Thin-Film Solar Cells [1 ed.] 9781617616549, 9781607418184

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Copyright © 2010. Nova Science Publishers, Incorporated. All rights reserved. Thin-Film Solar Cells, Nova Science Publishers, Incorporated, 2010. ProQuest Ebook Central,

Copyright © 2010. Nova Science Publishers, Incorporated. All rights reserved. Thin-Film Solar Cells, Nova Science Publishers, Incorporated, 2010. ProQuest Ebook Central,

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THIN-FILM SOLAR CELLS

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THIN-FILM SOLAR CELLS

Copyright © 2010. Nova Science Publishers, Incorporated. All rights reserved.

ABBAN SAHIN AND HAKIM KAYA EDITORS

Nova Science Publishers, Inc. New York

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Copyright © 2010 by Nova Science Publishers, Inc. All rights reserved. No part of this book may be reproduced, stored in a retrieval system or transmitted in any form or by any means: electronic, electrostatic, magnetic, tape, mechanical photocopying, recording or otherwise without the written permission of the Publisher. For permission to use material from this book please contact us: Telephone 631-231-7269; Fax 631-231-8175 Web Site: http://www.novapublishers.com

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NOTICE TO THE READER The Publisher has taken reasonable care in the preparation of this book, but makes no expressed or implied warranty of any kind and assumes no responsibility for any errors or omissions. No liability is assumed for incidental or consequential damages in connection with or arising out of information contained in this book. The Publisher shall not be liable for any special, consequential, or exemplary damages resulting, in whole or in part, from the readers‟ use of, or reliance upon, this material. Independent verification should be sought for any data, advice or recommendations contained in this book. In addition, no responsibility is assumed by the publisher for any injury and/or damage to persons or property arising from any methods, products, instructions, ideas or otherwise contained in this publication. This publication is designed to provide accurate and authoritative information with regard to the subject matter covered herein. It is sold with the clear understanding that the Publisher is not engaged in rendering legal or any other professional services. If legal or any other expert assistance is required, the services of a competent person should be sought. FROM A DECLARATION OF PARTICIPANTS JOINTLY ADOPTED BY A COMMITTEE OF THE AMERICAN BAR ASSOCIATION AND A COMMITTEE OF PUBLISHERS. LIBRARY OF CONGRESS CATALOGING-IN-PUBLICATION DATA Thin-film solar cells / editors, Abban Sahin and Hakim Kaya. p. cm. ISBN 978-1-61761-654-9 (Ebook) 1. Thin films. 2. Solar cells. 3. Crystals--Thermal properties. I. Sahin, Abban. II. Kaya, Hakim. TA418.9.T45T44 2009 621.31'244--dc22 2009038658

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CONTENTS

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Preface

ix

Chapter 1

Bridgman- Grown CuInSe2 C. H. Champness

Chapter 2

Review on Sulfur Compounds Deposited by Spray Pyrolysis Technique as Buffer and Absorber Thin Films for Solar Cells Karem Boubaker and Mosbah Amlouk

Chapter 3

Boubaker Polynomials Expansion Scheme BPESRelated Analysis of Low-Cost ZnS1-XSex Buffer Layers Thin Films Properties Karem Boubaker and Mosbah Amlouk

Chapter 4

Progress in Crystalline Silicon Thin-Film Solar Cells Armin G. Aberle

Chapter 5

Photothermal Analyses as an Up-To-Dated Guide of Zn-Doped Solar Cells Thin Films Evaluation Karem Boubaker and Mosbah Amlouk

Chapter 6

Formation and Characterization of C60- and Perylene-Based Bulk Heterojunction Solar Cells Takeo Oku, Atsushi Suzuki, Kenji Kikuchi, Nariaki Kakuta, Ryosuke Motoyoshi, Katsunori Nomura, Atsushi Kawashima, Yasuhiko Hayashi and Tetsuo Soga

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1

51

95 129

149

179

viii Chapter 7

Contents Prospective of Hydrogenated Amorphous Silicon (a-Si:H) Thin Films in Photovoltaics Sukti Chatterjee

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Index

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193 205

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PREFACE A thin-film solar cell (TFSC), also called a thin-film photovoltaic cell (TFPV), is a solar cell that is made by depositing one or more thin layers (thin film) of photovoltaic material on a substrate. This book deals with some physical properties of Sulfur binary and ternary thin films used as buffer and absorber layers in solar cells and prepared using economic spray pyrolysis technique. This book also investigates some thermal properties of Zn-doped binary thin films used as solar cells buffer layers and prepared using economic techniques. Other chapters in this book describe the development of diverging band gap amorphous silicon materials and their optoelectronic properties, the unique one ampoule Bridgman method, as well as the cleavage and twinning characteristics of the single crystals and how they are influenced by annealing, etching, deviation from stoichiometric starting proportions and by the addition of sodium. This book also investigates emerging trends that might lead to additional commercial c-Si thinfilm solar cells after 2010. Chapter 1 - The chalcopyrite semiconductor CuInSe2 is the basic compound underlying the quaternary alloy CuIn1-xGaxSe2, investigated for the p-type absorber layer in thin film photovoltaic solar cells. To examine the fundamental and material properties of the ternary compound, it is helpful to have it in bulk monocrystalline form. Single crystals of CuInSe2 can be grown from the vapour, from a liquid solution and from the melt of the synthesized compound. The present chapter describes the last of these processes, using a vertical Bridgman method. In this, the compound is first synthesized from the original elements copper, indium and selenium by heating them together above the compound melting point in a completely sealed quartz ampoule, to prevent loss of the volatile selenium above the melting point of CuInSe2 (986 oC). Then in a preferred process, the same sealed ampoule, containing the liquid compound, is slowly

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Abban Sahin and Hakim Kaya

lowered through a temperature gradient, resulting in a solidified ingot containing single crystals. The chapter describes the unique one ampoule Bridgman method used in this laboratory to grow ingots containing centimeter-sized CuInSe2 monocrystals with a minimum of the problems of ingot-ampoule adhesion, voids, cracks and other imperfections. It further describes the cleavage and twinning characteristics of the single crystals and how they are influenced by annealing, etching, deviation from stoichiometric starting proportions and by the addition of sodium. Chapter 2 - This chapter deals with some physical properties of Sulfur binary and ternary thin films used as buffer and absorber layers in solar cells and prepared using economic spray pyrolysis technique. First, a brief historic of the binary and ternary compounds is presented, then, the spray technique fundaments and experimental features are detailed. Finally, the physical characterization (XRD, SEM, AFM, optical, electrical as well as thermal measurements) results are presented showing the possibility of use of such films in the solar conversion devices. Chapter 3 - This chapter deals with sprayed-annealed ZnS1-xSex thin films used as buffer in solar cells and similar devices. First, a brief historic of the binary and ternary compounds is presented, then, the spray technique fundaments and experimental features are detailed. Then the physical characterization (XRD, SEM, AFM, optical, electrical as well as thermal measurements….) along with Boubaker Polynomials Expansion Scheme BPES investigations results for asgrown ZnS1-xSex thin films are presented showing the possibility of use of such films as buffer layers in the solar conversion devices. Finally, a new 3D AmloukBoubaker expansivity-Band gap Energy-Vickers microhardness  AB /Eg/Hv abacus is presented as a guide to solar-thermal-hybrid cells applications Chapter 4 - Photovoltaic (PV) modules based on crystalline Si (c-Si) wafer solar cells are a robust, proven and long-term stable technology. However, due to its material intensiveness, there are severe doubts as to whether a wafer-based c-Si technology will ever reach the low cost levels ($/kWh) required for a widespread application of PV. Given these cost constraints, the need for a less material intensive c-Si technology – a thin-film technology – exists. In this article, the most promising thin-film c-Si PV technologies that have emerged during the last 10 years are reviewed. Progress has been excellent and it now seems that half a century after the invention of the Si wafer solar cell the time has come where thinfilm c-Si solar cells can be transferred to industrial production. Furthermore, this article investigates emerging trends that might lead to additional commercial c-Si thin-film solar cells after 2010. It is expected that crystalline silicon thin-film

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Preface

xi

solar cells will lead to a massive reduction of about 35 % in the cost of gridconnected PV electricity by 2015. Such a cost reduction within the next 10 years seems very difficult to achieve with Si wafer based PV technologies. Hence, thinfilm photovoltaics finally seems ready to conquer the terrestrial PV market. Chapter 6 - Polymer bulk heterojunction solar cells were fabricated and the electronic and optical properties were investigated. C60 and perylene were used as n-type semiconductors, and copper phthalocyanine, zinc phthalocyanine and pentacene were used as p-type semiconductors. Energy levels of the molecules were calculated, and HOMO levels were localized around the main chains of the zinc phthalocyanine and pentacene. Nanostructures of the solar cells were confirmed as mixed nanocrystals by transmission electron microscopy and electron diffraction. Polymer/fullerene bulkheterojunction solar cells with poly [3hexylthiophene] (P3HT), poly [2-methoxy-5-(20-ethylhexoxy)-1,4phenylenevinylene] (MEH-PPV), and 6,6-phenyl C61-butyric acid methyl ester (PCBM) were produced and characterized. A device based on P3HT and PCBM provided better efficiency, fill factor, and short-circuit current compared to those of a device based on MEH-PPV and PCBM. The solar cell with P3HT and PCBM structure showed a higher photoresponse in the range of 400-650 nm. Energy levels of the molecules were calculated and discussed. Chapter 7 - For the consumer applications, amorphous silicon (a-Si:H) multijunction solar cell technology is the most popular photovoltaic technology for its low cost. Different band gap intrinsic layers (or carrier generation layers) are normally used in an a-Si:H multi-junction solar cell. In this article, the development of diverging band gap a-Si:H materials has been described. The optoelectronic properties of such materials were studied. Photovoltaic qualities of them are really promising.

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Chapter 1

BRIDGMAN- GROWN CuInSe2 C. H. Champness Electrical and Computer Engineering Department, McGill University, 3480 University Street, Montreal, Quebec, Canada, H3A 2A7

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ABSTRACT The chalcopyrite semiconductor CuInSe2 is the basic compound underlying the quaternary alloy CuIn1-xGaxSe2, investigated for the p-type absorber layer in thin film photovoltaic solar cells. To examine the fundamental and material properties of the ternary compound, it is helpful to have it in bulk monocrystalline form. Single crystals of CuInSe2 can be grown from the vapour, from a liquid solution and from the melt of the synthesized compound. The present chapter describes the last of these processes, using a vertical Bridgman method. In this, the compound is first synthesized from the original elements copper, indium and selenium by heating them together above the compound melting point in a completely sealed quartz ampoule, to prevent loss of the volatile selenium above the melting point of CuInSe2 (986 oC). Then in a preferred process, the same sealed ampoule, containing the liquid compound, is slowly lowered through a temperature gradient, resulting in a solidified ingot containing single crystals. The chapter describes the unique one ampoule Bridgman method used in this laboratory to grow ingots containing centimeter-sized CuInSe2 monocrystals with a minimum of the problems of ingot-ampoule adhesion, voids, cracks and other imperfections. It further describes the cleavage and twinning characteristics of the single

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C. H. Champness crystals and how they are influenced by annealing, etching, deviation from stoichiometric starting proportions and by the addition of sodium.

