CuInSe2-based photovoltaic devices. This has led to extensive research on the effects of this element on the polycrystalline material, with the ultimate objective of identifying the mechanism by which Na acts on the cells. However, much less research has been done on the effects of sodium on the monocrystalline form of this material. Such research could help to differentiate bulk from grain-boundary effects, as well as to identify reactions between the Na and the compound itself, or the individual elements within the compound. Therefore, in the present work, Na was added in varying quantities to quartz ampoules containing Cu, In and Se, in the atomic ratios of 1:1:2. The ampoules were evacuated and sealed before being put through a vertical-Bridgman procedure, resulting in ingots containing large, cm-size crystals. Electrical measurements on the ingot material revealed p-type conductivity for all material grown with stoichiometric proportions of the starting elements, without Na, but n-type conductivity for material grown with Na above a certain critical value. It was discovered that this critical value of Na increased when excess

Se, above stoichiometry, was also included in the ampoules. Further experiments confirmed the mechanism responsible for the conductivity type change to be a reaction between the Na and Se, in a 2:1 atomic ratio, corresponding to the chemical formula Na2Se, which starved the CuInSe2 of its share of selenium, rendering it Se- deficient and therefore n-type. Other effects of Na on the material are identified, including no detection of sodium within the ternary itself. As well, some photovoltaic cells were made, the best of which achieved an efficiency of 8.8 %.

i RÉSUMÉ

Le sodium améliore la performance des appareils photovoltaïques polycristallins à base de CuInSe2. Ce fait a déclenché beaucoup de recherches sur les effets de cet

élément sur les matériaux polycristallins ayant comme objectif primaire d’identifier le mécanisme par lequel le Na agit sur les cellules. Par contre, très peu de recherches

étudient les effets du Na sur la forme monocrytalline de ce matériel. De telles recherches pourraient aider à différencier les effets à l’intérieur du cristal des effets aux surfaces des cristaux, en plus d’identifier les réactions entre le sodium et la composition même, ou ses éléments individuels. Ainsi, l’expérimentation ci-présentée consiste à ajouter du Na en quantités variées à des ampoules en quartz contenant Cu,

In et Se en respectant le ratio atomique de 1 :1 :2. Elles ont été évacuées et scellées avant d’être introduites selon la procédure « vertical-Bridgman ». Ce qui en résulte est un lingot contenant un large cristal mesurant quelques cm. Des analyses

électriques sur le lingot ont révélé de la conductivité de type p sur tout le matériau créé avec des proportions stœchiométrique de l’élément originaire, sans Na, mais avec de la conductivité de type n sur tout matériau créé avec du Na au delà d’une certaine valeur. Il a été découvert que cette valeur critique de Na augmente lorsqu’un excédant de Se, au-delà de la stœchiométrie, est aussi incluse dans les ampoules. Des expérimentations ultérieures ont confirmé que le mécanisme responsable pour le changement de type de conductivité est une réaction entre Na et Se selon un ratio atomique de 2 :1, correspondant à la formule chimique Na2Se. Cela rend le CuInSe2 déficient en Se et donc de type n. D’autres effets du Na sur le matériau ont été identifié, incluant la non-détection de sodium dans le ternaire. En plus, certaines

ii cellules photovoltaïques ont été créées, la meilleure ayant atteint une efficacité de

8.8%.

iii ACKNOWLEDGEMENTS

The author wishes to express his gratitude to his supervisors, Prof. C.H.

Champness and Prof. I. Shih, for their guidance over the course of this work.

The author also wishes to acknowledge Prof. Mark Sutton, from the McGill

Department of Physics, for help in the analysis of the X-ray diffraction results. As well, the author would like to thank Don Pavlasek, Bob Thomson, from the McGill ECE

Design and Machining facility, and Monique Riendeau, from the McGill Mining and

Material Engineering Analytical laboratories, for their time and effort spent in the advancement of this work.

Further thanks are extended to graduate students Yue Hua Tan, Sriraman

Rajagopalan and Kiran Kumar, from the Department of Mining and Materials

Engineering, Jeanne-Louise Shih, Sunyoung Park, Blaise Frisson and Hieu Nguyen, from the Department of Electrical and Computer Engineering, and former students

Jack Wu, Han-Jen Yang and Yi Chen, as well as all of the members of the Nano

Electronic Devices and Materials laboratory.

