CdTe and Related Compounds; Physics, Defects, Hetero- and Nano-structures, Crystal Growth, Surfaces and Applications, 1st Edition

Elsevier Radarweg 29, PO Box 211, 1000 AE Amsterdam, The Netherlands Linacre House, Jordan Hill, Oxford OX2 8DP, UK First edition 2010 Copyright # 2010 Elsevier Ltd. All rights reserved. No part of this publication may be reproduced, stored in a retrieval system or transmitted in any form or by any means electronic, mechanical, photocopying, recording or otherwise without the prior written permission of the publisher. Permissions may be sought directly from Elsevier’s Science & Technology Rights Department in Oxford, UK: phone (þ44) (0) 1865 843830; fax (þ44) (0) 1865 853333; email: [email protected]. Alternatively you can submit your request online by visiting the Elsevier web site at http://elsevier.com/locate/permissions, and selecting Obtaining permission to use Elsevier material. Notice No responsibility is assumed by the publisher for any injury and/or damage to persons or property as a matter of products liability, negligence or otherwise, or from any use or operation of any methods, products, instructions or ideas contained in the material herein. Because of rapid advances in the medical sciences, in particular, independent verification of diagnoses and drug dosages should be made. Library of Congress Cataloging-in-Publication Data A catalog record for this book is available from the Library of Congress British Library Cataloguing in Publication Data A catalogue record for this book is available from the British Library ISBN: 978-0-08-096513-0 For information on all Elsevier publications visit our Web site at www.books.elsevier.com Printed and bound in Great Britain 10 11 12 10 9 8 7 6 5 4 3

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LIST OF CONTRIBUTORS

N. Audet 5N Plus Inc., 4405 rue Garand, Montreal, Canada H4R 2B4. K.W. Benz Freiburger, Materialforschungszentrum, Stefan-Meier-Strasse 21, D-79104 Freiburg, Germany. A.W. Brinkman Science Laboratories, University of Durham, South Road, Durham DH1 3LE, UK. J. L’Ecuyer 5N Plus Inc., 4405 rue Garand, Montreal, Canada H4R 2B4. M. Fiederle Freiburger, Materialforschungszentrum, Stefan-Meier-Strasse 21, D-79104 Freiburg, Germany. G.L. Herrit II-VI Incorporated, 375 Saxonburg Blvd., Saxonburg, PA 16056, USA. R.B. James Brookhaven National Laboratory, Upton, NY 11973, USA. Carl J. Johnson II-VI Incorporated, 375 Saxonburg Blvd., Saxonburg, PA 16056, USA. P.G. Kasherininov A.F. Ioffe Physico-Technical Institute, 26 Polytechnicheskaya Street, St Petersburg 194021, Russia. D. Lincot Laboratoire d’Electrochimie et de Chimie Analytique, UMR 7575, ENSCP-CNRS, 11 rue Pierre et Marie Curie, F-75231 Paris, France. J.-Y. Moisan 4 A Route Crech Argant, F-22730 Tregastel, France. E.R. Mueller Coherent Inc., DEOS, 1280 Blue Hills Avenue, Bloomfield, CT 06002, USA.

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x

List of Contributors

D. Shink 5N Plus Inc., 4405 rue Garand, Montreal, Canada H4R 2B4. V.N. Tomashik V.Ye. Lashkaryov Institute for Semiconductor Physics of National Academy of Sciences of Ukraine, 41 av. Nauki, Kyiv, Ukraine. Z.F. Tomashik V.Ye. Lashkaryov Institute for Semiconductor Physics of National Academy of Sciences of Ukraine, 41 av. Nauki, Kyiv, Ukraine. A.A. Tomasov A.F. Ioffe Physico-Technical Institute, 26 Polytechnicheskaya Street, St Petersburg 194021, Russia. R. Triboulet CNRS, GEMaC (Groupe d’Etude de la Matie`re Condense´e), 1 Place A. Briand, F-92195 Meudon Cedex, France. Ge Yang Brookhaven National Laboratory, Upton, NY 11973, USA.

FOREWORD

Thirty years after the remarkable monography of K. Zanio and the numerous conferences and papers dedicated since that time to CdTe and CdZnTe, after all the significant progresses in that field and the increasing interest in these materials for their extremely attractive fundamental properties and industrial applications, the editors have thought timely to edit a book on CdTe and CdZnTe, covering all their most prominent, modern, and fundamental aspects. The subject has become so wide and enriched during these 30 years that we have decided to call in well-known specialists and experts of the field. The editors would like to thank them deeply for their valuable contributions, with special acknowledgments to Dr Henri Mariette for his pertinent recommendations and his continued help and support. This part covers the topics Crystal Growth Technology and Surfaces, and Applications. The topics Physics, CdTe-Based Nanostructures and Semimagnetic Semiconductors, and Defects have been covered in Part I. R. Triboulet P. Siffert

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CHAPTER

I Crystal Growth and Surfaces

Contents

Ia. Technology: Purification of the Cadmium and Tellurium Elements 1. Introduction 2. Extraction and Initial Refining 2.1. Cadmium 2.2. Tellurium 2.3. Zinc 3. Final Refining 3.1. Cadmium 3.2. Tellurium 3.3. Zinc 3.4. Purification of the CdTe and CdZnTe Compounds 3.5. Oxygen Contamination 4. Concluding Remarks References Ib. CdTe and CdZnTe Growth 1. Introduction 2. Phase Equilibria in the Cd-Te System 3. Crystal Growth Versus Cd-Te Chemical Bond Characteristics 4. Synthesis 4.1. Liquid phase synthesis 4.2. Vapor phase synthesis 4.3. Solid state synthesis 4.4. Solution synthesis 5. Container 6. Crystal Growth 6.1. Melt growth 6.2. Solution growth 6.3. Vapor growth 6.4. Solid state recrystallization 7. Bridgman Growth Modeling and Interface Shape Determination

CDTE and Related Compounds

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4 4 5 5 7 9 9 9 12 14 14 15 16 16 19 19 19 23 27 28 30 30 30 35 36 36 41 44 48 49

2010 Published by Elsevier Ltd.

DOI: 10.1016/B978-0-08-096513-0.00001-7

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Crystal Growth Technology and Surfaces

8. CZT Properties 8.1. CZT properties at macroscopic and microscopic scale 8.2. Segregation 8.3. Solid-vapor equilibrium in the CdTe-ZnTe system 8.4. Industrial growth 9. Purity, Contamination, and Doping 10. Typical Structural and Electronic Properties of CdTe and CZT Crystals 11. Conclusions and Perspectives References Ic. Crystal Growth of CdTe/CdZnTe in Microgravity 1. Introduction: Crystal Growth Under Microgravity 2. Growth from the Vapour Phase 3. Growth by THM with a Rotating Magnetic Field 4. Bridgman Growth Using Dewetting Phenomenon 5. Summary and Outlook References Id. Heteroepitaxial Growth of CdTe Thin Films 1. Introduction 2. Overview of Deposition Methods 2.1. Molecular beam epitaxy 2.2. Hot wall epitaxy 2.3. Close space sublimation 2.4. Atomic layer epitaxy 2.5. Metal organic vapor phase epitaxy 3. Substrate Effects on CdTe Heteroepitaxy 3.1. Growth on Ge 3.2. Growth on Si 3.3. Growth on CdS 3.4. Growth on GaAs 3.5. Growth on NbSe2 4. Outline and Conclusions References Ie. Chemical Treatment of the CdTe and ZnxCd1
xTe Surfaces 1. Introduction 2. Bromine and Iodine Containing Etchant Compositions 3. Etchant Compositions Based on Nitric Acid 4. Etchant Compositions Based on Cr(VI) Compounds 5. Etchant Compositions Based on H2O2 6. Halogen-Evolving Etchant Compositions 7. Influence of Doping on Chemical Etching 8. Influence of Crystallographic Orientation on Chemical Etching

56 56 56 62 64 64 65 65 67 76 76 77 79 80 82 82 85 85 90 90 93 93 95 96 99 99 100 107 107 113 113 114 119 119 121 124 126 127 127 129 133

Crystal Growth and Surfaces

9. Chemical Etching of ZnxCd1–xTe Solid Solutions 10. Nanodimensional Formation on CdTe and Zn1–xCdxTe Surfaces at Chemical Etching 11. Conclusion References

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135 137 139 140

CHAPTER

IA Technology: Purification of the Cadmium and Tellurium Elements Jacques L’ Ecuyer, Nicholas Audet and Denis Shink

1. INTRODUCTION Suitable base elements, namely, cadmium, tellurium, and zinc, are required to synthesize and grow high-quality CdTe and CdZnTe crystals. Suitable in this instance implies pure to the extent required to avoid that the final device properties be dominated by extrinsic impurities coming from these base elements. Since the publication of the monograph on CdTe [1] there have been considerable improvements made in the purity of these base elements to the extent where 6Nþ (>99.9999%) base elements are now readily available from commercial vendors. Although denominated as 6Nþ, these are in many cases of even greater purity (9N) if one excludes the carbon, nitrogen, and oxygen content which typically range in the 50–500 part per billion atomic (ppba) level. Significant enough, these substantial improvements in the base element purity have not led to corresponding (expected) improvements in the purity of the CdTe and CdZnTe crystals largely because of remaining contamination issues resulting from the growth process itself [2, 3]. Breakthroughs in analytical techniques are perhaps the most important factor which have been responsible for the improvements in the base element purity over the last 3 decades. More specifically, the routine use of glow discharge mass spectrometry (GDMS) has enabled the direct measurement of purity down to ppba levels [4, 5] for a wide range of impurities in a rapid and reproducible way. This has enabled significant progress to be made in the purification process itself and highlighted the 5N Plus Inc., 4385 Garand Street Montreal, Canada H4R 2B4

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Technology: Purification of the Cadmium and Tellurium Elements

5

impact of environmental contamination resulting for example from inappropriate handling and storing practices of these base elements [6]. Commercial vendors and their customers alike are now well aware of such matters and have widely adopted the use of glove boxes, clean rooms, special packaging practices, and materials in an effort to maintain purity at the required levels. Cadmium, tellurium, and zinc, all being relatively low melting metals with high vapor pressures, are amenable to similar purification techniques, especially in the final stages. In most commercial practices these would include a combination of distillation (and/or sublimation) and zone-refining procedures which take advantage of differences in vapor pressure (distillation/sublimation), and changes in solubility between liquid and solid phases. Purity levels of 6Nþ can be obtained using such practices, as shown in Table 1, provided proper attention is given to the quality of the feedstock and appropriate steps are taken to minimize the impact of environmental contamination [7]. In what follows, we review the purification methods used for the base elements, namely, cadmium, tellurium, and zinc, and also examine the approach consisting in the purification of the compounds namely CdTe and/or CdZnTe.

2. EXTRACTION AND INITIAL REFINING Successful and economical purification of the base elements often depends on the suitability of the feedstock that is fed into the final purification stages. For this reason, an integrated purification scheme, involving control over both the initial and final refining stages, is preferred. In this way, it is possible to optimize the entire purification process so that impurities difficult to remove in the final purification stages may be removed earlier on.

2.1. Cadmium Cadmium is a relatively abundant element. World reserves are estimated at 6,000,000 MT [9], and approximately 20,000 MT are produced every year primarily as a by-product of zinc refining; although cadmium is also found in lead and copper ores. Typical cadmium concentrations in zinc concentrates are in the 0.1–0.5% range. Most of the cadmium is eventually recovered following roasting and leaching in sulfuric acid of the zinc concentrate. The resulting zinc sulfate solution, which typically contains up to 100 mg/l of cadmium, is cleaned using a zinc cementation process, whereby fine zinc powder is added to the electrolyte leading to the precipitation of

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Jacques L’ Ecuyer, Nicholas Audet and Denis Shink

Table 1

Typical purity levels of 6Nþ metals as determined by GDMS [8]

ppba

Cd

Te

Zn

ppba

Cd

Te

Zn

Li Be B C N O F Na Mg Al Si P S Cl K Ca Sc Ti V Cr Mn Fe Co Ni Cu Zn Ga Ge As Se

<3 <0.6 <0.7 460 35 220 <3 <0.9 <0.6 <0.5 <2 <0.6 <0.7 <4 <45 <4 <0.2 <0.1 <0.08 <0.4 <30 <25 <0.09 <0.9 <0.9 70 <0.5 <1 <10 <10

<9 <7 <7 200 65 400 <9 <5 <4 <2 <4 <4 <4 <4 <15 <30 <1 <0.7 <0.5 <3 <1 <1 <0.8 <2 <4 <12 <4 <9 <3 <20

<0.7 <0.6 <0.7 95 30 160 <5 <0.5 <0.4 <0.3 2 <0.4 <50 <0.7 <3 <4 <0.09 <0.08 <0.09 <2 <0.2 27 <0.05 <0.2 <60 Matrix <4 <2 <0.2 <5

Br Rb Sr Y Zr Nb Mo Pd Ag Cd In Sn Sb Te I Cs Ba La Ce Hf Ta W Pt Au Hg Tl Pb Bi Th U

<25 <0.2 <0.1 <0.1 <0.1 <0.6 <0.5

<2 <1 <0.8 <0.8 <0.7 <2

<5 <0.2 <0.1 <0.08 <0.07 <2 <0.4

<10 Matrix <45 <7 <0.4 <4 <0.9 <0.8 <1 <1 <2 <0.1

<5 <13 <2 <10 <55 Matrix <540 <4 <5 <0.2 <0.3 <1

<2 <3 <0.1 <0.4 <0.2 <1 <10 <30 <5 <0.1 <0.1 <0.07

<0.2 <1 <5 <2 <0.3 <0.2 <0.2 <0.7 <0.04

<1 <4 <10 <6 <2 <0.8 <1 <0.2 <0.2

<0.2 <1 <5 <2 <0.1 <0.06 <0.08 <0.02 <0.02

“<” denotes below the detection limit indicated.

cadmium and a corresponding reduction in the cadmium concentration of the solution down to 1 mg/l. The cadmium precipitate, which is also called a cadmium sponge, is relatively porous. Even after washing, it is still relatively impure containing high levels of sulfate and other impurities such as Cu (0.001%), Ni (<0.001%), Pb (0.002%), and Zn (0.4%). Subsequent operations typically involve melting of the sponge under a caustic soda slag, followed by casting. At this point, a good commercial quality metal 99.995% would result having as main impurities Al, Co, Fe,

7

Technology: Purification of the Cadmium and Tellurium Elements

Table 2

Composition of cadmium metal and slag following melting under NaOH [8]

(104 wt%)

Cd metal

Slag

(104 wt%)

Cd metal

Slag

Na Mg Al Si Ca Cr Mn Fe Ni

4 <0.02 0.3 <0.3 <0.1 <0.1 <0.01 <0.1 1

Matrix 944 18,820 8,046 211 3 43 430 3

Cu Zn Se Sn Sb Te Pb Bi Cd

1 <0.1 <0.4 <0.2 <0.4 <2 1 <2 >99.999%

<0.1 225 275 <0.2 3 <2 <0.5 <2 2%

“<” denotes below the detection limit indicated.

Ni, Cu, Zn, Sn, Tl, and Pb [7] as well as environmental contaminants such as N, O, C, Na, Si, and Ca. Other impurities that may also be present at significant levels include B, Cl, S, Cr, As, Se, Sb, Ag, Hg, and Bi. An alternative to zinc cementation for cadmium recovery involves electrowinning. This technique can lead to higher purity levels for the resulting commercial purity metal and has even been used as a means to attain purity levels of 99.999% or more [1]. Electrowinning is generally carried out in aqueous solutions of cadmium sulfate. The best results are obtained in solutions that have been previously purified using various hydrometallurgical techniques like selective chemical precipitation, ion exchange, and/or solvent extraction. To obtain a relatively pure cadmium deposit, all metals nobler than cadmium must be removed from the solution before electrolysis and in particular, the concentration of Cu, As, and Sb must be less than 1 mg/l. The cathodic cadmium deposit, which is generally in the form of sheets, dendrites, or small pieces, is melted under a layer of molten sodium hydroxide (NaOH) to improve purity by removing some of the main impurities as shown in Table 2, and also reduce losses of metal through volatilization.

2.2. Tellurium Tellurium is less abundant than cadmium and world reserves are not very well known [10]. Approximately 500 MT are produced every year primarily as a by-product of copper refining [11]. Tellurium is also found in sulfide ores of lead and zinc. Typical tellurium concentrations in copper concentrates are in the 0.01% range or lower as a significant portion of the tellurium present in the copper ore is lost in the tailings [11, 12]. Most of the tellurium in the concentrate eventually reports, in the form of intermetallic compounds of

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Jacques L’ Ecuyer, Nicholas Audet and Denis Shink

Ag, Cu, or Au, to the copper refinery slimes, which also contain most of the precious metals. Many approaches can be used to recover the tellurium from these slimes [13, 14], the most common involving leaching and decopperization of the slimes followed by copper cementation. The end product is a relatively pure copper telluride (Cu2Te), which is amenable to alkaline leaching and electrowinning. This step is typically carried out in alkaline solutions and involves reduction of tellurite ions (TeO32
) to elemental tellurium which plates on the cathode. The cathodic deposit is then melted and purified by fluxing with borax [15] or oxidizing salts such as NaOH. A relatively pure tellurium can be so obtained, in many instances of 99.999% purity, having as main impurities Na, Se, Pb, Ni, and Cu. To achieve this level of purity requires careful control of the electrolyte purity, the electrowinning conditions, and the fluxing process, the latter being very effective in removing many impurities, particularly those which form stable oxides, as shown in Table 3. Impurities that are not effectively removed by the fluxing process include the precious metals as well as Ni, Cu, Se, In, Hg, Tl, Pb, and Bi and efforts must therefore be taken to minimize their deportment to the tellurium cathodes. This is especially true for Se, which is very difficult to remove in the latter stages of the purification process (final refining), but which can be removed from the electrowinning solution using hydrometallurgical techniques [14, 15]. Table 3

a

Distribution of impurities in molten tellurium in contact with NaOH [8]

(wt%)

NaOH slag (%)

Te metal (%)

Be B Na Mg Al Si Sa Ca Ti V Cr Mn Fe Ni Cu Zn

88 99 98 93 98 84 >99 66 95 99 >99 >99 >99 42 1 94

12 1 2 7 2 16 0.01 34 5 1 0.03 0.1 0.05 58 99 6

As sulfate (SO4).

(wt%)

NaOH slag (%)

Te metal (%)

Ga Ge As Se Zr Mo Ag Cd In Sn Sb Hg Tl Pb Bi

98 93 >99 1 96 99 3 76 0.2 >99 99 0.2 11 83 79

2 7 0.1 99 4 1 97 24 >99 0.4 1 >99 89 17 21

Technology: Purification of the Cadmium and Tellurium Elements

9

2.3. Zinc Zinc is one of the most common metals with several million tons being produced every year (over 10 million in 2006). Today, most zinc metal is produced using the so-called hydrometallurgical process, which yields electrolytic-grade zinc, is typically of 99.995% purity (special high grade, SHG). Main impurities include Pb, Cd, Fe, Sn, Al, Cu, Tl, and In.

3. FINAL REFINING Distillation and zone refining of the base elements have been, and remain, the preferred approaches for final refining to reach the required purity levels (>99.9999%) for electronic applications [7]. Both are carried out in quartz, graphite coated quartz, or in high-purity graphite crucibles. Distillation is performed under vacuum to reduce the operating temperature. In order to evaluate the potential impact of distillation and zone refining, we show, in Figs. 1 and 2, the vapor pressure ratios (with respect to the base element of interest) and segregation coefficients for typical impurities, an approach adopted by Ali and coworkers [16]. These graphs should be interpreted with care as they do not account for possible intermediate phases, for interactions between impurities (an effect which can be significant even at low impurity concentrations [17]), and in a more general way for deviations from ideality. Furthermore, purification efficiencies for distillation, calculated from the vapor pressure ratios, and those for zone refining inferred from equilibrium segregation coefficients (K0), must also be corrected for kinetic effects that can significantly degrade the achievable calculated purification efficiencies (leading in the case of zone refining to the effective segregation coefficient (Keff)). Quantitative interpretation of the data presented in Figs. 1 and 2 is thus subject to these limitations.

3.1. Cadmium Most of the typical impurities found in cadmium can be effectively removed by distillation as illustrated in Fig. 1 because of the relatively high vapor pressure of cadmium. This has indeed been demonstrated by several groups [18–21] which have shown that high-purity cadmium can be produced by distillation. Some common impurities such as Na, Zn, As, and Hg are more difficult to remove however as they have comparable (or only marginally higher) vapor pressures than that of cadmium. For these impurities, rectification techniques have been proposed [22] as well as distilling in the presence of cadmium oxide which is claimed to act as a getter for impurities having a greater oxygen affinity, such as zinc [18].

0

5

−5

4

−10

3

−15

2

−20

1

−25

0

−30

–1

−35

–2

−40

–3

−45

–4

−50 Br

–5

I Hg

P S

Cs Rb Se As

K

Y C U V B Pt Zr Hf Li Sr Sb Ca Tl Ba Bi Pb In Sn Be Al Cu La Au Sc Pd Cr Fe Ni Co Cd Na Zn Te Mg Ge Si Ce Ti Th Nb Ta W Mn Ga Ag Mo

Vapor pressure ratio@450 ⬚C

K0 (from phase diagrams [49])

Figure 1 Segregation coefficients (K) and vapor pressure ratios in cadmium.

K0 [50]

Keff@30 mm/h[8]

Log (K)

6 Jacques L’ Ecuyer, Nicholas Audet and Denis Shink

log (vapor pressure ratio)

10

5

0

5

−5

4

–10

3

–15

2

–20

1

–25

0

–30

–1

–35

–2

–40

–3

–45

–4

–50 Br

–5

I Hg

P

S

Cs Rb Se As

K

Y Li Sr Sb Ca Tl Ba Bi Pb In Mn Ga Ag Sn Be Al Cu La Au Sc Pd Cr Fe Ni Co Cd Na Zn Te Mg

Vapor pressure ratio@650 ⬚C K0 (from phase diagrams [49])

Ti Th U V Ge Si Ce

Keff@35mm/h [34]

B

Pt Zr Hf

C Mo

Keff [1]

Nb Ta W

Log (K)

log (vapor pressure ratio)

6

Technology: Purification of the Cadmium and Tellurium Elements

5

Figure 2 Segregation coefficients (K) and vapor pressure ratios in tellurium.

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Jacques L’ Ecuyer, Nicholas Audet and Denis Shink

Se, Te, and S also have similar vapor pressures to that of cadmium but the actual distillation purification efficiency for these impurities is much larger than calculated because of the very low activity coefficients of these elements in liquid cadmium [23], which reduces their vapor pressure. Indeed, these elements form extremely stable compounds with cadmium, namely, CdTe, CdSe, and CdS, which all have relatively low vapor pressures. Figure 1 also illustrates that many of the critical impurities such as Cu, Fe, Ni, Co, Ag, and Au do not segregate very effectively and that on this basis alone, distillation should be preferred over zone refining. Further support for distillation over zone refining comes from the fact that zone refining of cadmium is not straightforward because of its relatively high thermal conductivity and low melting point [24]. Both create difficulties in controlling the width of the molten zone, and hence its speed at the interface where solidification occurs, and increase the likelihood of constitutional supercooling. Nevertheless, zone refining has been used extensively [7, 16, 24–29] and current commercial practices involve a combination of both treatments, typically carried out under hydrogen in the following sequence, distillation followed by zone refining [1, 7, 28]. In this way, it is generally possible to remove all impurities and attain the desired 6Nþ levels in a cost effective manner provided an adequate feedstock is selected. Adequate in this instance would imply of a minimum purity of 99.995% (preferably 99.999%) with low levels of Na, Zn, and Hg because both distillation and zone refining are not very effective in removing them. Some workers claim that the cadmium zone-refining purification efficiency can be significantly improved by the addition of minute amounts of tellurium or CdTe which increase the segregation coefficients of many impurities [24]. Variations to the standard zone-refining techniques have also been used or considered, such as the overlap zone-melting method (OZM) [29], and the use of an electric field [30]. Combining both distillation and normal freezing [31] or zone refining in a so-called in situ approach [16, 32] has been claimed to lead to higher levels of purity. In all likelihood, this claimed improvement in purity can be correlated with a reduction in the number of handling steps. Other work also revealed inprocess copper contamination resulting from the use of what is now known to be an inappropriate gas delivery system [33].

3.2. Tellurium As shown in Fig. 2, impurities tend to segregate more effectively in tellurium than cadmium. Furthermore, all impurities are more soluble in the liquid phase (segregation coefficient is less than 1), and therefore

Technology: Purification of the Cadmium and Tellurium Elements

13

none tend to accumulate at the leading end of the ingot. Tellurium is also a relatively poor conductor of heat and has a higher melting point, which facilitates zone stabilization and reduces the tendency for interface breakdown and constitutional supercooling. Zone refining would therefore be the preferred purification technique although commercial practices generally involve a combination of both vacuum distillation and zone refining under hydrogen, in much the same way as for cadmium. Over the last 3 decades, considerable work has been carried out on the distillation and zone refining of tellurium [6, 7, 24–26, 32, 34–40]. It has been shown that it is generally possible to remove all impurities and attain the desired 6Nþ levels in a cost effective way provided an adequate feedstock is selected. Adequate in this instance would imply of a minimum purity of 99.99% (preferably 99.999%) with low levels of Se and S because both distillation and zone refining are not very effective in removing these elements. Removal of selenium can be more easily carried out via a reaction with hydrogen to form hydrogen selenide (H2Se). The reaction is as follows: H2þSe ! H2Se, for which the H2Se to H2 ratio lies in the 0.7–0.99 range for typical operating temperatures [41]. Under typical zone-refining conditions, this reaction is limited by the saturation of H2Se so that the residual selenium concentration (C) can be calculated from the following equation: log

C ¼
5:22  Vð1Þ  MðgÞ; C0

where C0 is the initial selenium concentration, V(l) is the volume in liters of hydrogen circulated and M(g) the weight in grams of tellurium being purified. Graphite inclusions are claimed by some workers to be a major contaminant of commercial purity tellurium and they report that distillation is an effective means for removing them [42]. In much the same way as for cadmium, it is also claimed that the zone-refining purification efficiency of tellurium can be improved by the addition of minute amounts of cadmium or CdTe [24, 37],1 and that combining both distillation and zone refining leads to higher levels of purity [32]. The use of an electric field to improve the purification efficiency has also been considered [30, 43]. Song [44] has used selected portions of tellurium ingots subjected to a normal freeze treatment and shown a further improvement in purity.

Although the reported segregation coefficients for K (Keff ¼ 0.9), Ga (Keff ¼ 0.88), and As (Keff ¼ 0.77) are very close to unity.

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Jacques L’ Ecuyer, Nicholas Audet and Denis Shink

3.3. Zinc There has been considerably less work [32, 45–47] carried out on the purification of zinc than on cadmium or tellurium. Based on the vapor pressure ratios, one would expect a similar behavior to that of cadmium with impurities such as Na, Cd, As, and Hg being most difficult to remove by distillation. Zone refining of zinc is easier than that of cadmium because the thermal conductivity is lower and the melting point higher which facilitates control of the molten zone width and decreases the likelihood of constitutional supercooling. However, zone refining of zinc remains fundamentally subject to the same limitations as cadmium, namely, marginal segregation efficiency for many impurities,2 with even more impurities (Al, Cu, Pd, Ag, and Au) having segregation coefficients in excess of 1 [46]. Current commercial practices involve a combination of both distillation and zone-refining treatments under hydrogen, as for tellurium and cadmium. Properly conducted, these yield the desired 6Nþ purity levels. Combining both distillation and zone refining in the in situ approach [32] is claimed to lead to higher levels of purity. Zone refining of a cadmiumzinc eutectic at our facility has also been attempted in an effort to improve post processing of the material.

3.4. Purification of the CdTe and CdZnTe compounds Although purification of the base elements is relatively straightforward and can provide the required purity levels, it may be appropriate in some instances to purify the compounds in combination with, or to some extent instead of, these very same base elements. This is particularly justifiable in the case where specific impurities do not segregate very effectively in one of these base elements. The purification efficiency will be largely dependent on the processing conditions and, for example, whether crystallization occurs under tellurium rich conditions or not. We compare in Table 4, the segregation coefficients for Li, Ti, Fe, Cu, Ag, and Au in cadmium and zinc to those in CdTe and CdZnTe. As can be seen, these impurities segregate much more effectively in CdTe or CdZnTe than in cadmium or zinc. With the exception of Li, however, all of these impurities can be effectively removed by distillation in both cadmium and zinc. Although this would suggest that purification of the compound is somewhat less relevant, one should consider that it provides means for compensating inprocess and environmental contamination of the base elements and can therefore ultimately lead to higher levels of purity. This is especially true

For example, Pb: Keff ¼ 0.6 at 30 mm/h and 0.8 at 60 mm/h [8].

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Technology: Purification of the Cadmium and Tellurium Elements

Table 4

Li Ti Fe Cu Ag Au

Segregation coefficients for selected impurities in Cd, Zn, CdTe, and CdZnTe Cd [49]

Zn [45]

CdTe [50]

CdZnTe [51]

2.62 <0.1 <0.1 0.5 [8] 2.52 0.55

0.35 <1 0.95 [46] >1 2.5 >1

0.3 0.054 0.0003 0.029 0.003 0.04

0.088 <1 [33] 0.006 0.038 – –

when purification of the compounds is carried out using low-temperature processes such as the traveling heater method (THM) [2, 48].

3.5. Oxygen contamination All of the base elements are subject to oxygen contamination which promotes adhesion (sticking) to the quartz crucibles through the formation of silicates. For this reason, the use of graphite-coated quartz crucibles is preferred especially for zinc and cadmium, which are more difficult to rid of oxides than tellurium. High-purity graphite or boron nitride crucibles can also be used but both have drawbacks including cost and high thermal conductivity. Oxygen is also relatively soluble in the liquid base elements as shown in Table 5. To avoid excessive oxygen contamination all final purification steps are generally carried out under a flow of hydrogen, which acts both as a blanket gas preventing (or reducing the impact of) oxygen diffusion inside the reactor and as a reducing agent media for oxides and dissolved oxygen. Thermodynamic calculations indicate that both cadmium and tellurium oxide can be easily reduced by hydrogen (H2) at typical zonerefining temperatures, but much less so for zinc oxide. For example, the H2O(g)/H2(g) ratios at equilibrium with the base elements and their Table 5

Solubility of oxygen (104 at%)

Temperature ( C)

Cd [52, 53]

Te [52, 54]

Zn [55]

321 420 450 500 550 600

0.043 1 3.2 11 35 95

– –

– 0.007 0.021 0.097 0.37 1.2

23 51 105 197

16

Jacques L’ Ecuyer, Nicholas Audet and Denis Shink

respective oxides, at the base element melting point, are 4E
6, 15, and 80E þ6 for zinc, cadmium, and tellurium respectively. Palosz [56] has studied the kinetics of oxide reduction with both hydrogen and graphite for a number of metals and nonmetals including zinc, cadmium, and tellurium. His results indicate that reduction with hydrogen is faster and occurs at lower temperatures than with graphite. The reaction is still relatively slow, however,3 and hence all necessary precautions should be taken to avoid oxygen contamination, and therefore the need for its removal.

4. CONCLUDING REMARKS The purification technology has now reached a point where suitable quality precursors for growth of high-quality CdTe and CdZnTe can be produced on a routine basis. Proper control of the entire purification process, including the primary refining steps, as well as the use of appropriate facilities designed to minimize the impact of environmental contamination are essential in order to do so. Remaining contamination issues resulting from the CdTe and CdZnTe growth process itself and the impact of native defects now appear to be limiting factors.

REFERENCES [1] K. Zanio, Semiconductors and Semimetals, vol. 13, Academic Press, New York and London, 1978, pp. 38–52. [2] A. Mokri, R. Triboulet, A. Lusson, A. Tromson-Carli, G. Didier, J. Crystal Growth 138 (1994) 168–174. [3] R. Triboulet, A. Aoudia, A. Lusson, J. Electron. Mater. 24 (1995) 1061–1065. [4] N. Sanderson, E. Hall, J. Clark, P. Charalambous, Mikrochim. Acta 91 (1987), 275–290. [5] A. Mykytiuk, P. Semeniuk, S. Berman, Spectrochim. Acta Rev. 1 (1990) 1–10. [6] D. Prasad, C.h. Sudheer, N. Munirathnam, T. Prakash, Bull. Mater. Sci. 25 (2002) 545–547. [7] H. Hirsch, S. Liang, A. White, Semiconductor and Semimetals, vol. 18, Academic Press, New York and London, 1981, pp. 21–45. [8] 5N Plus Inc., Internal data. [9] W. Butterman, J. Plachy, US Geological Survey Open-File Report 02-238, 2002. [10] J. Hein, A. Koschinsky, A. Halliday, Geochim. Cosmochim. Acta 67 (2003) 1117–1127. [11] F. Ojebuoboh, Proc. EMC 2007 (2007) 571–585.

3 From the data of Palosz [56], the maximum reduction rate for solid CdO at 340  C is estimated to lie in the 0.1-3 mg range of atomic oxygen per cm2 per min. This equates to an oxygen removal rate, at or above saturation, of between 0.15  10
4 and 5  10
4 wt% per day in a typical 20 kg zone refined ingot of cadmium (four times more for a tellurium ingot at 645  C) subjected to a continuous flow of hydrogen. Below saturation, the oxygen removal rate should be lower as the activity of oxygen is decreased.

Technology: Purification of the Cadmium and Tellurium Elements

17

[12] D. Chizhikov, V. Shchastlivyi, Tellurium and Tellurides, Collets’s Ltd., London and Wellingborough, 1970, p. 45. [13] J. Jennings, in: W.C. Cooper (ed.), Tellurium, vol. 14, Van Nostrand, New York, 1971. [14] J. Hoffmann, JOM 41 (1989) 33–38. [15] R. Hisshion, R. Casauay, Proceedings of STDA’s 5th International Symposium, 1994, pp. 59–62. [16] S. Ali, R. Reddy, N. Munirathnam, C. Sudheer, G. Anil, T. Prakash, Separ. Purif. Technol. 52 (2006) 288–294. [17] R. Von Kujawa, Z. Phys. Chem. 232 (1966) 425–431. [18] S. Kovalevski, V. Kosyakov, I. Shelpakova, J. Crystal Growth 167 (1996) 208–211. [19] L. Kozin, E. Berezhnoi, K. Kozin, Russ. J. Appl. Chem. 71 (1998) 755–761. [20] S. Ali, J. Rao, K. Varma, T. Prakash, Bull. Mater. Sci. 25 (2002) 479–481. [21] S. Ali, N. Munirathnam, C. Sudheer, R. Reddy, T. Prakash, Mater. Lett. 58 (2004) 1638–1641. [22] L. Kozin, E. Berezhnoi, K. Kozin, Inorg. Mater. 35 (1999) 795–799. [23] R. Sharma, Y. Chang, J. Electrochem. Soc. 5 (1989) 1536–1542. [24] R. Bult, A. Bollong, R. Redden, US Patent No. 4690725, (1987). [25] V. Marychev, Russ. Metall. 4 (1985) 30–36. [26] R. Redden, R. Bult, A. Bollong, Warrendale, PA, TMS (1986) A86–60. [27] Y. Ishikawa, Y. Bailiang, K. Mimura, T. Tomizono, J. Min. Mater. Process. Inst. Jpn. 110 (1994) 1175–1178. [28] N. Munirathnam, D. Prasad, C.h. Sudheer, J. Rao, T. Prakash, Bull. Mater. Sci. 28 (2005) 209–212. [29] J. Wang, S. Song, Y. Ishikawa, M. Isshiki, Mater. Sci. Eng. 117 (2005) 271–275. [30] S. Dost, Y. Liu, J. Haas, A. Roszmann, S. Grenier, N. Audet, J. Crystal Growth 307 (2007) 211–218. [31] L. Holland, J. Crystal Growth 70 (1984) 280–286. [32] G. Feldewerth, A. Bollong, D. Bunnell, US Patent No. 5,513,834, (1996). [33] A. Bollong, G. Feldewerth, J. Tower, S. Tobin, M. Kestigian, P. Norton, F. Schaake, C. Ard, Adv. Mater. Opt. Electron. 5 (1995) 87–93. [34] A. Zaouir, M. Hage-Ali, J. Koebel, A. Bentz, P. Siffert, Phys. Status Solidi A 100 (1987) K139–K143. [35] A. Zaouir, M. Hage-Ali, J. Koebel, P. Siffert, Mater. Sci. Eng. B 3 (1989) 331–334. [36] A. Zaouir, K. Zahraman, M. Roumie, J. Charara, A. Fawaz, F. Lmai, M. Hage-Ali, Mater. Sci. Eng. B 131 (2006) 54–61. [37] A. Bollong, R. Bult, JOM 41 (1989) 39–41. [38] N. Munirathnam, D. Prasad, C.h. Sudheer, A. Singh, T. Prakash, Bull. Mater. Sci. 25 (2002) 79–83. [39] N. Munirathnam, D. Prasad, C.h. Sudheer, T. Prakash, J. Crystal Growth 254 (2003) 262–266. [40] S. Ali, D. Prasad, N. Munirathnam, T. Prakash, Separ. Purif. Technol. 43 (2005) 263–267. [41] J. Rawling, J. Toguri, Can. J. Chem. 44 (1966) 451–456. [42] M. Churbanov, V. Gerasimenko, V. Shiryaev, Inorg. Mater. 37 (2001) 1017–1020. [43] M. Lopez, L. Colombo, M. Brau, in: Proceedings of Symposium on Electromigration of Metals, (1984), pp. 35–42. [44] S. Song, J. Wang, M. Isshiki, J. Crystal Growth 236 (2002) 165–170. [45] B. Aleksandrov, Fiz. Metal. Metalloved 4 (1961) 588–595. [46] V. Grigor’yev, V. Vigdorovich, A. Yaroslavtsev, Russ. Metall. 4 (1973) 48–49. [47] L. Rowinska, L. Walis, W. Dalecki, M. Kusowski, Nukleonika 31 (1986) 151–162. [48] N. Audet, M. Cossette, J. Electron. Mater. 34 (2005) 683–686. [49] J. Drapala, L. Kuchar, Burkhanov G Inorg. Mater. 34 (1998) 114–127. [50] L. Kuchar, J. Drapala, J. Lunacek, J. Crystal Growth 161 (1996) 94–103.

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[51] M. Schieber, R. James, H. Hermon, A. Vilensky, I. Baydjanov, M. Goorsky, T. Lam, E. Meerson, H. Yao, J. Erickson, E. Cross, A. Burger, J. Ndap, G. Wright, M. Fiederle, J. Crystal Growth 231 (2001) 235–241. [52] S. Otsuka, Z. Kozuka, Trans. Jpn. Inst. Metals 22 (1981) 558–566. [53] H. Wriedt, Bull. Alloy Phase Diag. 8 (1987) 140–147. [54] H. Wriedt, Bull. Alloy Phase Diag. 8 (1987) 166–176. [55] S. Otsuka, Z. Kozuka, Met. Trans. 11B (1980) 119–124. [56] W. Palosz, J. Crystal Growth 173 (1997) 427–439.

CHAPTER

IB CdTe and CdZnTe Growth R. Triboulet

1. INTRODUCTION Since last 40 years and the encyclopedic and pioneering work of de Nobel [1], considerable efforts have been dedicated to the crystal growth of CdTe. In spite of its “moderate” melting temperature and partial vapor pressures of its components, the crystal growth of CdTe turns out to be extremely tricky, mainly its melt and vapor growth. The reproducible growth of large CdTe single crystals of stoichiometric composition, and then of high crystallographic quality, remains an issue even if spectacular breakthroughs have been achieved during the recent years in the reproducible growth of large oriented Te-rich CdTe and CdZnTe (CZT) single crystals of industrial size by seeded traveling heater method. In this chapter, the CdTe crystal growth is reviewed in close relationship with its properties, mainly those that come directly from the characteristics of the Cd-Te chemical bond.

2. PHASE EQUILIBRIA IN THE Cd-Te SYSTEM Crystal growth of CdTe requires precise knowledge of the existence regions of the solid, liquid, and gaseous phases with respect to temperature, pressure, and composition. For this purpose, the temperature versus composition, T-x, component pressure or total pressure versus temperature, p-T or P-T, and finally p-T-x or P-T-x diagrams have been experimentally determined and theoretically modeled. In order to control the stoichiometry of the crystals, the deviations from which can affect strongly their semiconducting properties, the knowledge of the extent of

CNRS GEMaC (Groupe d’Etude de la Matie`re Condense´e) 1 Place A. Briand F-92195 Meudon Cedex France

19

20

R. Triboulet

the homogeneity domain and of the position of the stoichiometric line inside the three-phase boundary is essential as well. The first liquidus data were obtained either from visual observation of the onset of initial freezing [1, 2], technique which is usually considered to have a certain uncertainty because of its subjective nature, or by thermal analysis [3, 4] and differential thermal analysis [5]. Such measurements as a function of a component partial pressure allowed to establish the first p-T diagrams [1, 4], which were investigated as well by measuring the optical density of the vapor to determine the component partial pressures [6]. P-T diagrams were first drawn using a regular associated solution model, which postulates the existence of stable CdTe complexes in the liquid phase and shows an interchange energy becoming infinite at equiatomic composition [7]. The T-x diagram was drawn from an improved approach in which liquid-mixing functions are expressed in an association model by using the Redlich-Kister polynomia form of species interactions [8]. The extent of the electrically active homogeneity region of CdTe was found from electrical measurements interpreted by simple point defect models to be roughly about 1017 cm
3 on both Cd and Te sides and the solidus to be of retrograde shape [1, 9–13]. Overestimated values of stoichiometry deviations, besides those obtained from electrical measurements, have been reported by several authors. The deviation from stoichiometry was suggested to be as large as 2  1019 cm
3 on the Cd side and 5.7  1019 cm
3 on the Te side from density measurements [14]. From calculations performed on the basis of the measurements of the Te vapor pressure above the vapor pressure for CdTe crystals with various limiting deviations from stoichiometry, the existence of a homogeneous region amounting to 1% at 800  C was stated [15]. More than 1020 cm
3 on both sides was estimated from a special mass spectrometry method and optical vapor density measurements [16]. From new calorimetric investigations, the enthalpy of formation of CdTe DfH0 and the standard entropy S0298 at 298.15 K were found to be
100,708 and 93.3 J/mol K, respectively [17]. The Cd-Te phase diagram was then revisited from total vapor pressure scanning experiments [18]. The determination of the composition XS, XV, and XL of the solid, vapor, and liquid phases in equilibrium allowed drawing solidus, vaporus, and liquidus curves of the Cd-Te system. Very precise P-T (Fig. 1) and T-X (Fig. 2) projections of the Cd-Te diagram were constructed, as well as several isobaric sections of the P-T-X phase diagram, allowing to follow the crystallization of nonstoichiometric CdTe under various P-T conditions from different matrices (liquid, vapor, or both) or annealing of the prepared material. Several significant conclusions can be drawn from the study of Greenberg: CdTe shows a congruent fusion (cf) with a maximum melting point of

103 / T K

CdTe and CdZnTe Growth

21

) Te

1 2 3 4 5 6 7 8 9

( LV

0.7

VSL SVL

0.8

s=v

d)

C

( LV

0.9

1.0 S

XS Max S=L

L=V

T

Tcf

V

1.3

L=V

1.2

XLMax

XV Max

S=L

TMax

Tcs

K

N M

1.4

L

Cd 10

100

L

V SL

S

X

Te

S

S=V

=V V S Q

VL

1.1

VL

V SL

P 1000

P, mm Hg

Figure 1 P-T projection of the Cd-Te diagram. The inset shows P-T and T-x projections of the CdTe melting region (from Ref. [18]).

1092  C; at Tcf, CdTe is Te-saturated; at Tcs (cs: congruent sublimation), crystalline CdTe is also Te-saturated; Tcs < Tmax by 41 K; stoichiometric CdTe is in equilibrium with a Te-rich melt and a virtually pure Cd vapor; there is no constant cs composition; the cs line lies on the Te-rich side, rather far from the stoichiometry; the maximum Te nonstoichiometry is of about 4  1018 cm
3, the maximum Cd nonstoichiometry of about 1018 cm
3. Most of the results of Greenberg were confirmed by Fang and Brebrick, from optical density measurements, who nevertheless found the solidus curve XTe-1/2 extending to 10
4 to compare with 1.5  10
4 in the case of Greenberg [19]. Their results throw a new light on the stoichiometric line determination: a singly ionizable Te antisite model gives a good fit of the high-temperature electrical measurements [20].

22

S=L

R. Triboulet

L= V

1400

1300

S=V

1200

1100

1000

Cd-solidus Te-solidus Liquidus Vaporus Congr.subl.

900

800

700

0

20

40 49.996 50 50.004

50.012

60

80

100

X, at. % Te

Figure 2 T-x projection of the Cd-Te diagram. Near-solidus region is given on an enlarged scale (from Ref. [18]).

Greenberg showed furthermore how to apply his P-T-X phase equilibrium data to prepare crystals with controlled composition, either stoichiometric or with a certain deviation from stoichiometry, according to different technologies, vapor phase growth, vertical, horizontal, and highpressure Bridgman (HPB) [21]. In vapor growth, the growth rate is linked to the composition of the vapor and the maximum rate corresponds to the

CdTe and CdZnTe Growth

23

stoichiometric composition. CdTe physical vapor deposition (PVD) is recommended at temperatures well below the maximum cs temperature Tcs ¼ 1324 K: the composition of the crystals becomes less temperature dependent below 1173 K, while at higher temperatures, small deviations in the vapor composition or small pressure variations will favor the vapor-liquid-solid (VLS) mechanism with detrimental consequences. For the conventional CdTe PVD temperature range (1123–1223 K), the composition of the crystals changes from 50.00083 to 50.0015 at.% Te. In the vertical Bridgman (VB) method, the growth occurs at the near-maximum melting temperature from the liquid þ solid state of the system. The crystals grown from the VLS equilibrium are Te rich, and moreover at the congruent melting point. To reduce the Te excess, the pressure should be increased above a pressure of 1.36 atm (up to about 2 atm according to Rudolph et al. [22]), either by using a Cd reservoir with independent temperature control in order to enrich the liquid in Cd or under an inert gas overpressure of over 100 atm [23]. In the horizontal Bridgman (HB) method, the growth occurs from the VLS state, allowing a flexible control of the composition. The three-phase VLS equilibrium being univariant, the crystallization temperature, as well as the composition of the crystals, become fixed when a vapor pressure is chosen. The vapor is essentially pure Cd. CdTe crystals can be grown from Cd-rich melts at vapor pressures higher than the one corresponding to TMAX, P(TMAX), or from Terich melts for pressures lower than P(TMAX). The results of Greenberg can be used as well to adjust postgrowth annealing conditions. Note that the so-called “modified” HB method is the only crystal growth technique allowing to grow perfectly stoichiometric crystals.

3. CRYSTAL GROWTH VERSUS Cd-Te CHEMICAL BOND CHARACTERISTICS The fairly high ionic character of the II-VIs chemical bond signifies not only most of the physical properties of these semiconductors but also their crystal growth. The higher the ionicity, the lower is roughly the thermal conductivity. The low CdTe thermal conductivity of 55 mW/cm
1/K
1 makes it difficult to control the solid-liquid interface in both melt and vapor growth. Pretransition phenomena on both sides of the melting point, in both liquid and solid states, have been found to occur in CdTe as a result of the ionic character of the Cd-Te chemical bond. This character of the bond makes CdTe melts highly associated close to the melting point, leading to the presence of highly organized particles influencing nucleation process and growth kinetics. The theory of associated solution model for liquid CdTe expresses the strong interaction

24

R. Triboulet

between Cd and Te atoms in the liquid state [24]. The melting process of CdTe was assumed to be a solid semiconductor to liquid semiconductor transition from experimental investigations of the melt parameters [25] and from the positive slope of the conductivity versus temperature in the region of the solid phase [26]. Above the melting point, the liquid-phase conductivity was found again to increase exponentially with temperature from eddy-current measurements [27]. Using the eddy current technique as well to measure the variation of the electrical conductivity versus temperature, transitions occurring between 1025 and 1050  C were suggested to be associated with allotropic phase transitions [28]. The mole fraction of associated species in the melt just above the CdTe melting point, or association coefficient, was estimated from calculations to be 0.908 [8], 0.95 [29], and 0.92 from the new calorimetric data [17]. The enthalpy of formation of the melts, determined from calorimetric measurements, shows very strong departures from ideality [30]. The authors explain this result by the existence of short-range ordering corresponding to CdTe associates. From neutron diffraction experiments, it was suggested that the structure of liquid CdTe could be roughly described by a four-coordinate random network, very different from the structure of group IV elements or III–V compounds [31]. From DTA measurements, endothermic effects above the melting point were observed (Fig. 3) [32, 33]. A high degree of structural ordering was concluded by the author, confirmed by molten CdTe viscosity measurements both on heating and cooling [34]. From DTA measurements, it was shown that the CdTe melt consists of two types of structures: clusters (with a short-range order close to the crystalline substance) within an amorphous matrix [35].

T,K 1373 K

1490 K

1379 K 1393 K

1391 K

1373 K 1373 K 1363 K

1363 K 1320 K

A

1363 K

B

1334 K

1379 K

C ΔT, K

Figure 3 Typical differential thermograms of the pure and doped CdTe melting process: (a) rapid heating to Tm (t ¼ 10–15 min) of the sample equilibrated at room temperature; (b) thermocycling in the 950–1150  C temperature range; (c) slow heating rate (Vh < 1 K/ min) of a CdTe sample previously excited by thermocycling (from Ref. [33]).

CdTe and CdZnTe Growth

25

40

supercooling ΔT –/K

30

20 Tm

V ΔT –

10

ΔT + Tm Tm

10

20

30

40

50

superheating ΔT +/K

Figure 4 Influence of the degree of superheating (DTþ ¼ Tm þ T) on the degree of supercooling (DT
¼ Tm
T) for stoichiometric CdTe measured with a thermocouple at the tip of vertical Bridgman ampoules moved with a constant velocity v ¼ 1 mm/min: D, supercooling of former high superheated melts subsequently cooled and kept for 5 h at the melting point Tm (from Ref. [36]).

From the activation energies of the crystallization process, the authors found that the clusters consist of about 66–98 atoms in pure molten CdTe. Because of the presence of such highly organized particles in the melt, no supercooling is observed at small superheating, while a large supercooling occurs at superheating values higher than 9–10 K (Fig. 4) [33, 36, 37]. Some duality supercooling-superheating is then suggested. At small superheating, parasitic nucleation occurs from the highly organized particles present in the melt: if a single crystal initial growth is often observed (Fig. 5a), grain boundaries and twins are originated during further growth. Higher superheating results in a polycrystalline first-to-freeze region followed by a single crystal part (Fig. 5b). However, such high superheating impedes the possibility of performing seeding: large superheating values are necessary for the destruction of associated melt complexes, but as a result seed melting becomes unavoidable. Furthermore, from an investigation of the electrical conductivity in liquid and solid CdTe near its melting point [38] the degree of supercooling was found to decrease with increasing Cd overpressure and to reach the lowest value at 1.6 atm. In solid state, the higher the ionicity, the higher is the tendency to the hexagonal structure and thus to twinning. The Cd-Te chemical bond shows a pretty high ionicity of 0.55. The tendency to twinning and to

26

R. Triboulet

3

2

1

A

B

Figure 5 Typical growth structures of (a) a high (DTþ ¼ 27 K) and (b) a low (DTþ ¼ 3 K) superheated unseeded CdTe crystal: 1, polycrystalline tip region; 2, twin lamella; 3, large grain boundary (from Ref. [36]).

hexagonal structure appears usually in such ionocovalent compounds when ionicity exceeds 0.5. Ionicity was already referred to to account for the high CdTe twinning tendency according to the classical inverse relationship found between stacking fault energy and ionicity [39]. The existence in CdTe crystals of a hexagonal structure at high temperature was reported by several authors [40–42]. A phase transition in the 893–920  C temperature range was pointed out [43]. This tendency of CdTe to twinning and to hexagonal structure appears also through the behavior of the CdMnTe alloys in which a transition in the solid phase was found between a phase of hexagonal structure at high temperature and cubic one at lower temperature [44, 45]. The transition temperature decreases with increasing Mn content in agreement with the increase of bond ionicity with the incorporation of Mn in CdTe [46]. The transition line converges at the CdTe melting point, suggesting that the CdTe structure could very easily oscillate between hexagonal and cubic phases. Intricate patterns of solid phase transformations were suggested by Albers [47]. So, the diagram of CdTe was suggested to be of eutectoid type with alloys differing slightly in composition from the 1:1 stoichiometry during cooling through the two-phase region. Furthermore, solid phase transformations in the 900–1092  C temperature range were reported from the electromotive force method and diffusion and annealing experiments [16]. Because of a stronger interaction between unlike particles arising from the considerable ionic contribution to the bond energy, the CdTe liquidus, like the other II-VI liquidus, shows a hyperbolic shape near the congruent melting point.

CdTe and CdZnTe Growth

27

It has been shown that the higher the ionicity of the bond, the smaller is the energy of formation of vacancies, which are the majority defects in II–VIs [48]. This accounts for the large II-VI nonstoichiometry, which is an additional complicating factor for their crystal growth. The higher the ionicity, the smaller is the energy of creation of dislocations and of stacking faults. This makes the highly ionic II–VI crystal lattice very sensitive to any strain, and hence highly defective. All these factors have a strong influence on the melt growth, making it difficult to obtain large crystals of high quality by this technique. As examples, the wide homogeneity range and the retrograde solidus necessitate a careful control of the stoichiometry to adjust the electronic properties of the crystals and to avoid the presence of precipitates. The highly associated melts lead to some duality between superheating and supercooling, making seeding extremely difficult. The low thermal conductivity makes it difficult to control the shape of the growth interface and to achieve Czochralski pulling. In order to overcome these difficulties, different mechanisms have been proposed: growth in a reduced temperature range to avoid the phase transitions in the solid state; growth under high superheating conditions; growth under forced convections regimes as accelerated crucible rotation technique (ACRT), micro-, or macrogravity conditions; growth under vibrational stirring, with the vibrations aimed at breaking the CdTe associates; growth under electric or magnetic field, and more.

4. SYNTHESIS In crystal growth the initial synthesis of the material is a mandatory step, which has been rarely studied [49]. It is typically the case for CdTe. Its synthesis can dramatically govern its subsequent crystal growth from the control of its stoichiometry and remains an issue as well. Compounding from the elements is achieved either in a separate ampoule (ex situ) or in the growth ampoule (in situ). Like its crystal growth, the synthesis of CdTe can be carried out as well from the melt, from solution, from the vapor, and even in the solid state. The enthalpy of formation DH and the Gibbs free energy DG for the formation of CdTe from the elements, calculated using the HSC chemistry code of Outokumpu Research as a function of temperature in the range 0–450  C, close to the Te melting point, vary respectively from
101.81 to
125.54 kJ and from
99.56 to
94.40 kJ. The Gibbs free energy of formation is negative, expressing that this exothermic reaction can occur from the left to the right. But indeed, even in the case of the reaction of very reactive elements like for instance hydrogen and oxygen, in spite of a highly negative variation of free energy, the combination of both gases

28

R. Triboulet

does not occur at RT, and the mixture persists indefinitely. The condition DG < 0 is necessary, but not sufficient. The reaction rate has also to be sufficient, otherwise the reaction will not occur even if it is thermodynamically possible. It can be necessary to have recourse to a “catalyst,” which will increase the reaction rate and activate the reaction. Energy under different forms, heat or mechanical energy for example, has to be provided to the system to start the reaction, which will be then often uncontrolled because of its highly exothermic character. The amount of energy, more generally heat, required by the system to start the reaction depends on numerous parameters, as will be seen later, and is not foreseeable. Here lie the main difficulties of synthesis: estimation of the energy to start the reaction and to control the reaction once it is started. Note that while crystal growth processes are commonly theoretically simulated, there are no such simulations in the case of synthesis because of its catalytic unpredictable character.

4.1. Liquid phase synthesis Synthesis in the liquid state from a stoichiometric mixture of the elements is probably the most commonly used technique. However, this technique faces most of the difficulties mentioned above. The enthalpy of formation DH and the Gibbs free energy DG for the reaction of formation of CdTe from the liquid elements from 450 to 1000  C calculated as above, vary respectively from
125.54 to
123.72 kJ and from
94.40 to
70.91 kJ. Here again, DG is negative, indicating that the reaction of formation of CdTe from the liquid elements should happen. The experience shows that the reaction rate is sufficient for the reaction to occur eventually from the melting point of Te without help of any catalytic heating. However, such a reaction is particularly difficult to control during the homogeneous heating of a Cd þ Te mixture in a vertical ampoule. The synthesis reaction starts locally in the 500–900  C temperature range and can spread to the whole liquid bath. The sudden increase in temperature and pressure, as a result of the high reaction heat DH, can lead to a great overheating of the material. The vapor pressure of the unreacted components can then increase to such values that the ampoule in which it is conducted may blow up. The reaction can occur locally, forming solid plugs that isolate liquid volumes in which the pressure increases with temperature. The strains can be suddenly released up to the breakage of the tube. Such a situation can also be depicted by the formation of an insoluble passive layer of CdTe developing around the reactants [50]. This solid layer isolates globules of the liquid metallic element from the remaining bath. Pressure increases in these little volumes with increase of temperature. The simultaneous rupture of these globules, at a critical

CdTe and CdZnTe Growth

29

temperature, leads to the breakage of the container. It is also to be stressed that, in addition to a high reaction heat, the CdTe formation is characterized by a volume expansion of about 13% as compared to the mixture of elements. The CdTe synthesis has been experimented in a horizontal vessel [51]. Under vacuum, the main synthesis has been found to start at a temperature some degrees lower than the Te melting point (a). The reaction heat has been estimated, from the evaluation of the decrease of the power supply of the furnace, to be close to 96 kJ/mol. After this first main reaction, other reactions of smaller intensity occur up to the melting temperature of CdTe. Under residual argon pressure (21 kPa at RT), the reaction occurs at about 800  C (b), to be compared with 450  C in vacuum. When the charge was previously stratified in the crucible (c), (Te being initially molten under hydrogen in the crucible and covered after cooling by Cd scraps), the synthesis began under vacuum at about 700  C and was less violent. In both cases (a) and (b), a significant fraction of the charge, about 5–10%, poured out the crucible and was lost. This effect was considerably minimized by using a stratified load. It turns out that the onset of the reaction of formation of CdTe from the molten components depends on numerous parameters that can be categorized as follows:  Geometrical parameters: vertical or horizontal configuration, size

and shape of the crucibles or ampoules

 Shape under which the Cd and Te elements are used: scraps, shots,

powder, or particle size

 Mixing features: how are the elements mixed in the crucible? Which

component, Cd or Te, is loaded first?

 The amount of materials  The residual atmosphere above the charge: vacuum or residual pres-

sure of a gas like Ar or H2

 Heating rate  Amount of heat  Stoichiometry of the charge

The previous descriptions of synthesis in both vertical and horizontal configurations show how hazardous the reaction can be in the liquid state. That is why the pressurization under an inert gas in an autoclave to compensate the high pressures developed during the reaction of formation has been proposed either originally without encapsulant [52] or more recently with B2O3 as an encapsulant [53]. This process makes use of a sophisticated and expensive apparatus and makes it uncertain to adjust precisely the stoichiometry by weighing large amounts of material. Furthermore, contamination has been shown to occur at the melting point of CdTe in such systems.

30

R. Triboulet

4.2. Vapor phase synthesis In order to overcome the difficulties of the melt synthesis, vapor phase synthesis has been proposed for a long time [54, 55]. This process is said to allow a more precise control of the deviation from stoichiometry. The enthalpy of formation DH and the Gibbs free energy DG of the reaction of formation of CdTe from the vapor elements calculated again as above from 500 to 1000  C vary, respectively, from
284.60 to
269.78 kJ and from
140.26 to
50.69 kJ. It has to be stressed that the reaction heat DH is more than twice of the one found when the reaction is achieved from the molten components and is thus not free of any risk of explosion. This means again that particular care has to be taken with the heating conditions. In this technique, Cd and Te are loaded in separate crucibles and the synthesis occurs by vapor transport of Cd into molten Te, containing Zn in the case of CZT [56]. Under isothermal conditions, CdTe forms in both reservoirs. A temperature of the Cd bath 20  C lower than the Te one is sufficient for the whole reaction to occur in the Te crucible.

4.3. Solid state synthesis Mechanical alloying to produce CdTe powder dedicated to the fabrication of solar cells by screen printing was developed in Matsushita Electronics [57]. Using a mixer mill, constituted by an agate grinding jar with agate balls, metallurgical (4N) Cd and Te powders were stirred in distilled water. CdTe powder of satisfactory purity grade, as assessed from X-ray diffraction patterns, was obtained after several hours. Such a ball milling process has been more recently reported [58] for the synthesis and growth of CdTe nanocrystals.

4.4. Solution synthesis Melt and vapor synthesis are difficult to control and can even become perilous. Furthermore, in both techniques, CdTe has then to be molten in order to get compact ingots. Its high melting temperature of 1092  C can be at the origin of contamination by the environment. In order to avoid high uncontrolled superheating and high contaminating temperatures, attempts have been made to achieve the synthesis in solution. Note that the heat of formation of CdTe in Te-rich solution is the same as from a stoichiometric mixture of the components. However, the excess Te has the beneficial assets to act as a heat sink, tempering too large superheating, to decrease the liquidus temperature, and to prevent the presence of any Cd excess generally at the origin of the appearance of excessive pressures. CdTe in such a solution synthesis process never reaches its melting point, which limits contamination dramatically.

CdTe and CdZnTe Growth

31

1 silica plug

Cd + Te mixture 2 T3

T1

3

T2

T4

temperature profiles

Figure 6

CdTe synthesis in a Te solution in a VB arrangement.

Solution synthesis has been achieved according to two different configurations, VB and traveling heater method (THM) ones. In the VB configuration, an ampoule filled by a charge of Cd þ Te in the 40:60 ratio, is placed in position 1 in a VB furnace, the temperature profile of which is schematized in Fig. 6. More details on the arrangement of such a furnace are given in Ref. [59]. In the first step, the temperature T1 is 480-500  C and T2 lies around 800  C (Fig. 6). The tube is slowly lowered in the steep temperature gradient at around 3-5 mm/h. The reaction of synthesis occurs in successive layers as each layer of liquid reaches the reaction temperature. The reaction does not spread to the whole volume because the steep temperature gradient ensures that the majority of the liquid is at too low a temperature to react. When the tube is in position 2, the synthesis is almost complete: T2 is then raised up to 970  C (T4) (at temperature a little bit higher than the liquidus associated with 40:60 Cd þ Te mixture according to the phase diagram of CdTe [18]) and T1 is simultaneously raised and adjusted at 550  C (T3). This temperature

32

R. Triboulet

leads to a Cd vapor pressure corresponding roughly to the thermodynamic equilibrium for the melting point of 970  C, as it can be determined from the CdTe three-phase line. The tube is subsequently slowly lowered into the furnace. If the rate is low, in the range 300 m/h and 1 mm/h, large crystals of CdTe can be obtained. For larger rates, up to 5–10 mm/h, a compact polycrystalline ingot is obtained. For excessive rates, the ingots show porosity and Te inclusions. An advantage of this process, also used for the ZnTe synthesis [60], is that the crystallization causes impurity segregation along the ingot. A very pure material, as assessed from photoluminescence spectroscopy experiments, is obtained. However, a drawback is that a fraction of the liquidus is described during the crystallization step, leading to a variation in the deviation from stoichiometry all along the ingot, according to the Cd-Te T-X phase diagram. A “drip” method has been provided by Redden et al. [61] for compounding, homogenizing, and consolidating CdTe and CZT. According to this kind of “controlled addition technique,” liquid Cd or liquid ZnCd, located in an upper “drip cup” is slowly poured at a controlled rate in a Te liquid bath placed below (Fig. 7). The whole arrangement is placed in a vertical reactor continuously flushed by hydrogen after being evacuated.

Figure 7 Schematic side cross-sectional views of the vessel apparatus showing the drip cup positioned in the ampoule, with 52 being the solute, Cd or CdZn, in the drip cup 96 and 51 the solvent, Te (from Ref. [61]).

CdTe and CdZnTe Growth

33

After the first step of fractional synthesis, the spongy material obtained is rapidly heated above the melting point of CZT for consolidating, maintained for 2 min at this temperature, and finally rapidly quenched for homogenizing. Note that the material has an uncontrolled stoichiometry, that is, it has to be heated up to its melting point with the classical subsequent contamination and that the Zn uniformity depends only on quenching. An ingot associated with a solidus composition depending on temperature can be obtained using the so-called cold traveling heater method (CTHM). According to the CTHM principle, a composite source material is constituted of a Cd rod of appropriate diameter surrounded by Te pieces and powder (Fig. 8) [62, 63]. The movement of the solvent zone through this composite charge, at a temperature of 750-950  C, induces the fractional and progressive synthesis of the compound in Te solution, silica plug

tellurium pieces cadmium rod thermal profile heater

molten zone

ΔT

Te solvent grown CdTe

Figure 8

Principle of cold traveling heater method (CTHM) (from Ref. [62]).

34

R. Triboulet

its growth, and purification. Traveling rates of a few millimeter per day can be used to synthesize ingots of several kilograms. In this method as in the VB in solution, the material never “sees” its melting point, which reduces significantly on contamination. Furthermore, significant purification results from the efficient zone-melting mechanism in Te solution. The same CTHM process has indeed been shown applicable to other tellurides like HgTe, ZnTe, and PbTe. The cs line S ¼ V, for which the total pressure over solid CdTe is minimum, PT ¼ PMIN, sometimes called for that the PMIN line, lies on the Te-rich side [1, 18]. THM growth temperatures can be correlated with cs temperatures, meaning that CTHM synthesized material can be congruently sublimated at given temperatures. Fig. 9 shows THM synthesis temperatures, or solidus compositions, versus cs temperatures calculated from the results of Greenberg. This graph will allow to choose a synthesis temperature for a sublimation planned at a given temperature, or inversely the sublimation temperature associated with a given synthesis temperature. The thermodynamic control of the stoichiometric deviation occurring from solution synthesis is obviously much more accurate than the one resulting from weighing in melt synthesis. Furthermore, CTHM can be easily extended to the synthesis and growth of homogeneous CZT ingots, as shown in Ref. [64]. This technique is now industrially exploited in 5N Plus for the synthesis of CdTe and CZT ingots of high purity and controlled stoichiometry weighing several kilograms [65, 66]. The average total impurity content, excluding carbon, nitrogen, and oxygen, has been found in the order of 100 ppb atomic, with Zn being the major contaminant and accounting for more than 40% of this value. 985

THM temperature (K)

980 975 970 965 960 955 950 945 1050

1100

1150 1200 1250 Congruent sublimation temperature (K)

1300

1350

Figure 9 THM molten zone temperature versus congruent sublimation temperature (from Ref. [49]).