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1. INTRODUCTION There has been a considerable increase in interest in thin film photovoltaic solar cells, due to their minimal use of semiconducting device material compared to that used in silicon cells. Of special interest here are cells based on the chalcopyrite compound CuInSe2, and more particularly, the quaternary alloy CuIn1-xGaxSe2, arising from its high optical absorption coefficient in the main range of solar wavelengths. In a cell with a layer structure of the form Mo-CuIn1-xGaxSe2-CdS-ZnO(i)-ZnO(c)-Al, a solar cell conversion efficiency approaching 20 % has been reached [1] at the time of writing. Here ZnO(c) is conducting zinc oxide and ZnO(i) is its high resistance form. In this type of device, the quaternary consists of a thin layer of about 2 µm thick of polycrystalline material, with crystal grains typically of micrometer dimensions. While some characteristics and properties of this material can be studied in thin polycrystalline form, fundamental properties are more conveniently determined in bulk single crystals. For this reason, efforts to grow monocrystals of the basic ternary compound, large enough for study, have been made in many laboratories. Single crystals of CuInSe2 have been grown from the vapour with iodine as a transport agent [2, 3], from solution using indium as the solvent [4, 5], and from the melt. The present chapter, however, describes work only on melt growth by the Bridgman method. This is because of extensive experience in this laboratory with this method and also because the Bridgman method has yielded single crystals large enough for transport measurements [6] and photovoltaic cells with a solar conversion efficiency of more than 10 % [7]. The compound CuInSe2 has a direct energy gap of 1.04 eV, and a chalcopyrite crystal structure with cell parameters of a = 5.789 Å and c = 11.612 Å for a c/a ratio of 2.006 [8]. The Cu2Se-In2Se3 pseudobinary phase diagram, according to Fearheiley [9], is shown in Figure 1. Above 810 oC, it has a cubic zincblende structure. In its growth from the melt, it is important to realize that the vapour pressure of selenium is very much higher than that of the constituent metals copper and indium. Thus, at 1100 oC, which is above the melting point of the compound (986 oC), the vapour pressure of Se is about 40 atmospheres. Because of this, crystals have been grown either under high argon pressure

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with B2O3 as a liquid encapsulant [10] or within a thick-walled sealed ampoule [11], able to withstand the Se vapour pressure. In the latter case, the procedure has often been [11,12,13] to synthesize the compound first in a sealed ampoule, followed by its removal and re-introduction in powdered form into a second sealed ampoule for crystal growth. If stoichiometric proportions of the starting elements are used in this procedure, the resulting ingots are either ntype or partly n-type and partly p-type. This inhomogeneity of conductivity type is believed to be due to the partial loss of selenium, resulting from the deposition of this element on the inner wall of the first container. However, if a single ampoule is used for both synthesis and crystal growth, with stoichiometric starting proportions of the elements, the resulting ingot is uniformly p-type, making this a preferred procedure. In earlier work with the one-ampoule method, problems were encountered in this laboratory with adhesion of the ingot to the inner wall of the quartz ampoule and with the existence of small spherical voids or cavities within the ingot. Nevertheless, both of these problems were completely solved with the introduction of a flame-hardened film of boron nitride [14] deposited within the ampoule at the start of the process.

Figure 1. Phase diagram of the Cu2Se-In2Se3 pseudobinary system according to Fearheiley [9].

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C. H. Champness

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The present chapter describes in detail the growth of ingots of CuInSe2, as developed in the present laboratory, using the one-ampoule method with a boron nitride inner ampoule coating. The resulting ingots were free of microcracks and voids and were uniform in composition and conductivity type. They contained monocrystals large enough for preliminary photovoltaic cells [7]. In this investigation, the effect on the crystals of varying the growth conditions was specifically studied. For the most part, stoichiometric starting proportions of the elements were used but ingots with nonstoichiometric proportions were also prepared, where the results have been already reported [15]. The predominant cleavage planes were found [16] to be {112} and {101} and twinning was observed to take place on {112} planes. Effects of annealing, etching and of the addition of sodium to the melt are also described.

Figure 2. Schematic diagram of the vertical Bridgman growth setup, showing also the temperature profile of the furnace.

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Figure 3. Photographs of the quartz ampoule used for the crystal growth, showing (a) the powdered boron nitride coating as deposited on its wall and (b) the ampoule transparent after the intense flaming treatment.

2. CRYSTAL GROWTH PROCEDURE Details of the crystal growth procedure, including the preparation of the ampoule using the boron nitride coating technique, are first described, starting with the growth apparatus.

2. (a) Crystal Growth Apparatus The crystal growth of CuInSe2, in this work, was carried out by a vertical Bridgman method, using a resistively heated, 4600 watt, two-zone furnace (Thermcraft Inc. model 23/238-24-2ZV -ST). The temperature of the two

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zones of the furnace was controlled with a Honeywell programmable digital controller (model UDC 5000) with a reported accuracy of about ±0.1 degree. A schematic diagram of the setup is shown in Figure 2, along with a measured temperature profile of the two zones. In order to obtain a high temperature gradient for the crystal growth, a ceramic baffle was placed between the top and bottom zones, which yielded a gradient of about 70 °C/cm. From the temperature profile, the solidification position of the compound was estimated to be at a depth of some 43 to 45 cm below the top surface of the furnace.

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2. (b) Ampoule Cleaning In this work, the CuInSe2 charge synthesis and crystal growth were carried out in the same ampoule. The ampoule was made from a quartz tube, having an inner diameter of 12 mm and an outer diameter of 16 mm. It was shaped using an oxygen-hydrogen flame, from one end to obtain a pointed tip resembling the letter V, in order to facilitate nucleation during early stages of the growth. The quartz tube, with only one end open, was cleaned with a mixture of HNO3 and HCl (1:1 by volume) for a period of about 24 hours and then rinsed thoroughly in de-ionized (DI) water. The ampoule was then soaked in acetone for a similar period of time to remove any organic contaminants. Finally, the ampoule was rinsed again in DI water, dried gently with the flame and was now ready for the BN coating.

2. (c) Boron Nitride Coating The boron nitride coating process of the ampoule was found to be an important step in the fabrication of adhesion- and void-free ingots of CuInSe2. It involves employment of the same ampoule for both the synthesis and growth of the compound, resulting in ingots which are uniform in conductivity type. Furthermore, it does not require a supporting crucible within the ampoule [17]. The boron nitride (BN) coating employed in this work was deposited as follows [14]. Boron nitride was obtained either in the form of a powder, of nominal purity 99.5%, from Johnson Matthey Ltd., or in the form of BN slurry, from the Carborundum Company. The BN slurry was obtained in two types: V-BN, consisting of BN mixed with silica, and S-BN, where the boron nitride was mixed with an alumina phase. A small quantity (less than 5 gram) of the BN source was thoroughly mixed with about 10 to 20 ml of acetone to

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create a suspension with a thick consistency. The mixture was then transferred to the ampoule and given a manual agitation for some 2 to 3 minutes. This action allowed the ampoule wall to be coated with BN particles. The excess quantity of the suspension was poured out of the ampoule. After introducing the BN, the coating was made more uniform by spreading it out with a special brush until the desired thickness was obtained on the wall. The coated area was localized at the bottom end of the ampoule by rinsing away the excess coating elsewhere using DI water. After this, the ampoule, coated with the appropriate amount of BN, was gently heated between 80 and 100 C in a small oven to drive out any residual acetone remaining behind. Figure 3 (a) shows a photograph of an ampoule indicating the appearance of the BN. At this stage, the BN particles were loosely bound to the quartz and were easily removed with rinsing. In order to harden the coating, it was necessary to heat intensely, using a flame, both the quartz and the coating until the quartz was almost white hot. In the case of the 99.5% purity BN, the coating became fully transparent inside the ingot, as shown in Figure 3 (b). In using the BN slurry as the source of BN, the coating was not completely transparent after the flame treatment; however, it was of superior hardness compared with powdered BN of equal thickness. In order to remove any residual BN particles, the ampoule was given a final rinse with DI water and dried gently with the flame. The intense flaming of the BN is an important step in the deposition of the coating whereby, if it is not done, the BN particles would be washed away by the molten charge and may cause contamination of the ingot.

2. (d) Preparation of Charge In most of the growth runs carried out in this study, a stoichiometric charge of the elements was used for the synthesis of the compound CuInSe2. The starting materials of copper, indium and selenium were acquired in pellet form, each pellet with a diameter of about 3 to 5 mm, having a nominal purity of 99.999%. The suppliers for these materials were CERAC Coating Materials Inc. for the copper, Metal Specialties Inc. for the indium and Noranda Technology Center (Pointe Claire) for the selenium. Stoichiometric proportions of the Cu, In and Se were calculated from their respective atomic weights, and the elements were weighed separately from a pre-etched supply of material. The pellets were etched prior to the weighing of the charge to avoid unwanted changes in the starting composition due to the etching action. The chemical etching of the Cu was carried out in dilute HNO3 (1:10 in DI

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water by volume), while the etching of the In and Se was done in dilute HCl (1:10 in DI water by volume). The pellets were rinsed thoroughly in DI water and then dried in a flow of nitrogen. Figure 4 (a) shows a photograph of the etched and cleaned elemental pellets, prior to introduction into the ampoule. After the etching, the charge was transferred to the BN-coated quartz ampoule which was attached, in a vertical position, to a vacuum pumping system (by Edwards Ltd.) equipped with a stainless steel manifold, having a 16 mm diameter ampoule inlet port. The system, with the ampoule in place, was then evacuated down to a pressure of about 4 x 10-7 Torr in some 2 to 3 hours. Using an oxygen-hydrogen flame, the ampoule was then sealed under vacuum from a position some 10 cm above the elemental charge pellets to avoid unwanted heating and possibly loss of selenium at this stage. As a precaution, the tip of the ampoule was kept cool in a water-soaked cloth. Prior to loading the ampoule for pumping, the position of sealing was marked and the quartz was thinned down to facilitate its melting. Using the flame, the thinned wall of the ampoule was gently melted, while the ampoule was rotated until the quartz wall had collapsed under the influence of vacuum, forming a sealed neck. Extra care is needed in the sealing-off process, especially in separating the ampoule from the remainder of the quartz tube, where the quartz may crack if too much force is used. After the sealing was complete, the ready ampoule, some 15 to 20 cm in length was allowed to cool down prior to the beginning of the next step.

