

electron diffusion length

106

electron lifetime

107

hole concentration

103

hole mobility

92

LVM frequencies

231

photoconductivity spectra

266

shallow donors

266

thermal stability

10

MOMBE growth

655

doping sources

666

growth rate

664

halogen assisted

632

sources

664

103

663 666

667

484

monolithic integrated circuits

926

monolithic phase shifter

929

MOS structures

478

479

MOSFET

479

785

MOVPE GaAs, see also MOCVD GaAs defect energy levels

387

electron lifetime

87

electron mobility

50

hole mobility

93

118

118

This page has been reformatted by Knovel to provide easier navigation.

654

993

Index terms

Links

MOVPE GaAs (Continued) photoluminescence spectra MOVPE growth

286 630

634

643

see also MOCVD growth chemical reactions

645

chlorine-assisted

631

flow dynamics

644

growth rate

648

precursors

648

reactor design

644

safety

648

starting compounds

646

surface chemistry

647

MSM detectors

699

multifunction circuits

822

multiple quantum well (MQW) lasers

864

632

645

648

863

893

341

376

559

376

559

see also lasers Multiple quantum wells (MQW) energy levels

862 893

multiple wafer growth

645

multiplexers

933

N Native defects

335

see also defects, deep levels, EL2, etc extended

335

micro

335

point

336

341

401

402

near bandedge absorption imaging see also reverse contrast imaging near-IR mapping

399

see also above negative conductance

827

negative photoconductivity

159

neutron irradiation effects

6

828 397

see also irradiation effects Neutron transmutation doped GaAs

338

photoconductivity spectra

267

photoluminescence spectra

304

shallow donors

267

nonpolar optical deformation potential

132

This page has been reformatted by Knovel to provide easier navigation.

994

Index terms

Links

nuclear particle detectors

942

nuclear stopping

770

O Ohmic contacts to GaAs

112

424

425

alloying treatments

513

Au

505

507

534

Au-Ge

510

511

513

Au-Mn

512

Au-Te

507

contact resistivity

513

high temperature

958

ion implantation effects

513

Ni-Au-Ge

511

Pd-Au-Ge

512

Pt-Au-Ge

512

WSix

517

oil industry electronics

958

949

OMVPE, see MOVPE and MOCVD one-sided diode I-V characterization technique

126

Optical absorption

201

202

342

343

amorphous GaAs

156

159

band edge

204

electric field effects

220

221

SI GaAs

342

343

wavelength dependence

207

222

optical data links

807

929

optical deep level transient spectroscopy (ODLTS)

368

optical deformation potential

165

optical displacement sensor

933

Optical functions

201

band edge

204

infrared

203

ultraviolet

205

visible

205

wavelength dependence

207

X-ray

205

166

204

207

160

179

207 204

Optical gap amorphous GaAs

156

optical interconnects

929

157

This page has been reformatted by Knovel to provide easier navigation.

159

160

220

995

Index terms

Links

optical phonon frequencies

203

optical pressure sensors

936

Optical receivers

803

bit error rate

807

optical sensors

933

optical transient current spectroscopy (OTCS)

367

optoelectronic device heterostructures

861

Optoelectronic integrated circuits

786

804

fabrication

786

787

laser-diode transmitter modules

925

optical receivers

804

376 925

optoelectronic integration

674

Opto-switches

693

927

operation

694

695

oscillators

822

826

832

833

Oxidation

463

464

470

477

489

487

489

490

477

482

487

ad-atom effects

491

atomic oxygen beam

491

492

ion-beam induced

492

493

laser assisted

486

low temperature

465

490

491

native oxides

463

464

471

optically induced

464

overlayer effects

491

plasma

464

465

482

thermal

464

470

thermodynamics

464

UV/ozone

464

465

vapour pressure

464

465

wet

477

oxidation promoters

473

oxidation rate

472

473

oxide layer growth

463

464

Oxide layer structures

463

amorphisation

463

composition

465

crystallization

463

466

P Pagers

785

parabolic band approximation

889

This page has been reformatted by Knovel to provide easier navigation.