CdTe and CdZnTe Growth

35

5. CONTAINER The material used for the container and the nature of an eventual coating has been shown of prime importance in the growth of CdTe, influencing both its structural and electronic properties. CdTe, with a melting point lower than the softening point of silica, can be grown in silica tubes that remain so far the preferred container for its growth. The silica tube technology involves several significant parameters: purity of silica, use of uncoated or coated tubes, with nature of the coating, graphite or boron nitride (BN). Besides the fact that the higher the thermal conductivity of the container material, the more concave becomes the solid-liquid interface, as reported further in VII, silica, although more favorable that graphite in this sense because of its lower thermal conductivity, has been shown to pose dramatic problems of purity. Contamination from the silica walls is usually observed. The purer the silica used, the less contaminated are the crystals. This justifies both the use of very pure silica tubes and graphite coating. As an example, strong contamination has been observed in CdTe crystals grown in uncoated silica tubes, leaded to semi-insulating compensated crystals without intentional doping, while purer crystals were obtained under the same conditions but in graphite coated tubes [67]. In the same study, as an experimental confirmation of the work reported in [68], related to the influence of the thermal conductivity of the container walls on the solid-liquid interface shape, single crystals were reported to be obtained in uncoated silica tubes and polycrystals in graphite-coated ones. Furthermore, contamination can also occur by diffusion through the silica walls which become porous at high temperature. Some “print” in the crystals of the furnace metallic heating elements, such as Ni, Fe, and Mn, can be found in the crystals according to the furnace technology. A reduction of the contamination from the container can arise in case of use of high-purity, high-density graphite or pyrolytic BN (pBN) crucibles. Nevertheless, when such crucibles are loaded in silica tubes, contamination from silica or through the silica walls is not completely eliminated. Influence of ampoule coatings on CdTe solidification has been demonstrated from dislocation density measurements [69]. The dislocation densities measured increased in the following order: BN-coated silica, carbon-coated silica, uncoated silica. A correspondence was seen between the mean etch pit density and the wetting property of the corresponding ampoule surface. The surface-tension and contact-angle characteristics of molten CdTe were determined using the sessile drop technique on the surfaces of silica, HF-etched silica, sandblasted silica, carbon-coated silica,

36

R. Triboulet

pBN [70], and glassy carbon [71]. The degree of wetting with CdTe increases in the order of pBN, glassy carbon, carbon-coated silica, sandblasted silica, HF-etched silica, and uncoated silica, while a small amount of Zn in CdTe (4%) was shown to decrease significantly the degree of wetting on glassy carbon. Two techniques of BN coating on fused silica ampoules have been proposed: chemical vapor deposition from borazine [72] and the use of an aqueous solution of boric acid dried in vacuum at 200  C and converted to transparent BN by heating in ammonia at 1000  C [73]. An improvement of the CZT crystal quality has been proposed to be achievable using graphite as a crucible. Etch pit density (EPD) and full width at half maximum (FWHM) values of HRXRD rocking curve of CZT crystals grown in graphite crucible were found to be much less as compared to CZT grown in carbon-coated silica ampoule [74]. The necessity of using ultra-pure, nonporous, high-density graphite is stressed by the authors.

6. CRYSTAL GROWTH Because of the difficulties mentioned above concerning the CdTe crystal growth, just about all the techniques of growth of semiconductor materials have been applied to CdTe and CZT, and even now, “novel” methods are proposed for their growth. The growth techniques of CdTe can roughly be divided into several classes: growth from “stoichiometric” (melt-growth) and off-stoichiometric melts (solution growth), growth from the vapor phase either by sublimation or by chemical vapor transport, and growth in the solid state. Given its melting point of 1092  C is lower than the softening point of silica, making it possible to use silica tubes, the melt-growth method has been very popular for the growth of bulk CdTe crystals.

6.1. Melt growth 6.1.1. Vertical Bridgman VB, generally in evacuated silica ampoules, is widely used according to different configurations. Without control of the Cd partial pressure, the growth from stoichiometric charges generally leads to Te-rich melts because of the loss of Cd by incongruent evaporation into the free ampoule volume over the charge [36, 75–82]. Some attempt to reduce the Cd loss with boric oxide as a liquid encapsulant was reported [83]. To control the deviation from stoichiometry, the charge composition can be varied and an excess Cd can be added to the charge to compensate for the Cd loss [36, 84, 85]. However, a better control of the stoichiometry can

CdTe and CdZnTe Growth

37

be achieved under a definite Cd partial pressure, using a Cd extra source, according to the so-called “modified” Bridgman or gradient freeze techniques [22, 56, 59, 86–96]. To minimize thermal stress, very small axial temperature gradients are now used, generally lower than 10 K/cm and even down to 1 K/cm, besides the larger gradients used in the past. Such small gradient led to the use of the vertical gradient freeze (VGF) technique, which does not require any mechanical movement of the charge. Using an axial gradient of 1 K/cm according to the VGF principle with control of the partial Cd vapor pressure, h111i-oriented 100 mm diameter CZT single crystals were grown [97]. According to a very similar arrangement called the “selfseeding technique,” in which an axial temperature gradient of about 3  C/cm opposite in direction to that typical of the classical VB technique allows to initiate the growth from the free-melt surface, 90 mm diameter h111i-oriented CdTe single crystals were grown [98]. Crystals of CdTe and Cd0.96Zn0.04Te were grown again by the Obreimov-Shubnikov method using the self-seeding technique and a special temperature schedule for the stepwise postgrowth cooling [99]. This postgrowth cooling process was aimed at taking into account a solid state transformation first suggested by Albers [47] and to minimize its harmful effects on the structural perfection of the crystals and the stability in time of their electrical properties. Crystals of high structural perfection with rocking curve FWHM very close to the theoretical value of 19 arcsec for the (110)reflection plane and stable electrical properties after storage for 2 years were reported to be obtained using this postgrowth cooling process. Cd1
x ZnxTe (x ¼ 0.04
0.1) single crystals with diameter up to 100 mm and height at most 40 mm were prepared by the VGF technique under full control of high Cd pressures (up to 4 bar) during the growth at growth temperature [100]. CdTe have been grown using a special multizone apparatus, called the universal multizone crystallizor, dedicated to crystal growth experiments on space shuttle and on board of the international space station [101]. The main advantage is very high stability and the elimination of all mechanical motions. From the early work of Raiskin and Butler [23], the VB technique under inert gas pressure, also called HPB has become very popular for the growth of semi-insulating crystals, mainly CZT, suitable for nuclear detection [95, 102–106]. Liquid encapsulation HPB using boron oxide as an encapsulant led to highly boron-doped semi-insulating CdTe crystals in which boron was suggested to play no role in the electrical resistivity [107]. A variant of the HPB technique called AHP-method, for axial heat flux close to the phase interface, in which a submerged heater is added to the VB configuration in the manner of the submerged heater method (SHM) (Fig. 10), is suggested to allow superior control of the growth interface shape and significant reduction in thermal stress [108, 109].

38

R. Triboulet

Hermetic case Graphite crucible Thermocouples

Melt

T2 T1

h d

T4 T3

Crystal

Figure 10 Schematic of the AHP crystallizer (from Ref. [109]).

This technique has been so far essentially submitted to computer simulations and mathematical modeling to assess its feasibility for growth of CZT (see paragraph VII). The influence of different factors on the basic Bridgman and gradient freeze processes has been studied. The influence of the nature of the residual atmosphere and of the residual gas pressure in the ampoules has been stressed. A residual hydrogen pressure in the growth tube was shown to act on the compensation state of the crystals [110], and in accordance, the resistivity of CdTe crystals was found to increase as the concentration of oxygen incorporated from the atmosphere into CdTe during preparation decreased [111]. The classical Bridgman process was extended by applying the ACRT for which many experimental studies have been conducted [112–118]. The overall effect of ACRT was shown to be an increase in size of the largest grains by a factor of between 2 and 5, a decrease in the small Te precipitate density, an absence of large precipitates, and a decrease of the X-ray line widths. A coupled vibrational stirring method was applied to the VB growth of CdTe [119]. The application of low-amplitude (0–100 Hz) mechanical vibrations to the ampoules during growth was found to move fluids rapidly without turbulence and at velocities significantly faster than ACRT. The results were an improved chemical homogeneity, but a degradation of the microcrystalline quality. An enlarged grain size was obtained when applying the heat exchanger method (HEM) to the VB growth of 3-in. diameter CdTe ingots in order to control the solid-liquid interface shape and to achieve controlled directional solidification [84]. Asymmetrical Bridgman, initially proposed for growing compositionally uniform HgCdTe

CdTe and CdZnTe Growth

39

single crystals [120], was found also to be very effective for the growth of large CdTe single crystals [121, 122]. In this method, the growth ampoule is held asymmetrically in the cylindrical furnace, allowing for flattening of the solid-liquid growth interface. Seeding in the CdTe Bridgman growth has not been very successful so far because of the duality supercooling-superheating addressed in Section 3. A smart way to overcome this obstacle has been proposed by Saucedo et al. [123]. In their “modified Bridgman technique,” the overheated melt is separated from the seed crystal. For this purpose, a turning furnace was designed and employed. Only after the furnace is rotated into its vertical growth position, the formerly superheated melt contacts the seed and the crystallization can proceed. Only a moderate positive effect of the superheating on the structural crystal quality and the crystallographic orientation of the large-angle grains is so far reported. Playing too with the effect of superheating, Su et al. [124] grew crystals of CdTe from first markedly superheated melts by the unseeded VGF method. The large crystals obtained by this method reached a resistivity of 1  107 O cm at room temperature. A very appealing technique called dewetting [125] has been applied to CdTe [126]. Its principle is to impose a gas pressure at the cold part of the crucible, approximately equal to the hydrostatic pressure of the molten material, to get a small meniscus between the solid-liquid interface and the crucible, and thus preventing crystal contact during the VB growth (Fig. 11). The structural quality of a crystal obtained from a polycrystalline seed is dramatically improved.

Crucible Gas Ph

Pc-Ph Controller Liquid

Gas Pc

Meniscus

Crystal

Crystal

Figure 11

Gas Pc

Principle of the method of “dewetting” (from Ref. [125]).

Crucible

Liquid

40

R. Triboulet

6.1.2. Horizontal Bridgman Horizontal Bridgman, historically one of the first methods used for the CdTe crystal growth, is generally achieved under Cd overpressure to adjust the stoichiometry through the control of the solid-liquid-vapor equilibrium [51, 55, 110, 127, 128]. Nevertheless, like in the VB case, the control of the melt composition through the Cd partial pressure control is sometimes neglected, and the composition of the crystals depends on the initial composition of the charge and of the free volume over the charge [129, 130]. Like in the VB case, different variants exist depending on how the temperature gradient is applied to the charge: relative movement of furnace and charge [130, 131], electronic displacement of the temperature gradient [51, 110], or “gradient freeze technique.”

6.1.3. Zone refining CdTe zone refining has been achieved either by horizontal zone refining or by a vertical sealed-ingot zone refining. Using silica crucibles as CdTe containers, the purification and growth of CdTe by horizontal zone refining was reported under controlled Cd vapor pressure [1]. After 40 passes, crystals with a residual impurity content of about 1 ppma were obtained from an initial charge containing about 0.01% impurities by weight. A maximum room temperature electron mobility of about 850 cm2/Vs was reached. From the use of graphite boats rather than quartz boats, room temperature electron mobilities of 990 cm2/Vs after 15 passes were measured [127]. The use of vitreous carbon as the most suitable crucible material for horizontal CdTe zone refining was suggested [132]. The transport of impurities through the vapor and back diffusion through the hot solid, which affect horizontal zone melting, are eliminated by the vertical sealed-ingot zone refining method [133], which allows furthermore to maintain a narrow molten zone without having to heat the solid and to realize a sharp temperature gradient between melt and solid [59, 134–136]. The effectiveness of zone-refining purification has been demonstrated from chemical analysis, photoluminescence, and electrical measurements. Electron mobilities exceeding 1.5  105 cm2/Vs at low temperature and room temperature mobilities of as high as 1100 cm2/Vs were measured [59]. It was suggested that the abnormally high mobility values measured at low temperature could be explained by the presence of microinhomogeneities and by donor-acceptor pairing [137]. Such zone-refined crystals were shown furthermore to be highly compensated and in a metastable state as a result of the high temperature gradients involved during the zone-melting process. Small crystals with low-etch pit density (10
3 cm
2) were obtained by floating-zone melting in space, the evaporation of Cd being controlled by adding excess Cd to the charge combined with heating the ampoule walls (850–950  C) [138].

CdTe and CdZnTe Growth

41

6.1.4. Czochralski pulling Attempts of CdTe Czochralski pulling under encapsulation (LEC) were not successful [139–142]. Besides severe heat transfer problems associated with the low CdTe thermal conductivity, difficulties are reported to arise from the insufficient ability of B2O3 to wet the molten CdTe, leading to a continuous loss of Cd, and to some solubility of CdTe in the encapsulant. Large ingots of 50 mm diameter were nevertheless reported to be grown by LEC, but exhibiting a high degree of polycrystallinity [142].

6.2. Solution growth In order to reduce the contamination resulting from high temperature growth, CdTe solution growth has been used to lower the growth temperature. The growth from nonstoichiometric melts has been essentially undertaken, especially from Te-rich melts. Using some kind of VGF configuration with a reticular diaphragm in the liquid dedicated to reduce convective currents and thus to stabilize growth conditions, small CdTe crystals were grown both from Cd-rich and Te-rich solutions, at 800–850 and 750–780  C, respectively [15]. Large CdTe crystals (45 mm diameter) of high purity were grown by VB by normal freezing of a slightly Te-rich solution with the growth interface kept at a fixed position by suitably lowering the furnace temperature versus crystallized length [143]. THM is essentially used for making g-detectors, an application that has become of great interest during the last decade. Different variants of the basic THM process, generally resting on the kind of source material used, are implemented. Two kinds of source material have been reported to be used for the THM growth of CdTe, either presynthesized ingots or a mixture of the constituting elements according to the CTHM principle. Given the fact that for crystal growth by THM, feed material with a density as close to 100% as possible is required, the preparation of dense feed material by casting of proper amounts of CdTe in form of an ingot by VB is generally used. This can be achieved either from a stoichiometric Cd þ Te mixture [144–146] or from a Te-rich solution, like in the ZnTe case [60, 147]. A purification of the presynthesized CdTe ingot is sometimes performed by vertical zone refining before THM growth [148–150]. The CTHM, as described in Section 4.4, is very suitable to obtain a well-controlled feed material. In order to increase the THM solution zone refining purification effect, a multipass THM (MTHM) process was developed [151]. In this technique, the starting Te solvent consists of a Te column, a certain fraction of which is taken up at each pass. This arrangement enables several passes to be carried out in the same ampoule without any handling of the material between each pass. Very pure crystals were reported to be obtained by

42

R. Triboulet

this MTHM technique, as expressed by electrical, chemical, and optical measurements. One of the main problems faced in the growth of CdTe single crystals by THM is the strong convection in the liquid zone, due to the high temperature gradients used, which adversely affects the quality of the grown crystals. The substantial contribution of convection in the THM growth of CdTe has been shown from matter transport studies in numerous papers [62, 152–154]. The application of an external stationary magnetic field is among the options to suppress the convective flow in the liquid zone. An external magnetic field aligned perfectly with the axis of the growth ampoule gives rise to a magnetic body force in the horizontal plane that balances the vertical gravitational body force, and consequently suppresses the convective flow. Not only static but also rotating magnetic fields (RMFs) have been implemented for this purpose. A threedimensional (3D) numerical simulation for the growth of CdTe crystals by THM under a static magnetic field has shown that a suitably applied field was very beneficial as it suppressed convection and led to more uniform concentration and temperature distributions and flatter growth interfaces for a prolonged and stable growth [155]. An applied magnetic induction of 3T has been shown to be beneficial to improve the growth interface and the microstructure of CZT crystals by reducing the amplitude of the growth front [156]. From a 3D computational analysis, an applied vertical static magnetic field intensity of 8 kgauss has been found to be the optimum level for which the convection was suppressed to a minimum and the growth interface was the flattest, while higher fields may lead to unstable transport structures in growth experiments [157]. In a further 3D numerical simulation [158], it was shown that in spite of the initially assumed axisymmetric boundary conditions, 3D transport structures develop as soon as the growth process begins. An optimum magnetic field level is determined for which the growth interface becomes slightly concave toward the liquid zone, giving rise to a favorable condition for growing a crystal with uniform composition and less inclusions. Application of the magnetic field suppresses the flow velocity in the solution and the magnitude of the maximum flow velocity decreases monotically with increasing magnetic field intensity. A forced convection regime was produced by a B ¼ 2 mT RMF, at 400 Hz under mg conditions, which generates in the solution zone a steady-state flow improving the radial and axial distribution of Te inclusions and mt products [159]. The influence of RMF on flow pattern and compositional uniformity in the solution zone of a THM system for growth of CdTe has been numerically investigated at the 10
6 and 10
1 g0, as representative of the space and groundprocessing conditions [160]. It is shown that under microgravity conditions, application of RMF can be used to overwhelm residual buoyancy-induced convection and to control the uniformity of solution-zone composition at

CdTe and CdZnTe Growth

43

the growth front without appreciable modification of the growth interface shape. At high-gravity levels, RMF is found not to be able to completely dominate buoyancy-induced convection. In this regime, for the range of field strengths studied, RMF is found to result in (a) complex flow structures in the solution zone, (b) enhancement of compositional non-uniformities at the growth front, and (c) increased convexity of the growth interface. ACRT has been used as well to impose a forced convection regime in the THM solvent zone. The use of ACRT, justified from a simple model confirming that convection is the dominant mechanism of matter transport in THM, has led to an enlargement of the size of CZT crystals grown by THM and to the possibility of increasing the growth rate [63]. The effects of various parameters of ACRT on the mixing are considered and a possible optimum cycle for high growth rate limits is given from detailed numerical calculations of the combined thermally driven and forced convection during THM growth of CdTe from a Te-rich solution [161]. CdTe boules of 25 cm3 containing 20 cm3 single crystals were obtained using a convection-assisted solution growth system in a VGF arrangement [162]. A sharp temperature gradient, inducing solution circulation and restricting Cd depletion at the interface, was used with a cold finger at the bottom of the ampoule to initiate nucleation. Growth proceeds with a convex solid-liquid interface due to the radial temperature gradient impressed by the cold finger. CdTe discs of 300 mm diameter dedicated to X- and g-ray detection have been grown by solvent evaporation from a Te-rich solution in an open tube maintained at a constant temperature [163]. In order to grow n-type crystals and to avoid the large deviation from stoichiometry of the Te-rich side, Cd was used as the solvent for the CdTe THM growth at 1000  C with a growth rate of 1mm/day [164], while the possibility of using such heterosolvents as In [165] or CdCl2 [166] was demonstrated. The growth from Cd-rich solutions was achieved as well by the solvent evaporation (SE) technique either by heating the solution from 970 to 1025  C at constant Cd pressure [167], or by reducing the Cd pressure isothermally at 1040  C [168]. In this last vertical SE technique, temperature instabilities were found to be due to Cd droplets falling back from the upper Cd reservoir in the solution. A modified SE technique was proposed, in which the positions of the CdTe/Cd solution and Cd reservoir were reversed [169, 170]. Small platelets and rods were obtained by flux growth at 900  C under the driving force of a temperature gradient using either Bi or Sn as the solvents [171, 172]. While millimeter-sized CdTe crystals were grown by the hydrothermal technique at about 350  C in (OH)
solutions [173], the growth of crystals from gels was unsuccessful because of the Te ion instability [174].

44

R. Triboulet

6.3. Vapor growth 6.3.1. Physical vapor transport In spite of a melting point compatible with easy melt growth in silica tubes, vapor phase growth has become very popular for the growth of CdTe crystals because of the difficulties met in the implementation of melt growth. Using a vertical configuration, CdTe crystals were grown by sublimation, or physical vapor transport (PVT), in closed tubes without seed either from a charge of composition adjusted by weighing before loading it in the growth tube [175, 176] or under controlled Cd or Te partial pressure to adjust a composition of the charge leading to optimum growth rate [177]. In classical PVT experiments, CdTe boules of limited thickness are obtained because the initial thermal conditions are lost after the growth of a 2-3 cm thick crystal due to the low CdTe thermal conductivity. In order to overcome this problem of limited thickness, a new technique called sublimation THM (STHM), in which the molten solvent zone is replaced by an empty space, was proposed [178]. By the movement of the charge relative to the heater, a temperature difference appears in STHM between sublimation interface and growth interface, and the empty space is made to migrate through the solid source material, allowing continuous growth over long distances. According to this STHM principle, CdTe crystals were grown without control of the pressure regime in the ampoule and without seed [178] or with seed using a monoellipsoid mirror furnace [179]. According to a modified STHM technique, a capillary affixed to the ampoule top and staying at room temperature allows the total pressure in the ampoule to be controlled in order to reach the conditions of minimum total pressure corresponding to cs (Fig. 12) [151]. PVT growth was also carried out in a vertical configuration in semiclosed systems, with a heat sink at the bottom of the ampoules constructed according to a scheme previously employed [180], either with a seed and the charge at the top of the ampoule, with either Ar or H2 or NH4I as the residual atmosphere [181, 182] or by in situ nucleation with the charge at the bottom of the ampoule under vacuum without heat sink [183]. In such a semiclosed system, the ampoule was left under Ar until the furnace reached operating temperature and then evacuated to 10
3-10
5 atm to initiate the growth on (111)A seeds [184]. This MarkovDavydov technique was reactivated [185] and used to obtain chlorinedoped, semi-insulating crystals [186]. The influence of the heat sink temperature on the characteristics of the crystals was shown. In such PVT growth, the source temperature ranges generally between 700 and 1040  C, with a 5–10  C difference between source and seed temperatures. CdTe crystals were grown as well according to a novel “multitube”

CdTe and CdZnTe Growth

45

Capillary

Source Material Empty space T + ΔT

Stationary heater

T Regrown CdTe

Direction of charge travel

Figure 12 Principle of sublimation traveling heater method (STHM) with capillary to control the total pressure (from Ref. [151]).

technique, close to the Markov-Davydov method, at a growth temperature of 700  C and a growth rate of 5 mm/day [187, 188]. In this multitube physical vapor transport (MTPVT) arrangement, source and growth regions are separated into separate vertical furnaces connected to a horizontal transport passage (Fig. 13). The transport passage incorporates a flow restrictor that regulates the mass flow and decouples it from the sourcesink temperature difference. This flow restrictor allows both noninvasive in situ observation of the growing crystal and real-time monitoring of vapor pressure and mass transport using a simple system based on optical absorption [189]. Growth takes place on a seed crystal after which the growth envelope is dynamically pumped to ensure that the transport is diffusionless. This system permits low-temperature growth under near-isothermal conditions and controls mass transport without the need of pulling.

46

R. Triboulet

flow restrictor

cadmium lamp and optics crossmember heating lamp

window (1 of 6) vacuum jacket

to pump

quartz “U” tube ground glass joints

upper heater zones

source crucible

growing crystal

mid heater zones

quartz block

lower heater zones gas inlet or pump to control system

to pump detectors

Figure 13 The “multitube” vapor growth system (from Ref. [188]).

The measurements of vapor pressure, mass transport, and growth temperature were used later on to control the growth process in real time after modeling the growth of CZT in a MTPVT system [190]. A model for the growth in terms of transport conditions has been further developed, which simulates the evolution of CdTe growth in the MTPVT system with time [191]. The interaction of viscous and molecular flows was investigated and the resulting stoichiometric variations calculated. CdTe crystals grown according to the MTPVT process have been shown to present etch pit densities comparable with melt-grown materials (6  104 cm
2) and relatively low strain in the crystals [192]. The level of strain and dislocation density was found reduced with growth direction, indicating an improvement in crystal quality with growth away from the CZT (4%) seed, where strain levels would be expected to be high due to the lattice mismatch. CdTe crystals of up to 50 mm in diameter have been produced by MTPVT with increased growth rates up to 12 mm/day [193]. Crystals several millimeters in thickness have been grown on commercially available 50 mm diameter (211)B GaAs seed plates using this MTPVT method with growth rates of 120 mm/h [194]. Double- and triple-axis X-ray

CdTe and CdZnTe Growth

47

diffraction gave resolution-limited FWHM values of 34 arcsec. Maps across an as-grown surface showed the FWHM to be less than 80 arcsec over the majority of the surface. PVT was employed as well in a horizontal configuration either without seed, under controlled partial pressure of Cd or Te [41, 195–200], or without control of component partial pressure without seed [201, 202] or with seed selection in a capillary by visual inspection [203, 204]. In an “improved” PVT, a three-step closing process of the ampoule was proposed to avoid any oxidation of the source and any condensation of volatile impurities evaporated from the silica tube while sealing it [205]. The influence of the source material composition on the growth rate was pointed out and studied from experimental measurements of the mass flux as a function of the source material composition [201]. The composition of the charge was adjusted by heat treatment in order to remain close to the cs condition and determined either by EDXRF and reflectivity in the region of free excitons [202] or by using an optical absorption technique [206]. Using seeded PVT, called also sublimation and PVT (SPVT), Cd1
xZnxTe crystals with x ¼ 0.2 and 40 g in weight were grown from a pretreated source material with excess Cd (or Cd þ Zn) in the ampoule (when desired) (Tsource ¼ 1000  C, DT ¼ 20–40  C, growth rate ¼ 1–4 mm/ day) [207] and CdTe crystals of 45–50 mm diameter 250–300 g in weight under Ar pressure (0.2 atm) in a semiclosed system allowing to reach the conditions of cs (Tsource ¼ 880  C, DT40  C, growth rate 0.8 mm/day) [208]. Single crystals shaped with (110) and (111) planes of approximately 10 mm size were obtained by an evaporation-condensation method on the surface of a polycrystalline source material at 820–900  C [209]. According to such a “contactless” vapor growth technique, Cd1-xZnxTe crystals with x  0.04 were obtained by this near-equilibrium process of self-selecting vapor growth (SSVG) [210]. While monocrystallinity of the product was confirmed, neither precipitation of a foreign phase nor twins were found, but—in contrast with pure CdTe—signs of undesired mosaicity emerged. Cd0.8Zn0.2Te single crystals of about 12 mm in size were produced by SSVG with an optimization of the growth system [211]. Variation in the ZnTe fraction of 0.2 was below 0.002 and rocking curve half widths around 30 arcsec were measured. No mosaicity was found.

6.3.2. Chemical vapor transport In an investigation of the vapor phase chemical transport of CdTe using I2 as chemical agent, it was shown that no transport is possible in closed tubes in the hot-cold direction, but only vapor phase transport controlled by the source sublimation [212]. Thin platelets were observed growing on the source under small temperature gradients. This growth was explained

48

R. Triboulet

in terms of a reverse (cold-hot) iodine transport associated with a reduced sublimation tendency of the facetted crystallized material with respect to the unfacetted powdered charge. CdTe single crystals were also grown in horizontal open systems either from Cd and Te2 vapors using H2 as a carrier gas [213], or from a polycrystalline charge using N2 or Ar as a carrier gas [214].

6.4. Solid state recrystallization Solid state recrystallization (SSR) appears as an attractive and promising alternative to the melt, vapor, and solution growth. In this process, strains intentionally introduced in a polycrystalline source are then released by high-temperature thermal annealing. The structure tends toward the state of lower potential energy, which is the single crystal through the motion of dislocations and grain boundaries. It works in the solid state at temperatures lower than the CdTe melting point, thus avoiding hightemperature post- and pretransition phenomena and also reducing high-temperature contamination from the environment. Furthermore, it avoids effects of the earth gravity field, which are uncontrolled in melt, solution, and vapor growth. It can be controlled to give a quasicontactless growth beside the crucible walls as in microgravity experiments. Seed-free growth of (111) CdTe and CZT crystals (Fig. 14) has been successfully achieved by SSR from polycrystalline sources obtained by fast sublimation [215].

Figure 14

CdTe single crystal grown by SSR under Cd vapor pressure (from Ref. [215]).

CdTe and CdZnTe Growth

49

7. BRIDGMAN GROWTH MODELING AND INTERFACE SHAPE DETERMINATION The difficulties met in the CdTe melt growth have prompted the idea that analytical modeling could be gainfully employed to the optimization and design of crystal growth processes. Several papers deal with numerical simulation of crystal growth in both vertical and horizontal configurations. The first analytical calculation of heat transfer in a VB-Stockbarger configuration was reported by Chang and Wilcox [216]. Identical heat conductivities in both melt and solid were assumed, and the effect of ends was neglected. Given these approximations, the isotherm shape was shown to depend on the temperature difference between the hot and cold plateau of the furnace, the container geometry, and the degree of heat transfer, taking into account the value of the Biot number Bi. The larger the Bi, the smaller the isotherm curvature, the narrower the transition region from hot to cold part of the crystal, and the steeper the temperature gradient at the growth axis are. This model was later extended [217] taking into account different heat exchange coefficients for hot and cold zones and inserting a layer of insulation between the heater and cooler. This adiabatic zone is shown not only to dramatically decrease the sensitivity of the interface shape to perturbations in the system, but also to decrease the radial temperature gradient inside the ampoule. Later on, a booster heater adjacent to this adiabatic zone was included [218]. Electrical analogs were used to model the thermal behavior of a Bridgman crystal growing system [219, 220]. This last study demonstrates that during the critical early stages of the growth, the isotherm shapes in the crystal/melt are influenced by the end losses from the ampoule. Changing the shape of the ampoule base or changing the conductivity of the ampoule stem should allow for adjusting the isotherm shape. The actual growth rate/extraction rate ratio is shown to depend on the furnace characteristics. To obtain a convex isotherm, a shallow temperature gradient should be used and end losses maximized [221]. In an analytical treatment of both axial temperature profile and radial temperature variations using a one-dimensional model, the effects of charge diameter, charge motion, thickness and thermal conductivity of a confining crucible, thermal conductivity change at the crystal-melt interface, noninfinite charge length, length of an insulating zone between hot and cold regions, thermal coupling between charge and furnace, and generation of latent heat were considered [222, 223]. A concave growth interface shape could be the result of the presence of a crucible for semiconductors. Use of quartz, with its relatively low thermal conductivity, is seen to have only a moderate effect when compared to BN. Demonstrating the power of the method of global calculation of heat transfer applied to the numerical simulation of crystal growth in a VB furnace, natural convection was shown to have little impact upon the

50

R. Triboulet

temperature field and the impact of material properties upon the global power input was clearly emphasized [224]. The same group used the steady-state finite element method (FEM) for the calculation of transient thermal flows and for the eventual calculation of the liquid-solid interface in the HB growth [225–227], and showed that 3D effects dominate the flow and that the release of latent heat of fusion has a major influence upon the shape of the interface. With the aim of getting a better understanding of fluid flow in both VB and HB, the interest of thermal gradients as low as possible was shown [228]. Favier furthermore showed that the HB configuration is the least stable from the fluid point of view, with the presence of free surfaces enhancing the instability of the flow because of the Marangoni convection, and, with respect to the VB configuration that appears more favorable from the fluid flow point of view, stressed the importance of furnace and crucible designs. Besides these general studies, numerous papers have been specifically dedicated to the Bridgman growth modeling of CdTe, which differs from many materials by its low thermal conductivity and its weak thermal conductivity difference between melt and solid, making it difficult the control of temperature distribution and interface shape. The FEM model was employed to calculate the temperature and stress fields in the CZT VB growth [229]. From their simulation, the authors recommended the following operating conditions: low cold-zone furnace temperature (1060–1065  C), low hot-zone furnace temperature (1110  C), and a moderate axial gradient (15 K/cm1). The effect of the temperature distribution in the growth ampoule on the crystallization process was studied using a finite element approach [230]. The experiments carried out with a predicted convex interface shape resulted in the growth of large crystals. From an FEM as well for the CZT VB growth, large radial gradients were shown to dominate the temperature field in the solid, while convection flattened the radial temperature distribution in the melt [231]. Concave interface shapes are predicted to arise from the thermal conductivity mismatch between solid and liquid. The shape of the solid-liquid interface is shown to be very sensitive to the growth rate due to the importance of latent heat release. The same FEM approach was used to account for Zn segregation [232]. The system is shown to be far from the diffusioncontrolled limit, and thus far from reaching a steady state. Lowering the growth rate in the system is predicted to increase axial segregation while decreasing radial segregation. From this FEM approach, the same authors [233] proposed an interrupted growth strategy in an unseeded VB system employing a series of relatively fast-growth periods followed by pauses. Providing a means to mix the Cd rejected at the growing interface into the bulk, this strategy allows the use of faster growth rates for grain selection while minimizing the risk of constitutional supercooling. Using the same FEM for heat transfer, melt convection, and interface position, Kuppurao

CdTe and CdZnTe Growth

51

and Derby [234] proposed an ampoule design with a shallow cone sitting upon a composite support made of a highly conducting core and a lessconducting outer sheath to promote axial heat transfer while inhibiting radial thermal flow. The result should be an interface shape convex toward the melt. The importance of the ampoule wall characteristics was stressed from a numerical simulation using a commercial computational code [68]. The three heat transfer modes, conduction, convection, and radiation, and a rather complex geometry of the entire system consisting of the enclosure wall, ampoule, and different graphite covers were taken into account. It was found that the graphite cover decreases the axial temperature gradient in the melt, this effect being intensified by the conical tip. Another effect of the heat conductive cover is to increase the growth rate at the onset of the growth process in the tip in such a way that, depending on the graphite thickness, this velocity can be several times larger than the ampoule translation rate; the growth rate reaches the stationary state after a time similar to the one needed in the uncovered ampoule case. Similarly, at the onset of the growth, the presence of graphite increases the concavity of the interface at the tip of the ampoule, this effect being reduced as the growth proceeds in the cylindrical zone (Fig. 15). 0.45 0.40

0 μm

0.35

1500 μm

250 μm

Concavity (ΔH/R)

0.30 0.25 0.20 0.15 0.10 0.05 0.00 0.08

0.10

0.12

0.14

0.16

0.18

0.20

Phase change isotherm position (m)

Figure 15 Concavity of the solid-liquid interface as a function of the 1365 K isotherm axial position (from Ref. [68]).

52

R. Triboulet

The importance of the tip geometry, conical or flat, and of the graphite cover thickness on the walls of the ampoule was stressed again by the same group [235]. Among several very significant conclusions, it was shown that a graphite cover makes the growth rate greater than the ampoule translation rate, increases the concavity of the liquid-solid interface mainly in the tip zone, reduces the axial thermal gradient in the melt near the interface, and increases the radial thermal gradient in the tip, producing a multicellular convective flow field in the melt. Besides a flat tip, a conical one was shown to increase the growth rate and the interface concavity produced by the graphite cover—mainly at the early stages of the growth—to retard the arrival of the axial temperature gradient in the melt to an asymptotic value as the growth progresses, to reinforce the radial thermal gradient when the interface is in the tip of uncovered ampoules, and to weaken this gradient in graphitized ampoules. The ampoule base shape has been also taken into account in VB crystal growth systems through the heat transfer between ampoule and furnace. It was found that both the flat base and the semispherical base systems produced slightly concave solid-melt interface, which is undesirable [236]. In a CZT VB growth simulation, a configuration in which the heat exchange of the ampoule with the environment occurred only by radiation was studied [237]. At the onset of the growth, heat removal through the ampoule bottom dominates, providing a convex shape of the crystallization isotherm. Portions concave toward the crystal may appear on convex isotherms as a result of the large bend height of the convex crystallization isotherms. Computer simulations were performed to assess the feasibility of a submerged heater system for growth of CZT according to the AHP method [108], which demonstrate that maintaining a constant gap width between the heater and growth interface is critical to achieving uniform axial segregation when using zone leveling. Radial segregation remains somewhat nonuniform, however, because of the poor mixing in the interface region. According to another mathematical modeling of problems of growth of CZT crystals at high-pressure inert atmosphere by the AHP method [109], it was shown that the change of the shape of the meltcrystal interface during CZT crystal growth is caused by a drastic change of the thermal resistance of the whole system. The changing interface shape is suggested to be the main reason for the inhomogeneity of the CZT crystals grown by the AHP method. Although the HB growth modeling has attracted less attention than the VB one, a study [238] deals with the understanding of CdTe and CZT growth according to the HB shelf growth process originally developed by

CdTe and CdZnTe Growth

53

4 2 0 0

5

10

15

Figure 16 Representative basic case showing temperature isotherms corresponding to a pull rate of 1 mm/h and the gradient zone positioned between x ¼ 16 cm and x ¼ 18 cm (from Ref. [238]).

Texas Instruments in the framework of the DARPA program entitled “infrared materials processing” (IRMP) [131]. A calculation using an FEM clearly shows that shelf growth naturally arises from simple heat transfer effects in the studied low-gradient system (Fig. 16). It is demonstrated as well that the shelf shape can be dramatically altered by process modifications like faster rates due to the increased rate of latent heat release, or the form and magnitude of the horizontal thermal gradient. An FEM analysis of thermal stresses presents as well the benefits of the shelf-growth configuration. The growing crystal is protected from adverse crucible sticking, and the low thermal gradients used to enhance shelf growth act to minimize thermal stresses. Besides these theoretical simulations, different strategies have been used to visualize the convection in the Bridgman technique and the liquid-solid growth interface. The influence of operating parameters on thermal convection in the Bridgman-Stockbarger technique was studied with a transparent furnace and melt [239]. Under stabilized thermal conditions, convection was found to be nearly absent in the melt near the interface, but always exists in the upper portion of the melt because of the temperature roll-off at the top of the heater. Under vertically destabilized thermal conditions caused by a short booster heater between the main heater and cooler, there is a significant convection throughout the melt. Numerous studies deal furthermore with the interface shape observation and calculation in the CdTe VB growth. 111 In was used as a radiotracing dopant for melt-solid interface shape demarcation [75]. The interface shape was found to be near optimal and varied from marginally concave to slightly convex, and attempts to modify them were successful but produced only slight changes. Surprisingly, the grain structure was not significantly improved. Actual crystal growth rates were found to be as much as 35% lower than the imposed mechanical translation rate, suggesting complex boundary conditions.

54

R. Triboulet

While decantation was sometimes used to show the interface shape [240], a quenching procedure was proposed to mark the interface shape [241]. Although the quenching conditions were not capable of freezing in the interface shape because of the low heat conductivity of CdTe, a small excess of Te was used to produce a constitutional supercooled zone in the melt making the interface shape to be sharply visible. The experimentally observed interface shapes were compared with those obtained from thermal modeling. The interface curvatures observed were shown to depend strongly on the axial position. Noticeable influence of the end effects was observed leading to the proposal of using longer ampoules in order to reduce end effects. In some studies, the equicomposition contour of Zn in doped crystals was used to reveal the solid-liquid interface shape [93, 242]. But assumptions that equated the isoconcentration contours of Zn with the shape of the solid-liquid interface were shown later to be in great error due to the large degree of radial segregation across the interface [231, 232]. Besides these indirect and sometimes questionable techniques used to reveal the interface shapes, an exploration of eddy current sensing of solid-liquid interfaces in real time during crystal growth was proposed [28, 240, 243]. Eddy current diagnostics is a nondestructive, noncontact, and remote electromagnetic technique sensitive to changes in electric conductivity. It is motivated by the difference between the electrical conductivities of most liquid and solid semiconductors. Even in the difficult case of CdTe, the electrical conductivity of which is low as compared to GaAs (which is an ideal material for such a technique), it was shown that the interface location could be determined to 1 mm and its curvature estimated with sufficient precision to be of use to characterize the VB growth process [244]. Concave interface shape in qualitative agreement with the shape of decanted ingots was found [240]. The initial melting and growth of solid Cd0.955Zn0.045Te has been successfully monitored by a two-coil eddy current sensor installed in a multizone VB furnace [245]. A finite element model was used to convert sensor impedance data to a liquid-solid interface’s position so that solid nucleation and growth velocity could both be deduced from the sensed signal. Large supercooling and unstable nucleation were revealed from in situ compounded melts. The normal growth conditions used for already precompounded charges were found to result in incomplete melt back, no supercooling, and a growth velocity less than that of the furnace translation rate. Such studies demonstrate the possibility of using in situ sensing to control seeded solidification processing. Laser ultrasonic sensing has been proposed and used as a potential noninvasive sensing methodology to monitor the solid-liquid interface during VB growth of CZT [246]. A laser ultrasonic system was used to monitor the time of flight of ultrasonic pulses that propagate across an

CdTe and CdZnTe Growth

55

ampoule during the melting and solidification of Cd0.96Zn0.04Te. The measurement principle is based upon the difference in the elastic stiffness of the solid and liquid phases of the alloy and its significant dependence upon temperature. As a result, a reduction of the longitudinal wave ultrasonic velocity was observed as the temperature was increased (from ambient temperature), followed by a 45% decrease upon melting. The data obtained indicated that the melting was slow and that solidification was accompanied by a 10–15  C undercooling. The VB growth of CZT using ACRT has been numerically analyzed by the finite difference method [247]. The forced convection resulting from ACRT and its effects on the position and shape of the melt-crystal interface were studied and the solute segregation evaluated. The authors conclude that the convection vortex in the melt appears, develops, declines, and disappears periodically following the periodic ACRT regime. As a result, growth rate, temperature, and concentration oscillate periodically. During the ACRT Bridgman growth of CZT, the interface depth increases markedly. Some suitable parameters for the ACRT Bridgman growth of CZT are proposed by the authors. Contrary to the assertions of these authors, who neglect the effects of thermal and solute buoyancy in their model, Yeckel and Derby [248] indicate from a theoretical simulation that thermal buoyancy has a dramatic effect on the flow even in a relatively small system at high rotation rate. Their computations indicate that, contrary to conventional wisdom, the ACRT rotation cycle, considered in their work for a small-scale growth system, actually suppresses mixing in the melt near the ampoule wall, resulting in diffusion-limited mass transport there. The authors stress the fact that the interaction of buoyancy and centrifugal forces is quite complicated and highly dependent on the system operating and design parameters. They state that clearly it is not possible to draw broad generalizations regarding the effect of ACRT on the VB system. The effects of the ACRT wave parameters on the solid-liquid interface concavity and the solute segregation of the crystal have been investigated later on by simulation [249]. The results show that ACRT can increase both the solid-liquid interface concavity and the temperature gradient in the melt in front of the solid-liquid interface. If the increase of the crucible maximum rotation rate can hardly improve the radial solute segregation of the crystal, the variation of the crucible acceleration time, the keep time at the maximum rotation rate, and the crucible deceleration time can dramatically affect the solute segregation. With suitable wave parameters, ACRT greatly decreases the radial solute segregation and even makes it disappear completely. However, bad wave parameters can increase notably both axial solute segregation and the radial one. As a result of this study, an excellent single crystal of which the majority part is with no segregation has been grown by suitably adjusting both ACRT wave parameters and crystal growth control parameters.

56

R. Triboulet

8. CZT PROPERTIES 8.1. CZT properties at macroscopic and microscopic scale It is essential to point out that CdTe is now being replaced for most and perhaps all the applications it gives rise to, among which are the epitaxial substrates and nuclear detectors, by CZT, which possess very specific properties besides CdTe. Zn is usually incorporated in CdTe for the CZT substrates to present a perfect lattice match to any HgCdTe composition by adjusting the Cd:Zn ratio. Furthermore, Zn is introduced in CdTe to produce ternary compounds of higher band gap, and thus higher resistivity and increased energy of defect formation, for making nuclear detectors with improved performance. The role of Zn in CdTe can be considered at both macroscopic and microscopic scale. Because of the Zn-Te bond length and ionicity smaller than the Cd-Te (but the Zn-Te binding energy is higher than Cd-Te), the CdTe lattice is strengthened by the incorporation of Zn, as shown by the increase of the shear modulus, which is a key signature of material stability [250]. This leads to the concept of increased stability of the Zn-containing alloys. As a result, the addition to Zn in CdTe was shown to reduce the density of both dislocations and subgrains [88, 251–253]. The CZT solution hardening was modeled and experimentally verified by plastic deformation and microhardness experiments [254, 255]. At the microscopic scale, the local bond length mismatch in the lattice leads both to structural distortion from ferroelectric origin [256] and to a miscibility gap in the solid for temperatures below 428  C as a result of the very strong repulsive mixing enthalpy in the solid. This last effect was theoretically predicted and experimentally verified in both thin films and bulk crystals (Fig. 17) [257, 258]. As a result, slowly cooled CZT crystals presenting phase separation show an extreme mechanical fragility. It was shown possible to avoid phase separation using sufficiently fast postgrowth cooling [258].

8.2. Segregation Besides the growth problems mentioned in Section 3 in the CdTe case, Zn segregation occurs in the CZT growth. The Zn segregation coefficient in CdTe was estimated to be in the 1.16–1.35 range [5, 259, 260] and this gives rise to inhomogeneous distribution along both axial and radial directions in the boules. Different ways have been proposed to avoid Zn segregation, which strongly affects the crystal uniformity required for any application, substrates, or nuclear detectors, and thus the production yield.

CdTe and CdZnTe Growth

57

(111)

A

200 nm

B Figure 17 (a) diffraction pattern of an as-grown Cd0.96Zn0.04Te crystal showing the (001) ZnTe//(111) matrix. Spots from ZnTe are circles. Extra spots are due to double diffraction. (b) Dark-field image of the superstructure spot indicated in (a), showing coherent ZnTe precipitates (from Ref. [258]).

The segregation coefficients of Zn and Se in CdTe being, respectively, larger than 1 (1.35) and lower than 1 (0.9, [261]), Zn and Se co-doped crystals grown by the VGF method were found to exhibit a very uniform lattice constant and high crystal perfection (Fig. 18) [259]. Under Cd/Zn partial pressure control, using a Cd þ Zn mixture with a Zn:Cd ratio determined from the thermodynamics law stating that the solute partial pressure is proportional to its concentration in the solvent, the axial Zn concentration was found to be uniform within 3%, as evaluated from X-ray diffraction and electron microprobe analysis, in Cd0.96Zn0.04Te ingots grown by the modified VGF method (Fig. 19) [93].

58

R. Triboulet

LATTICE CONSTANT (A)

6.470 b 6.465

a

6.460

6.455

c

0

.2

.4

.6

.8

1

g

Figure 18 Lattice constant of a Zn- and/or Se-doped CdTe crystal. Solid lines show the calculated results for different dopant concentrations: (a) 0.96 at % Zn and 3.6 at % Se; (b) 3.7 at % Zn; (c) 4.9 at % Se (from Ref. [259]). 4.0

Zn concentration

3.9 (b) Cd/Zn reservoir 3.8

2.6

(a) Cd reservoir

2.5 2.4

0

200

400

600

800

1000

Distance (μm)

Figure 19 Zn microsegregation in two CdZnTe crystals (from Ref. [93]).

Ingots showing a good Zn axial homogeneity were grown by a modified VB technique using a (Cd,Zn) alloy source in communication with the melt, whose temperature was gradually changed from 800 to 840  C during growth (Fig. 20) [262]. According to a close approach, an attempt was made to reduce segregation in the VB growth of Cd0.96Zn0.04Te, employing a replenish melt that was supplied from a second crucible immersed in the melt from which the crystal grew (Figs. 21 and 22) [263]. The replenish crucible had a long small-diameter passageway between the two melts to suppress diffusion. The melts were encapsulated with B2O3 under pressurized Ar to prevent evaporation. Cd0.96Zn0.04Te and Cd0.8Zn0.2Te ingots of 2 in. diameter and of excellent radial and axial uniformity (Fig. 23) were obtained by CTHM; these

Zn distribution (atom %)

7 6 5 4 No.2

3 2

No.1 0

20

A

40 60 Volume fraction (%)

80

100

40 60 Volume fraction (%)

80

100

Zn distribution (atom %)

6 5 4 3 No.3

2 1 0

B

20

Figure 20 Zn composition versus CZT volume fraction solidified: (a) constant source temperature; (b): source temperature gradually changed (from Ref. [262]). B2O3 Encapsulant Replenishing Melt,C Replenishing Crucible Long Melt Passageway (Spiral)

CRYSTAL, Co

Growth Melt, Co / k Growth Crucible

Figure 21

Schematic sketch of the replenish melt arrangement (from Ref. [263]).

60

R. Triboulet

Composition, mole % ZnTe

6 Cd1–xZnxTe 4

2 Present Technique Conventional Bridgman

0

0

4 Axial Distance, cm

8

Axial distribution profile of a CZT crystal (from Ref. [263]).

Figure 22

Zn concentration [%]

Longitudinal Profile 6 5.5 5 4.5 4 3.5 3 2.5 2 1.5 1 0

10

20

30

40 50 Distance 1(mm)

60

70

80

90

Radial Profile

Zn concentration [%]

20.5

20

19.5

19 0

5

10

15

20

25

30

35

40

45

Distance 1(mm)

Figure 23 Longitudinal and radial composition profile of a Cd0.80Zn0.20Te ingot grown by cold THM (from Ref. [63]).

50

CdTe and CdZnTe Growth

61

crystals were purer than the crystals grown by the Bridgman technique because of the purification effect induced by THM solution zone refining [64]. Both Cd and Zn rods of appropriate diameter to obtain the desired composition, surrounded by Te pieces and powder, were used as a source material. A geometry in which the feed material is a cylinder constituted of two cylindrical segments—one in CdTe and the other one in ZnTe— whose cross-sections are in a ratio corresponding to the desired alloy composition, has also proved to be successful as well for the growth of uniform CZT ingots by THM (Triboulet, unpublished work). A high degree of homogenization of Zn distribution in Cd0.96Zn0.04Te crystals was reported to be achieved on the first hand by solid state diffusion from thermal annealing occurring at three stages: (a) during the growth when the solidified crystal is near the melting temperature, (b) during a postgrowth annealing of the crystal at a high temperature, and (c) during the cooling down to room temperature and, on the other hand, by enhancing convective mixing of the melt during the initial growth stage through a proper choice of ampoule and furnace dimension [264]. In such a way, it has been possible to grow Cd0.96Zn0.04Te crystals, which have nearly 75% of their fraction within 1% Zn concentration variation (for ingots about 5 cm long) (Figs. 24 and 25).

5 CZT179

centre mid-radius edge

4 Zn concentration (%)

k = 1.17 3

Growth speed: 1mm/hr 2

35 mm

Thermal Gradient : 4-5 C°/cm Furnace dia. : 80 mm

35 mm

No annealing 1 25 mm

0

Figure 24

0.2

0.4 0.6 Fraction solidified (g)

0.8

Zn distribution in an as-grown crystal (from Ref. [264]).

1.0

62

R. Triboulet

6 CZT175

centre mid-radius edge

Zn concentration (%)

k = 1.1

4

Growth speed: 1mm/hr

35 mm

Thermal gradient : 2-3°C /cm 2

Furnace dia : 55 mm

32 mm

10 day annealing at 1000°C 18 mm 0

0

0.2

0.4 0.6 Fraction solidified (g)

0.8

1.0

Figure 25 Zn distribution in a crystal annealed for 10 days at 1000  C (from Ref. [264]).

8.3. Solid-vapor equilibrium in the CdTe-ZnTe system One of the main problems in bulk crystal growth of Cd1
xZnxTe is to obtain a prearranged composition of the alloy. It is then necessary to know functional dependencies of the crystal composition on the composition of the crystallizing matrix (melt, vapors, or both). This information is contained in the heterogeneous phase-equilibrium data. Knudsen cell mass spectrometry has been applied to investigate the solid-vapor equilibrium in the CdTe-ZnTe system at 900 K in the whole composition range [265]. At 900 K, a continuous solid solution with no phase separation is found. In the P-X section (Fig. 26) of the P-T-X phase diagram, no vapor pressure minimum was observed. Sublimation of CZT is then an incongruent process in the whole composition range. In particular, it is clear that for x < 20 mol % ZnTe, even small changes in the vapor composition lead to a significant variation in the composition of the conjugated solid phase, making it extremely difficult to control the composition of a CZT crystal in a vapor phase crystal technology without precise knowledge of the P-T-X diagram. In a further investigation, the same group determined both x and d of Cd1
xZnxTe1d by the high-precision vapor-pressure scanning method [266, 267]. From their experimental data for x ¼ 0.0, 0.05, 0.1, 0.15, 0.25, 0.5, 0.75, 0.8, 0.9, and 1, they constructed the complete Cd1
xZnxTe1d phase diagram. Their results were presented in projections of the solidus onto the T-X 3D space on the P-T plane (Fig. 27), and P(i)-T planes with

CdTe and CdZnTe Growth

Xv Xs

P*100, mm Hg

5

5

4

4

3

3

2

2

1

1

0 0

20

60 40 ZnTe, mol%

80

0 100

P-X section of the P-T-X phase diagram ZnTe-CdTe at 950 K (from Ref. [265]).

Figure 26

V) d(L

C

1 2 3 4 5 6 7 8 9 10

)

(LV

Zn

1000

P, mm Hg

63

C

S(

VL

) (LV Te

e) dT

100

e) nT

10

600

V SL 700

800

900

Zn Te (

Cd Te (

S=

V)

VL

S=

V)

Z S(

1000

1100 1200 1300

t, ⬚C

Figure 27 P-T projection of the Cd-Zn-Te system: (1) CdTe, (2) 5%, (3) 15%, (4) 50%, (5) 80%, (6) 90% ZnTe, (7) ZnTe, (8) Ref. [277], (9) Ref. [7], and (10) Ref. [4] (from Ref. [267]).

64

R. Triboulet

(i) being (Cd), (Te2), and (Zn). Formation of the Cd1
xZnxTe solid solution was shown to lead to an extension of the homogeneity range, especially at high temperatures. Increase of ZnTe in the solid solution results in a shift of the solidus toward Te, so that already for x ¼ 0.15 the solidus does not contain the stoichiometric plane. From a detailed study of the SV equilibrium, a special sublimation mode was observed in the region of the P-T-X diagram restricted by the cs curve of the binaries, S(CdTe) ¼ V and S(ZnTe) ¼ V. This region corresponds to the minimum vapor pressure, Pmin, in the system, and the characteristic feature of sublimation here is that the ratio between the sum of the metals and tellurium, (nCd þ nZn)/ nTe in the solid is equal to that in the vapor, whereas the ratio between the metals, nCd/nZn, in the two conjugated phases is different. This sublimation mode is said to determine the quasibinary behavior of the Cd1-xZnxTe1d solid solution. The essential knowledge of the described phase relationship and P-T-X phase arrangement of the solidus and vaporous surfaces for the crystal growth technology is well-illustrated in the case of the Cd0.95Zn0.05Te alloys. These studies and results, the knowledge of which plays a fundamental role in crystal growth of CZT alloys, are analyzed in detail in a review paper of Greenberg [268].

8.4. Industrial growth Only a few of the growth techniques described in Section 4 are actually used for the industrial production of CdTe and CZT crystals for substrates and mainly for nuclear detectors. Cl-doped CdTe crystals for nuclear detection are produced by THM using Te as the solvent (Eurorad, Acrorad), which is up to 4 in. diameter in Acrorad [269]. THM is used as well for the growth of 2-in.-diameter CZT single crystals according to a “proprietary” process in Redlen [270, 271]. On the other hand, the high-pressure VB (eV, Digirad) and the “modified” HB (Orbotech, formerly Imarad) techniques are, respectively, used for the production of CdTe and CZT nuclear detectors. For making substrates, HB according to the “shelf-growth process,” initially developed by TI (TI, II-VI, Inc., JME), and VB, either classical (AEG, GEC, Hughes Res. Lab., JME, Sofradir) or gradient freeze (Japan Energy), are essentially employed. Some CdTe substrates are also produced by SPVT (Eagle Picher).

9. PURITY, CONTAMINATION, AND DOPING From chemical analysis measurements by glow discharge mass spectrometry and Zeeman-corrected graphite furnace atomic absorption, the main elements detected in CZT substrates, grown in three companies (TI, II-VI,

CdTe and CdZnTe Growth

65

and JME) by their standard processes, the HB and VB techniques, were found to be Li, Na, Mg, Al, Si, Cl, Se, and Cu, and were attributed to the crystal growth processes more than to the raw materials, except for Se originating from Te [272]. Cd was also suspected to be the main source of Cu contamination for residual levels in the 1–40 ppba range [273]. In another study, such acceptors as Li, Na, K, As, Cu, Ag, Sb, and Bi, donors as Al, Cl, Ga, In, and Tl, and neutral elements as Si, Ca, Cr, Co, Sn, and Pb, were identified as classical residual impurities in as-grown p-type VB CdTe crystals [274]. Contamination by residual acceptors like Cu, Li, P, and Na, taking mainly its origin in silica, was shown from electrical and optical measurements to occur during the CdTe Bridgman growth [67]. The surface oxides on the high-purity elements, the residual gases, and the hydrocarbon vapors from the evacuation system were identified as well as sources of contamination from infrared absorption experiments [121]. Furthermore, it was found in the same study that both Zn and Se substitution enhances, and hydrogen purging reduces, the uptake of these contaminants. Although the electrical activity of impurities can be hampered by selfcompensation from native defects or other residual impurities and by gettering of impurities into Te precipitates, p-type and n-type doping of CdTe is easy to achieve; here, the elements of the first and fifth columns of the periodic chart act as acceptors and those of the third and seventh columns as donors. Mainly Al, Ga, In, Cl, and I have been used as donors and Li, Cu, Ag, N, P, and As as acceptors. Some of these elements present special behaviors in CdTe, depending on the site they occupy in the crystalline lattice, like the amphoteric behavior of Li or Ag [275]. A complete review of the impurity doping of CdTe was reported elsewhere [276]. The maximum doping levels achievable in bulk CdTe are of some 1017 cm
3 for holes (As, P, Li) and about 1018 cm
3 for electrons (Al, I, In), very close to the deviations from stoichiometry on both Cd and Te sides.

10. TYPICAL STRUCTURAL AND ELECTRONIC PROPERTIES OF CdTe AND CZT CRYSTALS Some typical structural and electronic properties of CdTe and CZT crystals grown by various techniques are displayed in Table 1.

11. CONCLUSIONS AND PERSPECTIVES The CdTe and CZT crystal growth is already a long story, but yet not ended in spite of numerous efforts for so many years. These efforts, dedicated as well to its materials science, are indeed a key for major industrial applications of these materials.

Table 1

Some typical structural and electronic properties of CdTe and CZT crystals grown by various techniques

66

Material

CdTe

Section

PVT PVT PVT HB VGF PVT HGF VMB VB VBACRT HPVB Mod. HB VB THM HB VB VB SPVT STHM VGF HPVB mod. HB THM SSVG VB VB (selfseeding) MTPVT

 36 mm  20 mm  23 mm  50 mm (sr)  30-60 mm  22 mm  4 cm (sr)  2 in.  47 mm  50 mm  100 mm 40  40 mm2  64 mm  32 mm  4 in. (sr)  37 mm  64 mm  45-50 mm  15 mm  100 mm  140 mm 45  45 mm2  15-25 mm 12 mm  40 mm  60-100 mm   50 mm

Resistivity (O cm) 10 109 108 1.0  107 7

R (mm/h)

G (K/cm)

4  10

4  10
5 1.5-25 30 2-3 1 1.5-2 1

<12 15

>1011 8.3  103 >108

5-10 8  108 30-37 1.7-4  1010 5  109 3  109 >104 5  108-5  109

FWHM (arcsec)


5

102

102

EPD (cm
2)

3-6 1 8  10
2 4-8 1 4  10
5 8  10
2 1.8

2 3-8 25 8 3-5 4.5 115 1

103-104 2  103-104 104-105 104 105 2-8  104 104 3-8  104 3.3-5.8  104

104

0.1-0.2

3-5

2.5  104

10-37 8-14 11.4 16-26 10-15 20-30 9.4

104 8  103

5  104 3-7  104

5.6 9.2 10-25

4-6  104

8-13

2-3  104

6  104

mth (cm2/V)


4


5

10

10

6  10
4

3  10
5

0.5-5  10
3 1-3  10
3 10
3

0.2-5  10
5

9.5-15 9.1

103-104 0.04 1

mte (cm2/V)

25-34 18-30  19

10
4-10
5

Reference [196] [203] [175] [129] [242] [204] [128] [89] [79] [117] [103] [110] [80] [147] [131] [278] [81] [208] [179] [97] [95] [279] [280] [211] [74] [99] [190]

R. Triboulet

CdTe CdTe CdTe CdTe CdTe CdTe CdTe CdZnTe CdZnTe CdZnTe CdZnTe CdZnTe CdTe CdTe:Cl CdZnTe CdTeSe CdZnTe CdTe CdTe:Cl CdZnTe CdZnTe CdZnTe CdTe:Cl Cd0.8Zn0.2Te Cd0.96Zn0.04Te Cd0.96Zn0.04Te

Technique

CdTe and CdZnTe Growth

67

CdTe has this unique feature of bringing together almost all the problems that can be faced in the bulk growth of semiconductors, but at a very high melting temperature. The addition of Zn does not simplify the question. In that sense, CdTe and CZT could be considered as privileged objects of study for students interested in the field of crystal growth. The difficult single-crystal growth of CdTe and CZT has given rise to not only numerous investigations of growth modeling and crystal growth, but also to the implementation of very sophisticated techniques of in situ monitoring, such as eddy current sensor techniques or laser ultrasonic sensing. While spectacular results have already been obtained in the growth of CdTe and CZT by seeded THM, several attempts have been proposed to control the convective flows occurring in the melt as a result of the influence of gravity, such as the ACRT or the use of static or RMFs. The growth by SSR is aimed as well to get rid of the uncontrolled influence of gravity. Dewetting dedicated to prevent crystal contact with the crucible walls during the VB growth appears, among the very exciting repercussions of microgravity experiments, as a very promising and appealing approach. All these efforts and research will undoubtedly open the way for the development of applications of bulk CdTe and CZT, mainly those associated with nuclear detection, such as gamma cameras for medical purposes that are boosting the interest for this attractive material. Furthermore, the very promising industrial development of CdTebased thin film solar cells open the way to a spectacular increasing production of CdTe raw material.

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CHAPTER

IC Crystal Growth of CdTe/CdZnTe in Microgravity K.W. Benz and M. Fiederle

1. INTRODUCTION: CRYSTAL GROWTH UNDER MICROGRAVITY At the beginning of the experiments of crystal growth in space, the explanations for the use of microgravity and the goals had been very enthusiastic [1]. The promises were better for crystals than earth-grown material. After many years of research on the influence of gravity on the crystal quality and many space missions, the microgravity is now used as a tool for studying the process of crystal growth and finding better growth conditions on earth [2, 3]. CdTe-based materials are excellent candidates for the crystal growth under microgravity. There is a strong need for high-quality crystals caused by the possible field of applications. However, the yield of single crystals with industrially relevant diameter is quite small. It is necessary to do extended research for a better understanding of the crystal growth of CdTe-based materials. Gravity is found to play a very important role in most of the various growth processes, like solution growth, growth from the melt and growth from the vapour phase. In general, heat and mass transport in the melt, in the solution or in the vapour phase, are all influenced by convective flows. These are normally created as a result of the effect of buoyancy. In the gravitational field, hot and less dense melt rises and the colder and denser melt sinks down. In turn, heat and mass transport influence the growth process and govern the crystallographic perfection of the grown crystal as well as the concentration and distribution of defects and doping elements.

Freiburg Material Research Center FMF, Albert-Ludwigs-Universitaet Freiburg, Stefan-Meier-Strasse 21, D-79104 Freiburg, Germany

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Crystal growth under reduced gravity in space reduces buoyancy convection drastically. Heat and mass transport may then be dominated by diffusion, and it has been theoretically demonstrated that this happens when the magnitude of the buoyancy convective flow is reduced below the microscopic growth velocity. Gravity also causes hydrostatic pressure in melts, and this in turn influences the shape of a liquid surface. Under microgravity conditions, the liquid shape is only determined by the surface tension. The shape of floating liquid zones is therefore likely to be modified in the absence of gravity. Also, the wetting behaviour of the melt on solid surfaces, such as seeds, crucibles or technical parts is modified by the absence of gravity and liquid menisci are formed, that markedly influence the crystal growth. In the following sections, an overview of the microgravity experiments dedicated to CdTe-based material is discussed. Crystal growth experiments of inorganic materials performed under microgravity in the past 25 years is summarised here [4]. The following text will describe are three different methods of the growth of CdTe-based crystals under microgravity: (a) Growth from the vapour phase (b) Growth by THM with external magnetic field (c) Bridgman growth using the Dewetting phenomenon

2. GROWTH FROM THE VAPOUR PHASE Crystal growth from the vapour phase is an alternative way to produce electronic grade semiconductor materials when growth from the melt gives no desirable results due to its inherent constraints, for example, high temperature of melt or stress due to contact with the ampoule wall. Vapour phase techniques are unique methods for processing of fragile materials, such as HgI2 and CdI2, and are also used for production of II–VI compounds (CdTe, CdHgTe, CdZnTe, ZnTe, ZnSe, etc.) [5–12]. Any imperfections in the crystal are detrimental for the performance of the device and all of them are caused by the non-optimally chosen parameters of the growth process. Due to complexity of the vapour transport and sedimentation of impurities, it is very difficult to find the growth regimes which ensure the desired properties of crystal. Experimentally it was evident that some of the difficulties in vapour growth are greatly initiated by the effect of gravity. Growth in a microgravity environment offers a unique opportunity to separate and study various components of the vapour transport and correlate them with the growth phenomena and crystal quality. In conditions when convective flow is eliminated, for example, direct evaluation of diffusion, sedimentation of impurities due to their own weight is possible. Therefore, important

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peculiarities of the growth process and interaction of the vapour with the solid surface can be revealed, studied and further used for the optimisation of the ground-based growth techniques. The first vapour grown CdTe crystal was obtained in space during the German Spacelab mission D1 (1985) by the travelling heater method (THM) with a moved vapour zone [13]. Significant local improvement of crystallinity, which is comparable with that produced by the phenomenon of dewetted growth from melt [2], has been achieved in this experiment. This early experiment with CdTe follows a series of HgI2 crystal growth experiments from the vapour phase. These experiments have been performed in the International Microgravity Laboratory 1, which pointed out a complexity of the interaction between vapour transport and surface kinetics. All HgI2 crystals grown at microgravity showed improved crystallinity and electronic transport properties compared to 1g reference crystals [14–16]. This was a clear evidence that convection has a primary role in the crystal growth from the vapour phase [17]. The obtained results were used to continue the experiments under microgravity with different semiconductors like (Hg,Cd)Te films [18], ZnSe [19] and CdTe [20]. The transport has been intensively studied in the growth process of ZnSe. In this case, firstly it was found that the vapour transport rate depends on the vapour phase stoichiometry. It has been observed that the gravity influences the distribution of impurities and native defects in the crystal. In the horizontal growth configuration, segregation of Si, Fe, Al and VZn has been observed along the gravity vector, while in the vertically configuration these point defects were distributed radial symmetrical [21]. These results are in a good agreement with the results obtained in the early 1990s, when growth of CdTe from the vapour phase has been conducted within the framework of the EURECA-1 mission. The main aim of these experiments was to study the effect of gravity on the compensation of CdTe with various dopants (V, Ti, Cl, Ga, etc.). High resistivity has been repeatedly obtained in all crystals, but a dramatic difference in resistivity distribution has been observed between the crystals grown in space and earth. This fact was attributed to the effect of gravity on the dopant distribution in the crystals [20]. The latest numerical study of ZnSe crystal growth by physical vapour transport demonstrated that the Stefan flux really dominates the system, but the gravitational induced flows, nevertheless, could be responsible for the observed inhomogeneous dopant distribution [21]. In the space experiments during the MIR and FOTON-11 missions, transport rates in the ZnSe vapour growth have been found to be considerably higher than the computed values for diffusion and Stefan flow only. It was supposed that other types of convection, which are deteriorated by thermal convection under 1g conditions, also participate in vapour transport at microgravity [22]. Nevertheless, despite amounting data on the presence of gravitational

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effects in vapour growth processes, it is a widely held belief that the mass transport rate in physical vapour transport is insensitive to gravity. These very limited experiments generated encouraging results on the understanding of the transport mechanism controlled by gravity. They are not complete and further experiments on growth of CdTe are needed to study systematically the influence of microgravity on the growth kinetic and transport processes.