In addition, the author would like to acknowledge the sources of financial support provided to him during the course of this work: the Department of Electrical and

Computer Engineering of McGill University, the Natural Science and Engineering

Research Council of Canada (NSERC), the Fond québécois de la recherche sur la nature et les technologies (FQRNT), and the Institute of Electrical and Electronics Engineers

(IEEE).

Finally, the author would like to thank his wife, Lisanne, for her patience and support throughout this work, and his daughter Eleanor, as well as his parents, his family and all of his friends. They are all a continuous source of inspiration and encouragement.

iv TABLE OF CONTENTS

ABSTRACT i RÉSUMÉ ii ACKNOWLEDGEMENTS iv TABLE OF CONTENTS v

1 INTRODUCTION 1

2 HISTORICAL REVIEW 8

2.1 INTRODUCTION 8 2.2 SUMMARY OF THE USE OF CuInSe2 AND Cu(In,Ga)Se2 IN PHOTOVOLTAICS 9 2.3 SUMMARY OF Na IN CIS AND CIGS 12 (a) The effects of Na on material characteristics and device performance 12 (b) Sodium incorporation mechanisms 15 2.4 RECENT ADVANCES IN CIGS PHOTOVOLTAICS 19

3 TRANSPORT THEORY 23

3.1 INTRODUCTION 23 3.2 EXPLANATION OF SEEBECK AND HALL EFFECTS 24 (a) Seebeck effect 24 (b) Hall effect 26 3.3 TRANSPORT FORMULAE 28 (a) Single carrier model 28 (b) Two carrier model 29 3.4 SIGN REVERSAL OF α AND RH 32 (a) Intrinsic temperature range inversion 32 (b) Inversion by doping at fixed temperature 34

4 EXPERIMENTAL PROCEDURE 40

4.1 INTRODUCTION 40 4.2 INGOT GROWTH PROCEDURE 42 (a) Ampoule and charge preparation 42 (b) Bridgman-growth 44 4.3 ANALYSIS EQUIPMENT 45 (a) Hot probe 45 (b) Sample holder for thermoelectric power measurements 46 (c) Sample holder for electrical resistivity measurements 49 (d) Hall coefficient measurements 50 (e) X-ray equipment 51 (f) SEM-EDX 52

v 4.4 ADDITIONAL MEASUREMENTS 52 (a) Heating experiments 52 (b) Cold powder mixing experiments 55 (c) Etching experiments 57 4.5 ADDITIONAL REMARKS ON GROWTH EXPERIENCE 58

5 EXPERIMENTAL RESULTS 67

5.1 INTRODUCTION 67 5.2 ELECTRICAL MEASUREMENTS 69 (a) Hot probe measurements on wafers 70 (b) Thermoelectric power measurements on filamentary samples 72 (c) Hall effect measurements on filamentary samples 74 (d) Electrical resistivity measurements on filamentary samples 75 (e) Carrier concentration and mobility 75 5.3 STRUCTURAL AND COMPOSITIONAL INVESTIGATIONS OF CuInSe2 AFTER GROWTH 77 (a) X-ray diffraction on bulk material 78 (b) SEM-EDX 79 5.4 CHARACTERIZATION OF DEPOSITS 81 (a) Red deposit 82 (b) White deposit 83 (c) Copper nodules 84 5.5 RESULTS OF EXTRA EXPERIMENTS 85 (a) Heating experiments 85 (b) Mixing experiments 87 (c) Etching experiments 88 5.6 DISCUSSION 90

6 PHOTOVOLTAIC CELLS MADE FROM MONO-CIS 111

6.1 INTRODUCTION 111 6.2 CELL FABRICATION PROCEDURE 112 (a) Wafer cutting, polishing, cleaning and surface etching 113 (b) Wafer annealing 115 (c) Cadmium sulfide layer formation by chemical bath deposition 116 (d) Back contact deposition 117 (e) ZnO deposition by sputtering 119 (f) Indium top contact deposition and cell mounting 120 (g) Testing 120 6.3 RESULTS ON CELLS 122 6.4 DIFFUSION LENGTH ESTIMATE BY PHOTOCURRENT-CAPACITANCE MEASUREMENTS 123

vi 6.5 DISCUSSION 124

7 DISCUSSION AND CONCLUSIONS 134

7.1 INTRODUCTION 134 7.2 MECHANISM OF SODIUM ACTION IN BRIDGMAN- GROWN CuInSe2 134 7.3 ORIGINAL CONTRIBUTIONS TO KNOWLEDGE 143

