P R O P E R T I E S O F

A m o r p h o u s S i l i c o n

a n d i t s A l l o y s

E d i t e d b y

T I M S E A R L E

U n i v e r s i t y o f S h e f f i e l d , U K

IEE I N S P E C

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

1998: The Institution of Electrical Engineers

This publication is copyright under the Berne Convention andthe Universal Copyright Convention. All rights reserved. Apartfrom any fair dealing for the purposes of research or privatestudy, or criticism or review, as permitted under the Copyright,Designs and Patents Act, 1988, this publication may bereproduced, stored or transmitted, in any forms or by anymeans, only with the prior permission in writing of thepublishers, or in the case of reprographic reproduction inaccordance with the terms of licences issued by the CopyrightLicensing Agency. Inquiries concerning reproduction outsidethose terms should be sent to the publishers at theundermentioned 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 theinformation and guidance given in this work is correct, allparties must rely upon their own skill and judgment whenmaking use of it. Neither the author nor the publishers assumeany liability to anyone for any loss or damage caused by anyerror or omission in the work, whether such error or omissionis the result of negligence or any other cause. Any and allsuch liability is disclaimed.

The moral right of the author to be identified as author of thiswork has been asserted by him/her in accordance with theCopyright, Designs and Patents Act 1988.

British Library Cataloguing in Publication Data

A CIP catalogue record for this bookis available from the British Library

ISBN 0 85296 922 8

Printed in England by Short Run Press Ltd., Exeter

Introduction

Research into amorphous silicon began, in England, nearly thirty years ago, but the majordevelopment, the demonstration of doping and the subsequent implementation of devices, wasmade in Scotland by Spear and LeComber. The history of their attempts to persuade UKindustry of the utility of the new material would make a casebook study of British managerialand financial short-termism. Since then the physics and applications of amorphous siliconhave been vigorously pursued in America, Asia and Europe, but only the Japanese haveestablished large scale device production.

Amorphous silicon is now, through the efforts of this worldwide community, a matureelectronic material, in the sense that most of its properties are well known, and reasonably wellunderstood. There are problem areas, of course: for example, there are uncertainties overfundamentals like the mobility edge, little understood results like the sign anomalies of the Halleffect, and detailed problems like the existence of exponential regions in the density of states orthe origin of the width and shape of the luminescence spectrum. Lists like these tend to growas you type, but nonetheless, there is an 'industry standard' model, generally accepted, thatqualitatively describes most aspects very well. The days of major new experimentalbreakthroughs are almost certainly over, except perhaps in the less well explored alloys.Amorphous silicon has found its applications role in large area devices: solar cells, xerographyand TFT arrays for matrix addressed displays.

This book follows two previous EMIS volumes on amorphous silicon, and has the sameintention as its predecessors: to provide a compact source of numerical information on differentaspects of the material. The detailed and integrated explanation of possible models was not themain aim, though all contributors have provided enough background for the reader tounderstand the data and its significance. All the Datareviews are new, though some of theauthors have contributed before. My aim as editor was to set out a structure for the book, andthen to persuade leaders in the various specialised areas that they could find the time in theirbusy schedules to distil their expertise into a Datareview. There are topics that I had hopedwould be covered, like novel methods of preparation, for which I was unable to find able andwilling authors in time to meet the publisher's schedule, but on the whole the original plan wasfulfilled. Thus, any lacunae are my responsibility. I believe that the assembled Datareviewsgive a useful and up-to-date summary of the state of knowledge of amorphous silicon.

Those familiar with the earlier editions will notice that the appearance has changed for thebetter, since text is not now restricted by the limitations of a character based screen, and for thesame reason it is now possible to include graphs.

Finally, I would like to thank all contributors to the book. It is their knowledge which gives itwhat value it has. I would also like to thank John Sears at the IEE for his help throughout, forsuggestions of names in areas less familiar to me, and for his continual prodding and harryingof contributors (and occasionally the editor) who were perhaps sometimes heard to wonder whythey had said "yes".

Tim SearleUniversity of SheffieldAugust 1998

G. Adriaenssens

F. Alvarez

W. Beyer

C. Bittencourt

R. Brueggemann

D.E. Carlson

A. Catalano

J.D. Cohen

U. Coscia

R. Durny

F. Finger

E.M. Fortunato

F. Giorgis

F. Giuliani

D. Goldie

T.A. Hayes

Contributing Authors

Katholieke Universiteit Leuven, Lab. Halfgeleiderfysica,Celestijnenlaan 200D, B-3001 Heverlee-Leuven, Belgium

Universidade Estadural de Campinus, Instituto di Fisica,Unicamp 13083-970, Campinus, Sao Paulo, Brazil

ISI-PV, Forschungszentrum Juelich, Juelich D-52425,Germany

Universidade Estadural de Campinus, Instituto di Fisica,Unicamp 13083-970, Campinus, Sao Paulo, Brazil

University of Abertay, Dept. Electrical Engineering, Bell St.,Dundee, DDl IHG, Scotland

Solarex, Thin Film Development, 826 Newtown-Yardley Road,Newtown, PA 18940, USA

MV Systems Inc., 17301 W. Colfax Avenue, Ste#3O5,Golden, CO 8041, USA

University of Oregon, Dept. Physics and MaterialsScience Institute, Eugene, OR 97403, USA

Polytechnic of Turin, Physics Dept. & INM,C. so Duca degli Abruzzi 24, 10129 Turin, Italy

Slovak Technical University, Dept. Physics,Fac. Electrical Engineering, Ilkovicova 3, Bratislava 812 19,Slovak Republic

ISI-PV, Forschungszentrum Juelich, Juelich D-52425,Germany

Uninova - Cemop, Quinta da Torre, 2825 Monte deCaparica, Portugal

Polytechnic of Turin, Physics Dept. & INM,C. so Duca degli Abruzzi 24, 10129 Turin, Italy

Polytechnic of Turin, Physics Dept. & INM,C. so Duca degli Abruzzi 24, 10129 Turin, Italy

University of Dundee, Dept. Applied Physics, Dundee,DDl 4HN, Scotland

Oregon State University, Dept. Mechanical Engineering,Corvallis, OR 97331, USA

4.2

3.8

1.3, 1.4

3.8

4.4

5.5

7.1

3.9

3.6

3.5

1.3, 1.4

6.1

2.6, 2.73.6, 3.7

2.6, 2.73.6, 3.7

4.1

8.2-8.5

D.I. Jones

M.E. Kassner

D. Kruangam

L. Ley

A.H. Mahan

C. Main

R. Martins

V. Nazdazy

H. Ohsaki

P.D. Persans

E. Pincik

CF. Pirri

V. Rigato

P. Roca i Cabarrocas

MJ. Rose

T.M. Searle

I. Shimizu

University of Dundee, Dept. Applied Physics, Dundee,DDl 4HN, Scotland

Oregon State University, Dept. Mechanical Engineering,Corvallis, OR 97331, USA

Chulalongkom University, Dept. Electrical Engineering,Bangkok 10330, Thailand

Universitaet Erlangen, Institut fuer Technische Physik,Erwin-Rommel-Strasse 1, D-91058 Erlangen, Germany

NREL, 1617 Cole Blvd., Golden, CO 8041, USA

University of Abertay, Dept. Electrical Engineering, Bell St.,Dundee, DDl IHG, Scotland

Uninova - Cemop, Quinta da Torre, 2825 Monte deCaparica, Portugal

Slovak Academy of Science, Institute of Physics,Dubravska cefta 9, Bratislava 842 28, Slovak Republic

Central Research Center, Asahi Glass Co. Ltd.,Hazawa-cho, Kanagawa-ku, Yokohama 221, Japan

Rensselaer Polytechnic Institute, Troy, NY 12180-3590,USA

Slovak Technical University, Dept. Physics,Fac. Electrical Engineering, Ilkovicova 3, Bratislava 812 19,Slovak Republic

Polytechnic of Turin, Physics Dept. & INM,C. so Duca degli Abruzzi 24, 10129 Turin, Italy