3. GROWTH BY THM WITH A ROTATING MAGNETIC FIELD The growth from solution can be used to grow semiconductors as well as organic or inorganic materials. It offers the possibility to grow nearly perfect single crystals due to a lower growth temperature in comparison with the growth from the melt. The formation of defects and inhomogeneities is significantly reduced. Several growth experiments had been performed to grow single crystals under microgravity from metallic solutions for different semiconductors like germanium, gallium arsenide or ternary compounds. In the D1 mission GaSb and InP were grown by the THM [23]. The objectives of the THM experiments under microgravity were to grow bulk single crystals to study the origin of dopant or compositional inhomogeneities (striations) and defects. Several experiments with CdTe and (Hg,Cd)Te had been carried out in close collaboration between European and Russian scientists using the FOTON satellite with the Zona 4 facility. [24–28]. The reduced gravity opens the possibility to study the influence of forced mixing by external fields [29]. The use of magnetic fields, for example, a static [30] or rotating magnetic field (RMF) in combination with microgravity, is a further tool for investigating the formation of defects and in-homogeneities [24]. This gives the possibility to control the flux in the solution and improve the material homogeneity in terms of structural defects and dopants. Theoretical calculations and experiments were carried out for different material transport regimes which were caused by diffusion (under microgravity), 1 g convection or forced convection [31, 32]. Under constant magnetic field the current is produced by the motion of the fluid in the magnetic field. The induced Lorentz force acts as a dumping force. The mixing of the material, and hence homogeneity, is reduced. RMF induces a current and the resulting force generates a forced convection. The mixing is increased and a homogenous distribution is achieved. THM experiments for CdTe, Cd(Se,Te) and (Cd,Zn)Te were performed in three Russian missions Photon 7–9 in the Zone 4 facility [28]. In all three configurations, a RMF was applied to the solution zone. In the case of CdTe compounds the solution zone consisted of tellurium. All crystals were doped with chlorine to obtain high-resistivity, detector-grade material. CdTe seeds were used to improve crystal quality.

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The results of the Photon missions demonstrated the improved homogeneity by a RMF. Remarkable results were obtained by the forced mixture due to the RMF in combination with the reduced gravity. The comparison between crystal grown with and without RMF showed an improvement of the homogeneity of the transport properties. The number of tellurium inclusions was reduced in the part grown under the magnetic field [26, 28].

4. BRIDGMAN GROWTH USING DEWETTING PHENOMENON The actual research of crystal growth under microgravity is focussed on contactless growth in a Bridgman configuration. This phenomenon is called dewetting or detached growth. Results of microgravity experiments have successfully transferred to use this method on earth for CdTe [3, 33], InSb [34] and Germanium [35, 36]. The interaction between melt and container wall is very crucial for CdTe and (Cd,Zn)Te. The standard growth configuration of CdTe and (Cd,Zn)Te is graphitised quartz ampoules. The graphite layer is necessary to avoid the sticking of the crystal during the cooling phase and a breaking of the ampoule. But the graphite is a source for secondary nucleation and a possible reason for the polycrystalline structure. A modification of the ampoule configuration is mandatory. The wetting behaviour of several semiconductors and their interaction with the crucible had been studied successfully under microgravity:  III–V in rough and smooth crucibles [37] in Spacelab flights and in

EURECA (in 1992)

 GaAs in boron nitride crucible on the MIR station (in 1994)  (Cd,Zn)Te in silica by Larson (in 1994) on space shuttle [38]  CdTe:V:Zn in silica ampoules mission STS-95 on space shuttle [39]

A specific phenomenon was observed, described in the literature as detached growth or dewetting: under microgravity the hydrostatic pressure can be neglected and the semiconductor melt is not forced to get in contact with the crucible wall. It appears that all semiconductor crystals grown in space were obtained without contact with the crucible wall, leading to a dramatic improvement of crystalline quality. The theoretical background of the detached solidification was discussed by Kota et al. [40]. A overview and discussion of stability, mass transport and influence of gravity is given by Popov et al. [41–43]. The influence of the shape of the interface of the dewetting behaviour is described in detail by Epure et al. [44]. Larson et al. [38] were the first, who applied the mechanism of dewetted growth to (Cd,Zn)Te. They achieved a partially contact-free growth of (Cd,Zn)Te. They could demonstrate that the reduced stress

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(due to the wall contact) led to a reduction of Etch Pitch Density (EPD) by a factor of 100. On the STS-95 mission [45, 46] the study of the dewetting behaviour was extended to the growth of high-resistivity CdTe. Two different ampoule designs had been chosen to clarify the role of the pressure difference to obtain dewetting. The results opened the possibility of controlled dewetting growth under hydrostatic pressure. The pressure difference, Pc – Ph, between the cold and cold part of the ampoule could be actively controlled by a vapour pressure system or by additional Cd source, which will increase the pressure. The material properties and the crystallinity of the crystals are comparable to high-quality CdTe:V crystals. The reference sample grown in the laboratory did not show the same quality as the microgravity-grown samples. The stability of the dewetting process was influenced by the growth conditions. Recently, space experiments were carried out aboard the Russian unmanned satellites FOTON-M in the POLIZON facility in September 2007 (mission M3). Two experiments have been carried out in space for the dewetted growth of CZT bulk crystals from the melt. The detailed description of the POLIZON facility is given in [47]. These experiments used the Dewetting configuration in combination with a RMF for the reduction of tellurium clusters in the melt [48]. High-resistivity Cd0.9Zn0.1Te:In feed materials have been synthesised by directional solidification. Two Cd0.9Zn0.1Te:In crystals were grown. At the end of the preliminary melting phase of one crystal, a RMF was applied in order to reduce the typical tellurium clusters. The other crystal was superheated with 20 K above the melting point before pulling. Most of the time, a detached crystal has some local attached areas, named as islands, that have a thickness of tens of micrometres and a length of hundreds of micrometres. Both crystals have an average 40–60 mm thickness gap, whereas the measurements on the ground sample reveal none. According to these measurement, the presence of a gap for both space crystal shows that dewetting was successfully obtained in space. The space crystal grown with the RMF shows the better crystalline quality because of the presence of a large single grain from the bottom to the top. It appears that the space crystal grown after the application of the magnetic field shows untypical inclusions shape. They are rounded and are not visible by transmitted light in three dimensions within the slices. These secondary phases are considered as carbon secondary phases coming from the graphite felts. The series of dewetting experiments under microgravity expanded the understanding of the dewetting phenomenon. First successful crystal growth experiments under earth condition could be successful performed. A major issue is the stabilisation of the dewetting condition and reproducibility to grow single-crystal CdTe.

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5. SUMMARY AND OUTLOOK The crystal growth of CdTe-based compounds under microgravity yields remarkable and very important results for the improvement of the quality of the materials. The successful transfer of results obtained from microgravity to the dewetting growth on earth are very promising examples for the future of microgravity-related material science. The recent success of the THM-grown CdTe and (Cd,Zn)Te favours future experiments under microgravity. More studies in combination of THM and external fields are necessary to open the possibility for larger homogeneous CdTe-based crystals. The quality of CdTe-based crystals needs still improvement, which strongly depends on the better understanding of stability, heat and mass transfer. Reduced gravity is a perfect way to increase the knowledge and to reduce the gap of understanding in the growth of these semiconductors. A strong disadvantage of all experiments under microgravity is the limitation of flight opportunities. The numbers of experiments performed is very small compared to earth growth experiments. The growth of CdTe compounds takes drastically longer than the growth of silicon or III–V semiconductors. The duration of these experiments is quite long and is only possible for satellite or shuttle missions with extended energy resources. For the melt growth of CdTe, a minimum duration of 31 h is required. For a successful growth of a large CdTe crystal using THM several days are required. An excellent new tool for a series of long-term experiments is the International Space Station (ISS). New furnaces have been developed and will be installed in the European Material Science Laboratory (MSL) in the near future. The ISS will give the opportunity to grow CdTe-based crystals in high-technology furnaces with a series of experiments.

REFERENCES [1] P. Goldsmith, An introduction to the symposium on heat and mass transfer on earth and in microgravity, J. Crystal Growth 79 (1986) 37–42. [2] M. Fiederle, T. Duffar, J.P. Garandet, V. Babentsov, A. Fauler, K.W. Benz, P. Dusserre, V. Corregidor, E. Dieguez, P. Delaye, G. Roosen, V. Chevrier, J.C. Launay, Dewetted growth and characterisation of high-resistivity CdTe, J. Crystal Growth 267 (2004) 429–435. [3] N. Chevalier, P. Dusserre, J.-P. Garandet, T. Duffar, Dewetting application to CdTe single crystal growth on earth, J. Crystal Growth 261 (2004) 590–594. [4] G. Seibert, et al., in: B. Fitton, B. Battrick (Eds.), A World without Gravity ESA Publication Division, 2001. [5] M. Schieber, M. Roth, W.F. Schnepple, J. Crystal Growth 65 (1983) 353. [6] N.V. Sochinskii, V.N. Babentsov, M. Fiederle, J. Crystal Growth 262 (2004) 191 and references therein.

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[7] J.B. Mullin, C.A. Jones, B.W. Straughan, A. Royle, J. Crystal Growth 59 (1982) 135. [8] M. Laasch, T. Kunz, C. Eiche, M. Fiederle, W. Joerger, G. Kloess, K.W. Benz, J. Crystal Growth 174 (1997) 696. [9] S.J.C. Irvine, J.B. Mullin, J. Crystal Growth 55 (1981) 107. [10] C.-H. Su, M. Dudley, R. Matyi, S. Feth, S.L. Lehoczky, J. Crystal Growth 208 (2000) 237. [11] T. Feltgen, J.H. Greenberg, A.N. Guskov, M. Fiederle, K.W. Benz, Int. J. Inorg. Mater. 3 (2001) 1241. [12] H.K. Sanghera, B.J. Cantwell, A.W. Brinkman, J. Crystal Growth 1741 (Pt 3) (2002) 237–239. [13] P. Siffert, B. Biglari, M. Samimi, M. Hage-Ali, J.M. Koebel, R. Nitzsche, M. Bruder, R. Dian, R. Scho¨nholz, Characterization of CdTe crystals grown under microgravity conditions, Nucl. Instr. Methods Phys. Res. A283 (1989) 363–369. [14] L. Van Den Berg, W.F. Schnepple, Nucl. Instr. Methods Phys. Res. Section A283 (1989) 335. [15] S. Leon-Gits, E. Fries, Nucl. Instr. Methods Phys. Res. Section A285 (285) (1989) 513. [16] B. Coupat, J.P. Badaud, J.P. Fournier, S. Er-Raji, R. Cadoret, A. Magnan, J. Crystal Growth 141 (1994) 465. [17] M. Piechotka, E. Kaldis, G. Wetzel, A. Flisch, J. Crystal Growth 193 (1998) 90. [18] H. Wiedemeier, Y.-R. Ge, M.A. Hutchins, Y.-G. Sha, J. Crystal Growth 146 (1995) 610. [19] C.-H. Su, Proc. SPIE Int. Soc. Opt. Eng. 3123 (1997) 7. [20] K.-W. Benz, M. Laasch, T. Kunz, M. Fiederle, W. Joerger, Proc. SPIE Int. Soc. Opt. Eng. 3123 (1997) 2. [21] C.-H. Su, S. Feth, D. Hirschfeld, T.M. Smith, Ling Jun Wang, M.P. Volz, S.L. Lehoczky, J. Crystal Growth 204 (1999) 41. [22] N. Ramachandran, Ching-Hua Su, S.L. Lehoczky, J. Crystal Growth 208 (2000) 269. [23] K.W. Benz, G. Mu¨Iler, GaSb and InSb crystals grown by vertical and horizontal travelling heater method, J. Crystal Growth 46 (1979) 35. [24] A.S. Senchenkov, I.V. Barmin, A.S. Tomson, V.V. Krapukhin, THM. Seedless growth of CdxHg1-xTe (x þ 0.2) single crystals within rotating magnetic field, J. Crystal Growth 197 (1999) 552–556. [25] C. Eiche, W. Joerger, M. Fiederle, D. Ebling, M. Salk, R. Schwarz, K.W. Benz, Characterization of CdTe:Cl crystals grown under microgravity conditions by time dependent charge measurements (TDCM), J. Crystal Growth 166 (1996) 245–250. [26] M. Fiederle, C. Eiche, W. Joerger, M. Salk, A.S. Senchenkov, A.V. Egorov, D.G. Ebling, K.W. Benz, Radiation detector properties of CdTe0.9Se0.:Cl crystals grown under microgravity in a rotating magnetic field, J. Crystal Growth 166 (1996) 256–260. [27] P. Rudolph, A. Engel, I. Schentke, A. Grochocki, Distribution and genesis of inclusions in CdTe and (Cd,Zn)Te single crystals grown by the Bridgman method and by the travelling heater method, J. Crystal Growth 147 (1995) 298–304. [28] M. Salk, M. Fiederle, K.W. Benz, A.S. Senchenkov, A.V. Egorov, D.G. Matioukhin, CdTe and CdTe09Se0.1 crystals grown by the travelling heater method using a rotating magnetic field, J. Crystal Growth 138 (1994) 161–167. [29] Y.L. Gelfgat, Electromagnetic Field Application in the Processes of Single Crystal Growth under Microgravity, Acta Astronaut. 37 (1995) 333–345. [30] Y. Wang, K. Kudo, Y. Inatomi, R. Ji, T. Motegi, Growth interface of CdZnTe grown from Te solution with THM technique under static magnetic field, J. Crystal Growth 284 (2005) 406–411. [31] Y. Liu, S. Dost, B. Lent, R.F. Redden, A three-dimensional numerical simulation model for the growth of CdTe single crystals by the travelling heater method under magnetic field, J. Crystal Growth 254 (2003) 285–297. [32] I.V. Barmin, A.S. Senchenkov, I.Ch. Avetisov, E.V. Zharikov, Low-energy methods of mass transfer control at crystal growth, J. Crystal Growth 275 (2005) e1487–e1493.

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[33] H. Zhang, D.J. Larson Jr., C.L. Wang, T.H. Chen, Kinetics and heat transfer of CdZnTe Bridgman growth without wall contact, J. Crystal Growth 250 (2003) 215–222. [34] W. Jianbin, L.L. Regel, W.R. Wilcox, Detached solidification of InSb on earth, J. Crystal Growth 260 (3–4) (2004) 590–599. [35] M. Schweizer, M.P. Volz, S.D. Cobb, L. Vujisic, S. Motakef, J. Szoke, F.R. Szofran, Stability of detached-grown germanium single crystals, J. Crystal Growth 237–239 (2002) 2107–2111. [36] W. Palosz, M.P. Volz, S. Cobb, S. Motakef, F.R. Szofran, Detached growth of germanium by directional solidification, J. Crystal Growth 277 (2005) 124–132. [37] T. Duffar, P. Boiton, P. Dusserre, J. Abadie, Crucible de-wetting during Bridgman growth in microgravity II. Smooth crucibles, J. Crystal Growth 179 (1997) 397–409. [38] D.J. Larson, M. Dudley Jr., H. Chung, B. Raghethamachar, Characterization of Zn-alloyed CdTe compound semiconductors processed in microgravity on USML-1 and USML-2, Adv. Space Res. 22 (8) (1998) 1179–1188. [39] T. Duffar, P. Dusserre, M. Giacometti, K.W. Benz, M. Fiederle, E. Dieguez, G. Roosen, J.C. Launay, Dewetting and structural quality of CdTe:Zn:V grown in space, Acta Astronaut. 48 (2–3) (2001) 157–161. [40] A.K. Kota, G. Anand, S. Ramakrishnan, L.L. Regel, W.R. Wilcox, Influence of oxygen, hydrogen, helium, argon and vacuum on the surface behavior of molten InSb, other semiconductors, and metals on silica, J. Crystal Growth 290 (2006) 319–333. [41] D.I. Popov, L.L. Regel, W.R. Wilcox, Detached solidification. 1. Steady-state results at zero gravity, J. Mater. Synth. Process. 5 (4) (1997) 283–297. Plenum, USA. [42] D.I. Popov, L.L. Regel, W.R. Wilcox, Detached solidification. 2. Stability, J. Mater. Synth. Proces. 5 (4) (1997) 299–312. [43] D.I. Popov, L.L. Regel, W.R. Wilcox, Detached solidification. 3. Influence of acceleration and heat transfer, J. Mater. Synth. Process. 5 (4) (1997) 313–335. [44] S. Epure, T. Duffar, L. Braescu, Comparison between analytical and numeric determination of the interface curvature during dewetted Bridgman crystal growth, J. Crystal Growth 310 (2008) 1559–1563. [45] T. Duffar, P. Dusserre, M. Giacometti, K.W. Benz, M. Fiederle, E. Dieguez, G. Roosen, J.C. Launay, Dewetting and structural quality of CdTe:Zn:V grown in space, Acta Astronaut. 48 (2–3) (2001) 157–161. [46] M. Fiederle, T. Duffar, V. Babentsov, K.W. Benz, P. Dusserre, V. Corregidor, E. Dieguez, P. Delaye, G. Roosen, V. Chevrier, J.C. Launay, Dewetted growth of CdTe in microgravity (STS-95), Cryst. Res. Technol 39 (6) (2004) 481–490. [47] A. Senchenkov, A.V. Egorov, I.V. Barmin, P. Sickinger, Automatic POLIZON facility for space experiments on the Russian FOTON satellite, in: Proceedings of the 1st International Symposium on Microgravity Research and Applications in Physical Sciences and Biotechnology, Sorreto, 2000. [48] L. Sylla, A. Fauler, M. Fiederle, T. Duffar, E. Dieguez, L. Zanotti, A. Zappettini, G. Roosen, Dewetting during the crystal growth of (Cd,Zn)Te:In under microgravity, IEEE Trans. Nucl. Science Vol. 56, 4 (Part 1), 2009-10-20.

CHAPTER

ID Heteroepitaxial Growth of CdTe Thin Films Daniel Lincot

1. INTRODUCTION Epitaxial growth is an essential aspect for achieving high-performance optoelectronic devices by providing the possibility of formation of a single crystalline material with low structural defect concentration. Most devices are based on the formation of thin-film heterostructures where epitaxial relationships are provided from one layer to the next one, even if the materials are chemically and structurally different. This is remarkably exemplified in the field of III-V heterostructures, where record efficiency solar cells are based on three junctions with different band gaps covering most of the solar spectrum, involving up to 32 epitaxial layers from a few tens of nanometers to the micron thickness range grown on single crystalline germanium substrates. In the field of CdTe and II-VI devices, such an approach is used for third-generation infrared imaging systems, like focal plane arrays (FPAs), in combination with Cd1-xHgxTe layers grown on various substrates (GaAs, Si, sapphire, etc.). The CdTe layer serves as an intermediate layer in between them. Superlattices where composition modulation is achieved at the nanometer scale are also superb examples of epitaxial heterostructures (Fig. 1; [1]). Even if the number of layers is more limited than for III-V compounds, preparing such devices has faced immense challenges in terms of epitaxial growth and quality optimization. The initial condition for epitaxial growth is the adaptation of both structures at the interface between the substrate and growing layer. It is in first order analyzed in direct relationship with lattice parameters and

Institute of Research and Development of Photovoltaic Energy (IRDEP), UMR 7174 CNRS-EDF-Chimie ParisTech, 6 Quai Watier, 78401 Chatou Cedex, France

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HgTe CdTe

5 nm 50 nm

CdTe Buffer

Figure 1 Example of an epitaxial CdTe/HgTe superlattice grown from Si (211)B by MBE. It consists of 200 periods (total thickness of 1.8 mm), with an intermediate CdTe epitaxial buffer layer of 7 mm (adapted from Ref. [1]).

crystallographic structures of both materials. The key parameter is the structural adaptation between two crystallographic planes facing each other, which corresponds to the minimization of the free energy of the interface. Figure 2 shows the relationships between CdTe and other II-VI materials, and conceivable substrate materials of face-centered cubic structure belonging to the IV or III-V groups [2]. Note that excellent matching exists with HgTe, explaining the interest of cadmium mercury

4.0 3.5

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Ge GaAs

ENERGY GAR IN eV

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InSb

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HgS

HgSe HgTe

–0.5

5.4 5.6 5.8 6.0 6.2 6.4 6.6 LATTICE PARAMETER IN ANGSTROM UNITS

Figure 2 Lattice parameters and band gap of II-VI compounds with respect to lattice parameters of main III-Vs, Si and Ge substrates for heteroepitaxial growth (adapted from Ref. [2]).

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telluride (CMT) alloys for infrared detectors. For the growth on III-Vs and Si, it immediately appears that lattice-matching ranges are almost perfect with InSb with larger mismatches toward GaAs (14.4%) and Si (19%). In spite of these mismatches, integration of CdTe or II-VI devices on GaAs and especially on Si substrates is of strategic importance due to their better availability and the necessary coupling with silicon-based technologies. This trend has dominated the evolution of CdTe epitaxial growth studies for 30 years. Beside structural aspects, chemical aspects play a great role in the interface formation and control of the epitaxy; chemical bonding may or may not promote epitaxy, depending on the affinity of the chemical elements of the substrate with Cd and Te. Chemical derivatization with other elements like Zn, Te, As, Sb, or H can modify subsequent growth. In the growth process, these local structural and chemical interactions have to be combined with initial steps of nucleation and growth, leading to layer-by-layer 2D growth, 3D growth, or sequential 2D-3D growth known as the Stranski-Krastanow growth mode; this can create in-plane mosaicity effects. After coalescence, the resulting layer will tend to be faulty due to some structural mismatch at the interface leading to strains and the presence of extended defects such as dislocation loops or stacking faults. These defects will tend to decrease as a function of thickness, leading to a relaxed layer with increasing quality. This is well-illustrated in Fig. 3 depicting stacking faults in the case of the growth of (111) CdTe by molecular beam epitaxy (MBE) on misoriented (100) Si [3]. Another example is shown in Fig. 4 in the case of the epitaxial growth

100 nm

Figure 3 Transmission electron micrograph of (111) CdTe in [110] orientation grown on misoriented (100) Si showing falling off in density of lamellar twinning defects away from the CdTe/Si interface (not shown) (adapted from Ref. [3]).

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250 EPD of CdTe/Si EPD of CdTe/GaAs

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106 150 105 100

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104 50

103

FWHM of CdTe/(211)misorientated Si FWHM of CdTe/(211) standard Si FWHM of CdTe/(211) standard GaAs

102

0 0

5

10 15 Thickness (mm)

20

Figure 4 Dependence of the FWHM value of (422) XDRC and EPD of CdTe epitaxial layers grown by MBE on GaAs and Si substrates as a function of the epitaxial layer thickness (adapted from Ref. [4]).

by MBE of CdTe films on Si and GaAs substrates under different conditions [4]. The structural quality is indicated by the full width at half maximum (FWHM) of a characteristic X-Ray Double crystal Rocking Curve (XDRC) peak on the left vertical axis and by the etch pit density (EPD) on the right axis. It appears that the thickness of the defected layer is about 5 mm under optimized conditions for both GaAs and Si, and more than 10 mm under nonoptimized conditions. The FWHM value of 50 arcsec is close to what is expected for a single crystal. EDP values of about 106 per cm2 are at the lower limit of what is needed for high-efficiency devices. However, EPD values closer to 105 per cm2 have been obtained on Si (211) under optimized procedure [5]. Another consequence of the lattice mismatch between the substrate and film is the presence of stress and strain gradients as shown in Fig. 5 in the case of the growth of CdTe by MBE on 4% mismatched Cd0.096Zn0.04Te substrate [6]. Detailed characterization of the subsequent defects is the aim of many papers, using for instance highresolution transmission microscopy for direct structural characterization or photoluminescence in order to study their direct impact on the optoelectronic properties. The fundamental analysis of their mechanism of formation is an issue for acting on specific parameters, like the growth ones, to reduce their concentration and extent in thickness. The last zone is the growing interface where additional matter is added from the adjacent phase. It is obviously the most challenging step and also

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100 80 60 40 20 0

0

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2

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6

7

8

Figure 5 Dependence of in plane Stress (in MPa) as a function of thickness of CdTe layers grown by MBE on 4% Cd0.96Zn0.04Te substrates (from Ref. [6]), by XRD and optical measurements. Dotted line corresponds to the 1/l expected variation.

the most fascinating scientific area. At this interface, growth precursors have to be supplied near the surface by evaporation or hydrodynamic transport in a carrier gas. They have then to adsorb to the surface and diffuse along it to reach a growth center where atoms are included in the lattice, involving possible chemical reactions, desorption reactions, dynamic surface reconstruction, etc. The growth process is thus a complex balance between hydrodynamics, thermodynamics, and kinetics, depending on the composition of the vapor phase, impinging rate, temperature, external excitation like light, HF, and initial interface conditions. The interface will thus develop in a self-assembling mode, giving rise to preferential orientations, shapes, etc. The possibility of characterizing, controlling, and monitoring these phenomena is clearly a major challenge. Main deposition methods for CdTe epitaxial films are MBE and organometallic vapor phase epitaxy (OMVPE), the first one being based on directional physical evaporation of the components in the elemental or molecular form, the second one on the chemical reaction of precursor molecules on the substrate. Other methods can be pointed out, such as hot wall evaporation (HWE), evaporation, atomic layer epitaxy (ALE), and close space sublimation (CSS). In the liquid phase, epitaxial electrodeposition of CdTe has also been demonstrated.

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In situ methods of characterization have been used, like XPS, RHEED, XRD, ellipsometry, reflectometry, and mass spectrometry. They can be combined with ex situ local probes like STM. The objective is to improve the control of the growth process, either from semiempirical correlations or by theoretical modeling of the deposition mechanism. This chapter contains two main parts: the first one devoted to deposition methods and related growth mechanisms, the second one to the effect of substrates on the epitaxial growth of CdTe.

2. OVERVIEW OF DEPOSITION METHODS 2.1. Molecular beam epitaxy MBE involves the evaporation from effusion sources of molecular beams of different combinations of CdTe, Te, and Cd under high-vacuum conditions, resulting in the formation of CdTe films in a highly controlled manner. CdTe has the asset of congruent sublimation [7], meaning that basically a CdTe source alone is sufficient [8, 9]. However additional Cd or Te sources offer the possibility of fine tuning the deviations from stoichiometry. Doping can be achieved adding In (n-type) or As (p-type). As a consequence, MBE is a highly directional process controlled by the theory of gases under high-vacuum conditions. The growth is controlled by the fluxes of the elements, which are further controlled separately using shutters and also by the substrate temperature. It is thus a highly versatile technique. Baseline pressure lies from about 5.10
11 to 10
9 torr, while working pressure can be raised up to 10
10-10
8 torr. Growth rates can be varied from a few tens of nanometer per hour to a few micrometer per hour. Deposition temperatures can be varied from room temperature to 500  C. The high-vacuum environment allows the implementation of in situ techniques, especially electron diffraction, like RHEED, LEED, and photon spectroscopy like XPS, AES, ellipsometry, and SIMS. Heteroepitaxial deposition of CdTe by MBE has been intensively investigated since the 1980s by various groups. On can quote as a first example the homoepitaxial growth of CdTe on CdTe (100) and (111) by MBE, reported for the first time in 1981. Epitaxial growth was succeeded from room temperature on CdTe (111) and from 80  C on CdTe (100) [8]. The heteroepitaxial growth on InSb (100), which is a perfectly matched substrate, has been then explored at 25-220  C, at a growth rate of 0.5 mm/ h (fixed by a single CdTe source at 650  C) [9]. The transition between polycrystalline and epitaxial growth takes place at 50  C. Optimal-quality twin-free (100) films were obtained at 220  C, attested by FWHM of 56 arcsec on the (400) peak for a 0.8-mm-thick film. Such a value corresponds to a perfect crystal at this thickness. Since 1982, this prompted the growth

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on sapphire, which is a key substrate for further applicative developments on which the epitaxial growth was demonstrated [10]; it was then carried out on GaAs [11–14], Ge [15], and Si [17] with lattice mismatches of 14%, 13.5%, and 19%, respectively. The growth was also performed from separated Cd and Te effusion sources at 290  C and 405  C, respectively, giving a stoichiometric beam with an atomic flux on the substrate of 2.9  1015 Te and Cd atoms/cm2s [15, 16]. Later on, photoassisted MBE (PMBE) was introduced using a broad-band argon laser (514–528 nm), with an impinging flux of 150 mW/cm2. PMBE was shown to improve markedly the n-type doping, using In as a dopant, and the photoluminescence properties [18]. Related surface processes are associated with lightinduced surface reconstruction effects and increased desorption rate of Te [19]. Instead of effusion sources, laser-induced evaporation, also called laser-assisted deposition (LADA), has also been developed for the growth of CdTe by MBE [20–22]. Contrary to PMBE, the CdTe source evaporation is light assisted in LADA. A Nd:YAG laser has been used with 100 ms pulses of 106-108 W/cm2. The variation of the beam energy allows changing the evaporated Te species from Te2 (low-energy pulses) to Te (high-energy pulses). This leads to an increase of the growth rate in the high-temperature deposition range (300  C–400  C), showing that atomic tellurium is more reactive than molecular tellurium, as expected. Mechanistic studies of the growth of CdTe by MBE have been the focus of several studies, mainly on the basis of in situ RHEED experiments [22–25]. Under stoichiometric Cd and Te fluxes, the growth rate shows a plateau at low temperatures, where sticking coefficients are close to unity and tend to decrease between 250  C and 400  C due to the desorption of Te at the onset of sublimation, as shown in Fig. 6 [25]. Surface reconstruction between (2  1) Te-rich surface to c(2  2)
(2  1) Cd-rich surface on (100) CdTe, as visualized in Fig. 7 [26], is observed as the temperature increases, as shown in Fig. 8 [23]. Theoretical models, both on analytical and computational grounds, allow to explain the experimental data using elementary surface steps and give the grounds for MBE growth modeling of CdTe, as shown in Fig. 6 [24, 25]. Using STM, the growth at atomic scale on (100) surfaces has been investigated. 2D islands are observed on c(2  2) Cd-rich surface with the incorporation of Te under the Te2 form. Vicinal surfaces with steps are indicated developing smooth layers at the atomic level [27]. Theoretical modeling of surface reconstruction of flat (100) CdTe surfaces has been reported later on [26]. One should also point out that the growth processes at the atomic level are also valid for co-evaporation. This more general term does not restrict to the epitaxial growth on single crystal substrates. It is also used for epitaxial CdTe films either in high vacuum (10
7 torr) [28] or ultrahigh vacuum (<2.10
11 torr) [29], and is widely used for the growth of polycrystalline thin films of CdTe on glass or plastic substrates, especially in the area of thin-film solar cells [30].

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growth rate (ML/s)

0,7

520

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560

580

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Figure 6 Growth rate of CdTe by MBE as a function of temperature. Closed squares are experimental results, stars are calculated values from Monte Carlo simulation, and the continuous line is from an analytical model. Inset: grounds of the analytical model: (A) desorption of a Te adatom, (B) diffusion of a CdTe pair, (C) to CdTe pairs meet and form a Te adatom dimer (adapted from Ref. [25]).

A

B

C

Figure 7 Sketches of the reconstruction of CdTe (100) surfaces which takes place under MBE and ALE growth. Cd atoms are shown in filled circles and Te atoms as open circles. The gray rectangles show the surface unit cell. The (110) axis is aligned horizontally. “a” and “b” correspond to Cd-terminated reconstruction, panel a shows the (2  2) Cd, panel b the (2  1) Cd. C corresponds to a Te-terminated reconstructed (2  1) Te surface with a full monolayer of Te (adapted from Ref. [26]).

Heteroepitaxial Growth of CdTe Thin Films

Transition temperature (°C)

10 Intensity (a.u.)

360⬚C 355⬚C 335⬚C 325⬚C 0

0

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c(2×2)−(2×1) (2×1)

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2 Cd/Te ratio

Time (s)

Figure 8 A RHEED intensity oscillation for CdTe growth on a type B vicinal surface (tilted 2 around [110]). The CdTe/Te ratio is 4 and the low temperature growth rate is 0.6 ML/s. b: Transition between Te-rich reconstructed (100) surface and Cd-rich reconstructed surface as a function of Cd/Te ration in the MBE and the deposition temperature (adapted from Ref. [23]).

2.2. Hot wall epitaxy Hot wall epitaxy (HWE) is also based on the evaporation of a CdTe source in a vertical closed tube under high vacuum. The bottom plate (source), the walls between the bottom plate and top plate, and the top plate where the substrate is situated are brought at different temperatures T1-T2-T3, with T1 > T2 > T3. The wall temperature prevents the formation of deposit on the walls. This method was introduced in 1978 for the epitaxial growth of CdTe [31]. An additional effusion source placed into the chamber was used for stoichiometric control, doping [32], or alloy formation [33]. Example of deposition conditions are T1 ¼ 410  C, T2 ¼ 400  C, T3 ¼ 320  C, resulting in growth rates from 0.3 to 1 mm/h [34], with a background pressure of 2.10
7 torr. An extension of HWE called multitube physical vapor transport (MTPVT) is reported to allow very fast growth rates up to 150 mm/h for the formation of bulk CdTe single crystals up to several mm thick by epitaxy even on foreign substrates such as GaAs [35, 36].

2.3. Close space sublimation CSS resembles HWE, but the source and substrates are brought much closer at a typical distance of 1-5 mm, either under vacuum or in presence of gaseous species like nitrogen, argon, or hydrogen [37]. The source and substrate temperatures are fixed at a wide range of temperature from 200  C to 700  C. Very high growth rates are attainable using CSS, for

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deposition rate (μm/hr)

100

sublimation limited diffusion limited

Simulation conditions: 10

Source temperature = 600°C Substrate temperature = 590 and 500°C Pressures = Variable Source substrate distance = 0.87mm Gases used: He and Ar

Experimental conditions: Source temperature = 600°C Substrate temperature = 500 °C Pressures = Variable Source substrate distance = 0.87mm Gas used: Ar

1 0.01

0.1

1 10 Pressure in chamber (torr)

HeΔT = 10°C

ArΔT = 10°C

Experimental

HeΔT = 100°C

100

1000

ArΔT = 100°C

Figure 9 Growth rate of CdTe films by CSS as a function of experimental parameters. Comparison with simulation (adapted from Ref. [45]).

instance 5 mm/min for a source at 650  C and a substrate at about 600  C [38]. Figure 9 shows the variation of the growth rate of CdTe in CSS experiments as a function of the deposition parameters. Two main regimes, a kinetic one (from source sublimation) and a diffusion one, are observed. This explains the use of CSS for industrial production of CdTe solar cells due to the high throughput of the method [39, 40]; this has been recently employed for the elaboration of thick epitaxial layers on GaAs, which is up to 120 mm, with FWHM of 300 arcsec for growth rates from 5 to 20 mm/h [41]. Epitaxial growth was not a primary objective of the CSS method, even if homoepitaxy was noted in earlier studies [38]. In presence of hydrogen, the growth proceeds via two competing reactions, the first one being the sublimation reaction and the other one, at lower temperature, a chemical transport with the formation of hydrogen telluride [42], explaining an early denomination of the method as close-space vapor transport (CSVT). Detailed published parametric studies are rather rare [43–45]. Nevertheless a simple sublimation/diffusion model has been proposed (Fig. 8) [45].

Heteroepitaxial Growth of CdTe Thin Films

95

2.4. Atomic layer epitaxy ALE is based on the sequential supply of precursor species reacting on the substrate. The growth is then controlled only by surface reactions and structure. It is thus an ideal tool for epitaxial growth. It can be performed by either using elemental Cd and Te beams in ultrahigh vacuum (UHV) as in MBE [46–48], or organometallic species like in OMCVD [49]. Typical conditions for UHV-ALE are 200-400  C for the substrate temperature, with typical pulse times ranging from 1 to 8 s and purge times from 0.1 to 2 s. High-quality epitaxial films were obtained with growth rates corresponding to one atomic monolayer (0.324 nm for (100) CdTe), down to 0.5 coverage in well-defined temperature windows. The growth is monitored very precisely using in situ RHEED measurements [47, 48, 50–52], as shown in Fig. 10. The efficient combination with STM allowed further relating precisely RHEED features to pulse-dependent surface reconstructions [48, 50–52]. Similar oscillations of surface structure are observed in the case of classical MBE with simultaneous supply of reactants, as shown in Fig. 8. The advantage of ALE is to force the time synchronization of surface structure variations. Similar changes of surface structures are anticipated, whatever the deposition method used

Step 1

Step 2

[110] 50 nm [110]

Te

Cd

Te

Step 3

Step 4

Te RHEED Cd

Cd Te

Figure 10 Typical RHEED specular spot intensity variation during two ALE cycles of CdTe (100) growth. Each step is imaged by STM and illustrated at the atomic scale in cross section (adapted from Ref. [52]).

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(MOVPE, electrodeposition, etc.). In the case of Fig. 10, the periodicity of the surface reconstruction process needs two Te-Cd pulse sequences with an intermediate half monolayer coverage, in agreement with the growth rate determination. This intermediate coverage has the maximum roughness, as explained by surface atomic structure. In case of using surfaces 2 misoriented around [110], the intermediate Cd-Te rough states disappear due to the transition between an island growth mode toward a step flow one from the edges of the created terraces [50]. These effects can be used for the formation of QD nanostructures [51, 52]. Surface reconstruction processes, such as those occurring in ALE, have been further modeled [53, 54].

2.5. Metal organic vapor phase epitaxy MOVPE is, as MBE, a major deposition method for epitaxial CdTe films. Epitaxial growth was first demonstrated at 500  C on sapphire using dimethylcadmium (DMCd) and dimethyltelluride (DMT) [55]—DMT being preferred to H2Te since it is more stable— and then diethyltelluride (DET) on InSb, InP, GaAs, and sapphire [56–59]. MOVPE is highly versatile since the precursor species are stable and can be used and transported in vacuum up to atmospheric pressure, and the deposition reaction only takes place on the substrate by temperature activation. Growth temperatures range from about 200 to 600  C. The advantage of MOVPE resulting from the use of stable precursors at low temperature has as a counterpart an increase of the deposition temperature as compared to MBE. Depending on the concentrations of precursors, growth rate can be varied from low values up to tens of m/h. In order to reduce the temperature, telluride precursors with less stability, such as diisopropyltelluride (DIPT) and di-N-propyltelluride (DNPT), with a stability order DMT > DET > DNPT > DIPT, have been used. Vapor pressures are fixed by the liquid-vapor equilibrium of a liquid source at a fixed temperature, namely 12, 5, 6, and 2 torr for DET, DIPT, and DNPT, respectively, at 30  C [60]. Typical partial vapor pressures in the deposition chamber are 10
5-10
4 atm. To reduce the deposition temperature for preventing interdiffusion at interfaces, light-assisted MOVPE has been introduced, using UV light from a mercury lamp at 2 W/cm2 [61]. Epitaxial films have been obtained at 250  C with a remarkable improvement of the quality as observed by luminescence [62]. Precracking of the DET source at 650  C has been used also to reduce the deposition temperature of CdTe [63]. This reduction can be attributed to the photolytic decomposition of the precursor molecules and in particular of tellurium, as reviewed in 1994 [64]. DMT has been introduced as a new source for the same purpose. It shows decomposition at 275  C, about 150  C lower than DET [65]. In presence of DMCd, the formation of CdTe takes place

Heteroepitaxial Growth of CdTe Thin Films

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already at 250  C due to an adduct formation. New Te alkyl molecules like ditertiary butyl telluride (tBu2Te), with decomposition temperature starting at 240  C in presence of hydrogen, were also used for further decreasing the deposition temperature of CdTe to the same level as MBE [66]. Plasma-enhanced MOVPE of CdTe from DMCd and DET has been reported [67]. An overview of the state-of-the-art MOVPE of CdTe in 1990-1991 is given in Refs. [68, 69]. The evolution of the technique was to include the use of an impinging jet in order to get better control of the hydrodynamic conditions [69]. A review of MOVPE of (Cd,Hg) Te is given in [64], which highlights in situ monitoring techniques like laser reflectometry, spectroscopic ellipsometry, and reflection difference spectroscopy (RDS). Detailed parametric studies of the effect of partial pressures of DET (0.28–0.58 torr) and of DMCd (0.28–1.04 torr) at temperatures around 450  C on the growth rate and epitaxial quality (FWHM) of CdTe films have been carried out [70]. Similar studies using DIPT have been carried out [71]. Rapid thermal MOVPE has been introduced in 1997 for CdTe in the case of DMCd and DET or DIPT, at temperatures of 350–480  C, allowing growth rates up to 60 mm/h while maintaining high structural quality (205 arcsec) [72]. Studies dedicated to the growth mechanisms of epitaxial CdTe films by MOVPE are scarce. On can mention the work of Nemirovsky et al. [70] dealing with the MOVPE from DMCd and DET at a total pressure of 300 torr under a wide range of precursor pressures and for two temperatures, namely 430  C and 480  C. The substrate was lattice matched (111)B CdTe or CdZnTe. A crucial result is the effect of temperature. At 430  C, the structural quality of the layer is strongly degraded when increasing the growth rate (from 2 to 24 mm/h), while it maintains as low as 100 arcsec at 480  C, independent of the concentration of the precursors. Excess Te pressure at low temperature ensures formation of a smooth surface. This question of surface roughness, associated with hillock formation, has been addressed for (100) CdTe at 410  C by considering surface diffusion and step-motion elementary processes and their critical dependence with the Te:Cd ratio in the vapor phase [73]. Modeling of MOVPE growth of CdTe has been achieved from in situ laser reflectometry measurements. Figure 11 shows results obtained in the case of DMCd with either DMT (A) or DIPT (B) precursors. The simulation is carried out by using a kinetic model involving the thermally activated Te desorption step and decomposition reaction [74]. While much less developed than in the case of III-V compounds, extrinsic doping of CdTe is an important issue. Insertion of impurities in the epitaxial film is also the result of surface reactions, and thus sensitive to the crystallographic and polarity state of the surface. This has been exemplified in the case of As doping of MOVPE CdTe films as shown in Fig. 12 [75].

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GR Å/s

10

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0.1 0.0012

0.0014

0.0016 1/T K

A

0.0018

GR Å/s

10

1 GR (data) GR (theory) 0.1 0.0012

0.0014

0.0016 1/T K

B

0.0018

Figure 11 Growth rate of CdTe by MOVPE as a function of temperature, using dimethyl cadmium as a precursor and diethyl telluride (A) or disopropyl telluride (B). Continuous lines correspond to the results of modeling assuming kinetic limitations (adapted from Ref. [74]).

A

(11

1)B

B )B

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20° 10°

0°

10°

20°

30°

Tilt Angle

Figure 12 Crystallographic orientations and polarity dependence of As incorporation in MOVPE epitaxial layers of CdTe on GaAs (adapted from Ref. [75]).

Heteroepitaxial Growth of CdTe Thin Films

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3. SUBSTRATE EFFECTS ON CDTE HETEROEPITAXY One of the main applications of CdTe layer is its use as a buffer layer for the epitaxial growth of CMT layers used as IR detectors on various substrates. Cadmium zinc telluride (CZT) is a classical substrate which tends to be replaced by more available substrates as GaAs, Ge, and Si. The 1999s review of Irvine et al. [76] gives an overview of substrate layer relationships in the growth of II-VI compounds. It presents some aspects of the growth of CdTe layers on sapphire, GaAs, InSb, and Si in relation with other systems (ZnSe, CMT, etc.). It indicates that the quality of the epitaxy is not only related to the lattice mismatch, but also related to the polarity of the surface and its chemical effects. The growth of CdTe on InSb is the best example where combining the almost perfect matching and avoiding the antimony depletion at the surface leads to a record epitaxial quality of 20 arcsec! In the case of the GaAs substrate, the large mismatch provokes the formation of defects in the CdTe layer. Also when grown on (100) GaAs, (111) CdTe with A or B polarity together with (100) CdTe are formed, as from both MOCVD and MBE. The surface preparation is thus a key factor to prevent this formation of domains, as for instance thermal soaking, removal of oxides, and extent of prereaction of tellurium at the surface. In the following section, we will enter into more details by considering different substrates.

3.1. Growth on Ge There are only a few papers dealing with the epitaxial growth of CdTe on Ge. A detailed study of the growth of CdTe on (100) and (111) Ge by MBE was reported in 1985 [77]. No specific attempts to prepare the Ge surface were indicated. Large strain was observed for the growth of (100) CdTe films, which tends to relax toward thicknesses about 5-6 mm. FWHM at a growth temperature of 350  C was 1.8 for 1.5 mm thickness to 0.3 for 5 mm. On the contrary, films grown on (111) Ge presented a much better value of 0.25 for 0.8-mm-thick films, which became constant and equal to 0.18 for larger thicknesses. This was attributed to an effect of strong twinning on (111) surfaces, which was not present on the (100) ones. Smoother films were obtained on (100) Ge, making them more suitable for further devices. In 1997, direct growth of CdTe was also considered on (100) Ge by MBE [78], but by using a misorientation of 8 around the (110) direction. The growth was carried out at 350  C at a rate of 25 nm/min. The authors observed that the growth of CdTe corresponded to a (331) in-plane orientation, 4 off the (110) toward (111). A Te surface polarity was found. Twinning with parallel (111) planes along two variants a and b, one becoming dominant with the increase of thickness, was found. The FWHM of an 8-mm-thick layer was 170 arcsec. Further work [79] extended the study of

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the influence of the Ge orientation. Contrary to this, multidomain (111) CdTe was obtained on Ge (100) [70]. On (111) Ge, (111) CdTe is formed, but with a poor quality due to twinning. The quality is improved by setting and increasing a tilt up to a few degrees around the (110) direction. Best results were obtained by using directly (211) Ge, leading to a FWHM of 49 arcsec (13 mm). An important aspect reported in [79] is the control of polarity via surface modifications. Direct growth of CdTe results invariably in a Cdterminated surface (211)A, while a Te-terminated surface is needed for further growth of CMT. This is attributed to the preferential reaction of Te with Ge surface atoms. To force the B polarity, a two-step treatment was carried out starting with an arsenic derivatization of the surface, considered as passivating the Ge dangling bonds, followed by a zinc treatment at monolayer level completing the chemical coverage. Only in the case of step one, Cd polarity is still observed. One can presume that the Ge-As-Zn sequence determines the polarity. The layers are completely twin free with a well-defined interface presenting a regular misfit dislocation array of 60 type along (011). Further work shows the effect of thickness on the FWHM, with a linear decrease from 110 to 85 arcsec for 5.4-6.9 mm [80]. Jacobs et al. [81] also presented a study of the growth using (211) oriented Ge. In this case, As treatment is also employed for Ge. FWHM values are about 130 arcsec for a 7.4 mm film, with EPD of a few 107 per cm2. Thermal stress and strain effects are addressed in detail.

3.2. Growth on Si Studies on the epitaxial growth of CdTe on Si started in the 1980s with direct deposition by both MBE and OMCVD [82, 83] and then witnessed a rapid development with intensive competition between different groups. It appears that direct growth on Si introduced high-defect densities due to the large lattice mismatch. Intermediate layers made of graded CaF2-BaF2 were introduced on Si (111) [84], starting with only 0.6% mismatch for CaF2 up to 14% for BaF2, with a remaining 5% gap between BaF2 and CdTe. The use of these intermediate layers resulted in good-quality (111) CdTe films. A similar approach was used later on for the growth of (100) CdTe on (100) Si [85]. Another approach was using GaAs buffer layers for the MOCVD growth of CdTe (100) on Si (100) [86, 87] and CdTe (111) on Si (111) [88]. Note that in Ref. [87], (100) Si was tilted 4 off toward h011i. Sporken et al. [88] have investigated further the epitaxial growth of CdTe by MBE on Si (100) substrates. They showed that the direct growth resulted in (111)B CdTe while the use of a ZnTe intermediate buffer layer resulted in the growth of (100) CdTe. In the case of direct growth, two domains rotated by 90 were observed. A misorientation of 8 leads to a single domain structure. This was attributed to a change of the surface reconstruction of Si to single domain (2  1), and a

Heteroepitaxial Growth of CdTe Thin Films

101

reduction of the lattice mismatch to 3.4% in the [211] direction. The average FWHM achieved was 514 arcsec for a 6-mm-thick layer [89]. Ge has also been introduced as a buffer layer for the growth of CdTe (100) by MOVPE on tilted (100) Si substrates [90]. The growth of 1 mm Ge by a two-step low T/high T process resulting in a FWHM of only 140 arcsec was followed by a ZnTe layer growth. This procedure allowed improving further the quality of the (100) CdTe film reaching a FWHM of 260 arcsec for a 4-mm-thick film. The ZnTe layer was found necessary to achieve the best quality. A major improvement is related to the use of (112) oriented Si substrates. This orientation is more favorable for twinand hillock-free growth of CdTe. The effect of misorientation toward h111i, h010i, and h001i was examined. A 1-mm-thick ZnTe buffer layer was used again. The best results were obtained for a misorientation of 5 toward h111i, with a FWHM of 72 arcsec for an 8-mm-thick film. The EPD decreased abruptly from 3.107 cm
2 to 2.106 cm
2 for thicknesses from 2 to 4 mm and stabilized beyond [91]. At the same time the direct growth on misoriented (100) Si substrates was further improved with FWHM routinely less than 100 arcsec (best 78). In addition to the use of misoriented substrates, a two-step procedure was introduced for the growth of CdTe, including postannealing under Te flux at 360  C, leading to a strong recrystallization of the layer. This attractive approach is similar to the concept of recrystallized amorphous deposits for relaxed epitaxy [92] used in modern growth methods for chalcogenide layers. Strong reorganization of the layers is provoked by changes in deposition or postdeposition conditions (CdTe, CuInSe2). The complete fall-off of twinning takes place for only 2.5 mm beyond the CdTe/Si interface. A detailed study of the growth of CdTe (111)B directly on (100) Si by MOVPE shows new findings related to the effect of a pregrowth Te adsorption step on Si surfaces whose temperature controls antiphase domain formation [93]. This effect is also dependent on the misorientation. Twinning is found to depend on the Te step temperature and also on the ratio VI/II in the vapor phase with a minimum for a value of 8. Twinning was almost suppressed for 1.5-mm-thick films. This dependence of twinning on the VI/II ratio shows that twinning is not only due to the interface CdTe structure; the authors indicate that it is probably due to multinucleation on CdTe terraces. The tilting of (111) planes with respect to the (100) plane was also evidenced. The growth of CdTe by MBE on Si (211) was studied using a treatment of the Si surface with arsenic in order to react with dangling bonds and prevent the reaction with Te [94]. Then a 60 nm ZnTe buffer layer was classically deposited. The growth of CdTe benefited from the combination of growth and annealing steps at 370  C, as shown before, but with successive layers (IMAC process). The idea is to block the defect spreading from the interface into the growing layer. A clear improvement is observed and FWHM of 100 arcsec was achieved for 8 mm thickness

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with a best EPD of 5  106 cm
2. Similar multistep growth of CdTe, with an intermediate annealing under Te pressure, was also used in a subsequent study of MBE growth of CdTe (111) on Si (100) [95]. A 1 off-cut toward (110) suppressed the double domains, while increasing the azimuthal angle suppressed twin formation. Residual twinning disappeared within 2 mm from the surface. The As treatment was again reported but surprisingly followed by a Te treatment to treat the Si surface. The buffer layer was also annealed under Te pressure before the growth of CdTe. This resulted in a remarkably low FWHM (78 arcsec for 10 mm thickness) and a record EDP as low as 1.5  105 cm
2 for a 5-mm-thick film [96]. The quality of CdTe (111) directly grown on Si (100), despite the large lattice mismatch, has raised fundamental questions concerning the nature of interfacial reactions at the early stages of growth. One is that the atomic structure of the Si surface is not controlling the epitaxial orientation of the CdTe layer, and the lack of dislocations at the interface indicates that it is incommensurate [97]. This led to a hypothesis of weak interactions during growth, explained by the formation of a weakly bonded Te matrix during the initial modification of the surface, which then serves as an insertion matrix for the formation of CdTe by Cd insertion. This concept named graphoepitaxy is also supported by ab initio calculations [97]. In 1999, a new study for CdTe (111) on Si (100) by MBE gave new insight in the interfacial effects, thanks to STM measurements [98]. A key point was to treat the Si surface with Te vapor making a Te derivatized Si layer at 500  C. Then a CdTe layer of (111)B orientation was grown at 300  C showing excellent quality with FWHM of 100 arcsec. The surface structure was observed by STM also after removing the CdTe by sublimation at 500  C. The results were interpreted by a model involving the (2  1) reconstruction with the presence of Te dimmers. This is an important step towards the understanding of the growth mechanism, which has been completed by modeling the CdTe (111)B growth processes from such a surface structure. Previous models were also proposed in the case of the growth on GaAs substrates [99]. It is reported that by using a Te derivatization process at lower temperature, CdTe (111)A layers are formed preferentially. In that case a (1  1) structure of the Si surface is maintained. The growth of CdTe (111)A has been also studied on As derivatized (1  1) Si (111) surfaces in presence of a ZnTe buffer layer by means of RHEED, as shown in Fig. 13 [96, 100]. The advantage of the system is that single-domain films can be achieved even in the presence of steps. The effect of postannealing of the ZnTe layer under Te2 flux is shown to release the mismatch strain. Remarkable value of EPD: 2  105 cm
2 and FWHM (56 arcsec) attest the crystal quality, comparable to those on (100) surfaces or even better.

Heteroepitaxial Growth of CdTe Thin Films

A

B

C

D

E

F

103

Figure 13 Surface engineering procedure for the growth of high-quality CdTe (111)B on Si (111) as seen from RHEED measurements. (A) after As treatment in As4 atmosphere: Si (111) As, (B) after the growth of 20 nm ZnTe buffer at 220  C, (C) after annealing at 310  C under Te2 pressure, (D) at the onset of CdTe growth, (E) after 2 mm, (F) after 10 mm, onset of c(8  4) reconstruction (adapted from Ref. [100]).

Attempts to remove the use of the ZnTe layer for the growth on (111) Si have been the aim of specific studies. The surface was treated either with Te or As. In the first case this resulted in a poor quality (111)A film, while good quality single domain (111)B films were obtained with As [101]. An interface atomic model explains the results as shown in Fig. 14. New advances have been obtained by HWE in 2003 [102] for the direct growth of CdTe (111) on Si (111), with 200 arcsec for a 1.5 mm thickness, down to 118 arcsec for 5 mm. Contrary to MBE processes, as prepared H terminated Si surfaces were used. The presence of H is indicated to relax the strong bonding and to reduce the dislocation density. A two-stage process with two temperatures was used for the growth. A detailed study of the effects of HWE parameters was presented. In a subsequent study [103] it is shown that twin formation is strongly reduced when a first layer of CdTe is grown at a low growth rate of 0.01 mm/h followed by a second phase with a growth rate of 0.5 mm/h. The analysis showed that for a 5-mm-thick layer grown at 350  C the FWHM is 120 arcsec. This demonstrated the importance of the nucleation layer on the epitaxial quality. The direct growth of CdTe (211) on Si (211) by MOVPE is reported if Si surface is treated at high temperature in presence of GaAs in an H2 atmosphere [104]. Without this treatment the CdTe layers were polycrystalline with poor adhesion. The MOVPE was achieved using DMCd and DETe in a first stage at 415-450  C and then at 560  C, at growth rates of 2 mm/h and then 10-12 mm/h up to 100 mm thickness. The FWHM was 525 arcsec for 5 mm thickness down to 170 arcsec for 18 mm.

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B

Cd-Te

A C

As-Si Si

A

B A

C A B C B A

As

B

Si

Cd Te

D

Figure 14 Direct growth of CdTe (111)B on Si (111) showing atomic contrast image of the As passivated interface and stacking sequences for the two possible epitaxial configurations (adapted from Ref. [101]).

A study of the MBE growth of CdTe on tilted Si (211) was reported in [105]. No real evolution is given with respect to previous procedure. The Si (211) surface was prepared with an As pretreatment followed by the deposition of a ZnTe layer of 2.5 nm, nucleated at 200  C and annealed under Te pressure at 380  C and 15 min. As a result, 4.3-mm-thick layers were deposited. The FWHM presented a value of 83 arcsec for a deviation of 2.7 toward (111). The original point is that the authors also presented an analysis of strains in the CdTe layers as a function of the angle. Jacobs et al. [81] presented a systematic study of the growth using (211) oriented Si, again based on the established As derivatization and use of a ZnTe buffer layer. Stress relaxation was achieved by incorporating periodic anneals at 560  C, resulting into films with FWHM less than 100 arcsec. Dislocation reduction by about two orders of magnitude using periodic annealing was also demonstrated in another paper according to a procedure pictured in Fig. 15 [106], based on the same grounds as in a previous study [94]. The direct relation with defect reduction was clearly evidenced from EPD measurements and from luminescence studies, as shown in Fig. 16 [106]. The progressive control of the growth on Si is now used to grow highly sophisticated structures like HgTe/CdTe superlatices on CdTe/Si (211)B [1].

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Figure 16 Influence of structural defects on photoluminescence of CdTe epitaxial films. (A) CdTe (100) films grown by HWE on GaAs (100) as a function of thickness. The three main emission centers in the 1.43-1.5 eV range are interpreted as related to structural defects (including excitation/defect bounding in lamellae twins) (adapted from Ref. [138]). (B) CdTe deposited on Si, showing the influence of the number of introduction of in situ thermal annealing sequence (adapted from Ref. [106]) as shown in Fig. 15.

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CdTe/GaAs 4.2K Tsou = 743K Twall = 753K Deep Band Tsub = 623K Tres = 488K

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3.3. Growth on CdS The growth of CdTe on CdS is of key importance in the field of thin film solar cells, since CdS/CdTe solar cells are now industrially produced at the lowest production cost among all solar cell technologies, using the CSS technique and may be in the future MOVPE [107]. Even if the solar cell devices are based on polycrystalline layers, epitaxial growth is relevant at local levels. However the amount of work on this system is very limited. The epitaxial growth of hexagonal CdTe (0002) was obtained by HCl vapor transport at 500  C on to (0002) CdS as early as 1965 [108]. Later on, Simons et al. reported the epitaxy of CdTe films on (0001) and (0116) oriented CdS with A and B polarities by MOVPE at 325  C from DMCd and diPTe [109]. Polycrystalline growth occurred on A surfaces while epitaxial growth took place on B surfaces with (111) CdTe/(0001)B CdS. A large number of twins were present parallel to the surface. Best results were obtained on (0116)B CdS. This marked differences as a function of polarity were also observed for homoepitaxial growth on (111) CdTe or on (111) GaAs. Epitaxial growth of (111) CdTe has been obtained by electrodeposition on (0001) CdS from a chemical bath on (111)B InP single crystals [110]. Further experiments have been carried out to study the onset of nucleation of CdTe on single crystalline CdS substrates [111]. It was shown that a Te derivatized surface was formed preferentially on (0001)B CdS.

3.4. Growth on GaAs The epitaxial growth of CdTe on GaAs substrates has been investigated much more than on any other substrate, by both MOVPE and MBE. There is an abundant literature ranging from 1981 up to the end of the 1990s. This reflects the availability of good quality GaAs susbtrates and the fact that the lattice mismatch of 14%, even if still high, is significantly lower than on Si. Another remark is that contrary to Si, MOVPE growth is as much used as MBE, and even a little bit more. Early work reported the growth of CdTe (100) on GaAs (100) by MOVPE [112], while the growth of CdTe (111) on GaAs (100) by LADA was demonstrated as early as 1983 [19]. In 1984, CdTe (100) was grown again on GaAs (100) by MOVPE, illustrating immediately a key fact related to the versatility of the epitaxial mode [113]. Similar results were obtained by MBE [17]. The role of Te adsorption, at the level of 1-2 monolayers as shown by RHEED measurements, was found to be the cause of the (111) growth mode [14]. The role of Te was immediately contested from HRTEM studies and a direct growth was postulated on clean surfaces. In the case of (100) growth mode, the presence of a thin oxide layer, due to insufficient cleaning of the GaAs surface, was found to be the cause

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of this growth habit [114]. These hypotheses were again contested in a study where no evidence of interfacial oxide was found for the growth by PLD of (100) CdTe on (100) GasAs [115]. However, the hypothesis of the control of the (111) growth by a Te interfacial monolayer was again supported, with the formation of a Ga-As-Te superficial phase. Depending on composition variations, (111) or (100) growth can be initiated [116]. Reno et al. [117] also reported that high temperature backing of GaAs (100) at 560  C induced the formation of CdTe (100). In fact a surface chemical model stating the bonding of Te on the surface with two different atomic structures leading either to (111) or (100) allowed to put light on this hot question and explained the experimental results reported so far [99]. However, the quality of the layers was limited by twinning, inherent to this substrate orientation [118]. The interface structure was revealed at the atomic level by HRTEM studies; an abrupt interface with periodic dislocation arrays was observed for (100) growth, while for (111) defects seemed to play a significant role. Substrate temperature control was pointed out [119]. The growth of CdTe (111) on (111)B GaAs was reported in 1987 [120, 121]. The use of buffer layers was then introduced using the MBE deposition of a (100) ZnSe layer. This resulted in the growth of (111) CdTe [122]. The use of (100) GaAs tilted 6 along [110] allowed removing twinning by suppressing multipositionning effect inherent to flat surfaces. It was shown that the growth of (111) CdTe was related to the flow along terraces. Surface reconstruction c(4  4) was obtained by high-temperature annealing under As pressure. A model at microscopic scale allowed explaining and discussing further the beneficial effect of misorientation [123]. At this stage, polarity control of the (111) growth was not established. This has been done using Se or Zn stabilized ZnSe (100) buffer layers on (100) GaAs [124]. The Zn or Se termination controls the selective initial adsorption of Te or Cd respectively, leading to A or B (111) CdTe growth. The quality of epitaxial layers on (100) and (111) GaAs was reported through FWHM measurements. Values of 750 arcsec (432) were obtained for 0.5 (1.5) mm thick (111) CdTe [125]. The growth by HWE on GaAs resulted to (111) for 580  C and (100) below 520  C with FWHM of 4800 on (111) planes for 3.8-mm-thick layers [126]. The role of a Te-adsorbed layer in fixing the (100) preferential orientation was further discussed by arguing that starting with DMCd the (111) growth of CdTe was favored on (100) GaAs [127]. The use of (211) GaAs substrates was then introduced [128], showing also a dual epitaxy effect for CdTe (211) and (133). On the basis of HRTEM results an atomistic model was proposed to explain this effect [129]. The use of a (100) ZnTe buffer layer was then described for the growth on GaAs (100). Contrary to CdTe, ZnTe shows only parallel epitaxy. This was found to force the growth of (100) CdTe [130]. The effect of surface microstructures on the (100) growth of CdTe on (100) GaAs was studied in [131]. While the

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growth of (111) CdTe gave more faulty layers, either on (111) or (100) GaAs, the growth of (100) CdTe was reported to depend on the microstructure of the superficial (111) CdTe layers. Yin et al. studied in more detail the dual epitaxy effect on GaAs (211) using B and A oriented surfaces [132]. They found that the (133) orientation was favored below 280  C for GaAs (211)A and (211) for temperatures higher than 280  C. The transition temperature increased slightly to 295  C for (211)B. FWHM for 2-mm-thick films was found of about 100 arcsec denoting a very good structural quality. The authors developed an analysis of the origin of these results in terms of interfacial energy changes between the different orientations. Mazzer et al. investigated further the effect of the ZnTe buffer layer and demonstrated its effectiveness in blocking spreading dislocations from the GaAs substrate. They obtained 900 arcsec for CdTe films of 1 mm [133]. Patriarche et al. confirmed these findings and found an additional beneficial effect in introducing ZnTe layers inside the CdTe layer, according to a superlattice like configuration [134]. This resulted in interdiffused CdZnTe layers, like in the interdiffused multilayer process (IMP). The density of dislocations was in the range of 104-105 cm
2 which is compatible with the application requirements. Nishino et al. showed that the deviations from stoichiometry in the ZnTe layer had a role in fixing the quality of the CdTe layer. Best results were obtained with a Te/Zn ratio of 0.6 [135]. The deposition of CdTe on (100) GaAs in an atmospheric MOVPE system was studied using DMCd and DETe at 360  C and 420  C in presence of H2 [136] with the VI/II ratio changing from 0.5 to 5. Normal growth with simultaneous introduction of Cd and Te precursors and Te stabilized growth conditions occurred with only DETe starting at 420  C for 3 min. The authors also considered the influence of prebaking the surface at 620  C for 10 min. It was clearly demonstrated that low temperatures below 390  C gave preferentially (100) CdTe with both normal or Te stabilized initiation, while higher temperatures could lead to (111). Prebaking led to the formation of (111) CdTe even at low temperature. In the case of normal growth, (111) was favored at high temperature by a high VI/II ratio. This was attributed to the chemistry changes at the surface, especially the Ga/As ratio. Lowering the As content promoted the growth of (111) orientation. Te binding also controlled the growth in agreement with former studies. The authors even refined the nucleation step by introducing an ALE growth sequence at 420  C starting with DMCd and then DETe. As shown in Fig. 17 the (111) growth was then completely suppressed and the FWHM decreased from 800 arcsec down to 200 arcsec. This demonstrated that the nucleation step is very important to control. An investigation by HRTEM of the dual epitaxial effect of CdTe on (100) GaAs has been reported [137]. CdTe layers were grown by MBE at different substrate temperatures. The parameters investigated for the GaAs substrates were the effect of a

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Figure 17 X-Ray diffraction study of the effect of surface treatment on Dual Epitaxy of CdTe (111)/(100) on GaAs (100) by MOCVD at 420  C. In presence of a nucleation layer of CdTe the (111) orientation is completely suppressed (adapted from Ref. [136]).

chemical treatment, which affected the flatness, after vacuum baking at 600  C for 5 min for removing the oxides. The transition between (100) and (111) growth was clearly shown when increasing the temperature with a steep transition temperature of 20  C at 290  C as shown in Fig. 18. In the transition range both modifications were observed in different zones. The balance between them was dependent on the chemical state of the surface, with or without chemical treatment or with residual oxides. Lattice mismatch dependent strain has been studied from HWE growth of (100) CdTe on (100) GaAs [138]. There was no specific surface treatment of GaAs. The oxides were removed by baking at 853 K for 30 min. The authors showed that compressive strain is present in the film up to about 5 mm thickness due to lattice mismatch. FWHM decreased (to less than 100 arcsec) as a function of the increase of the film thickness in direct relation with the improvement of the structural quality. Photoluminescence spectra were dependent on the structural quality and most of the low-energy peaks were related to structural defects (Fig. 16).