2. (e) Reaction of Charge The ampoule, with the charge inside, was placed in a horizontal position in a brick furnace, where the initial reaction between the elements was carried out. Here, the temperature was first raised slowly from room temperature to about 300 C, at about 2 C/min to allow the exothermic reaction between the indium and the selenium, near the temperature of 220 °C, to take place. If the rise in temperature is too fast, an explosion would occur. To ensure a complete reaction, the temperature was maintained at 300 C for some 24 hours. At this stage, the only reactants were the In and the Se, while the Cu remained in pellet form. This can be seen in Figure 4 (b), where the contents of a run, interrupted at a temperature of about 300 C, are shown, clearly indicating the presence of unreacted elemental Cu. Next, the sealed ampoule was transferred to the top zone of the vertical crystal growth furnace, where the remainder of the reaction was carried out. To do this, the ampoule was attached, from its top end, to a quartz

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rod some 60 cm in length. The rod was connected to the mechanical lowering mechanism of the assembly and, with the furnace turned off, the ampoule was positioned in the center of the top zone of the furnace, at a depth of about 35 cm. Next, the temperature was set to 300 C and the furnace was allowed to stabilize for about 1 hour. The temperature was then raised at a rate of about 5 C/min up to 1100 C, that is some 100 C above the melting point of CuInSe2. The reaction with the copper was complete however, before the melting point of this element (1083 C) was reached. This is demonstrated in Figure 4 (c), where the contents of a run, interrupted at around 700 C, were powdered and are shown to reveal the absence of any Cu pellets. In order to ensure homogenization of the melt, the ampoule was maintained at 1100 C typically for 24 to 48 hours and given a thorough manual agitation several times during this soaking period.

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2. (f) Crystal Growth Procedure After the reaction and soaking of the charge was complete in the top furnace zone, the temperature of the lower zone was set to about 700 C and the bottom furnace was allowed to stabilize for about 1 to 2 hours. This resulted in a temperature gradient of about 70 C/cm between the top and bottom zones. The directional freezing of the melt was then begun. The ampoule, in the top zone, was slowly manually displaced downwards to a depth of about 40 cm, nearer to the solidification position, where the temperature was still well above the melting point of the compound. The motorized lowering mechanism was then activated and the control was set to a lowering rate of 2 to 10 mm/hr. The growth was continued for 24 to 48 hours until the ingot was completely solidified, well into the cooler zone. The temperature of the top zone was then reduced to that of the bottom zone and the entire furnace cooled at a rate of about 20 to 30 C/hr. The resulting ingot is shown in Figure 5 (a), where it is clearly seen to be loose inside the ampoule, due to the effect of the BN coating and was thus easily extracted from the quartz ampoule. In contrast, an ingot exhibiting severe adhesion and residue is shown in Figure 5 (b) of this figure, for comparison, where no BN was used. Here, the ingot is stuck within the ampoule and deposits are evident on the ampoule wall. Table 1 summarizes many of the crystal growth runs reported in this chapter with a few additional runs of special interest. Table 1. Summary of Bridgman Growth Conditions for CuInSe2 Ingots Thin-Film Solar Cells, Nova Science Publishers, Incorporated, 2010. ProQuest Ebook Central,

10

TC4 (°C )

Lowering rate (mm/hr)

Coolin g rate (°C/hr)

None

Soa k tim e3 (hr) 72

650

2

27

Stoich Stoich Stoich Stoich Stoich Stoich Stoich Stoich Stoich Stoich Stoich Stoich

Pwdr Pwdr Pwdr Pwdr Pwdr S-BN Pwdr Pwdr Pwdr V-BN V-BN V-BN

48 48 24 48 24 24 48 96 20 10 22 48

600 600 650 700 650 560 550 520 520 850 850 RT

2 1 18 3 10 3 2 7 2 4 4 700

105 107

Stoich Stoich

V-BN Pwdr

68 2

850 600

6 -

25 25 25 19 27 23 22 22 22 18 11 Quenc h 20 -

110 111

Stoich Stoich

S-BN Pwdr

24 48

620 660

6 5

26 25

112

Stoich

S-BN

24

650

7

25

113

Stoich

Pwdr

26

660

6

31

115 116 119 125 137 138

Stoich Stoich Stoich Stoich Stoich Cu excess Se deficient In excess Se excess Cu excess Stoich In deficient In excess Stoich

V-BN V-BN V-BN S-BN Block V-BN

12 12 48 68 48 47

800 900 680 670 600 670

7 8 2 2 2 3

27 30 28 27 25 28

Pwdr

24

670

3

28

Pwdr Pwdr

20 24

670 640

3 3

28 53

Pwdr

21

670

3

50

Pwdr Pwdr

28 48

900 900

3 6

37 30

Block Pwdr

43 24

800 700

3 6

33 30

Run No.

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Starting charge1

BN coatin g2

21

Stoich

36 41 43 55 63 64 76 83 86 91 95 101

140 141 142 143 144 145 146 149

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Comments

Sticking and voids present Intermediate growth rate Large {101} cleavage Large {112} cleavage Large {110} cleavage Fast growth rate Si seed dissolved during soaking Al2O3 single crystal seed used SiO2 single crystal seed used Quartz ampoule diameter was 2.5 cm Inverted growth5 Inverted growth5 Large monocrystal Slow growth rate

EPMA growth runs

11

Bridgman-Grown CuInSe2

Run No.

BN coatin g2

Table 1. (Continued)

Soa k tim e3 (hr) 24 72

TC4 (°C )

Lowering rate (mm/hr)

Coolin g rate (°C/hr)

Comments

Stoich S-BN 700 277 50 Cu S-BN 700 6 Quenc excess h 156 Cu S-BN 48 700 7 Quenc Interrupted growth excess h runs 157 Cu S-BN 48 700 7 Quenc excess h 1 Elemental charge used at the start of the growth (Stoich = stoichiometric). 2 Type of boron nitride used for the ampoule coating, Pwdr = powdered boron nitride of 99.5% purity obtained from Johnson Matthey Ltd., V-BN = slurry of boron nitride and silica obtained from Carborundum, S-BN = slurry of boron nitride and alumina obtained from the Carborundum Company and Block = Block of pyrolytic boron nitride. 3 Time allowed for homogenization of the molten charge at a soaking temperature of 1100 °C. 4 Temperature of the second cooler zone of the furnace. 5 Soaking temperature was 1000°C. Copyright © 2010. Nova Science Publishers, Incorporated. All rights reserved.

151 155

Starting charge1

Figure 4. Photographs showing (a) the elemental charge before loading into the ampoule, (b) the charge after treatment at about 300 °C, showing residual copper pellets and (c) the charge reacted at a temperature near 700 °C, after powdering.

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3. INGOT GROWTH RESULTS

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The results of applying the procedure to grow ingots of CuInSe2, described in the previous section, are now presented. Included here are the effects of the boron nitride coating, growth rate variation and compositional uniformity in the ingot crystallites.

Figure 5. Photographs showing (a) ampoule pre-treated with BN, yielding an ingot completely loose inside and (b) ampoule without BN, where the ingot exhibited adhesion, cracks and residue on the quartz wall.

3. (a) Effect of Boron Nitride The boron nitride coating was clearly found to eliminate the adhesion of the CuInSe2 ingot to the quartz ampoule. This was a major breakthrough in this work, where for the first time in this laboratory, ampoules were routinely grown without any difficulties with cracking or adhesion. It was also found to reduce significantly the amount of voids in the interior of the ingot. Because of this, the growth of larger monocrystals was made possible and crystals with significant cleavage areas were obtained. This is demonstrated in Figure 6 (a), where two halves of a cleaved piece from run 64, Table 1, using BN, is shown, yielding complete cleaved surfaces of greater than 40 mm2 in area. This is compared with run 21 (Table 1), where no BN coating was used, and where the resulting ingot exhibited a large density of voids (Figure 6 (b)). Another beneficial aspect of the BN coating was that it allowed the ingot to develop facets on the exterior wall, particularly when thicker BN coatings were used.

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In some cases, this yielded information about the orientation of monocrystalline regions inside the ingot. While this was useful from the point of view of crystal growth, it was found that the BN did not have to be in physical contact with the melt to work. An experiment was carried out whereby the BN was in the form of a block located inside the ampoule but well above the melt. The result was again an adhesion-free ingot. Thus, this showed clearly that the BN did not act just as a physical barrier.

Figure 6. Photographs showing the interior of ingots (a) a run where boron nitride was used in the growth revealing the two halves of a cleaved single crystal and (b) a run where BN was not used, showing high void density.

3. (b) Lowering Rate/Crystallite Size

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As indicated in Table 1, most of the high quality CuInSe2 ingots were grown with vertical ampoule lowering rates between 0.2 and 1 cm/hr. This rate was determined from a special set of runs where the lowering rate was varied over two magnitudes from 70 to 0.2 cm/hr. For this work, pieces, obtained from the ingots, were first abrasively lapped and then polished with 0.3 micron alumina powder, followed by etching in a 5 % brome-methanol solution to reveal the crystallites. From these, the grain size could be determined. Figure 7 shows samples grown at 70, 1.8 and 0.2 cm/hr and Figure 8 shows a plot of average grain size against ampoule lowering rate. Here it is evident that to obtain centimeter-size monocrystals, the growth rate should be at or below 1cm/hr.

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3. (c) Compositional Uniformity Localized composition in the ingots was determined by electron probe micro analysis (EPMA). This was done by first slicing each ingot along its length into two halves. One of these was polished with first 1 micron and then 0.05 micron alumina powder to obtain a flat mirror-like finish. Following this, the surface was thoroughly cleaned with D.I. water, dried in nitrogen and coated with a thin layer of carbon to prevent charge-up during the EPMA examination. Figure 9 (a) shows the variation of Cu, In and Se along the length of stoichiometrically-prepared ingot number Z144 and Figure 9 (b) shows the variation across the width of the same ingot. Good uniformity of the elements is apparent. The average atomic proportions for this ingot were determined to be 24.6: 25.4: 49.1 at % for Cu: In: Se respectively, prepared with the starting stoichiometric proportions of 25: 25: 50 at %.

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Figure 7. Photographs of samples with variation of growth rate at (a) 700 mm/hr, (b) 18 mm/hr, (c) 2 mm/hr, where the samples were polished and etched to reveal grain boundaries.

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Figure 8. Double logarithmic plot of average grain size against ampoule lowering rate. The error bars represent the minimum and maximum grain size estimates obtained for each sample.

4. MONOCRYSTALLINITY 4. (a) Laue Diffraction

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In order to verify that the mirror-like cleavage facets did in fact correspond in each case to one crystal, Laue back-reflection pictures were taken sequentially at three displaced locations on a cleaved plane of a stoichiometrically grown ingot number119. This was done by moving the sample linearly, with no rotation, perpendicular to the X-ray beam to the spots indicated in Figure 10 (a). The clear single circular superimposed spots of the resultant 3-exposure Laue photograph in Figure 10 (b) confirms monocrystallinity over the cleaved area, with no evident streaking or splitting that would otherwise indicate a misorientation between adjacent grains. Further, the three-fold symmetry of the Laue spots confirms that the cleaved surface was a {112} plane.

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Figure 9. Plot of elemental atomic concentration (EPMA) in the stoichiometric ingot 144 (a) against distance along ingot and (b) against distance across ingot.