Polytechnic of Turin, Physics Dept. & INM,C. so Duca degli Abruzzi 24, 10129 Turin, Italy

Laboratoire de Physique des Interfaces et des CouchesMinces, Ecole Polytechnique (UPR 258 du CNRS),F-91128 Palaiseau Cedex, France

University of Dundee, Dept. Applied Physics, Dundee,DDl 4HN, Scotland

University of Sheffield, Dept. Physics, Hicks Building,Sheffield, S3 7RH, UK

Tokyo Institute of Technology, The Graduate School,4259 Nagatsuta, Midori-ku, Yokohama City 227, Japan

4.3

8.2-8.5

6.3, 7.5

3.1,3.2

2.1

4.5

6.1

3.4

8.1, 8.6-8.8

6.2

3.5

2.6, 2.73.6, 3.7

2.6

1.1, 1.2

7.2

5.1-5.4

7.4

S. Shirai

M. Stutzmann

K. Suzuki

R. Swanepoel

Y. Tatsumi

P.C. Taylor

I. Thurzo

E. Tresso

D.L. Williamson

S. Zandolin

Canon Ecology Research & Development Center, Canon Corp.,411 Kizugawadai, Kizu-cho, Souraku-gun, Kyoto 619-02,Japan

Technische Universitaet Muenchen, Walter Schottky Institut,Am Coulombwall, D-85748 Garching, Germany

Toshiba Corporation, Res. Lab. 1, Materials & Devices Labs.,33, Shin Isogo-cho, Isogo-ku, Yokohama 235, Japan

Rand Afrikans University, Dept. Physics, PO Box 524,Johannesburg 2006, South Africa

Shinshu University, Dept. Physics, Inst. Higher Education,Nishi-Nagano, Nagano 380, Japan

University of Utah, Dept. Physics, Room 201,115 South 1400 East, Salt Lake City, UT 84112-0830, USA

Slovak Academy of Science, Institute of Physics,Dubravska cefta 9, Bratislava 842 28, Slovak Republic

Polytechnic of Turin, Physics Dept. & INM,C. so Duca degli Abruzzi 24, 0129 Turin, Italy

Colorado School of Mines, Dept. Physics, Golden,CO 80401, USA

Polytechnic of Turin, Physics Dept. & INM,C. so Duca degli Abruzzi 24, 10129 Turin, Italy

7.4

2.3-2.5

6.4, 7.3

8.9

8.1

3.3

3.4

2.6, 2.73.6, 3.7

2.2

2.6

Abbreviations

The following abbreviations are used in this book:

ACASAXS

BBLBCBISBM

CBCBTCCDCFCFSYPSCMOSCPMCxVCVDCW

DBDBRDCDLTSDMRDOSDRAMDSCDV

ECRELERDERDAESREXAFS

FFFWHM

GDGD

HDHMC

alternating currentanomalous small angle X-ray scattering

bottom blocking layerbond-centrebremstrahlen isochromat spectroscopyblack matrix

conduction bandconduction band tailcharge coupled devicecolour filterconstant field state photoemission yield spectroscopycomplementary metal oxide semiconductorconstant photocurrent methodcurrent-voltage measurementschemical vapour depositioncontinuous wave

dangling bonddistributed Bragg reflectordirect currentdeep level transient spectroscopydeuteron magnetic resonancedensity of statesdynamic random access memorydifferential scanning calorimetrydirect view

electron cyclotron resonanceelectroluminescenceelastic recoil detectionelastic recoil detection analysiselectron spin resonanceextended X-ray absorption fine structure

fill factorfull width at half maximum

gas dischargeglow discharge

high dilution (with hydrogen)heterojunction-monitored capacitance

HOMOCVDHTIHWHWCVD

IBSICTSIPEIRITOIxV

LALCLCDLEDLESRLEYSLOLPCVD

MBEMISMOSMPCMPG

NMR

ODMRODOSOEICOMSOPC

PAPASPCVDPDOSPDSPEPECVDPEPPESPIDPJPLPPES

homogeneous chemical vapour depositionhot carrier tunnelling injectionhot-wirehot-wire chemical vapour deposition

ion beam sputteringisothermal capacitance transient spectroscopyinternal photoemission spectroscopyinfraredindium tin oxidecurrent-voltage measurements

longitudinal acousticalliquid crystalliquid crystal displaylight emitting diodelight-induced electron spin resonancelow energy yield spectroscopylongitudinal opticallow pressure chemical vapour deposition

molecular beam epitaxymetal-insulator-semiconductormetal oxide semiconductormodula ted photocurrentmodulated/moving photocarrier grating

nuclear magnetic resonance

optically detected magnetic resonanceoccupied density of statesoptoelectronic integrated circuitoptically modulated spectroscopyorganic photoconductor

photoinduced absorptionphotoacoustic spectroscopyplasma chemical vapour depositionpartial density of statesphotothermal deflection spectroscopyphotoelectronplasma enhanced chemical vapour depositionphotolithography and etching processphotoelectron spectroscopyphoto-induced dischargeprojectionphotoluminescencephoto-pyroelectric spectroscopy

PTTOFPVPVD

QFRS

RRBSRFRFRPRT

SANSSASSAWSAXSSCLCSCL-TOFSIMSSPSRSSPCSSPGSW

TATBLTCOTEMTFLEDTFPDTFTTOTOFTPCTSCTSCAPTSD

UHVUPSUV

VBVB-DOSVBMVBT

post transit time of flightphotovoltaicplasma vapour deposition

quadrature frequency resolved spectrum

RamanRutherford backscatteringradio frequencyreflective moderemote plasmaroom temperature

small angle neutron scatteringsmall angle scatteringsurface acoustic wavessmall angle X-ray scatteringspace charge limited currentspace charge limited time of flightsecondary ion mass spectrometrysputteringspectral responsesteady state photoconductivitysteady state photocarrier gratingStaebler-Wronski

transverse acousticaltop blocking layertin copper oxidetransmission electron microscopythin film light emitting diodethin film photodiodethin film transistortransverse opticaltime of flighttransient photoconductivitythermally stimulated currentthermally stimulated capacitancethermally stimulated desorption

ultra high vacuumultraviolet excited photoelectron spectroscopyultraviolet

valence bandvalence band density of statesvalence band maximumvalence band tail

VFPVGAVHFVHNVL

XESXPSXPS

voltage-filling pulsevideo graphic arrayvery high frequencyVickers hardness numbervacuum level

X-ray emission spectroscopyX-ray photoelectron spectroscopyX-ray photoemission spectroscopy

v This page has been reformatted by Knovel to provide easier navigation.

Contents

Introduction ............................................................................................................ vii

Contributing Authors .............................................................................................. viii

Abbreviations ......................................................................................................... xi

1. Preparation .................................................................................................... 1 1.1 Growth of Undoped a-Si:H by PECVD ............................................................... 3 1.2 Growth of Doped a-Si:H by PECVD ................................................................... 13 1.3 Growth of a-Si:Ge:H Alloys by PECVD Gas Sources, Conditions in the

Plasma and at the Interface ................................................................................ 20 1.4 Growth of a-Si:Ge:H Alloys by PECVD Optimization of Growth

Parameters, Growth Rates, Microstructure and Material Quality ....................... 30

2. Structural and Vibrational Properties ......................................................... 37 2.1 Structural Information on a-Si:H from IR and Raman Spectroscopy .................. 39 2.2 Structural Information on a-Si:H and Its Alloys from Small Angle Scattering

of X-Rays and Neutrons ..................................................................................... 47 2.3 Data on Hydrogen in a-Si:H from IR and Raman Spectroscopy ......................... 56 2.4 Data on Hydrogen in a-Si:H from NMR .............................................................. 61 2.5 Data on Hydrogen in a-Si:H from Diffusion and Effusion Studies ....................... 66 2.6 Structural Information on a-SiC:H from IR and Raman Spectroscopy ................ 74 2.7 Structural Information on a-SiN:H from IR and Raman Spectroscopy ................ 85