996

Index terms

Links

Passivation

112

227

268

295

409

447 479

463 651

464

467

471

715

722

defect

409

hydrogenation

112

409

651

surface

324

447

493

pattern formation

629

632

707

PC-LEC growth

568

see also Czochralski growth, LEG etc. Peierls barrier

421

Penn model

155

Permittivity

824

frequency dependence

825

826

temperature dependence

825

826

persistent hole conductivity

260

persistent photoconductivity

250

Phase diagrams

608

260

see also thermodynamics Ga-As

559

Ga-As-O

465

phase equilibria

557

phase extent

561

466

see also thermodynamics phonon mobility limit

78

Photocapacitance quenching recovery rate

349

350

350

photocathodes

620

photochemical washing

477

Photoconductive switches

695

dark current

696

responsivity

698

time response

696

Photoconductivity

259

699 260

doping dependence

259

epilayers

262

far-IR

262

IR

259

260

line narrowing

262

263

LPE GaAs

264

magnetic field effects

262

MBE GaAs

266

267

This page has been reformatted by Knovel to provide easier navigation.

262

997

Index terms

Links

Photoconductivity (Continued) MOCVD GaAs

265

266

MOMBE GaAs

266

notch effect

263

NTD GaAs

267

pressure effects

262

263

quenching

269

260

shallow donors

262

SI bulk GaAs

259

260

temperature dependence

259

262

VPE GaAs

264

265

Photodetectors

933

266

349

see also photodiodes dark current

933

integrated circuits

927

photoresponsivity

933

photodiodes

673

photoelectric-Compton crossover

942

photoenhanced oxidation

486

photo-induced transient spectroscopy (PITS)

367

photoionisation absorption spectra

246

Photoluminescence

928 804

866

273

298

308

320

carrier concentration dependence

321

322

doping effects

320

epitaxial layers

320

group II doped

298

group IV doped

302

group VI doped

308

growth rate effects

294

intensity

322

KP lines

291

LEG GaAs

273

LPE GaAs

278

LTMBE GaAs

295

296

MBE GaAs

291

303

MOCVD GaAs

286

NTD GaAs

304

pressure dependence

287

294

299

308

315

316

rare earth element doped

317

room temperature

320

SI GaAs

273

309 327

This page has been reformatted by Knovel to provide easier navigation.

311

998

Index terms

Links

Photoluminescence (Continued) stress effects

299

substrates

320

surface damage effects

274

temperature dependence

287

transition metal doped

310

twofold excitation modulated

274

VPE GaAs

282

zero-phonon lines

310

304

294

311

313

Photoluminescence homogenization

587

593

AlGaAs

323

324

defect monitoring

323

324

doping measurement

322

323

surface recombination effects

324

transients

324

126

320 326

excitation sources

325

sample preparation

325

SI GaAs

328

substrates

327

328

photoluminescence quenching

273

296

photon recycling

87 920

88

photon recycling LED

883

884

photoquenching

337

338

photoreceivers

927

928

photovoltaic effect

159

Piezoelectric constants

223

326

see also optical absorption, photoluminescence, etc.

224

54

Piezoelectric sensors

125

325

collection and detection components

MBE GaAs

313

299

photoluminescence decay time

Photoluminescence mapping

299

937

938

acoustic

938

pressure

937

938

temperature

937

938

piezo-optic effect

164

plasma anodisation

482

483

plasma edge

240

241

plasma energy

236

see also plasma oxidation

This page has been reformatted by Knovel to provide easier navigation.

350

136

999

Index terms

Links

Plasma etching

719

damage

721

etchants

720

profiles

719

rates

719

selectivity

721

plasma frequency

235

Plasma oxidation

482

720 236

240

241

373

559

see also plasma anodization ECR-plasma

483

surface structure

483

Plasmon-phonon coupling

235

n-type GaAs

236

p-type GaAs

236

plastic deformation

421

236

422

see also dislocations Pockelseffect

217

Point defects

336

341

see also defects, impurities, EL2, etc. chemical impurities

373

deep levels

376

low temperature GaAs

684

native defects

376

559

position sensitive detectors

946

947

power amplifiers

794

power devices, HEMT

816

Power MESFET

794

associated gain

794

output power

794

power added efficiency

794

power sensors

938

Precipitates

591

amorphous

373

Pressure sensors

935

high pressure

936

optical

936

prismatic loops

371 936

423

see also dislocations proton irradiation effects

397

proximity annealing

774

pseudomorphic growth

604

775

This page has been reformatted by Knovel to provide easier navigation.