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4. (b) Rocking Curve Diffraction Using a 5 x 5 mm2 {112} cleaved surface from ingot number 125, a diffraction rocking curve was obtained at the Advanced Technology Laboratory of Bell Northern Research Ltd, Ottawa. In this measurement, the cleaved sample was placed in an X-ray diffractometer with its cleavage parallel to the surface of the mounting stage, making an equal angle with the incident and reflected beams. The angle of incidence was then fixed at the Bragg angle corresponding to this (112) plane and the diffracted X-ray intensity was measured with small deviation from the Bragg angle. The rocking curve obtained is shown in Figure11 and indicates a full width half maximum (FWHM) of about 200 arc seconds for this crystal. This value is much smaller than that reported by Tiwari et al [18} of about 680 arc seconds for their heteroepitaxially-grown thin film CuInSe2 structures but larger than the value of some 70 arc seconds reported by Cheuvart et al for CdTe [19] and certainly larger than values observed for present-day high quality Si and GaAs single crystal wafers for devices.

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Figure 10. (a) Photograph showing the positions of three Laue X-ray exposures on the cleaved {112} plane of ingot 119 and (b) the resulting superimposed image indicating the region to be one crystal.

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Figure 11. Rocking curve on a {112}cleaved sample of ingot 125, indicating a full width half maximum of about 200 arc seconds (taken at Bell Northern Research Ltd.).

Figure 12. X-ray diffraction peaks of a powder sample of CuInSe2 obtained from a stoichiometric growth run, showing the presence of all of the chalcopyrite structure peaks. Thin-Film Solar Cells, Nova Science Publishers, Incorporated, 2010. ProQuest Ebook Central,

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4. (c) Chalcopyrite Structure Confirmation X-ray diffraction was determined on powdered samples from stoichiometrically prepared CuInSe2 ingots and Figure 12 shows a typical diffractogram, obtained using 1.5405 Å X-rays from a Cu target. The pattern is consistent with JCPDS data for reflections for this chalcopyrite compound and the presence of (hkl) peaks, where l is odd, such as 101,103 and 105, is consistent with the presence of the chalcopyrite and not the sphalerite phase [20].

5. CLEAVAGE AND TWINNING

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5. (a) Cleavage After removal from the ampoule, the ingots frequently showed a common reflection over the outer curved surface formed by the ampoule inner surface, such as that indicated in Figure 13, where the light reflection extends over about 3 cm of the ingot. By subsequent cleavage, it was shown that this arises out of reflection from a common crystallographic plane. In cases where an external cavity existed in an ingot, in which growth unimpeded by the ampoule walls was possible, observation by microscope at the bottom of the cavity often showed faceting. Figure14 shows triangular faceting, characteristic of {112} planes in such a cavity.

Figure 13 . Reflection of light from a grown ingot of CuInSe2 indicating a single crystallographic plane.

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Figure 14. Triangular growth facets, in a {112} plane, as seen at the bottom of a side cavity in a stoichiometric CuInSe2 Bridgman-grown ingot.

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Table 2. Summary of Cleavage in Bridgman-Grown CuInSe2 Ingots1 Run No. 36 41 55 63 64 76 83 86 91

Location in ingot2 L C C C L U L L L

{hkl} plane 112 101 112 112 101 112 112 112 110

Area3 (cm2) 0.3 0.4 0.6 0.8 0.5 0.1 1.0 1.9 2.0

Max. dim.4 (cm) 0.6 0.8 1.2 1.5 1.2 0.4 1.3 2.2 2.0

Comments

Largest {101} cleavage. Largest {112} cleavage Only ingot with {110} cleavage plane.

95 L 101 0.8 1.0 105 L 112 0.8 1.1 110 L 101 0.1 0.6 119 L 112 1.3 1.9 125 L 112 0.7 1.5 149 L 101 0.7 1.0 1 Summary of cleavage planes observed in all the ingots prepared in this work. 2 Location of the cleavage in the ingot: L = lower region, C = central region and U = upper region. Total length of a typical ingot was 4cm. 3 Estimate of the area of the cleavage plane. 4 Length of maximum dimension of cleaved region.

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Figure 15. Chalcopyrite unit-cell schematic diagrams, indicating the three principal cleavage planes: (from left) {112}, {110}, and {101}.

Parts of a number of ingots were cleaved, using a knife edge, into several pieces at room temperature, generating surfaces, which were identified crystallographically by visual step patterns and Laue X-ray diffraction. Figure 15 illustrates the 3 low index crystal planes {112}, {110} and {101} planes in CuInSe2. Table 2 lists all the cleavage planes obtained on the ingots, including their dimensions. Here, it is seen that the maximum cleaved surface average diameters ranged from 0.4 to 2.2 cm on the predominant {112} and {101} planes. It should also be noted that most of the larger cleavage planes were located at the bottom end of the ingot, which was the first region to freeze in the growth process. Figure 16 shows a {112} cleavage plane, with an area exceeding 1 cm2, on ingot number 119 (stoichiometric), located just above the ingot tip. The surface is parallel to the ingot axis and thus also to the growth direction. Other ingots have indicated this same orientation, which is apparently different from that reported by Parkes et al [11], where the growth axis was stated to be “within a few degrees of” a direction, that is perpendicular to a {112} plane. This discrepancy, apart from a possible error in description, could arise from the fact that the angle between consecutive {112} planes is 70.5o, which differs from a right angle by only19.5 o. However, Baldus and Benz [17] reported, by contrast, Bridgman growth in a direction. Table 3 shows the values of angles measured between cleavage planes from the ingots, together with angles calculated assuming a perfect chalcopyrite lattice with a c/a ratio of exactly 2. The frequency of observation

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of these angles, in the last column, indicates that the most prevalent cleaved surfaces in the Bridgman-grown crystals were the {101} planes. This plane is indicated in Figure 15.

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Figure 16. Photograph of the bottom half of ingot 119, exhibiting a {112} cleavage plane of over 1 cm2 in area.

Table 3. Angles between principal cleavage planes in CuInSe2. Cleavage planes1 {HKL}--{hkl}

Angles between planes (degrees) Calculated2

Frequency of angle5

Observed

{112}--{112} {112}--{101}

70.53 70.5 1 39.23 39.4 2 75.04 75.0 3 {112}--{110} 35.26 35.03 1 90.00 90.1 2 {101}--{101} 53.13 53.3 2 78.46 78.6 3 {101}--{110} 50.77 50.8 2 {110}--{110} 90.00 -4 0 1 Only the {112} {101} and {110} chalcopyrite crystal planes are listed here. 2 Assuming the c/a ratio to be exactly 2 in the chalcopyrite unit cell for CuInSe2. 3 Observed between the {110} surface and the {112} microcleavage ridges. 4 A second {110} cleavage was not observed, so that this angle could not be checked. 5 Number of samples observed with angle.

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Figure 17. (a) Scanning electron micrograph of a cleaved {110} surface of CuInSe2 and (b) the corresponding diagram, indicating the {112} microcleavages.

As noted in Table 3, while {112} and {101} facets were often observed in cleaved and as-grown surfaces, {110} planes were less frequently seen. An SEM photograph of a {110} cleaved surface is shown in Figure 17 (a). Here, together with the diagram in Figure 17 (b), it is seen that the surface consists of microscopic parallel V-shaped grooves, oriented along a direction, where the sides of the grooves are {112} planes. The angle between the overall {110} plane and the {112} planes is 35o (35.23o calculated). Thus, on a microscopic scale, such apparent {110} planes are really not true cleavage surfaces, although they appear to be so on a macroscopic scale. Cleavages in the {112} planes correspond to {111} cleavages in cubic diamond-structured Si and Ge but the {101} cleavages in CuInSe2, on the other hand, appear to be unique to chalcopyrites and have not been reported in zincblende crystals, such as GaAs, where {110} is the dominant cleavage plane. A discussion of why {101} planes are predominant cleavage planes in CuInSe2 has been given elsewhere [21].

5. (b) Twinning Twinning lines are often visible on the surface of freshly-grown ingots but these can be made more prominent after chemical or thermal etching. Figure 18 (a) shows a sample surface etched in 5 % brome-methanol solution and Figure 18 (b) shows a sample after annealing for 2 hours in vacuum at 300 oC.

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Figure 18. Twin bands seen in polished samples as revealed by (a) chemical etching in 5% brome – methanol solution and (b) thermal etching under vacuum at a temperature of about 300 °C for 2 hours.

Twinning occurs in a {112} plane as illustrated in Figure 19 (a) and (b) for a (110) plane in the plane of the page, where crystal B is the twin of crystal A. Note that the two indicated directions make an angle of 54.74 o to the cleavage plane. Crystal B is the same as crystal A but rotated through 180 o about a [221] axis perpendicular to the (112) cleavage plane. Thus, it is an orthotwin [22] of the rotation type and not a reflection or paratwin. A further view of the twin plane is shown by a photograph of a crystal model in Figure 20 (a) and the corresponding diagram in Figure 20 (b).

Figure 19. Photomicrograph showing the {112} twinning plane boundary between two crystals having a common {110} cleavage plane (plane of the page). The ridges in the photograph (a) point in directions in both crystals and make an equal angle of about 54.7° with the {112} twinning plane, as shown in the schematic (b).

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Figure 20. (a) Crystal model of a twin in CuInSe2 in a {112} plane; (b) explanatory diagram of twin, where {110} is the plane of the page.

Figure 21 (a) shows a 150o angle observed between two adjacent cleaved {101} surfaces, viewed perpendicular to their intersection line. However, in a single chalcopyrite crystal, the only possible angles between [101} planes are 53,13o or 78.46o (Table 3). This 150o angle however, arises from a twinned region, as illustrated in Figure 21 (b), whereby the {112} twinning plane makes an angle of 75o (actually 75.04o) to the {101} plane of crystal A and an angle also of 75o to the other {101} plane of crystal B. Thus the angle between the two {101} cleavage planes comes to twice 75o or 150o.

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(a)

(b) Figure 21. (a) Photograph of a twinned CuInSe2 sample, viewed perpendicular to the intersection of a pair of cleaved {101} planes from twinned crystals A and B, and (b) schematic showing the twinned region in this sample.

6. DEVIATION FROM STOICHIOMETRY 6. (a) Composition in Main Part of Ingot In a first set of experiments with nonstoichiometry, 7 Bridgman ingots were prepared with different deviations from stoichiometric starting proportions of the three elements involved plus one stoichiometric ingot for

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comparison purposes. Compositions were determined by EPMA in the manner described in section 3 (c) above in which each ingot was first cut lengthwise into two halves, with one half prepared for the measurements. The results are shown in Table 4 for the 8 growth runs. It is noted that after growth, the EPMA compositions are nearer to stoichiometry than the original starting compositions. This fact is brought out more strikingly in the ternary composition diagram of Figure 22, where the solid squares indicate starting proportions and the open squares the EPMA-determined compositions after synthesis and growth [15]. Furthermore, it is apparent that the conductivity type was determined more by the starting composition, rather than the aftergrowth composition (column 12, Table 4).

Figure 22. Ternary composition diagram showing a plot of the starting compositions of the ingots used in this study along with their final compositions, as measured after growth.