3. Electronic Structure ..................................................................................... 91 3.1 Conduction and Valence Band Density of States of a-Si:H

Photoemission, Inverse Photoemission and Core Level Absorption Spectroscopy ...................................................................................................... 93

3.2 Band Tails of a-Si:H Photoemission and Absorption Data .............................. 113 3.3 Information on Gap States in a-Si:H from ESR and LESR ................................. 139 3.4 Information on Gap States in a-Si:H from Thermal Defect Spectroscopies ........ 143 3.5 Information on Gap States in a-Si:H from Photoinduced Absorption .................. 151 3.6 Information on Gap States in a-SixC1-x:H from ESR, LESR, Constant

Photocurrent and Photothermal Deflection Spectroscopies ............................... 161

vi Contents

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3.7 Information on Gap States in a-Si1-xNx:H from ESR, LESR and Photothermal Deflection Spectroscopies ............................................................ 168

3.8 Valence Band Offsets of a-Si1-xCx on c-Si and a-Si:H ........................................ 174 3.9 Electronic Structure of a-Si:Ge:H ....................................................................... 180

4. Electronic Transport ..................................................................................... 189 4.1 Dark Conductivity in Undoped a-Si:H Deposited by Plasma-Enhanced

CVD Methods ..................................................................................................... 191 4.2 Mobilities in a-Si:H .............................................................................................. 199 4.3 Thermoelectric Power and Hall Effect in a-Si:H .................................................. 209 4.4 Steady State Photoconductivity in a-Si:H and Its Alloys ..................................... 217 4.5 Transient Photoconductivity in a-Si:H and Its Alloys .......................................... 227

5. Recombination of Excess Carriers ............................................................. 235 5.1 Luminescence of a-Si:H ..................................................................................... 237 5.2 Luminescence of a-Si:N:H .................................................................................. 245 5.3 Luminescence of a-Si:C:H .................................................................................. 252 5.4 Luminescence of a-Si:Ge:H ................................................................................ 259 5.5 Light-Induced Defects and the Staebler-Wronski Effect in a-Si:H ...................... 264

6. Junctions and Thin Film Transistors .......................................................... 271 6.1 Schottky a-Si:H Devices ..................................................................................... 273 6.2 a-Si:H/a-Si:X:H Multilayers and Evidence for Quantum Confinement ................ 284 6.3 Electroluminescence from a-Si:H p-i-n Junctions and a-Si:X:H

Heterojunctions ................................................................................................... 293 6.4 a-Si:H Thin Film Transistors ............................................................................... 305

7. Photoelectronic Devices .............................................................................. 311 7.1 Amorphous Silicon Solar Cells ........................................................................... 313 7.2 Amorphous Silicon Photodetectors .................................................................... 319 7.3 Amorphous Silicon Large Area Displays ............................................................ 325 7.4 Amorphous Silicon Xerographic Applications ..................................................... 331 7.5 Amorphous Silicon Alloy LEDs ........................................................................... 337

8. Macroscopic Data ......................................................................................... 347 8.1 Density of a-Si, a-SiNx, a-SiC and a-SiGe .......................................................... 349 8.2 Elastic Constants of a-Si and a-Si:H .................................................................. 359 8.3 Hardness and Wear of a-Si and a-Si:H .............................................................. 363 8.4 Intrinsic Stress in a-Si and a-Si:H Films ............................................................. 367 8.5 Thermal Expansion Coefficient of a-Si and a-Si:H ............................................. 370 8.6 Specific Heat of a-Si, a-Si:H and a-SiNx ............................................................. 372

Contents vii

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8.7 Thermal Conductivity of a-Si, a-SiNx and a-SiC ................................................. 376 8.8 Melting Point of a-Si and a-Si:H ......................................................................... 383 8.9 Optical Functions of Amorphous Silicon ............................................................. 386

Index ..................................................................................................................... 405

CHAPTERl

PREPARATION

1.1 Growth of undoped a-Si:H by PECVD1.2 Growth of doped a-Si:H by PECVD1.3 Growth of a-Si:Ge:H alloys by PECVD - gas sources,

conditions in the plasma and at the interface1.4 Growth of a-Si: Ge:H alloys by PECVD - optimisation of

growth parameters, growth rates, microstructure andmaterial quality

1.1 Growth of undoped a-Si:H by PECVDP. Roca i Cabarrocas

August 1997

A INTRODUCTION

Hydrogenated amorphous silicon thin films have been the subject of extensive research in thepast thirty years, boosted by applications and the challenging fimdamental issues related to thisdisordered semiconductor (structure, doping, stability etc.) [1,2]. Here we focus on the growthprocesses of a-Si:H, which determine, to a large extent, its optoelectronic properties and theperformance of related devices. The optimisation of a-Si:H deposition conditions implies acomplete understanding of the processes involved in its growth as well as the correlationbetween the deposition conditions and the optical, structural and transport properties [2].Because of the disordered nature of a-Si:H, the detailed characterisation of the film propertiesand the correlation between the structure and the optoelectronic properties is still in progress[I]. As a consequence, the optimisation of a-Si.H has been mostly achieved by trial and error,and supported by fundamental studies which offer a better understanding of the growthmechanisms and the necessary framework for a further improvement of this material. FIGURE1 schematically describes such a framework, which can be decomposed into four steps: (i) thedissociation of the gas precursors; (ii) the plasma physics and chemistry, which determine theflux and nature of reactive species to the substrate; (iii) the plasma-surface interactions; and

Plas

ma

proc

esse

sSu

rfac

e an

dBu

lkPr

oces

ses

Substrate TemperatureFIGURE 1 Schematic representation of the processes involved in a-Si:H deposition.

Growth zoneBulk

Surface mobilityChemical equilibrium

Ions, Radicals, Photons,...

Primaryreactions

IncreasingRF power,Pressure,Geometry,Flow

Secondaryreactions

Clusters, PolymersPowder

SiH4FlOw Pump

Electrical PowerElectron density

Energy distribution

(iv)the reactions taking place in a growth-zone where cross-linking reactions result in theformation of the film.

B DEPOSITION METHODS

The search for a-Si:H with improved properties (low defect density, higher carrier mobility,enhanced stability, etc.) has led researchers to explore a large number of deposition methodsand, within each of them, the effects of each process parameter. Amorphous silicon filmsproduced by evaporation of a silicon target or by sputtering in the absence of hydrogen, have ahigh density of defects which render them useless for electronic applications. In contrast, thefilms produced by the dissociation of hydride gases have a low defect density which allowsdoping [3]. It took a few years to recognise the fundamental role of hydrogen in the passivationof silicon dangling bonds and thus in reducing the density of defects in the gap of thesemiconductor. Today, the role of hydrogen is largely recognised and a-Si:H is considered insome aspects as a hydrogen glass [4] in which hydrogen plays a key role during growth as itdetermines the structure of the film [5]. Different methods have been used to dissociate the gasprecursors, all variants of a CVD process: HOMOCVD, PECVD, PHOTOCVD etc. [6-8].Among PECVD methods, different excitation modes and geometry of the reactor have beenexplored [6,9]. RF (13.56 MHz) glow discharge is the most widely used deposition techniquebecause it combines low temperature operation, thanks to the plasma dissociation, and thepossibility of scaling-up the size of the substrates. Even though the use of higher excitationfrequencies has been studied as a way to increase deposition rate [10], high rates are alsoachieved at 13.56 MHz, the standard frequency for industrial applications. As a matter of fact,whatever the technique, the increase of the gas dissociation will favour secondary reactions(FIGURE 1), leading to the formation of polymers, clusters and powder, which will be in mostcases the limiting factor to the increase of the deposition rate [H]. In the following we willfocus on RF glow discharge processes even though most of the discussion can be applied todeposition by other techniques.