Heteroepitaxial Growth of CdTe Thin Films

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Figure 18 Illustration of temperature threshold on Dual Epitaxy effects for the growth of CdTe on (100) GaAs. (A) HRTEM image of (100) CdTe/(100) GaAs for a growth at 310  C after cleaning and prebaking at 600  C for 5 min before MBE growth, but not chemically etched. (B) HRTEM image of (111) CdTe/(100) GaAs for a growth at 280  C after the same treatment (adapted from Ref. [137]).

The growth of CdTe films was carried out on GaAs (100) substrates covered with a 0.5-mm-thick ZnTe layer prepared by MOVPE using a hydrogen-assisted VPE method and controlling CdTe source and substrate temperatures [139]. Detailed parametric results are given for the growth rate as a function of the source and substrate temperatures. The process is thermally activated, with an activation energy of 33 kcal/mol at a source temperature of 827  C. A maximum growth rate of 10 mm/h is obtained. The curves were modeled. Very smooth films were obtained, with FWHM of 60 arcsec for a 27-mm-thick film and 500 arcsec for 5 mm thickness. Instead of using the ZnTe buffer layer approach to limit the competitive formation of (111) and (100) CdTe domains, a sulfur treatment has been proposed [140]. After etching its surface in the H2SO4: H2O2:H2O solution, the GaAs substrate is dipped in an ammonium sulfide solution for 3 min at room temperature. The substrates are then loaded in the MBE chamber and heated at 450  C for 10 min to remove excess sulfur. An initiation CdTe layer of 100 nm thickness was grown at 280  C and then a 300-nm-thick CdTe layer at 320  C. After this process the RHEED pattern was faint and broad (2  1). The temperature was then ramped up to a postannealing temperature of 360  C for 30 min under slight Cd overpressure to prevent Cd losses by sublimation. A clear improvement of the structural quality displayed through RHEED patterns and photoluminescence was obtained. The growth of CdTe on GaAs (100) with Zn pretreatment illustrates the variety of approaches for surface modification of the GaAs surface [141]. The aim was to make waveguides with CdMnTe. The method developed for the epitaxy of

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ZnSe on GaAs is adapted, where a treatment of the GaAs surface was achieved using a Zn predeposition. This allowed preventing the formation of a Ga2Se3 layer by direct contact between GaAs surface and Se vapor. An amorphous Zn layer was formed at 60  C and controlled by RHEED. The growth of CdTe proceeded at 265  C. The structure is CdTe (1mm)/ZnTe (1nm)/GaAs (2  1) reconstructed. Using XRD measurement the FWHM and the twin content were detected. The improvement of the FWHM was from 2.2 arcmin to 0.7 arcmin and from 0.2% to 0 for the twin density. This beneficial effect is related to the absence of Ga2Te3 formation at the interface which tends to act as initiating threading dislocations and stacking faults. Surprisingly, the use of Sb-treated GaAs surface led further to opposite conclusion [142]. It was found that, thanks to an intermediate layer of GaSb which has an intermediate lattice parameter between GaAs and CdTe, the epitaxial growth of (100) CdTe by MOVPE was obtained. The smooth surface was attributed to a surfactant effect of Sb. Note that previous studies have been carried out on the deposition of (111) CdTe on (111)A GaSb substrates [143], with a lattice mismatch of 6.4%. CdTe was deposited from an evaporation process from the compound. The interface layer is Ga-Te. The films grown through a Stransky-Krastanov mode with Sb segregation were not as smooth as reported in Ref. [132]. A new comprehensive study of the thickness effect in the MBE CdTe growth on (211)B GaAs in comparison with (211) Si has been reported [4]. The lowest FWHM, observed for an off angle of 2 , has a value of 50 arcsec from about 5 mm thickness (Fig. 1). There has been an emphasis for the growth of mm thick epitaxial CdTe layers on GaAs substrates for X- and gamma-ray detector applications. The growth of thick epitaxial films of CdTe by MBE has been considered by Sorgenfrei et al. [144]. They have used (100) oriented GaAs substrates to grow (100) CdTe layers up to 100 mm at about 10 mm/h. A hightemperature back up of the GaAs surface at 610  C for 10 mn was used to prevent the dual epitaxy with additional (111) orientation. Jiang et al. [145] report a systematic study of the epitaxial growth of CdTe layers on (100), (111)A and (211)B GaAs substrates from a modified VPE method, named MTPVT specially designed for fast deposition of thick CdTe layers (>100 mm/h). On (100) GaAs the growth of the CdTe can follow the dual epitaxy effect and (100) and (111) CdTe domains were observed. Concerning the formation of (100) oriented CdTe films up to 100 mm thick, the FWHM was about 0.35 . On (111)A GaAs, the growth of pyramidal surface shaped CdTe (111)A was observed with a FWHM of 250 arcsec for a 10-mm-thick film. In the case of (211)B GaAs, the epitaxial quality was found to be dependent on the etching treatment, HCl treatment giving superior properties as compared to HF with FWHM of 0.15 . The strain and FWHM values have been measured as a function of the

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thickness of the layer up to 20 mm thickness. Both decreased markedly with the increase of thickness. Despite the results are less good in terms of structural quality of the thinnest layers, this method gave indications on the growth of thicker epitaxial layers. The epitaxial growth of thick epitaxial layers has also been addressed using the CSS method on GaAs substrates [35, 146]. Dual epitaxy was observed on (100) GaAs. The best results have been obtained on (211)A CdTe grown on (211)A GaAs with FWHM of 320 arcsec. The growth rate was about 120 mm/h and FWHM on thick layers was 35 arcsec, close to the one of single crystals. This opens the route toward integrated X-ray detectors on GaAs substrates.

3.5. Growth on NbSe2 As a last example of heteroepitaxial system let us consider the case of NbSe2 substrates [147]. As reported before, the main approach for epitaxy is strong bonding between the chemical elements of the substrate and those of the layer. Another route for the epitaxial growth is to use substrates with low bonding properties with the epilayer. This can be realized with layered materials and is called Van der Walls epitaxy. The growth of CdTe (111) layers on (0001) NbSe2 layers has been studied, in spite of a lattice mismatch of 34%. The growth was performed by MBE under Cd and Te elemental fluxes with Te excess at a substrate temperature in the range 280-330  C. A two-step process was used starting first under Te flux only at 100  C followed by the progressive increase of temperature to 300  C and then the introduction of the Cd flux. The growth was followed in situ by RHEED. The results show that very smooth and relaxed 2D (111) CdTe films are grown within plane epitaxial relationship with the surface lattice of Se atoms, with a commensurate matching in the ratio 3:4.

4. OUTLINE AND CONCLUSIONS Heteroepitaxial growth of CdTe has been the subject of continuous studies for about 30 years. This review has recalled the different steps in the advance of research. After an impetuous growth of studies in the 1980s and 1990s, where different key methods like MBE and MOVPE have been introduced for this material together with the use of foreign substrates like GaAs and Si, basic knowledge of the growth processes has been mastered, allowing specific applications in infrared detection technologies. CdTe was in that case mostly used as an intermediate buffer layer for MCT active layers deposition on GaAs or Si substrates. A high level of mastering has been reached, as illustrated by the control of superlattices formation with HgTe, ZnTe, and MnTe. However, this limited role of CdTe has also reduced the attention paid to CdTe epitaxy and material

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properties per se, such as doping control. Even if an important amount of knowledge has been gained for the growth on GaAs and Si substrates, with the specific effects of surface modifications and specific growth procedures, the level of control and knowledge of CdTe and II-VI heterostructures formation and properties beyond the combination with HgTe is still limited as compared to III-V heterostructures. This comes from the more limited applications of II-VI—mainly as infrared detectors for military applications—as compared to III-V used for a wide range of optoelectronic applications from the UV to IR radiations, up to now. It seems also that CdTe and II-VIs are more difficult to control than Si or III-V compounds. As softer materials they often need in situ annealing treatments, as shown all along this review, which are more difficult to control. This is also highlighted by annealing treatment of CdTe films for solar cells in presence of cadmium chloride which improves dramatically the electronic properties. May be further work on film growth of CdTe will have to move away from classical growth methodologies, based on “rigid materials” concepts, as developed for III-V compounds, and involve “soft material” new approaches. Such approaches are already mastered for the elaboration of chalcopyrite absorber Cu (In,Ga) Se2 for high-efficiency thin-film solar cells, as known as multistage processes [148]. This may pave the way for the elaboration in the future of very high efficiency II-VI heterostructures and multijunction solar cells, including CdTe layers, either in single crystalline form or in multicrystalline/ microcrystalline one.

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CHAPTER

IE Chemical Treatment of the CdTe and ZnxCd1
xTe Surfaces V.N. Tomashik and Z.F. Tomashik

1. INTRODUCTION The development of modern semiconductor materials devices is closely related with the advancement in processing and preparing of semiconductor surfaces. The development of high-quality surfaces of semiconductor materials, which have perfect structure and geometry as well as homogeneous chemical composition, has exceptional importance for the manufacturing of different semiconductor devices. One of the methods that can be used for preparing the surfaces with necessary properties is chemical wet etching [1, 2]. Chemical etching is widely used in modern semiconductor technology as a method for examination of bulk single crystals and thin epitaxial films applied for the manufacturing of various microelectronic devices. The dissolution of semiconductor materials, which includes the breaking of surface interatomic bonds, is based on the difference in the reactivity of the material’s constituents toward the etchant. The factors determining this process are the thermodynamic parameters of the material, the nature of the solvent, the process temperature, the way the etchant is supplied to the dissolving surface, the nature and type of bonding between the structural parts of the crystals, and the crystallographic orientation of the surface [3]. In spite of that, chemical etching is a relatively simple process for all practical purposes for obtaining high-quality polishing and structurally perfect and defectless surfaces of the semiconductor materials, while retaining the needed geometrical parameters, which are associated with complex scientific and technical problems [1]. Many aspects of

V. Ye. Lashkaryov Institute of Semiconductor Physics, National Academy of Sciences of Ukraine, 03028 Kyiv, Ukraine

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developing structurally perfect and defectless surfaces of the semiconductor substrates are not completely investigated and are of utmost importance at the present time. In 1975 [4], it was noted that suitable etchants for the chemical treatment of the semiconductor materials were selected using a trial-and-error method and chemical polishing was often considered an art rather than a science. Since the last more than three decades, such definition of chemical polishing practically does not hold true; now, the success in the achieving the formulated aim very often depends on the experimentalist skill. According to the data in Ref. [2], 60% of the failure of semiconductor devices depends on defects of the phase boundaries and about 40% is due to the surface treatment of the semiconductor working elements. Chemical etching, based on the chemical dissolution processes, is one of the main technological operations in the chemical treatment of semiconductor single crystals and thin films, which are widely used in the manufacturing of different semiconductor devices and integrated circuits. The knowledge of kinetic behavior, mechanism, and nature of semiconductor dissolution is the most important factor for the selection of corresponding solution composition for polishing, anisotropic or selective etching, and chemical cutting. High resolving power of some etchants allow their usage on different phases of the substrate treatment, but for this purposes, it is necessary to develop the etchants corresponding to material removal, surface roughness, and some other parameters [1]. For silicon, technological methods of substrate preparation including electric spark or wire cutting of the ingots, decreasing thickness of the obtained disks using abrasive powders, and polishing with the help of diamond pastes, appear unusable for the treatment of the plastic semiconductors. Cadmium telluride and solid solutions based on it are examples of such materials. At present, the methods for substrate manufacturing of the plastic semiconductor materials using only wet chemical etching are developed [5, 6]. The technological scheme of the process includes chemical cutting when the etchant is supplied to the cut using a wire, disc thickness decreasing with the help of the device for chemical mechanical polishing, and a final chemical treatment. It is necessary to develop the etchants for each of these procedures. The developed etchant must ensure a suitable rate of material removal and that the roughness of the obtained substrates does not corrode the wire and the polishing materials. The rate of semiconductor dissolution is the quantitative characteristic of the etching process and one of the main etchant properties. Therefore, it is necessary to investigate the concentration dependence of semiconductor dissolution for all etchant compositions. Such dependencies for three-component etchants can be represented as the surfaces of equal etching rate.

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The investigation of the processes occurring at the “semiconductor/ etchant” interfaces is of utmost importance, as according to the data of numerous studies, films are formed on the surfaces of the chemically treated single crystals [7]. The composition, structure, and thickness of these films depend on the etchant composition and etching conditions. For example, the film formed on the CdTe surface as a result of dissolution in nitric acid is composed of three layers. The outer layer with a thickness of 2–3 nm contains 20 at% Cd, the near-surface layer consists of a mixture of TeO2 and Te, and the innermost layer is enriched by Te.

2. BROMINE AND IODINE CONTAINING ETCHANT COMPOSITIONS Solutions of elemental bromine in organic and inorganic solvents (ethanol, methanol, dimethylformamide, hydrobromic acid, etc.) are most frequently used for the surface etching of CdTe and ZnxCd1
xTe solid solutions [8–14]. Such mixtures possess polishing properties and their etching rate is limited by the diffusion stages of the heterogeneous interaction. They are characterized by rather high dissolution rates of CdTe, and their components are volatile and toxic substances. The treated surface is, however, Te rich, slightly oxidized, and contaminated with bromine. Polishing with Br2-methanol also results in haze or orange-peel-like surfaces and corrodes polishing machines [15]. According to the data of Refs. [10, 16], the dissolution of CdTe in the bromine-containing etchant compositions includes two conjugated mechanisms: chemical dissolution with the formation of CdBr2 and TeBr4, and electrochemical dissolution with the formation of elementary tellurium. As a result of the CdTe treatment in the dark using 0.5% solution of Br2 in CH3OH, a layer with a high quantity of cadmium vacancies [(1–4)  1014 cm
3] is formed on the surface [17], which could be removed electrochemically [18]. The treatment in such solution at the illumination leads to the detection of subgrain boundaries for all surfaces in both CdTe and Cd1–xZnxTe solid solutions [19]. A tellurium layer with the thickness of 1-4 nm is formed on the CdTe surface after etching using 5% solution of Br2 in CH3OH [20, 21]. The surface layers are also enriched by Te when etched of CdTe and Zn0.05Cd0.95Te in a 2% solution of Br2 in methanol [22–25]. Similar results were obtained in Refs. [26, 27], where the determining ratio [Cd]:[Te] is equal to 0.43 after 2 min of etching in the solution containing 1% Br2 in CH3OH. The treatment of the CdTe surface using 0.02% solution of Br2 in CH3OH leads to the formation of the amorphous tellurium layer with the thickness of 0.5–0.7 nm [28, 29]. A rinsing in the solution containing NaBH4 diminishes the layer thickness by up to 0.25 nm. The thickness

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˚ depending on the bromine of this tellurium layer could be reduced to 40 A content in the solution and the time of the chemical treatment [26]. A mixture of Te and TeO2 was obtained on the CdTe surface after chemical treatment for 30 s in a solution containing 50 vol% of Br2 in CH3OH [17]. The quantity of TeO2 in such films could be diminished by heating the surface in vacuum at up to 450  C. According to the data of Ref. [30], the layer, which is formed as a result of CdTe chemical treatment in 0.5% solution of Br2 in CH3OH, is enriched by cadmium ([Cd]:[Te] ¼ 1.42). Adding of AgNO3 to the bromine-methanol etchant, diminishes the value of [Cd]:[Te] to 1.1-1.2. According to the data of Ref. [31], the surface layers forming on CdTe as a result of chemical etching in the solution containing 12% Br2 in CH3OH, are enriched by Cd ([Cd]:[Te] ¼ 1.58). The thickness of such a layer is ˚ , but on exposure to air, it increases according to the equation equal 10 A d ¼ a þ b ln t, where t is the time of the exposure, “d” is thickness, “a” and “b” are constants. This dependence indicates that the thickness of the surface layer increases linearly with time. Increasing the tellurium content on the surface with time could be explained by its diffusion from the bulk. Cadmium washing out occurs at the chemical etching of CdTe in the Br2-methanol and Br2-butanol solutions as a result of Cd2þ and Br
interactions [32]. Tellurium clusters and possibly pores are formed when the Cd2þ concentration reaches a critical value. The layer of TeO2 is formed after treatment in such solutions in the presence of water. The rates of etching, oxidizing, and tellurium enrichment for A and B are not the same due to the difference in concentrations of Te2
on these surfaces. The rinsing of the etched in the bromine containing solutions CdTe surfaces in methanol for 20 min completely removes Br2 and the ratio of [Cd]/[Te] decreases from 0.7 to 0.5 [33]. The results of the X-ray photoelectron spectroscopy indicated the formation of TeBr4, which is poorly soluble in such solutions. The bromine adsorbed on the surface could penetrate deep into the film. The reaction products formed on chemical etching of CdTe can be also dissolved in hydrobromic acid. The composition of 8% Br2 in HBr for the chemical polishing of CdTe was proposed in Ref. [9]. It was established that the dissolution of CdTe in Br2-HBr solutions is similar to that of HNO3-HBr solutions at [HNO3]/[HBr] < 1 [10]. At low Br2 content, the etching rate is limited by the bromine diffusion to the surfaces, and at high Br2 content it depends on the dissolution rate of the forming CdBr2 film. Adding of lactic acid to the bromine-methanol etchant compositions improves the properties of the X-ray and g-ray detectors based on Cd1–xZnxTe solid solutions [34]. The ingots were etched in a 5% solution of Br2 in CH3OH [35] or in a 2% solution of Br2 in 20% solution of lactic acid in ethylene glycol [36] after mechanical treatment by Al2O3 abrasive with grain diameter of 0.5 mkm and ultrasonic cleaning in acetone

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or methanol. The etching rate of n-type Cd1–xZnxTe solid solutions at x ¼ 0.1–0.2 in 2-10% solutions of Br2 in methanol was equal to 50 mkm/min at room temperature [37]. The etching time or the surface layer removing depth is also a significant parameter in the manufacturing of detectors of ionization radiation [38]. It was established that chemical etching by bromine solutions in ethylene glycol for 6 min is optimal and significantly increases the yield of the quality detectors [39]. The supplementary surface states, which can be removed by electrochemical treatment of the surface, strongly influences the characteristics of the obtained surfaces by chemical etching in such solutions [18]. The etchant composition, containing 4.15 g KI þ 0.5 g I2 þ 12.5 ml HBr þ 25 ml ethylene glycol þ H2O up to a volume of 50 ml, can be used for the chemical treatment of CdTe, HgTe, and CdxHg1
xTe solid solutions [40]. The composition of the surface layers formed after the treatment in such solution is similar to that obtained by treatment using bromine-containing solutions (0.1 vol% Br2 in HBr). On determining the dependence of the etching rate on the time of solution ageing, the HI-based etchants were found to be more stable than Br2-based solutions [15]. A small drop of the etching rate with time of solution storage was observed after tens of hours and then the rate remained essentially unchanged (Fig. 1). 10.0

Etching velocity (μm/min)

9.5 (111)A (111)B (110) (100)

9.0

8.5

8.0

7.5

7.0

0

100

200 Time (hours)

300

400

Figure 1 Dependence of CdTe etching rate on the time of solution storage at 300 K for 8 vol% H2O2 þ 65 vol% HI þ 27 vol% citric acid solution (the rotation speed of disc at 84 rpm).

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3. ETCHANT COMPOSITIONS BASED ON NITRIC ACID Nitric acid as well as its mixture with CrO3, K2Cr2O7, and H2SO4 is often used as an oxidant in the etchant composition for the treatment of CdTe and based on it solid solutions. This acid is an oxidant at a concentration of more than 2 M [5]. According to the data of Ref. [41], the thickness of a layer forming on the CdTe surface after chemical treatment with HNO3 for 2 min is equal to about 50 nm and can be removed by etching it in bromine-methanol solution. Rinsing in concentrated nitric acid leads to the formation of a film on the CdTe surface consisting of TeO2 and Te [42]. Cadmium was not observed in the surface layer after the CdTe chemical etching in 15% HNO3 [43], and the thickness of such layer reaches up to 60–80 nm [30]. Selective etching is used to determine the structural defects in single crystals of CdTe and solid solutions based on it [44]. Dislocation on the CdTe (110) surface can be determined by chemical treatment in HNO3 for 2 min [9]. The surface layer forming on the p-CdTe surface is characterized by resistance that is two orders of magnitude higher than that of the bulk material, and its thickness is equal to 500 nm [36]. Etching in solution containing 1 p. 65% HNO3 þ 79 p. 85% H3PO4 þ 29 p. H2O þ 1 p. 100% HNO3 þ 85 p. 85% H3PO4 þ 33 p. H2O leads to the formation of a surface enriched by Te; this also leads to an increase in the sensitivity of solar cells based on the structure of CdTe-Sb/Au by up to 12.5%, which is better than that obtained by the treatment using Br2-methanol solutions [45]. Density of the etching pits on the CdTe surface can be determined using chemical treatment in the solution with a composition of 60 vol. p. H2O þ 1 vol. p. chromic acid þ 1 vol. p. concentrated HNO3 þ 1 vol. p. concentrated HF [46]. The surface of the ingots was previously mechanically polished by Cr2O3 powder and then chemically polished in the bromine-methanol solution containing 1% Br2. According to the data of Ref. [47], chemical etching of CdTe in Na2Cr2O7 solution in 3 M HNO3 diminishes the surface-state concentration. The high quality of the CdTe surface with roughness of Rz 0.07 mkm can be obtained by chemical-mechanical polishing using a solution containing 2–6 p. K2Cr2O7 þ 8–15 p. HNO3 þ 90–80 p. H2O [48]. According to data of Ref. [49], one of the most frequently used solution for CdTe surface chemical polishing is solution E, which is composed of 20 ml HNO3 þ 10 ml H2O þ 4 g K2Cr2O7. The CdTe dissolution rate decreases from 12–14 [50] to 4-7 mkm/min [30] on the addition of K2Cr2O7 to nitric acid. Refs. [51, 52] indicate the formation of an adsorbed layer of chromium on the CdTe surface after its treatment using the E

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125

solution. The ratio of [Cd]/[Te] in the surface layer is equal to 0.08 and the tellurium film layer thickness reaches 30–60 nm [30]. It was determined that a ratio of concentration peaks of [Te]/[Cd] in an XPS study is more than 9, and tellurium exists in the near-surface region in form of Te4þ ions [53]. The ratio of [Te]/[Cd] in this region is equal to 1.6, and in the layer with a thickness 0.7 nm it reaches up to 3.5 [54]. It is possible to determine the (111) crystallographic orientation of n-CdTe [55, 56] and to define A and B planes [57] using such etchant solutions. A surface layer forming on the CdTe surface on chemical etching with the E solution is composed of Te and TeO2 [58–60]. On adding 0.5 or 10 mg AgNO3 to the E solution, EAg I and EAg II etchant compositions for determining the thin crystal structure could be obtained. But Refs. [9, 61– 63] indicate that neither of these two etchants helps in determining a real dislocation density. The formation of etching pits by the action of EAg I and EAg II solutions is governed by the fact that upon addition of Agþ to the E solution, the latter becomes an anisotropic etchant [49, 64]. The plane (111) is etched slowly in the EAg I etchant, and on increasing the Agþ concentration (EAg II etchant), the plane (1 1 1) is characterized by the slowest etching rate. The etching of CdTe in these solutions is diffusion limited [2], and the influence of the diffusion stages is so significant that a solution stirring can generate effects comparable with those caused by increasing the concentration of Agþ [65]. That means that without supplementary investigation of the kinetics and mechanism of CdTe treatment using EAg solutions, it is impossible to interpret unambiguously the etching pits morphology. EAg II can be also used for determining the dislocation density on the plane (111) of the Cd0,96Zn0,04Te solid solution [49]. Instead of EAg I etchant, solutions containing 3 p. HNO3 þ 2 p. HF þ 1 p. H2O [66] or 3 p. HNO3 þ 1 p. HF þ 5 p. 0.5% solution of AgNO3 [60] can be used for an identification of the CdTe polar plane. The etching of CdTe in such solutions lead to the formation of triangular (pyramidal) pits on the (111) Te plane, while the (111) Cd plane is mirror smooth. It is necessary to note that the etching results obtained by the using such solutions are well-reproducible. Adding HF to HNO3 leads to a surface enrichment of tellurium, which can be removed by rinsing in solutions containing 7 p. K2Cr2O7 þ 3 p. H2SO4 [67]. The solution with a composition of 1 p. HF þ 1 p. HNO3 þ 2 p. H2O can be used for the formation of an electroconductive tellurium layer on the CdTe surface during the manufacturing of devices based on Te/CdTe heterostructures [68]. According to the data of AES, the CdTe surface etched in the solution containing 1 p. HF þ 1 p. HNO3 þ 2 p. CH3COOH is enriched by cadmium and [Cd]/[Te] ¼ 1.42 [30, 22, 69]. The etching rate at 300 K is equal to 60 mkm/min.

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To increase the CdTe etching rate, it is necessary to add effective complexing agents to nitric acid, and hydrohalogenic acid (HF, HCl, HBr, and HI) are the most frequently used as such agents [11, 70, 71].

4. ETCHANT COMPOSITIONS BASED ON Cr(VI) COMPOUNDS Chromium (VI) compounds, which are strong oxidants, are widely used in etchant compositions for the treatment of both elementary and complex semiconductors [72, 73]. Kinetic regularities of CdS, ZnS, ZnTe, and CdTe dissolution in mixtures containing 1 p. 2 N K2Cr2O7 and 1 p. 11.3 N HCl were investigated in Refs. [74, 75]. The correlation between the etching rates in the temperature range 298–328 K is as follows: VZnS < VCdS < VZnTe < VCdTe. The etching rate of ZnS is independent of solution stirring when the etching rate of CdS, ZnTe, and CdTe increases in the order of magnitude 0.2–0.4. Aqueous solution containing 20-50 mass% CrO3 þ 32-37 mass% HCl þ 0.08–0.15 mass% NH4F was developed for chemical polishing of II–VI semiconductor compounds [76]. The mixture of 70 vol% of K2Cr2O7 saturated solution and 30 vol% H2SO4 is also used for CdTe surface treatment [77]. A low quantity of tellurium is present on the treated surface, which can be removed by a hot solution of NaOH þ Na2S2O3. After CdTe chemical polishing in the solution containing 7 g K2Cr2O7 and 3 g H2SO4, the ratio of [Cd]/[Te] is equal to 0.9 and the depth of the Cd˚ [30]. The CdTe surface after chemical etching depleted layer is 50–100 A in aqueous solutions of K2Cr2O7-H2SO4 was investigated [33]. It was established that by using solutions with a high K2Cr2O7 content, TeO2 is formed on the surface and on increasing of the solution acidity, this oxide is absent. The depth of the Cd-depleted layer depends on the K2Cr2O7 content and the etching time, and can reach up to 800 nm. The kinetics and mechanism of CdTe etching in solutions of K2Cr2O7H2SO4-H2O system was investigated [78]. It was noted that a limiting stage of the dissolution process is the tellurium layer dissolution, which is formed as a result of the CdTe chemical interaction with the etchant. The CdTe etching rate increases with the increasing of the K2Cr2O7 content in the etchant. Investigations [79] indicate a high adsorption of Cr2O72
ions. The maximum quantity of chromium (1.51017 cm
2) is adsorbed on the cleavage plane; the minimum Cr quantity (3.51016 cm
2) was determined on surfaces after removal of a disturbed mechanical treatment layer by chemical pretreatment. It was shown that after removing the layer with a thickness of 50–100 mkm (6–9)  1015 Cr atoms were fixed on 1 cm2. The presence of a high chromium levels after chemical etching can be explained by the fast Cr diffusion into the ingot.

Chemical Treatment of the CdTe and ZnxCd1
xTe Surfaces

127

5. ETCHANT COMPOSITIONS BASED ON H2O2 Hydrogen peroxide as a strong oxidant is often used for CdTe chemical etching in the mixture with HF as a strong complexing agent. A film with ˚ enriched by tellurium ([Te]/[Cd] ¼ 5) is formed a thickness of 600-800 A on the CdTe surface as a result of the chemical treatment in the solution containing 2 p. H2O2 þ 3 p. HF þ 1 p. H2O [30]. The etchant with the composition of 2 p. H2O2 þ 3 p. HF þ 2 p. H2O can be used for the identification of (111) CdTe polar planes and for the dislocation detection on the (111) Cd plane [19, 61]. The solution containing 0.5 p. H2O2 þ 4 p. HF þ 2 p. H2O is used for the detection of structural defects on the (111) Te [59] plane. The growing steps on the (110) plane of the Cd0,96Zn0,04Te single crystals were detected by chemical etching in the solution containing 200 ml HF þ 400 ml H2O2 þ 400 ml H2O [80]. According to the data of ellipsometry, the surface layers forming on CdTe by chemical treatment in solutions containing H2O2 are not a mixture of CdO and TeO2 oxides, but a chemical compound, probably CdTeO3 or Cd2TeO5 [28]. The solubility of the oxidation products strongly depends on the acidity of the solution. Therefore, it is possible to govern the layer composition by changing the pH of the etchants. The crystallographic orientation (111) of the n-CdTe surface can be determined using the etchant with a composition of 3 p. HF þ 2 p. H2O2 þ 1 p. H2O þ aqueous solution of 1 mg AgNO3 in 100 ml of H2O [81]. According to the data of SIMS, the CdTe surface treated in a hot solution containing 4 p. H2O2 and 1 p. NH4OH contains some quantities of TeO, TeO2, and TeO3, which indicates its strong oxidation nature [31]. ˚ and its refractory coefficient is The layer depth can reach up to 700–800 A equal to approximately 2, which is very close to that of TeO2. The next oxidizing step does not increase the layer depth, but increases the refractory coefficient up to 2.2. The etchant, which is composed of 1 p. HF, 1 p. H2O2, and 1 p. CH3COOH, is used for the detection of twinning, inclusion, and dislocation on the CdTe surfaces [22].

6. HALOGEN-EVOLVING ETCHANT COMPOSITIONS During the last two decades in the V. Ye. Lashkaryov Institute of Semiconductor Physics, a scheme for investigating the chemical etching of semiconductor materials was developed. This scheme includes a mathematic planning experiment using simplex methods, physicochemical simulation by minimizing the system free energy, kinetic investigations, and determination of the composition and the structure of the forming surface layers [5, 6]. Such a scheme has been successfully used for the

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elaboration of the active media with high etching rates for chemical cutting, and with medium and slow etching rates for thinning and polishing of the II-VI semiconductor materials. Bromine- and iodine-containing etchants are often used as the oxidants. Therefore, it is practical to use as these etchant compositions for the chemical treatment of semiconductor materials in liquid active media, in which the halogens are formed as a result of the chemical interaction of initial components of the etchant composition [6]. For example, if HNO3, H2O2, and K2Cr2O7 are oxidants and HBr is a halogen-containing compound, bromine can be formed as a result of chemical reaction: HNO3 þ 3HBr ¼ Br2 þ NOBr þ 2H2 O H2 O2 þ 2HBr ¼ Br2 þ 2H2 O K2 Cr2 O7 þ 14HBr ¼ 3Br2 þ 2CrBr3 þ 2KBr þ 7H2 O

ð1Þ ð2Þ ð3Þ

It is necessary to note that hydrogen peroxide has the biggest oxidizing potential and the lowest ionization constant among the used oxidizing agents [2]. H2O2 exhibits weak acidic properties in aqueous solutions and in combination with hydrohalogenic acids can generate etchants for CdTe and based on it solid solutions with small etching rates and good polishing properties. The use of nitric acid as an oxidant in bromine-evolving etching mixtures based on the HNO3-HBr-solvent systems yields etchants with high dissolution rates. The substitution of H2O2 for oxidant in Br2evolving etchants usually abruptly decreases the etching rate. Chlorine and iodine are evolved as a result of the chemical interaction of the above-mentioned oxidants with HCl or HI. As an additional solvent, some organic acids, ethylene glycol, or dimethylformamide are used and the evolving halogens are partly dissolved in them. Moreover, such additional solvents can regulate the halogen generated, decrease the etching rate, and facilitate the dissolution of the forming surface reaction products. One of the main characteristics of the etchant is the etching rate. The concentration dependence of etching rates of the semiconductor compounds using three-component etchants “oxidant-HHal-solvent” can be presented as surfaces of equal etching rates. Graphical construction of such dependencies gives a possibility of determining the concentration regions of polishing, unpolishing, and selective etchants in each investigated system. Using such diagrams, it is possible to choose in the given system a polishing etchant with necessary polishing rate or to conclude that the polishing etchant with necessary properties does not exist in this system. In the latter case, it is necessary to change one of the etchant components and to investigate such new ternary system to obtain an etchant with necessary properties.

Chemical Treatment of the CdTe and ZnxCd1
xTe Surfaces

129

However, at the chemical treatment of semiconductor surfaces, it is necessary to know not only the concentration dependence of the etching rate but also a roughness of the obtained surface, the surface contamination by the etchant components and the chemical interaction products, the deviation from stoichiometry in the surface layers, and some other characteristics. Therefore, concentration dependencies for all these parameters must be constructed in the form of such diagrams. One can choose the best etchant for a given semiconductor if we construct diagrams and compare them with each other. These diagrams may also provide the definition of the mechanism of semiconductor interaction with etchant by chemical etching. One can conclude that the dissolution of cadmium telluride and solid solutions based in it, in solutions of HNO3-HCl-citric acid is limited by the interaction of tellurium, which is formed on the surface with etchant components, if we compare the corresponding surfaces of equal etching rates (Fig. 2). Such conclusion is based on the similarity of the obtained surfaces of equal etching rate for CdTe, Zn0.03Cd0.97Te, and Cd0.2Hg0.8Te solid solutions and Te in the investigated etchant compositions [87]. The diagrams “etching rate of semiconductor—etchant composition” of the system containing H2O2-HBr(HI)-organic acid, in which the bromine- and iodine-releasing etchant compositions are formed, and H2O2-HNO3(HCl)-organic acid were constructed by us with the help of mathematical planning of the experiment. Such diagrams help in the comparison of different etchant compositions on their etching rates and the selection of the best etchant for the given semiconductor. We constructed more than 100 such diagrams for the investigation of chemical etching of undoped and doped II-VI semiconductor compounds and its solid solutions in different liquid solutions. Some of them are given in Table 1. One can note that by using surfaces of equal etching rates for interaction of solid solutions with the given etchant, it is possible to define the influence of solid solution composition on the mechanism of its dissolution.

7. INFLUENCE OF DOPING ON CHEMICAL ETCHING It is necessary to emphasize that doping of semiconductors leads to changing of not only the etching rate but also the range of polishing solutions in each investigated system. Therefore, at chemical treatment of semiconductors, it is necessary to establish the nature of the doping element and to develop the polishing etchant compositions for different dopants, which can be used for the modification of semiconducting properties of a given semiconductor.

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C

C 1 5

19

2327 31

9

17 13

13 17 21

9

15 13

15

5 11

15

17

1

17

7

21 1

A

A

20

40

60

80

B

C

3

A

20

40

vol.%

60

80

B

vol.%

C

C

3 5 7 9 11 15

23 19 4 19 16 13 10 7

15 7 1319

4

7

1

B

A

20

40

60 vol.%

11 9 7

80

B

D

5

3

1

9

A

20

40

60

80

B

vol.%

Figure 2 Surfaces of equal etching rates (mkm/min) of (A) Te, (B) Cd0.2Hg0.8Te, (C) Zn0.03Cd0.97Te, and (D) CdTe in HNO3-HCl-citric acid solutions. Ratio of HNO3:HCl: citric acid in A, B, and C (vol%): A, 10:90:0; B, 20:20:60; C, 90:10:0.

It was determined that the etching rate of undoped CdTe in all cases is practically lower than that of doped cadmium telluride with Ga, Ge, Sb, Sn, and As þ Cl [97, 99, 104, 107, 117, 118]. The field of the polishing solutions practically does not change. However in some cases, inverse effect takes place, when depending on the etchant composition, the etching rate of doped CdTe can both insignificantly increase or decrease [99, 100] The dissolution rate versus concentration curves for CdTe samples, undoped or doped with Ga, Ge, Sb, Sn, and As þ Cl, are shown in Fig. 3 [117]. These curves are alike for all materials under investigation. The following tendency is observed: an increase in the H2O2 content in the mixture from 2 to 10 vol% leads to a systematic increase in the dissolution rate for all materials. For undoped CdTe, the rate increases from 4 to 13 mkm/min. Doped CdTe dissolves somewhat rapidly in the same solutions: the highest etching rate reaches 18 mkm/min. It is noteworthy

Chemical Treatment of the CdTe and ZnxCd1
xTe Surfaces

Table 1 Halogen evolving etchant compositions for the chemical polishing of the CdTe and ZnxCd1
xTe surface Semiconductor Etchant composition

CdTe

ZnxCd1
xTe

HNO3-HCl-H2O HNO3-HCl-citric acid HNO3-HCl-tartaric acid HNO3-HCl-lactic acid HNO3-HBr-H2O HNO3-HBr-tartaric acid HNO3-HI-H2O HNO3-HI-tartaric acid K2Cr2O7-HCl-H2O H2O2-HCl H2O2-HCl-tartaric acid H2O2-HBr-H2O H2O2-HBr-lactic acid H2O2-HBr-tartaric acid H2O2-HBr-citric acid H2O2-HBr-ethylene glycol H2O2-HI (2.5–15 vol% H2O2) H2O2-HI-citric acid H2O2-HI-(citric acid/ ethylene glycol ¼ 1:1) H2O2-HI-ethylene glycol H2O2-HI-tartaric acid H2O2-HI-lactic acid H2O2-HI-oxalic acid HNO3-HCl-acetic acid HNO3-HCl-citric acid HNO3-HCl-(citric acid/ ethylene glycol ¼ 1:1) HNO3-HCl-tartaric acid HNO3-HCl-lactic acid HNO3-HI-tartaric acid H2O2-HBr-H2O H2O2-HBr-lactic acid H2O2-HBr-tartaric acid H2O2-HBr-citric acid H2O2-HBr-ethylene glycol H2O2-HI (2.5–15 vol% H2O2) H2O2-HI-tartaric acid H2O2-HI-oxalic acid

Etching rate (mkm/min)

References

1–14 1–37 1–33 2–98 5–150 10–178 0.5–6.5 0.1–15 3–7.5 up to 2.3 1–3.5. 1–18 1–24 1.5–19 5–19 1–23 10–23 6.5–24 2.8–23.5

[83–86] [87] [88, 89] [90] [91, 92] [89] [71] [93] [74] [94] [95] [96] [97, 98] [99–101] [102, 103] [104–106] [107, 108] [15, 109] [110]

4.2–33 8–22 2–23 7–22 2–14 1–33 1–17

[111] [112, 113] [107] [114] [115] [87, 115, 116] [115]

1–16 2–82 1–14 1–22 2–20 3–21 6–22 2–22 9.5–23 6–22 6.8–23

[88, 115] [90, 115] [93] [96] [98] [101] [103] [105, 106] [108] [112] [114]

131

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V.N. Tomashik and Z.F. Tomashik

v, μm/min 16

1

12

2

8 4 0 0

10 20 30 40 50

16

1

12

3

8 4 0 0

10 20 30 40 50

16 1

12 4 8 4 0 0 10 20 30 40 50 16 1 12 5 8 4 0 0

10 20 30 40 50

16 1 12 6

8 4 0 0

10 20 30 40 50 H2O2, vol %

Figure 3 Etching rates (mkm/min) versus concentration plots for (1) CdTe, (2) CdTe:Ga, (3) CdTe:Ge, (4) CdTe:Sn, (5) CdTe:Sb, and (6) CdTe:(As þ Cl).

Chemical Treatment of the CdTe and ZnxCd1
xTe Surfaces

133

that in almost all the solutions the dissolution rate of undoped CdTe is somewhat lower than the dissolution rate of doped CdTe. It is found that the etching rate of the materials under investigation in the solutions containing 10 vol% H2O2 increases in series CdTe ! CdTehSbi ! CdTehSni ! CdTehGai ! CdTehAs þ Cli ! CdTehGei. The treatment of all investigated samples in solutions containing 2–10 vol% H2O2 results in a mirror surface. Further increase in the H2O2 content in the etching mixture from 10 to 50 vol% leads to a decrease in the etching rate to 0.5–1.5 mkm/min for undoped CdTe and to 1.8–2.5 mkm/min for doped samples. The quality of the treated surface depends on the H2O2 content in the etchant.

8. INFLUENCE OF CRYSTALLOGRAPHIC ORIENTATION ON CHEMICAL ETCHING Behavior of the single crystal surface at chemical etching depends on its crystallographic orientation [3]. As a rule, by using bromine-methanol solutions the etching rate of the anionic surface is higher than that of the cationic surface; the difference between the etching rates of A and B decreases with a decreasing of the bond ionicity [119]. The same behavior in increase of etching rates takes place at CdTe dissolution in iodine-methanol etchants with various I2 content (Fig. 4): the plane (110) is characterized by the smallest etching rate and the etching rate of the (100) plane is the largest. The cationic plane (111) A is dissolved at a slower rate than the anionic plane (111) B [120]. When I2 content decreases, a line slope decreases also and on using diluted iodine 4

v, mkm/min

3

6 5

2

4 3

1

2 1

0 (110)

(111)A

(111)B

(100)

Surface orientation

Figure 4 Dependence of the CdTe etching rate on the surface orientation in solutions containing (1) 3, (2) 5, (3) 8, (4) 10, (5) 12, and (6) 14 mass% I2 in methanol.

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V.N. Tomashik and Z.F. Tomashik

v, mkm/min 16

7

14 12

6

10 8

5

6

4

4 2 0

3 2 1 (111)B

(110)

(100)

Surface orientation

Figure 5 CdTe dissolution rate versus surface orientation for dimethylformamide solutions containing (1) 2, (2) 5, (3) 10, (4) 12, (5) 15, (6) 20, and (7) 25 mass% I2.

solutions, the difference between the etching rates of different faces becomes less appreciable. In dilute solutions I2-dimethylformamide containing 2–5 mass% I2, the dissolution rate decreases in the order of (100) > (110) > (111) B > unoriented surface (Fig. 5) [118]. As the oxidizer concentration in the etchant is increased, this order changes to (111) B > (110) > unoriented surface > (100) at [I2] ¼ 20–25 mass%. Therefore, the dissolution rate of CdTe crystals depends not only on the crystallographic orientation of the dissolving surface but also on the concentration of the active component in the etching solution. Note that the etching of various CdTe planes with I2-dimethylformamide solutions is rather slow and yields a high-quality polish. The best polish is achieved with dimethylformamide containing 15 mass% I2. Addition of citric acid to the etchant solution brings about a reduction of the etching velocity because of the low value of its ionization constant and the enhancement of etchant viscosity. In addition, citric acid promotes more effective removal of reaction products from the CdTe surface. The etching rate of CdTe in aqueous solutions of the H2O2-HI-citric acid system was observed to depend on crystal orientation [15]. The highest etching rate was mostly observed for the (111) B orientation and the lowest rate for the (111) A one. The next dependence of etching rates of the oriented planes V(111)B > V(110) > V(111)A > V(100) was obtained at CdTe chemical etching in solutions containing H2O2-HI-(citric acid/ethylene glycol ¼ 1:1) [110]. Given consequence is not rigorous and in some cases can change with the changing of etchant composition and the condition of the experiment

Chemical Treatment of the CdTe and ZnxCd1
xTe Surfaces

+20.00

135

nm

Height

+13.75 +7.50 +1.25 –5.00

A

0 +6.00

50

100 Distance

150

mkm

50

100

150

mkm

nm

Height

+3.00 0 –3.00 –6.00

B

0

Distance

Figure 6 Profile lines of CdTe single crystal surface after its treatment using two solutions of H2O2-HI-(citric acid/ethylene glycol ¼ 1:1): (A) 10 vol% H2O2 þ 62 vol% HI þ 28 vol% (citric acid/ethylene glycol) (B) and 5 vol% H2O2 þ 40 vol% HI þ 55 vol% (citric acid/ethylene glycol) (B).

being carried out [110, 111]. The crystallographic orientation influences the concentration regions of the polishing, selective, and unpolishing solutions in this system. The profile lines of the CdTe single crystal surface after its treatment using two solutions of H2O2-HI-(citric acid/ ethylene glycol ¼ 1:1) are shown in Fig. 6. One can see that the surface roughness does not exceed 40 nm at chemical etching by mixtures containing high H2O2. Addition of the mixture of citric acid/ethylene glycol to the etchant may lead to surfaces with low roughness.

9. CHEMICAL ETCHING OF ZnxCd1–xTe SOLID SOLUTIONS The formation of solid solutions based on given semiconductors leads to the modification both the etching rate and the range of polishing composition in each investigated system, but these modifications are small in comparison to the dopant influence. This can be explained by the fact that solid solutions are formed by chemical elements with similar chemical and physical properties and different chemical elements are used for semiconductor doping.

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V.N. Tomashik and Z.F. Tomashik

v, mkm/min

6 1

5

2 4 3 3 0

2

4 6 mol. % ZnTe

8

10

12

Figure 7 Dependence of the polishing rate of the ZnxCd1–xTe solid solutions single crystals in the solutions, containing HNO3, HCl and citric acid (in vol. %), respectively: 32.5; 52.5 and 15 (1); 33.75; 43.25 and 22.5 (2); 36.25; 26.25 and 37.5 (3).

The dependencies of the etching rate for ZnxCd1–xTe solid solutions single crystals in polishing solutions of HNO3-HCl-citric acid from ZnTe content are shown in Fig. 7 [116]. It can be seen that the etching rate increases with the increase of x. Using the obtained values, it is possible to calculate the etching rates of the solid solutions for a given x, and these values help to choose the etchant for chemical polishing of the solid solutions at various ZnTe contents. The minimum dissolution rate of CdTe and ZnxCd1–xTe in the etchant compositions H2O2-HI-tartaric acid and the solution regions with minimum dissolution rate shifts gradually to the solutions enriched by tartaric acid at the transition from CdTe to Zn0.04Cd0.96Te and Zn0.2Cd0.8Te solid solutions [112]. It is interesting that the region of polishing solutions in H2O2-HI-tartaric acid changes at transition according to the abovementioned semiconductor materials. Increasing of Zn content in the ZnxCd1–xTe solid solutions leads to an increase of the etching rate [103, 105]: the etching rate of the Zn0.04Cd0.96Te solid solution in H2O2-HBr-ethylene glycol etchant compositions changes in the range of 2–18 mkm/min; for the Zn0.2Cd0.8Te solid solution, it is equal to 4–22 mkm/min [105]. The concentration dependence of the CdTe, Zn0.04Cd0.96Te, and Zn0.2Cd0.8Te etching rates in aqueous solution H2O2-HBr is given in Fig. 8 [121]. It is shown that for all above-mentioned materials, the maximum value is obtained by using the mixture containing 10 vol% H2O2 in HBr. This can apparently be explained by the fact that at this component ratio, maximum quantity of bromine evolves, which dissolves in excess of hydrobromic acid. It is necessary to note the CdTe etching rate is lower than that for ZnxCd1–xTe solid solutions.

Chemical Treatment of the CdTe and ZnxCd1
xTe Surfaces

137

24 22 20 18

v, mkm/min

16 14 12 10

8 6 1 3 2

4 2 0 0

10

20

30 40 H2O2, vol. %

50

60

Figure 8 Concentration dependences of the CdTe (1), Zn0.04Cd0.96Te (2), and Zn0.2Cd0.80Te (2) etching rates in aqueous solutions of the H2O2-HBr system (T ¼ 291 K, g ¼ 86 rpm).

10. NANODIMENSIONAL FORMATION ON CdTe AND Zn1–xCdxTe SURFACES AT CHEMICAL ETCHING As a result of experimental studies of CdTe and ZnxCd1
xTe etching in solutions of H2O2-HBr-lactic acid (ethylene glycol) systems, it has been determined [122] that the microstructure of the treated surfaces is characterized by high quality and high luster. The roughness (Rz) does not exceed 0.05 mkm, thus meeting the requirements of surface preparation in the manufacturing of various semiconductor devices. However, a more detailed study of the formed polished surfaces has shown that in most cases the polished surfaces are not perfectly smooth. Figure 9 depicts the surface microstructure of Zn0.04Cd0.96Te solid solutions after two-stage treatment: chemical-mechanical and chemicaldynamic polishing with H2O2-HBr-ethylene glycol polishing etchant compositions (a, b, c); as well as the similar pictures for nominally undoped CdTe (100) polished by H2O2-HBr-lactic acid polishing etchant compositions (d, e); and for Zn0.2Cd0.8Te solid solution (f) after corresponding interstage rinsing [122]. It is evident that the obtained CdTe, Zn0.04Cd0.96Te, and Zn0.2Cd0.8Te surfaces are characterized by a

nm

nm –433.06 0.132

mm

0.000

mm

mm

0.000 0.352

D

0.000

mm

+538.68

+568.71

nm

nm

–206.93 0.264

–350.00 0.264

mm

B

0.000

mm

mm

0.000 0.352

E

0.000

mm

0.000

mm

0.000 0.176

0.000 0.352

+496.14

+811.67

nm

nm

–263.33 0.132

–72.70 0.132

mm

C

0.000 0.176

mm

F

0.000

mm

0.000 0.176

Figure 9 Surface microstructure after treatment with polishing etchant compositions: H2O2-HBr-ethylene glycol for Zn0.04Cd0.96Te (A–C) and H2O2-HBr-lactic acid for CdTe (D, E) and Zn0.2Cd0.80Te (F).

V.N. Tomashik and Z.F. Tomashik

–170.63 0.264

A

138

+755.76

+627.72

Chemical Treatment of the CdTe and ZnxCd1
xTe Surfaces

139

high luster and high polishing quality. However, aperiodic nano-sized, needle-shaped formations are revealed on polishing surfaces, which appeared in the course of chemical polishing. The size of these nanoneedles varies between 10 and 80 nm, depending on the semiconductor material, chemical treatment methods, and the polishing etching composition. It follows from the surface roughness that the arithmetic average profile deviation Ra for undoped CdTe and ZnxCd1
xTe solid solutions single crystals after the chemical two-stage treatment, including chemicalmechanical and chemical-dynamic polishing with H2O2-HBr-ethylene glycol and H2O2-HBr-lactic acid polishing etchant compositions, varies within the limits of 11-30 nm and the irregularity profile height is of 8–50 nm. The phase and elementary compositions of such needle-shaped formations have not been studied. However, it is possible to assume that these formations must be enriched with tellurium since CdTe dissolution in acid solvents is limited as a rule by the dissolution of tellurium sublattice [84]. Obviously, such aperiodic, nanosized, needle-shaped formations could appear also in the course of chemical surface polishing of other semiconductors. This should be taken into account for chemical polishing, since their presence could influence the physical properties of the formed metal/semiconductor contacts and for the manufacturing of device working elements.

11. CONCLUSION Having vast experience in the field of chemical treatment of semiconductor compounds, we can summarize some problems that need to be resolved in the future [82]. The first problem is the influence of chemical polishing on the electrophysical properties of semiconductors. Our experimental results aimed at the determination of a potential barrier height (fb) at the chemical polishing of CdTe surface by solutions of HNO3-HCl-citric acid and HNO3-HBrtartaric acid systems on exposure of Au-p-CdTe structures from both the Au and CdTe sides have been indicated that the potential barrier height can be change from 0.2 to 0.8 eV depending on the etchant composition used. Some of the other problems are summarized as follows:  Determination of doping influence by different impurities on chemical

etching of semiconductors;

 Creation of a technological pattern of the total cycle of semiconductor

substrate preparation in etchants, containing the same components but in different proportions and having a broad spectrum of etching rates and introducing a minimum quantity of surface pollution after each manufacturing operation;

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 Development of etchants with too small etching rates of semiconductor

 





materials and high quality of polished surfaces to treat thin crystals, epitaxial layers, and films; Development of etchants with high etching rates for chemical cutting of semiconductors; Development of etchants for new semiconductor materials to introduce them in the industry (usually etchants developed for the treatment of other materials are used for chemical etching of new semiconductors); Development of etchant and rinse compositions that form special layers on surfaces (e.g., passivation) and could have practical significance in the manufacturing of working elements; Usage of less toxic, cheaper, and practically feasible components in etchant compositions.

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[93] Z.F. Tomashik, Y.e.O. Bilevych, V.N. Tomashik, Optoelektronika i Poluprovodn. Tekhnika (in Russian), Vyp. 41 (2006) 108–111. [94] V.N. Tomashik, E.M. Lukiyanchuk, Z.F. Tomashik, Optoelektronika i Poluprovodn. Tekhnika (in Russian), Vyp. 38 (2003) 204–209. [95] Z.F. Tomashik, E.M. Lukiyanchuk, V.N. Tomashik, Semicond. Phys., Quantum Electron. Optoelectron. 7 (2004) 452–455. [96] I.I. Hnativ, Z.F. Tomashik, V.N. Tomashik, I.B. Stratiychuk, Fiz. i Khim. Tverd. Tila (in Ukranian) 9 (2008) 357–362. [97] I.B. Stratiychuk, Z.F. Tomashik, P.I. Feychuk, V.N. Tomashik, Kondensir. Sredy i Mezhfaz. Granitsy (in Russian) 7 (2005) 154–160. [98] Z.F. Tomashik, V.N. Tomashik, I.I. Hnativ, I.B. Stratiychuk, Neorgan. Materialy (in Russian) 42 (2006) 949–953. [99] I.B. Stratiychuk, V.N. Tomashik, Z.F. Tomashik, P.I. Feychuk, Novi Tekhnologhii (in Russian) 3 (6) (2004) 29–33. [100] Z.F. Tomashik, I.B. Stratiychuk, V.N. Tomashik, Fiz. i Khim. Tverd. Tila (in Ukrainian) 6 (2005) 99–103. [101] V.N. Tomashik, I.I. Hnativ, Z.F. Tomashik, I.B. Stratiychuk, Vopr. Khim. i Khim. Tekhnol. (in Russian) 5 (2006) 47–51. [102] V.N. Tomashik, I.B. Stratiychuk, Z.F. Tomashik, P.I. Feychuk, Vopr. Khim. i Khim. Tekhnol. (in Russian) 1 (2005) 43–47. [103] Z.F. Tomashik, I.I. Hnativ, V.N. Tomashik, I.B. Stratiychuk, Zhurn. Neorgan. Khimii 52 (2007) 1234–1238. [104] Z.F. Tomashik, I.B. Stratiychuk, V.N. Tomashik, P.I. Feychuk, L.P. Shcherbak, Neorgan. Materialy (in Russian) 41 (2005) 775–781. [105] Z. Tomashik, I. Stratiychuk, V. Tomashik, S. Drygybka, Visnyk L’viv. Un-tu. Ser. Fiz., Vyp. 38 (2) (2005) 441–448. [106] Z.F. Tomashik, I.I. Hnativ, V.N. Tomashik, I.B. Stratiychuk, Zhurn. Neorgan. Khimii 51 (2006) 1406–1409. [107] O.R. Gumenyuk, Z.F. Tomashik, V.N. Tomashik, P.I. Feychuk, Kondensir. Sredy i Mezhfaz. Granitsy (in Russian) 4 (2002) 242–246. [108] O.R. Gumenyuk, Z.F. Tomashik, V.N. Tomashik, Optoelektronika i Poluprovodn. Tekhnika (in Russian), Vyp. 37 (2002) 147–149. [109] Z.F. Tomashik, O.R. Gumenyuk, V.N. Tomashik, P.I. Feychuk, Kondensir. Sredy i Mezhfaz. Granitsy (in Russian) 5 (2003) 248–252. [110] V.G. Ivanitskaya, Z.F. Tomashik, V.N. Tomashik, P.I. Feychuk, P. Moravec, Y.a. Franc, Kondensir. Sredy i Mezhfaz. Granitsy (in Russian) 9 (2007) 47–52. [111] Z.F. Tomashik, V.G. Ivanitskaya, V.N. Tomashik, P.I. Feychuk, Nauk. Visnyk Cherniv. Un-tu, Vyp. 307. Khimiya (2006) 136–141. [112] O.R. Gumenyuk, Z.F. Tomashik, V.N. Tomashik, Fiz. i Khim. Tverd. Tila 4 (2003) 127–132. [113] Z.F. Tomashik, O.R. Gumenyuk, V.N. Tomashik, Proc. SPIE. 5065 (2003) 241–245. [114] Z.F. Tomashik, G.M. Okrepka, V.N. Tomashik, O.R. Gumenyuk, Vopr. Khim. i Khim. Tekhnol. (in Russian) 6 (2007) 55–58. [115] Y.e.O. Bilevych, V.N. Tomashik, Z.F. Tomashik, V.K. Komar, S.G. Danylenko, Optoelektronika i Poluprovodn. Tekhnika (in Russian), Vyp. 36 (2001) 118–124. [116] Z.F. Tomashik, S.G. Danylenko, P. Siffert, V.N. Tomashik, P.I. Feychuk, Y.e. O. Bilevych, Optoelektronika i Poluprovodn. Tekhnika (in Russian), Vyp. 35 (2000) 57–62. [117] I.B. Stratiychuk, Z.F. Tomashik, V.N. Tomashik, P.I. Feychuk, Zhurn. Neorgan. Khimii 49 (2004) 2095–2100. [118] Z.F. Tomashik, V.G. Ivanitskaya, V.N. Tomashik, P.I. Feychuk, L.P. Shcherbak, Zhurn. Neorgan. Khimii 50 (2005) 1765–1768.

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CHAPTER

II Applications

Contents

IIa. Photorefractive CdTe 1. Introduction 2. The Photorefractive Properties 2.1. The photorefractive effect 2.2. Carriers photogeneration quantum yield and wavelength sensitivity range 2.3. Trapping level in the photorefractive crystal 3. Experimental Results 3.1. Photorefractive properties of CdTe:V 3.2. Effect of Zn content on photorefractive properties 4. Applications 4.1. DPCM experiments 4.2. DPCM between single-mode fibres 5. Conclusion References

148 148 149 149

IIb. Cadmium Telluride-Based Solar Cells 1. Introduction 2. State of the Art 3. Device Properties 3.1. Back contact effects 3.2. Inter-diffusion at the CdTe/CdS junction 4. Fabrication of Cells 4.1. Deposition of the absorber CdTe layer 4.2. Deposition of the CdS window layer 4.3. Post-growth annealing in chlorine 4.4. Back contacts 4.5. Transparent conducting oxide front contacts 4.6. Alternative structures 5. Manufacture of CdS/CdTe Modules 6. Degradation Mechanisms 7. Use of CdS/CdTe Modules in Large-Scale Power Generation

187 187 189 189 189 193 196 196 197 197 199 201 202 206 207

CDTE and Related Compounds

#

151 152 153 153 169 178 178 182 183 185

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2010 Published by Elsevier Ltd.

DOI: 10.1016/B978-0-08-096513-0.00002-9

145

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Applications

8. Concluding Remarks References

210 211

IIc. Applications of CdTe, CdZnTe, and CdMnTe Radiation Detectors 1. Introduction 2. National Security and Nonproliferation Inspections 3. Medical Imaging 3.1. Gamma (g)-camera 3.2. Digital mammography 3.3. X-ray computed tomography (CT) 4. Space and Astrophysics 5. Nature and Development of CMT Detectors 6. Summary and Future Work Acknowledgments References

214 214 215 217 217 222 223 228 233 236 237 237

IId. Electro-optic Modulator Applications 1. Introduction 2. Practical Configurations 3. Issues and Limitations 3.1. Mechanical 3.2. Optical 3.3. Electrical 4. Successful and Contemplated Deployments 4.1. Laser Q-Switching 4.2. Laser Cavity Dumping 4.3. Optical Pulse Shaping 4.4. Free-Space Optical Communications 4.5. Optical Frequency and Phase Modulators 4.6. Laser Intracavity Modulation and Mode Locking References

239 239 240 241 241 242 243 245 245 246 246 246 247 247 247

IIe. Optical Detectors Based on CdTe Pure Crystals for High-Efficiency Optical Computers 1. Introduction 1.1. Optical computers based on semiconductor structures 1.2. Optical registering media in contemporary optical processors on mis structures 1.3. Fast optical registering media on semiconductor M(TI)S-nanostructures 2. Processors for Digital Optical Computers Based on n-p(TI)M Nanostructures of CdTe

248 248 248 249 250 250

Applications

3. Processors for Analog Optical Computers of Incoherent Light on n-p(TI)M Nanostructures on CdTe 4. Optoelectronic Image Correlator of Incoherent Light Based on Analog Optical Processors 5. Conclusion References

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252 254 255 255

CHAPTER

IIA Photorefractive CdTe J.-Y. Moisan

1. INTRODUCTION For optical telecommunication networks, optical switching systems have been studied, and some systems using integrated optics have been proposed, but a spatial holographic interconnect is also an attractive solution for switching of high bite rate channels. Holographic gratings can be used to steer the optical beams, emerging from an input matrix of single-mode optical fibres to an output matrix of single-mode optical fibres. Two characteristics have to be fulfilled in such a system: it must be active at the telecommunication signal wavelength, that is, 1.3 and 1.5 mm, and must be managed as large a number of channels as possible. Photothermoplastic devices have been proposed [1] and, in such an optical configuration, two recording beams are used in the visible range (their wavelength depending on the sensitivity of the recording material) and their reading beams, at 1.3 or 1.5 mm, are deflected by the recorded gratings. In this case, the photothermoplastic device is not sensitive to the signal wavelength. With photorefractive materials, it is possible to imagine an optical system where the signal beam is active itself. Thus, in two-wave mixing (TWM) experiments, which are commonly used to estimate the properties of photorefractive crystals, one can consider that the studied material is active to the wavelength used. In an optical switching system used in an optical network, it is essential that the photorefractive crystals are sensitive to the communication wavelengths; this is the first requirement. The second requirement, concerning aberrations in optical configurations, is that the single-mode fibres could be used as an input and output signal source, with little loss.

4 A Route Crech Argant, F-22730 Tregastel, France

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First, the photorefractive effect will be presented and the properties discussed. Next results, obtained with CdTe materials, will be given and discussed. And finally, an optical configuration will be presented and the first results of a beam-steering system presented.

2. THE PHOTOREFRACTIVE PROPERTIES 2.1. The photorefractive effect The photorefractive effect, which can be described as the combination of two properties, namely, photoconductivity and electro-optic effect, is explained in Fig. 1. Two interfering beams are focused on the crystal; so carriers are photogenerated in the illuminated zones of the material (Fig. 1A). These carriers diffuse into the non-illuminated zones where they are trapped (Fig. 1B), and a non-homogeneous density of carriers y –

–

–

–

I(z)

z

Λ z

A rsc(z) +

B

+

+ –

–

z

+ –

–

s⬘(z) ~ Esc(z) ~ ∫Psc(z)dz z

C fg Δn(z) ~ –Esc(z)

D

z

Figure 1 (A) Illuminating energy, following the interference pattern on the crystal, and generation of carriers. (B) Inside density of carriers? (C) ESC, space charge field built up by the carrier density. (D) Index modulation created in the crystal and phase shifted from the interference pattern.

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reproduces the interference pattern. Because of the modulated density of trapped carriers, a space charge field occurs (Fig. 1C), which is phase shifted by p/2 from the interference pattern. When the electro-optic coefficient of the crystal is no longer negligible, a modulation of the refractive index can be observed (Fig. 1D). So, a phase grating shifted from the recording interference pattern can be read from the recording beams themselves or from a reading beam at a wavelength not active for the material. The basic equation that describes the index modulation Dn in the crystal is Dn0 ¼

1 3 n r41 ESC 2

ð1Þ

where n0 is the refractive index, r41 the electro-optic coefficient depending on the crystal orientation and ESC the space charge field. The most commonly used optical experiment for studying the photorefractive behaviour of materials is TWM, and the configuration for this is shown in Fig. 2. Two beams interfere on the crystal: one – the pump beam – with a higher energy than the other – the signal beam. Because of the grating recorded inside the crystal, the pump beam ‘sees’ this index grating and part of its energy is deflected in the signal beam direction; if this deflected energy is higher than the losses due to the signal beam absorption and the deflected energy of the signal beam, an amplification gain can be measured with the detector. Note that, in this case, the two beams are active and are used as recording and reading beams. From Eq. (1), it is clear that n0 and r41 are the intrinsic properties of the crystal and depend only on the chosen material and for a given wavelength, and that ESC depends on the experimental parameters (and also on the crystal, briefly listed on Table 1). Mirror

Pump beam

Signal beam

Detector PRC

Mirror BS

Figure 2 Two-beam coupling configuration: PRC, photorefractive crystal; BS, beam splitter.