779

684

1000

Index terms

Links

pseudomorphic HEMT

813

pseudomorphic layer growth

861

p-type conversion

592

816

Q Quantum box laser

866

867

quantum-confined Stark effect

869

870

quantum confinement effect

267

quantum dots

426

quantum size effects

887

quantum well electro-optic phase modulator

870

459

604

605

Quantum well growth CBE

863

MBE

862

MOVPE

863

Quantum well heterostructure lasers

864

887

band filling

902

903

carrier collection

899

design rules

898

emission wavelength

896

897

MQW

864

893

900

phonon assisted stimulated emission

903

square well model

888

SQW

864

865

893

868

906

Quantum well infrared photodetectors bound-to-bound

908

bound-to-continuum

908

bound-to-quasibound

908

dark current

907

908

detectivity

911

912

imaging arrays

912

light coupling

912

noise

910

p-doped

915

responsivity

908

Quantum wells

862

863

band bending

888

889

Debye length

888

density of states

894

energy levels

887

915

911

895

This page has been reformatted by Knovel to provide easier navigation.

887

900

648

1001

Index terms

Links

Quantum wells (Continued) shallow donor spectra

266

square well model

888

wave functions

891

892

897

426

459

604

Radar applications

785

929

Radiation detectors

942

quantum wires

267

605

648

139

140

R

bulk

944

charge collection efficiency

944

energy resolution

943

epitaxial

942

radiation length

942

radiative recombination Radical-beam-ion-beam etching (RBIBE)

944

946

87

88

135

751

759

damage

761

etchants

760

profiles

759

rates

759

760

selectivity

759

761

rapid thermal annealing

774

775

Reactive-ion-beam etching (RIBE)

751

damage

754

etchants

753

plasma sources

752

profiles

754

rates

752

selectivity

754

Reactive ion etching (RIE)

755

725

ECR-RF

741

magnetron

725

731

RF-only

725

736

Read-type devices

855

Recombination mechanisms

135

732

Auger

137

138

radiative

135

139

140

Shockley-Read-Hall

137

139

140

recycling cofactor

136

reflectivity coefficient

240

241

This page has been reformatted by Knovel to provide easier navigation.

1002

Index terms

Links

Refractive index

201

202

204

157

159

207

218

537

538

221 amorphous GaAs

156

band edge

204

electric field effects

221

infrared

204

temperature dependence

204

wavelength dependence

207

regrowth after epitaxial lift-off

673

remote plasma etching

716

resonant cavity LED

882

883

responsivity

698

699

reverse contrast imaging

401

402

RF-only reactive ion etching (RF-RIE)

725

736

damage

728

729

etchants

736

morphology

727

process chemistries

729

profiles

726

727

rates

725

726

selectivity

727

728

Richardson velocity

218

736

187

S Sacrificial overlayers

467

SAGM avalanche photodiode

867

SAM avalanche photodiode

867

satellite communications systems

811

SAW sensors

938

scattering mechanisms

123

Schlieren imaging techniques

403

132

133

Schottky barrier heights Ag/GaAs

530

Al/GaAs

499

521

annealing effects

523 540

524 542

532

Au/GaAs

505

507

534

clean surfaces

522

531

535

541

etched surfaces

521 539

522

530

534

high temperature

958

959

This page has been reformatted by Knovel to provide easier navigation.