C GAS PHASE PROCESSES

Silane is the most common source in a-Si:H deposition even though other hydrides (Si2H6,Si3H8) [12], fluorides (SiF4, Si2F6) [13], chlorides (SiCl4, SiH2Cl2) [14], SiH3F and SiH2F2 [15]source gases have been used. The claimed benefits associated with the use of these gases(higher deposition rates, or improved stability) are however a subject of controversy [12].Indeed, many reactions are involved in the growth process and it is not straightforward toimpute an improvement in the a-Si:H properties to the gas precursor.

The production of reactive species can result from either primary or secondary reactions(FIGURE 1). By primary reactions one understands inelastic collisions between the electronsand the gas molecules. The energy of the electric field is coupled to the electrons through fourbasic mechanisms: (i) secondary electron emission, (ii) stochastic heating, (iii) wave-riding, and(iv) Joule heating [9]. The results of the primary reactions are the radicals and ions responsiblefor a-Si:H deposition [16]. At low pressure, low silane dissociation and temperatures above2000C primary reactions alone are sufficient to describe a-Si:H deposition. Under theseconditions, SiH3 radicals can account for a large fraction of the deposition [17] even though this

has been a subject of controversy [18]. Understanding of plasma chemistry has been achievedthrough plasma diagnostics:

(i) Langmuir probes and electrostatic analysers [19,20] which provide information on theelectron and ion densities and distribution functions.

(ii) Mass spectrometry which provides information on the reactive species: radicals (SiH3,SiH2, SiH, Si), as well as the positive and negative ions reaching the substrate [21].

(iii) Optical techniques such as optical emission spectroscopy, actinometry, opticalabsorption, and laser induced fluorescence, giving information on the neutral and excitedspecies of the discharge [21]. Light scattering techniques have been recently developedto study particles in plasmas [22].

Thanks to these diagnostic techniques, a set of cross sections for the different reactions as wellas the constant rates have been tabulated [23,24]. They are essential for modelling silanedissociation and a-Si:H deposition processes [25]. Nowadays, modelling a-Si:H depositionunder conditions dominated by primary processes is satisfactory and allows prediction of a-Si.Hdeposition rates [26]. However, applications push towards high deposition rates and/or lowsubstrate temperatures. In those cases, secondary reactions are no longer negligible and theymay completely determine the film quality. In contrast to the primary reactions, secondaryreactions are less well characterised and therefore the predictions of a-Si:H deposition modelsbecome less reliable when secondary reactions are dominant.

The effect of secondary reactions has been mainly addressed through the formation of powders.Powder formation in silane plasmas was recognised from early studies as a factor limiting filmquality [23,27], even though low defect density a-Si:H films can be deposited at high ratesunder powder conditions [H]. More recently, dusty plasmas have been the subject ofworkshops [28]. Even though the detailed mechanisms leading to powder formation in silaneplasmas are not completely known, strong evidence has been given for anions being theprecursors of powder [29]. The effect of powders on film properties has not been completelyelucidated. Large powders are negatively charged and therefore confined in the plasma untilother forces (thermophoresis, gravity, ion drag, gas flow, etc.) can overcome the electrostaticconfinement [30]. When powders overcome the electrostatic confinement, they can beincorporated in the a-Si:H thin film and produce macroscopic defects like pin-holes. To avoidthese problems, schemes like a progressive decrease of RF power, total pressure, or an increaseof the gas flow have been suggested in order to sweep the powders out of the discharge, andimpede them from falling on the substrate [9]. Rather than avoid powder formation, there is atendency to live1 with powders because their presence favours the coupling of the RF powerinto the discharge and therefore the achievement of high deposition rates [31]. Finally, recentreports focus on the role of small particles on a-Si:H properties. While some reports suggestthat these nanoparticles will have a negative effect [32], there is experimental evidence that theincorporation of these nanoparticles in a-Si:H films can result in nanostructured silicon thin filmswith improved properties compared to a-Si:H [33,34].

This short review of the gas phase processes highlights the complexity of the plasma processesinvolved in a-Si:H growth. Indeed, depending on the process conditions (FIGURE 1), the

growth of a-Si:H will be a consequence of primary reactions, secondary reactions, or both ofthem, resulting in a wide variety of a-Si:H materials.

D SURFACE AND BULK PROCESSES

Gas phase processes are important because they determine the nature of the reactive speciescontributing to a-Si:H deposition. However, a-Si:H growth cannot be considered just as acondensation of radicals on the substrate. Indeed, while the radicals and ions responsible fora-Si:H deposition are highly hydrogenated, a-Si:H films contain 10% hydrogen. Thereforecross-linking reactions must take place at the film surface or in a growth-zone [23]. One couldask whether the a-Si:H properties are determined by the film precursors (gas phase reactions) orby the reactions of these precursors on the film surface. As for the plasma processes, diagnostictechniques have been developed to monitor in-situ the growth processes and the optical andelectrical properties of the films/devices. UV-visible ellipsometry provides real-time informationon the optical properties of the films, and through the use of effective medium theories allowsone to get a clear picture of the effects of the different process parameters on the growthmechanisms (homogeneous growth, nucleation, interaction with the substrate) and the evolutionof the film properties during growth [35,36]. The extension of ellipsometry to the infraredwavelengths [35,37] or the use of other IR techniques [38] provides information on thehydrogen bonding and therefore on the cross-linking reactions resulting in the formation of thefilm. The transport properties can also be measured in-situ by the use of other techniques suchas the Kelvin probe which provides information on the effects of doping and the evolution of thepotential at interfaces [39], optogalvanic photoemission which gives information on the workfunction of the layer [40], and time resolved microwave conductivity which can provideinformation on the majority carrier mobility and lifetime [41].

While the use of these techniques helps the optimisation of the deposition process and gives adeep insight into the growth mechanisms, the important parameters for the surface reactions arenot always directly accessible. Among them are the substrate temperature, the surface mobilityand the sticking coefficient of the film precursors.

The effects of substrate temperature on the growth and properties of a-Si:H thin films have beenwidely studied. However, the temperature at the surface of the growing film cannot be easilymeasured and is affected by the plasma conditions and reactor geometry [42,30]. As aconsequence, and because substrate temperature affects not only surface reactions but also theplasma processes [43,44], different and/or opposite trends can be found for the deposition rateor in a-Si:H properties (see FIGURE 2).

The mobility of the radicals at the growing a-Si:H surface is often considered an importantfactor for the deposition of high quality a-Si:H films [45-47]. The combination of surfacemobility and hydrogen coverage can account for example for the usually observed optimumsubstrate temperature around 2500C. At lower substrate temperature, surface mobility wouldbe too low for an optimum growth to occur, while at higher temperature the depletion ofhydrogen at the film surface would produce surface defects and a decrease of the surfacemobility. This concept of surface mobility goes along with the sticking coefficient of theradicals. The sticking coefficient of SiH3 has been measured to be 0.2 - 0.3 [48,49], which isequivalent to a high surface diffusion. In contrast, because they have more than one unsatisfied

bond, a sticking coefficient close to 1 is attributed to SiH2, SiH and Si radicals; i.e. they havelow surface mobility. Based on these ideas, it is commonly accepted that SiH3 is the gooda-Si:H precursor. This distinction between radicals with a low sticking coefficient (high surfacemobility) and radicals with high sticking coefficient (low surface mobility) has also been used todistinguish between CVD-like and PVD-like processes [50].