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Figure 23 is a ternary diagram, showing the present determined EPMA compositions (circles) over a larger scale, together with results (triangles) of Neumann and Tomlinson [23], who studied deviations from stoichiometry by annealing their CuInSe2 samples under minimum and maximum selenium pressure. Noting the limited accuracy of EPMA determinations of  0.5%, the combined points indicate that n-type behaviour occurs mainly to the right of a line corresponding to 25 at % of indium. In other words, n-type conductivity in Bridgman material takes place essentially when [In] > 25 at %, where the square brackets indicate atomic percentage. Thus, by this criterion, having an excess of Se to convert from n-type to p-type only operates when In is not in excess.

Figure 23. Ternary composition diagram showing the present results (circles) plotted along with those of Neumann and Tomlinson [23] (triangles).

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Table 4. EPMA results from single-phase region of runs with deviation from stoichiometry.

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Run No.

1

138 140 141 142 143 144 145 146

Final Composition1

Starting Composition

Comments3

Cu (at%)

In (at %)

Se (at %)

Cu/I n Rati o

Metal/S e Ratio

Cu (at%)

In (at% )

Se (at %)

Cu/In Ratio

Metal/S e Ratio

Conduc t. Type2

28.5 27.0 20.0 24.0 31.0 25.0 28.0 22.0

23.8 27.0 30.0 24.0 23.0 25.0 21.0 31.0

47.7 46.0 50.0 52.0 46.0 50.0 51.0 47.0

1.20 1.00 0.67 1.00 1.35 1.00 1.33 0.71

1.09 1.17 1.00 0.92 1.17 1.00 0.96 1.13

26.2 25.1 22.6 25.0 26.5 24.6 25.5 24.5

25.1 25.4 27.1 26.0 25.2 25.9 25.4 25.8

48.7 49.5 50.3 49.0 48.3 49.5 49.1 49.7

1.04 0.99 0.83 0.96 1.05 0.95 1.01 0.95

1.05 1.02 0.99 1.04 1.07 1.02 1.04 1.01

p n n p p p p n

Cu excess Se deficiency In excess, Cu deficiency Se excess Cu excess Stoichiometric Cu and Se excess, In def. In excess, Cu and Se def.

Representative average compositions of the data points for each ingot, as determined by EPMA, for the homogeneous single-phase region. 2 Conductivity type as determined from the hot probe method. 3 Elemental excess or deficiency in the starting composition with reference to proportions 25, 25 and 50% for Cu, In and Se, respectively.

32

C. H. Champness Table 5. (a) Copper and copper-indium phases at end of ingot.

Atomic conc. (%)1 Cu In

Elemental Possible Compound Present in run ratio composition2 known no. Cu/In 63.7 36.3 1.75 Cu7In4 Yes 140, 141, 146 68.9 31.1 2.22 Cu9In4 Yes 140 79.7 20.3 3.93 Cu4In Yes 143 62.7 37.3 1.68 Cu5In3 No 146 56.5 43.5 1.3 Cu4In3 No 146 ≥ 90.0 ≤ 10.0 ≥ 9.0 Cu + Cu4In? Yes 138,143 1 Atomic concentrations, determined by EPMA, averaged over several data points and rounded-off. The selenium concentration was below the detection limit of the electron micro probe, where the resolution was about ± 0.5%. 2 Chemical formula having the nearest elemental ratio.

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Table 5. (b) Selenium and selenide phases at end of ingot Atomic conc. (%)1

Elemental ratio

Cu 50.8 BD

In BD 51.0

Se 49.2 49.0

In/Se 1.04

BD BD

58.2 BD

41.8 99.9

1.39 -

1

Cu/Se 1.03 -

Possible composition2

Compound known

Present in run no.

CuSe InSe

Yes Yes

-

In7Se5 Se

No Yes

142, 145 140, 141, 146 141, 146 142, 145

Atomic concentrations, determined by EPMA, averaged over several data points and rounded off. BD indicates concentrations below the detection limit of the electron micro probe. The error in the measurements is estimated to be about ± 0.5%. 2 Chemical formula having the nearest elemental ratio.

6. (b) Composition along Ingot EPMA compositions determined with distance along and across the ingots showed general uniformity, except for the last zone to freeze (LZTF). For example, Figure 24 shows data for ingot number 143, prepared with excess Cu, where the first 80 % of the ingot had uniform elemental compositions but, in the last 20 %, which contained multi-phases, nonuniformity was evident. In ingots with excess copper, precipitation occurred at the very end of the ingot in the form of copper or copper alloy fibres with a diameter of about 30 microns, as indicated in Figure 25 corresponding to the XRD pattern in Figure 26. Figure 27

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shows photographs of some of the other secondary phases observed in the LZTF of the different ingots and Tables 5 (a) and (b) list the possible compounds detected. Thus, the action of Bridgman growth for nonstoichiometric deviations from the ratio 1:1:2 for respectively Cu:In:Se, is that some binary compounds plus excess elements are formed in the liquid but are gathered up and swept to the LZTF, leaving the main gradient-frozen region with near-stoichiometric ternary proportions.

Figure 24. Plot of elemental concentration against distance along the single-phase region of ingot 143, which was prepared with excess copper.

Figure 25. Copper fibers with a diameter of about 30µm precipitated at the last surface to freeze of one of the excess copper ingots.

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Figure 26. X-ray diffraction pattern of the copper fibres shown in Figure 25, together with 4 diffraction lines for pure copper.

7. ANNEALING AND ETCHING 7. (a) Heat Treatment of p-Type Samples Stoichiometrically prepared p-type material heated under vacuum above 250 to 300 oC for just 8 minutes has been found sufficient for conversion to n-type conductivity [24]. This has been ascribed to out-diffusion of selenium, which is evident particularly for a sample in a heated ampoule with a temperature gradient, where the Se can deposit on the cooler part or, in a tube closed at one end, and pumped on from the other end. Such type conversion can also occur even at room pressure under argon at 500 oC with annealing times of days [25], as indicated in Table 6. Here, in originally p-type samples, the junction depth was found by hot probing and etching, in a solution [26] of HNO3 + HF +H2O in the ratio of 1:1:1, whereby a red stain was deposited on a p-type region within each wafer heated for more than 72 hours, leaving an n-type outer layer with the thicknesses indicated in column 6 of Table 6. These unstained thicknesses, interpreted as junction depths, d, are plotted against the square root of the annealing time, t, in Figure 28, yielding an empirical relationship of the form:

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(d  d 0 )(mm)  0.4 t (days) , where d0 is a constant distance of just over 0.1 mm. The form of this relationship is consistent with the variation illustrated in Figure 29, for assumed simultaneous outdiffusion of acceptors and donors, with different diffusion coefficients DA and DD respectively. With simplifications, such as DA»DD, diffusion theory [25] leads to a formula of the type: N D0 d  d0  D A t , which is the same form as that found experimentally, N A0

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where N D0 and N A0 are respectively the original pre-outdiffusion donor and acceptor concentrations. Extrapolating the diffusion data from the work of Tell and Bridenbaugh [27] from 300 oC to 500 oC, yields a DA value of 7.3 x 10-7 cm2s-1, assumed to correspond to Se. Comparison of the experimental and theoretical relations then yields a NDo/NAo ratio of 0.09 for the samples. In other words, the original material was partially compensated, containing by this calculation, about 8.3 at % donors and 91.7 at % acceptors. Outdiffusion of selenium was confirmed from Auger profiles (not shown) indicating a 2 to 4 % reduction of this element in p-type CuInSe2 samples, annealed in argon for 2 hours at 350 oC [25]. The work of Tell and Bridenbaugh [27] suggests that metals, such as copper, zinc and cadmium, deposited on p-type material act as a sink for outdiffusing selenium. This is consistent with work in our own laboratory, where heated indium and bismuth depositions produced homojunctions [24] with clear photovoltaic characteristics.

7. (b) Chemical Etching No extensive and definitive work on chemical etching of CuInSe2 monocrystals has been reported. However, preliminary tests with some etches used in this laboratory are set out in Table 7. The chromic-sulphuric solution was found to bring out triangular etch pits on {112} faces (see Figure 30). The reaction with the nitric-hydrofluoric etch is very strong and in the 1:1:1 proportion with water, as mentioned above, it gave a red stain on p-type material under illumination. Brome-methanol is often mentioned as a moderate etch. Thermal etch pits can also be observed and Figure 31 shows triangular pits on a {112} polished surface, after heating under vacuum for 4 hours at 600oC.

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C. H. Champness Table 6. Samples, originally p type, in quartz tube open at both ends through which argon flowed. Heat treatment

Conductivity type

Sample No.

Temp. (°C)

Time (h)

Before anneal

After anneal

Z119-4 G1-1 G1-2 G1-3 G1-4 G1-5 G1-6 G1-7 G1-8

450 500 500 500 500 500 500 500 500

24 2 22 29 72 120 168 240 288

p p p p p p p p p

p p p p n/p n/p n/p n/p n/p

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p = p type

n = n type

Depth uncoated region after etching (mm) … … … … … 0.8 1.0 1.2 1.3

Conclusion

Type conversion from p to n type can occur at atmospheric pressure (and, in the present conditions, began to take place after more than one day at 500°C in flowing argon)

n/p = outer part n, inner part p

Table 7. Chemical Etching Characteristics of Bridgman-grown CuInSe2 Crystals Etchant K2Cr2O7 + H2SO4 (1:9 by wt.) Br + methanol (1 vol. %) HCl + HNO3 (1:1 by vol.) HF + HNO3 (1:3 by vol.) HF + HNO3 + H2O (1:1:1 by volume)

Action moderate

Etch pits yes

moderate

yes

strong

yes

very strong

not clearly defined

very strong

not clearly defined

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Etch pit shape Equilateral triangles on {112} elongated sectorshaped on {101} Small irregularly-shaped on {112} and {101} Large and circular on {112} and {101} Darkened and coarse surface Yields a red stain on ptype under illumination

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Figure 27. Photographs of secondary phases observed in the last region to freeze in (a) ingot 140 prepared with a deficiency of selenium, (b) ingot 143 prepared with excess copper, (c) ingot 145 prepared with a deficiency of indium, (d)-(f) in ingot 146 prepared with an excess of indium.

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Figure 28. Plot of depth (d) of etched pn boundary against square root of annealing time (t) at 500 °C in Ar.

Figure 29. Schematic of outdiffusion concentration profiles assumed after annealing for two different times t1 and t2, with

N A0  N D0

and DA > DD.

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Figure 30. Triangular etch pits on a {112} plane revealed after etching for 5 minutes with K2Cr2O7 + H2SO4 (1:9 by weight).

Figure 31. Thermal etch pits obtained on polished {112} surfaces, treated under vacuum, showing triangular features; sample annealed at 600°C for 4 hours.