535

1003

Index terms

Links

Schottky barrier heights (Continued) high temperature stability

526

hydrogenation effects

535

ideality factor

517

521

526

530

534

527

528

533

535

531

532

539 interfacial layer effects

525

Mn/GaAs

533

MoAlx/GaAs

526

pressure dependence

540

Pt/GaAs

539

surface reconstruction effects

525

526

TiPtAu/GaAs

539

958

W/GaAs

541

WNx/GaAs

543

WSix/GaAs

517

WSiN/GaAs

543

segregation coefficients, see various growth techniques Selective area epitaxy

630

AlGaAs

633

GaInAsP

633

InGaAs

633

InP

633

self aligned gate (SAG) process

542

543

959

self-compensation

562

781

782

870

928

929

total internal reflection

876

877

critical angle

876

877

see also various growth techniques self electro-optic effect device Semiconductor/air interface

Sensors

933

bio

939

gas

939

magnetic

933

mass flow

936

optical

933

piezoelectric

937

938

93

8

power pressure

935

separate confinement heterostructure

864

shallow donor identification

262

shallow dopant passivation

410

This page has been reformatted by Knovel to provide easier navigation.

1004

Index terms

Links

shear deformation potential shear modulus

165

177

18

Shockley-Read-Hall recombination Shubnikov-de Haas effect

137

139

40

74

SI GaAs, see Czochralski, LEG growth, etc. absorption quenching

350

351

carrier concentrations

596

597

defect density

371

dislocations

372

electron mobility

46

infrared absorption spectra

342

infrared imaging

399

native defect structure

335

photoconductivity spectra

259

photoluminescence mapping

328

photoluminescence spectra

273

Sidegating

836

single mode LED

884

single quantum well lasers

864

Single quantum wells

862

47

260

865

see also quantum wells energy levels

893

sliding boat technique

609

small lattice relaxation model

251

smart pixels

929

Solar cells

620

918

concentrator solar cells

921

conversion efficiency

918

irradiation damage

922

modeling

919

920

space applications

918

921

stacked cells

922

923

terrestrial use

918

920

thin film

921

922

space applications

816

918

space communication

929

948

22

27

922

solubility of dopants, see various growth techniques

Specific heat classical

29

liquid

28

solid

28 This page has been reformatted by Knovel to provide easier navigation.

29

921

922

1005

Index terms

Links

Specific heat (Continued) temperature dependence

27

Sputter etching

722

damage

722

profiles

722

rates

722

selectivity

722

723

square well model

888

901

Stacking faults

420

422

867

868

16

179

adiabatic

16

17

doping dependence

18

isothermal

17

pressure dependence

17

temperature dependence

19

423

see also transmission electron microscopy staircase superlattice Stiffness constants

223

18

stoichiometry, deviations

559

560

strain tensor

162

Stranski-Krastanow growth

499

stress tensor

162

163

substrate/LPE interface degradation

279

280

Substrates

601

602

epi-ready

602

656

etch pit density

602

step structure

602

surface backgating

841

passivation

843

surface cleaning

714

Surface passivation

324

493

adverse effects due to interface states

447

448

Fermi level pinning

449

interface state density

447

non-traditional methods

451

selection criteria

447

612

see also Czochralski, HB, LEG growth, etc.

Surface passivation of devices

843

690

691

690

691

surface processing

466

467

surface reconstructions

455

647

LT GaAs

This page has been reformatted by Knovel to provide easier navigation.

447

679

680

1006

Index terms

Links

Surface structure

455

(100) surface

455

(110) surface

457

(111) surface

458

459

high index surfaces

459

460

phase diagrams

455

458

Telecommunication networks

785

929

thermal conduction current

949

T

Thermal conductivity

32

alloying effects

34

doping effects

33

InGaAs

35

temperature dependence

32

Thermal diffusivity

854 34

34

35

temperature dependence

34

35

Thermal expansion coefficient

8

9

doping effects

25

epitaxial

26

low temperature

26

stoichiometry effects

25

temperature dependence

23

Thermal oxidation

22

470

crystal orientation dependence

470

dopant effects

473

high temperature

472

low temperature

470

temperature dependence

472

thermal oxide films

471

473

Thermal resistivity AlGaAs

936

Thermal stability

109

111

see also low temperature GaAs MOCVD GaAs

109

MOMBE GaAs

109

thermally stimulated current (TSC) spectroscopy

376

Thermodynamics

557

three terminal dual colour optical emitter

902

time-of-flight technique

369

370

82

126

81

This page has been reformatted by Knovel to provide easier navigation.