REFERENCES 147

APPENDIX A INGOT GROWTH CONDITIONS 157

APPENDIX B RESULTS OF ASSAY ON 164 ELEMENTAL SODIUM

APPENDIX C RESULTS OF ELECTRICAL AND 166 EDX ANALYSIS OF INGOTS

APPENDIX D PHOTOVOLTAIC CELL 169 FABRICATION PROCEDURES

vii Chapter 1

Introduction

Copper indium diselenide (CuInSe2) is a compound semiconductor with a bandgap of 1.04 eV and an optical absorption coefficient of the order of 105 cm-1 over the visible spectrum, the primary application of which is as a photovoltaic absorber.

Of the various photovoltaic absorber materials used currently, CuInSe2 is of a type used in “thin-film” device configurations, meaning devices that are made very thin, with absorber layers of the order of a few microns thick but which still absorb a majority of incoming solar radiation, a result of the high absorption coefficient. In this capacity, CuInSe2 is almost exclusively used in polycrystalline form, and alloyed with gallium, forming the quaternary Cu(In1-yGay)Se2 (sometimes written as CIGS);

Ga is used to fine-tune the bandgap of the material and thereby improve the energy collection efficiency of the layer. The highest performing cell to date which uses

CIGS as the absorber layer has an energy conversion efficiency of 20.3 % [1.1]. In comparison, the record efficiency for a single-junction crystalline silicon device is

27.6 % (obtained with a 92x solar concentrator) [1.2]. There may be several reasons for this disparity, including an unbalanced understanding of silicon compared to

CIGS (the author speculates that humans know more about silicon than about any other substance in the universe), as well as the respective material processing and fabrication technologies.

The current work aims to address the former of these by contributing to the scientific community’s fundamental understanding of the underlying mechanisms governing the material properties and characteristics of CuInSe2. Included in this understanding is the beneficial role of sodium on the performance of CIGS photovoltaic cells. It is well established that Na acts to improve the open-circuit voltage and fill factor of CIGS devices, and thereby increase their energy conversion efficiency [1.3-1.5]. However, aspects of the mechanism behind this enhancement are unclear, including whether the sodium somehow acts on the material itself, or if the effects are merely due to changes at the grain-boundaries.

Therefore, the primary motivation of this work is the elucidation of the role of sodium on the performance of thin-film CuInSe2-based photovoltaic cells. However, rather than studying polycrystalline thin-films directly, this work seeks to differentiate between the bulk and grain-boundary effects of Na by investigating the effects of this element on monocrystalline CuInSe2, grown by a vertical-Bridgman method. This method allows for the starting compositions of growth melts to be easily controlled, and results in ingots with cm-sized single crystals. The ability to take measurements within individual grains is therefore greatly facilitated as compared to thin-films, where grains are only one or two microns in diameter. Note that in this work, the gallium has been omitted in order to avoid questions of Ga distribution associated with the quaternary, thereby eliminating a variable, as Na has

2 been shown to act to improve the performance of devices made from CuInSe2 as well as CIGS.

Pioneering work on monocrystalline CuInSe2 has been carried out for more than two decades in this laboratory [1.6-1.12]. Some of the earlier work focused on improvements in the growth process, with the aim of obtaining a solid, void-free ingot [1.8]. Once this was achieved, principally by using a boron nitride coated ampoule [1.13-1.14], work progressed on the effects of variations of stoichiometry.

Research here has investigated the fabrication of ingots grown from stoichiometric proportions of copper, indium and selenium (CuInSe2), as well as deviations from stoichiometry involving the copper (CuyInSe2, where y = 1.1, 1.2 or 1.3) or the selenium (CuInSe2+x, where x = ± 0.2 [1.15] and x = -0.3, 0.08 and 0.17 [1.16]).