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8. OXYGEN IN COPPER As described in section 2(c), the practical remedy for avoiding ingot sticking and voids was found to be the use of a boron nitride flamed-in coating within the ampoule, prior to compound synthesis and growth. Arising from this, a study was made [28] to find out if the sticking came from oxygen and if so, from whence its source, the quartz or starting elements In, Se or Cu? These elements were stated to be of 99.999 % purity in respect of metallic impurities. The quartz and selenium first were eliminated from having an effect. Pre-heating the indium pellets for 48 hours at 1100 oC in a closed ampoule with boron nitride before preparing the compound still gave the undesired sticking and voids but doing the same thing, with the copper pellets yielded an ingot free of these defects. Accordingly, in a special experiment, the nominally five nines purity copper pellets (melting point 1086 oC) were heated alone under vacuum at 1000 oC, while the gaseous agents present were determined with a quadrupole gas analyzer, which indicated several peaks as a function of molecular mass number. Two of these corresponded to mass numbers of 16 and 32 for oxygen (Figure 32 (b)). For comparison, heating an empty ampoule under similar conditions showed only very small oxygen peaks (Figure 32 (a)), thus excluding the quartz from suspicion. Hence, the two strong peaks pointed to the copper as the oxygen source. Next the copper was again heated alone under vacuum, without boron nitride, at 1030 oC, with continuous pumping for up to 4 days. During this time, Figure 33 shows that the two oxygen peaks progressively decreased in magnitude over the first 3 days with continuous pumping. Using the four-day vacuum-annealed copper pellets, with the other untreated In and Se pellets, to synthesize CuInSe2, resulted in ingots completely free of adhesion to the ampoule and free of cracks, deposits and voids. Further, using commercial special oxygen-free copper (from Otokumpo Ltd.) without boron nitride, also yielded an ingot which was again found to be free of the undesirable characteristics. Thus, this work clearly showed that oxygen in the original copper was the source of the ingot sticking and voids.

9. ADDITION OF SODIUM It is well established experimentally that the presence of sodium in CuInSe2based thin film solar cells is beneficial to device performance. In monocrystalline Bridgman material, the addition of this element to the melt was

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shown by Wang et al [29] to cause a sign conversion from p- to n-type for amounts greater than 0.05 at % Na. They also found the same type conversion if the sodium was diffused into an already-grown p-type sample. Recent work [30] confirmed this but showed that the conversion did not occur if sufficient excess selenium over stoichiometry was also present A possible explanation of the sign inversion is that the sodium forms on the surface, or in a thin surface layer, of the crystal and there acts as a sink for outdiffusing selenium. This enhanced loss of Se, increases the concentration of Se vacancies, which act as donors, causing the change from p- to n-type. Some precipitation of copper was also observed in the experiments with added Na, possibly due to enhanced outdiffusion of Cu as well. This would result in an increase in the Cu vacancies and hence in an acceptor concentration sufficiently large to prevent type conversion.

Figure 32. Gas analyzer partial pressure peaks as a function of molecular mass number recorded during pumping and heating at about 1000 °C of (a) ampoule containing only the quartz crucible, (b) ampoule containing quartz crucible plus five-nines-purity copper. (DPO: diffusion pump oil).

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Figure 33. Gas analyzer partial pressure peaks as a function of molecular mass number for ampoule containing five-nines purity copper pumped-on at about the melting point of copper. Scans were taken after 1, 2 and 3 days of pumping.

10. CONCLUSION The main obstacle in the growth of CuInSe2 by the Bridgman method, in this work, has been the adhesion problem. To a considerable extent, this prevented large monocrystals from being obtained and used for other studies. The boron nitride coating method, developed in this study, was successful in eliminating the adhesion completely and enabling adhesion-free and void-free ingots to be obtained routinely. This also allowed, for the first time, the cleavage of surfaces of area greater than 1 cm2 to be obtained on CuInSe2. The inclusion of graphite in the ampoule, in the work of L.S. Yip et al [31], was found to be equally

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effective in eliminating the adhesion problem. However, they reported that the graphite, when deposited as a film on the inside of the ampoule wall, flaked-off and was invariably incorporated into the melt. This did not occur with the boron nitride coating used in this work, since the intense flaming step carried out, caused a hardening of the BN layer. Thus, no BN was observed to be incorporated, macroscopically, into the melt. The BN coating method does not require the BN to be in the form of a crucible, which alone, without a containing quartz ampoule, would be more difficult to seal off. Since the BN coating can be deposited to a controlled thickness onto the ampoule wall, it eliminates the need for a crucible. In the work of Bachmann et al [32] on CuInSe2, Yasuda et al [33] on CdTe and Omino and Suzuki [34] on ZnSe, using a pre-reacted charge, a BN crucible was employed inside a quartz ampoule. In the one ampoule method adopted in that work, where an elemental charge was used, a BN crucible can sometimes crack inside the quartz ampoule due to accumulation of reaction products between the crucible and ampoule walls. Thus, the coating method offers a less troublesome alternative for using the BN for the crystal growth, if the one ampoule technique is used. As far as adhesion goes, the BN does not have to be in contact with the melt in order to prevent the sticking. It has been shown that the use of a block of BN inside the ampoule, away from the charge, has a similar action. This is also consistent with the results of L.S. Yip et al [31] with the use of graphite inside the quartz ampoule. In a special study already described [28], it was shown that the boron nitride acts as a getter in removing oxygen efficiently from the commercial "high purity" copper. While the advantage of having BN in the quartz ampoule is clear, it posed, nevertheless, a possible risk of microscopic contamination of the CuInSe2 charge. Analysis of CuInSe2, grown in an ampoule with BN, was carried out with SEM/EDX. However, boron was not detected. In the work of Yasuda et al on CdTe [33] and Omino and Suzuki [34] on ZnSe, where BN crucibles were used, no measurable amount of boron was found in the ingots. In the work of Ciszek [10], on solid solution growth of compounds, where a B203 encapsulant was used, the boron trace in CuInSe2 was below the detection limit of 30 ppm (by weight). However, boron contamination was reported by Yoshida et al [35] for ZnSe if unpurified BN crucibles were used. In any case, further analysis for CuInSe2 would be required, especially if boron is shown to have a doping role in CuInSe2. The BN method made possible the growth of ingots with large monocrystals having dimensions comparable to the size of the containing quartz ampoule. However, the vertical Bridgman crystal growth procedure, in this laboratory, was by no means fully optimized and many parameters still need

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to be carefully studied in order to understand their effect on the ingot solidification. These include: the soaking temperature above the melting point of CuInSe2, the temperature of the bottom zone and also the temperature gradient at the compound freezing mark. In this work, the gradient of temperature was fixed at about 70°C/cm, which was believed to be high enough to avoid constitutional supercooling. However, this needs to be verified by experiment. The absence of ampoule rotation during the soaking and the growth did not seem to hinder severely the growth of large monocrystals; however, it would appear to be beneficial from the point of view of thermal homogenization of the ingot, particularly during the solidification stage. In the work of Weng et al [36], an accelerated crucible rotation technique (ACRT) was employed for the vertical Bridgman growth of CuInSe2. The upper furnace zone temperature, in this work, was kept at 1100 °C in almost all the runs. While this parameter may not seem to be critical, it was reported [37] that, such an elevated temperature, above the melting point of the compound during the growth, may cause the solid-liquid interface shape to become concave, as viewed from the solid, which is undesirable. Some interrupted growth experiments, using extra copper in the melt [38], not described, showed that the apparent solid-liquid interface shape in these ingot was concave with respect to the solid. However, it is possible that the quenching action itself influences the concavity and thus, the surface indicated could misrepresent the actual shape of the interface. This is because, in the quenching, heat is lost more rapidly from the walls than from the interior of the ingot, which in turn causes the excess copper to segregate in the central part of the ingot as the walls solidify first. Be this as it may, the relative degree of concavity could indeed increase with growth, as evidenced by the presence of larger monocrystals near the bottom of the ingot compared with those in the last region to freeze. The results of the experiment on the variation of growth rate indicate that slower lowering rates give larger crystallites. Here, it would appear that a growth rate of some 10 mm/hr or less is sufficient for obtaining ingot-sized monocrystals, at least using the present growth conditions of temperature gradient and ampoule size. For example, using a larger diameter ampoule may require a much slower growth rate in order to maintain the thermal equilibrium needed. The quartz ampoule used, had a wall thickness of about 2 mm, which seemed to be strong enough to withstand the very high saturated vapour pressure of the selenium of more than 40 atmospheres at the soaking temperature of 1100 °C. This soaking temperature was also apparently high enough to synthesize

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completely the CuInSe2 compound, since no other elemental or binary phases were observed in the ingots, after the growth. While the X-ray diffraction pattern on a powdered sample indicated only the presence of the chalcopyrite phase, the presence of the sphalerite phase of CuInSe2 cannot be completely ruled out, particularly since the second zone temperature of the furnace was set to about 600 to 700 °C, in order to obtain a high temperature gradient at the freezing mark, which was less than the 810 °C sphalerite-to-chalcopyrite transition temperature. This could have hindered the phase transformation, rendering it incomplete. A more careful study is therefore needed to verify if the sphalerite phase was indeed present in these ingots. While no microscopic cracks were observed in any of the ingots, very limited macroscopic cracking was observed in some runs where the cooling rate was too fast and particularly evidenced in the quenched ingots. Thus, after growth, slower cooling rates below those used in this study, of around 25 to 30 oC/hr may be recommended to reduce the possible incidence of cracks in these crystals [10]. A reduced ampoule lowering rate may also be important, since Baldus and Benz [17] used a rate of only 0.33 mm/hr to avoid cracking. These workers believed the transition from zincblende to chalcopyrite at 810 oC was the cause of cracking. It should be clearly pointed out here that examination of cut and polished surfaces of the grown CuInSe2 under the scanning electron microscope indicated, apart from twinning, featureless photomicrographs (not shown) with no evidence of the inclusions of the type reported by Mullan et al [39]. This was the case with all parts of the present stoichiometrically prepared ingots and also even with the first part of nonstoichiometic ingots, where only the last zone to freeze contained multiphases. The dominant cleavage in the Bridgman-grown CuInSe2 crystals was along {011} or {101} planes, which was also confirmed to be the case in solutiongrown crystals [5]. This is quite different from zincblende crystals, where dominant cleavage occurs along {110} planes. However, in CuInSe2, the observed {110} planes were found to consist microscopically of parallel Vshaped grooves, where the side-walls were {112} planes. The second most prevalent cleavage planes were {112}, which correspond to {111} planes in cubic crystals. Twinning, which was often observed in the Bridgman crystals, was the rotation type and occurred along {112} planes. The experiments carried out on annealing of stoichiometric CuInSe2 indicated outdiffusion of selenium with heating at sufficient temperatures under reduced pressure and even at room pressure. With sufficient heating time, p-type

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samples were converted to n-type. This result is consistent with the concept that selenium vacancies act as donors. Deposition of metals, like bismuth, zinc, cadmium and copper on CuInSe2 appears to accelerate the selenium outdiffusion and, on p-type material, it creates a pn-homojunction, rather than a metalsemiconductor Schottky junction. Chemical etching is very important in the processing of the compound for measurements and device fabrication but very little has been published on this and more work needs to be done. In the present studies on nonstoichiometry in the compound, it was found that excess of indium over stoichiometry (i.e. [In]> 25 at %) produced n-type material. Excess copper produced higher resistance p-type material along with precipitated copper. Excess selenium produced also p-type material with precipitated selenium. Despite deviations from stoichiometry in the starting melt compositions, the main part of the resulting ingot after compound synthesis and crystal growth was always near to a stoichiometric composition. Thus, the process was “forgiving” of initial deviations from the 1:1:2 composition. In ingots prepared with intended composition deviation, binary compounds were also found in the last-zone-to-freeze, but the first approximately 80 % of the ingot was always single phase chalcopyrite, again an example of “tolerance” in the process. In respect of the effect of sodium, current monocrystalline research is being carried with this element added to the melt of CuInSe2 prior to crystal growth. At the present time, it found that the sodium forms on the surface of the crystal but does not enter it beyond a very thin layer. Here, it appears to act as a sink for the outdiffusing selenium atoms. In originally p-type material, this changes it to ntype. However, if excess selenium is present in the melt, the semiconductor remains p-type. It is well known that sodium incorporated into thin polycrystalline solar cells increases their photovoltaic performance but this effect has not, so for, been demonstrated in cells using monocrystalline CuInSe2 absorbing substrates. Finally, attempts were made in the present laboratory, at seeding [38] (not described), where a small single crystal of Al2O3, SiO2 or Si was placed within the ampoule tip, prior to growth. However, this and also changing the ampoule taper angle did not yield larger CuInSe2 monocrystals.