CVD-like processes usually lead to low deposition rates. As the plasma parameters are changedin order to increase the deposition rate, it is generally found that the substrate temperature hasto be increased to obtain low defect density a-Si:H [47,51]. Indeed, in the framework ofmodels based on the mobility of the film precursors, the increase of the deposition rate shouldbe accompanied by the supply of more energy to the growing surface in order to enhance thediffusion of the film precursors and favour cross-linking reactions. Alternative models in whichchemical reactions take place in a growth-zone have been proposed. Under the hypothesis thatgrowth reactions take place under equilibrium conditions, Winer proposed that optimal growthoccurs when the deposition rate is equal to the rate of hydrogen diffusion from the growth-zone[51]. As in the case of the surface models in which only the mobility of the film precursors istaken into account, these equilibrium models account for the shift of the optimum substratetemperature to higher values when the deposition rate increases. These equilibrium modelshave been further elaborated by taking into account the evolution of the chemical potential ofhydrogen in the plasma and in the growing film [5]. Models based on surface mobility or onchemical equilibrium reactions in a growth-zone give a good description of the optimumsubstrate temperature for a-Si:H growth, but are limited to growth from SiH3. However, lowdefect density a-Si:H films have been deposited even at 500C [52] and/or under conditions ofhigh pressure and RF power, under which other radicals or even silicon nanoparticles contributeto the growth [33,34]. Why can low defect density a-Si:H films be deposited under suchconditions? The weakness of surface models lies in the underlying hypothesis: growth fromSiH3 radicals and process controlled by substrate temperature and deposition rate. Indeed,other parameters than the substrate temperature can modify the diffusion rate of hydrogen [53].As an example, besides radicals, positive ions also contribute to a-Si:H deposition and eventhough the ratio of ions to radicals is generally smaller than 0.1, their effect on film properties isvery important, in particular under non-optimised deposition conditions [54]. As a matter offact, ion bombardment has been found to have effects comparable to an increase of the surfacetemperature [55].

E OPTIMISATION OF DEPOSITION CONDITIONS

The optimisation of a-Si:H growth has been one of the main subjects in this field. The effects ofdifferent plasma conditions (FIGURE 1), reactor geometry, excitation frequency and dilution ofthe silane on a-Si:H properties have been widely studied [7]. The difficulty of optimising a-Si:Hgrowth is related to the fact that these parameters are correlated. Indeed, the important aspectsfor the growth are the rate of production of reactive species, their interaction during theirdiffusion to the walls (substrates) and their reactions at the growing surface.

Increasing the RF power, for example, increases the dissociation rate and therefore thedeposition rate. However, the higher dissociation may result in a change in the species involvedin the growth because of an increasing importance of secondary reactions [56].

An increase of the pressure may produce similar effects: enhanced dissociation because of theincreased electron-silane collisions, enhanced secondary reactions and deposition rate, alongwith a change in the nature of film precursors [12,57].

The dilution of silane in inert gases (argon, helium, xenon) can affect the electron distributionfunction, the rate of silane dissociation and therefore the deposition rate [58]. While inert gasesdo not play a role in the cross-linking reactions taking place at the film surface or in the growth-zone, the use of hydrogen dilution strongly modifies them. The use of hydrogen dilution hasbeen reported as an excellent way to improve a-Si:H properties [59]. As a matter of fact, theinteraction of hydrogen in the growth-zone allows a change from a-Si:H to microcrystallinesilicon growth [60].

Changing the reactor geometry, in particular the inter-electrode distance, strongly affects theflux of particles to the substrate [24,61,62]. To some extent, the increase of inter-electrodedistance is equivalent to an increase of the total pressure (Paschen's law [19]).

The substrate temperature, whose effects have been mainly discussed in terms of surfacereactions [45-47], also has a strong effect on plasma reactions [44], in particular on theformation of powders [43].

As an example of the interdependence of the deposition parameters, FIGURE 2 shows thedeposition rate as a function of the substrate temperature for a-Si:H films deposited in the samereactor [42] under different plasma conditions.

Dep

ositi

on ra

te (A

Js)

Substrate temperature (0C)FIGURE 2 Effect of the substrate temperature on the deposition rate of a-Si:H films deposited under different

plasma conditions. Series A: pure silane, 40 mtorr, low RF power (0.5 W) and inter-electrode distance, d,of 2.8 cm. Series B: 40% silane in helium gas mixture, 550 mtorr, 10 W, d = 12 mm. Series C: same

conditions as series B but d = 28 mm. Series D: same conditions as series C but 15 W.

Series A (open circles) correspond to a-Si:H films produced by the dissociation of pure silane atlow pressure (40 mtorr) and low RF power (0.5 W), with the reactor having an inter-electrodedistance of 28 mm. Under these conditions there is no formation of powder in the reactor. The

increase of the deposition rate in this so-called a-regime [57] has been explained by thetemperature dependence of gas phase reactions [44],

Series B (filled circles) corresponds to a-Si:H films produced by the dissociation of 40% silanein helium gas mixture at high pressure (550 mtorr) and high RF power (10 W)9 with the reactorhaving an inter-electrode distance of 12 mm. We observe the same dependence of thedeposition rate on the substrate temperature as for series A, except for the factor of 4 increasein deposition rate, which can be explained by the higher RF power and the smaller inter-electrode distance.

Deposition under the same plasma conditions as in series B but with an inter-electrode distanceof 28 mm (open squares, series C) completely changes the temperature dependence of thedeposition rate. In this case the plasma is in the so-called y-regime [12,57] in which the RFpower is more efficiently coupled to the plasma because of the formation of powders. Theincrease of deposition rate between 1500C and 2500C can be attributed to the reduction ofpowder formation and therefore to a decrease in the losses to the pumps [56]. Above 2500C,the further reduction in powder formation makes the plasma switch to the a-regime. Finally,under the same conditions as in series C, but with an RF power of 15 W (series D, filledsquares), the deposition rate shows a similar trend but with higher absolute values.

These changes in deposition rate are accompanied by changes in the optical, structural andelectronic properties of the films [11,53]. FIGURE 3 shows the defect density as a function ofthe substrate temperature for the above series of samples. For each series there is a temperatureabove which the defect density sharply decreases. The higher the deposition rate, the higher thetemperature threshold, in qualitative agreement with models taking into account the mobility ofthe radicals at the film surface or the diffusion of hydrogen in the growth zone. However, thefilms of series B have about the same threshold temperature as the films of series A, despite thefactor of 4 higher deposition rate, while the samples of series C have a higher thresholdtemperature despite a smaller deposition rate. Therefore, other factors than deposition rate andtemperature determine the optimal growth conditions. This is emphasised by the two stars at100 and 500C corresponding to films deposited at similar growth rate to series A, but at lowerpressure (1000C) or from a silane-hydrogen mixture (500C). In both cases the lower defectdensity at these temperatures can be explained by the enhanced ion bombardment whichincreases the surface temperature and avoids the polymerisation at the film surface [52].Therefore plasma and/or surface polymerisation appear as limiting factors for the growth of lowdefect density a-Si:H at low substrate temperatures. Indeed, while it is not clear whetherpolymerisation takes place in the gas phase or at the film surface, it is well established thatsilicon polymers are thermally dissociated above 2000C [63], which could explain why theoptimal growth temperature is often reported to be above 2000C as well as the sharp decreasein defect density (FIGURE 3). While SinH1n polymers result in highly hydrogenated anddisordered films with a high defect density, the growth of a-Si:H films under conditions close tothe formation of powders can result in low defect density films with enhanced microstructureand stability [33,34]. As an example, the samples of series C and D in FIGURE 3 weredeposited under powder conditions. However, except for the films deposited below 2000C,corresponding to the temperature necessary to avoid polymers, the defect density is low.However the films have a larger microstructure as can be defined from the hydrogen bondingfor example [53]. Under these conditions of high deposition rate the models based on surfacemobility and chemical equilibrium reactions still apply. Indeed, as the deposition rate increases

Substrate temperature (0C)FIGURE 3 Effect of substrate temperature on the defect density of a-Si:H films produced

under different plasma conditions. Same symbols as in FIGURE 2.

F CONCLUSION

The growth of a-Si:H thin films by the dissociation of gas precursors in a glow discharge is acomplex process in which gas phase reactions as well as surface reactions have to beconsidered. In-situ diagnostic techniques of the plasma phase and of the growing film providevaluable information to understand the growth mechanisms and help to optimise the depositionconditions. These in-situ techniques are complemented by ex-situ characterisation methodsproviding complementary information on the optical and electronic properties of the films.Deposition rate and substrate temperature are the main parameters which determine theproperties of a-Si:H. While many models based on the surface mobility of the film precursors oron the existence of chemical equilibrium reactions at the growing film surface describe thegrowth under conditions where SiH3 radicals are the main film precursors, the ability to depositlow defect density a-Si:H films down to 500C or under conditions where other radicals thanSiH3 or even nanoparticles contribute to deposition suggests that ion bombardment andsecondary reactions taking place in the gas phase also play a crucial role in a-Si:H quality.