Work has been done on the effects of including gallium in the melts of Bridgman- grown material (CuIn1-zGazSe2) [1.17, 1.18], and on the ordered-defect compound

CuIn3Se5 [1.19]. This latter work, performed by H.P. Wang, also involved the effects of sodium on the material, including, for the first time, the addition of sodium to the melts of CuInSe2, before compound synthesis, in the concentrations of 0.05, 0.25 and

2 at. %.

The present work expands on the work of Wang [1.19] by including sodium in the melts of CuInSe2 in varying concentrations, ranging from 0 to 3 at. %.

Furthermore, the works of Du [1.15] and Wang [1.19] have been combined so as to explore the effects of an excess of Se (i.e. CuInSe2+x, where x = 0.005, 0.02, 0.05, 0.1,

3 0.2, 0.3, or 0.4) on the material grown in the presence of Na up to 11 at. %.

Additionally, the effects of a deficiency of selenium in the growth melt on the resulting grown material, first documented by Du [1.15], were confirmed and expanded here by fabricating an ingot from a starting composition of CuInSe1.7, in order to draw comparisons between this sample and the samples grown with sodium.

In this work, all samples were grown by a vertical, one-ampoule Bridgman method

[1.20-1.21], whereby the charge is loaded into cleaned ampoules of fused-quartz, which are evacuated and sealed. In this case, the charge consisted of weighed quantities of copper, indium and selenium pellets, and sodium in either elemental form or as Na2Se. The ampoules remained sealed during the entire heating and crystal growth process, and were broken open to retrieve the ingots only after growth was completed and the system had been cooled to room temperature

Investigation of the effects of adding Na to the growth melts was done by first visually examining the ampoules upon their removal from the furnace. It was found that the physical appearance of the ampoules, and specifically, the presence and quantity of deposits found within the ampoules after growth, was affected by the composition of the melts. The ampoules were then broken open in order to obtain the ingots themselves. The electrical characteristics of the ingots were investigated by first determining the conductivity type of the bulk material, done by hot-probe, and then by measuring the thermoelectric power, Hall coefficient and electrical resistivity.

This was more difficult for some samples than for others as the brittleness and cohesiveness of the ingots, as will be explained, was also dependent on the melt

4 composition. Structural and compositional measurements were also made on the ingots, and deposits, by X-ray diffraction and scanning electron microscopy – energy- dispersive x-rays. Some additional experiments were also performed on pre- fabricated material (grown by the author over the course of this work), and used in conjunction with the results of electrical and structural measurements in the development of a model used to describe the action of the Na in monocrystalline,

Bridgman-grown CuInSe2. This, along with a summary of the entire work, is presented in Chapter 7.

The contents of each chapter of the thesis are now described. Chapter 2 contains a review of the use of CuInSe2 in photovoltaic cells and devices. The effects of Na on the material, and on devices using the material, as discovered elsewhere, are also given in this chapter, as well as an explanation of the various Na incorporation mechanisms developed over time. The chapter then concludes with a brief summary of some of the more recent advancements made with regards to CuInSe2-based photovoltaics.

In Chapter 3, the equations governing the electrical characteristics measured on the material in this work, being the thermoelectric power and Hall coefficient, are described. The chapter further describes how these characteristics are affected by changes in the number of carriers in the material, brought upon either directly, by doping, or through changes in temperature. These are used to understand the

5 electrical characteristic changes seen in the material grown in this work as the composition of the melts are varied.

Chapter 4 contains the descriptions of the experimental procedures used in the investigation of the effects of Na on the characteristics of the bulk material. Included in this chapter is an explanation of the growth procedure, the electrical and structural measurements, and other experiments aimed at enhancing the understanding of the effects of Na on CuInSe2. The results relating to the work described in Chapter 4 are then presented in Chapter 5. As will be seen, a unique relationship was discovered between the critical amount of sodium and the excess amount of selenium added to the melt.

As well, some studies have been made on photovoltaic devices using

Bridgman-grown CuInSe2 as the absorber layer. Solar cells were fabricated in this work from wafers cut from the ingots obtained following the growth procedure described in Chapter 4. These devices were then tested and, in one case, used to estimate the diffusion length of the CuInSe2 absorber. The experimental procedures, results and discussions concerning specifically these studies are confined to Chapter

Finally, the entire work is summarized in Chapter 7, and a model that aims to explain the mechanisms by which Na acts on the material is given. The thesis then

6 concludes with a list of the specific original contributions to science and advancements of knowledge made by this research.