1007

Index terms

Links

Transferred electron devices

830

AlGaAs/GaAs

832

efficiency

831

832

GaAs

831

834

hot electron injection

831

832

InP

831

834

output power

831

832

phase noise

833

planar structure

832

sandwich structure

830

tuning range

833

transferred electron effect

830

transimpedance amplifiers

804

transistor breakdown

953

Transmission electron microscopy

418

833

809

bulk GaAs

419

films

424

heterostructures

425

426

metal/GaAs system

424

425

specimen preparation

419

trimers

458

tunneling injection laser

866

TUNNETT diode

853

Two-dimensional electron gas

44 71

459

61 890

electron mobility

61

modulation doping

70

71

two-dimensional hole gas

70

71

twofold excitation modulated photoluminescence spectroscopy

274

type I to type II behaviour crossover

547

type II alignment in narrow quantum wells

547

type conversion

434

67

890

U Unity gain transistor cut-off frequency technique up conversion

82

126

293

Urbach energy amorphous GaAs

157

158

UV/ozone oxidation

487

489

This page has been reformatted by Knovel to provide easier navigation.

490

68

70

1008

Index terms

Links

V Vacancies

338 687

339

355

895

898

see also deep levels, defects, etc. vacancy generation parameter

442

vacancy loops

420

valence band discontinuity AlGaAs/GaAs

894

see also below Valence band offset

545

AlGaAs/GaAs

547

AlGaInP/GaInP

550

GaInP/AlInP

551

GainP/GaAs

549

550

Van der Waals bonding

673

674

vapour levitation epitaxy

50

629

551

vapour mixing epitaxy

629

varactors

828

829

833

Varshni coefficients

148

151

204

576

VB GaAs dislocation distribution

372

growth

565

575

photoluminescence mapping

327

328

thermal environment control

576

wetting control

576

VCZ growth

568

see also Czochralski growth Vegard’s law

10

velocity overshoot

66

Vertical boat grown GaAs photoluminescence mapping

854

565 328

vertical boat growth

565

vertical zone melting

575

329

Vertical gradient freeze (VGF) GaAs dislocation distribution electron mobility

372 47

growth

565

575

576

photoluminescence mapping

327

330

331

magnetic field effects

576

thermal environment control

576

wetting control

576 This page has been reformatted by Knovel to provide easier navigation.

559

686

1009

Index terms

Links

vertical-cavity surface-emitting lasers

112

926

complexes

412

413

stretching

412

413

wagging

413

927

Vibrational modes of dopant-hydrogen

VLSI

836

voltage comparators

821

voltage controlled oscillator

833

volume defects

372

847

849

373

VPE GaAs, see also MOVPE GaAs etc. defect energy levels

385

EL2 concentration

342

electron mobility

50

magnetoluminescence

284

285

photoconductivity spectra

264

265

photoluminescence spectra

282

shallow donors

264

VPE growth

265

625

chloride

626

halogen transport

625

hydride

626

low pressure

627

low temperature

627

metal-trichloride source

629

vapour levitation epitaxy

629

vapour mixing epitaxy

629

628

630

W Wafer annealing

586

device performance

593

electrical properties

592

multiple-step

591

optical properties

593

waveguide coupling

673

Wet etching

707

AlGaAs

711

electrolytic

711

etch rates

707

etchants

708

InGaAs

711

non-electrolytic

707

591 593

712

This page has been reformatted by Knovel to provide easier navigation.

629

959

1010

Index terms

Links

Wet etching (Continued) profile

707

selectivity

710

trenching

710

711

Wet oxidation

477

anodic

478

479

chemical

477

478

X X-ray astronomy

942

X-ray detection

942

946

Z Z-contrast imaging

419

Zener tunneling rate

853

854

82

126

zero-phonon lines

245

246

zero-threshold laser

884

zero field time-of-flight technique

This page has been reformatted by Knovel to provide easier navigation.

947