As a conclusion, a-Si:H is a new thin film which, pushed by its applications, has experiencedextraordinary development in the last thirty years. The understanding acquired of the plasmaand growth processes allows us to foresee further developments of this material.

Def

ect

dens

ity

(cm

"3 )

(series D versus series C), the threshold substrate temperature shifts to higher temperature.However, the sharp transition between high and low defect density films suggests that also thenature of the film precursors changes. Therefore, deposition under powder conditions canresult in a-Si:H films with low defect density. This is further validated by recent studies on thegrowth of a-Si:H films under conditions where silicon nanoparticles as well as radicalscontribute to the growth [33].

ABCD

REFERENCES

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1.2 Growth of doped a-Si:H by PECVDP. Roca i Cabarrocas

September 1997

A INTRODUCTION

Doping of amorphous semiconductors was unexpected as it is against Mott's 8-N rule [1]stating that atoms in an amorphous solid should have a coordination of min[N, 8-N], where N isthe number of valence electrons. Thus boron and phosphorus should be threefold coordinated,i.e. non-dopant. However in 1975 Spear and Le Comber reported on the substitutional dopingof amorphous silicon films, produced by the decomposition of silane [2]. Of course, thequestion as to why it is possible to dope a-Si:H drew a large amount of attention raising thepossibility of various applications. Review articles on the doping of a-Si:H can be found in[3-6].

B DOPING EFFICIENCY AND DOPING MECHANISMS

Even though the predictions of non-dopability of a-Si:H were demonstrated to be wrong [2], itappeared that most of the dopants in a-Si:H are threefold coordinated, as expected from the8-N rule, i.e. doping efficiency (r\) is very low. Indeed, only extended X-ray absorption finestructure (EXAFS) [7] measurements on a-Si:H films doped with either P or As gave, at themost, a 20% fraction of fourfold coordinated atoms, while nuclear magnetic resonancemeasurements on boron-doped a-Si:H revealed only threefold coordinated boron [8]. Detailedstudies by Street using a charge sweep-out experiment [9] confirmed this low doping efficiency(IO"3 < T] < 10"1), which is inversely proportional to the square root of the concentration ofdopants in the gas phase [6,10]:

(1)

where k is a constant and Cg is defined by the ratio of the dopant gas flow rate to the silane flowrate. The low values of r\ suggest that most of the dopant atoms satisfy the 8-N rule and only asmall fraction of them are fourfold coordinated.

Different models have been developed to explain doping in a-Si.H. In particular, the formationof valence-alternation-pairs has been suggested as the mechanism responsible for doping[11,12]. However, it was the extension of the 8-N rule to charged impurity states by Street thatallowed for a better understanding of doping of a-Si:H [10]. According to Street, doping byphosphorus can be described by a solid state chemical equilibrium reaction:

(2)

where the equilibrium between threefold and fourfold coordinated phosphorus atoms isbalanced by defects (silicon dangling bonds). This reaction, first proposed to take place at the

film growing surface, was extended to describe bulk dopant equilibrium [13]. As a directconsequence of this model doping of a-Si:H is accompanied by an increase of the defect density.Moreover the application of the law of mass action to EQN (2) leads to the square rootdependence of the doping efficiency (EQN (I)) and is in agreement with most of theexperimental results concerning doping in a-Si:H and a-Ge:H [3].

Surprisingly, the deposition conditions are not taken into account in EQN (2); i.e. doping isconsidered independently of the plasma parameters. While the importance of solid-phasereactions is supported by a large number of studies concerning thermal equilibrium andmetastability in a-Si:H thin films [6], some experimental results cannot be explained, inparticular the presence of neutral defects and the fact that the dopant efficiency is dependent onthe gas phase concentration and not on the solid phase concentration [14]. This has led someauthors to extend Street's approach to specific chemical reactions taking place at the a-Si:Hsurface during growth [15,16]. Further progress in understanding doping of a-Si:H will beobtained by a better understanding of a-Si:H growth processes [14].

C DOPANT SOURCES

Interestingly, with the growth of a-Si:H by PECVD dopants can be mixed with silane in acontrolled way to achieve the desired doping level. Indeed, even though other techniques suchas ion implantation or diffusion from a solid source have been used [17,18], the direct additionof the dopant gases to the silane remains the easiest way to achieve doping. While phosphineand diborane are the most common dopants, other gases have been used to reduce the toxicityor to improve the doping efficiency. In particular, doping with diborane is difficult to controlbecause of the thermal CVD at low temperature [19]. Trimethylboron has attracted muchattention because of its higher stability and lower toxicity [20]. Moreover, doping withtrimethylboron is accompanied by the incorporation of carbon in the a-Si:H network [21] whichcan be exploited for the growth of low absorption layers. Boron trifluoride [22],trimethylgallium [20,23], trimethylaluminium, and triethylboron [24] have been studied asp-type dopants, while liquid organic sources have been used for n-type doping [25]. Impuritiesat low concentration have also been reported to act as dopants in a-Si:H, with particular reportsof the donor-like effects of oxygen [26,27] and nitrogen [28]. In the following we will focus onthe doping of a-Si:H by phosphine or diborane.

D EFFECTS OF DOPANTS ON a-Si:H GROWTH

Few studies have been devoted to the effect of dopants on the plasma chemistry. However,both B2H6 and PH3 have ionisation potentials below 11 eV and may modify the plasmacomposition in a similar way to Ar and Kr dilution [29]. Mass spectrometry studies have shownthat while silane and silane-phosphine discharges are similar, the addition of diborane results ina modification of the discharge because of the higher dissociation of diborane and the formationof diboron-type ions [30].

The effects of dopants on a-Si:H growth have been widely studied. It is well established thatthe addition of diborane results in an increased deposition rate for deposition from silaneprecursors [31], while the addition of phosphine has a small effect or produces a reduction in

the deposition rate [32]. The enhancement of the deposition rate by diborane has beendiscussed in terms of a catalytic effect of diborane [33]. However, when a-Si:H is obtainedfrom halogenated silicon reactants (SiCU and SiF4), the addition of diborane decreases thedeposition rate, while phosphine increases it [34]. This opposite effect of dopants on thedeposition rate has been explained by the effect of the surface band-bending on thechemisorption of silicon radicals.

In-situ ellipsometry studies have been used to monitor the growth of a-Si:H films, the effects ofdoping with diborane on the initial stages of deposition [35], and the thermal CVD of diborane[19]. Moreover, the in-situ Kelvin probe has been proved an excellent technique to controlchanges in the Fermi-level position during a-Si:H doping with either phosphine or diborane[36].

The studies of a-Si:H growth clearly indicate that the addition of a small amount of dopant tothe discharge results in dramatic changes in the growth processes, especially in the case ofdiborane.

E EFFECTS OF DOPANTS ON a-Si:H PROPERTIES

The main effect of doping is the change of the Fermi level position within the gap of thesemiconductor. In a-Si.H, p-type doping allows the Fermi level to be moved down to 0.3 eVfrom the valence band, while n-type doping allows the Fermi level to be moved up to 0.2 eVfrom the conduction band. These changes in the Fermi level position are accompanied by achange of more than eight orders of magnitude in the dark conductivity [2]. However, contraryto crystalline silicon, degenerate doping is not observed in a-Si:H. This is due to the existenceof band-tails and deep defects in the gap of a-Si:H, and to the creation of midgap defects alongwith doping.