7 Chapter 2

Historical Review

2.1 INTRODUCTION

An extensive review of CuInSe2 (CIS) and Cu(In,Ga)Se2 (CIGS) has been written by the author, and can be found in the precursor to this current work [2.1].

Included in that review was a history of the use of CIS and CIGS in photovoltaic devices, leading to the discovery of the beneficial effect sodium has on device performance. This discovery led several groups to begin investigating the effects of

Na on the material, and finally to the development of models aimed at explaining the action of the Na and the cause of these effects.

This review will be summarized in this chapter, with an emphasis on the effects of Na on the characteristics of the material, which is updated from the work mentioned above. Following this will be an explanation of the various mechanisms used to incorporate Na into the material. The chapter will then be concluded with a brief look at some of the most recent advancements made in CIGS photovoltaic technology.

8 2.2 SUMMARY OF THE USE OF CuInSe2 AND Cu(In,Ga)Se2 IN

PHOTOVOLTAICS

While the ternary compound was first synthesized by Hahn, Frank, Klinger,

Meyer and Störger in 1953 [2.2], the first definitive copper indium diselenide solar cell was constructed at Bell Telephone Laboratories by Wagner, Shay, Migliorato and

Kasper in 1974 [2.3]. Fabricated using melt-grown monocrystalline material, it had an energy conversion efficiency of 5 %. One year later, in 1975, they had managed to increase the efficiency to 12 % by forming the junction on the (112) plane, and with the use of an anti-reflection coating [2.4]. The active area of that cell was quite small by 2011 standards, 0.79 mm2, due to the presence of microcracks in the CIS wafer.

At this point, interest had shifted towards CuInSe2 as a polycrystalline material, due to the cost advantages over monocrystalline wafers, as well as the ability to make films very thin and even potentially flexible (this was finally accomplished in 1995 by J.M. Marino et al [2.5] using a Kapton substrate specifically for space applications). Nevertheless, work continued to be carried out on improving the performance of single-crystal devices. Kazmerski and Sheldon [2.6] experimented with Bridgman-grown CuInSe2 and indium-tin-oxide (ITO), rather than CdS, achieving an efficiency of 8.5 % in 1978. Grindle et at [2.7], in 1980, fabricated a device using CuInSe2 deposited by molecular beam epitaxy, with a CdS window layer, achieving an efficiency of 4.7 %. A short time later, in 1983, Arya et al [2.8] reported an efficiency of 5.9 % for a 4.5 mm2 device made using a multicrystalline

9 wafer of CuInSe2, also with CdS. The efficiency was later increased to 10.6 % by reducing the cell area to 0.8 mm2.

More than ten years after the result obtained by Arya et al, work on monocrystalline material was carried out by Shukri [2.9], using a CdO window layer, rather than the traditional CdS. Despite the considerable lattice mismatch between the CIS and CdO, Shukri nonetheless managed to obtain an efficiency of 6 %. It was found in that work that annealing the wafers prior to the CdO deposition improved the performance of the cells. At around the same time, Yip [2.10] achieved an efficiency of 11 % on a device with an active area of 8 mm2 made from a monocrystalline,

Bridgman-grown wafer using CdS-ZnO window structure. This result was later improved upon by Du et al [2.11] for a device with a layer structure of Au-CuInSe2-

CdS-ZnO-In. The effective area conversion efficiency for this device was 12.5 %, as confirmed by NREL, without an anti-reflection coating. This achievement is even more impressive considering that the active area of the cell was 12.6 mm2 (total area

13.8 mm2).