151

Photorefractive CdTe

Table 1 Figures of merit of photorefractive materials where the r values are r41 for the crystals, except for organic materials Material

n0

r (pm V1)

n03r (pm V1)

l (mm)

LiNbO3 BaTiO3 Bi12SiO20 Bi12GeO20 GaAs:EL2 GaAs:Cr InP:Fe CdTe:V CdTe:V Polymer Organic crystal

2.26 2.36 2.54 2.55 3.48 3.5 3.29 2.82 2.82 1.56 1.7

31 1640 5 3.5 1.43 1.2 1.34 5.5 5.5 2.5 24

360 21500 82 58 60 51 48 123 120 9.5 118

0.633 0.546 0.633 0.633 1.06 1.06 1.06 1.06 1.52 .514 .676

2.2. Carriers photogeneration quantum yield and wavelength sensitivity range The absorption of a CdTe:V crystal is reported in Fig. 3. It is clear that the semiconductive crystals can be used only at wavelength corresponding to lower energies than the band gap, where the crystals are sufficiently transparent. In Fig. 3, the 1.3 and 1.5 mm wavelengths are indicated by vertical lines, and the residual absorption, which is said to come from doping agent V, is observed. A similar spectrum has been published [2] for ZnTe:V, and also the photoconductivity studied in this way; the 1.0

0.75

0.5

0.25

0.0 12000

10000

8000

WAVENUMBERS

6000

CM–1

Figure 3 Absorption spectrum of CdTe:V. The ordinate is absorbance. The thickness was 5 mm. Vertical lines indicate 1.3 and 1.5 mm wavelengths.

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J.-Y. Moisan

formation time of the holographic grating was measured by following the total intensity of the two beams. It clearly appears that the formation time obeys the same law as the photocurrent in the material. This means that the photorefractive effect is linearly related to the photogeneration quantum yield. Therefore one can expect from the spectrum in Fig. 3 that CdTe: V will be sensitive to the 1.3 and 1.5 mm wavelengths, which has already been demonstrated [3]. So, two aspects have to be emphasised: the generation quantum yield governs the response time of the material, but the number of trapped carriers will govern the space charge field.

2.3. Trapping level in the photorefractive crystal Because it comes from the trapped carrier density, a large space charge field will be observed in two conditions: when the carriers are trapped in deep traps, and when the number of traps is large. For II–VI crystals, it is well known [4] that doping with V, Ti or Ni can induce deep levels, especially in CdTe and ZnTe. The dopant solubilities are larger in II–VI material than in III–V [5]. This means that the photocurrent will be larger and the response time shorter but, above all, that the number of trapped carriers will be larger and, as a consequence, the space charge field larger. Consider now, in Fig. 1, the diffusion of carriers; when some of them are trapped in the dark zones, the space charge field begins to increase. However, a conduction current in the opposite direction appears, and equilibrium will limit the space charge field when the conduction current becomes equal to the diffusion current. In these conditions, the electric field inside the crystal is ESC ¼

2p kB T l e

ð2Þ

where l is the wavelength used, kB is the Boltzmann constant and T the temperature. Different solutions have been proposed to exceed this limit. 1. A continuous electric field is applied to the crystal [6], but the phase shift between the interference pattern and the electric field modulation must be preserved. So another limit will now appear. ESC ¼

eLNA 2p e e0

ð3Þ

where L is the grating period, NA the number of effective traps and e the dielectric constant of the material. Following the grating period, the effective limit can come from the resistivity of the material, limiting the external applied electric field, especially in semiconductor crystals.

Photorefractive CdTe

153

To avoid too large a dark current, highly resistive materials are needed (108–109 Ocm). 2. Two others solutions have been proposed to exceed the limit in Eq. (3): (a) an applied d.c. field and a moving grating [7]; (b) an applied a.c. electric field, of sinusoidal or square wave nature [8]. However, the same requirement of high resistivity in needed for both these conditions. 3. Note that the resonance mechanisms have also been proposed to increase the gain in TWM experiment: a Franz-Keldysh effect near the band gap [9], and an intensity-dependent resonant behaviour depending on the temperature of the material [10].

3. EXPERIMENTAL RESULTS 3.1. Photorefractive properties of CdTe:V Due to its high electro-optic factor of merit, CdTe:V is very attractive [11, 12]. At 1.32 mm, TWM gains larger than 10 cm1 have been obtained by Ziari [11], with a large amplitude (23 kV/cm) high-frequency (23 kHz) square-shaped electric field at 75 mW/cm2 pump intensity. Similar results have been observed with a CdTe:V sample [13, 14] called DAV31. In order to present and discuss photorefractive properties of CdTe crystals, three crystals (called DAV30, DAV31 and S2105) will be studied. It allows observing the different measurement parameters effect on TWM gain presented previously.

3.1.1. Experimental details The crystals were grown by the Bridgman technique, at CNRS Bellevue for DAV30 and DAV31 and at CNRS Strasbourg for S2105. In the melt 4% and 1% Zn was added respectively, which is expected to improve the mechanical properties of the crystals. The main characteristics are given in Table 2. The crystals were cut in 5  5  5 mm3 cubes with edges aligned along the h110i, h111i and h112i crystallographic directions. In TWM experiments, the polarisation of the two beams lies within the plane of incidence; these beams form a grating along the h111i direction. They are expanded to provide uniform illumination on the sample. Three wavelengths were used: 1.05, 1.32 and 1.54 mm. When necessary the crystal temperature was stabilised, using a Peltier effect device, from 275 to 305 K. The electron–hole competition factor x was measured, without external field, using the experimental procedure described in Ref. [18].

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Table 2

Main characteristics of crystals DAV31 and S2105

[Zn] (%) Starting [V] (cm3) [V] (cm3) r (O cm) a (cm1) (1.048 mm) (1.32 mm) (1.535 mm) x (1.048 mm) (1.32 mm) (1.535 mm) Neff (1015 cm3) (1.048 mm) (1.32 mm) (1.535 mm)

DAV30

DAV31

S2105

4 1019 0.5  1017 (SIMS)a 1.25  1010 1.90 1.23 0.86 0.12 (e)

4 2  1019 1.7  1017 (SIMS)a 1010 4.88 2.24 1.85 0.08 (e) 0.68 (hþ) 0.87 (hþ) 3 1.37 1.74

1 9.4  1018 4.5  1017 (A.A.)b 5  109 4.82 2.94 2.62 0.69 (e) 0.35 (e) 0.05 (hþ)

0.84 (hþ) 0.26 0.84

a

SIMS: secondary ion mass spectrometry. A.A.: atomic absorption.

b

3.1.2. TWM gains without an external electric field

In the electro-optic configuration given above, the TWM gain G is calculated, assuming low-intensity (cw) illumination, one deep-level model and one small grating period L [15] as G¼

Ed Eq 2p 2 xn0 3 r41 Edþ Eq l0 cosy √3

ð4Þ

or, in more succinct, form G ¼ Ax

Ed E q Edþ Eq

ð5Þ

where l0 is the wavelength, y the half-angle of the two beams in the TWM setup, n0 the refractive index, Ed the diffusion field and Eq, the maximum space-charge field. Ed ¼

k B Tk K; e

Eq ¼

e Neff ; eKNeff

Neff ¼

nT0 pT0 nT0þ pT0

ð6Þ

Here K (2p/L) is the grating vector, and Neff is the number of effective traps. The factor x [10] is sp pT0  sn nT0 ð7Þ x¼ 0 ðsp pT0 þ sn 0 nT0 Þð1 þ Id =Io Þ

Photorefractive CdTe

155

where Id ¼

en th nT0  ep th pT0 sn 0 nT0 þ sp 0 pT0

ð8Þ

and I0 is the mean illumination, sp0 and sn0 are the optical emission rates, epth and enth are the thermal emission rates of holes and electrons, respectively, nT0 and pT0 are the number of occupied and empty deep traps, and e is the dielectric constant. Figure 4 shows the TWM gain as a function of the incident intensity for the DAV30 crystal at the 1.32 and 1.535 mm wavelengths. 0.001

0.01

0.1

1

10

0.01 0.1 1 Io, Intensity (mW/cm2)

10

DAV30

Γ (cm–1)

0.2

Wavelength : 1.32 mm Period : 1.72 mm β : 100

0.1

0 DAV30

Γ (cm–1)

0.2 Wavelength : 1.535 mm

0.1

0 0.001

Period : 2.00 mm β : 100

Figure 4 TWM gain as a function of the incident intensity: the crystal is DAV30 sample and the experimental conditions are indicated in the figure. The points denote the experimental results, and the curves represent the fits according to G ¼ G0/(1 þ Id/Io), with G0 ¼ 0.240 cm1, Id ¼ 75 mW/cm2 at the 1.32 mm wavelength and G0 ¼ 0.260 cm1, Id ¼ 155 mW/cm2 at the 1.535 mm wavelength, respectively.

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J.-Y. Moisan

The experimental results are fitted to the equation G ¼ G0/(1 þ Id/Io). Best fits are reported in the figure caption. The grating periods are 1.72 and 2.00 mm, respectively. The Id values are in good agreement with previously published results [3]. A lowest gain (0.06 cm1) is obtained at 1.048 mm. The sign of the dominant photorefractive carriers was found by separate measurements of the signs of the electro-optic coefficient and TWM gain [17]. Results are also given in Table 2. A change in the sign of carriers (from electrons to holes) is observed between 1.048 and 1.32 mm wavelengths. The same phenomenon was reported with another CdTe:V crystal [16]. Figure 5 represents the plot of (K/G) against K2. The electron–hole competition factor x is related [18] to the intercept ordinate axis Or by ð9Þ

x ¼ e=Or kB Tk A

We evaluated x for the three wavelengths (see Table 2). The following two points are to be noted: (1) At 1.32 and 1.535 mm, positive x values with relatively large absolute values are found, as previously published [3]; this absolute value seems to be larger for a higher V content. (2) On the other hand, the negative but small x values at 1.048 mm provide a smaller gain. (Note that the increase of x versus wavelength can be explained by the variation of the electron and hole

17 DAV30 15

K/Γ

13

11 Wavelength : 1.535 mm Ip : 10 mW/cm2

9

7

β : 100

0

10

20

30

K2

Figure 5 Experimental plot of K/G versus K2. K is the grating wave vector, and G is the TWM gain. The deduced value x is 0.72  0.01, and the Neff value is 9.0  1014 cm2.

Photorefractive CdTe

157

photoionisation cross sections sn0 and sp0 of the V-related level. This result is presented by Bre´mond et al. [21], and provides a full explanation.) This behaviour is different from that of the sample of CdTe:V used in ref [16], for which the electron–hole competition was stronger at the 1.32 mm than at the 1.06 mm wavelength. For the S2105 crystal, a change of the sign of the dominating photorefractive carriers is experimented between 1.32 and 1.54 mm (see Table 2). If one assumes that, in this crystal, grown by the same technique as DAV30 and DAV31, holes are still the thermally generated carriers, we must now expect to observe an intensity resonance at 1.32 mm wavelength. The theory [10] predicts that this effect can be observed when E0  Ed and when the pump intensity is comparable with Id. This is verified in Fig. 6 (the optimum pump intensity being 1 mW/cm2 at room temperature). The same resonant effect is still expected at 1.05 mm and is still displayed in Fig. 6. Note that in this case, the optimum intensity is much lower (more than one decade than at 1.32 mm). 4 3.6

TWM Gain (cm–1)

3.2 S2105 2.8 2.4 Ip/Is : 100 2 1.6 1.2 0.8 0.4 0 0.01

0.1 290 K - 1.048 mm

1

10

100

Intensity (mW/cm2)

300 K - ... 310 K - ... 290 K - 1.32 mm 300 K - ... 310 K - ...

Figure 6 TWM gain versus intensity with S2105 crystal at 1.048 and 1.32 mm. The electron–hole resonant effect is observed at 1.048 mm (for less than 0.1 mW/cm2) and 1.32 mm (for about 1–2 mW/cm2).

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J.-Y. Moisan

Using the same plot as shown in Fig. 5 and taking into account the slope P of the best fit, we can evaluate Neff as e ð10Þ Neff ¼ PA x e The Neff values are reported in Table 2. As is generally reported for CdTe: V crystals [3, 16], Neff is much smaller, approximately 1015 cm3, than the effective V content in the material, 5  1015 cm3, measured by secondary-ion mass spectroscopy. The variation of Neff with respect to the wavelength could indicate that the one-level model is not perfectly valid. On the other hand, the experimental determination of Neff is less precise than the one of x. Indeed, with theses samples it is not possible to use grating periods short enough to show a clear maximum of the curve G(L) curves (which is related to Neff), and on the same reason the slope of the K/G against K2 is not well defined, whereas the uncertainty regarding the value of the ordinate at the origin is not so large. The use of samples oriented for contra-directional TWM could help to solve this difficulty.

3.1.3. TWM gain with an external electric filed 3.1.3.1. Continuous electric field The electric field is applied along the

h111i direction, through gold electrodes, on the crystal between crossed polarisers. The high resistivity of the samples allows us to apply electric filed as large as 15 kV/cm. The electric field distribution between the electrodes was visualised by an electro-optic technique in the presence of uniform illumination at 1.32 or 1.54 mm. This distribution appears to be homogeneous for the voltage/interelectrode distance ratio to as much as 10 kV/cm and for light intensities to as much as 10 mW/cm2. Figure 7 shows the dependence of the TWM gain on the applied electric field at different grating periods. The gain value always remains smaller than 1 cm1 at the 1.32 mm wavelength, which is smaller than the absorption coefficient. But a large transient effect is observed, which indicates that much higher gains can probably be obtained by the usual techniques to maintain a p/2 phase shift between the interference pattern and the index grating, such as the moving-grating method or the periodic-field method. We also studied gain versus intensity in the presence of a d.c. field to detect a resonant intensity–temperature effect similar to that observed in S2105 crystal or in InP:Fe [10]. Figure 8 shows that, at 1.32 mm and with a grating period of 12 mm (Ed ¼ 0.135 kV/cm), the gain versus intensity characteristics in the presence of a 2 kV/cm field have the same nonresonant form as the zero-field gain. One therefore can conclude that, because at 1.32 mm the dominant photorefractive carriers are holes, the thermal carriers in these samples are also holes. These factors suggest that a temperature-dependant resonance effect may be seen at 1.048 mm because the dominant photorefractive carriers, as

Photorefractive CdTe

1 Period : 10 mm Period : 4 mm Period : 2 mm

0.8

DAV31

Wavelength : 1.32 mm lo : 5 mW/cm2

Γ (cm–1)

0.6

β : 100 0.4

0.2

0

0

2

4

6

8 10 Eo (kV/cm)

12

14

16

Figure 7 TWM gain in the DAV31 crystal versus a continuous electric field for different grating periods.

0.25

Wavelength : 1.32 mm Period : 12 mm 0.2

Γ (cm–1)

DAV30 0.15

0 kV/cm 0.1

2 kV/cm Fit

0.05 β : 100 0 0.001

0.01

0.1

1

10

lo, Intensity (mW/cm2)

Figure 8 Gain as a function of the incident intensity with the DAV30 crystal. The curve represents the best fit with G0 ¼ 0.180 cm1, Id ¼ 20 mW/cm2, as used in Fig. 4.

159

160

J.-Y. Moisan

3 Wavelength : 1.048 mm DAV31

T : 300 K

Γ (cm–1)

2

1 β : 100 Eo : 5 kV/cm Period : 10 mm 0 0.1

1

10

100

Io, Intensity (mW/cm2)

Figure 9 Gain as a function of the incident intensity with the DAV31 crystal.

indicated in Table 2, are electrons. This is illustrated in Fig. 9, in which a clear resonant behaviour appears with an optimum intensity of approximately 1.5–2 mW/cm2 at 300 K.

3.1.3.2. Periodic square-shaped electric field A square-shaped electric field is applied with a variable frequency F in the low-frequency regime, starting from a quasi-d.c. field regime (F ¼ 1 Hz) to as much as a few hundred hertz. When F ¼ 1 Hz, the amplified signal level at the end of each half of the electric field period is the same as a continuous electric field. In an intermediate regime (to as much as approximately 40 Hz in the conditions given in Fig. 10) the signal intensity exhibits an oscillatory behaviour. In this regime G is defined from the signal amplification reached at the end of each half-period of the electric field. For higher frequencies the transient effect becomes negligible. In Fig. 10, the gain (measured as discussed above) is plotted versus the frequency for different square-shaped electric field amplitudes at a wavelength of 1.32 mm. Gains as high as 10 cm1 are achieved with E0 ¼ 14 kV/cm. An optimum in frequency (slightly decreasing when field amplitude decreases) is clearly seen near 30–100 Hz. At 40 Hz the amplified signal residual oscillation is approximately 5%; whereas this oscillation becomes negligible at 100 Hz. The same behaviour occurs at 1.535 mm, as illustrated in Fig. 11. In this case the gains are lower, for example, 7 cm1 instead of 10 cm1 at 14 kV/cm. A practical difficulty arises when one considers the experimental influence of the frequency of a square-shaped field. Indeed, it is well known [11, 19, 20] that, in the presence of a relatively small departure

Photorefractive CdTe

161

10 14 kV/cm 12

DAV31

10 8

Wavelength : 1.32 mm

8

Period : 10.2 mm Ip : 10 mW/cm2 β : 1000

6

Γ (cm–1)

6

5 4 4 3 2

2 1

0

1

10

100 Frequency (Hz)

1000

Figure 10 Gain versus frequency of the square-shaped alternative electric field (in kilovolts per centimetre), as indicated, with the DAV31 crystal.

from the ideal square-shaped electric field, the gain may be significantly reduced. For a given high-voltage amplifier, the field shape and its frequency are correlated. Thus, one may suspect that the frequency dependence shown in Figs. 10 and 11 could be due to this effect rather than to a more fundamental cause. Therefore the results achieved with two different high-voltage amplifiers (with slew rates of 25 and 105 V/ms, respectively) are compared. The result is shown in Fig. 12 for a 6 kV/cm electric field: there is no significant effect of the shape of the electric field for frequencies lower than 300 Hz. Unless otherwise specified, the results reported here were obtained with the 25V/ms amplifier. It appears that the field shape can be considered to be square to as much as approximately 150 Hz for a 12 kV/cm electric field and to roughly 1 kHz for a 2 kV/cm electric field. These factors clearly mean that the experimental optimum in frequency (generally situated between 10 and 100 Hz) is not due to the slew rate of the electric field. This frequency effect must be explained by the intrinsic properties of the CdTe:V samples.

8

Wavelength : 1.535 mm Ip : 10 mW/cm2 Period : 10 mm β : 1000

14 kV/cm

DAV31

12 kV/cm

6

10 kV/cm

Γ (cm–1)

8 kV/cm

4

6 kV/cm 5 kV/cm 4 kV/cm

2

3 kV/cm 2 kV/cm 1 kV/cm

0

1

10

100

1000

Frequency (Hz)

Figure 11 Gain versus frequency of the square-shaped alternative electric field, with the DAV31 crystal. The conditions are the same as in Fig. 10, except for the wavelength of the incident light. 5 TWM Gain (cm–1) 4

3

6 kV/cm - 1.32 mm

2

Slew-rate : 105 V/ms Slew-rate : 25 V/ms 1

0

5 mW/cm2

1

10

100

1000

Frequency (Hz)

Figure 12 TWM gain versus frequency, with the DAV30 Crystal, with two high-voltage amplifiers with different slew rates (105 and 25 V/ms).

Photorefractive CdTe

163

10 5 mW/cm2

DAV31

10

8

20 2 1

6 Γ (cm–1)

0.5 0.2 0.1

4

2

0

β : 1000 Eo : 10 kV/cm Period : 15 mm Wavelength : 1.32 mm 1

10

100

1000

Frequency (Hz)

Figure 13 The light intensity is indicated for each curve. The decrease of the gain, for 10 and 20 mW/cm2, is probably due to a non-homogeneous pump beam intensity.

Figure 13 presents the G versus frequency curves at 1.535 mm for different light intensities with an external field of 10 kV/cm. In the 1–20 mW/cm2 range the optimum frequency F0 was found to vary as F0 ¼ 13.5 I00.6, where I0 is given in mW/cm2 and F0 is given in hertz. The optimum frequency is also highly sensitive to the grating period, as demonstrated in Fig. 14. Experiments taking into account different values of the pump-to-signal intensity ratio b were conducted. It appears (see Fig. 15) that the optimum frequency decreases when b increases. The gain at optimum frequency increases regularly with b as plotted in the inset of Fig. 15. This behaviour resembles that reported by Ziari et al. [11] under rather different conditions. By making an appropriate choice of the experimental parameters, it is possible to optimise the gain. The highest value (11 cm1) was achieved with the DAV31 sample under the following conditions (see Fig. 16): E0 ¼ 15 kV/cm, F ¼ 40 Hz, L ¼ 15 mm, l0 ¼ 1.32 mm, I ¼ 5 MW/ cm2, b ¼ 1000. Approximately the same value was reported [11], also with a square-shaped field but under quite different conditions (E0 ¼ 23 kV/cm, F ¼ 230 Hz, L ¼ 7.5 mm, l0 ¼ 1.32 mm, I ¼ 75 MW/cm2, and b ¼ 10,000).

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10 DAV31

15.9mm

22.9mm

8.9mm 8

Γ (cm–1)

5.4mm 6

4

Wavelength : 1.32 mm Io : 1 mW/cm2 Eo : 10 kV/cm β : 1000

2

0

1

10

100

1000

Frequency (Hz)

Figure 14 Gain versus frequency of the 10 kV/cm square-shaped alternative electric field, with the DAV31 crystal. The grating period is indicated for each curve.

15 Γ (cm–1)

10

DAV31

Γ (cm–1)

10

Wavelength : 1.32 mm Ip : 10 mW/cm2 Eo : 10 kV/cm Period : 10 mm

10,000

5

0

1

10

100

10,000 β

5000 1000 500

5 100 10

0

1

10

100

1000

Frequency (Hz)

Figure 15 Gain versus frequency of the 10 kV/cm square-shaped alternative electric field, with the DAV31 crystal for different beam ratios b. A plot of the gain for each optimum frequency versus b is shown in the inset.

Photorefractive CdTe

165

15

DAV31 1.32 mm Γ (cm–1)

10

Ip : 5 mW/cm2 Fo : 40 Hz Period : 15 mm β : 1000

5

0

0

4

8

12

16

20

Eo (kV/cm)

Figure 16 Gain versus the square-shaped electric field amplitude E0 obtained with the DAV31 crystal at the 1.32 mm wavelength.

3.1.3.3. Discussion on the frequency dependence of the gain Inadequacy of the standard one-level model In a photorefractive crystal, assuming thermal and optical transitions between one deep centre and both valence and conduction bands, the evolution of the space charge field E1 in the presence of an external d.c. or square-shaped field E0 follows the equation tg

dE1 þ E1 ¼ mESC dt

ð11Þ

where ESC is the saturated space charge field and tg is the grating formation constant. This equation (11) is valid when T and tg are much larger than the other time constants involved in the space charge field formation. In most cases this condition is fulfilled when the carrier lifetime tR is much smaller than tg. The most common method for enhancing the TWM gain in a photorefractive crystal – based on the application of a periodic square-shaped field with a period T – satisfies the conditions tR  T  tg

ð12Þ

In this case, Stepanov and Petrov [22] showed that the space charge field reaches a constant value that does not depend on the period of the external field frequency. The results presented in Section 3.1.3.2 were obtained with T values that are near tg. Therefore, they do not contradict the prediction of Stepanov and Petrov. Moreover, Bylsma et al. [23] noticed, by applying a square-shaped electric field to an InP:Fe crystal,

166

J.-Y. Moisan

that the gain may be slightly improved by operation at a frequency near the inverse grating lifetime; they suggest that some resonance occurs near 1/tg, but they did not give any mathematical evidence of this assumption. However, Mathey et al. [24] observed a very pronounced external field frequency dependence of the gain in some Bi12GeO20 crystals when T is shorter than tR and near the carrier drift time. They showed that one can interpret this result by considering second order terms in the space charge field differential equation, which cannot be neglected under these conditions. However, the operating conditions for the results presented in Section 3.1.3.2 are completely different from these conditions. In the frequency range used, experiments were carried out to know if the electric field built up inside the crystal can be considered to be instantaneous, in contrast to some reported complex dynamic effects observed with InP:Fe crystals larger than 10 kV/cm [25]. Once we remember that the effect of the high-voltage amplifier slew rate cannot explain the results given above, all the above remarks indicate that the space charge field dynamics in these CdTe:V samples must be described by a second order differential equation, even at low external field frequencies. This description should involve a physical mechanism different from that given in [24]. The issues raised in the preceding section lead us to consider a twolevel model to explain, at least qualitatively, the periodic field behaviour of these samples. Different variants of two-level models have been presented to give a more accurate description of the characteristics of different photorefractive materials. For instance, the sub-linear photorefractive response time of BaTiO3 [26] and the temperate dependence of the photorefractive effect in InP:Fe [28] have been clarified with such models. A shallow trap may also have an effect through trap limited mobility [28]. A two-level model In this section, are shown the results obtained for the TWM gain in the presence of a low-frequency square-shaped external field by considering a two-level model (see Fig. 17) of the photorefractive effect with a principal deep level situated near the mid-gap and optically (and thermally) coupled to both the conduction and the valence bands. Because the main goal is to explain the frequency response at 1.32 and 1.54 mm, where the main (optical and thermal) carriers are holes, a secondary level is introduced, situated somewhere between the valence band and the principal level that is expected to have an effect on the hole transport. Furthermore, the optical intensity is assumed to be moderate enough that one can consider the dominant emission process for this level can be the thermal emission of holes. This thermal emission is supposed to be much more important than the emission of the deep level; this second level is supposed to be filled by electrons in the dark.

Photorefractive CdTe

Conduction band

167

μn

σ0n enth

Cn

nT0 – – – – – – – – – – –

pT0 NT0 σ0p

Cp n′T0 p′T0 – – – – – – – – – – – – – – – e′p C′p Valence band

epth N′T0

μp

Figure 17 Energy diagram for the two-level model with the different parameters used for the calculations.

Figure 18 shows, for a given set of material parameters, the calculated curve of a TWM gain as a function of the square-shaped frequency for different values of the second level concentration N0 T. The parameters relative to the principal deep level are sn0 ¼ 2  10–17 cm2, sp0 ¼ 2  1016 cm2, cn ¼ 108 cm3/s, cp ¼ 5  108 cm3/s (where cn and cp are the capture coefficients of electrons and holes, respectively), mn ¼ 1000 cm2/Vs, mp ¼ 80 cm2/Vs (Where mn and mp are the mobilities of electrons and holes,

two wave mixing gain (cm–1) 20

NT′ 0 3×1012

15

1×1013

10

3×1013 5 1×1014 3×1014

0 1

10

103 100 FREQUENCY (Hz)

104

105

Figure 18 TWM gain as a function of the square-shaped field frequency for different values of the second level concentrations N0 T (in cm–3): l0 ¼ 1.32 mm, L ¼ 10 mm, I0 ¼ 1 mW/cm2, E0 ¼ 10 kV/cm.

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respectively), r41 ¼ 5.5 pm/V, er ¼ 10.4 (where er is the relative dielectric constant), nT0 ¼ 1.6  1016 cm3, and pT0 ¼ 8  1015 cm3. For this level the thermal emissions of both types of carriers are neglected. The characteristics of the level are described by two parameters: its hole thermal emission coefficient e0 p (¼ 2000 s1) and its hole capture coefficient c’p (¼ 1.5  104 cm3/s). The parameters of the principal level (supposedly related to the V doping) are not precisely known. The values of the photoionisation cross sections are issued from measurements performed by deep-level optical spectroscopy [29]. All these characteristics start from the d.c. field value and then present some oscillations that correspond to the oscillations of the continuous case grating formation. Then G increases to a maximum that may be more or less wide and finally decreases until reaching a limit value. With negligible second level concentration, the limit value is equal to the maximum, and there is no decrease of gain with frequency (as long as F is much smaller that the inverse carrier lifetime). This result corresponds to the behaviour described by Pauliat et al. [30]. When the second level concentration increases, the limit decreases considerably, whereas the maximum value also decreases, but not to the same extent. The curves in Fig. 18 (particularly the curve corresponding to N0 T ¼ 1013 cm3) look very similar to the experimental curves obtained with the faster highvoltage amplifier (Fig. 12). Note that, with the set of material parameters used in Fig. 18, the d.c. field gain value remains practically unchanged, whereas the periodic field results are strongly dependent on the second level concentration. It has been verified that, when the concentration becomes comparable with that of the principal level, the d.c. field gain becomes significantly reduced, in agreement with the predictions of Rana et al. [27]. The freedom to choose parameters is considerable; however, a quantitative agreement by use of the present parameters (see Fig. 18) is obtained for some important experimental characteristics, as given in Table 3: for optical absorption, the electron hole competition factor and photoconductivity, the agreement is quite good. As shown, the difference between the Table 3 Theoretical DAV31 characteristics deduced from the two-level model parameters, as used in Fig. 17, compared with experimental values as given in Table 2

a x Ip (10 kV/cm) Neff

DAV31 (theoretical values)

1.32 mm to 1 mW/cm2 (experimental values)

1.9 cm1 þ0.67 45 mA/cm2 5.3  1015 cm3

2.24 cm1 þ0.68 55 mA/cm2 1.37  1015 cm3

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theoretical end experimental values of Neff presented in Section 3.1.2 cannot be explained by the introduction of a second level with moderate concentration (see Fig. 18). This difference can be explain if one takes into account two facts: the lack of accuracy of the experimental determination of Neff as discussed previously and the uncertainty of the photoionisation cross section values which results from preliminary measurements.

3.2. Effect of Zn content on photorefractive properties In the previous section, results on photorefractive behaviour of the different CdTe:V crystals has been presented, especially the effects of the different experimental parameters on the TWM gain. In this section, physico-chemical properties of the different crystals will be analysed and taken into account to explain the photorefractive behaviour of CdZnTe:V crystals: two classes of material have been seen and are presented now.

3.2.1. Sample preparation Six crystals have been studied: CTV2, DAV31, DAV35, DAV36, DAV37 (grown in CNRS-Bellevue) and S2105 (grown in CNRS-Strasbourg). The starting content of V in the melt is in the 1019 cm3 range. Three of them (DAV31, DAV35 and DAV37) contain about 4% Zn, S2015 about 1% and the other contain no Zn (see Table 4). Apart from DAV35 and S2105, the cadmium losses (due to the high Cd vapour pressure inside the ampoule during the growth of the crystal) are compensated for: so, CTV2, DAV31, DAV36 and DAV37 are assumed to be ‘stoichiometric’ crystals. On the other hand, DAV35 and S2105 are supposed to have some cadmium vacancies ([VCd] in Table 4).

3.2.2. Measurements The results are presented by separating the crystals in two classes, following their behaviour in the presence of an external continuous electric field. All the measurements are carried out at 1.32 mm wavelength. With no field applied, an electron hole competition factor is obtained for each crystal. Its sign reflects the dominating photocarriers and are precised in Fig. 19. Other values at lower (1.048 mm) and upper (1.535 mm) wavelengths had been presented in previous sections [31, 32].

3.2.2.1. Class I crystals Three crystals are in this class: DAV31, DAV35 and DAV37. The left part of the Fig. 20 shows the TWM gain (G) versus the intensity (I0) with a continuous 10 kV/cm electric field and a 10 mm grating period, for theses crystals. The gain curves present the expected S-shape due to the saturation of the space charge field inside the crystal in the limit of equilibrium between photoionisation of the carriers and re-capture process. As previously discussed, the optically and thermally

170 J.-Y. Moisan

Table 4

Main physico-chemical specifications of the V-doped CdTe and CdZnTe crystals studied

DAV31 DAV35 DAV37 S2105 DAV36 CTV2

[V]SIMS (cm3)

% Zn

[VCd]

[Acc.]

x (l ¼ 1.32mm)

[V2þ]EPR (cm3)

[V3þ]EPR (cm3)

3  1017 4  1017 1  1017 2  1017 7  1017 3  1017

4 4 4 1 0 0

Stoichiometric Non-stoichiometric Stoichiometric Non-stoichiometric Stoichiometric Stoichiometric

High High High Low High Low

þ0.68 þ0.23 þ0.22 0.37 0.57 0.60

3  1017 3.6  1017 1  1017 1.7  1017 Not seen Not seen

Not seen Not seen Not seen Not seen 5  1015 7  1014

Electron/hole Competition Factor

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1 h+

DAV31

0.6 ⴛ #

0.2

DAV35 DAV37

Wavelength (μm)

–0.2 S2105 ⴛ

–0.6 –1

1

1.2

DAV36 CTV2

1.4

e− 1.6

1.8

Figure 19 Electron–hole competition factor x against wavelength for the six studied crystals. The results are mainly obtained for the 1.32 mm wavelength. The main optically emitted carriers are holes in the upper part of the graph, and electrons in the lower part.

main emitted carriers have the same sign (holes), so no electron hole resonant effect is expected. This also means that the electron/hole competition factor x, which is equal, in a one deep-level model and in the limit of short grating period, to x¼

sp 0 pT0  sn 0 nT0 sp 0 pT0 þ sn 0 nT0

ð13Þ

is positive. Above equation can now be written as sp 0 pT0 > sn 0 nT0

ð14Þ

The x values are reported in Table 4.

3.2.2.2. Class II crystals The right part of Fig. 20 reports the same experiments for class II crystals (CTV2, DAV36 and S2105). An electron/hole resonant effect is clearly seen at about 1 mW/cm2 intensity for this set of three crystals, as discussed previously. If one considers that semiinsulating crystals are residual p-type materials [33] one can conclude that in this class II, the main optically emitted carriers are electrons. It means that the x parameter is negative, thus sp 0 pT0 < sn 0 nT0

ð15Þ

The experimental results are given in Table 4. For the given 1.32 mm wavelength, the major difference between the two classes of crystal is therefore a change of the dominant photoionised carrier. In a model of only one deep centre induced by V doping (i.e. with constant photoionisation cross section sn0 and sp0 for all crystals) the reverse of the sign would be explained by the reverse in the pT0/nT0 (or V3þ/V2þ) ratio. However the EPR analysis will show that this hypothesis is not consistent.

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TWM Gain (cm–1)

TWM Gain (cm–1)

DAV35 300K 10kV/cm 0.4 1,32mm Ip / Is: 100 Period: 10mm

2 S2105 5kV/cm 1 1,32mm Ip/Is: 100 Period: 10mm

0.2

A

D

0

0.4

DAV37-6 300K 10kV/cm 1,32mm Ip / Is: 100 Period: 10mm

2

E

DAV31 300K 10kV/cm 0.2 1,32mm Ip / Is: 100 Period: 10mm

0

0.2 CTV2 4kV/cm 1,32mm Ip/Is: 100 Period: 10mm

0.1

0

C

6kV/cm 1,32mm Ip/Is: 100 Period: 10mm

1

0

–0.2 0.001

0 DAV36-8b DAV36-5b

0.2

B

310° 300° 290°

0.01

0.1

1

Intensity (mW.cm–2)

10

0 0.001

100

F

0.01

0.1

1

10

100

Intensity (mW.cm–2)

Figure 20 Two wave-mixing gain versus pump beam intensity under a continuous electric field. (A), (B) and (C) are respectively DAV35, DAV37–6 and DAV31. Note the S-shape in all cases. (D), (E) and (F) are respectively for S2105, DAV36 and CTV2 crystals. An electron–hole resonant effect is clearly seen in this case. For convenience G has been plotted as positive but the x factor is clearly measured as negative. Crystals with different numbers have been cut from different place in the ingot (numbered from the top to the tail of the ingot).

3.2.3. EPR results

EPR results have been carried out on both classes of crystals. V3þ has been already observed in CdTe:V crystals [33] by EPR analysis and the content [V3þ] can be evaluated in this way. In CdTe:V crystals (DAV36 and CTV2) V3þ is clearly identified with the experimental setup. Results are reported in Table 4. But V3þ has never been seen in CdZnTe alloy

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173

crystals. This could be due to the low degree of compensation since these crystals at room temperature p-type conductive with hole concentration of 1015 cm3 only. However, from Eq. (15) and with the photoionisation cross-section values given by Bre´mond et al. [33], these crystals are expected to have a V3þ content which would lie well above the limit of detection of the used EPR apparatus. This limit is estimated to be around 3  1014 cm3 for CdTe:V crystals and about one decade higher if one considers a reasonable line-width increase due to the Zn alloying. On the other hand, for Zn alloyed crystals, a signal attributed to V2þ is detected [33]. In all the crystals where such a low C2v symmetric related signal has been detected, the maximum of the spectrum follows the segregation of V measured by SIMS. This rules out the possible attribution to, for example, a V-related complex. So a quantitatively analysis has been tentatively performed. First one sample (DAV37), where a weak resonance signal of an effective mass shallow donor is observed, is proceeded. The hypothesis was done that in these crystal the [V2þ] is equal to [V] as determined by SIMS experiments. EPR results estimate a nearstoichiometric V charge state ([V]SIMS  [V2þ]EPR) for the four crystals with a Zn content (see Table 4). Even though the non-observation of the V2þ spectrum in CdTe:V crystals has received possible interpretations [35] it is still a mater of discussion. The fact that S2105 crystal belongs to the class II crystals, makes this sample an intermediate crystal which will be discussed below.

3.2.4. Discussion Following their photorefractive properties, two classes of crystals of CdTe:V have been defined and the main difference is obviously expressed by Eqs. (14) and (15). The main optically emitted carriers, at 1.32 mm wavelength, are holes (hþ) in the class I crystals and electrons (e) in the class II ones. In both classes, holes are the thermally emitted carriers. It comes out that the resonant e/hþ effect is only observed in the class II crystals (i.e. when x is negative). In a first approach, if one considers that with or without a continuous electric field, the photorefractive properties are explained – in all crystals – by one deep trap level due to the V doping effect, one can conclude the following (see Fig. 19): 1. The x value increases as the wavelength increases; so the ratio sp0/sn0 increases with the wavelength. This is confirmed by previous deep-level spectroscopy (DLOS) experiments [35]. 2. The ratio V3þ/V2þ is larger for class I crystals than for class II ones. This is not confirmed by EPR results (see Table 4). The problem is therefore to find out what material parameters can explain this behaviour.

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3.2.4.1. Effects of physico-chemical parameters In Table 4 some physicochemical parameters are also summarised. The total V content evaluated by SIMS ([V]SIMS) has been measured in each crystal. [Acc.] is the content of acceptor like impurities (Li, Na, K, etc.) which could – as for Cd losses – influence the charge state of V (i.e. V3þ or V2þ) dopant, for which the concentrations have been evaluated by EPR experiments. The results of SIMS measurements on the six studied crystals are given in Fig. 21. Absolute contents are only given for [V] concentrations which lie between 1017 and 7  1017 cm3. For all the other elements only the relative concentration is indicated.

1.×1018 DAV31 DAV35

1.×1017

DAV37 S2105

1.×1016

DAV36 CTV2

1.×1015

1.×1014 1.×1013 V

Zn

Relative Concentrations (Arb. Unit.)

DAV31 DAV35 DAV37 S2105 DAV36 CTV2

Li

Na

K

Al

Cl

Figure 21 SIMS results: in the upper part, absolute concentrations of V are given and Zn contents are relative values as in the lower part for different impurities. Except for Li, no significant differences are clearly seen for the two crystal classes.

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The two DAV36 and CTV2 crystals have a Zn content as a residual impurity, but in S2105, [Zn] is about four times lower than in the class I crystals. No clear differences appear in the Al and the Cl content – which are donor-type impurities – between the two classes of crystals; the same feature for K (an acceptor type impurity). The Li content appears to vary in a large range: one decade higher crystals than in the others. From Table 4, it appears that the main difference between the class I and class II crystals remains the Zn content. So, one can conclude that [Zn] is the only parameter which can explain the difference between the two classes. If one now suppose that we always involve the same deep trap level (with the same photoionisation cross sections), it follows that increasing [Zn] should lead to a drastic increase in pT0 or [V3þ], which is not observed in EPR analysis. DLOS has been already obtained on such crystals [34]. Then major result is the characterisation of a main deep level in the band gap of V-doped CdTe or CdZnTe crystals, which is proposed to be the main trap responsible for the photorefractive behaviour of these materials. This level is located at an activation energy of 0.95 eV below the bottom of the conduction band. Its concentration seems to be proportional to the V content introduced in CdZnTe crystals, as determined by SIMS experiments. Moreover, DLTS experiments on an un-doped CdTe crystal [29] have revealed a deep trap level signature at 1.08 eV but in two decade lower concentrations range (1013 cm3). Actually, DLTS gives a position energy for the trap which is the sum of the thermal ionisation energy and possible thermal activation of capture cross section; explaining the apparent discrepancy for the 0.95 eV level attributed to V and the value usually given in the literature (Ecb – 0.7 eV). Its photoionisation cross section sn0(hn) and sp0(hn) have been measured in absolute scale (better than 50% of uncertainties) by electrical and optical DLTS, respectively. A band gap change is expected with the [Zn] increase: about 7 meV [36] have been experimented for 1% of Zn. One can assume that this increase of the band gap has only a weak effect on the photoionisation cross sections. So the largest reasonable changes of this rates for an increase of 28 meV (4% of Zn), does not provide full explanation for the differences observed for the crystals with and without Zn.

3.2.4.2. Effects of a second deep trap level DLOS has shown [37] that another deep trap level appears in CdTe:V crystals with a low (or no) Zn content. This level is situated at an activation energy of 0.78 eV below the bottom of the conduction band and can appear in a higher concentration than for the 0.95 eV level. Similar results have been confirmed for DAV36 crystal by PICTS experiments; but DLTS characterisations on this crystal are not available. So no quantitative evaluation of the deep level concentrations are known in this class II crystal. Note that this second

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deep level was also present in DAV31 crystal [34] but in lower concentration than for the 0.95 eV level. Based on these results, a model (see above) has been developed with two traps in the band gap of the material: an acceptor type second deep level is supposed. This means that the Fermi level pinned near the main 0.95 eV deep trap. DLOS experiments [37] indicate than one can neglect sp20 with regard to sn0 at 1.32 mm wavelength. Even though material parameters have been deduced by DLTS and DLOS experiments on converted n-type annealed samples, these results are expected to be also valid for as-grown high resistivity samples. One can see in Fig. 22 some numerical computation results for the G versus I0 curves for such a two deep traps model. Material parameters are precised on the legend of the figure. They are mainly extracted from DLTS results of DAV31 [34] and S2105 [37] crystals. As the total concentration of the second deep trap level (NT2) increases from a negligible or low value (like in DAV31 or type I crystal) to a dominant value (like in S2015 or type II crystals) a change is observed in the sign of the x (and also G) with an emergence of the resonant effect. The maximum of this resonant effect is situated around an intensity of 1 mW/cm2 which pleasantly agrees with experimental results (see Fig. 20).

NT2

GAIN (cm–1)

.5

0

0

5×1015

–.5

7×1015 5×1016 1×1016

–1. –1.5. –2. –2.5 10–3

10–2

10–1

1.

10

102

103

INTENSITY (mW.cm–2)

Figure 22 TWM gain as a function of total intensity for different values of the second deep level concentrations (NT2). Experimental parameters: l ¼ 1.32 mm, E0 ¼ 5 kV/cm. Material parameters for 0.95 eV (A) and 0.78 eV (B) deep levels: er ¼ 10.4, r41 ¼ 5.5 pm/V, nTA0 ¼ 3  1015 cm3, pTA0 ¼ 8  1015 cm3, mn ¼ 1000 cm2/Vs, mp ¼ 1000 cm2/Vs, snA0 ¼ 2.4  1017 cm–2, spA0 ¼ 7.4  1017 cm2, CnA ¼ 2  106 cm3/s, CpA ¼ 3  108 cm3/s, ethnA ¼ 0.038 s1, ethpA ¼ 0.56 s1, spB0 ¼ 3  1016 cm2, snB0 ¼ 1018 cm2, CnB ¼ 2  106 cm3/s, CpB ¼ 108 cm3/s, ethnB ¼ 0.01 s1, ethpB ¼ 0.1 s1.

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177

.8

GAIN (cm–1)

.6

a b

.4

NT2: a : 1×1014 b : 2×1015 c : 4×1015 d : 5×1015 e : 7×1015

c

.2

d

0

e

–2. –.4

10–3

10–2

10–1

1.

10

102

103

INTENSITY (mW.cm–2)

Figure 23 Same numerical computations as in Fig. 22 but for E0 ¼ 10 kV/cm and material parameters closer for DAV31 crystal: nTA0 ¼ 7  1015 cm3, pTA0 ¼ 6.7  1015 cm3, CnA ¼ 6  105 cm3/s, CpA ¼ 1.5  108 cm3/s, ethnA ¼ 0.23 s1, ethpA ¼ 0.56 s1, snB0 ¼ 3  1016 cm2, spB0 ¼ 1018 cm2, CnB ¼ 2  106 cm3/s, CpB ¼ 108 cm3/s, ethnB ¼ 0.01 s1, ethpB ¼ 0.1 s1.

In Fig. 23, some material parameters deduced from DLTS results of DAV31 crystal [34] are used. One can see that for NT2 lower than for S2105 crystal but not negligible, the G changes sign for low I0 values depending on NT2. This effect has been experimentally detected (Fig. 20C) and had been also theoretically predicted [38] when more than one trap level is inserted in the photorefractivity theory. This phenomenon is due to an intensity dependence of the x factor, x ¼ x(I0) and would explain the difficulties to deduce precise material parameters with analytical expressions based one model with only one deep trap. Further analytical developments based on two deep traps both in thermal and optical interaction with both the valence end conduction band have to be carried out. Apart from these results, some thermal spectroscopy was also carried out of the resonance electron–hole effect for both crystals from class I (DAV37) and class II (DAV36 and S2105). The measurements have both been realised at l ¼ 1.06 and 1.32 mm wavelengths in order to get a resonance effect also for class I crystals (see Fig. 19). If one holds the hypothesis of crystals with a dark conductivity due to holes (i.e. dominated by thermal emission of holes), Arrhenius plots of the optimum intensity (at the centre of the resonance) versus 1/Tk should allow one to extract energies of activation of the deep level which dominate the photorefractive effect in each crystal. As it is expected, one can see in Fig. 24 that crystals of class I and class II are not ruled by the same deep trap. Due to its energy of activation of only 0.67 eV, S2105 is confirmed to be intermediate origin which could be due to a content of

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J.-Y. Moisan

–7 –8

DAV36.5 (1.32mm) Eact = 0.81eV

DAV37.6 (1.06mm) S2015 (1.06mm)

–9 Eact = 0.34eV –10 –11 –12 –13 –14 –15 2,6

Eact = 0.54eV

Eact = 0.67eV

Eo = 5kV/cm Ip/Ip = 100 Grating period = 10mm 2,7

2,8

2,9

3

3,1 1000/ T

3,2

3,3

3,4

3,5

3,6

(K–1)

Figure 24 Arrhenius plots of the optimal intensity (at the centre of the resonance) for crystals DAV37, DAV36 and S2105. From the slope of each strait line one can deduce an energy of activation of deep traps which thermally dominates the photorefractive effect in each crystal. Note that class I and class II crystals seem not to be dominated by the same deep levels.

only 1%. Note that energies of activation come from thermal emission of holes; so they are referred to the maximum of the valence band. With an estimated band gap, at 0 K, of Eg ¼ 1.6 eV, – the effect of the Zn in the range 0 to 4% is neglected – it has been illustrated by a photorefractive method the general tendency of crystals governed by different deep levels which was deduced from DLTS and PICTS measurements.

4. APPLICATIONS One of the greatest interest for CdTe:V is its high sensitivity at telecommunication wavelengths (1.3 and 1.5 mm). This has been largely demonstrated in the previous sections. With these materials, double-phase conjugated mirrors [39] (DPCMs) could become interesting for free space bi-directional optical links, owing the self-aligning properties of phase conjugation [40–42].

4.1. DPCM experiments The same crystals (8 mm long) as studied above were used (DAV31L A and B). High TWM gain is obtained under an alternating square-wave electric field. Antireflection coatings have been deposited onto both

Photorefractive CdTe

179

optical faces. Their reflectivity is smaller than 0.3% from 1.2 to 1.6 mm. The absorption at 1.54 mm of the two samples is, respectively, 1.35 and 1.22 cm1. TWM characterisation and DPCM experiments were performed with a 20 mW diode-pumped solid state laser. A TWM gain G of 6 cm1, reported above in another un-coated sample from the same ingot, was confirmed with these antireflection coated crystals under identical experimental conditions. The analytical expression [43] for the DPCM threshold in the presence of absorption a predicts for equal pump intensities a threshold value of G  6 cm1 for the two 8-mm-long crystals, which is the value that is experimentally obtained for 10 kV/cm. One therefore expects to observe the threshold electric field for DPCM near 10 kV/cm. The experimental setup is presented in Fig. 25. In some cases cylindrical lenses are used to focus the two beams in the sample. The conversion efficiency  (defined by the ratio of the diffracted beam intensity to the total transmitted one) measured as a function of the applied field frequency for the two crystals is shown in Fig. 26. With the high slew rate (>105 V/ms) high-voltage amplifier used in all the experiments reported here, the field shape can be assumed to be perfectly square for a frequency up to 500 Hz. Note that for the measurements reported in Figs. 26–28,

camera

screen Imaging lens for far field pattern observation collimated input beam

cylindrical lenses H.V.

Beam splitter Cd Te Cystal

Optical detector

Figure 25 Experimental setup used in all the experiments reported here. Cylindrical lenses allow for the elimination of conical diffraction and were removed to yield the experimental data of Figs. 26 and 27 pump depletion of collimated beams. HV, high voltage.

180

J.-Y. Moisan

14

Conversion efficiency (%)

12 DAV31L - A

10

DAV31L - B

8 6 4 2 0

1

10

100

1000

Field frequency (Hz)

Figure 26 Experimental DPCM conversion efficiency versus applied field frequency for the two 8-mm-long crystals. The grating period is L ¼ 5.5 mm. The intensity of each pump is 10 mW/cm2 and E0 ¼ 10 kV/cm. 14 12 Conversion efficiency (%)

Ip = 20 mW/cm2 Ip = 15 mW/cm2

10

Ip = 10 mW/cm2 Ip = 5 mW/cm2

8 6 4 2 0 3000

4000

5000

6000 7000 8000 Applied field (V/cm)

9000

10000

Figure 27 Experimental DPCM conversion efficiency as a function of the applied field for DAV31LA. The field frequency value (70 Hz) is optimised for pump intensity of 20 mW/cm2. Ip, pump intensity.

Photorefractive CdTe

181

24

Conversion efficiency (%)

Eo = 10kV/cm

18

9

12 8

7

6

6 0

2

4

6

8

10 12 14 16 Grating period (mm)

18

20

22

Figure 28 Spatial bandwidth characterisation of DPCM for different applied fields E0. The field frequency is 70 Hz, and the pump intensity is 20 mW/cm2.

collimated beams was used so that the conical diffraction was not eliminated and  was measured from pump depletion. Figure 26 illustrates the influence of the frequency dependence under a square shape external field. Because the TWM gain at high field frequency is markedly decreased, the DPCM efficiency can be severely reduced (DAV31LA) or even quenched (DAV31LB) because in the latter case the product G  L becomes lower than the threshold value for the DPCM. Even though such a difference could be explained by inhomogeneities of the second level, no definitive explanation is given for this behaviour for the two samples operating from different parts of the same ingot. In any case, one cannot explain these by taking the absorption into account, because the highest efficiency is obtained with the most absorbing sample. Pump depletion of the DPCM versus the applied field with different pump intensities is illustrated in Fig. 27. All the experiments were performed at a field frequency of 70 Hz. This frequency enables to avoid fluctuations of the diffracted beam intensity with the highest illumination (20 mW/cm2 for each pump). At 70 Hz, the best result is obtained with the highest illumination. The results obtained at low illumination could in principle be improved by reduction of the frequency [12, 14] with no increase in the intensity fluctuations, as the response time of the photorefractive effect is inversely proportional to the incident illumination.

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J.-Y. Moisan

The threshold determined by the appearance of conical diffraction occurs at E0 ¼ 6 kV/cm for all intensities. Note that for electric field lower than this threshold, the experimental values in Fig. 27 are due to a fanning effect. It is believed that the unpredicted low threshold value arises from an underestimation of the maximum gain in TWM experiments, especially for long crystals, because of the so-called large-signal effect. Indeed, with a TWM gain of 6 cm1 a beam ratio of 500 at the input face of the crystal is reduced to 4 at the output, leading to strong nonlinearities of the space charge field. The TWM gain versus intensity dependence, for operation at an optimum external field frequency [14], may also lead to some overestimation of the threshold field for the DPCM. Indeed, in a DPCM configuration, the crystal is illuminated from both sides, so that variations of illumination are much smaller than in the TWM experiments. Because of the good photosensitivity at 1.54 mm a DPCM is obtained by using only a 5 mW/cm2 of power for each of the pumps in CdTe:V, as shown in Fig. 26. Figure 28 shows the experimental conversion efficiency versus grating period for different applied fields. The maximum efficiency is obtained for a grating period of 8 mm. A larger optimum grating period (15 mm) has been observed with a sample from the same ingot in TWM experiments [14]. This phenomenon, already known for InP:Fe [42, 44], could be related to nonlinearities of the space charge field at high fringe modulation (as is the case in DPCM experiments) and occurs mainly when the small signal gain is the highest (i.e. near 15 mm). A maximum conversion efficiency of 22% for an applied field of 10 kV/cm is reported. This leads to a maximum diffraction efficiency (calculated by inclusion of the optical absorption) of 7.4% for crystal DAV31LA.

4.2. DPCM between single-mode fibres To realise true phase conjugation (i.e. without conical diffraction) the insertion of aberrators or cylindrical lenses [45] was proposed. In the latter geometry, two cylindrical lenses of 20 cm focal length focus the beams in the incident plane, making diffraction impossible outside that plane (Fig. 25). A beam splitter is inserted into the path of one of the collimated pump beam to be viewed without astigmatism by a camera. In Fig. 29, the far-field pattern is represented with and without cylindrical lenses. The resolution of the imaging optics is too low to ensure that the device is diffraction limited. Insertion of single-mode fibres for more precise evaluation is being considered. In a configuration with cylindrical lenses the input beams overlap inside the crystal in a range across the electrodes but in a thin region in a perpendicular direction (h112i axis here). One possible application of such linear arrays of a DPCM is the adaptative coupling of the beams within output single-mode fibres into a

Photorefractive CdTe

183

ZOOM X3

0′

4°3

4°3

0′

ZOOM X4

2°

5′

15

′

2°1

Figure 29 Experimental far-field pattern image (with an 300-mm lens) on a Vidicon camera analysed with a PC-monitored beam analyser. The upper figure represents the cross section of the conjugated beam without cylindrical lenses. Conical diffraction is clearly seen; the lower figure shows its successful elimination with two cylindrical lenses. Note the different scales for the two parts of the figure.

vector–matrix crossbar switch. Indeed, this architecture, in which the different beam directions (corresponding to one output) are situated in a single plane, is well matched to one-dimensional phase conjugation. Single-mode fibres could be used in such an architecture, leading to reduce insertion losses and thus to high switch capacities and bit rates, even though cross-talk problem have to be characterised first.

5. CONCLUSION This study was performed to demonstrate that optical beam steering could be possible using photorefactivity properties of crystals. The first part of this chapter has explained the photorefractive effect in a semiconductive crystal. Because in a optical beam steering system, optical beams are issued from optical fibres at telecommunication wavelengths (i.e. in near-infrared range of wavelengths) the used crystals must have high

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photorefractive figure of merit. Two semiconductive materials have such properties: InP:Fe and CdTe:V as largely presented in the literature for the first one. In the last part of this chapter, demonstration of beam steering is reported: an optical system is presented and the conditions for such an optical function seems to be fulfilled. The environmental conditions for the used CdTe:V crystal have been previously studied and applied: they were obtained through the measurements of the optical gain in a TWM optical experiment. It appeared that the following materials conditions have to be fulfilled:  The V content that creates a trap level in the band gap, as it is well

known.

 The Zn content: Zn is known for its effect on crystal synthesis and

is generally added for this reason. But it appears that Zn content has a drastic effect on the photorefractive properties following the wavelengths range. To manage as precisely as possible the photorefractive properties of the CdTe:V material, it is clear that it is necessary to manage the semiconductive behaviour of this material as precisely as possible. This is the object of the second part of this chapter. So it has been demonstrated that CdTe:V crystal is a good candidate for optical systems in the near infrared wavelengths range, but the conditions are to fulfil some chemical parameters (V and Zn contents) and then the semiconductive behaviour of the material.

ACKNOWLEDGEMENTS The author would like to thank the following colleagues from his laboratory (CNET, now Orange Labs): P. Gravey, N. Wolffer, G. Martel, V. Vieux for the optical set ups and G. Picoli, B. Lambert and M. Gauneau for the study of the properties of the materials. The author also gratefully acknowledges the contributions of colleagues from several different laboratories who participated in this study: R. Triboulet, Y. Marfaing and co-workers from the CNRS Lab of Meudon, P. Siffert, M. Hadj-Ali, J.M. Koebel, and co-workers from the CNRS Lab of Strasbourg, G. Bremond and co-workers from INSA Lyon and M. Pugnet from INSA Toulouse. This collective work, coordinated via a formal work group, has resulted in the material presented here, thanks to a fruitful collaboration via friendly and open discussion. It is thus a collective work that is presented here, but is conducted under a common project that explains the diversity of the technical domains treated and the interest found therein by each of the participants.

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REFERENCES [1] P. Gravey, J.Y. Moisan, Proc. Soc. Photo-Opt. Instrum. Eng. 1507 (1991) 239–246. [2] M. Ziari, W.H. Steier, P.M. Ranon, S. Trivedi, M.B. Klein, Appl. Phys. Lett. 60 (9) (1992) 1052. [3] A. Partovi, J. Millerd, A.M. Garmire, M. Ziari, W.H. Steier, S. Trivedi, M.B. Klein, Appl. Phys. Lett. 57 (9) (1990) 846. [4] J.M. Langer, Phys. Rev. B 38 (11) (1988) 7723. [5] W. Giriat, J.K. Furdyna, Semiconductors and Semimetals, vol.25, Academic Press, 1988. [6] N.V. Kukhtarev, Ferroelectrics 22 (1979) 949. [7] G.C. Valley, J. Opt. Soc. Am. B 1 (1984) 868. [8] S.I. Stepanov, M.P. Petrov, Opt. Commun. 53 (1985) 292. [9] A. Partovi, E. Garmire, J. Appl. Phys. 69 (10) (1991) 6885. [10] G. Picoli, P. Gravey, C. Ozkul, V. Vieux, J. Appl. Phys. 66 (8) (1989) 3798. [11] M. Ziari, W.H. Steier, P.M. Ranon, M.B. Klein, S. Trivedi, J. Opt. Soc. Am. B 9 (1992) 1461. [12] Y. Belaud, P. Delaye, J.C. Launay, G. Roosen, Opt. Commun. 105 (1994) 204. [13] J.Y. Moisan, P. Gravey, N. Wollfer, O. Moine, R. Triboulet, A. Aoudia, in: Topical Meeting on Photorefractive Materials, Effects and Devices, PRM’93, Kiev, 1993, pp. 283–286. [14] J.Y. Moisan, N. Wolffer, O. Moine, P. Gravey, G. Martel, A. Aoudia, E. Repzka, Y. Marfaing, R. Triboulet, J. Opt. Soc. Am. B 11 (1994) 1655. [15] J. Strait, J.D. Reed, N.V. Kukhtarev, Opt. Lett. 15 (1990) 209. [16] J.C. Launay, V. Mazoyer, J.P. Zielinger, Z. Guellil, P. Delaye, G. Roosen, Appl. Phys. A 55 (1992) 33. [17] A.M. Glass, M.B. Klein, G.C. Valley, Electon. Lett. 21 (1985) 220. [18] M.B. Klein, G.C. Valley, J. Appl. Phys. 57 (1985) 4901. [19] K. Walsh, A.K. Powell, C. Stace, Y.J. Hall, J. Opt. Soc. Am. B 7 (1990) 288. [20] N. Wolffer, P. Gravey, Ann. Phys. (NY) 16 (1991) 143. [21] G. Bre´mond, et al., E-MRS Spring Meeting 1994, Strasbourg, Opt. Mater. 4 (1995) 246. [22] S.I. Stepanov, M.P. Petrov, Opt. Commun. 53 (1985) 292. [23] R.B. Bylsma, A.M. Glass, D.H. Olson, Electon. Lett. 24 (1988) 360. [24] P. Mathey, G. Pauliat, J.C. Launay, G. Roosen, Opt. Commun. 82 (1991) 101. [25] K. Turki, G. Picoli, J.E. Viallet, J. Appl. Phys. 73 (1993) 8340. [26] P. Tayebati, D. Mahgerefteh, J. Opt. Soc. Am. B 8 (1990) 1053. [27] R.S. Rana, D.D. Nolte, R. Steldt, E.M. Monberg, J. Opt. Soc. Am. B 9 (1992) 1614. [28] P. Nouchi, J.P. Partanen, R.W. Hellwarth, in: Photoconductive Materials, Effects and Devices, vol. 14 of 1991 OSA Technical Digest Series, Optical Society of America, Washington, DC, 1991p. 236. [29] G. Bre´mond, Personal communication, (1993). [30] G. Pauliat, A. Villing, J.C. Launay, G. Roosen, J. Opt. Soc. Am. B 7 (1990) 1481. [31] J.Y. Moisan, P. Gravey, G. Martel, N. Wollfer, A. Aoudia, Y. Marfaing, R. Triboulet, M.C. Busch, M. Hadj-Ali, J.M. Koebel, P. Siffert, Opt. Mater. 4 (1995) 219. [32] J.Y. Moisan, N. Wollfer, O. Moine, P. Gravey, G. Martel, A. Aoudia, E. Repka, Y. Marfaing, R. Triboulet, J. Opt. Soc. Am. B 7 (1994) 1665. [33] B. Lambert, M. Gauneau, G. Grandpierre, M. Shoisswohl, H.J. Von Bardeleben, J.C. Launay, V. Mazoyer, A. Aoudia, E. Rzepka, Y. Marfaing, R. Triboulet, Opt. Mater. 4 (1995) 267. [34] G. Bre´mond, A. Zerrai, G. Marrakchi, A. Aoudia, Y. Marfaing, R. Triboulet, M.C. Busch, J.M. Koebel, M. Hadj-Ali, P. Siffert, J.Y. Moisan, Opt. Mater. 4 (1995) 246. [35] R.N. Schwartz, M. Ziari, S. Trivedi, Phys. Rev. B 49 (1994) 5274. [36] S. Adachi, T. Kimura, J. Appl. Phys. 32 (1993) 3496.

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[37] A. Zerrai, G. Bre´mond, G. Marrakchi, J.Y. Moisan, G. Martel, M. Gauneau, B. Lambert, P. Gravey, N. Wollfer, A. Aoudia, Y. Marfaing, R. Triboulet, J.M. Koebal, M. Hadj-Ali, P. Siffert, in: E-MRS Spring Meeting, Strasbourg, France, Elsevier Amsterdam, 1995. [38] P. Tayebati, J. Opt. Soc. Am. B 9 (1992) 415. [39] S. Weiss, S. Sternklar, B. Fischer, Opt. Lett. 12 (1987) 114. [40] R.B. Bylsma, A.M. Glass, D.H. Olson, M. Cronin-Golomb, Appl. Phys. Lett. 57 (1990) 1968. [41] P.L. Chua, D.T.J.H. Liu, L.J. Cheng, Appl. Phys. Lett. 57 (1990) 858. [42] V. Vieux, P. Gravey, N. Wolffer, G. Picoli, Appl. Phys. Lett. 58 (1991) 2880. [43] N. Wolffer, P. Gravey, J.Y. Moisan, C. Laulan, J.C. Launay, Opt. Commun. 73 (1989) 351. [44] J.E. Millerd, E.M. Garmire, M.B. Klein, J. Opt. Soc. Am. B 9 (1992) 1499. [45] M.P. Petrov, S.L. Sochava, S.I. Stepanov, Opt. Lett. 14 (1989) 284.

CHAPTER

IIB Cadmium Telluride-Based Solar Cells A.W. Brinkman

1. INTRODUCTION With a near-optimal room temperature band gap of 1.52 eV and a high absorption coefficient (>105 cm1) [1, 2] above the band edge, CdTe is in principle an ideal candidate for use as a thin film solar cell material. There is a correspondingly large literature on various aspects of CdTe-based cells and the present review makes no pretence to be comprehensive. The aim instead, is to introduce some of the background, review the state of the art, to briefly examine obstacles to the realisation of potential performance and to consider the degradation mechanisms that determine operational lifetime. The first reported CdTe-based solar cells were shallow p-n homojunction devices that had conversion efficiencies of 4–6% [3–5]. Eventually, efficiencies of up to 13% were achieved [6], but it was already clear that a combination of high absorption in the shallow surface layer, excessive surface recombination and problems in making a shallow p-n junction meant that the homojunction was unlikely to be successful and the focus of attention shifted to heterostructures, principally the n-CdS/p-CdTe heterojunction [7]. Other heterostructure devices were briefly investigated, for example, the ITO-CdTe [8], but the CdS/CdTe cells ultimately have proved to be superior. The CdS/CdTe cells are fabricated in the superstrate configuration (Fig. 1): a multilayered structure where layers of CdS and CdTe are deposited in sequence on transparent conducting oxide (TCO) – typically a doped tin oxide film (e.g. SnO2:F, InSnOx (ITO), etc.) – coated glass and a

Department of Physics, University of Durham, Durham DH1 3LE, United Kingdom

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Back contact CdS Window Layer

CdTe Absorber layer

Transparent contact (ITO)

Glass substrate

Sunlight

Figure 1 Superstrate configuration.

suitable back contact is applied to the CdTe layer. The wide band-gap CdS acts both as a widow layer admitting visible and infrared (l > 520 nm) light to the strongly absorbing CdTe and as the n-type limb of the n–p junction [7]. Studies based on the detailed balance Schockley-Queisser model [9], suggest a maximum efficiency for an ideal CdS/CdTe cell of about 29% (AM1.5G) [10]; however, these calculations do not take into account the polycrystalline nature of the CdS and CdTe layers with all the losses due to grain boundaries, voids etc. (Fig. 2), and in practice, only just over half of the theoretical best performance has been achieved in the laboratory. Ideally, the n-CdS layer is kept to the minimum ( 100 nm) commensurate with proper junction formation and made as conducting as possible – nþp one-sided junctions are more tolerant of conduction band discontinuities. [10]. The CdTe side of the junction is typically a few micrometre thick and frequently both the CdS and the CdTe layers are subjected to some post-deposition treatment, usually involving CdCl2 and/or heating in air. For a comprehensive review of the material issues see Durose et al. [11]. Contact

Grain boundaries

Voids

CdTe

CdS TCO Glass

Figure 2 ‘Real’ polycrystalline cell.