Besides doping, the effects of dopants on the growth processes are also reflected by changes inthe structure and properties of a-Si:H films. Those changes have been found to strongly dependon the doping level. For Cg < 10"3 (doping regime) the addition of PH3 or B2H6 to silane resultsin an increase of the hydrogen content and of the optical gap of the films, while for Cg > 10"3(alloy regime) the addition of dopants leads to a decrease of the hydrogen content and of theoptical gap [37,38]. These effects have been mostly studied in the case of boron doping[39-42]. In particular, an optimum doping level in the range OfB2H6ZSiH4 flow rates 10"4 -10"3 has been observed and correlated with an improvement in the open circuit voltage of solarcells [43].

Structural changes are also often reported, in particular for boron doped a-Si.H films whereinan inhomogeneous morphology has been observed [44]. The presence of microstructuralinhomogeneities in boron doped films, supported by the observation of boron clustering [45], isalso inferred from hydrogen evolution measurements which show that boron, contrary tophosphorus, results in a low temperature hydrogen evolution [46-48]. This effect is alsosupported by annealing studies which show large changes of the optical and electricalproperties. As annealing temperature is increased the hydrogen content decreases, darkconductivity increases, and hydrogen accumulates in internal voids as H2 molecules, even forannealing temperatures below the deposition temperature [49]. This evolution of hydrogen at

low temperatures can be responsible for the thickness dependence of the properties of borondoped films [5].

The addition of dopants in a-Si:H results in an increase of the defect density of the films asshown by the decrease of photoluminescence [6] and measured by a wide range of techniques[3,50,51]. Moreover, it is found that the density of defects is about a factor often higher thanthe concentration of fourfold coordinated dopant atoms [52]. The changes in defect density arealso reflected in strong changes of the photoconductivity, transport and deep trapping [53-57].Overall, doped a-Si:H films have high defect density and poor transport properties, which limitstheir use as active layers in a device. However they are widely used as contacts with intrinsica-Si:H films.

F EFFECTS OF GROWTH PARAMETERS

As discussed above, the doping efficiency is inversely proportional to the square root of theconcentration of dopants in the gas phase. However, this refers to films deposited under thesame plasma conditions [14]. Change of the discharge conditions affects a-Si:H deposition andthe chemical reactions at the film surface, and should therefore affect doping.

The substrate temperature has been one of the most studied parameters as it strongly affects theproperties of intrinsic a-Si:H. For both phosphorus and boron doped films, the decrease of thesubstrate temperature results in a decrease of the dark conductivity [5,58,59]. Nevertheless,highly conductive boron and phosphorus doped films have also been obtained at 500C [60].Moreover, changing substrate temperature also affects the incorporation of hydrogen and ofdopants in the film [59,61], which makes it difficult to correlate process conditions and dopingefficiency.

The effect of ion bombardment has attracted particular attention in the study of anode versuscathode deposited samples [39,46] or through the effect of a bias voltage applied to the anode.In the case of intrinsic a-Si:H, ion bombardment improves the quality of the material [62].However, quite different effects have been reported on doped films [63,64].

As for the substrate temperature, the plasma parameters such as pressure, gas flow rate, poweretc. affect both dopant and hydrogen incorporation [61]. As a matter of fact, the simultaneouschanges in hydrogen and dopant concentration suggest that hydrogen incorporation is governedby the Fermi level position [65]. Moreover, hydrogen plays a crucial role in determining thedoping efficiency in a-Si:H, particularly in boron doped films where the formation of boron-hydrogen complexes, as observed in crystalline silicon [66], may result in the decrease of thedoping efficiency [67]. The passivation of boron by hydrogen has also been observed in a-Si:Hthrough in-situ studies [68] and is supported by nuclear magnetic resonance measurementswhich indicate that about half of the boron and phosphorus atoms in a-Si:H form H-dopantcomplexes [69]. Moreover, the importance of hydrogen in the metastability of phosphorus andboron doped a-Si:H films has also been studied [70,71]. However, while light-soaking ofphosphorus doped a-Si:H films produces a decrease of the fourfold coordinated P atoms[72,73], light-soaking in boron doped a-Si:H results in an increase of the film conductivity [74],which suggests that in this case light favours the formation of fourfold coordinated boronatoms. This difference is in agreement with recent models which show that, contrary to

phosphorus doping, the doping efficiency in boron doped a-Si.H is limited by H passivation[75]. Light-induced dopant activation has also been observed in the p-layer of p-i-n solar cells,which results in an increase of the open circuit voltage of the devices [76].

G CONCLUSION

Substitution^ doping in a-Si:H is against the 8-N rule of bonding in amorphoussemiconductors. Nevertheless, a small fraction of dopants in a-Si:H (IO"2 - 10"3) violates thisrule and is fourfold coordinated. However, doping is accompanied by the creation of defectsand results in highly defective films. Models based on chemical equilibrium reactions takingplace at the growing surface or in the bulk of the material account for most of the observedeffects of doping. This contrasts with the poor knowledge of the effects of dopants in theplasma chemistry and growth processes.

Incorporation of phosphorus or boron in a-Si:H has quite different effects on the materialproperties. While phosphorus doping is quite well understood, doping with diborane still needsfurther research as it results in large structural changes of the a-Si:H matrix, which arecorrelated to the changes in the incorporation of hydrogen forming boron-hydrogen complexes.

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1.3 Growth of a-Si:Ge:H alloys by PECVD - gas sources,conditions in the plasma and at the interfaceF. Finger and W. Beyer

September 1997

A INTRODUCTION

The preparation of amorphous hydrogenated and/or fluorinated silicon germanium alloys(a-Si:Ge:(H,F)) with a large variety of deposition techniques has been given great attentionsince it was shown that by admixture of Ge and Si the optical gap can be shifted continuouslybetween the value for a-Si:H (1.7 - 1.8 eV) and that for a-Ge:H (1.1 - 1.2 eV) [I]. This makesthe material an interesting candidate for meeting the requirements of, for example, stacked solarcells and optoelectronic devices where a certain bandgap, a variety of different bandgaps orother material properties are needed. Si-based thin film solar cells including a-Si:Ge:H layershave the highest stable efficiency of 13% [2], and other applications like multispectral coloursensors have been successfully developed [3]. Research into preparation techniques fora-Si:Ge:H alloys is necessary because together with alloying and the required shift of the opticalgap other material properties are also influenced - often in a way which is detrimental fortechnical applications. Research activities have thus concentrated on investigating the growthprocess and proposing and testing alternative deposition techniques in order to control thematerial properties according to specified optimisation parameters, believed to be responsiblefor or correlated with desirable performance for applications.

The understanding and interpretation of deposition processes of a-Si:Ge:H is not conclusive.An obvious complication is that optimised growth conditions on the Si- and the Ge-rich side,respectively, might be quite different. While most studies approach from the Si-side, somereports start on the Ge-side. Several seemingly partly contradictory requirements for optimisedgrowth conditions have been postulated or concluded including (among others):

(1) choose only neutral radicals with long lifetime [4], reduce short lifetime radicals throughreaction with H2 or avoiding depletion OfGeH4 [5];

(2) enhance surface diffusion of Ge related radicals by hydrogen coverage [6];

(3) use an alternative Ge dangling bond terminator [7,8];

(4) need of no [4,9,10], versus moderate [11], versus high energy ion bombardment (or highion bombardment at high deposition rate [12]) to reduce heterogeneities and obtainoptimum performance [13-15];

(5) use higher electron temperatures in the PECVD process to obtain favourable dischargechemistry [15];

(6) need of low [ 16,17] versus high [18,19] deposition rate.

Previous Page

In this Datareview the concepts of different PECVD based preparation techniques for a-Si:Ge:Hare summarised. The reader is referred to review articles and data collections on a-Si:Ge:Hpreparation techniques and properties. These review articles have summarised in the pastcertain stages of development or special preparation techniques and related material issues ofa-Si:Ge:H. Most of these review articles contain a large amount of additional literature. Theauthors represent some of the important research groups involved with a-Si:Ge:H.

Early work is summarised in [20] and [21].