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11. ACKNOWLEDGMENTS The author would like to thank his graduate research students for their contributions to this work, particularly Z. Shukri, along with G.I. Ahmad, H. Du, L.S. Yip and H. Myers. He would also like to thank Dr. I. Shih for discussions and to acknowledge the partial support of the research by the Natural Sciences and Engineering Research Council of Canada. He also wishes to thank Dr. I. Bassignana of Bell-Northern Research for the rocking curve measurement.

12. REFERENCES [1]

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[2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14]

Repins, I; Contreras, M; Romero, M; Yan, Y; Metzger, W; Li, J; Johnston, S; Egaas, B; DeHart,C; Scharf, J; McCandless, BE; Noufi. R. Conference Record of the 33rd IEEE Photovoltaic Specialists Conference. San Diego, CA, 2008. Paorici, C; Zanotti, L; Romeo, N; Sberveglier, G; Tarricone, L. Solar Energy Materials, 1979, 1, 3-9. Cisek, TF. J. Cryst. Growth, 1984, 70, 405-410. Miyake, H; Sugiyama, K. J. Cryst. Growth, 1982, 125, 548-552. Lyahovitskaya, V; Richter, S; Frolow, F; Kaplan, L; Manassen, Y; Gartsman, K; Cahen, D. J. Cryst. Growth, 1999, 197, 177-185. Champness, CH; Cheung, T; Shih, I. Solar Energy Materials & Solar Cells, 2007, 91, 791-800. Du, H; Shih, I; Champness, CH. J. Vac. Sci. Technol. 2004, A22, 10231026. Vahid Shahidi, A; Shih, I; Araki, T; Champness, CH. J. Electronic Materials., 1985, 14, 3, 297-310. Fearheiley, ML. Solar Cells, 1986, 16, 91-100. Ciszek, TF. J. Electronic Materials, 1985, 14, 451-460. Parkes, J; Tomlinson, RD; Hampshire, MJ. J. Cryst. Growth, 1973, 20, 315-318. Haupt, H; Hess, K. Ternary Compounds, 1977. Inst. Phys. Conf. Series No. 35, Edinburgh, 1977, 5-12. Arya, RR; Warminski, T; Beaulieu, R; Kwietniak, M; Loferski, JJ. Solar Energy Materials, 1983, 8, 471-481. Shukri, ZA; Champness, CH; Shih. I. J. Cryst. Growth, 1993, 129, 107110.

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[15] Shukri, ZA; Champness, CH. J. Cryst. Growth, 1998, 191, 97-107. [16] Shukri, ZA; Champness, CH. Acta Crystallographica, 1997, B53, 620630. [17] Baldus, A; Benz. KW. J. Cryst. Growth, 1993, 130, 37-44. [18] Tiwari, AN; Blunier, S; Zogg, H; Zelezny, V; Schmid, D; Schock, HW. Proceedings of the 12th European Photovoltaic Solar Energy Conference, Amsterdam, 1994, 625-628. [19] Cheuvart, P; El-hanani, U; Schneider, D; Triboulet. R. J. Cryst. Growth, 1990, 101, 270-274. [20] Kavcar, M; Carter, MJ; Hill. R. Solar Energy Materials and Solar Cells, 1992, 27, 13-23. [21] Shukri, ZA; Champness, CH. Surface Review and Letters, 1998, 5, 419422. [22] Holt, DB. J. Phys. Chem. Solids, 1964, 25, 1385-1395. [23] Neumann, H; Tomlinson. RD. Solar Cells, 1990, 28, 301-313. [24] Shih, I; Champness, CH; Shahidi, AV. Solar Cells, 1986, 16, 27-41. [25] Champness, CH; Ahmad, GI. J. Vac. Sci. Technol., 2000, A18(2), 693696. [26] Tell, B; Wagner, S; Bridenbaugh, PM. Appl. Phys. Lett., 1976, 28, 454455. [27] Tell, B; Bridenbaugh, PM. J. Appl. Phys., 1977, 48, 2477-2480. [28] Shukri, ZA; Champness, CH. J. Cryst. Growth, 1993, 166, 708-711. [29] Wang, HP; Shih, I; Champness, CH. Conference Record of the 28th IEEE Photovoltaic Specialists Conference. Anchorage, AK, 2000, 642-645. [30] Champness, CH; Myers, HF; Shih, I. Thin Solid Films, 2008, 517, 21782183. [31] Yip, LS; Shih, I; Champness; CH. J. Cryst. Growth, 1993, 129, 102-106. [32] Bachmann, KJ; Fearheiley, M; Shing, YH; Tran, N. Appl. Phys. Lett., 1984, 44, 407-409. [33] Yasuda, K; Iwakami, Y; Saji, M. J. Cryst. Growth, 1990, 99, 727-730. [34] Omino, A; Suzuki, T. J. Cryst. Growth, 1992, 117, 80-84. [35] Yoshida, H; Fujii, T; Kamata, A; Nakata, Y. J. Cryst. Growth, 1992, 117, 75-79. [36] Weng, WS; Yip, LS; Shih, I; Champness, CH. Can. J. Phys., 1989, 67, 294-297. [37] Vere, AW. Crystal Growth, Principles and Progress; Plenum Press: New York, NY, 1987, 76.

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[38] Shukri, ZA. Bridgman growth of CuInSe2 monocrystals for photovoltaic cell research; Ph.D Thesis; McGill University: Montreal, QC, 1996. [39] Mullan, CA; Kiely, CJ; Casey, 8.M; Imanieh, M; Yakushev. MV; Tomlinson, RD. J. Cryst. Growth, 1997, 171, 415-424.

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Chapter 2

REVIEW ON SULFUR COMPOUNDS DEPOSITED BY SPRAY PYROLYSIS TECHNIQUE AS BUFFER AND ABSORBER THIN FILMS FOR SOLAR CELLS Karem Boubakera and Mosbah Amloukb* Unité de Physique des Dispositifs à Semiconducteurs (UPDS), Campus Universitaire 2092 El-Manar Tunis/ 63 Rue Sidi Jabeur 5100 Mahdia, Tunisia. b Faculté des Sciences de Bizerte & Unité de Physique des Dispositifs à Semiconducteurs (UPDS),Campus Universitaire 2092 El-Manar Tunis, Tunisia.

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a

ABSTRACT This chapter deals with some physical properties of Sulfur binary and ternary thin films used as buffer and absorber layers in solar cells and prepared using economic spray pyrolysis technique. First, a brief historic of the binary and ternary compounds is presented, then, the spray technique fundaments and experimental features are detailed. Finally, the physical characterization (XRD, SEM, AFM, optical, electrical as well as thermal

*

Corresponding authors: : e- mails: boubaker_karem@ yahoo.com and [email protected]

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Karem Boubaker and Mosbah Amlouk measurements) results are presented showing the possibility of use of such films in the solar conversion devices.

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I. GENERAL INTRODUCTION According to Kyoto Protocol cumulative and despite the price per kilowatt hour of PV module, The PV system installation has been accepted as a public commitment in the world. This is one reason that PV market is successfully expanding every year. For terrestrial applications I-III-V2 and especially copper indium gallium diseleneide CIGS and based thin film technology is well acknowledged to have high potential to attain lower production cost then relatively crystalline silicon as total output power and efficiencies are approaching the requirements for the power market. Further improvements have led to a new generation process from many advanced companies (Shell Solar Industries, Würth Solar GmbH and Global Solar Energy…) demonstrating champion efficiencies of over 13% for 60x90 cm2 large area modules in a new pilot line [1]. The goal of production research on GIGS based thin film PV technology is to fill the gap between the research of small area solar cells focusing on the demonstration of high potential on the performance by targeting a 20% [2] as solar efficiency and the production of large-area modules competitive enough to the performance of polycrystalline Si solar modules. The issues that should be completed during the production research are understood as follows: 1. To develop and enhancement the deposition apparatus capable to the large surface production and reduce the production cost of both absorber/buffer and transparent conductive oxide window. 2. To benefit from low-cost and abundant materials given high reproducibility and possible recycle of them. 3. Solar power does not lead to any harmful emissions during operation, but the production of the panels leads to some amount of pollution which it must be reduced and avoided. As possible alternatives to CIGS, CuSnS based semiconductors ( Cu2SnS3, Cu3SnS4, Cu2ZnSnS4.... ) are also expected to make good solar cell absorber layers [3]. The recourse of the international community to these new materials is due to

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the relatively low abundance and production rate of Gallium and Indium ( 80 tonnes and 510 tonnes respectively, reported by U.S. Geological Survey 2008 ). A few buffer sulfur materials have been tested in solar cells (ZnIn2S4, ZnS, Zn(O, S, OH)x [4-6]). Among them indium sulphide (In2S3) grown by spray pyrolysis is used as buffer for CuInS2 solar cells with efficiencies in the range 79.5 % [7, 8]. The physical properties of these materials have not been extensively studied especially electric ones and only little has been published on the optical and electrical ones which is of interest for optoelectronic modelling of the complete device structure and constitutes the basic motivation for future works. However, as researchers in the photovoltaic domain and have been working since 1987 on the growth of sulphur materials as buffer and absorbers sprayed thin films [9-36], we hope that this review study provides some valuable information for scientist community working in photovoltaic solar cells over the world.

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II. WHY SULFUR MATERIALS? Among the chalcogeneide compounds, the sulfur ones have been extensive studied during the last two decades because of the numerous physical and chemical properties of sulfur element:    



The sulfur is an abundant in nature and a relatively cheap element. The sulfur materials show relatively high chemical and thermal stabilities. For ecological reasons it is desirable to work on sulphides than on other VI elements such as selenides and tellurides. With this element, it can be made buffer layers with various band gap ( 2.2 for InS to 3.7 eV for ZnS ) as well as large absorber compounds which have various gap energy ( from 1.15 for Cu2SnS3 to 1.9eV for AgInS2 ). The efficiency of Cu(Ga,In)Se2 devices has been improved due to the incorporation of sulfur in the film, the band gap can be changed gradually and a better alignment at the interface buffer/ absorber results. Those alloys Cu(Ga,In)(Se,S)2 show an increased open circuit voltage and efficiency compare to free-sulfur ones.