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2. STATE OF THE ART In principle, an efficiency of about 18% [12], (short circuit current density, JSC ¼ 27 mA/cm2, open circuit voltage, VOC ¼ 0.88 V and fill factor, FF ¼ 76%) should be possible for a polycrystalline CdS/CdTe cell. As listed in the latest solar cell efficiency tables [13], the best conversion efficiency to date for a laboratory-scale device is that recorded by Wu et al. [14] in 2001, which had an efficiency of 16.5% (JSC ¼ 25.9 mA/cm2, VOC ¼ 0.845 V, FF ¼ 75.5%, aperture area ¼ 1.03 cm2). The preceding decade had seen steady if slow improvements, with a 16% cell reported by Aramoto et al. [15] in 1997, and in 1993 by Britt and Ferekides [16] of a 15.5% efficient cell. The incremental nature of these improvements suggest that, either the performance of the Wu cell is close to what is in fact attainable, or alternatively, there are some materials issues that are not fully understood; for example, the role of post-deposition treatment in Cl/O2 heat treatments necessary for the fabrication of efficient cells, and which are discussed later. According to the same solar cell efficiency tables the best module efficiency is 10.7% for a module produced by BP Solarex [17] with an aperture area of 4874 cm2. The test module delivered 3.205 A at an output voltage of 26.21 V with a fill factor of 62.3%. Although lifetime and performance in the field have yet to be established (a field life of 25 years is generally accepted as being a necessary requirement), modules are being manufactured with production exceeding 6 MW in 2004 and expected to rise above 20 MW by 2006 [18].

3. DEVICE PROPERTIES 3.1. Back contact effects Producing good injecting contacts to p-CdTe is difficult. The combination of its large electron affinity (w ¼ 4.28 eV) and moderately wide band gap (Eg ¼ 1.52 eV) means that in the Schottky limit a simple metal contact would require a material with a work function greater than about 5.8 eV (Eg þ w) [19] to obtain the valence band flat-band condition; no such metal exists. Observed barrier heights (e.g. by internal photoemission, CV etc.) are almost always larger than this due to pinning by interface states that create a dipole, in effect changing the electron affinity [20, 21], and making the problem even more severe. As a result alternative contact ‘recipes’ have to be adopted, such as doping the back surface to produce a narrow barrier through which carriers can tunnel or introducing buffer layers to accommodate the contact potential. Although a large variety of back contact regimes have been developed – the technologies are discussed

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in detail later – there is invariably some sort of residual barrier, which depending on the magnitude will adversely affect performance. By confining carriers at the contact–CdTe interface, barriers increase surface recombination losses. For a given surface recombination velocity Sbc, the recombination current density JR qnSbc, that is, JR is proportional to the electron density, n, at the rear surface. For a small barrier, < 0.2 eV, n will also be small and so will JR even for sizeable values of Sbc. Consequently, the impact on efficiency is insignificant, but for larger potential barriers the reverse is true.

3.1.2. Rollover When the current is constrained by a barrier the J–V characteristics often display ‘rollover’, where the current saturates at bias voltages just above VOC and frequently ‘crossover’ as well, where the dark and illuminated J–V curves intersect in the first quadrant (Fig. 3) [22]. Rollover, formally defined [23] as the case when the second derivative of the J–V curves evaluated at VOC is negative, is generally attributed to the back contact barrier acting as a reversed biased diode in series with the output [24] (Fig. 3, inset). The majority hole current is then limited by the reverse saturation current (JSbcd) of the back contact diode (BCD) (Fig. 3, curve b). 0.030 Back contact

J (a) (d)

Solar cell

Current density (J)

0.015

V (b)

0

(c) Rollover

Crossover –0.015

0

0.4

0.8

1.2

Voltage (V)

Figure 3 Current–voltage characteristics illustrating rollover and crossover: (a) dark J–V with no BCD; (b) dark J–V with BCD; (c) photovoltaic response with BCD but no contribution from minority electron current; (d) as for (c) but with electron current contribution. Inset shows equivalent circuit of solar cell including BCD.

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This will occur under both illumination and in the dark. Often rollover is observed to be temperature dependent, becoming increasingly apparent as the temperature is reduced, because JSbcd is proportional to the thermal electron velocity [24]. Niemegeers and Burgelman [22] have shown that the dark current will start to saturate when the bias voltage reaches:
 AkB T JSbcd ð1Þ ln VSDark  JSpn q where JSpn is the reverse saturation current density of the main CdS/CdTe junction (with ideality factor, A), and KB, T and q are the Boltzmann constant, temperature and electron charge, respectively. Under illumination current saturation sets in at a higher bias voltage (determined by JL/JSbcd:JL is the photo-generated current):
 AkB T JL ð2Þ VSLight  VSDark þ ln 1 þ JSbcd q but the total current should still saturate at JSbcd (Fig. 3, curve c). The corresponding perturbation in the output characteristics can lead to a significant deterioration in the fill factor if the barrier height is large and/or the saturation current is very small. As a series component, the BCD should not, in principle, affect VOC nor should it significantly limit JSC. Under short circuit conditions (and in the absence of other losses), hole current flow through the back contact should induce an equal potential, DV, of opposite sign to that generated across the primary junction, thus the BCD is forward biased and no longer limiting:
 kB T JL ð3Þ ln 1 þ DV ¼ JSbcd q The fill factor, FF0, for an ideal cell (i.e. one in which series and shunt resistive losses can be neglected) is a function of the open circuit voltage. For voltages below the maximum power point, the total device current is effectively constant and therefore so is the bias voltage across the BCD, DV. Consequently, the current–voltage characteristic for the non-ideal diode can be approximated by a translation of the ideal J–V curve by an amount DV. The non-ideal fill factor would then be approximated by the ideal FF00 at (VOC – DV); but the actual open circuit voltage, of course, in the non-ideal diode is VOC and not (VOC – DV). Thus the non-ideal FF is given by FF00 , reduced by a factor dependent on the degree to which the BCD is limiting performance: 
 0 kB T JL ð4Þ ln 1 þ FF  FF0 1  JSbcd qVOC

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Clearly, as long as JL/JSbcd 1 (i.e. JL not limited by JSbcd) any reduction in the FF will be small and unimportant compared to other factors, but that as the ratio increases, FF loss starts to become significant.

3.1.2. Open circuit voltage In principle, CdS/CdTe cells should have high values of VOC, (a consequence of the relatively large band gap energy), and one reason for the high theoretical efficiency [10]. Predicted VOC values are seldom observed in any of the solar cell systems, but the VOC deficit (Eg–qVOC), as it is termed, is large in CdS/CdTe devices compared to other cells [20]. The deficit can arise from a variety of causes, although a potential barrier at the back contact is in general involved. The interaction of carrier lifetime and a back contact barrier is a case in point [23]. Where the back contact is not dominant, but carrier lifetimes are short, space charge region recombination is high, and VOC is reduced. Longer lifetimes imply less space charge region recombination, and consequently an increased value of VOC. If instead the back contact barrier is significant, then VOC, FF and hence efficiency are affected irrespective of the carrier lifetime. Problems may also arise in thin cells if the space charge regions associated with the BCD and the main p-n junction overlap [25]. The two independent diodes model [24] breaks down and the cell starts to behave as a reach-through diode [26], where at some threshold voltage, VRT, the CdTe becomes fully depleted. The current is no longer controlled by either the main junction or the BCD and if this occurs in the fourth quadrant, the output voltage at open circuit will be VRT, that is, VOC is in effect reduced. Although for thick cells, with small barriers, reach-through is unimportant, the need to cut costs leads inevitably to a reduction in CdTe thickness with a concomitant increase in the probability of reach through effects. More importantly, back contacts are not laterally uniform and the magnitude of the associated potential barrier will vary with position. Local points of weakness can arise creating reach-through ‘microdiodes’, even though the CdTe layer as a whole has not become depleted. Except at short circuit, the microdiodes are forward biased and constitute shunt current paths, reducing VOC in the process. Clearly, this is more significant for thin CdTe layers where the incidence of microdiodes is likely to be greater. Reach-through effects are contingent on the magnitude of the back contact potential barrier, demonstrating once more the need to reduce this as much as possible.

3.1.3. Crossover Crossover is less well understood and there is still some debate [27] as to whether it is associated with the front part of the cell or with the back contact. Niemegeers and Burgelman [22] calculated the electron surface

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recombination current (due to photo-excitation) by solving the transport equation in the neutral zone (i.e. the region bounded by the edges of the p-n junction and BCD space charge regions). They showed that if the diffusion length (Ln) was sufficiently long in relation to the width of the field free region (d), then the electron current could make a significant contribution to the total current. The total current would still saturate, but at a higher value than in the dark (Fig. 3, curve d). Beier et al. [28] extended the model, to allow for the voltage and wavelength dependence of the contact saturation current. In deriving the hole current, Jp, the contact saturation current had been assumed to be a constant. In real devices it is often found to be dependent on the voltage and the illumination. Electron transport across the neutral region is given (in the dark) by the term exp (–d/Ln), that is, the probability that an electron will travel from the edge of the junction space charge region to the edge of the BCD depletion width. Under forward bias, (reverse bias for the BCD), the BCD space charge region is increased, hence reducing the width of the neutral region, and resulting in pronounced voltage dependence. The intensity and spectral content of any illumination can also affect the electron current due to barrier lowering by absorption at interface acceptor states. Crossover has also been attributed to the effects of illumination on deep levels in the CdTe. Ko¨ntges et al. [27] modelled the case where there were a high density (1016 cm3) of deep acceptor states with dissimilar capture cross sections for holes and electrons of 1018 and 1013 cm2, respectively. The low acceptor concentration (Na < 5  1014 cm3) in CdTe solar cells and the low hole mobility ( 50 cm2/(V s)), favour a diffusion model for current transport over a back contact barrier less than 0.5 eV. If under illumination acceptor traps are ionised, then the effective doping at the back contact region is increased. In the diffusion model for transport over a barrier, current is proportional to Na½, thus photoionisation of the acceptor traps in the vicinity of the back contact will increase JSbcd, with crossover in the J–V characteristics as a result.

3.2. Inter-diffusion at the CdTe/CdS junction Inter-diffusion across the CdS/CdTe junction is inevitable, creating intermediate ternary compounds of the form CdSxTe1–x [29–31]. Ohata et al. [32] were the first to publish the phase diagram for the CdS–CdTe pseudobinary system from high temperatures down to 700  C. Nunoue et al. [33] extended the range to 650  C. The two binaries are not miscible in all proportions except at high temperature and Ohata and co-workers were only able to obtain mixed crystals over almost the complete compositional range by annealing mixtures of CdS and CdTe at temperatures of 1000  C

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and quenching. Using XRD, the crystal structure of CdSxTe1–x was found to be zincblende when x < 0.2 and wurtzite at higher values of x; it was concluded that at 1000  C the miscibility gap was narrow. The liquidus and solidus were measured by differential thermal analysis and found to converge to the same minimum temperature of 1071  C at x ¼ 0.2. In the solid phase they observed that the miscibility gap increased rapidly with decreasing temperature, which they attributed to lattice strain due to the large difference in the atomic radii of the group VI constituents. At the lower temperature of 700  C, the CdSxTe1–x is zincblende for x < 0.2 and wurtzite for x > 0.9 with a mixture of phases in between. Jensen et al. [29] estimated the solubility of S in CdTe to be 5.8% at a temperature of 415  C (a typical temperature for Cl doping in electroplated cells, though significantly below temperatures encountered in close space sublimation). As a result, the interface between the CdTe and the CdS tends to form two (or more) intermediate ternary compounds, a Te-rich layer (CdTe1–xSx) and a S-rich layer (CdTeyS1–y). Nakayama et al. [34] reported a structure comprising two Te-rich zincblende layers (y ¼ 0.97 and y ¼ 0.96) adjacent to the CdTe side of the junction and a S-rich wurtzite layer (x ¼ 0.99) on the CdS side in screen printed cells. Similar results were obtained by Rogers et al. [35]. They used synchrotron radiation to measure the lattice parameter variation in electroplated CdTe/CdS cells as a function of CdTe layer thickness and anneal time at 450  C (with and without chlorine), and found considerable intermixing. For example for a chlorinated structure where the thicknesses of the CdS and CdTe layers were 80 nm and 400 nm respectively, the CdS was completely ‘consumed’ after only two minutes anneal to form CdS0.93Te0.07 which did not appear to change with further annealing. Over the same two minute period, cross-diffusion of S into the CdTe layer created a non-stoichiometric layer of composition, CdTe0.98S0.02, adjacent to the original junction. Eventually after sufficient annealing time ( 15 min), the average composition of the 400 nm CdTe layer changed to CdTe0.95S0.05 in agreement with the solubility studies of Jensen et al. [29], the S being derived from the non-stoichiometric layer. Lane and co-workers have also measured the variation in the CdTe lattice parameter as a function of depth [36] and concluded that the CdTe layer must have been under an in-plane stress of 140 MPa near the interface, sufficient to introduce a significant level of structural defects. On annealing the lattice parameter was found to follow a decreasing trend, which was attributed to the in-diffusion of S, to give a Te-rich mixed crystal layer with y ¼ 0.96, in good agreement with Nakayama et al. [34]. The interfacial compounds play an important role in accommodating the 10% mismatch between CdS and CdTe [37] in particular reducing the concentration of recombination centres. Carrier lifetime should be

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increased and, as discussed above [23], rollover with its detrimental effect on FF should be less severe. Inter-diffusion will also affect the band line-ups [38]. As already noted, conversion of CdS and CdTe to the ternary compounds ‘consumes’ parts of both. However, since the thickness of the CdS layer is only 100 nm, while that of the CdTe is typically several mm, the effect on the CdS is disproportionately greater. The residual CdS layer is sandwiched between the nþ TCO and the smaller band gap ternary, creating a ‘hump’ in the energy bands. Discontinuities in the band lineups on the CdTeyS1y side produce barriers to minority hole injection across the junction. Holes generated by absorption of light in the CdS, will drift into the hump, which constitutes a potential minimum, from where they may be emitted over the barrier or trapped in the deep acceptor states of which there will be a high density. The trapped holes will neutralise some of the negative charge, thereby reducing the barrier height and increasing emission over the barrier. When externally biased, the potentials throughout the cell will self-adjust to allow the same current to flow through all series connected parts of the device. Superposition no longer holds and crossover will occur in the characteristics. The spectral response (variation of quantum efficiency (QE) with wavelength of illumination, QE(l)) of CdS/CdTe solar cells displays a windowtype characteristic, bounded at the short and long wavelength ends by the band gaps of the CdS window layer and the CdTe absorber layer. The optical band gap energy of CdSxTe1–.x is given by the empirical expression [39]: Eg ðxÞ ¼ 2:4x þ 1:51ð1  xÞ  1:8xð1  xÞ

ð5Þ

This differs from the expression given in Ref. [40], which underestimates the energy gap of CdS. Equation (5) is strongly bowed, with a band gap energy less than that of CdTe for 0 > x 0.51 and a minimum value at x ¼ 0.3 of 1.4 eV, corresponding to a wavelength of 884 nm. Interdiffusion may be expected to modify QE(l). Compared with the absorption spectrum of a pure CdTe film (band edge at l 820 nm), a small amount of S diffusion will cause a red shift in the long-wavelength end of the QE(l) response. For CdTe1–xSx with x 0.05, that is, close to the solubility limit [29], the band gap would be reduced to 1.47 eV (l 842 nm); this has been widely reported [29, 34, 41]. There is also a corresponding shift to longer wavelengths at the blue end of the spectral response, as a consequence of the diffusion of CdTe into CdS. This is generally held to be undesirable [11], since it reduces transmission into the absorbing CdTe. Absorption in the CdS side of the junction is not efficient, and it is better if the incident light is absorbed in the CdTe; this dictates that the CdS layer be as thin as possible, and higher

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values of JSC are indeed obtained [37, 42] for thinner layers of CdS. However, excessive shunting will result if the CdS is too thin due to pinhole formation.

4. FABRICATION OF CELLS One of the singular advantages of the CdS/CdTe solar cell, is that it may be produced by a wide variety of low cost and scaleable techniques [43, 44], including physical vapour deposition and chemical bath methods well suited to industrial scale manufacture of modules with m2 dimensions. Other deposition processes that could be used for module production, for example, atomic layer deposition, spray pyrolysis and screen printing etc., have not so far proved useful, either because of cost or the poor performance of the resulting modules.

4.1. Deposition of the absorber CdTe layer The CdTe layers in the most efficient cells [14–16] have all been deposited by close-spaced sublimation (CSS). Originally developed by Nicoll [45] over 40 years ago for the heteroepitaxial growth of GaAs on Ge, CSS is a physical vapour deposition process, where, as the name implies, the gap between the sublimating source and substrate is very small; just a few millimetre. At this proximity, the substrate will necessarily be at a relatively high temperature given that dissociative sublimation of CdTe takes place at temperatures above 500  C in vacuum. High-performance cells are typically deposited at substrate and source temperatures in the range 400–700  C in a vacuum of 102 to 5  103 Pa, conditions under which sublimation is diffusion limited [16]. It is a comparatively rapid process with growth rates in the order of 1–2 mm/min. The as-deposited grain size in CSS films is usually 3–5 mm somewhat larger than is the case with other growth techniques. Layers grown at higher temperatures are dense with a pronounced {111} orientation [37]. Cadmium telluride may also be deposited by electroplating [2, 46–48]. Although once the preferred technique, it is less common now as CSS has consistently proved to give superior devices. The CdTe is cathodically deposited from aqueous solutions containing Cd and Te ions in concentrations of 104 mol/l. Plating is carried out at temperatures of 80–90  C and the pH is typically adjusted to 2. Electrodeposited CdTe layers are n-type and resistive, and must be type converted to form a proper p-n junction. This is typically done by an empirical procedure where the cells are heated in air at 400  C for a few minutes. It is frequently carried out simultaneously with the Cl treatment discussed below, resulting in some confusion in the literature between the respective roles of the Cl and air

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heating processes. Grain sizes in as-grown electroplated CdTe layers are relatively small, <1.5 mm, with a columnar microstructure.

4.2. Deposition of the CdS window layer Chemical bath deposition (CBD) and CSS are generally used for the CdS layer, although in the Aramoto cell [15], the CdS was deposited by metal organic chemical vapour deposition (MOCVD), probably not well suited to large area production in spite of the authors’ claims. CBD of CdS [49] entails the heterogeneous reaction of thiourea (CS(NH2)2) and a Cd salt in a heated (60–90  C) basic aqueous solution [50, 51]. Ammonia is commonly used to control the pH of the solution, although its volatility is a problem and less volatile alternatives such as ethylenediamene are sometimes used instead [51, 52]. Growth rates tend to be low, 1–20 nm/min, and grain sizes very small, 15–100 nm. The as-deposited CdS layer is randomly oriented and often reported as being in the meta-stable cubic sphalerite rather than the thermodynamically stable hexagonal wurtzite phase, or sometimes a mixture of phases and polytypes [53].

4.3. Post-growth annealing in chlorine Efficient cells invariably require some form of post-deposition heat treatment or activation, generally by heating the layer in air after exposure to CdCl2 [11, 54, 55]. There are essentially three techniques for introducing the CdCl2: (i) to deposit a layer by briefly immersing the cell in a solution of CdCl2 in methanol before heating it for 20 min in air at 400  C; (ii) to deposit the CdCl2 layer (60–100 nm thick) directly by vacuum evaporation followed by the same heating step as (i); (iii) to heat the CdTe in CdCl2 vapour at the annealing temperature. Note the latter does not apparently involve air/oxygen, although residual O2 may be present following the deposition process. The CdCl2 process is not fully understood, but comparative studies of treated and untreated cells have confirmed that its use is critical [42, 56, 57]. In these studies the Cl treatment was found to increase the efficiency from 1.5–10% in nominally identical cells. Heating in the absence of Cl resulted in a small increase of about a factor of two to 3%; it is noteworthy that the most efficient cells [14–16] were all subjected to a CdCl2 procedure. The treatment does increase grain size in those cases where the as-deposited grain size is small (i.e. in electro-plated cells), is often accompanied by some degree of recrystallisation, and promotes interdiffusion. It appears to be a necessary step even where the CdTe has been deposited by CSS, that is, at high temperature, when inter-diffusion is inevitable [37] and grain sizes are already large.

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The microstructure of as-deposited CSS-grown CdTe is normally columnar with a pronounced {111} texture. Recrystallisation disrupts this, producing a more random structure [42] that results in higher currents. Durose et al. [11] have suggested that this somewhat counterintuitive observation could arise because the incidence of twinning parallel to the interface is greater in a columnar microstructure and lamellar twins of this type may be more prejudicial to the current flow. The CdCl2 treatment has also been reported to promote inter-diffusion, and some have suggested that it is a necessary requirement [42]. The CdCl2 process changes the principal dark current transport mechanism from interface tunnelling/recombination to space charge recombination [57, 58] implying a reduction in the density of interface states or possibly grain boundary states. Heating in the absence of Cl had no effect. In addition to changing the principal dark current transport mechanisms, the Cl treatment modifies the density and nature of the deep defect levels. Deep level transient spectroscopy (DLTS) of CdCl2 processed cells [57] indicated that the use of Cl introduced a deep level at EV þ 0.64 eV, tentatively assigned either to a doubly ionised Cd vacancy (VCd)2 or more probably a singly ionised Cd–Cl complex (VCdCl) . In a systematic DLTS study intended to elucidate the effects of post-deposition processing on the distribution of deep levels, Lourenc¸o et al. [59] found the distribution to be dominated by a hole trap at 0.46–0.49 eV above the valence band, which was present in all samples, irrespective of the post-deposition processing, although the concentration was much greater in the CdCl2-treated cells. Post-deposition processing was found to introduce a number of other metastable levels, with properties that depended to some extent on previous sample history. A comparative DLTS and admittance study by Versluys et al. [60] of the effects on deep levels of carrying out the CdCl2 activation in two different ways; comparing CdCl2 activation in vacuum and in air, also found a level with an activation of 0.45 eV, which they assumed to be a barrier at the back contact. Additionally, they observed a range of shallow and deep levels in both air and vacuum annealed cells, including what appeared to be the chlorine A-centre (VCd–ClTe) at EV þ 0.113 eV in air-annealed cells and a similar, though less pronounced, level in the vacuum-annealed devices. Optical beam induced current (OBIC) studies have also shown that the CdCl2 treatment improves uniformity [61]. In untreated cells, the majority of the cell area was inactive at the wavelength of the He-Ne laser used for the investigation, whereas CdCl2 treated devices displayed almost no visible structure, suggesting a high degree of uniformity. Replacing the laser by a well focused (<100 mm2) monochromated light source indicated that a good window response was only observed for CdCl2 treated cells. When the light was directed onto poor regions of the cell, the response peaked at the CdTe band gap and decreased with decreasing wavelength, behaviour indicative of a buried junction.

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4.4. Back contacts As discussed above, the fabrication of an injecting, low resistance and durable back contact to CdTe has proved difficult. The ensuing problems of rollover and crossover in the J–V characteristics have already been discussed at length. It is important that the barrier height and width should not cause the localisation of carriers, since this would result in excessive back surface recombination. In practice, this places an upper limit on the barrier height of 0.4 eV, unless the barrier is particularly thin. Contact regimes fall into two generic groups (see review by Fahrenbruch [20]): CdTe-metal or more realistically CdTe-dipole-metal types; and a group where some appropriate intermediate or buffer layer is incorporated between the CdTe and the metal to form a heterogeneous contact. Whatever the contact regime the CdTe surface is invariable etched, typically using a Br/methanol solution or the so-called N–P etch (an aqueous solution of HNO3 and H3PO4) both of which preferentially remove Cd leaving a conducting Te-rich surface. The etching step is crucial; not only is the Te layer formed over the surfaces of grains, it is also produced along the grain boundaries where it creates a conductive bridge between grains, increasing lateral conductivity [62]. However, the NP etch is aggressive and excessive etching will widen the gaps between grains. A brief etch of a few seconds in a stronger NP solution appears to be preferable to a longer etch in a weaker solution. In the latter case, grain boundary broadening appears to be more significant allowing diffusion related shunt leakage paths to develop with a consequential loss of performance [62]. The impact of Fermi level pinning at the CdTe–metal interface on the barrier height is clearly demonstrated in the study undertaken by Dharmadasa et al. [63]. They measured the electrical properties of a large number of metal contacts (to single crystal (110) n-CdTe) and with the exceptions of Mn, Cr and V, found a constant barrier height of 0.72 eV independently of the metal. A barrier of this magnitude is too great for efficient cell performance, and apart from cases where the objective of the study was to model contact barrier effects explicitly [22, 25], simple metal-CdTe contacts, for example, evaporated Au [47] are seldom used. More commonly, the surface is heavily pre-doped with an acceptor such as Cu to lower the resistivity before depositing the contact metal, for example, Au, graphite, Mo, Ni [54], although the predisposition to self compensation in CdTe makes it difficult to achieve the necessary levels of doping ( 1018 cm3). The second generic group offer greater flexibility, but at the cost of an additional junction and potentially another barrier. However, the judicious choice of materials can lead to an optimum situation where the buffer layer is highly or even degenerately doped minimising any barrier to the outer metal. XPS and related techniques have demonstrated that on

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exposure to air, the CdTe surface will be oxidised to TeO2 [64, 65], but the oxide layer is very thin (it is readily removed by briefly sputtering in UHV) and would not seriously impede quantum mechanical tunnelling of carriers. The elemental Te layer is several nm thick with a valence band offset, DVB, between Te and CdTe reported by Niles et al. [64] to be 0.26 and 0.5 eV by Kraft et al. [65]. As observed by Kraft, this determines the lower limit for any back contact involving Te. The most commonly used contacts are based on the use of Cu. A Te-rich surface is prepared as above, followed by the application of a Cu-loaded graphite paste [37]. A brief heat treatment diffuses the Cu into the CdTe surface where it reacts with the Te to form non-stoichiometric Cu2–xTe (x 2); a degenerate semiconductor (due to the excessive density of Cu vacancies, 1021 cm3 [66]) with a band gap of 1.04 eV. A layer of Cu2–xTe may also be deposited by co-evaporation of Cu and Te [67] or elemental Cu may simply be deposited by evaporation [68]. Surface science studies [67] of Cu-based contacts have shown that the band alignment between CdTe and Cu differs from that between CdTe and Cu2–xTe, with corresponding barrier heights of 1.0 eV and 0.8 eV, respectively. Notwithstanding such large barriers, efficient contacts can be readily made using Cu2–xTe. This paradox may reflect the fact that UPS and XPS measurements made in UHV are not representative of contacts made in ordinary ambient. Alternatively, a very high density of surface doping is in fact achieved in these contacts, with the result that the barrier is sufficiently narrow to allow tunnelling. Wu et al. [69] have shown that control of the Cu2–xTe phase is critical. They noted that although the resistivity of the chalcopyrite phase (x ¼ 0) can be as low as 104 O cm, this does not always make the best contact. The most efficient of their test cells used a mixed phase consisting of CuTe þ Cu1.4Te which gave an efficiency of 12.9%. The degenerate character of the Cu2–xTe means that a great variety of metals may be used as the outer contact, since barriers at this interface will be particularly narrow. Although widely used, Cu2–xTe contacts are unfortunately short lived. Copper diffuses into the semiconductor, depleting the contact and thus increasing its resistance, while simultaneously decreasing the effectiveness of the main junction [20]. Copper will in a relatively short period of time diffuse through the entire CdTe, creating defect states near the p-n junction [69] with a consequential reduction in the photo-generated carrier life-time. Similar results have been found with HgTe and ZnTe:Cu [70] based contacts. Alone among the possible contenders, ZnTe has a favourable valence band alignment with CdTe, with a negligible DVB of only 0.05 eV [71] and when doped with N, is potentially a good Cu-free contact with high doping levels, 1020 cm3 [72]. Baron [73] has suggested that the doping efficiency of N is higher in ZnTe than in CdTe, because in the tetrahedral

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cell, NTe is more stable when surrounded by Zn atoms than by Cd atoms. They quote X-ray diffraction data which suggest the Zn–NTe bond length ˚ close to metal–nitrogen bond lengths of 2.1 A ˚ . Nitrogen, therefore, is 2.2 A should not diffuse so rapidly as Cu into CdTe. However, ZnTe and CdTe are miscible across their compositional range and inter-diffusion at the CdTe/ZnTe interface to form a graded junction is to be expected. The extent of any such inter-diffusion and its potential effects (positive or otherwise) has not been systematically studied. Amin and co-workers [72] investigated the effects of a Cd1–xZnxTe layer at the back contact of a thin film of CdTe, but this was for x ¼ 0.5, larger than would be expected for any cross-diffusion, moreover the Cd1–xZnxTe layers were Cu doped. Nitrogen-doped ZnTe layers can be grown by vacuum evaporation of ZnTe using a RF plasma source for the N-doping [70–72]. The thickness of ZnTe:N is kept below 1 mm. Although the work function of ZnTe is 5.27 eV [70] and thus a high work function material such as Au or C must be used as the outer contact, this is not such an important issue as it is with CdTe, due to the high doping levels. The metal/ZnTe:N barrier is therefore very narrow and carriers can tunnel through the potential barrier relatively easily. Evidently, back contacts must be stable over periods of many years if CdS/CdTe modules are to be viable sources of solar energy. Antimonybased contacts have so far proved to be among the most stable. A 100 nm thick film of Sb is deposited onto freshly etched CdTe (NP etch) either by vacuum evaporation [62] or as a sputtered film of Sb2Te3 [74] grown at 150  C. Metallisation is provided by Au or Mo preferably deposited without breaking vacuum. Using Sb2Te3/Mo contact efficiencies of 12.5% [74] were achieved, although other researchers have found only modest performance [20]. Antimony telluride is a small band gap semiconductor ( 0.3 eV), but since the electron affinity is not known the valence band line up cannot be estimated. Current–voltage characteristics typically show roll over [62] suggesting that there is a barrier in the order of 0.4 eV at minimum. Ba¨tzner et al. [62] have suggested the better stability of Sb-based contacts (as compared with Cu-based contacts), may be ˚ for Cu and related to the relative sizes of the atomic radii – 1.28 and 1.45 A Sb, respectively, the smaller atom being able to diffuse more rapidly.

4.5. Transparent conducting oxide front contacts In comparison with, say the problems encountered with the back contact, the TCO/CdS interface is not widely regarded as being important as a limiting factor in cell performance [11, 75]. There are, nevertheless, some issues to be considered. Conventionally, SnOx often doped with F, has been used as the TCO in CdS/CdTe cells. It is a wide band gap (between 3.6 and 4.3 eV), direct

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gap semiconductor which is strongly n-type or even degenerate, due to oxygen vacancies. Photoemission studies [75] imply that the conduction band minimum in the SnOx coincides with that in the CdS, making it a low-resistance interface. In theory SnOx should be an ideal TCO, but the comparatively high sheet resistance 10 O/□ and modest optical transmission of 80%, represent some loss of efficiency. With a resistivity r 2  104 O cm (some 3–4 times lower than SnOx) and a greater transmission 80% over the relevant spectral range, ITO ((In2O3)0.9(SnO2)0.1) gives better performance. A variety of doped In2O3 materials have been suggested including In2O3:Ge [76, 77] and In2O3:F [76], both of which are comparable to ITO – r 2  104 O cm and transmission 85%. Although the performance of ITO and related materials is quite acceptable, they are expensive, due to the high price of In. This has inspired a search for alternative In-free TCOs, of which the most promising are Cd2SnO4 (CTO) and Zn2SnO4 (ZTO) [78]. Cadmium stannate films have resistivities (r 1.5  104 O cm) some 2–6 times lower than SnOx, a peak transmission > 90% at l ¼ 500 nm and are an order of magnitude more smooth. Zinc stannate (ZTO) is much more resistive (r 102 O cm) [78], but offers several advantages when used as an integrated buffer layer between a low resistive TCO (i.e. CTO) and CdS. During the process of fabricating cells, cross-diffusion will take place at both the CdTe/CdS and ZTO/CdS interfaces, consuming CdS from both sides in the process. Zinc cadmium sulphide has a larger band gap than CdS, moving the short wavelength cut-off deeper into the blue, increasing JSC as more of the available spectrum is transmitted through the window layer into the absorbing CdTe. In addition, this allows the initial deposition of thicker CdS films, reducing the incidence of pinholes and short circuits. Zinc stannate also improves layer adhesion, a particular problem with the CdCl2 process where extended treatment can lead to blistering and delamination [78]. Diffusion between CdS and ZTO appears to relieve the interfacial stress, allowing greater latitude in optimising the CdCl2 process.

4.6. Alternative structures 4.6.1. Lightweight flexible substrates For space applications, CdS/CdTe cells fabricated onto glass substrates offer no weight advantage compared to high-performance single-crystal Si or GaAs cells, and although these are expensive and susceptible to radiation damage, traditionally they have been used to provide the electrical power. In low earth orbits, cell degradation due to radiation damage is not a significant problem, but to provide global coverage for, say communications, would require a large number of satellites. Far fewer would be needed in high Earth orbits, but here cell degradation is a problem and the design of the solar panels becomes complicated [79].

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However, the production of high specific power (kW/kg) thin film modules on thin metal film foils, if it can be achieved, would bring considerable benefits. Since CdS/CdTe cells are only 10 mm thick the support foils need only be 50 mm thick, and support structures can be kept to the minimum. In addition, unlike single crystal devices, the CdS/CdTe cell is already quite highly defected (due to lattice and thermal expansion coefficient mismatch etc.), and in consequence it is expected that it should be less susceptible to radiation damage. Finally, foil mounted cells maybe folded (in principle) into any shape. There are in essence two principal device configurations for a foil mounted cell: with the CdTe adjacent to the foil (referred to as ‘back wall’) or with the CdS next to the foil (front wall). Front wall cells require a light transmitting contact to the CdTe, and at that time no satisfactory transparent contacts to p-CdTe were known. Although, a grid contact structure could have been used, the high sheet resistance of CdTe meant the grid would need to be dense and inefficient. Of course, back wall devices also suffer from the problems of making reliable contacts to p-CdTe, as discussed earlier. Mathew et al. [80] have reviewed the deposition of CdTe by a variety of techniques including CSS onto several metal films, including stainless steel (SS), Mo, Ni and Cu. Molybdenum was considered to be a potentially good substrate due to the close match in thermal expansion coefficient with CdTe – important for the high temperatures involved in CSS. However, energy band analysis indicated that the Mo/p–CdTe contact would be rectifying and the CdTe at the interface would need to be heavily doped to promote efficient quantum tunnelling through the barrier (see Section 4.4). Virtually all the studies reviewed related to metal/n-CdTe Schottky devices and were aimed at determining charge transport processes, trapping levels and barrier heights etc. With the exception of the study by McClure et al. [79], little work had been reported, at that time, on the fabrication of heterojunction devices on flexible substrates. McClure et al. focused on the back wall configuration, because of the problems with the front wall design. Initially, they used low carbon steel substrates, as iron has a relatively high work function. The films adhered well with a low contact resistance, but proved to be n-type, probably due to Fe3þ substituting for Cd2þ. They also studied Ni and Mo. Of the common metals Ni has one of the highest work functions, but they found there were adhesion problems following annealing. Ultimately Mo proved the most successful. To circumvent the contact problems, 50 nm layers of Cu and Te were deposited in sequence to produce an inter-layer between the CdTe and the Mo. This procedure gave contact resistances of less than 1 O cm2 due, it was assumed, to the formation of a thin CuxTe layer at the interface. A thin 0.5 mm top layer of CdS was deposited by

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evaporation, but subsequent annealing steps resulted in the inter-diffusion of S and Te to form a p-type layer of CdSxTe1–x. Since this had consumed most of the n-CdS the device did not function correctly and another layer of CdS had to be deposited after the anneal. The structure was completed with reactively sputtered ITO giving a sheet resistance of 10 O/□. Although only a few cells were fabricated and tested, they reported values of VOC and JSC of 580 mV and 12 mA/cm2, respectively. Using ‘lift-off’ processes Romeo and co-workers [81] have recently developed efficient cells on flexible polymer substrates. They have been able to produce cells in both the superstrate and substrate configurations, by supporting the thin films on conventional soda-lime glass through the deposition and annealing stages. Although most transparent polymers are not stable at the high temperatures (450–550  C), some polyimides retain sufficient transparency for use in solar cells. Initially, a thin buffer layer of NaCl was evaporated onto the glass. A 10 mm film of polyimide, which had been developed in-house, was spin-coated onto this and after curing at 430  C was coated by a layer of ITO by RF magnetron sputtering. CdS and CdTe layers were then deposited and activated by CdCl2 by their standard process. When complete, the cell was rinsed in water to dissolve the NaCl and separate the device on its flexible polymer substrate from the glass carrier. These cells yielded high efficiencies of 11%, (VOC ¼ 842 mV, JSC ¼ 18.5 mA/cm2, FF ¼ 70.9%). The relatively low JSC was due to absorption in the polyimide. Clearly a practical manufacturing process will require a low cost, commercially available polymer, and after evaluating a number of commercially available polyimide films, an UpilexTM 10 mm thick film was found to be sufficiently transparent for use in CdS/CdTe solar cells. Cells produced on the Upilex films were marginally more efficient, 11.4%, (VOC ¼ 765 mV, JSC ¼ 20.9 mA/cm2, FF ¼ 70.9%) due probably to a slightly increased transparency. They also studied the inverse ‘substrate’ configuration, where the TCO was deposited onto the NaCl. After removing the cell from the glass support by rinsing in water as before, the cells proved to be less efficient. Two TCOs, ZnO:Al and FTO, were investigated and exhibited efficiencies of 6% and 7.3%, respectively. In both cases, the FFs were very low as were the open circuit voltages. However, the short circuit currents were greater than for the superstrate case, since loss due to absorption in the polyimide (now at the back of the device) was reduced. The high efficiency polyimide-based flexible cells demonstrated a specific power potential of 2 kW/kg. If the celllevel laboratory-based process can be transferred to the factory, use of commercially available UpilexTM is an attractive option for an in-line manufacturing process, while replacing the polyimide with a suitable metal foil, for example Mo, makes this a promising contender for civil space applications.

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4.6.2. Bifacial CdTe solar cell It has been suggested that with the use of multi-junction configurations, the next generation of CdTe-based cells could realistically have efficiencies of 25% [82]. Such devices can be either mechanically stacked (three or four terminals) or two terminal monolithic structures. In addition to the usual transparent conducting top contact, both arrangements require a conducting, transparent back contact to allow the transmission of light to the underlying cell; i.e., a bifacial configuration. The principal challenge is to produce a good transparent contact to the CdTe. Simply sputtering ITO onto the CdTe (after a Br/methanol etch) produced a transparent contact (>85%) with a low sheet resistance of 10 O/□ [83]. However, as-deposited cells proved to be inefficient,  2.5%, though annealing in air at 350  C and after light soaking the efficiency could be more than doubled. The deposition of a thin Cu layer just before sputtering the ITO improved the efficiency, but accelerated lifetime tests produced results that depended on the thickness of the Cu. Cells with 3 nm Cu inter-layer degraded over a four year equivalent time, but were comparatively stable thereafter, while the performance of cells with <0.5 nm Cu improved slightly with accelerated ageing. Cells with thicker, 2.5 mm, absorber (CdTe) layers performed better when illuminated from the front ( ¼ 10.3%) but less well under back illumination ( ¼ 2.1%). When the absorber layer was reduced to 1 mm the efficiency under front illumination was reduced to  ¼ 8.6%, but under back illumination was increased to 3.2% largely as a result of an increase in JSC. Marsilac et al. [84] investigated even thinner cells, CdS ¼ 0.13 mm and CdTe ¼ 0.68 mm. These ultra-thin cells followed the same general trend, with efficiencies of 5.7% and 5% when illuminated from the glass and back contact sides, respectively. The cells were fabricated on SnOx: F-coated soda lime glass with a reactively sputtered ZnTe:N and ITO back contacts to produce a glass/SnOx:F/CdS/CdTe/ZnTe:N/ITO structure. The external QE of front and back contact spectra were quite different. The profile of the former was similar to that of thicker cells, but with a reduction in QE from 600–800 nm due to the reduced absorption in the thin CdTe. Light incident from the back contact side must first pass through the ITO and the ZnTe:N which has a band gap of 2.2 eV, and although it was very thin (70 nm), the QE was less than expected at l 560 nm, implying there was little or no carrier collection at these wavelengths. It was suggested that at these wavelengths, photons are absorbed near the back contact where the electric field is weak and carrier separation and collection is as a result poor. At wavelengths in the range 400– 800 nm, photons are increasingly absorbed near the principal CdTe/CdS junction, leading to a progressive increase in QE with increasing wavelength up to the band edge.

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5. MANUFACTURE OF CdS/CdTe MODULES Commercial high-volume module production requires the conversion of laboratory-based batch processes into continuous or at least semicontinuous assembly lines. Transfer of laboratory expertise and experience to the factory is not merely a change of scale, but a paradigm shift. The magnitude of the necessary investment dictates radically different perspectives. Questions of yield and price transcend the search for ultimate performance. Birkmire and Eser [43], in their review of polycrystalline cells, outlined the challenge facing manufacturers of thin film polycrystalline modules. They noted that a factory manufacturing 10 MWp (i.e. the maximum power the module can deliver under 100 mW/cm2, AM1.5 insolation) CdTe modules (1.2  0.6 m2) per annum, assuming 100% yield and 83% process uptime would need to produce one module every 2½ min, working 3 shifts/day, 250 days/year. Moreover, this must be achieved at a cost that the market will accept. In spite of the challenges, a number of companies are beginning to manufacture CdS/CdTe modules on the MW per annum scale. By far the largest, First Solar [85], was created in 1999 with the purchase of Solar Cells, Inc. (SCI) and started production of commercial modules in 2002 with a 25 MW plant at its Perrysburg, Ohio facility. By the end of 2007 production capacity had been expanded to 307 MWp with increased production at its Ohio base and four more lines in Germany. Capacity is planned to reach 1 GWp by the end of 2009 with the construction of additional plants in Malaysia. While the module design is essentially a standard CdS/CdTe superstrate configuration, First Solar is particularly secretive about their manufacture process and very little is known about it. However, it is believed to be a refined version of that used by its predecessor SCI [86]. The SCI manufacturing process is described in an annual subcontract report to NREL in 1993 [87]. In the SCI process SiO2/ SnOx-coated glass is passed in sequence through a series of four heated vacuum chambers on ceramic rollers. After the glass superstrate has been loaded into the first chamber, the entry valve is closed and the chamber is evacuated. While it is being pumped down, the superstrate is radiatively heated to 600  C before it is transferred to the second chamber for deposition of the CdS layer. This is carried out in a single pass through elemental vapours from a crucible containing a powder source at 700  C located above the superstrate. Control is exercised by admitting nitrogen into the chamber at a pressure of 1 torr. The superstrate is then passed into the CdTe zone after which it is subjected to what appears to be a standard Cl-based heat treatment at 400  C before being quenched using N2. Cadmium sulphide and cadmium telluride deposition typically take only 10 and 40 s, respectively. Nickel is used as the CdTe contact backed by a thicker layer of Al. Interconnects are formed by laser scribing and the

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module is completed by laminating a second sheet of glass on to the back using a layer or of ethyl vinyl acetate. The SCI process differs from the usual CSS methods, in at least two significant respects: the superstrate is pre-heated before it is admitted into the deposition zones and secondly the crucibles containing the source powders are located above the superstrate and face away from it. At 600  C the glass is close to its softening temperature and this arrangement will allow it to be supported by more closely spaced rollers. In an enterprising innovation First Solar developed the first prefunded module collection and recycling programme in the PV industry. Typically modules can deliver 75Wp at a conversion efficiency of 10.6%, with a manufacturing cost of $1.14 per Wp, though to compete against conventional fossil fuels, manufacturing costs will have to be reduced to $0.7 per Wp [86]. The principal European manufacturer is ANTEC Solar GmbH, who have established a line intended to manufacture [88] 100,000 m2 of modules per annum (standard size 60  120 cm2), corresponding to a power generating capacity of 10 MWp. The plant, which has been in operation since 2002, is based on the CSS deposition of both CdS and CdTe in an automatic production line, giving a high throughput of about 1 m/min. The projected net saving in CO2 is 16,000 kg/m2 over a 30 year life span.

6. DEGRADATION MECHANISMS Clearly for CdS/CdTe modules to be practical, they must be stable in the field and under load for many years. Laboratory-based studies of individual solar cells tend to focus on improving cell performance and acquiring an understanding of the device physics. Apart from being time consuming, field trials involve other extrinsic factors such as the integrity of packaging, temperature cycling, humidity etc. not normally carried out in laboratory-based studies. It is important therefore, to distinguish between external influences and inherent ageing effects [89]. Due to obvious time constraints, degradation studies are usually carried out in accelerated conditions where the unit under test is maintained in a climate chamber at high levels of temperature and humidity (e.g. 85  C and 85%, respectively) under 1 sun (or greater) illumination and in open circuit. More sophisticated trials utilise a range of stressing protocols that aim to mimic in some degree actual use, for example, cycling between maximum power, open circuit and short circuit conditions [90]. The loss of performance is generally manifested as a reduction in VOC and FF (Eq. (4)), indicative of an increase in series resistance and/or the development of a blocking contact [68]. Usually, cells and modules stressed under open circuit or in reverse bias show greater power loss

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than those tested under short circuit or at maximum power [90]. This observation implies that degradation processes originate in material and behavioural changes within the active components of modules. The principal cause of instability is probably the back contact, and there is considerable evidence for the diffusion of back contact materials under heat and near junction field stressing [62]. In a comparative study, Ba¨tzner et al. [74] demonstrated that even for rather short periods of heat stress, 200  C for 30 min in air, considerable loss of performance was observed for all contact regimes except those based on Sb or Sb2Te3 buffer layers and Mo metallisation. Secondary ion mass spectroscopy analysis of cells that had been subjected to more rigorous accelerated life time tests, where the cells were light soaked at 65  C under 1 sun in open circuit conditions in air, revealed that Cu and Au had diffused through the CdTe and accumulated within the CdS and at the CdS/TCO interface. There was also some Sb, but this was an order of magnitude less and only observed from Mo/Sb bi-layer contacts. No Sb was detected from Mo/Sb2Te3 contacts, nor was any Mo. A blocking contact can also develop due to oxidation of the back contact resulting in an insulating layer at the back surface of the CdTe [91] again leading to roll over in the J–V characteristics and a reduction in the FF. The oxidation process will be accelerated in conditions of high humidity, emphasising the need for proper encapsulation of modules. Outdoor field trials [92] of sixteen standard modules (1.2  0.6 m2) with Sb-based contacts showed an average of 12–13% relative degradation in power output after 18 months outdoor exposure. This was following an initial increase of 4% over the first 5 months, which correlated with a rise in ISC over the same period before saturating. The loss of output thereafter was a consequence of the progressive reduction in VOC and FF. Small area samples were cut from modules that had suffered more than 10% degradation in an attempt to distinguish the role of interconnects in the loss of output. Detailed measurement and analysis indicated that the interconnects did not contribute significantly to the degradation and it was concluded that the reduction in maximum power was partly due to changes in the cells and partly to the front contact TCO. This is consistent with the findings of Ba¨tzner referred to earlier.

7. USE OF CdS/CdTe MODULES IN LARGE-SCALE POWER GENERATION With the exception of large hydroelectric systems, renewable sources of energy do not readily integrate into large grid-based distribution networks. The electrical grid was developed in response to the need for reliable power; if one station went out of commission the supply to users

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was not interrupted. If wide-scale implementation of photovoltaic (PV) power is to be practical, then ways must be found to scale up from the few tens of Watts typical of individual modules by orders of magnitude to tens or hundreds of MW at a single location that is comparable to a medium scale thermal or nuclear facility. Another problem with PV power generation is the variability of output power. There is obviously a diurnal variation, but that can be managed. Much more difficult is the short-term variation due to, for example, the passage of clouds, and this dictates that ideally any large-scale PV station should be sited in a desert. More than that, sandy deserts are not really suitable, due to the damaging effects of sandstorms; only rocky or gravel deserts are suitable. There are in principle, three important parameters. The energy payback time (EPT), the net CO2 emission rate (gm C/kWh) and the generating cost (price per kWh). The EPT is simply the time taken to recover the total energy investment in building and running the station over the life cycle using its own net energy production. The net CO2 emission rate is the ratio of the total CO2 emission released by the station over its life span to the total energy generation over the lifetime. It is a useful index in determining the effectiveness of the plant’s contribution to global warming. It is obvious that the generation cost must not be unduly out of line with other generating systems. An important factor which must be included in any cost analysis is the ‘balance of system’ (BOS) costs; that is, the additional expenditure in construction of the facility, connection to the grid, maintenance, dismantling etc. Very large-scale PV stations ( 100 MW) have yet to be built, but there have been a limited number of smaller, but still substantial stations delivering up to a few tens of megawatts, such as the Springfield, Arizona plant [93]. This has an installed PV capacity of 4.6 MWp in a mix of multicrystalline (mc) Si (3.5 MWp) and thin film (1.1 MWp) modules. The output is used to power the water pumps at the Springfield Coal Fired station (10 MWac-al), and when the water pumps are not operating the capacity is fed into the grid for general use. In the case of the Springfield station, the EPT was 0.21 years for the BOS, which equates to about 0.37 years for average US conditions. The life cycle green house gas emissions (BOS) are estimated to be 29–31 kg CO2/m2. Larger facilities, for example First Solar’s 40 MW Brandis solar farm in Germany, are beginning to be built, usually with government subsidised programs, as in Germany [86]. Currently First Solar is only selling panels for use in solar farms and commercial rooftop installations. It is not selling to the public at large as all its output is being used to meet the larger scale commercial demands. There have been a few simulation studies of larger stations, for example, Ito 100 MW [94]. This was a comparative study of all the main solar cell systems based on current performance, and included all equipment, transport, operational, maintenance costs as well as transmission losses

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over a distance of 100 km. For CdTe modules the EPT and CO2 emission rates were estimated to be 1.9 years and 12.8 g-C/kWh, respectively. Corresponding values for Si modules were 1.5 years and 9.4 g-C/kWh. An important conclusion was that module efficiency was important for energy production and CO2 reduction, because the build costs – foundations, cabling, troughs etc. – were reduced. This was more important in ‘cold’ deserts than in ‘hot’ deserts, for example, the Sahara, where thin film systems were a viable proposition due to their lower price ($/Watt).

8. CONCLUDING REMARKS The increasingly large volumes of Si required for Si wafer solar cells place Si-based PV in direct competition with conventional Si-microelectronics for semiconductor grade material, currently in short supply. There is as a result, an opportunity for thin film technologies in general and CdS/CdTe cells in particular to gain an increased market share [95]. That is not to say that there are not supply issues with thin film PV. The Cu-chalcogenide cells utilise indium a rare element, subject to competitive pressures from flat screen displays and LEDs. Tellurium is also classified as a scattered or precious element, and also has other competing uses in metal alloying and thermoelectric applications. Rare elements of this kind are only economically viable as a by product of some other large volume mineral extraction process, for example, Te is a by product of Cu-refining. Cadmium (which is plentiful) is a waste product from the mining and refining of Zn. An additional concern for the widespread deployment of CdS/CdTe solar cells is the toxicity of Cd. The module collection and recycling programme instituted by First Solar is an attempt to find a solution to the problem and pre-empt further legislation restricting the use of Cd, as is the case in the EU, where the sale of consumer products containing more than 0.1% by weight Cd was banned in 2006. PV modules are exempt for the present, but the exemption will be subject to review every four years. However, there is a strong argument that the sequestration of Cd from the waste of Zn mining into PV modules where it can be both used and controlled makes environmental sense. The toxicity of Cd and scarcity of Te both serve as drivers for more efficient manufacturing processes and the development of extraction technologies from CdTe/CdS scrap and spent modules [96]. In conclusion, thin film CdS/CdTe solar cells are emerging as a potentially viable source of PV power generation. There are outstanding questions to be resolved, not least: the scarcity of Te, public acceptability of Cd containing products and their long term durability. The increasing production and use of CdS/CdTe modules in the field may provide answers to these questions.

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CHAPTER

IIC Applications of CdTe, CdZnTe, and CdMnTe Radiation Detectors Ge Yang and R.B. James

1. INTRODUCTION For more than four decades, CdTe has been known as a good candidate material for radiation detection. Its high atomic number ensures effective radiation-atomic interactions, while delivering good sensitivity and energy resolution. Furthermore, CdTe detectors offer high conversion efficiency of photons to electronic charge carriers as compared to scintillators; consequently, for high-quality CdTe detectors, the energy resolution is expected to be considerably better than scintillators. More importantly, CdTe radiation detectors operate at room temperature, thereby obviating the need for a complicated cooling system and, accordingly, affording greatly enlarged application fields. However, CdTe sometimes displays polarization effect, wherein the counting rates or peak positions change with time [1]. A relatively high leakage current also has proven to be a limiting factor during the development of largevolume CdTe radiation detectors. In addition, CdTe detectors are relatively expensive, because it is difficult to grow large-volume CdTe crystals of acceptable quality and yield. On the basis of the utility of CdTe, CdZnTe (CZT) was explored carefully, because the introduction of Zn increases the band gap and reduces the leakage current (and noise) of radiation detectors. It also is less expensive and easier to produce large-volume single crystals with CZT than with CdTe. Since the first practical CZT gamma-ray detector reported in 1992 [2], there have been many advances in the performance Brookhaven National Laboratory, Upton, NY 11973, USA

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of the devices. The high-resistivity CZT crystals required for radiation detection initially were grown using the high-pressure Bridgman method (HPB). In recent years, other methods were employed, such as the travelling heater method (THM) and the modified Bridgman method (MB). Today, large-volume CZT single crystals up to hundreds of cubic centimeters can be produced due to improvements in the growth techniques [3]. At the same time, the performance-limiting factors of CZT detectors have been better clarified with modern characterization methods, such as synchrotron radiation X-ray mapping, time-of-flight secondary ion mass spectrometry (Tof-SIMS) technique, Pockels effect measurements, deep-level transient spectroscopy, X-ray diffraction methods, and others [4–6], thereby reciprocally helping CZT researchers enhance the availability of these large-volume crystals. Such progress in bettering the quality of the CZT material for CZT radiation detectors and electron-transport-only device designs, as well as improving the related electronics greatly have enlarged their application fields and accelerated their commercial marketing. CdMnTe (CMT) is a relatively novel material for radiation detection. Burger et al. reported the first investigation of its potential for such applications in 1999 [7]. They proposed that CMT has two main advantages over CZT, namely, better homogeneity and the lower amount of Mn needed to reach the desired band gap, so making it a good competitor to CZT. Most problems with CMT radiation detectors center on the poor quality of the crystals. Until now, it was difficult to obtain CMT single crystals with high resistivity and acceptable carrier transport properties. The limited availability of CMT crystals to researchers also inhibits its development as a radiation detector. As a result, they remain at the stage of laboratory development and have not been incorporated practically in commercial detection systems. However, some recent progress suggests the possibility of a breakthrough in this field, leading to large improvement in the performance of CMT radiation detectors [8, 9]. In this chapter, we focus primarily on the applications of CdTe and CZT radiation detectors. The nature of and future trends in CMT radiation detectors will also be discussed.

2. NATIONAL SECURITY AND NONPROLIFERATION INSPECTIONS In undertaking national security and nonproliferation inspections, radiation detectors mainly serve to locate and identify special radioactive nuclear materials. These applications especially demand portable, lightweight radiation detectors with high resolution. CdTe and CZT radiation detectors have proven well-suited for these purposes, because they are

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compact, robust, and with low maintenance and low-power consumption, while simultaneously providing good energy resolution and high detection efficiency per unit detector volume. The International Atomic Energy Agency (IAEA) has used CdTe and CZT detectors for over three decades [10, 11]. Recently, their usage significantly increased, stimulated by the improvement of material properties, their improved availability, and the development of improved detector-specific, low-noise integrated circuits for electronic readouts. The IAEA employs hemispheric CdTe/CZT detectors to verify irradiated nuclear material, that is, spent-fuel assemblies; generally, they must measure high count rates and high gamma energies. In such cases, a shielded and collimated detector normally must be placed very close to an irradiated item to obtain its specific signature because accessible space often is limited. Therein, the small size and high efficiency of CdTe/CZT detectors are essential prerequisites. A good example is the miniature CZT detection probe with its integrated preamplifier that the IAEA employs for light-water-reactor assemblies and collimators; the probes’ volumes range between about 6 and 60 mm3 and their resolution between 1% for 137Cs (7 keV for 662 keV) and about 3% (18 keV) [12]. In addition, the long-term stability of hemispheric CdTe/CZT detectors generally is very good; they have survived for several weeks after burn-in tests. For some applications, the detectors operated over months without noticeable changes in their spectral parameters. Large-volume CZT detectors also are needed in nonproliferation inspections, because they potentially reduce the measurement time significantly. This is especially important in verifying un-irradiated nuclear material where count rates are low and typical gamma lines of uranium or plutonium gamma spectra must be resolved. However, the large-volume CdTe/CZT detectors required for these inspections presently are limited by the lack of CZT single crystals whose active volume is more than about 500 mm3. The dearth of larger volume CdTe/CZT single crystals must be resolved before their widespread application in nonproliferation inspections. Recently, scientists at Brookhaven National Laboratory (BNL) developed a hand-held gamma-ray spectrometer for nonproliferation inspections based on virtual Frisch-grid CZT detectors. The whole system achieves an effective detection volume of 19.2 cm3, that is, 10 times larger than commercial co-planar grid (CPG) CZT detectors. Consequently, detection efficiency is improved significantly. The system employs an 8  8 virtual Frisch-grid CZT detector array (Fig. 1); each detector is 5  5  12 mm3. By using front-end application-specific integrated circuits (ASICs) developed at BNL, this spectrometer has a small profile and high energy-resolution. Further, its relatively simple configuration greatly lowers the cost. This achievement has allowed us to build an inexpensive,

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Figure 1

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Virtual Frisch-grid CZT radiation detectors developed at BNL.

large-volume detector array with high energy resolution and high detection efficiency, so affording wide application potentials in national security and nonproliferation inspections. Moreover, the detector modules are scalable to address a larger range of efficiency requirements.

3. MEDICAL IMAGING The unique advantages of CdTe and CZT detectors over other materials in medical-imaging applications rest on their long-term stability, direct digitization, and spectrometer-mode imaging, among other benefits. Especially, because these detectors are made from tiny individual pixels, they are relatively small, thereby ensuring high spatial resolution and good energy resolution. Therefore, the arrays of small CdTe/CZT detectors possess good image quality with low noise, as invariably needed in medical imaging to precisely localize possible lesions. Furthermore, the room temperature operation of CdTe/CZT detectors effectively simplifies the structure of medical-imaging equipment and reduces maintenance costs.

3.1. Gamma (g)-camera The research and development on CdTe-based gamma cameras have been largely determined by the progress of both low-noise dedicated integrated chips, as well as the pixelization of the detectors, like the “small pixel effect.” The first attempt to use a CdTe array for nuclear cardiac imaging goes back to 1991: a small probe developed by Scheiber et al. [13], including 12 independent 10  10 mm2 detectors. The first full imaging system was presented in 1996 by Eisen et al. [14], who developed

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a small field-of-view camera equipped with 40  32 detectors (each 4  4 mm2). Since then, great efforts have been undertaken by several groups worldwide to develop clinical systems using CdTe or CZT detectors. For example, CZT has been used in integrated pixel detector structures by Barber et al. [15], Polichar et al. [16], and Doty et al. [17]. Most of these earlier references are reported in a paper by Scheiber and Giakos [18]. The NUCAM mobile camera developed by Eisen et al. [14, 19] incorporating 1280 detectors is shown in Fig. 2. The contrast resolution due to scatter rejection was proven superior in the CdTe NUCAM camera, when compared to a scintillator-based Anger camera. Nearly during the same period, in the frame of a European program, a consortium “BIOMED II” developed a heart-devoted camera with 2304 pixels using dedicated I.C. low-noise chips [20]. Subsequently, extensive researches were devoted to attain the best compromise between spectrometric performance and detection efficiency to improve the feasibility of a g-camera based on CdTe/CZT. They included preventing the early recombination of the holes, optimizing the electrodes’ geometry and the size and shape of the detector elements, as well as assuring signal acquisition by the associated specific integrated circuits.

Figure 2 A photograph of the first version of a moveable gamma camera, NUCAM, based on large arrays of CdTe detectors. (Source: Ref. [19], reprinted with permission from Elsevier.)

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In 2003, Siemens reported the imaging performance of a prototype CZT g-camera system comprising 15 CZT modules, as depicted in Fig. 3 [21]. This camera had a 12  20 cm2 active area, comprising 3 columns each with 5 rows of modules. Each module had a square array of 16  16 pixels. The pixels’ pitch was 2.46 mm. Figure 4 compares the higher quality images of a brain phantom acquired with the CZT camera with those poorer ones taken on a standard commercial NaI(Tl) gamma camera.

Figure 3 Photograph of the CdZnTe prototype camera showing the 3  5 arrangement of the modules. The low-energy, ultra-high resolution (LEUHR) collimator is behind it. (Source: Ref. [21], reprinted with permission from Elsevier.)

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Collaborations between Verger et al. [22] and Mestais et al. [23] generated the design of a type of new CZT g-camera based on measuring both pulse height and fast rise time [22, 23]. This method brings the advantages of high scatter rejection while supporting high detection efficiency. In 2004, Verger et al. described their latest developed CZT g-camera and discussed the performance of a small CZT imager of 256 discrete detectors in an array of four platforms [24]. They compared and evaluated the planar imaging performance of this small CZT detector imager with that of a standard NaI(Tl) g-camera using two different phantoms. Figure 5 shows the detector. Figure 6 demonstrates that the CZT camera’s image is superior to the NaI(Tl) one for all cold rod diameters. Figure 7 also clearly reveals that the better detector energy resolution of CZT noticeably improves the image contrast in a high-scatter environment for the same system spatial resolution. In addition, Konstantinos et al. developed and tested the photoncounting CdTe g-camera shown in Fig. 8 [25]. It is built of eight individual detector hybrids, each consisting of a pixel CdTe detector of 22  11 mm2, solder bump bonded to a photon-counting, custom-designed ASIC. The effective pixel size (image pixel pitch) is 0.5 mm. The current full active imaging area of the CdTe g-camera covers 44  44 mm2. The camera operates both in the real-time imaging mode with a maximum speed of 100 frames/s, and in the accumulation mode with user-adjustable counting time; its dynamic range is 1:14,000,000. It exhibits excellent sensitivity.

Figure 5 Small 256 CZT detector imager. (Source: Ref. [24], reprinted with permission from, reprinted with permission from IEEE, # 2004 IEEE.)

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Figure 8 Photograph of the developed gamma/X-ray camera. (A) Detector board with eight detector hybrids. (B) Camera housing with the 0.5-mm pitch collimator (150-mm septal thickness, 350-mm openings, and 1:15 aspect ratio). (Source: Ref. [30], reprinted with permission from Elsevier.)

3.2. Digital mammography Digital mammography offers the potential for improving image quality and, subsequently, the possibility of better detecting breast cancer, particularly in women with dense breasts, where current screen-film mammography often is inadequate. Tu¨mer et al. developed hybrid pixel detector arrays with 50  50 mm2 pixel sizes for use in digital mammography with different detection materials [26]; CZT and CdTe pixel detectors gave the best results. The images from CZT and Si pixel detectors are compared in Fig. 9, which shows a finger phantom with an embedded human bone. Detailed bone structures are visible in both. Although the silicon detector’s thickness was 1 mm, much larger than that of CZT, the quality of the image from the latter is much higher, reflecting the lack of angle blurring and the higher detective quantum efficiency (DQE) achieved with the CZT pixel detector, even though only holes are collected. Further, the results from a standard mammography test and calibration phantom, shown in Fig. 10, also demonstrate the improved contrast

Figure 9 Images of a human finger phantom from a 0.15-mm thick CdZnTe detector (left) and a 1-mm thick Si detector (right). (Source: Ref. [26], reprinted with permission from Elsevier.)

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Figure 10 Comparison of partial phantom images from a breast model taken with a CdZnTe detector (left) with that of a commercial first-generation digital mammography unit (right). The phantom is the standard mammographic model RMI 156 with only the wax insert. (Source: Ref. [26], reprinted with permission from Elsevier.)

of CZT pixel detectors as compared to first-generation digital mammography systems. Additionally, Fig. 11 compares images of a small mosquito fish (Gambusia affinis), 20 mm long, taken with a 0.15 mm CZT detector (top) and a 0.15 mm CdTe detector (bottom). Both images clearly reveal its bone structure. The CdTe image seems of slightly poorer quality than the CZT one, which might be due to the polarization effect in the former material that causes a higher background noise.

3.3. X-ray computed tomography (CT) In 1989, Zelenina et al. demonstrated the advantages of using CdTe detectors for medical X-ray CT according to their calculations for four main types of statistical noises in CdTe metal-semiconductor-metal

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Figure 11 Two images of a 20-mm-long mosquito fish using a 0.15-mm thick CdZnTe detector (top), and a 0.15-mm thick CdTe detector (bottom). (Source: Ref. [26], reprinted with permission from Elsevier.)