An important stage in optimisation with hydrogen dilution and triode systems is reported byMatsuda and Tanaka [22].

Bullot [23], and in much detail Tsuo and Luft [24] and Luft and Tsuo [25], compare anddescribe different preparation techniques for a-Si:Ge:H and collect results obtained from thedifferent techniques. The latter two are the most comprehensive works found to date onpreparation techniques for a-Si:Ge:H. A detailed study of the plasma chemistry insilane/germane and disilane/germane mixtures is reported by Doyle et al [26]. Also [27]contains data on the preparation and the properties of a-Si:Ge:H alloys.

Among the reports which deal with the end-point of the alloy system (a-Ge:H) summaries aregiven in [28,29].

Data on the alloying effect on the density of states and the electronic transport andrecombination in a-Si:Ge:H material is reported in [30-35].

Materials from different preparation techniques ( different source gases) are also compared in[36,37], emphasising the importance of microstructure, and the direction for future research isdiscussed in [37].

Finally, the widest collection of selected properties (photosensitivity, slope of the opticalabsorption edge and defect density) as a function of the optical gap in a-Si:Ge:H prepared bydifferent techniques is given by Ichikawa and Sasaki [38]. We use this data collection as areference for the present Datareview (see FIGURES 1 to 3 in Datareview 1.4 in this book). Byfar the most data found for photosensitivity, slope of the optical absorption edge and defectdensity in the literature falls in the envelope spanned by this collection. We want to stress thatfor the present Datareview more than 400 articles in journals, conference proceedings andbooks dealing with a-Si:Ge:H dating back to 1977 have been searched. Among them about 200report on details of the preparation techniques. Thus the data collections given in [24], [25] and[38], together with the individual data points quoted in this Datareview, should give a completeoverview.

B DEPOSITION TECHNIQUES

The simplest 'original1 deposition technique for a-Si:Ge:H is an RF (13.56 MHz) glow dischargein a diode type reactor with gas mixtures of SiH4 and GeH4. Starting from this basic technique,where results have been described in the early review articles [20,21], the following alternativesto the standard PECVD process have been used and described in the literature:

(1) Hydrogen dilution(2) Fluorinated gases(3) Disilane(4) Variation of deposition reactor

(i) Triode reactor(ii) DC plasma(iii) Cathodic deposition

(5) Helium, argon and other dilution gases(6) Microwave, ECR, remote plasma, VHF(7) Special gases(8) Other techniques and combinations(9) Substrate temperature

In addition, photo-CVD, sputtering, thermal CVD, catalytic CVD and evaporation have beenused, which will not be considered here. References on these methods are to be found in[24,25].

Bl Hydrogen Dilution

The concept of hydrogen dilution of the process gases SiH4 and GeH4 for deposition ofa-Si:Ge:H was first introduced by Matsuda et al in 1986 [6]. Hydrogen dilution was also usedin earlier work but not systematically to improve the material properties (see early reviewarticles). The concept is based on the ideas that (1) Ge related radicals might have smallersurface diffusion coefficients on the film growing surface than Si-related radicals, and (2) Hatoms bonding with Ge are thermally evolved at lower temperatures compared with H on Si,which causes a higher density of free bonds on the growing surface of Ge-rich a-SiGe alloys [6].By adding hydrogen in the discharge it was believed that the H radical density increases, leadingto a better surface coverage and thus a higher surface diffusion coefficient. Furthermore, the H2dilution was thought to reduce the density of highly reactive species GeHy (y = 0 - 2).

Although it was found later that in fact upon hydrogen dilution the H radical density decreasesin an SiH4ZH2 discharge [39], and thus the role of hydrogen dilution had to be reinterpreted [5],the H2-dilution method was very successful in improving the electronic properties at a givenoptical gap. H2-dilution is also the method most widely used to date in combination with othermethods (triode reactor, fluorinated gases, disilane, microwave frequencies etc.). Generally theimprovements reported by the numerous studies confirm the trends already reported in theoriginal work [6]. Data are collected in [25,38]. An effect considered as a disadvantage fortechnical applications is the decrease in deposition rate upon dilution.

As an offshoot of the hydrogen dilution method one can consider the hydrogen plasmaannealing method where a thin deposited a-Si:Ge:H layer is exposed to a hydrogen plasma andthis step is repeated many times. The idea is to relax the SiGe network and to passivatedangling bond states. Marked improvement over other techniques was not achieved [40-42].

Most recent reports using hydrogen dilution of process gases are given in [12,15,17,42-47].

B2 Fluorinated Gases

On the basis of the preferential attachment effect of hydrogen to silicon in SiH4-GeH4 mixturebased a-Si:Ge:H alloys, Paul et al [7,20] proposed in 1981 the use of fluorinated gases (or theincorporation of oxygen) for a-Si:Ge:H deposition. Fluorine (and oxygen) was considered tosaturate Ge dangling bonds more effectively than hydrogen. The use of fluorine had alreadybeen proposed earlier for the improvement of a-Si films [48]. a-SiGe film deposition usingfluorinated gases was first reported by Nozawa et al [8] in 1983 who characterised the filmquality predominantly by photo- and dark conductivity measurements. While a mixture of SiF4and GeF4 did not result in an enhancement of photoconductivity compared to hydrogenatedreference samples, a considerable improvement was reported for gas mixtures OfGeF4, SiF4 andH2 as well as GeH4, SiF4 and H2. These results were confirmed by other researchers [49].Subsequently, it was recognised by the Shimizu group [50] using IR absorption and XPSmeasurements that little (less than 1%) fluorine was incorporated into their good quality SiGefilms suggesting that the improvement of the fluorinated films compared to the hydrogenatedsamples was not due to fluorine incorporation. Oda et al [50] attributed the improvement toplasma chemical effects. Based on the reported improvement of photoconductivity and on thefact that fluorinated gases are much safer and easier to handle, several groups startedprogrammes to investigate SiF4(SiHt)-GeF4-H2 based alloys [49,51-53]. While mostly RFglow-discharge was applied, Aljishi et al [53] reported results for both RF and DC glow-discharge deposition. In a comprehensive article, Mackenzie et al [54] compared the propertiesof (SiF4, GeF4, H2) and (SiH4, GeH4) based plasma grown a-Si:Ge:H,F films with Eg = 1.4 - 1.5eV. They noted in the substrate temperature range between 200 and 35OC a much smallerinfluence of Ts on hydrogen content and optical gap for fluorinated material than forhydrogenated a-SiGe. The higher photoconductivity (about an order of magnitude) influorinated material was attributed to a changed microstructure rather than to the saturation ofdangling bonds by fluorine. A particularly wide range of parameters was investigated by Morinet al [55,56]. For gas mixtures of purely fluorinated gases with H2 they reported somedisadvantages: a rather low deposition rate (e.g. r = 0.6 A/s at Eg = 1.25 eV), a strongpreferential incorporation of Ge, which had also been reported by Oda et al [50] and can giverise to sample inhomogeneity [54], and a tendency for the films to become microcrystalline.Better results (deposition rate of 3 - 5.5 A/s, no preferential Ge incorporation and no tendencytoward microcrystallinity) were obtained for SiH4-GeF4-H2 mixtures yielding device-gradematerial for Eg > 1.5 eV according to defect density and photosensitivity. Using the initialdefect density as a figure of merit, best films were grown at a process gas pressure between0.05 torr and 0.1 torr, a substrate temperature of between 300 and 35OC, an RF power (13.56MHz) between 7.5 and 10 W, an H2 flow between 5 and 20 seem, a GeF4 flow between 2 and10 seem, an SiH4 + SiF4 flow of 6.8 seem and an SiF4 flow below 1 seem [55]. Guha et al [57]reported high quality a-Si:Ge:H,F material according to PDS and SCLC measurements usinggas mixtures of Si2H6, GeH4, SiF4 and H2 at substrate temperatures exceeding 225C and withflow ratios of Si2H6-H2 > 1:10.

B3 Disilane

An advantage of using Si2H6 instead of SiH4 for a-Si