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It was for this reason and from the economical point of view as well as environmental interest, the synthesis of thin films based on sulfur element via a cheap process is essential in order to consolidate photovoltaic energy.

II.1. Sulfur and Thiourea II.1.1. Sulfur Sulfur element, considered as the tenth most abundant element in the universe, has been known since 1777 by the French scientist Antoine Lavoisier who convinced the rest of the scientific community that sulfur was an element.

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Table 1. Physical characterization of sulfur element Atomic Number Atomic Weight Melting Point Boiling Point Density Color Phase at Room Temperature Electron configuration

16 32.065 388.36 K 717.75 K 2.067g/cm-3 a pale yellow Solid non-metal 1s22s22p63s23p4

Sulfur element has chalcogen as group name in periodic table of elements. This element has some physical properties listed in Table 1. From its electron configuration, the oxidation states are: +6, +4 and -2. The majority of the sulfur produced today is obtained from underground deposits, usually found in conjunction with salt deposits, with a process known as the Frasch process. Sulfur is a component of many common minerals and binding with some metals ( PbS, FeS2, HgS and so on), but, nearly 25% of the sulfur produced today is recovered from petroleum.

II.1.1. Thiourea During the pyrolysis process, the deposited sulfur materials is generally obtained via endothermic reactions in which the sulfur element is provided by

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the thiourea as S2- anion . Thiourea formulate is SC(NH2)2 and it has a molar mass of 76.12g. It is a planar molecule where bond distance C=S is 1.60±0.1 Å and has two tautometric forms:

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Thiourea is commercialed as a white powder and its global annual production is around 10,000 tons. It has a relatively great solubility in water and it emits oxides of carbon, sulfur and nitrogen when heated to decomposition in air. It has a density of the order of 1.4g/cm3, a boiling point of the order of 150°C and a melting one of 175°C.

III. EXPERIMENTAL SETUP AND FUNDAMENTS III.1. Spray Principe This technique is used commonly in heavy coatings and to prepare oxides and chalcageoeides thin films materials. The material to be deposited (here sulfur compounds) arise from a liquid solution which contains precursors as salts melted and dissolved in it. Afterwards, solution formulated by dissolving salts is entrained in air or inerter gas jet from the atomizer or nozzle (Figure 1.a). A conical jet of droplets is then formed and deposited onto a heated substrate. The substrate is generally glass, but other substrates can be used as well. Finally, endothermic reaction gives then in principle the desirable material. Because of the toxicity of volatile gas, the global spray apparatus must be lying in a hood with laminar flux. For all experiments, the solution and gas flow rates were varied in 2 -15 cm3.min-1 and 2-4 l.min-1 domains respectively. The deposition state can be manipulated through spray parameters and can be used to significantly control coating properties, such as molarities and the rate flux of solution and carrier gas. Indeed, nitrogen gas is advised to ovoid and

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minimized chemisorptions of oxygen during the thin films of sulfur material products. The deposited spray materials arise from endothermic reactions between precursors, e.g.: ZnCl2 + SC(NH2)2 + 2 H2O → ZnS + CO2 + 2NH4Cl CuCl2 + InCl3 + 2SC(NH2)2 + 4H2O → CuInS2 + 2CO2 + 4NH4Cl

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Other important parameters from which the spray depends on are distance between atomizer and the heat substrate as well as the exposure time will be detailed in the following section Application costs are significantly less than Epitaxial coating products and slightly higher than chemical bath deposition (CBD) ones. The mostly disadvantage of this process is the loss of non negligible precursor amount ( ≥50% ) during the deposition process and the non-uniformity of the thin film thickness. For example, Figure 1.b shows a global perturbed surface of Cu2SnS3 sprayed thin film via a 2D profilometry scan of 2x2mm2. Nevertheless, further work was reported in 2008 which refined this technique in order to overcome to this problem by using a piezoelectric transducer working at 1.72 MHz which oscillates to nebulize the starting solution [5].

III.2. Enhancement of the Spray Parameters In a common spray pyrolysis setup (Figure 2a), the main parts are: the spraying unit, the liquid feeding unit, and the temperature control one. As in common spray deposition process at ambient temperature, the precursor solution arrives at the substrate in the form of very small droplets. Droplet constituents react and form a chemical compound onto the substrate. The chemical reactants are selected such that substances other than desired compound are volatile at the temperature of deposition and before final solidification.

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(a)

(b) Figure 1.(a). The spray pyrolysis set up (b). 2D profilometry scan of a Cu2SnS3 sprayed thin film.

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(a)

(b) Figure 2. (Continued)

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tr

Tb 550

N=3

Tg (K)

N=4 500

N=5

T 450

400

Too 350

300 0

1

2

3

4

5

6

time (s)

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(c)

(d) Figure 2. (a) Schematic diagram of a common Spay apparatus. (b) Uniform layer model geometrical features. (c) Glass-deposit surface temperature rise Tg versus time. (d) Variations of relaxation time tr versus Lmean

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Karem Boubaker and Mosbah Amlouk

Uniform deposition (Figure 2b) is better ensured by moving adequately nozzle or substrate support. In most cases substrate support is fixed due to heating source size. Nozzle relative position versus targeted sample is thus held constant. Solution anion to cation ratio, spray rate, substrate temperature, ambient atmosphere and carrier gas properties are invariable and thus considered first as exogenous parameters. Deposited solution is supposed to be uniformly spread on targeted glass layer. Problem can hence be considered as onedimensional (on z-axis). Heat equation inside glass layer medium (g) and deposited layer (s) is expressed by:

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with : Tg Ts Dg Ds Pb Ps kg ks

  ²T g ( z, t ) 1 T g ( z, t ) 1   .(Pb  Ps )  Dg t kg  z ²    ²Ts ( z, t )  1 Ts ( z, t )  1 .P s  z ² Ds t ks 

: absolute temperature inside glass medium (in K) : absolute temperature inside deposited layer (in K) : glass medium thermal diffusivity (in m2.s-1) : deposited layer thermal diffusivity (in m2.s-1) : power transmitted from bulk to glass (in Wm-3) : power transmitted from glass to layer (in Wm-3) : glass medium thermal conductivity (in W.m-1.K-1) : deposited layer thermal conductivity (in W.m-1.K-1)

As ambient air mean thermal conductivity is very weak comparatively to used material, fact verified experimentally by non-relevant temperature rise beyond deposited material free surface; boundary conditions concern mainly temperature distribution continuity at median plane (z=-H) and glass-layer contact plane (z=0). According to bulk size and thermal supply, lower heat conduction toward glass layer (z=-H) can be considered as issued from an infinite source under constant temperature Tb. Heat is considered to be transmitted integrally from bulk to deposited material (PsPb). Under this assumption, a solution to heat equation inside the glass layer could be proposed through the as-defined Boubaker polynomials Bm(x) defined by:

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 B0 ( x )  1  B ( x)  x  1   B2 ( x)  x ²  2  Bm ( x)  x.Bm 1 ( x)  Bm  2 ( x)

61

for : 2  m

The solution is illustrated by a t-dependent temperature profile (Figure 2c) : A condition of convergence of the yielded result was:  2 D g . .

 2.( z  H )   .H (z  H ) 4

¨ 1

This condition couldn‟t be verified for values of height z close to –H despite of being valid for z close to zero. Mathematical thorough investigation led to a particular value of height or critical depth z0, under which model is not to be applied.

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z  z 0  H  4 2Dg . ².H

In this context, the first enhancement concerns firstly targeted glass layer dimension. In fact it was demonstrated that glass layer thickness can be lessened to the lower limit –z0, with no weighty change in thermal response and consequently in deposited material evolution. This enhancement, corroborated by experiment is important at the economic level, as long as raw material consumption would be perceptibly reduced. The second enhancement concerns the nozzle-substrate distance L . In fact, from empirical relation between h and L: b L² with : a  1.58 10 6 and b  6.06 10 8 m 3

h  a.L 

a convergence condition leads for z=0, to:  2 Dg .

 2.   2 H3

1

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Since values of  depend on h and consequently on L, for each value of , the synthetic parameter Lmean can be calculated as mean value of distances L() that satisfies convergence condition. In the same time, for each value of , the relaxation time tr as duration of surface temperature transient evolution back to bulk temperature Tb can be evaluated. Variations of tr versus Lmean are presented in (Figure 2d): Analysis of tr variations versus Lmean show that for low or high values of distance between spraying nozzle and targeted surface relaxation time falls down. This can be explained either through meaningless deposit width for correspondent values of L, or by bulk thermal inertia that causes substrate quick return to bulk temperature. Lopt is the value of L that maximizes relaxation time. Its value was (Lopt  25.3 cm).

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IV. STRUCTURAL RESULTS ON BUFFER AND ABSORBER SULFUR SPRAYED THIN FILMS IV.1. Structural Results on Absorber Sulfur Sprayed Thin Films Absorber sulfur sprayed thin films are principally ternary compounds such as CuInS2, AgInS2 and Cu(Zn)SnS as well as Zn-free ones. These, under studies over the past two decades are considered as an alternatives to old ones (CuIn(Ga)X2, with X=Se,Te,.) reported for the first time as a potential absorber material by Lofersky et al [41,42]. By 1996, thin film cells based on this material (ZnO/CuInS2) has been reported by Braunger et al to show efficiency as high as 11.4 % [43]. On the contrary, other ternary compounds, obtained by replacing indium and gallium by tin have not been yet tested in the solar conversion. In Table 2 are listed the experimental conditions of sprayed deposited sulphur ternary materials. First, we note (except for AgInS2) that the substrate temperature used for the preparation of thin films of these compounds is lower than 400°C. Second, the solution composition of metal elements in solution is upper than the solid composition in the film one i.e. in our laboratory, successful deposition of polycrystalline CuInS2 ternary compounds by the spray process was achieved thanks to a careful control of the spray solution composition via a concentration ratio [Cu+]/[In3+] equal to 1.1 [ 11, 12 ]. This phenomena related to

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some precipitation problems in the spray solutions is also observed on other ternary compounds reported previously [25,26]. Figure 3 shows the X-ray diffraction spectrum of a CuInS2 sprayed thin film which crystallizes in the chalcopyrite structure characterized by (112), (200), (220) and (116) principal orientations. As seen from this figure, the (112) reflection peak has an especially very strong intensity and the full width at half maximum of this line is of the order of 0.3° indicating a crystalline state with privileged orientation along this direction. (TS = 320 °C, Cu: In: S = 1.1: 1: 4 ) Among I-III-VI2 ternary compounds, silver indium disulphide AgInS2 based chalcopyrite semiconductor is also expected to make good solar cell absorber layer. 1000

(112) 800

I(a.u)

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600

400

(220)

200

(116)

(200) 0 10

20

30

40

50

60

70

2 (deg) Figure 3. X-ray diffraction spectrum of a sprayed CuInS2 thin film

According to AgInS2 phase transition [44], the orthorhombic form is stable at high temperature (>620°C) and the chalcopyrite one is rather stable at low temperature (