(MSM) structures [27]. In 1992, a prototype tomography machine, able to scan 0.5-m-size high-density objects, was realized by Glasser et al. with 25 CdTe detectors (25  15  0.8 mm3) [28]. It produced good-quality 1024  1024 tomographic images. More recently, Sueki et al. developed a monochromatic X-ray CT using a photon-counting 256-channel CdTe array detector that offers several advantages [29]. First, there is no beam-hardening effect. Second, the CT value has a linear attenuation coefficient. Furthermore, a subtraction image can be obtained using dual monochromatic energy X-rays. The energy resolution of this CdTe X-ray CT system is quite adequate for its purposes. Figure 12 is a schematic diagram of this system; Fig. 13 is a CT image of the head phantom obtained with it. Collimator Target

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Figure 12 Schematic diagram of the monochromatic X-ray CT system. (Source: Ref. [29], reprinted with permission from Elsevier.)

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Figure 13 X-ray CT image of the head phantom obtained with the X-ray CT system. (Source: Ref. [29], reprinted with permission from Elsevier.)

In 2004, Konstantinos et al. also manufactured a real-time X-ray imaging sensor with high-resistivity p-type CdTe suitable for small-field CT and similar applications [30]. To reduce the dark current and to prevent afterglow, a serious problem in real-time imaging, they included a rectifying indium anode contact. Figure 14 depicts its structure along with a photograph of this imaging sensor. The pixel size (pitch) is 100 mm, and the number of pixels is 506  508. The CdTe crystal is 0.75 mm thick. The sensitivity and resolution of this CdTe X-ray imaging sensor is excellent. Figure 15 shows the X-ray images of a printed circuit board with mounted components acquired at a 70-kV tube voltage and 40-mA tube current. The board’s multilayered structure is clearly apparent. Because of the high sensitivity of the sensor voids in the balls of a ball grid, the array is visible even in single-frame (20 ms) magnified images. The National Institute of Radiological Sciences of Japan also developed a dual-energy X-ray CT with CdTe array [31] consisting of 64 elements, 0.8 mm wide by 5 mm high by 5 mm deep, which are aligned side by side at intervals of 0.1 mm. This CdTe array detector registers

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CdTe pixel detector Bump bonds

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the energy of incident photons as well as the photon number and incident position. Therefore, it is not necessary to tune a monochromator twice to produce two monochromatic X-rays. This method is extendable to an advanced approach using polychromatic X-rays produced by a conventional X-ray tube. Recent progress has proven that CZT has potential for use in making a combined imager for X-ray CT and single photon-emission computed tomography (SPECT). By contrast with using SPECT alone, the simultaneous measurement by X-ray CT and SPECT yields structural and functional correlations, and improves the quantification and localization of the radionuclide. However, fabricating such a combined imager is challenging, as the performance demands of these measurements are very different. The X-ray CT detector must have a linear response across a wide dynamic range at high count rates, while SPECT requires high detection efficiency with good energy resolution at low count rates. CZT detectors are the most promising candidates for this application, because they satisfy both requirements while possessing better g-ray energy resolution

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Figure 15 Images of a printed circuit board (top) and a ball grid array (bottom). The images on the left are single-frame images (20 ms). The integration time of the images on the right side is 10 s (500 added frames). (Source: Ref. [30], reprinted with permission from Elsevier.)

and higher count-rate capabilities than scintillation detectors. Furthermore, the acceptable cost and the avoidance of cryogenic cooling help to promote the use of CZT in a combined CT/SPECT system. In 2003, William et al. produced a simultaneous CT/SPECT imager using a single CZT detector [32], so allowing the capture of structural and functional information at the same time. Figure 16 shows the CT and SPECT images thus collected from a hot lesion phantom containing 28.6 mCi of 99mTc sodium pertechnetate. The photo-peak efficiency and energy resolution for 140 keV gamma rays of 70% and 10%, respectively, are constant for a fluence rate up to 103 cps. The smallest lesions visible in SPECT and in CT are, respectively, 9 mm and 4.5 mm in diameter. Count rates less than 103 cps are sufficient for radionuclide studies, and the energy resolution of the SPECT images is comparable to current clinical systems.

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Figure 16 CT (left) and SPECT (right) images of a cylindrical phantom with hole diameters ranging from 4.5 to 38 mm. The left image is a map of attenuation coefficients derived from detecting X-rays (CT). The radionuclide image on the right was reconstructed using the CT-generated attenuation map to compensate the 99mTc-emission data for photon attenuation. (Source: Ref. [32], reprinted with permission from Elsevier.)

4. SPACE AND ASTROPHYSICS Both CdTe and CZT detector arrays have crucial applications for space exploration and astrophysics investigations. Operating typically in the 10-500 keV range, CdTe/CZT detectors possess obvious advantages compared with Ge detectors and scintillators. Unlike Ge detectors, the ability of CdTe/CZT detectors to operate at room temperature obviates a complex cooling system, a feature that is especially important for low-power radiation detection systems used in space. Furthermore, the energy resolution of CdTe/CZT detectors is also superior to that of scintillators. In this section, we discuss the latest progress of CdTe/CZT radiation detectors in space exploration and astrophysics. The International Gamma-Ray Astrophysics Laboratory’s (INTEGRAL) ISGRI imager of the European Space Agency (ESA) is the first space instrument using good spectral resolution CdTe; it was launched in 2002 by the Russian PROTON launcher. The ISGRI comprises 16,384 CdTe detectors of 4  4  2 mm3 [33, 34], representing a sensitive area of 2621 cm2. Figure 17 is a view of the detection plane of the ISGRI camera formed with eight independent modules [33]. The ISGRI mainly is devoted to detecting and precisely measuring celestial gamma-ray photons between 15 keV and 1 MeV. It supports research on violent processes occurring near black holes, neutron stars, or in supernovae. Figure 18 illustrates the quality of the spectacular images obtained with a 128  128 pixels array during a test at CEA-Saclay [34]. The practical performance of ISGRI is illustrated in Fig. 19 with a

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Figure 17 View of the 8 ISGRI MDUs (white) at the bottom of the passive shield well (black) after integration in the IBIS detection system, 2004. (Source: Ref. [33], reprinted with permission from EDP Sciences.)

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Figure 18 Gammagraphy of a discobole statuette recorded by ISGRI detector array during a test at CEA-Saclay. (Source: Ref. [34], reprinted with permission from Elsevier.)

picture of the Cygnus region in the energy range 15-40 keV, where at least three sources are clearly visible. This is one of the finest images obtained so far in the soft gamma-ray domain. The results from the INTEGRAL/ ISGRI imager show that CdTe stability is better than expected and its internal background is comparable to that of scintillators. Meanwhile, the

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Figure 19 ISGRI image of the Cygnus region in the 15–40 keV energy range. (Source: Ref. [33], reprinted with permission from EDP Sciences.)

spectroscopic degradation of this imager in space is slow, with a lifetime of about 40 years in an eccentric orbit. As the first CdTe space gamma-ray imager in the world, the ISGRI has verified the possibility of employing a huge quantity of CdTe crystals in space applications. Presently, the Burst Alert Telescope (BAT) onboard the Swift gammaray burst (GRBs) explorer is the largest CZT gamma-ray imager in the world. It was launched on November 20, 2004 as one of NASA’s mediumclass explorer programs. The BAT instrument is equipped with CZT semiconductor detectors beneath a D-shaped coded mask. The 32,768 semiconductor detectors (4  4 mm2 area, 2 mm thick), which form a 5243 detector plane, are built in 16 blocks. The smallest unit encompasses 128 individual CZT wafers. A single XA1 ASIC reads out the signals from each unit. In orbit, the CZT detectors operate at a nominal temperature of 20  1  C. The nominal bias voltage is 200 V, where the designed maximum is –300 V. Figures 20 and 21, respectively, show the structure diagram of BAT and a photograph of the CZT detector module [35, 36]. BAT is designed to detect sources of GRBs and bright transient X-rays to determine their positions with an accuracy of 1-4 arc-min. The BAT’s energy range is from 15 to 150 keV, with an energy resolution of 6 keV (FWHM) at 122 keV. Figure 22 shows reconstructed images demonstrating the instrument’s burst imaging capabilities [36]. Still several other CdTe/CZT X-ray telescopes have been loaded on balloons for astrophysics investigations. The High Energy Focusing Telescope (HEFT) and the International Focusing Optics Collaboration for m-Crab Sensitivity (InFOCmS) are two good examples, both of which incorporate pixel-based CZT detectors for imaging.

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Figure 20 The Burst Alert Telescope (BAT) onboard the Swift spacecraft. The BAT has a 3 m2 D-shaped coded aperture mask with 5 mm pixels. The CZT array is 5243 cm2 with 4  4 mm2 detectors. (Source: Ref. [35], reprinted with permission from Elsevier).

Figure 21 The BAT detector module (DM). Two sub-arrays of 8  16 pieces of CZT tile lie on the top surface of the DM. (Source: Ref. [36], reprinted with permission from Springer.)

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Figure 22 Image of the letters “BAT”, demonstrating the burst imaging capabilities of the instrument, taken by irradiating with gamma rays from a 57Co radioactive source on to the BAT’s detector plane through a coded mask in a preflight ground calibration test. (Source: Ref. [36], reprinted with permission from Springer.)

HEFT was developed in a program led by Caltech in collaboration with Columbia University and the Danish Space Research Institute. Figure 23 illustrates the HEFT CZT focal-plane detector system [37]. An individual sensor consists of a 1.2  2.4 cm2, 2-mm-thick CZT crystal, with the anode contact segmented into pixels with 500 mm pitch.

Figure 23 Photo showing a HEFT focal plane detector system. Two CdZnTe/ASIC hybrid pixel sensors are mounted side-by-side. (Source: Ref. [37], reprinted with permission from Springer.)

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Gold stud/epoxy interconnects couple the sensor pixel’s contacts to a custom low-noise readout chip. HEFT has been flown once in Spring 2005, demonstrating the angular resolution of 1.5 arcmin (HPD) and spectral resolution of 1 keV (FWHM) at 60 keV. The Goddard Space Flight Center, NASA, leads the InFOCmS project, whose focal plane consists of a CZT detector with a 12  12 array of 2 mm pixels. The telescope was flown twice, in 2001 and 2004, and achieved an angular resolution of 2.2 arc-min (HPD) and spectral resolution of 4 at 32 keV. The good performance of the CZT detectors facilitated the successful detection of the astrophysical source Cyg X-1, even with an on-target observation time of only about a minute. Additionally, Japanese researchers are proposing to develop hard Xray telescopes (HXI) for the Non-thermal Energy eXploration Telescope (NeXT) mission with CdTe detectors [38]. The current goal for the CdTe detector in the HXI is a pixel detector with both a fine-position resolution of 200-250 mm and a high-energy resolution of better than 1 keV (FWHM) in the energy range from 5 to 80 keV. In addition, designers will use CdTe pixel detectors and a stack of 24 Si DSSDs to assemble a semiconductor Compton Telescope. Figure 24 illustrates the prototype of the CdTe pixel detectors and the spectra acquired by them. Efforts are underway to improve the performance of CdTe pixel detectors as thick as 5 mm. Furthermore, the Danish National Space Institute (DNSC) is developing a CdTe/CZT detectors program to fabricate the miniature X- and gamma-ray sensor (MXGS) for the ESA-supported Atmosphere-Space Interactions Monitor (ASIM) mission, which is expected to be launched to the International Space Station (ISS) in 2011.

5. NATURE AND DEVELOPMENT OF CMT DETECTORS CMT is a diluted magnetic compound semiconductor, previously used as Faraday rotators, optical isolators, solar cells, lasers, magnetic field sensors, and infrared detectors. In 1999, Burger et al. first investigated the potential application of this material as radiation detector. CMT displays some advantages over CZT, making it a good candidate to compete with the latter in radiation-detector applications. First, the segregation coefficient of Mn in CdTe is nearly equal to unity in all directions, while that of Zn in CdTe has a coefficient of 1.35 [39]. This difference is reflected in the nearly uniform concentration of Mn in CdTe, compared to the high variation of Zn concentration in CdTe. The superior compositional homogeneity of CMT potentially enhances the yield of crystals suitable for detectors; ultimately, this might lower the costs of producing large-area arrays. Secondly, CMT has greater tunability of the band gap due to the large compositional influence of Mn. The addition of Mn increases

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the room-temperature band gap at a rate of 13 meV/[%Mn], that is, more than twice as large as the increase after adding Zn to CdTe [40]. Therefore, the band gap in the range 1.7-2.2 eV, which proved ideal for assuring optimal signal/noise ratio in X-ray and gamma-ray detectors [41], can be attained by adding relatively less Mn. This merit diminishes many alloying related problems. However, several material properties must be improved before CMT can be practically employed for X-ray and gamma-ray detections. First, compared to CZT, the bond ionicity of CMT is higher, entailing a greater tendency for crystallization into a hexagonal structure, but not in the expected zinc-blende structure [42, 43]. Also, higher ionicity can generate

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twins easily in crystals. Second, the resistivity of CMT crystals must be improved. Normally, CMT crystals grown by the Bridgman methods are p-type materials due to their high concentration of cadmium vacancies (VCd) the dominant acceptors, and the resistivity of as-grown crystals can be as low as 10–103 O cm, thus not satisfying the resistivity requirement of X-ray and gamma-ray radiation detectors. Recently, BNL and the Institute of Physics, Poland Academy of Sciences (PAS) cooperatively addressed some of these issues using their comprehensive material characterization techniques and improved crystalgrowing capabilities. Significant progress was made, and better results were reported [8]. Detector-grade CMT crystals were grown, and the first CMT detector was fabricated (Cd0.94Mn0.06Te doped with Vanadium 5  1016 cm3), as shown in Fig. 25. This figure also shows the energy spectra obtained with a sealed Am-241 source. Accordingly, the mtproduct of electrons is 2.1  104 cm2/V. Furthermore, Fig. 26 is the X-ray map from a CMT detector measured at BNL’s National Synchrotron Radiation Source (NSLS). It demonstrates that Synchrotron X-ray mapping

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can be used to correlate how microscale defects, for example, twin boundaries, Te inclusions, and dislocations, affect the CMT detector’s performance. However, compared with the performance of CZT detector, there still is much room to improve CMT detectors before they can to be applied as practical radiation detectors.

6. SUMMARY AND FUTURE WORK As promising materials for radiation detection, CdTe/CZT have high stopping power for energetic photons, good sensitivity and energy resolution, and excellent room-temperature operation capacity. This chapter discussed the primary applications of CdTe and CZT radiation detectors, including national security and nonproliferation inspections, medical imaging, and space exploration and astrophysics investigation. The last decade saw great progress in all these fields. However, material issues, for example, improving the availability of large-volume CZT crystals, remain the main limiting factor for further developing these radiation detectors. The combination of modern characterization methods and modified growth techniques is providing a deep understanding of CdTe/CZT material properties, which might resolve these problems and accelerate the related applications in the near future. CMT radiation detectors are in the elementary stages of development compared with CdTe/CZT detectors, and as yet have no practical application in commercial

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radiation detection due to their poorer material properties. Nevertheless, recent progress possibly predicts a breakthrough of this promising radiation material.

ACKNOWLEDGMENTS We would like to thank the authors of the references and the Institute of Electrical and Electronics Engineers, Inc. (IEEE), the International Society for Optical Engineering (SPIE), Elsevier, Springer, John Wiley & Sons, Inc. (Wiley), Materials Research Society (MRS), American Institute of Physics (AIP), and E´dition Diffusion Presse Sciences (EDP), among others for permission to reprint the figures in their works.

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[20] C. Scheiber, B. Eclancher, J. Chambron, V. Prat, A. Kazandjian, A. Jahnke, R. Matz, S. Thomas, S. Warren, M. Hage-Ali, R. Regal, P. Siffert, M. Karman, Nucl. Instr. Meth. A 428 (1999) 138. [21] D.J. Wagenaar, S. Chowdhury, J.C. Engdahl, D.D. Burckhardt, Nucl. Instr. Meth. A 505 (2003) 586. [22] L. Verger, J.P. Bonnefoy, F. Glasser, P. Ouvrier-Buffet, J. Electron Mater. 26 (1997) 738. [23] C. Mestais, N. Baffert, J.P. Bonnefoy, A. Chapuis, A. Koenig, O. Monnet, P. Ouvrier Buffet, J.P. Rostaing, F. Sauvage, L. Verger, Nucl. Instr. Meth. A 458 (2001) 62. [24] L. Verger, M.C. Gentet, L. Gerfault, R. Guillemaud, C. Mestais, O. Monnet, G. Montemont, G. Petroz, J.P. Rostaing, J. Rustique, IEEE Trans. Nucl. Sci. 51 (2004) 3111. [25] K. Spartiotis, A. Leppa¨nen, T. Pantsar, J. Pyyhtia, P. Laukka, K. Muukkonen, O. Mannisto, J. Kinnari, T. Schulman, Nucl. Instr. Meth. A 550 (2005) 267. [26] T.O. Tu¨mer, S. Yin, V. Cajipe, H. Flores, J. Mainprize, G. Mawdsley, J.A. Rowlands, M.J. Yaffe, E.E. Gordon, W.J. Hamilton, D. Rhiger, S.O. Kasap, P. Sellin, K.S. Shah, Nucl. Instr. Meth. A 497 (2003) 21. [27] N.K. Zelenina, S.M. Ignatov, V.P. Karpenko, L.V. Maslova, O.A. Matveev, D.A. Popov, A.I. Terent’ev, A.A. Tomasov, Nucl. Instr. Meth. A 283 (1989) 274. [28] F. Glasser, G. Thomas, M. Cuzin, L. Verger, Nucl. Instr. Meth. A 322 (1992) 619. [29] S. Baba, K. Ohmori, Y. Mito, T. Tanoue, S. Yano, K. Tokumori, F. Toyofuku, S. Kanda, Nucl. Instr. Meth. A 458 (2001) 262. [30] K. Spartiotis, J. Havulinna, A. Leppa¨nen, T. Pantsar, K. Puhakka, J. Pyyhtia¨, T. Schulman, Nucl. Instr. Meth. A 527 (2004) 478. [31] Y. Ohno, M. Torikoshi, T. Tsunoo, K. Hyodo, Nucl. Instr. Meth. A 548 (2005) 72. [32] W.C. Barber, K. Iwata, B.H. Hasegawa, P.R. Bennett, L.J. Cirignano, K.S. Shah, Nucl. Instr. Meth. A 505 (2003) 595. [33] F. Lebrun, J.P. Leray, P. Lavocat, J. Cr´etolle, M. Arque`s, C. Blondel, C. Bonnin, A. Boue`re, C. Cara, T. Chaleil, F. Daly, F. Desages, H. Dzitko, B. Horeau, P. Laurent, O. Limousin, F. Mathy, V. Mauguen, F. Meignier, F. Molinie´, E. Poindron, M. Rouger, A. Sauvageon, T. Tourrette, Astron. Astrophys. 411 (2003) L141. [34] O. Limousin, Nucl. Instr. Meth. A 504 (2003) 24. [35] G. Sato, A. Parsons, D. Hullinger, M. Suzuki, T. Takahashi, M. Tashiro, K. Nakazawa, Y. Okada, H. Takahashi, S. Watanabe, S. Barthelmy, J. Cummings, N.l. Gehrels, H. Krimm, C. Markwardt, J. Tueller, et al., Nucl. Instr. Meth. A 541 (2005) 372. [36] S.D. Barthelmy, L.M. Barbier, J.R. Cummings, E.E. Fenimore, N. Gehrels, D. Hullinger, H.A. Krimm, C.B. Markwardt, D.M. Palmer, A. Parsons, G. Sato, M. Suzuki, T. Takahashi, M. Tashiro, J. Tueller, Space Sci. Rev. 120 (2005) 143. [37] F.A. Harrison, F.E. Christensen, W. Craig, C. Hailey, W. Baumgartner, C.M.H. Chen, J. Chonko, W.R. Cook, J. Koglin, K. Madsen, M. Pivavoroff, S. Boggs, D. Smith, Exp. Astron. 20 (2005) 131. [38] T. Takahashi, K. Nakazawa, S. Watanabe, G. Sato, T. Mitani, T. Tanaka, K. Oonuki, K. Tamura, H. Tajima, T. Kamae, G. Madejski, M. Nomachi, Y. Fukazawa, K. Makishima, M. Kokubun, Y. Terada, J. Kataoka, M. Tashiro, Nucl. Instr. Meth. A 541 (2005) 332. [39] A. Tanaka, Y. Masa, S. Seto, T. Kawasaki, J. Crystal Growth 94 (1989) 166. [40] D.J. Olego, J.P. Faurie, S. Sivananthan, P.M. Raccah, Appl. Phys. Lett. 47 (1985) 1172. [41] J.T. Toney, T.E. Schlesinger, R.B. James, IEEE Trans. Nucl. Sci. A 428 (1999) 14. [42] S. Perkowitz, R. Sudharsanan, J.M. Wro´bel, B.P. Clayman, P. Becla, Phys. Rev. B 38 (1998) 5565. [43] R. Triboulet, Mater. Forum 15 (1991) 30.

CHAPTER

IID Electro-optic Modulator Applications Carl J. Johnson, Gray L. Herrit and Eric R. Mueller

1. INTRODUCTION Since the discovery of infrared (IR) lasers in the mid-1960s, the potential of CdTe as an electro-optic (EO) modulator material has been wellrecognized. Experiments conducted at 1.0 mm in 1967 and 1970, at 3.39 mm in 1969, at 10.6 mm in 1969 and 1971, and at 23 and 28 mm 1968 demonstrated CdTe’s modulation capability over a wide range of IR wavelengths and generated the measurements of its EO coefficient as shown in Table 1. With the emergence of instrument- and industrial-grade CO2 lasers around 1970, commercial interest in 10.6 mm modulators increased and motivated considerable efforts to overcome problems impeding the deployment of CdTe at that time. Rapid progress by material and device manufacturers to some degree mitigated a variety of ingotsize, ingot-cracking, optical-absorption and scatter, electrical-resistivity, crystal-fabrication, antireflection (AR)-coating, electroding, and high-cost limitations. By 1973, a reasonable capability existed to provide 10  10  50 mm3, >108 ohm-cm, >99.5% transmissive, affordable CdTe EO-modulator devices and components. Early applications that required substantial numbers of modulator crystals and components included the CO2 Laser Fusion Program at the Los Alamos Scientific Laboratory in the United States, space-communication development programs at both Hughes Aircraft Company in the United States and the European Space Agency, and the US Department of Defense CMAG Program.

Carl J. Johnson is with II-VI Incorporated, Saxonburg, PA 16056 USA Gary L. Herrit is with II-VI Incorporated, Saxonburg, PA 16056 USA Eric R. Mueller is with Coherent, Incorporated, Bloomfield, CT 06002 USA

239

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Carl J. Johnson, Gray L. Herrit and Eric R. Mueller

Table 1

EO coefficients (r41) of CdTe versus wavelength

Wavelength (mm)

1.0 3.39 10.6 23 28

EO coefficient r41 (m/V 1012)

2.2 5.3 6.8 6.8 6.2 5.5 5.0

Radiation source

Reference

Monochromator

[1] [2] [3] [3] [4] [5] [5]

HeNe laser CO2 laser CO2 laser Water vapor laser Water vapor laser

The EO effect in CdTe depends on the fact that an electric field applied in certain crystallographic directions causes lattice distortions which, in turn, induce changes in the indices of refraction or birefringence encountered by polarized light traveling through the crystal. Thus, it is possible to electrically modulate the phase, polarization, or intensity of laser radiation using properly oriented CdTe crystals. Several derivations of the formulas that characterize EO modulation in zinc-blende structure crystals, such as CdTe, have been presented by others [6–8]. In summary, this crystal type gives rise to a “transverse” EO effect, where the electric field is applied normal to the direction of light propagation and the induced changes in refractive index vary as Dn / n30 r41

V ; d

ð1Þ

where n0 is the unperturbed refractive index, r41 is the EO coefficient, V is the applied voltage, and d is the electrode separation.

2. PRACTICAL CONFIGURATIONS Amplitude modulation (AM) is the most commonly used type of modulation. The configuration generating the maximum AM effect is shown in Fig. 1(A), where light propagates in the [ 110] direction, the electric field is applied in the [110] direction, and the linear polarization of the incident light is also oriented along the [110] direction. In this case, the amount of relative phase retardation in radians, Grel, experienced by the light at the exit of the CdTe modulator is Grel ¼

2pl 3 V n r41 ; l 0 d

ð2Þ

where l is the crystal length and l is the free-space wavelength of the laser radiation. For extra-cavity AM of a laser beam, with a polarization analyzer that is crossed with the input polarization following the crystal, the intensity, I, of the light passing through the analyzer is

241

Electro-optic Modulator Applications

E = V/d

E = V/d

Analyzer

Analyzer d

d

l

l

[111]

[110] [110]

Polarizer

[001]

A Figure 1

Polarizer

[110] [112]

B Configurations for generating (A) maximum AM and (B) FM using CdTe crystals.


 I 2 Grel ; ¼ sin 2 I0

ð3Þ

where I0 is the intensity of the incident light. With a ¼-wave plate placed in series with the modulator crystal, the amplitude modulated output intensity is
 I Grel p þ ¼ sin2 ; ð4Þ 2 I0 4 and good small-signal linearity results. For frequency modulation (FM), the configuration shown in Fig. 1(B) is most commonly utilized. Light propagates in the [110] direction with both the light polarization and applied electric field oriented in the [111] direction. In this case, the amount of phase shift, G, produced is pffiffiffi 3pl 3 V n0 r41 : ð5Þ G¼ l d

3. ISSUES AND LIMITATIONS 3.1. Mechanical CdTe EO modulator crystals can be manufactured in a variety of sizes. The grain sizes obtainable, at least in the case of Bridgman growth, limit modulator cross-sections to about 10  10 mm2; however, the most common cross-sections utilized in CO2 laser applications fall between 2  2 and 6  6 mm2. Grain-size limitations also restrict modulator lengths to about 50 mm and because these are transverse EO modulators, it is desirable for the crystals to be as long as possible. As shown in the previous section, the modulation per volt obtainable with a CdTe modulator is proportional to the length of the crystal and

242

Carl J. Johnson, Gray L. Herrit and Eric R. Mueller

inversely proportional to the electrode spacing. In theory, it is desirable to make the cross-section as small as possible; however in practice, the minimum cross-section is limited by the physical strength of the material. CdTe is relatively soft and readily cleaves parallel to the (110) family of crystal planes, thus the manufacturing of small cross-section crystals presents a serious challenge. Historically, 0.8  0.8  30 mm3 (aspect ratio ¼ 37.5) AM-cut crystals have been fabricated; however, a more practical size for the deployment of both AM- and FM-cut modulators is in the range of 2  2  50 mm3 (aspect ratio ¼ 25.0).

3.2. Optical

Top

1 Grd

2

. El

A

Phase Shift (deg.) –25 0 25 50

Phase Shift (deg.) –25 0 25 50

EO-grade CdTe has excellent optical properties in the mid- and far-IR regions. The highest quality material displays minimal near-IR optical scatter and can be used at wavelengths as short as 1 mm. Some ingots having excellent transmission properties at 10 mm cannot be used below 5 mm because of Te-precipitate induced, near-IR scatter in the bulk material. The generation of stress can occur during the cutting, grinding, and polishing of CdTe crystals. Herrit and Reedy [9] showed that certain fabrication techniques can induce or reduce stress birefringence in CdTe modulator crystals. In a later paper [10], they showed that residual stress birefringence from the fabrication process, along with other crystal defects, can affect the operation of a modulator. Their measurements were accomplished by focusing a CO2 laser beam through a CdTe modulator, scanning the beam across the aperture, and measuring the phase shift at each position. Figure 2(A) shows a phase-shift plot indicating a notable amount of residual birefringence in a CdTe modulator having coarse-ground sides. Figure 2(B) shows the corresponding plot indicating lesser amounts of birefringence for a modulator having highly polished sides. The 10.6 mm bulk-absorption coefficient of EO-grade CdTe is typically 0.0005 cm1. In high-power CO2 lasers, bulk absorption results in heating

ect

3

rod

e

4

1

2 e Sid

3

4

B

Top

1 Grd . El 2 3 ect rod e

4

1

2 e Sid

3

Figure 2 Stress birefringence induced by (A) course-ground and (B) highly polished sides.

4

Electro-optic Modulator Applications

243

of the optical material, so low-bulk absorption is highly desirable. Unfortunately, the low-bulk absorption of CdTe is offset by its low thermal conductivity. In other words, although the material absorbs very little laser power, the small amount of heat generated is difficult to conduct away, which leads to some practical power limitations for CdTe modulators. Numerous laboratory experiments and end-user applications have indicated that the practical, continuous-wave (CW), power limit for an air-cooled CdTe modulator is about 20 watts (W). If the modulator package is water cooled, then this power handling limit can be increased to about 50 W. Generally, the material does not damage at these power levels, but thermal lensing of the laser energy can be quite severe. Through good thermal management and engineering, some exceptions to the above limits have emerged over the years. Most notable is the work done by DeMaria et al. in producing a modulator package that is capable of withstanding significantly higher laser intensity levels [11]. Their work was documented in US Patent 5,680,412. In this work, they sandwiched the CdTe crystal between two ZnSe windows, as shown in Fig. 3. The windows were AR coated on one side and uncoated (UC) on the opposite side, and the CdTe modulator was uncoated on both ends (UC/UC). The UC face of a ZnSe window was optically contacted to the respective UC faces of the modulator. Since, in many high-power laser applications, surface damage is the ultimate failure mechanism, the fabrication technique of DeMaria et al. greatly extended the range of use of these modulators. Figure 3 shows the concept of this assembly method, wherein 10 is the UC/UC modulator and 12 and 14 are the UC/AR windows.

3.3. Electrical Electrically, CdTe modulators possess high resistivity (>108 ohm-cm) and low capacitance (5-10 pf). The material has a very high intrinsic breakdown voltage, but more practically, the modulator packages usually breakdown first due to surface conduction along the faces or sides of

8⬘

LASER INPUT

102⬘

104

12⬘

Figure 3

10

LASER OUTPUT

14⬘

Use of UC/AR end-plates to reduce damage to CdTe modulator crystals.

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Carl J. Johnson, Gray L. Herrit and Eric R. Mueller

A

Phase Shift (deg.) –25 0 25 50

Phase Shift (deg.) –25 0 25 50

the crystals. This often occurs at field strengths of about 1 kV/mm of electrode spacing at atmospheric pressure. At lower ambient pressures, air breaks down at lower voltages, so care must be taken to ensure that this does not happen. Most CdTe modulators are made with an electrode spacing and a crystal length such that the ½-wave voltage is about 1 kV/mm of electrode spacing. For example, a 5  5  50 mm3 crystal will have a ½-wave voltage of about 5 kV. Since the cross-section of the crystal is 5  5 mm2, the electrodes are spaced 5 mm apart. This means that the air breakdown voltage and the ½-wave voltage for the crystal are roughly the same, thus it is advisable to avoid using this modulator in an application where ½-wave modulation is required. These rules do not apply when the voltage pulse is very fast (<1 ms) or if the modulator package is sealed and a dry, inert gas is used to pressurize the package. Certain electrically active impurities can give rise to deep-level traps in the bulk material of a CdTe modulator. When a voltage is applied, these deep-level traps can distort the electric-field distribution, resulting in field-strength nonuniformities across the clear aperture and nonuniform modulation across the laser beam. Figure 4 displays phase-shift plots for a CdTe modulator. Figure 4(A) indicates the pattern of residual birefringence in the crystal at rest, that is, with no voltage applied. Figure 4(B) shows the phase-shift distribution with the ¼-wave voltage applied to the modulator. Ideally, this plot should be a flat plane at 90 of phase shift. Charge trapping typically results in a higher phase shift near the positive electrode and a lower phase shift near the ground electrode. Trapping is primarily a problem that occurs when the user wants to apply a DC voltage or long-duration voltage pulse to the modulator. Since it takes on the order of 1/10 of a second for the charge traps to fill, any voltage pulse that is much shorter than this will not cause such nonuniform behavior. Applications where a DC bias is required will suffer from this phenomenon.

Top

1 Grd . El 2 3 ect rod e

3 4

1

4

2 e Sid

B

Top

1 Grd . El 2 3 ect rod e

4 3 4

1

2 e Sid

Figure 4 Nonuniformity in phase retardation caused by charge trapping in CdTe material.

Electro-optic Modulator Applications

245

4. SUCCESSFUL AND CONTEMPLATED DEPLOYMENTS CdTe modulator technology can be combined with high-reliability RFexcited CO2 laser technology to enable a wide range of applications. Over the years, many techniques utilizing CdTe modulators have been realized in the laboratory, including laser Q-switching, laser cavity dumping, optical pulse shaping, frequency/phase modulation, laser mode locking, and optical free-space communications. To date, only the first two of these have found their way into commercial products; however, a number of research groups continue to investigate wider deployment of some of the other applications.

4.1. Laser Q-switching This application typically involves the Q-switching of an RF-excited, waveguide, CO2 laser by using the CdTe as a voltage-tunable birefringent element as shown in Fig. 5. Initially the modulator (EO QS) voltage is set to 0 V. In this state, the cavity Q is very low because the light leaving the gain region and making a round-trip through the EO QS and reflectivephase retarder (RPR)-mirror pair is rotated in polarization by 90 and reflected out of the cavity by the thin film polarizer (TFP). The EO QS voltage is then rapidly switched to the ¼-wave voltage. In this state, the cavity has a typical lasing Q value, since the energized EO QS cancels the polarization rotation effect of the RPR-mirror pair. In this configuration, the lasing action will build-up and an output pulse will exit the cavity via the partially reflecting (partial R) output coupler. With this configuration, the termination of the light pulse is adjustable via the timing for returning the EO QS voltage to 0 V. At that point in time, any optical energy remaining in the cavity will be “dumped out” by reflecting from the surface of the TFP. The resulting optical pulses consist of an initial temporal pulse in the order of 150 ns FWHM, followed by a lower energy tail whose length can be adjusted up to a few microseconds. When combined with folded-cavity gain sections, practical commercially packaged lasers with average powers in the range of 50 W and pulse repetition rates of 100 kHz have been realized. The optical output from this type of laser has found application in material micromachining.

Partial R

Gain Region (often folded)

TFP

Output

Figure 5

Optical arrangement for a Q-switched CO2 laser.

EO QS RPR-Mirror

246

Carl J. Johnson, Gray L. Herrit and Eric R. Mueller

4.2. Laser cavity dumping In this application, the optical configuration is quite similar to that shown in Fig. 5, with the exception that the partial R is replaced by a 100% reflective mirror and the output is reflected from the surface of the TFP. The operation is also similar to the Q-switched operation with the exception that once the circulating power in the laser builds up, the EO QS voltage is very rapidly returned to 0 V, thus dumping the optical energy out of the cavity. This results in a quasi-Gaussian temporal optical pulse with a FWHM in the order of 15 ns. Practically sized, cavity-dumping lasers can have average output powers in the order of 30 W. The Coherent EOM-10, pictured in Fig. 6, is a scientific laser of this type having a pulse width of 12 ns and an average power of >10 W. To date, the majority of applications for this laser have been in scientific research, but there are potential commercial applications in development.

4.3. Optical pulse shaping The use of CdTe in optical pulse shaping again takes advantage of its voltage-tunable birefringence. The CdTe crystal is used in conjunction with polarization-selective passive optical elements and fast-drive electronics, to either tailor the shape of an existing optical pulse (e.g., making its rise- or fall-time shorter), or obtain pulses from a CW laser beam. Material processing is the primary area of interest for this modulator embodiment. To date, however, the cost of such a CdTe-based pulse shaper has prevented its widespread adoption in commercial applications.

4.4. Free-space optical communications With the availability of high-reliability, single-frequency compact CO2 lasers and the fairly good transmission of 10 mm light through the atmosphere, optical free-space communications might be feasible.

Figure 6 Photograph of a Coherent EOM-10.

Electro-optic Modulator Applications

247

At this point in time, there are a few research organizations pursuing this possibility. Most of the communication approaches under development utilize CdTe as an amplitude modulator, but there are communication strategies under investigation, which utilize FM as well.

4.5. Optical frequency and phase modulators As described in earlier sections, depending on the cut of the crystal, one can realize a voltage-dependent optical phase shift using CdTe. If a timevarying voltage waveform is placed on such a crystal, then frequency or phase modulation can be realized. At this point, the only applications for this embodiment have been in research.

4.6. Laser intracavity modulation and mode locking Intracavity AM and FM of CO2 lasers, outside of the mode-locking regime, have been utilized in various remote-sensing research programs, typically sponsored by government agencies. These programs are exploratory in nature and there are no commercially significant applications at the present time. The use of CdTe for the mode locking of a CO2 laser has been realized in both AM and FM configurations. While applications in this area have thus far been entirely in research, the mode-locked or more exotic mode-locked, Q-switched, cavity-dumped format may find its way into material processing applications in future.

REFERENCES [1] O.M. Stafsudd, F.A. Hack, K. Radisavljevic, Appl. Opt. 6 (1967) 1276. [2] V.S. Bagaev, T.Ya. Belousova, Yu.N. Berozashvili, D.Sh. Lordkipanidze, Sov. Phys. Semicond. 3 (1970) 1418. [3] J.E. Kiefer, A. Yariv, Appl. Phys. Lett. 15 (1969) 26. [4] I.V. Nikolaev, M.M. Koblova, Sov. J. Quantum Electron. 1 (1971) 158. [5] C.J. Johnson, Proc. Inst. Electr. Eng. 56 (1968) 1719. [6] C.S. Namba, J. Opt. Soc. Am. 51 (1961) 76. [7] J.E. Pankove, Optical Processes in Semiconductors, Prentice-Hall, Englewood Cliffs, NJ, 1971. [8] A. Yariv, Introduction to Optical Electronics, second ed., Holt Rinehart and Winston, New York, 1976. [9] G.L. Herrit, H.E. Reedy, J. Appl. Phys. 65 (1) (1989) 393–395. [10] G.L. Herrit, H.E. Reedy, Electro-Opt. Mater. Switches, Coatings, Sensor Optics, Detectors, SPIE 1307 (1990) 1509–1515. [11] A.J. DeMaria, J.T. Kennedy, R.A. Hart, Apparatus for Improving the Optical Intensity Induced Damage Limit of Optical Quality Crystals, US Patent No. 5,680,412 (1997).

CHAPTER

IIE Optical Detectors Based on CdTe Pure Crystals for HighEfficiency Optical Computers P.G. Kasherininov and A.A. Tomasov

1. INTRODUCTION At present, CdTe crystals with high electrical, optical, and electro-optical characteristics are available. They are widely utilized as optical, roentgen, and nuclear radiation detectors. Pure (noncompensated) crystals with high electrical resistance (r ¼ 107–108 O cm), low impurity densities (Nt < 1013 cm3), and high carrier mobilities and lifetimes are of special interest in the field of detectors. Detectors based on such a materials at present are the main spectrometric detectors of nuclear radiation functioning without cooling. Such detectors are MSM—structures produced with cold evaporation of metal electrodes (M) on the crystal surface (S). They function at high applied voltages and posses large working volumes, low dark currents, and low noises. They are not polarized during the radiation detecting. At the last time, it was discovered that MSM structures of this type based on pure p-CdTe crystals with high electrical resistance are promising as a new type of fast optical registering media for high-efficiency optical computers.

1.1. Optical computers based on semiconductor structures During the last 20 years, the efficiency of modern electronic computers is increasing due to the redoubling of transistor density in super large integrated microchips (SLIC), which takes place in each year and a half.

A.F. Ioffe Physico-Technical Institute, St Petersburg 194021, Russia

248

Optical Detectors Based on CdTe Pure Crystals

249

Technology of integrated schemes has certain physical limits. These limits result in the stagnation of microelectronic systems’ efficiency. Contemporary electronic computers exhaust their development. There is a need of computers functioning on a new principle with higher rate. One of the substitutions of modern electronic computers is to create so-called optical computers. The information in these computers is transported with a light flux. Optical calculations are divided in digital optical calculations and analog ones. Digital optical calculations use light signals to perform operations of digital logic. They are targeted at the class of applications, which are performed currently with electronic computers. The advantage of optical computers is their ability to transport the information at the speed of light and to substitute wires with wireless optical connections. Analog calculations include analog operations above images. Such a processor is able to perform action with two-dimensional pictures at one step; meanwhile, the computer instruction might also be a picture. The development of analog two-dimensional optical processor is now of maximum interest and allows one to look forward at producing of super efficient computers with an operation rate more than 1012 operations/s and processor operating frequency of 106 cycles/s. Such processors are still not realized.

1.2. Optical registering media in contemporary optical processors on MIS structures Modern optical processors is based upon the registering media, which are metal (M)-insulator (I)-semiconductor(S) structures (MIS structures) with a thick (1 mm) insulating layer opaque to the charge carriers [1–3]. The external bias voltage is applied to such structures, and the recorded image is projected on the surface of it with a light flux of a fixed exposure time. The image is recorded in structure as a two-dimensional electric charge distribution with the density following distribution of brightness of the image on a surface of structure. In such MIS structures, the recorded charge cannot leak out of crystal through the insulating dielectric layer after the recording light is off. Low operation rate of these structures and the devices based on them (n ¼ 102–103 cycle/s) is predefined by the necessity to carry out the erasing of charge. It takes up most of the time of information recording cycle (t ¼ 102 103 s). Meanwhile, information reading and recording of images takes only microseconds. The main contemporary optical processors are based on such a structure: spacetime optical modulators (PROM), liquid crystal light modulators (MIS-LC), etc. [1–3].

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P.G. Kasherininov and A.A. Tomasov

1.3. Fast optical registering media on semiconductor M(TI)S-nanostructures At present, devices based on MIS structures with a thin nanosize dielectric layer (TI) (M(TI)S nanostructures) are proposed as fast optical registering media [4–14]. The thickness of a dielectric is 2–5 nm. It was established that photocurrent occurs in such M(TI)S nanostructures under illumination with “proper” light. The current is proportional to the illumination intensity. Meanwhile, in crystal at the boundary of dielectric layer (TI) the free photo carriers are created. Their density is proportional to the illumination intensity. This charge exists under the recording light only. At switching on (off) the recording light the charge establishes (dissipates) at the scale of microseconds. It is not necessary to erase the existing charge in order to record new image in such structures. It is established that real contacts of metal-semiconductor in MSM structures which are produced by covering of real crystal surface with metal appeared to have the following structure: metal (M)-thin insulator (TI) semiconductor (S)-thin insulator (TI)-metal (M). It occurs due to presence of a thin dielectric (2–5 nm) at the crystal surface. They are shortly called M(TI)S(TI)M nanostructures [10–14]. Such M(TI)S(TI)M nanostructures based on pure high-resistive CdTe crystals are proposed to be used in fast optical registering media to realize high efficient optical computers. It is convenient to use a modification of such M(TI)S(TI)M nanostructure on pure CdTe crystal, which presents n-p transition with inverse bias at one side of the crystal and thin nanosize dielectric (TI) layer at the opposite side (n-p (TI)M-nanostructures) [10–14].

2. PROCESSORS FOR DIGITAL OPTICAL COMPUTERS BASED ON N-P(TI)M NANOSTRUCTURES OF CdTe Processors of optical digital computers are intended for realization of operations of digital logic and represent the photon keys executed as n-p(TI)M nanostructure, illuminated with two light streams: controlling (I2) and information (I0) [13]. Figure 1 depicts a optical shutter (A) and commutator (1  2) (C) realized upon these structures. The functioning of photon switches is based on reversible change of electric field strength in such nanostructures under illumination with light stream. Photon switches functioning is based on transversal electro-optical effect. It is n-p(TI)M nanostructure with high resistive electro-optical CdTe crystal with a thin nanosize insulating layer (TI) with a thickness 2–5 nm. The nanostructure is placed between to crossed polarizers. It works in a following way: narrow

λ=0.82mkm

Optical Detectors Based on CdTe Pure Crystals

I2

+

I01

M1 I0

251

n-CdTe

λ=1.3mkm

p-CdTe M2

TI

A

2

1

– 3

I, arbitrary units

0.4 0.3 1 2 34

4 3 2 1

0.2 0.1 0

0

50

100 150 t, mks

=0.82mkm

B

I2

200

Il +

I0

M1

Ill

n-CdTe

λ=1.3mkm

p-CdTe 3

C

1

TI

2

M2 –

Figure 1 Photon switches on semiconductor n-p(TI)M nanostructures utilizing electrooptical crystals (CdTe). (A) Principle diagram of an optical gate and illumination geometry: 1 and 3, polarizers; 2, n-p(TI)M nanostructures. (B) Oscillogram of transmission of commutated (passing through) stable light beam (I0, l ¼ 1.3 mm) under the illumination of the structure with rectangle pulses of controlling light (I2, l ¼ 0.82 mm) with intensities I2 (mW/cm2): 1-0.5; 2-5; 3-7; 4-9 (U0 ¼ 400 V). (C) Principle diagram of light-controlled optical commutator (1  2): 1, polarizer; 2, n-p(TI)M nanostructure; 3, polarization-sensitive prism (Glan prism).

information light beam (I0) is let through the area of volume charge of n-p transition in the crystal. It does not undergo absorption in crystal. Controlling beam of “proper” light illuminates the structure surface in direction parallel to the electric field. Controlling light induces formation of electric charge at the boundary of insulating layer (TI). This charge

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changes the electric field strength distribution in crystal on the way of light beam. At the entry of the structure the intensity of information light beam rapidly changes due to transversal electro-optical effect. Deep impurity levels in crystals determine the rate of such structures. The following parameters are achieved in n-p(TI)M structures with pure CdTe crystals (r ¼ 108 O cm, Nt < 1013 cm3) [13]: Spectral range of the switched light, l (mm)

1–2

Spectral range of the controlling light, l (mm)

0.82–0.83

Cycle time, t (s)

105–106

Intensity of controlling light, I (W/cm2)

101

Working voltage, U0 (V)

400

Modulation depth (%)

>90

3. PROCESSORS FOR ANALOG OPTICAL COMPUTERS OF INCOHERENT LIGHT ON n-p(TI)M NANOSTRUCTURES ON CdTe Analog optical processor is targeted at the treatment and comparison of the images. It might be designed on a basis of n-p(TI)M nanostructures with CdTe [11, 12, 14]. Figure 2(A) represents the basic circuit of the processor and geometry of illumination. “Proper” light pulse containing of recording image is projected at the surface of such a processor from the side of inverse biased photosensitive n-p transition. On an opposite surface of the processor the pulse of the reading “proper” light homogeneously distributed on a surface of a crystal or carrying the image of compared object is projected. Under recording light stream (from the side n-p transition) the twodimensional electric charges are created in crystal. The density of the charges is proportional to the brightness of image. Reading the recorded information in such processor is made by registering a magnitude of a photocurrent on an output of the processor from action of a reading light stream (with simultaneous illumination of processor with recording light). A photocurrent thus contains the information on magnitude and a configuration of an electric charge written down in a crystal (information of recorded image). Reading of the recorded image in such processors is performed by a one-step process, and the instruction (reading light beam) itself can be an image. Light and photoresponse pulses are shown in Fig. 2B and C [11, 12]. Frequencies of recording and reading light are equal, pulses are

Optical Detectors Based on CdTe Pure Crystals

τ0=20μs I0=10mW/cm2

253

τ1=50μs I1=4mW/cm2 M n p Tl + –

I, mW/cm2

10 5 0

J, mA/cm2

A

0.2 0.1 0

τ=40μs

0

20

I, mW/cm2

B 10 5 0

J, mA/cm2

τ=10μs

C

0.6 0.5 0.4 0.3 0.2 0.1 0

0

20

R

60

80

40 60 t, μS

80

40

60 80 t, μS

40 60 t, μS

80

100

Figure 2 Optical processor based on n-p(TI)M nanostructure with CdTe targeted at registering and processing of image signals (A), structure and illumination geometry (B, C). Time dependence of recording and reading light pulses and pulses of photoresponse at the outlet of processor (l ¼ 0.82 mm). The delay time between the pulses: (B) 40 ms (C) 10 ms.

shifted respectively toward each other with a delay time t ¼ 40 ms (Fig. 2B) and 10 ms (Fig. 2C). At the overlap (t ¼ 10 ms, Fig. 2C) the stationary electric field strength distribution establishes in the structure. It corresponds to the recording image. Photoresponse of reading light pulse reaches maximum at overlap of recording and reading image. After the recording light pulse, photoresponse of reading pulse diminishes to

254

P.G. Kasherininov and A.A. Tomasov

its minimum value. When recording light pulse contains recording image and reading light intensity is homogeneous over the sample surface, photocurrent pulse at the outlet of processor is proportional to the recording image square (Fig 2C). When reading light pulse contains compared image the processor functions as image correlator.

4. OPTOELECTRONIC IMAGE CORRELATOR OF INCOHERENT LIGHT BASED ON ANALOG OPTICAL PROCESSORS Such analog computers can be applied to the development of an optoelectronic correlator aimed at image recognition [12]. For the benefit of it the surface of a processor containing n-p transition is illuminated with reference image of a necessary object. The opposite surface is illuminated with the image that needs to be compared. Further, it is performed with their mutual shift along both axes and their adjustment in scale and angular orientation. At the moment the signal at the outlet of a processor is defined by the overlap integral of these images, that is, proportional to the correlation function (CF) of reference and recognized images. At coincidence of reference image of the object and current image in size and angular orientation, the CF value reaches its maximum. The decision about the presence of a necessary object in the field of the system vision is made after the CF reaches a fixed threshold value. The time of calculation of correlation integral in this correlator is to be defined by its operation rate and is expected to be 1 ms, while the calculation process itself is finished at one step independently on dimension of images. Thus, the rate of the correlator is to be defined by the input time of the images at the surface of processor and not by time of calculating the correlation integral. As a result the processor is to work at the rate of image input, that is, real-time rate. The mutual shift of reference and recognized images, their scaling, and rotating are to be performed by means of known electronic methods, used in television and optical image correlators. The suggested image correlator is purposed to the usage in intellectual systems of technical vision, functioning of which is related to the necessity of image recognition. It can be used to solve the following problems: automatic assembly at the assembly line, face recognition, fingerprints, credit card identification via the photographs and fingerprint of the owner, prevention of the undesirable access into the special areas, navigation and astro-navigation of the flying objects in accord with ground signs and stars, automatic connecting of spacecrafts, and so on.

Optical Detectors Based on CdTe Pure Crystals

255

5. CONCLUSION Thus, pure crystals of CdTe with low impurity density (Nt < 1013 cm3) are very prospective as a material for fast optical registering media. Registering media with a cycle time t ¼ 106 s, spatial distribution 5–10 couples lines/mm, and sensitivity 102 W/cm2 are realized on the basis of CdTe. The optical, digital, and analog computers and image correlators can be realized on this basis.

REFERENCES [1] A.A. Vasilev, D. Kasasent, I.N. Kampaneez, A.V. Parfenov, Spatial Light Modulators, Radio and Communication, Moscow, Russia, 1987. [2] M.P. Petrov, I.S. Stepanov, A.V. Homenko, Fotosensitive Electrooptical Medias in Holography and Optical Information Processing, Nauka, Leningrad, Russia, 1983. [3] A.V. Boroshnev, N.F. Kovtonyuk, Russ. Appl. Phys. 6 (2000) 5–10. [4] J. Sewchun, A. Waxman, G. Warfield, Solid State Electron. 10 (12) (1967) 1165–1186. [5] W.E. Dahlke, S.M. Sze, Solid State Electron. 10 (8) (1967) 865–873. [6] M.A. Green, J. Shewchun, Solid State Electron. 17 (4) (1974) 349–365. [7] A.A. Gutkin, V.E. Sedov, Semiconductors 9 (9) (1975) 1155–1158. [8] A.Ya. Vul’, S.V. Kozyrev, V.I. Fedorov, Semiconductors 15 (1) (1981) 83–86. [9] A.Ya. Vul’, A.V. Sachenko, Semiconductors 17 (8) (1983) 865–875. [10] P.G. Kasherininov, A.V. Kichaev, A.A. Tomasov, Semiconductors 29 (11) (1995) 1092–1099. [11] P.G. Kasherininov, A.V. Kichaev, A.N. Lodygin, V.K. Sokolov, Proc. SPIE 5381 (2004) 292–301. [12] P.G. Kasherininov, A.N. Lodygin, V.K. Sokolov, Proc. SPIE 5066 (2003) 273–280. [13] P.G. Kasherininov, A.V. Kichaev, A.A. Tomasov, V.K. Sokolov, Proc. SPIE 6594 (2007) 65941G. [14] P.G. Kasherininov, A.V. Kichaev, A.N. Lodygin, A.A. Tomasov, V.K. Sokolov, Proc. SPIE 6251 (2006) 625112–625124.

AUTHOR INDEX

A Abadie, J., 80 Abe, S., 103 Abu Shama, J., 205 Adachi, S., 175 Adhiri, R., 37 Afifi, H., 93–94 Agata, Y., 103 Agostinelli, G., 195 Ahlquist, C.N., 36–37 Ahmed, M.U., 99 Ahr, M., 91, 92f, 95–96 Aitken, N.M., 44–47 Akawa, M., 107–110 Akira, T., 64 Akutagawa, W., 47, 66t Al-Allak, H.M., 197, 198 Albers, W., 26, 37 Aleksandrov, B., 14, 15t Alikhanian, A.S., 62, 63f Ali, S., 9, 12–13 Allred, W.P., 40, 66t Almeida, L.A., 85–88, 99–100 Alov, D.L., 37 Amin, N., 200–201 Amir, N., 96–97, 99–100 Amirtharai, R.M., 121–122 Amirtharaj, P.M., 36–37 Amman, J., 64 Amzil, A., 23–25 Anandan, M., 38–39, 64–65 Anand, G., 80 Anderson, G.B., 107–110 Anderson, R.J., 214–215 Ando, K., 111–112 Andrews, C., 99–100 Anil, G., 9, 12 Anthony, T.C., 187 Aoudia, A., 4, 35, 38–39, 64–65, 153, 169, 171, 172–173, 175–176, 177, 181–182 Aparo, M., 216 Apotovsky, B.A., 37, 66t, 217–218 Appell, J., 25–26

Aramoto, T., 189, 196, 197 Archbold, M.D., 197 Ard, C.K., 12, 38–39, 54, 64–65, 66t Argent, B.B., 20, 23–25 Arita, S.T., 189, 196, 197 Arlt, R., 216 Arnold, M., 204 Arnoult, A., 91, 93f Arque`s, M., 228, 229f, 230f Artemov, V.M., 37, 66t Arwin, H., 121–122, 127 Asadov, M.M., 127 Asahi, T., 37, 66t Ashen, D.J., 96, 107–110 Asher, S., 189, 196, 197, 200, 201, 202 Aspnes, D.E., 121–122, 127 Astles, M.G., 65 Audet, N., 12, 13, 14–15, 34 Augustine, F.L., 217–218 Austin, R.F., 107–110 Avetisov, I.Ch., 79 Awadalla, S.A., 64 Ayoub, M., 66t Azoulay, M., 36–37, 54, 57, 58f B Baba, S., 224, 224f, 225f Babentsov, V., 76, 77–78, 81 Babentsov, V.N., 77 Bachem, K.H., 112–113 Bach, P., 41 Badano, G., 99–100 Badaud, J.P., 78 Baffert, N., 220, 226f Bagaev, V.S., 240t Bagai, R.K., 36, 38–39, 61f, 62f, 64–65, 66t, 125, 127 Bailiang, Y., 12 Bailly, F., 93–94, 102, 107–110, 187 Bai, Y.-C., 47 Bajaj, J., 97, 98f Bakken, D.W., 30, 36–37

257

258

Author Index

Bak-Misiuk, J., 47 Balasubramanian, R., 35–36 Ballet, P., 99–100 Ballingall, J.M., 107–110 Ballutaud, D., 97, 98f BaMol, B.M., 196–197, 199 Barbe, M., 93–94, 102, 107–110, 187 Barber, H.B., 217–218 Barber, W.C., 226–227, 228f Barbier, L.M., 231f, 232f Barczy, P., 37, 39 Barmin, I.V., 79, 81 Baronenkova, R.P., 121, 122, 124, 125 Baron, T., 200–201 Barrett, H.B., 217–218 Barrioz, V., 107 Barthelmy, S., 230 Barthelmy, S.D., 231f, 232f Baruch, P., 155, 192 Barz, R.U., 43 Basu, A., 44–47, 93–94, 112–113 Ba¨tzner, D.L., 91, 124, 195, 199, 201, 207–208 Baudry, X., 99–100 Baumgartner, W., 232–233, 232f Bayban, M., 111–112 Baydjanov, I., 15t Bean, R.C., 100–102 Becker, C.R., 90–91, 107–110 Becker, U., 36–37 Becla, P., 25–26, 234–235 Beier, J., 192–193 Belas, E., 23–25, 36–37 Belaud, Y., 153, 181–182 Bellisent, R., 23–25 Bell, R.O., 41, 42–43, 214 Bell, S.L., 36–37, 56 Belousova, T.Ya., 240t Bennett, P.R., 226–227, 228f Bentz, A., 11f, 12–13 Benz, K.-W., 42–43, 44–47, 62–64, 63f, 66t, 76, 77–78, 79, 80, 81 Berchenko, N.N., 121 Berezhnoi, E., 9 Berger, H., 20, 66t, 121–122 Bergman, C., 23–25 Berkenblit, M., 28–29 Berman, S., 4–5 Berozashvili, Yu.N., 240t Besombes, L., 95–96, 95f Bettridge, V., 43 Betz, J., 196–197 Bezalel, E., 208

Bhat, I., 96, 107–110, 110f Bhat, I.B., 95–96 Bichara, C., 23–25 Bicknell, R.N., 90–91, 100–102, 107–110 Bicknell-Tassius, R.N., 90–91, 107–110 Biehl, M., 91, 92f, 95–96 Biglari, B., 77–78 Bilevych, Y.e.O., 129, 131t, 136 Birkmire, R.W., 193–194, 195–196, 197, 198, 206 Bissoli, F., 37 Black, D.R., 36–37, 66t Black, M., 214–215 Blackmore, G.W., 96–97 Bland, L.G., 36–37, 66t Blank, Z., 43 Bloedner, R.U., 38–39 Blondel, C., 228, 229f, 230f Blunier, S., 90–91, 100–102 Boeck, T., 64–65 Boggs, S., 232–233, 232f Boiton, P., 80 Bok, J., 37 Bollong, A.B., 12–13, 14, 64–65 Bolotnikov, A.E., 214–215, 235–236, 235f, 236f Bolotnikov, G., 64 Bonnefoy, J.P., 220, 226f Bonnet, D., 93–94, 187–188, 189, 198, 207 Bonnin, C., 228, 229f, 230f Boone, J.L., 47, 66t Borle, W.N., 38–39, 64–65, 125, 127 Boroshnev, A.V., 249 Bosio, A., 197, 201–202 Boue`re, A., 228, 229f, 230f Boukerche, M., 90–91, 100–102, 107–110 Bowers, K.A., 37, 214–215 Boyd, P.R., 36–37, 40 Boyle, D.S., 197 Boyn, R., 36–37 Boynykh, N.M., 121, 122, 124, 125 Braescu, L., 80 Brandon, S., 50–51, 54 Brandt, G., 36–37 Brau, M., 13 Brebrick, R.F., 20, 21, 23–25, 47, 56, 90 Brellier, D., 43 Bre´mond, G., 37, 167–168, 175–176, 177 Brenner, W., 43 Brice, J.C., 38–39 Bridenbaugh, P.M., 107–110

Author Index

Brill, G., 104, 105f, 106f Brillson, L.J., 189–190 Brinkman, A.W., 44–47, 66t, 77, 93–94, 112–113, 196, 197, 198, 208 Britt, J., 189, 196, 197 Broetz, J., 199–200 Browning, N.D., 85, 86f, 102, 103, 103f, 104, 104f Brown, P.D., 107–110 Bruder, M., 36–37, 66t, 78 Brunet-Jailly, A., 41 Brunet, P., 38–39, 40, 66t Brunett, A., 37 Brunett, B.A., 56 Brun-Lecunff, D., 121–122 Bryant, A.W., 20, 23–25 Bube, R.H., 187 Buckley, D.J., 44–47 Buck, P., 47, 66t Bult, R., 12–13 Bunnell, D., 12–13, 14 Burckhardt, D.D., 219, 219f Burdette, H., 36–37, 66t Burgelman, M., 190–191, 192–193, 195, 198, 199 Burger, A., 15t, 47, 122–123, 124–125, 214–215 Busch, M.C., 169 Butler, J.E., 22–23, 37 Butler, J.F., 37, 66t, 214–215, 217–218 Butterman, W., 5 Bykova, S.V., 37–38, 38f, 52 Bylsma, R.B., 165–166, 178 C Cadoret, R., 78 Cahen, D., 65, 208 Caillot, M., 40 Cajipe, V., 222, 222f, 223f, 224f Camarda, A.E., 64 Camarda, G.S., 214–215, 235–236, 235f, 236f Campo, M., 199–200 Canevari, V., 201–202 Cantwell, B.J., 44–47, 66t, 77, 93–94, 112–113 Cantwell, G., 47, 66t Caporasso, A.J., 43 Capper, P., 38–39, 49, 66t, 193–194, 196 Cara, C., 228, 229f, 230f Carini, G.A., 214–215 Carini, G.S., 64 Carles, J., 44–47

259

Carlson, F.M., 23–25, 54 Carlsson, L., 36–37 Carlsson, T., 208 Casagrande, L.G., 36–37, 66t Casauay, R., 7–8 Castanet, R., 23–25 Castro, C.A., 40 Catos, H.C., 124–125 Caudano, R., 102 Chaleil, T., 228, 229f, 230f Chambron, J., 217–218 Chami, A., 107–110 Chang, C.E., 49 Chang, S.L., 107–110 Chang, W.-M., 40 Chang, Y., 12 Chapuis, A., 220, 226f Charalambous, P., 4–5 Charara, J., 12–13 Charleux, M., 95–96 Chattopadhyay, K., 122–123, 124–125, 215 Chavada, F.R., 124–125 Chee, K.T., 90–91, 107–110 Chen, A.-B., 56 Chen, C.M.H., 232–233, 232f Cheng, L.J., 178 Chen, H., 64, 122–123, 124–125, 215 Chen, H.A., 64 Chen, K.-T., 47 Chen, L., 85–88, 88f, 104, 112 Chen, R., 215 Chen, S.D., 107–110 Chen, T.H., 80 Chen, W., 107–110 Chen, Y., 104, 105f, 106f Chen, Y.P., 85–88, 87f, 100–102, 103f Cherkaoui, K., 37 Chernov, M., 37 Cheung, J.T., 90–91 Cheuvart, P., 29, 40 Chevalier, N., 39, 76, 80 Chevrier, V., 76, 77–78, 81 Chew, N.G., 96–97 Che`ze, I., 97, 98f Chieux, P., 23–25 Chi, M.F., 85, 86f, 104 Chizhevskaya, S.N., 23–25 Chizhikov, D., 7–8 Choi, B.W., 23–25, 54 Chonko, J., 232–233, 232f Chopra, K.L., 197 Choubey, A., 44–47, 93, 112–113

260

Author Index

Chou, K.S., 100–102, 107–110 Chouraqui, P., 217–218 Chou, R.L., 100–102 Chowdhury, S., 219, 219f Chow, P., 90–91 Christensen, F.E., 232–233, 232f Ch. Steer, 50 Chua, P.L., 178 Chudakov, V.S., 37, 66t Chung, H., 80–81 Churbanov, M., 13 Chu, S.N.G., 96–97 Cibert, J., 91, 92f, 107–110 Cirignano, L.J., 214–215, 226–227, 228f Clark, A., 38–39 Clark, E., 202–203 Clark, J., 4–5 Clauws, P., 198 Clayman, B.P., 25–26, 234–235 Coates, W.G., 38–39 Cobb, S., 80 Cobb, S.D., 80 Coche, A., 40 Cockrum, C.A., 100–102 Cohen-Solal, G., 93–94, 102, 107–110, 187 Cola, A., 111–112 Cole, H.S., 96 Cole, S., 25–26 Collins, E.E., 47 Collins, W.E., 122–123 Colombo, L., 13, 40 Compaan, A.D., 189–190, 192, 199, 205 Comtois, R., 40 Conibeer, G.J., 193–194 Contreras, M.A., 113–114 Cook, W.R., 232–233, 232f Cordes, H., 121–122 Cornet, A., 40, 41 Corregidor, V., 76, 77–78, 81 Corsini-Mena, A., 48 Cortes, R., 95–96, 107–110 Cossette, M., 14–15, 34 Coupat, B., 78 Coutts, J., 205 Cowache, P., 197 Craig, W., 232–233, 232f Crestou, J., 37, 56, 57f Cr´etolle, J., 228, 229f, 230f Crimes, G., 96–97 Crochet, M.J., 49–50 Cronin-Golomb, M., 178 Cross, E., 15t

Cruz-Campa, J.L., 93–94, 94f Cui, Y., 214–215, 235–236, 235f, 236f Cullis, A.G., 96–97 Cummings, J.R., 230, 231f, 232f Cunningham, D., 189 Currie, M.C., 36–37 Cuzin, M., 223–224

D Dal’Bo, F., 85–88, 89f Dalecki, W., 14 Daly, F., 228, 229f, 230f Danaher, J.W., 122, 126 Danylenko, S.G., 131t, 136 Das, B.N., 107 Datta, S., 107–110 Daudin, B., 121–122 Davies, K., 189 Davydov, A.A., 44–47 Dean, B., 40, 66t Dean, B.E., 40, 66t De Carolis, M., 216 Degrave, S., 193, 198 DeHart, C., 113–114, 189, 196, 197 Delaye, P., 76, 77–78, 81, 153, 156, 157, 158, 181–182 del Cueto, J.A., 207–208 de Lyon, T.J., 100–102 DeMaria, A.J., 243 Demtsu, S., 200 Derby, J.J., 37–38, 50–51, 52–53, 53f, 54, 55 Dereniak, D.L., 217–218 Desages, F., 228, 229f, 230f Destefanis, G., 99–100 de Vos, A., 155, 192 Dharmadasa, I.M., 121–122, 199 Dharmasena, K.P., 54 Dhar, N., 104, 105f, 106f Dhar, N.K., 100–102, 103, 104, 104f Dhere, R., 205 Dhere, R.G., 189, 196, 197 Dian, R., 41, 77–78 Didier, G., 4, 14–15, 31–32, 33–34, 33f, 41, 42–43, 60f Dieguez, E., 39, 76, 77–78, 80, 81 Dierre, F., 43 Di Marzio, D., 36–37, 66t Dinan, J.H., 100–102, 104 Ding, R.J., 85–88, 88f, 112 Dobbyn, R.C., 36–37

Author Index

Do¨beli, M., 201, 207–208 Dobrotvorskaya, M.V., 121 Dobson, K.D., 208 Domagala, J., 47, 121–122, 125, 127 Dong, C.H., 23–25 Dost, S., 12, 13, 42–43, 79 Doty, F.P., 37, 66t, 214–215, 217–218 Doty, M., 196–197 Doumae, Y., 47 Dragnev, T., 216 Drapala, J., 10f, 11f, 15t Drayton, J., 192, 199 Drigo, A.V., 107–110 Driver, M.C., 36–37, 66t Druilhe, R., 56 Drygybka, S., 131t, 136 Dubowski, J.J., 90–91, 107–110 Duda, A., 189, 196, 197, 200, 201, 205 Dudley, M. Jr., 36–37, 39, 66t, 77, 80–81 Duffar, T., 35–36, 39, 39f, 76, 77–78, 80, 81 Dupont, S., 49–50 Dupret, F., 49–50 Durose, K., 44–47, 107, 121–122, 125, 127, 155, 195–196, 197, 198, 201 Durst, F., 42–43 Dusserre, P., 35–36, 39, 39f, 76, 77–78, 80, 81 Dutton, D., 38–39, 66t Duvaut, P., 99–100 Dzitko, H., 228, 229f, 230f E Ebina, A., 124–125 Ebling, D.G., 79, 80 Eclancher, B., 218 Edwards, K., 52–53, 53f Edwards, P.R., 155, 195–196, 197, 198, 201 Egaas, B., 113–114 Eger, D., 44, 66t Egorov, A.V., 42–43, 79, 80, 81 Eguchi, K., 103 Ehsani, H.E., 95–96 Eiche, C., 44–47, 77, 79, 80 Eisen, Y., 217–218, 218f Eissler, E., 56 El Hanany, U., 29, 40, 66t Elli, M., 48 El Mokri, A., 33–34, 43, 60f Emen, H., 36–37 Emery, K., 189 Engdahl, J.C., 219, 219f Engel, A., 79, 121–122, 125

261

Epure, S., 80 Erickson, J., 15t, 122–123 Er-Raji, S., 78 Eser, E., 196, 206 Eymery, J., 91 F Fahrenbruch, A.L., 187, 189–190, 192, 199, 200, 201 Fang, R., 21, 47 Fanning, T., 36–37, 66t Farag, B.S., 200 Farrell, S., 104, 105f, 106f Farrow, R.F.C., 90 Faschinger, W., 95–96 Fauler, A., 76, 77–78, 81 Faurie, J.P., 85–88, 87f, 90–91, 100–102, 107–110, 233–234 Favier, J.-J., 50 Fawaz, A., 12–13 Feichouk, P., 23–25 Feichuk, P.I., 129, 131t Feigelson, R.S., 36–37, 38–39, 53, 125 Feldewerth, G., 12–13, 14 Feldman, R.D., 96–97, 107–110 Feltgen, T., 77 Fenimore, E.E., 231f, 232f Ferah, M., 20, 23–25 Ferekides, C., 189, 196, 197 Ferekides, C.S., 93–94, 194–196, 197, 200 Ferret, P., 99–100 Fesh, R., 37 Fesh, R.N., 124–125, 126 Feth, S., 77, 78, 124–125 Feuillet, G., 95–96, 107–110 Feychouk, P., 23–25 Feychuk, P.I., 121, 123, 124–125, 126, 130–135, 131t, 136 Fiederle, M., 15t, 42–43, 44–47, 62–64, 63f, 76, 77–78, 79, 80, 81, 112–113 Fischer, A.G., 29 Fischer, B., 178 Fischer, R., 100–102 Flege, S., 199–200 Flisch, A., 78 Floeder, W., 100–102 Flores, H., 222, 222f, 223f, 224f Fomin, A.V., 120, 121, 122, 124, 126, 127–128 Fontaine, C., 85–88, 89f

262

Author Index

Foreman, B.A., 47, 66t Foucher, C., 56 Fougeres, P., 37, 66t Fournier, J.P., 78 Franc, J., 23–25, 36–37, 121, 123, 131t, 134, 137–139 Franc, Y.a., 131t, 134–135 Fries, E., 78 Fritsche, J., 200–201 Froment, M., 95–96, 107–110, 197 Fthenakis, V.M., 209, 210 Fu, F., 85, 86f, 104 Fukazawa, Y., 233, 234f Fulop, G., 196–197 Furdyna, J.K., 152 Fu, T.W., 49 Fu, X.L., 85–88, 88f, 112 G Gafni, G., 36–37, 54, 57, 58f Gaillard, J.P., 85–88, 89f Galassini, S., 205 Galazka, R.R., 44–47 Galkina, O.S., 121 Gallet, J., 41 Galloway, S.A., 197, 198 Garandet, J.-P., 39, 76, 77–78, 80 Garg, A.K., 36, 66t Garmire, A.M., 151–152, 155–156, 158 Garmire, E.M., 153, 182 Garstein, E., 96–97, 99–100 Gasgnier, M., 56, 57f Gaspard, J.P., 23–25 Gaugash, P., 122, 124–125, 126, 127 Gauneau, M., 171, 172–173, 175–176 Gautam, M., 124–125 Gehrels, N., 230, 231f, 232f Geibel, C., 47, 66t Gelfgat, Y.L., 79 Gentet, M.C., 220, 220f, 221f Gentile, P., 91 Genzel, C., 36–37, 66t George, M.A., 47 Gerasimenko, V., 13 Gerfault, L., 220, 220f, 221f Gessert, T.A., 189, 196, 197, 200, 201, 202, 205 Geyling, F.T., 49–50 Ge, Y.-R., 78 Ghaddar, C.K., 42–43 Ghandhi, S.K., 96

Ghijsen, J., 124–125 Giacometti, M., 80, 81 Giacometti, N., 39, 39f Giakos, G.C., 217–218 Gibson, P.N., 197 Giess, J., 96–97 Gilath, C., 217–218 Gille, P., 38–39 Gilles, B., 95–96, 95f Gilles, D.C., 42–43 Gilles, N.C., 90–91, 107–110 Girard-Franc¸ois, A., 107–110 Giriat, W., 152 Glagoleva, N.N., 23–25 Glasov, V.M., 23–25 Glasper, J.L., 96–97 Glass, A.M., 156, 165–166, 178 Glasser, F., 220, 223–224 Glass, H.L., 30, 36–37 Glenn, C., 64 Gloekler, M., 190–191, 192, 194–195 Gobi, Y., 107–110 Goffart, J., 23–25 Golacki, Z., 44–47, 66t Goldsmith, P., 76 Golyshev, V.D., 37–38, 38f, 52 Gomagala, J.Z., 47, 66t Goodlet, G., 197 Goorsky, H., 37 Goorsky, M., 15t, 56 Goorsky, M.S., 122–123 Gordon, E.E., 222, 222f, 223f, 224f Go¨ro¨g, T., 29, 37 Gorska, M., 44–47 Gosney, J.J., 49 Gosney, J.J.G., 38–39 Gossla, M., 195 Goto, H., 111–112 Gough, J.S., 96–97 Grammond, L., 189 Grandpierre, G., 171, 172–173 Grasza, K., 44–47, 121, 127 Gravey, P., 153, 154, 157, 158, 160–161, 169, 175–176, 178, 179, 181–182 Gray, A., 100–102 Grebenyuk, N.N., 121 Greenberg, J.H., 20, 21f, 22–23, 22f, 31–32, 34, 62–64, 63f, 77 Greenlaw, D.K., 90–91 Green, M.A., 189, 210 Greiffenberg, D., 112–113 Grein, C.H., 85, 86f, 104

Author Index

Grenier, S., 12, 13 Grigor’yev, V., 14 Grill, R., 23–25, 36–37 Grindlay, J., 64 Grochocki, A., 79 Grodzicka, E., 44–47 Groenert, M., 99–100 Gronkowski, J., 107–110 Groza, M., 214–215 Gryshchuk, V., 216 Guellil, Z., 156, 157, 158 Guergouri, K., 56 Guillemaud, R., 220, 220f, 221f Gumenyuk, O.R., 130–133, 131t, 136 Gunshor, R.L., 107–110 Gupta, S.C., 124–125 Guskov, A.N., 77 Guskov, V.N., 34, 62–64, 63f Guskov, V.S., 62, 63f Gwynn, P.J., 123 H Haas, J., 12, 13 Hack, F.A., 240t Hadj-Ali, M., 169, 175–176, 177 Ha¨drich, M., 195 Hage-Ali, M., 11f, 12–13, 37, 50, 66t, 77–78, 122, 127, 214, 218 Hahnert, I., 124 Ha¨hnert, I., 66t Hailey, C., 232–233, 232f Hails, J.E., 107–110 Hall, E., 4–5 Halliday, A., 7 Halliday, D.P., 44–47, 155, 195–196, 197, 198, 201 Hall, Y.J., 160–161 Halsted, R.E., 40, 125 Hamilton, W., 217–218 Hamilton, W.J., 222, 222f, 223f, 224f Hanafusa, T., 189, 196, 197 Han, M.S., 122–123 Hansen, T., 209 Harris, J.E., 38–39, 66t Harrison, F.A., 232–233, 232f Harsch, W.C., 47, 66t Hartmann, J.M., 95–96 Hart, R.A., 243 Hasanen, T.O., 32–33, 32f Hasegawa, B.H., 226–227, 228f Hase, T.P., 93

Hase, T.P.A., 197 Hassani, S., 37, 48, 48f Havulinna, J., 222f, 225, 226f, 227f Hawkey, J.E., 40, 66t Hayer, E., 23–25 Hay, K.A., 100–102 Healy, M., 194 Hearne, S., 193–194 Hegedus, S.S., 195–196, 197, 198, 200, 207–208 Hein, J., 7 Heinrich, H., 93 He, L., 85–88, 88f, 107–110, 112 Hellman, C., 217–218 Hellwarth, R.W., 166 Hemmat, N., 41 Hemmi, T., 193–194 Henneberger, F., 107–110 Hermon, H., 15t Hermon, R.B., 37 Herrit, G.L., 242 Hettich, H.L., 36–37 Heumann, F.K., 40 Heurtel, A., 20, 23–26, 43 He, X.Z., 107–110 Hibino, K., 189, 196, 197 Higuchi, H., 189, 196, 197 Hirata, K., 36–37 Hiroyuki, S., 64 Hirschfeld, D., 78 Hirsch, H., 5–7, 9, 12–13 Hishikawa, Y., 189 Hisshion, R., 7–8 Hlidek, P., 36–37 Hnativ, I.I., 131t, 136, 137–139 Hobgood, H.M., 41 Ho¨chst, H., 199–200, 201–202 Hodes, G., 197, 208 Hoffmann, J., 7–8 Hoffmann, N., 107–110 Hoke, W.E., 96–97 Holland, A.J., 198 Holland, L., 12 Holzberg, F., 91 Homenko, A.V., 249 Homewood, K.P., 198 Ho¨mmerrich, U., 30 Hong, A., 64 Hooper, S., 40, 66t Horch, R., 30 Horeau, B., 228, 229f, 230f Hori, S., 43

263

264

Author Index

Horning, R.D., 107–110 Ho¨schl, P., 23–26, 36–37, 47, 121, 123, 131t, 134 Hossain, A., 214–215, 235–236, 235f, 236f Huang, C., 44–47 Huang, Q., 107–110, 111–112 Huang, W., 32–33, 32f Hullinger, D., 230, 231f, 232f Hutchins, J.W., 100–102 Hutchins, M.A., 78 Hu, Z., 47 Hyodo, K., 225–226 I Ichimiya, T., 126 Ickert, L., 49 Ido, T., 111–112 Igaki, K., 47 Ignatov, S.M., 223–224 Ikegami, S., 30, 194, 195 Imhoff, D., 43, 56 Inano, S., 38–39 Inatomi, Y., 42–43, 79 Indenbaum, G.V., 121, 122, 124, 125 Inoue, M., 124–125 Inuischi, Y., 41 Ipser, H., 23–25 Ireland, P.J., 122–123 Irvine, S.J.C., 77, 96–97, 98f, 99, 107–110 Ishikawa, Y., 12, 47, 106f, 107–110 Isselin, S., 99–100 Isshiki, M., 12, 13, 47, 103, 106f, 107–110 Ito, M., 209–210 Ivanits´ka, V.G., 121, 123, 131t, 134 Ivanitskaya, V.G., 130–135, 131t Ivanov, A., 47 Ivanov, A.A., 37 Ivanov, V.Y.u., 121–122, 125, 127 Ivanov, Yu.M., 20, 26, 36–37, 66t Iwanaga, H., 125 Iwase, Y., 41, 66t, 125 Iwata, K., 226–227, 228f J Jacobs, K., 22–23, 36–37, 64–65, 107–110 Jacobs, R.N., 99–100 Jaegermann, W., 93–94, 107, 199–201 Jahnke, A., 218 Jaime-Vasquez, M., 99–100 James, M.S., 37

James, R., 15t James, R.B., 56, 64, 122–123, 214–215, 233–234, 235–236, 235f, 236f James, R.W., 56 Janik, E., 20, 23–25 Jasinski, T., 49 Jedrzejczak, A., 44–47 Jeffers, K.S., 96–97, 107–110 Jennings, J., 7–8 Jensen, D.G., 193–194, 195 Jianbin, W., 80 Jiang, Q., 44–47, 93–94, 112–113 Ji, R., 42–43, 79 Joerger, W., 44–47, 77, 78, 79, 80 Johnson, C.J., 40, 56, 66t, 240t Johnson, D., 90–91 Johnson, D.R., 197 Johnson, R.L., 111–112 Johnson, S.M., 100–102 Johnston, S., 200, 201 Jones, C.A., 43, 77 Jones, C.L., 38–39, 49, 66t Jones, E.W., 107 Jones, G.R., 90 Jordan, A.S., 20, 23–25, 63f Jost, J.M., 63f Jung, H.S., 85, 86f, 104 Juza, P., 95–96 J.Y. Moisan, 179 K Kaitasov, O., 37, 56, 57f Kaldis, E., 78 Kaliszek, W., 215 Kamae, T., 233, 234f Kampaneez, I.N., 249 Kampmann, A., 95–96, 107–110 Kanda, S., 224, 224f, 225f Kanevsky, V.M., 37, 66t Kang, T.W., 107–110, 111f, 122–123 Kany, F., 95–96 Karman, M., 218 Karpenko, V.P., 223–224 Karpov, V.G., 192, 199 Kasap, S.O., 222, 222f, 223f, 224f Kasasent, D., 249 Kasherininov, P.G., 250–253, 254 Kataoka, J., 233, 234f Katayama, S., 38–39 Kato, E., 193–194 Kato, K., 209–210

Author Index

Katty, A., 35–36, 38–39, 40, 66t katz, J., 100–102 Kaur, I., 197 Kawasaki, T., 56, 57, 58f, 233–234 Kawayoshi, H., 90–91, 107–110 Kazandjian, A., 218 Keane, J.C., 189, 196, 197 Kelly, M.K., 189–190 Kendall, E.J.M., 25–26 Kennedy, J.J., 36–37 Kennedy, J.T., 243 Kenworthy, I., 38–39 Kestigian, M., 12, 64–65 Khan, A.A., 40, 66t Khanin, E., 96–97, 99–100 Khattak, C.P., 36–37, 38–39 Kheng, K., 95–96, 95f Khodier, S.A., 200 Khoshnevisan, M., 90–91 Khrypunov, G., 91, 204, 205 Kibel, M.H., 123 Kichaev, A.V., 250–253 Kichimi, T., 209–210 Kiefer, J.E., 240t Kikuchi, K., 106f, 107–110 Kilday, D., 189–190 Kim, H.C., 209 Kim, J.M., 58, 59f Kim, J.S., 58, 59f Kim, M.D., 107–110, 111f Kim, S.U., 58, 59f Kim, T.W., 122–123 Kimura, T., 175 King, D.E., 202 Kinnari, J., 220 Kirste, L., 112–113 Kiseleva, K.V., 20, 41 Kisker, D.W., 85–88, 86f, 96–97, 107–110 Klein, A., 93–94, 107, 199–201 Klein, M.B., 151–152, 153, 155–156, 158, 160–161, 163, 182 Klevkov, U.V., 20 Klevkov, V.Yu., 30, 40 Klimenko, J.A., 121 Kloess, G., 44–47, 77 K. MaOek, 121, 123, 131t, 134 Kobayashi, M.Z., 20 Koblova, M.M., 240t Kochanowskab, D., 215, 235–236, 235f, 236f Koebal, J.M., 169, 175–176, 177 Koebel, J., 11f, 12–13

Koebel, J.M., 66t, 77–78 Koebel, J.-M., 37, 50 Ko, E.I., 96–97 Koenig, A., 220, 226f Koglin, J., 232–233, 232f Kohman, K.T., 214–215 Koike, K., 111–112 Kokubun, M., 233, 234f Kolb, E.D., 43 Kolchin, A.A., 37 Kolesnikov, N.N., 37 Koliwad, K., 100–102 Kolodziejski, L.A., 107–110 Komar, V.K., 121, 131t Kometani, T.Y., 96–97 Komisarchik, M.S.h., 126 Komoto, K., 209–210 Konagai, M., 200–201 Konak, C., 25–26, 47 Kondon, N., 121–122 Konkel, W.H., 36–37, 66t Ko¨ntges, M., 192–193 Koo, B.H., 106f, 107–110 Kopach, O., 23–25 Kore-eda, T., 124–125 Koschinsky, A., 7 Kosyakov, V., 9 Kota, A.K., 80 Kotina, I.M., 122 Koul, S.K., 124–125 Kou, S., 58, 59f, 60f Kovalevski, S., 9 Kovtonyuk, N.F., 249 Kowalczyk, L., 47 Kowalski, B.J., 124–125 Koyama, A., 37, 66t Koziejowska, A., 125 Kozin, K., 9 Kozin, L., 9 Kozuka, Z., 15t Kraft, D., 199–200 Krapukhin, V.V., 79 Kravetskiy, M.Y.u., 121, 122 Krevs, V.E., 121 Krimm, H., 230 Krimm, H.A., 231f, 232f Kropp, E., 64–65 Kuchar, L., 10f, 11f, 15t Kudo, K., 42–43, 79 Kukhtarev, N.V., 152, 154 Kulwicki, B., 20 Kumar, V., 36, 42–43, 66t

265

266

Author Index

Kumazawa, S., 189, 196, 197 Kunz, T., 44–47, 77, 78 Kuppurao, S., 50–51, 54 Kurdesau, F., 91, 204 Kuroda, S., 113 Kurokawa, K., 209–210 Kusowski, M., 14 Kutcher, S.W., 215 Kuwamoto, H., 44–47 Kwon, M.S., 107–110, 111f Kyle, N.R., 36–37 L Laasch, M., 44–47, 66t, 77, 78 Lacroix, S., 39, 39f Lajzerowicz, J., 37, 66t Lakeenkov, V.M., 52 Lakus-Wollny, K., 93–94, 200 Lalev, G.M., 103 Lamakina, G.A., 187 Lamb, D., 107 Lambert, B., 171, 172–173, 175–176 Lam, T., 15t Landolt, O., 121, 122–123 Landsberg, P.T., 155, 192 Landwehr, G., 90–91, 107–110 Lane, D.W., 193–194 Lange, M.D., 100–102, 107–110 Langer, J.M., 152 Larson, D.J. Jr., 36–37, 66t, 80–81 Latorre, B., 56 Laudise, R.A., 43 Laukka, P., 220 Laulan, C., 179 Launay, J.C., 76, 77–78, 80, 81, 153, 156, 157, 158, 166, 168, 171, 172–173, 179, 181–182 Laurent, P., 228, 229f, 230f Lavine, M.C., 124–125 Lavocat, P., 228, 229f, 230f Lay, K.Y., 40, 66t Lebrun, F., 228, 229f, 230f Lee, C., 214–215 Leech, P.W., 123 Lee, C.K., 42–43 Lee, J.Y., 107–110, 111f Lee, M., 64 Lee, M.B., 36–37, 66t Lee, S.B., 58, 59f Lee, T.S., 58, 59f Le Gouge, Y., 56

Legros, R., 43 Lehoczky, S.L., 37, 39, 47, 77, 78 Lemonias, P.J., 96–97 Lent, B., 42–43, 79 Leo, G., 107–110 Leombruno, R., 41 Leon-Gits, S., 78 Leppa¨nen, A., 220, 222f, 225, 226f, 227f Leray, J.P., 228, 229f, 230f Le Si Dang, E., 107–110 Le Si dang, Y., 85–88, 89f Leung, C., 90–91, 107–110 Levi, D.H., 189, 196, 197, 202 Lewandowski, R.S., 40 Leyarovska, N., 189–190 Leybov, V.A., 25–26 Liang, S., 5–7, 9, 12–13 Liao, P.K., 40, 52–53, 66t Liaw, I.R., 107–110 Li, C., 39 Lie, C., 37 Ligeon, 107–110 Li He, J., 104 Li, J., 107–110 Li, L., 50, 214–215 Limousin, O., 228–230, 229f, 230f Lincot, D., 93–94, 95–96, 107–110, 121, 122, 187, 197 Lin, L., 107–110 Lin, M.S., 100–102 Li, Q.-F., 36–37, 127 Li, S., 111–112 Lischka, K., 93 Liu, C.H., 196–197 Liu, D., 122–123, 124 Liu, D.T.J.H., 178 Liu, H.C., 47 Liu, J., 107–110 Liu, J.C., 55 Liu, X., 189–190 Liu, Y., 12, 13, 42–43, 79 Li, X., 199–200, 201 Li, Y.J., 85–88, 88f, 112 Lloyd, G., 93 Lmai, F., 12–13 Lodygin, A.N., 250, 252–253, 254 Lomakina, G.A., 187 Long, G.G., 36–37 Longo, M., 107–110 Lopez, M., 13 Lopez Otero, A., 93 Lordkipanidze, D.Sh., 240t

Author Index

Lorenz, M.R., 20, 40, 43, 63f, 125 Lorenz, N., 195 Lourenc¸o, M.A., 198 Loutts, G., 30 Lovergine, N., 107–110, 111–112 Lo, Y., 90–91, 100–102 Lubin, E., 217–218 Lu, F., 214–215 Luft, B.D., 119–120 Lu, J., 107–110 Luke, J.S., 64 Lukiyanchuk, E.M., 131t, 139 Lunacek, J., 10f, 15t Lunn, B., 43 Luo, C.P., 107–110 Lu, P.Y., 96–97 Lusakowska, E., 47 Luschtiz, J., 93–94 Lush, G.B., 202–203 Lusson, A., 4, 14–15, 33–34, 35, 37, 38–39, 43, 48, 48f, 60f, 64–65, 97, 98f Lu, Y.C., 38–39, 125 Lyakhovitskaya, V., 65 Lynch, R.T., 48 Lyons, L.E., 122, 124, 126 Lyubchenko, A.V., 126 Lyubomirsky, I., 65 M MacDougal, M.H., 197, 198 Mackett, P., 38–39, 66t Mack, S., 195 Madejski, G., 233, 234f Madsen, K., 232–233, 232f Maeda, K., 125, 127 Maekawa, T., 56 Magee, T., 90–91 Magee, T.J., 107–110 Magee, T.L., 90–91 Magnan, A., 78 Magnea, N., 85–88, 89f, 91, 95–96, 200–201 Mahadik, N., 99–100 Mahajan, S., 96–97 Mahavadi, K.K., 90–91, 100–102, 107–110 Mahgerefteh, D., 166 Maier, H., 36–37, 47, 66t Mainprize, J., 222, 222f, 223f, 224f Majewski, J., 44–47 Makhnyuk, V.I., 121, 122 Makishima, K., 233, 234f Makowski, J., 44–47

267

Maksimovskii, S.N., 30, 40 Manasevit, H.M., 96 Manchini, A.M., 107–110, 111–112 Mannisto, O., 220 Mansour, B.A., 200 Marbeuf, A., 20, 23–25, 56 Marchal, J.M., 49–50 Marchenko, M.P., 37–38, 38f, 52 Marfaing, Y., 38–39, 41–42, 44, 45f, 56, 97, 98f, 153, 169, 171, 172–173, 175–176, 177, 181–182 Margaritondo, G., 189–190 Mar, H.A., 90–91, 107–110 Mariano, A.N., 124–125 Mariette, H., 91, 95–96, 95f Marinskiy, D., 194–196, 197, 200 Marinsky, D., 93–94 Markov, E.V., 44–47 Markunas, J., 99–100 Markwardt, C.B., 230, 231f, 232f Marrakchi, G., 37, 175–176, 177 Marsal, L., 95–96, 95f Marsilac, S., 205 Martel, G., 153, 169, 175–176, 181–182 Martinez-Thomas, C., 35, 51, 51f, 52 Martrou, D., 91, 95–96 Maruyama, K., 56 Marychev, V., 12–13 Marychurch, M., 122, 126 Marzio, D.D., 36–37, 66t Masa, Y., 56, 57, 58f, 125, 233–234 Maslakovets, Yu.P., 187 Maslova, L.V., 223–224 Mason, J.E., 209 Masset, G., 100–102 Masumoto, K., 36–37, 103 Matherson, K.L., 217–218 Mathews, N.R., 203 Mathew, X., 203 Mathey, P., 166 Mathieu, J-C., 23–25 Mathot, G., 102 Mathy, F., 228, 229f, 230f Matioukhin, D.G., 42–43, 79, 80 Matsumoto, H., 194, 195 Matsumoto, K., 36–37, 41 Matsumura, N., 90–91, 99–100, 107–110 Matveev, O.A., 20, 40, 223–224 Matyi, R., 77 Matz, R., 218 Mauguen, V., 228, 229f, 230f Mawdsley, G., 222, 222f, 223f, 224f

268

Author Index

Ma, X., 122–123, 215 Maximovsky, S.N., 20 Mazoyer, V., 156, 157, 158, 171, 172–173 Mazzamuto, S., 197 Mazzer, M., 107–110 McCandless, B.E., 193–194, 195–196, 197, 198, 200, 207–208 McClure, J.C., 202–203 McDevitt, S., 40, 66t McLean, A.B., 121–122, 199 Mc Lean, E.O., 36–37 McMahon, T.J., 207 Medvedev, S.A., 20, 30, 40, 41 Meerson, E., 15t Meignier, F., 228, 229f, 230f Meiling, G.S., 41 Meinhardt, J., 44–47 Merle D’Aubigne´, S., 85–88, 89f Mestais, C., 220, 220f, 221f, 226f Metallkd, Z., 121–122 Metzger, W.K., 200 Metzner, H., 195 Meyers, P., 196–197 Meyers, P.V., 206–207 Meyers, T.H., 90–91 Meyers, T.J., 217–218 Migal, V.P., 121 Miki, T., 47 Millerd, J., 151–152, 155–156, 158 Millerd, J.E., 182 Million, A., 85–88, 89f, 90, 99–102, 104 Milnes, A.G., 122, 124–125, 126, 127 Mimila Arroyo, J., 93–94 Mimura, K., 12 Minami, K., 113 Min, J., 122–123, 124 Minoru, F., 64 Mitani, T., 233, 234f Mito, Y., 224, 224f, 225f Mitzuma, K., 126 Miura, N., 93 Miyamoto, T., 201–202 Miyazaki, K., 36–37 Mizetskaya, I.B., 124–125, 131t, 139 Mizuno, M., 201–202 Mochizuki, K., 36–37, 41, 47, 125 Mo¨ck, P., 64–65 Mohan, G., 125, 127 Moine, O., 153, 169, 181–182 Moisan, J.Y., 153, 169, 175–176, 177, 181–182 Mokili, B., 95–96, 107–110, 197 Mokri, A., 4, 14–15

Molinie0 , F., 228, 229f, 230f Mo¨ller, M.-O, 36–37, 66t Monberg, E.M., 49–50, 168 Mondoloni, C., 56 Monfroy, G., 90–91, 100–102 Monnet, O., 220, 220f, 221f, 226f Montemont, G., 220, 220f, 221f Montgometry, H.C., 125 Mopas, E., 189 Moravec, P., 23–25, 36–37, 121, 123, 131t, 134–135, 137–139 Morel, D.L., 93–94, 194–196, 197, 200 Morgan, S.H., 122–123 Moriarty, T., 200, 201 Morimoto, I., 125 Moritz, R., 36–37 Morkoc, H., 100–102 Morris, G.C., 124 Morris, G.D., 122, 126 Motakef, S., 42–43, 80 Motegi, T., 42–43, 79 Mozer, W., 40 Mu¨hlberg, M., 23–25, 25f, 26f, 36–37, 54, 64–65, 66t, 125 Mu¨Iler, G., 79 Mullin, J.B., 41, 43, 77, 96–97, 107–110 Mullin, S.J.C., 96 Mullins, J.T., 44–47, 93–94, 112–113 Munirathnam, N., 4–5, 9, 12–13 Munoz, V., 35, 51, 51f, 52 Murotani, T., 90–91 Murozono, M., 189, 196, 197 Muukkonen, K., 220 Mycielski, A., 47, 215, 235–236, 235f, 236f Myers, T.H., 100–102 Mykytiuk, A., 4–5 N Nair, J., 208 Nakagawa, K., 125, 127 Nakajima, M., 189, 196, 197 Nakamura, Y., 107–110 Nakano, A., 194, 195 Nakayama, N., 194, 195 Nakazawa, K., 230, 233, 234f Nalivayko, D.P., 121 Namba, C.S., 240 Narula, R.C., 36, 61f, 62f, 66t Natarovsky, A.M., 62, 63f Naumann, R.J., 49 Naumov, G.P., 187

Author Index

Ndap, J., 15t Ndap, J.-O., 37, 122–123, 215 Nellist, P.D., 102 Nemirovsky, Y., 96–97, 99–100 Neubert, M., 64–65 Neu, G., 93–94 Neugebauer, G.T., 53 Ng, K.K., 192 Nichols, D., 40, 66t Nicoll, F.H., 196 Nieke, H., 30 Niemegeers, A., 190–191, 192–193, 199 Niimi, T., 126 Nikolaev, I.V., 240t Nikolaev, O.V., 187 Niles, D.W., 199–200, 201–202 Niraula, M., 103 Nishijima, Y., 96–97 Nishino, H., 96–97, 100–102, 107–110 Nishio, J., 189, 196, 197 Nishitani, K., 90–91 Nitsche, R., 41, 47, 66t Nitzsche, R., 77–78 Nobel, D., 19, 20, 34, 40 Noda, K., 103 Nolan, J.F., 206–207 Nollet, P., 192–193, 198 Nolte, D.D., 168 Nomachi, M., 233, 234f Norman, P., 90–91 Norton, P.W., 12, 64–65 Nouchi, P., 166 Noufi, R., 113–114, 200, 205 Nouhi, A., 100–102 Nunoue, S.Y., 193–194 O O’Connor, N., 189 Oda, O., 36–37, 66t Ohashi, N., 47 Ohata, K., 193–194, 195 Ohmori, K., 224, 224f, 225f Ohmori, M., 41, 66t, 125 Ohnishi, H., 103 Ohno, R., 41, 66t Ohno, Y., 225–226 Ohshima, T., 90–91, 99–100, 107–110 Ohyama, H., 189, 196, 197 Ohyama, T., 47 Ojebuoboh, F., 7–8 Okada, Y., 230

Okamoto, N.L., 85, 86f, 104 O’Keefe, E., 38–39, 66t Okhata, R., 90–91 Okrepka, G.M., 131t Olego, D.J., 233–234 Olsen, R., 214–215 Olsen, R.W., 214–215 Olson, D.H., 165–166, 178 Omeltchouk, A.R., 47 Omeltchuk, A.R., 121–122, 125, 127 Omura, H., 189, 196, 197 Oonuki, K., 233, 234f Orlova, G.M., 126 Orlov, Y.u.F., 126 Orlowski, B.A., 124–125 Otsuka, N., 107–110 Otsuka, S., 15t Ouimette, D.R., 214–215 Ou, S.S., 197 Ouvrier-Buffet, P., 220, 226f Ouyang, H., 52 Ozaki, T., 125 Ozkul, C., 153, 154, 157, 158 ¨ zsan, M.E., 194, 196–197 O P Painter, J.D., 194 Palekis, V., 93–94, 194–196, 197, 200 Palik, E.D., 187 Palmer, D.M., 231f, 232f Palosz, W., 15–16, 47, 80, 121, 127 Panchouk, O., 23–25 Panchuk, O., 23–25 Panchuk, O.E., 124–125, 126 Panchuk, O.O., 124, 126, 131t Pandya, D.K., 197 Pan, J., 190–191, 192, 194–195 Pankove, J.E., 240 Pantsar, T., 220, 222f, 225, 226f, 227f Paorici, C., 29, 37, 47–48 Parfeniuk, C., 50 Parfeniuk, C.L., 30, 36–37 Parfenov, A.V., 249 Parikh, V.Y., 205 Parisi, J., 192–193 Park, I.H., 58, 59f Park, M.J., 58, 59f Parrot, J.E., 155, 192 Parsons, A., 230, 231f, 232f Partanen, J.P., 166 Parthier, L., 64–65, 107–110

269

270

Author Index

Partovi, A., 151–152, 153, 155–156, 158 Pashaev, E.M., 37, 66t Patriarche, G., 56, 107–110 Patsekina, G.V., 122 Patterson, M.H., 199 Patterson, R., 124 Patterson, V.H., 121–122 Pauliat, G., 166, 168 Pellegrino, J., 99–100 Pelliciari, B., 41, 43 Pelosi, C., 47–48 Pelosini, L., 48 Pennycook, S.J., 102, 103, 104f Perevoshchikov, V.A., 119–120, 125, 128 Perkowitz, S., 25–26, 234–235 Perlins, C.L., 113–114 Perry, D.L., 217–218 Peter, L.M., 197 Petrov, M.P., 153, 178, 249 Petroz, G., 220, 220f, 221f Peyla, P., 91, 92f Pfeiffer, M., 54 Pham Van, K., 33–34, 33f, 42–43 Picca, F., 39, 39f Picoli, G., 153, 154, 157, 158, 166, 178, 182 Piechotka, M., 78 Pimpinelli, A., 91, 92f Piotrowski, J., 93, 107–110 Pivavoroff, M., 232–233, 232f Plachy, J., 5 Plevachuk, Yu., 23–25 Poindron, E., 228, 229f, 230f Polichar, R., 217–218 Pollac, F.H., 121–122 Polyakov, A.N., 37, 66t Ponce, F.A., 107–110 Pong, Z., 107–110 Ponpon, J.P., 121, 126 Popov, D.A., 223–224 Popov, D.I., 80 Potter, M.D.G., 44–47 Powell, A.K., 160–161 Prakash, T., 4–5, 9, 12–13 Prasad, D., 4–5, 12–13 Prat, V., 218 Praus, P., 37 Prete, P., 111–112 Price, S., 36–37 Prokof’ev, S.V., 20, 40 Puhakka, K., 222f, 225, 226f, 227f Puzhevich, B.K., 121, 122 Pyyhtia¨, J., 220, 222f, 225, 226f, 227f

Q Qadri, S.B., 36–37, 99–100 Qian, Y., 122–123, 124 Qiao, Y., 107–110 Queheillalt, D.T., 54–55 Queisser, H.J., 155 R Rabenhorst, H., 187–188 Raccah, P.M., 233–234 Radhakrishnan, G., 100–102 Radhakrishnan, J.K., 61f, 62f Radisavljevic, K., 240t Raftery, N.A., 124 Raghothamachar, B., 37, 39, 80–81 Raiskin, E., 22–23, 37 Rajavel, D., 100–102 Rakhshani, A.E., 187, 195, 196–197 Ramachandran, N., 78 Ramakrishnan, S., 80 Raman, R., 61f, 62f Rambaud, P., 99–100 Rana, R.S., 168 Ranon, P.M., 151–152, 153, 160–161, 163 Rao, J., 9, 12 Rawling, J., 13 Redden, R., 12–13 Redden, R.B., 64 Redden, R.F., 32–33, 32f, 42–43, 79 Reddy, R., 9, 12 Reed, J., 217–218 Reed, J.D., 154 Reed, M.D., 37 Reedy, H.E., 242 Regal, R., 66t, 218 Regel, L.L., 35–36, 40, 80 Reineke-Koch, R., 192–193 Reislo¨hner, U., 195 Reisman, A., 28–29 Reno, J., 107–110 Repins, I., 113–114 Repka, E., 169 Repzka, E., 38–39, 153, 171, 172–173, 181–182 Rhiger, D., 222, 222f, 223f, 224f Rhiger, D.R., 36–37 Ricco, A.G., 125 Ricco, A.I., 124–125 Richter, H., 198 Rinas, U., 22–23, 36–37 Ringel, S.A., 197, 198

Author Index

Rioux, D., 201–202 Rioux, J., 25–26, 56 Rit, C., 66t Rivie`re, J.P., 107–110 Robbins, D.J., 96–97 Robinson, M., 44–47 Rodot, H., 36–37 Rodway, D.C., 91 Rogalski, A., 93 Rogers, K.D., 193–194 Rohatgi, A., 197, 198 Rohsenow, W.M., 49 Rolands-Jones, R.L., 107 Rolland, G., 99–100 Romanato, F., 107–110 Romani, S., 193–194 Romelue‘re, J.F., 97, 98f Romeo, A., 91, 124, 199, 201, 204, 205, 207–208 Romeo, N., 197, 201–202 Romero, M., 205 Rongalski, A., 107–110 Roosen, G., 76, 77–78, 80, 81, 153, 156, 157, 158, 166, 168, 181–182 Rose, D., 47, 121–122, 125, 127 Rosemeier, R.D., 215 Rosen, G.J., 23–25, 54 Rossmann, H., 107–110 Rostaing, J.P., 220, 220f, 221f, 226f Roszmann, A., 12, 13 Rothemund, W., 36–37 Roth, M., 36–37, 54, 57, 58f, 77 Rotter, S., 36–37, 54, 57, 58f Rouger, M., 228, 229f, 230f Roumie, M., 12–13 Roussillon, Y., 192, 199 Route, R.K., 36–37, 38–39, 53 Rouvie`re, J.L., 95–96, 95f Rowinska, L., 14 Rowlands, J.A., 222, 222f, 223f, 224f Royle, A., 43, 77, 96–97 Ruault, M.-O., 37, 56, 57f Rubcich, M., 189 Rubenstein, M., 43 Rudolph, P., 22–25, 25f, 26f, 36–37, 39, 42–43, 44, 49, 64–65, 66t, 79, 125 Rujirawat, S., 85–88, 100–102, 103, 103f, 104f Russell, G.J., 44, 107–110 Rustique, J., 220, 220f, 221f Ryckmans, Y., 49–50 Ryoichi, O., 64

271

S Sabhapathy, P., 43 Sadeghi, M., 189, 197 Saito, T., 56, 96–97 Saji, M., 107–110, 121–122 Salamanca-Riba, L., 100–102 Salansky, N., 90–91, 107–110 Salcudean, M., 43 Salk, M., 42–43, 79, 80 Saller, E.J., 100–102 Salomon, M.B., 214–215 Salviati, G., 107–110 Samaraserka, I.V., 50 Samimi, M., 77–78 Saminadayar, K., 200–201 Sanderson, N., 4–5 Sanghera, H.K., 44–47, 66t, 77 Sang, L., 122–123, 124 Sangwal, K., 125 Sangwall, K., 119, 133 Sanin, K.V., 23–25, 40 Saraie, J., 90–91, 107–110, 193–194, 195 Sarai, J., 99–100 Sassenberg, U., 121–122 Sato, G., 230, 231f, 232f, 233, 234f Saucedo, E., 39 Sauvage, F., 220, 226f Sauvageon, A., 228, 229f, 230f Sava, A.A., 121, 122, 124–125, 126, 131t, 139 Saxena, A.N., 122, 127 Schaake, F., 12 Schaake, H.F., 64–65 Scha¨ffler, R., 192–193 Scharf, J., 113–114 Schaub, B., 41, 43 Scheiber, C., 217–218 Schenk, M., 20 Schentke, I., 79 Schetzina, J.F., 37, 90–91, 96, 100–102, 107–110, 214–215 Schevzov, C., 50 Schiber, M., 122–123 Schieber, H., 37 Schieber, M., 15t, 56, 77 Schirato, R., 217–218 Schlesinger, T.E., 37, 56, 122–123, 233–234 Schmid, F., 36–37, 38–39 Schmid-Fetzer, R., 121–122 Schmitt, R., 36–37, 47, 66t Schneider, D., 29, 38–39, 40, 66t Schnepple, W.F., 77, 78 Schoenholz, R., 41

272

Author Index

Scho¨nholz, R., 77–78 Schulman, T., 220, 222f, 225, 226f, 227f Schwartz, H.-J., 36–37, 66t Schwartz, R.N., 173 Schwarz, R., 44, 66t, 79 Schweizer, M., 80 Schwenkenbecher, K., 42–43 Scott, S.E., 37 Sebastion, P.J., 203 Segall, B., 125 Segmu¨ller, A., 91 Seibert, G., 77 Sellin, P., 222, 222f, 223f, 224f Selvaraj, P., 93–94, 194–196, 197, 200 Semeniuk, P., 4–5 Senchenkov, A.S., 42–43, 79, 80, 81 Senin, R.A., 37, 66t Sen, S., 36–37, 56, 66t Senturina, N.N., 20, 41 Serreze, H.B., 214 Seth, G.L., 38–39, 61f, 62f, 64–65, 125, 127 Seto, S., 56, 57, 58f, 233–234 Sewell, P.B., 90–91 Shachna, A., 44, 66t Shah, K.S., 222, 222f, 223f, 224f, 226–227, 228f Shao, S.-Y., 36–37, 127 Shapkin, P.V., 30, 40 Sharma, R., 12 Sharma, S.R., 36–37, 66t Shaw, D., 20 Shaw, J.L., 189–190 Shaw, N., 43 Sha, Y.-G., 47, 78 Shchastlivyi, V., 7–8 Shcherbak, L., 23–25, 24f Shcherbak, L.P., 121, 123, 124–125, 126, 130–134, 131t Shchukarev, A.V., 122 Sheldon, P., 189, 196, 197, 199–200, 201, 202 Shelpakova, I., 9 Sher, A., 56 Shetty, R., 35–36, 54 Shiau, J.-J., 38–39 Shibata, N., 125 Shibutani, T., 189, 196, 197 Shi, D., 122–123 Shih, C.K., 56 Shinya, K., 64 Shiozawa, L.R., 63f Shirafuji, J., 41 Shiryaev, V., 13 Shi, W., 122–123, 124

Shmatov, N.I., 52 Shockley, W., 155 Shoisswohl, M., 171, 172–173 Shor, A., 217–218, 218f Shvydka, D., 192, 199 Shyy, W., 52 Sickinger, P., 81 Sides, P.J., 96–97 Sieber, B., 43 Siegrist, T., 91 Siepchen, B., 107 Siffert, P., 11f, 12–13, 23–25, 37, 40, 41, 50, 66t, 77–78, 122, 127, 131t, 136, 169, 175–176, 177, 214, 215, 217–218 Silk, J., 20, 23–25 Siminadayar, K., 107–110 Simons, M.Y., 107 Simpson, W.I., 96 Singh, A., 12–13 Singh, V.P., 202–203 Sites, J.R., 190–191, 192, 194–195 Sitter, H., 36–37, 93, 95–96 Sivananthan, S., 85–88, 87f, 90–91, 100–102, 103, 103f, 104f, 233–234 Sivapathasundaram, D., 197 Skinner, D., 189 Sklyarchuk, V., 23–25 Skronne, B.J., 100–102 Smith, A.W., 197, 198 Smith, D., 232–233, 232f Smith, D.J., 85–88, 87f, 100–102, 103f Smith, D.S., 85–88 Smith, F.T.J., 20 Smith, T.M., 78 Snyder, D.W., 96–97 Socha, A.J., 30, 36–37 Sochava, S.L., 178 Sochinskii, N.V., 77 Sokolov, V.K., 250–253, 254 Soldner, S.A., 37 Sone, S., 107–110 Song, J.H., 58, 59f Song, S., 12, 13 Song, W.-B., 36–37 Sorgenfrei, R., 112–113 Sou, I.K., 107–110 Sowinska, M., 215 Spartiotis, K., 220, 222f, 225, 226f, 227f Spa¨th, B., 200–201 Spicer, W.E., 56 Sporken, R., 90–91, 100–102, 103, 104f, 107–110, 111–112

Author Index

Sredin, V.G., 121 Srighton, H.S., 124–125 Srivastava, M., 36, 61f, 62f, 66t Stace, C., 160–161 Stadelmeier, H.H., 100–102 Stafford, A., 99 Stafsudd, O.M., 197, 240t Stahle, C.M., 125 Staudenmann, J.L., 107–110 Steier, W.H., 151–152, 153, 155–156, 158, 160–161, 163 Steigerwald, M.L., 96–97 Steininger, J., 20, 56, 57 Steldt, R., 168 Stepanov, I.S., 249 Stepanov, S.I., 153, 165–166, 178 Sternklar, S., 178 Stevenson, R., 206–207, 209 Steward, V., 43 Stoller, R.E., 56 Stollwerck, G., 190–191, 192 Stolyarova, S., 96–97 Strait, J., 154 Stratiychuk, I.B., 130–133, 131t, 136, 137–139 Straughan, B.W., 41, 43, 77 Strauss, A.J., 20, 56, 57, 90 Strehlow, W.H.J., 133 Stuck, R., 214 Stuk, R., 122, 127 Su, C.-H., 37, 39, 47, 77, 78, 124–125 Sudharsanan, R., 25–26, 234–235 Sudheer, Ch., 4–5, 9, 12–13 Sugiura, Y., 107–110 Suh, S.H., 58, 59f Sukkivan, A.S., 63f Sumah, P., 216 Summa, D., 197 Summers, J.G., 196–197 Sunderseshu, B.S., 38–39, 61f, 62f, 64–65 Suzuki, M., 230, 231f, 232f Sverdlin, I.A., 119–120 Svob, L., 97, 98f Swanson, B.W., 41 Swiatek, K., 47, 66t Sylla, L., 81 Szadkowski, A., 44–47, 215 Szapiro, S., 44 Szczerbakow, A., 44–47, 66t, 121–122, 125, 127 Szeles, C., 36–37, 66t Sze, S.M., 192 Szofran, F.R., 80 Szoke, J., 37, 39, 80

273

T Tabarrok, B., 42–43 Taguchi, T., 41 Tahashi, M., 111–112 Tai, H., 43 Tajima, H., 233, 234f Takahashi, H., 103, 230 Takahashi, T., 230, 231f, 232f, 233, 234f Takajanagi, S., 124–125 Takamura, H., 125 Takeuchi, S., 125, 127 Takeuchi, T., 124–125 Takigawa, H., 56 Takita, K., 113 Tamura, K., 233, 234f Tanaka, A., 54, 56, 57, 58f, 66t, 107–110, 125, 233–234 Tanaka, T., 111–112, 193–194, 195, 233, 234f Tan, G.L., 30 Taniguchi, Y., 37, 66t Tanner, B.K., 44–47 Tanoue, T., 224, 224f, 225f Tao, Y., 58, 59f, 60f Tapfer, L., 111–112 Tashiro, M., 230, 231f, 232f, 233, 234f Tatarenko, J., 85–88, 89f Tatarenko, S., 91, 92f, 107–110, 121–122 Tayebati, P., 166, 177 Taylor, R.E., 36–37, 66t Teeter, G., 200 Tegal, R., 217–218 Temple, D., 30 Tenne, R., 36–37, 54, 57, 58f Terada, Y., 233, 234f Teramoto, J., 124–125 Terent’ev, A.I., 223–224 Terheggen, M., 201–202, 207–208 Terry, J., 189–190 Tetali, B., 93–94, 194–196, 197, 200 Teubner, T., 107–110 Theret, G., 99–100 Thiry, P.A., 111–112 Thissen, A., 199–200 Thomas, G., 223–224 Thomas, J.E., 47, 66t Thomas, R.N., 41 Thomas, S., 218 Thompson, G., 202–203 Thompson, G.W., 203 Thompson, J.E., 23–25, 54 Tighe, S.J., 36–37, 66t Tinjod, F., 95–96, 95f

274

Author Index

Tiwari, A.N., 91, 100–102, 124, 199, 201, 204, 205, 207–208 Tkach, V.N., 131t To, B., 113–114 Tobin, S.P., 12, 64–65 Toguri, J., 13 Tokumori, K., 224, 224f, 225f Toman, J., 93 Tomashik, V.M., 121, 123, 131t, 134, 137–139 Tomashik, V.N., 120, 121, 122, 124–125, 126, 127–128, 129, 130–135, 131t, 136, 139 Tomashik, Z.F., 120, 121, 123, 124, 126, 127–128, 129, 130–135, 131t, 136, 137–139 Tomasov, A.A., 223–224, 250–252 Tomizono, T., 12 Tomizuka, A., 125 Tomson, A.S., 79 Toney, B., 37 Toney, J.E., 56 Toney, J.T., 233–234 Torikoshi, M., 225–226 Toth, A.L., 36–37 Tourrette, T., 228, 229f, 230f Tower, J.P., 12, 64–65 Toyofuku, F., 224, 224f, 225f Traczwski, R., 96 Tranchart, J.-C., 41, 56 Triboulet, R., 4, 14–15, 25–26, 27, 29, 31–32, 33–34, 33f, 34f, 35–37, 38–39, 40, 41–43, 44, 45f, 48, 48f, 51, 51f, 56, 57f, 58–61, 60f, 64–65, 66t, 96–97, 153, 169, 171, 172–173, 175–176, 177, 181–182, 234–235 Trivedi, S., 151–152, 153, 155–156, 158, 160–161, 163, 173, 215 Tromson-Carli, A., 4, 14–15, 33–34, 38–39, 40, 43, 48, 48f, 56, 60f, 66t Trumbly, T., 189 Tsabarim, M., 217–218 Tsen, S.C.Y., 85–88, 87f, 100–102, 103f Tsuboya, I., 36–37 Tsuji, A., 189, 196, 197 Tsunoo, T., 225–226 Tuck, B., 119–120 Tueller, J., 230, 231f, 232f Tukhkonen, L.M., 122 Tu¨mer, T.O., 222, 222f, 223f, 224f Turjanska, L., 23–25 Turkevych, I., 23–25 Turki, K., 166 Turner, A.K., 196–197

U Uchida, W., 106f, 107–110 Uda, H., 194, 195 Ufimtsev, V.B., 52 Ullal, H., 189 Uzura, S., 111–112 V Vaillant Roca, L., 201–202 Valley, G.C., 153, 156 Vandekerkof, J., 41 Van Den Berg, L., 78 van Gemmeren, T., 111–112 Vannuffel, C., 99–100 Van Schaftingen, J.J., 49–50 Van Scyoc, J.M., 56, 122–123 Vanyukov, A.V., 25–26 Varma, K., 9 Vasilev, A.A., 249 Vedel, I., 121, 122 Vedel, J., 95–96, 107–110, 197 Velumani, S., 203 Vengel, P.F., 121 Venzon, J.E., 217–218 Vere, A.W., 25–26, 43 Verger, L., 37, 66t, 220, 220f, 221f, 223–224, 226f Ve´rie, C., 27 Versluys, J., 198 Viallet, J.E., 166 Vieux, V., 153, 154, 157, 158, 178, 182 Vigdorovich, V., 14 Vilensky, A., 15t Villing, A., 168 Visoly-Fisher, I., 208 Viswanathan, V., 93–94, 194–196, 197, 200 Viturro, R.E., 189–190 Vodakov, Yu.A., 187 Volkmann, T., 95–96 Volz, M.P, 78, 80 Von Bardeleben, H.J., 171, 172–173 Von Kujawa, R., 9 von Roedern, B., 189, 207–208 Vozmilova, L.N., 119–120 Vujisic, L., 80 W Waag, A., 90–91, 107–110 Wadley, H.N.G., 23–25, 54–55

Author Index

Wagenaar, D.J., 219, 219f Wald, F., 41 Wald, F.V., 42–43 Waligura, A., 216 Walis, L., 14 Walker, R.C., 197 Wallace, J.P., 54 Wallis, D.J., 102 Walsh, D., 42–43 Walsh, K., 160–161 Wang, C.L., 80 Wang, J.F., 12, 13, 36, 103, 106f, 107–110 Wang, L., 122–123, 124 Wang, L.J., 78 Wang, W., 210 Wang, W.S., 95–96, 100–102, 107–110, 110f Wang, Y.Z., 42–43, 79, 85–88, 88f, 104, 112 Wanqi, J., 55 Ward, J.S., 205 Warekois, E.P., 124–125 Warren, S., 218 Warta, W., 189 Watanabe, K., 93 Watanabe, S., 230, 233, 234f Watson, A., 20, 23–25 Weibel, H., 100–102 Wei, H.Y., 100–102 Weinberg, F., 50 Weinstein, M., 107 Wei, S.H., 200 Weiss, S., 178 Wendt, R., 124, 199, 201, 207–208 Wermke, B., 36–37, 125 Werthen, J.G., 187 Wetzel, G., 78 Whelan, R.C., 20 White, A., 5–7, 9, 12–13 White, M.S., 125 Wiame, F., 102 Wiedemeier, H., 47, 78 Wienecke, M., 20, 124 Wijewarnasuriya, P., 104, 105f, 106f, 107–110 Wilamowski, Z., 47 Wilcox, W.R., 23–25, 35–36, 40, 49, 53, 54, 80 Wilde, L., 107–110 Williams, D.F., 90–91 Williams, D.J., 25–26, 43 Williams, G.M., 90 Williams, H., 124 Williams, L.M., 96–97 Williams, R.H., 121–122, 199

275

Wilshaw, P.R., 198 Witkowska, B., 215 Witkowska-Baranb, M., 215, 235–236, 235f, 236f Witt, A.F., 49 Witthuhn, W., 195 Wojnar, P., 215 Wolfenden, J.M., 217–218 Wolffer, N., 153, 160–161, 169, 175–176, 178, 179, 181–182 Wolf, G.A., 107 Wolf, M., 36–37, 53 Woodbury, H.H., 40, 96 Wood, C.E.C., 100–102 Woodcock, J., 196–197 Wood, D.A., 193–194 Woods, J., 44, 107–110 Wouters, P., 49–50 Wriedt, H., 15t Wright, G., 15t Wright, J.M., 100–102 Wrobel, J.M., 25–26 Wro´bel, J.M., 234–235 Wu, J., 36–37, 66t, 85–88, 88f, 104, 112 Wu, N.Q., 30 Wu, W.-H., 36–37 Wu, X., 189, 196, 197, 200, 202, 205 Wu, Y., 85–88, 88f, 104, 112 Wu, Y.S., 90–91, 107–110 X Xiaohua, L., 55 Xia, Y., 122–123, 124 Xie, Q., 107–110 Xin, Y., 102, 103, 103f, 104f Xu, J.Z., 107–110 Xu, X.N., 85–88, 88f, 112 Xu, Z.Y., 107–110 Y Yadava, R.S.D., 38–39, 64–65 Yaffe, M.J., 222, 222f, 223f, 224f Yakubtsov, O.A., 121, 122 Yamada, A., 200–201 Yamashita, T., 194, 195 Yang, B., 47 Yang, C.Y., 44–47 Yang, G., 214–215, 235–236, 235f, 236f Yang, J.R., 85–88, 88f, 112 Yang, N.J., 20, 23–25

276

Author Index

Yanka, R.W., 90–91, 107–110 Yano, M., 111–112 Yano, S., 224, 224f, 225f Yan, Y., 200, 202 Yao, H., 15t Yaohe, Z., 55 Yao, H.W., 122–123, 214–215 Yariv, A., 240, 240t Yaroslavtsev, A., 14 Yasuda, K., 103, 107–110 Yeckel, A., 37–38, 52, 55 Yellin, N., 44, 66t Yermolayeva, T.P., 126 Yew, Y.K., 198 Ye, X., 42–43 Yin, J., 107–110, 111–112 Yin, S., 222, 222f, 223f, 224f Yodogawa, Y., 90–91, 99–100, 107–110 Yokota, K., 38–39 Yoon, H., 56, 122–123 Yoon, J., 37 Yoshiie, T., 125 Yoshikawa, M., 56 Yoshikawa, T., 38–39 Yound, I.M., 90 Young, D.L., 205 Young, E.T., 217–218 Young, M.R., 200, 201 Young, T.L., 124 Yuan, S., 107–110 Yue, A.S., 44–47 Yu, F.-L., 36–37, 127 Yu.F. Schelkin, 52 Yu Ivanov, M., 25–26 Yu, J., 107–110 Yukio, A., 64 Yu, M.F., 37, 85–88, 88f, 104, 107–110, 112 Yu, M.-Y., 36–37 Yu Rud’, V., 40 Yu, T.-C., 23–25 Yu, V., 47 Yu.V. Klevkov, 20, 41 Yu.V. Rud’, 20, 23–25

Z Zaets, W., 111–112 Zahraman, K., 12–13 Zanatta, J.P., 99–100 Zanio, K., 4, 7, 11f, 12, 20, 43, 47, 66t Zanio, K.R., 100–102 Zanotti Fregonara, C., 107–110 Zanotti, L., 29, 37, 81 Zaouir, A., 11f, 12–13 Zappettini, A., 29, 37, 81 Zelenina, N.K., 223–224 Zerrai, A., 37, 175–176, 177 Zha, B.J., 127 Zha, M., 29, 37 Zhang, H., 80 Zhang, Q.Y., 85–88, 88f, 112 Zhang, Y., 47 Zhao, B-J., 36–37 Zhao, S.N., 44–47 Zharikov, E.V., 79 Zheng, J.G., 30 Zhong, J., 107–110 Zhou, J., 107–110, 111–112, 200 Zhu, S.-F., 36–37, 127 Zhu, X.-H., 36–37, 127 Ziari, M., 151–152, 153, 155–156, 158, 160–161, 163, 173 Zielinger, J.P., 156, 157, 158 Zimmermann, H., 36–37 Zitter, R.N., 124 Zogg, H., 90–91, 100–102, 124, 199, 201, 204, 205, 207–208 Zozime, A., 56 Zubia, D., 93–94, 94f Zuccalli, G., 29 Zumbiehl, A., 66t Zupp, R.R., 20, 63f Zuzga-Grasza, U., 44–47 Zvara, M., 36–37 Zweibel, K., 189 Zwerger, A., 112–113

SUBJECT INDEX

A Absorption coefficient, 158, 187, 242 Accelerated lifetime tests, 205, 208 ANTEC Solar GmbH, 207 Aperture area, 189 Astrophysics, 228–233 B Back contact, 188–193, 198–201, 205, 208 Back contact diode (BCD), 190–193 Back contact effects, 189–193 Back wall, 203 Balance of system (BOS) costs, 209 Band line-ups, 195 Barrier lowering, 193 Bifacial CdTe solar cell, 205 Bridgman, 22–23, 25, 36–40, 49–55, 61, 65, 77, 80–81, 153, 215, 235, 241 Br/methanol solution, 124, 133, 199 Buffer layers, 86, 99–102, 104, 108–109, 111, 113, 189, 199, 202, 204, 208 Buoyancy convective flow, 77 C Cadmium purification, 4–16 Capture cross sections, 175, 193 CdMnTe (CMT), 26, 111, 214–237 Cd2SnO4, 202 n-CdS/p-CdTe heterojunction, 187 CdTe, 4–5, 12–16, 19–67, 76–82, 85–114, 119–140, 148–184, 187–210, 214–237, 239–255 CdTe-metal, 199 CdTe:V DPCM (Double phase conjugated mirror), 178–184 photorefractive effect, 153–178 TWM (Two-wave mixing), 153–169 CdZnTe, 4–5, 14–16, 19–67, 76–82, 172, 175, 214–237 Cd0.9Zn0.1Te:In, 81

CdZnTe:V DPCM (Double phase conjugated mirror), photorefractive effect, 169 TWM (Two-wave mixing), Chemical bath deposition (CBD), 197 Chemical etching, 119–120, 122–127, 129–140 Chemical polishing, 120, 122, 124, 126, 131, 136, 139 Close-spaced sublimation, 89, 93–94, 107, 113, 194, 196–198, 203, 207 Collection and recycling programme, 207, 210 Compound purification, 5, 14–15 Crossover, 190, 192–193, 195, 199 Crystal structure, 125, 194 Cu, 6–9, 12, 14, 65, 114, 199–201, 203, 205 Cu2-xTe, 200 D Deep levels, 152, 154, 166–168, 171, 173, 175–178, 193, 198, 215, 244 Deep level transient spectroscopy (DLTS), 175–178, 198, 215 Degradation mechanisms, 187, 207–208 Deposition methods, 89–98 Detached, 80–81 Detectors, 41, 56, 64, 79, 87, 99, 112–114, 122–123, 150, 214–237, 248–255 Dewetting, 39, 67, 77, 80–82 Diffusion length (Ln), 193 Diffusion model, 94, 193 Distillation, 5, 9, 12–14 Diurnal variation, 209 E Electron affinity, 189, 201 Electron current, 190, 193 Electroplating, 196 Encapsulation, 37, 41, 208 Energy payback time (EPT), 209–210 Etchant composition, 120–131, 134, 136–140

277

278

Subject Index

Etching rate, 120–123, 125–126, 128–130, 132–137, 139–140 EURECA, 78, 80 F Fabrication of cells, 30, 189, 196–205, 239, 242 Fermi level pinning, 199 Fill factor (FF), 189, 191 First Solar, 206–207, 209–210 Flexible polymer substrates, 204 Flexible substrates, 202–204 FOTON–11, 78 FOTON-M, 81 Front wall, 203 G Gamma (g)-camera, 217–222 Generating cost, 209 Grain boundaries, 25, 48, 188, 199 Growth from the melt, 76–77, 79 Growth mechanism, 90, 97, 102 H Halogen evolving etchants, 127–129, 131 Heteroepitaxy, 99–113 Heterogeneous contact, 199 HgI2, 77–78 High-volume module production, 206 Hole current, 190–191, 193 Hole mobility, 193 (Hg,Cd)Te films, 78 Hydrogen peroxide (H2O2), 127–128, 133–135 I Ideality factor (A), 191 Image correlator, 254, 255 Impurity segregation, 32 Injecting contacts, 189 Interconnects, 148, 206, 208, 233 Inter-diffusion, 193–198, 201 Intermediate ternary compounds, 193 ITO, 187, 202, 204–205 L Large-scale power generation, 208–210 Lateral conductivity, 199 Lattice strain, 194 Layer adhesion, 202 Legislation restricting, 210

Lift-off processes, 204 Light soaking, 205 M Mammography, 222–223 Maximum efficiency, 182, 188 Maximum power point, 191 Medical imaging, 217–228, 236 Microgravity, 42, 48, 67, 76–82 Microstructure, 42, 108–109, 137–138, 197–198 MIR, 78, 80 Miscibility gap, 56, 194 Mixed crystals, 193–194 Modulator cavity dump, 245–247 CdTe, 240–245 CO2 laser, 240–242, 245–247 electro-optic, 239–247 mode lock, 245, 247 Q-switch, 245–247 Module efficiency, 189, 210 Multi-junction configurations, 205 N Nanodimensional formation, 137–139 Nanostructure, 96, 250–254 Net CO2 emission rate, 209 Nitric acid, 121, 124–126, 128 Non-ideal fill factor, 191 Nonproliferation, 215–217 N–P etch, 199 O Observed barrier heights, 189 Open circuit voltage (VOC), 189–192, 204, 207–208 Optical beam induced current (OBIC) studies, 198 Optical computer, 248–255 Optical information recording, 249 Optical transmission, 202 Oxygen contamination, 15–16 Oxygen solubility, 15 P Phase diagram, 20, 31–32, 62–63, 193 p-n homojunction, 187 POLIZON, 81 Polycrystalline nature, 188 Post-growth annealing in chlorine, 197–198

Subject Index

279

R

T

Radiation, 51, 114, 123, 194, 202–203, 214–237, 240, 248 Radiation damage, 202–203 Reach-through diode, 192 Reach-through microdiodes, 192 Recrystallisation, 197–198 Reversed biased diode, 190 Rinsing, 121–122, 124–125, 137, 204 Rollover, 190–192, 195, 199 Rotating magnetic field, 42, 79–80

Tellurium purification, 4–16 TeO2, 121–122, 124–127, 200 Te-rich surface, 91, 199–200 Thin films, 56, 85–114, 120, 204 Thin film solar cell, 67, 91, 107, 114, 187 THM. See Traveling heater method Toxicity of Cd, 210 Transparent conducting oxide (TCO), 187, 195, 201–202, 204, 208 Traveling heater method (THM), 15, 31, 34, 41–44, 60–61, 64, 67, 78, 82, 215 Tunnel, 189, 198, 200–201, 203 Two independent diodes model, 192 n-Type limb, 188

S Sb, 7, 65, 87, 112, 124, 130, 132, 201, 208 Sb2Te3, 201, 208 Schockley–Queisser model, 188 Schottky limit, 189 Segregation coefficient, 9–12, 14–15, 56–57, 233 Self compensation, 199 Semiconductor surface, 119, 129 Series resistance, 207 Sheet resistance, 202–205 Short circuit current density (JSC), 189, 191, 196, 202, 204–205 Short-term variation, 209 Shunt current paths, 192 Simulation studies, 209 Single crystal growth, 67 Solar farms, 209 Solution growth, 36, 41–43, 48, 76 Space applications, 202, 204, 230 Space charge region recombination, 192 Space exploration, 228, 236 Spacelab mission D1, 78 Specific power, 203–204 Spectral response, 195 Stable in the field, 207 State of the art, 97, 187, 189 Stressing protocols, 207 Structural defects, 79, 85, 106, 110, 124, 127, 194 Superstrate configuration, 187–188, 206 Surface recombination losses, 190 Synthesis, 27–34, 81, 184

U Under load, 207 V Vapour phase, 76–78 Variability of output power, 209 VOC deficit, 192 Voids, 188, 225 W Widow layer, 188 Window-type characteristic, 195 Work function, 189, 201, 203 Wurtzite, 194, 197 X X-ray computed tomography, 223–228 Z Zincblende, 194 ZnSe, 77–78, 108, 112, 243 Zn2SnO4, 202 ZnTe, 32, 34, 41, 47, 57, 61–64, 77, 100–104, 108–109, 111–113, 126, 136, 151–152, 200–201, 205 Zona 4 facility, 79 Zone refining, 5, 9, 12–14, 40–41, 61