Meanwhile, while this work was taking place, other work was being carried out by L. Kazmerski on polycrystalline films made by co-evaporation of powdered

CuInSe2 with selenium. Polycrystalline CIS cells were constructed by depositing this layer on Au-covered glass, followed by a layer of CdS, reaching an efficiency of 5.7

% in 1976 [2.12] and 6.6 % in 1978 [2.13]. Two years later at Boeing Aerospace

Company, Michelsen and Chen had started experimenting with simultaneous

10 elemental evaporation for the deposition of polycrystalline CuInSe2, achieving 5.7 %

[2.14]. This cell comprised of a molybdenum layer sputtered onto an alumina substrate, followed by a gold layer, and then two layers of CuInSe2 deposited at different substrate temperatures. The CdS layer was also deposited in a two step process, the second of which was a co-evaporation with indium to reduce the series resistance, with an aluminum grid acting as the top contact. By 1984, the gold layer was completely omitted, leaving only molybdenum as the back contact, and zinc was alloyed into the CdS in order to increase the open-circuit voltage of the devices

[2.15], resulting in an efficiency of 11 % [2.16]. An anti-reflection coating, grown by heating in oxygen after fabrication, contributed to the high performance of the cell.

Refinements of this anti-reflection layer pushed the efficiency up slightly to 12.4 %

[2.17].

In time, other researchers discovered the beneficial effects of using gallium to increase the bandgap of the material in 1990 [2.18], forming the now-standard quaternary CuIn1-yGaySe2, also written as Cu(In,Ga)Se2 or simply CIGS, as well as the addition of highly resistive layer of ZnO immediately above the CdS, and a layer of ZnO doped with aluminum above that to reduce the resistivity near the top contact in 1993 [2.19]. However, the development in the fabrication process of CuInSe2 solar cells that is the most relevant to this work is the change of substrate material from alumina to soda-lime glass in 1989 [2.20].

11 2.3 SUMMARY OF Na IN CIS AND CIGS

Soda-lime glass (SLG) was initially considered as a substrate for CuInSe2 due to its lower cost over other substrates, such as borosilicate glass and alumina [2.21].

However, apart from being merely a suitable substitute, devices fabricated on SLG consistently outperformed those of other substrates [2.20, 2.22], demonstrating higher open-circuit voltages and higher fill factors. Initially, it was thought that the performance enhancement was due to a better matched thermal expansion coefficient of the SLG to CuInSe2 than other substrates. However, in 1994, Bodegård et al [2.23] and Holz et al [2.21] sought to isolate the effect of the sodium in the SLG on the material and on device performance. Bodegård’s team did this by depositing CIS on various substrates not containing sodium, as well as SLG with and without a diffusion barrier. The material grown on SLG without the diffusion barrier exhibited an increase in (112) orientation, and devices fabricated on that substrate performed better than those fabricated on the Na-free substrates and the SLG with the diffusion barrier.

Holz’s team investigated the electrical effects of Na on the material by depositing

CuInSe2 on Na-free and Na-implanted sapphire, observing an increase in conductivity with Na, which they attributed to an increase in p-type doping. Thus, Na was shown to do more to CuInSe2 than simply affect the structure and orientation of the films.

2.3 (a) The effects of Na on material characteristics and device performance

As mentioned above, sodium has been observed to cause an increase in the

(112) preferred orientation in thin-film CIS and CIGS [2.22-2.26], as well as to result

12 in an increase in average grain size [2.22, 2.25-2.28]. Sodium also acts to change the electrical properties of the films. Observers have noted a decrease in the resistivity of the films with Na [2.21, 2.25, 2.28], and an increase in acceptor concentration [2.29-

2.33] by capacitance measurements. Despite these changes, Na is generally accepted to collect at defects and grain boundaries, and not to enter intact surfaces, as observed by Lyahovitskaya et al [2.34], although more recently, Lei et al [2.35] found no evidence of oxidation or Na buildup at grain boundaries using EDS. As well,

Cojocaru-Mirédin et al [2.36] found dilute concentrations of Na (~0.002 at. %) distributed homogenously within the grains of CuInSe2 deposited on soda-lime glass substrates using atom-probe tomography.

Shown in Figure 2.1 [2.37] is a ternary phase diagram of the Cu-In-Se system, with dots indicating stable phases of the compounds, including those found on the tie- line connecting the binaries Cu2Se and In2Se3. CuInSe2 can be found on this line, along with various other phases known as ordered-defect compounds (ODC), which include CuIn3Se5, CuIn4Se7 and CuIn5Se8; note that the ternary α-phase containing excess copper is not permitted, according to the phase diagram shown in Figure 2.2

[2.38]. These compounds are made of different combinations of an acceptor-like pair

- 2+ of copper vacancies, 2VCu , coupled with an indium on copper site, InCu , the latter of which acts as a double donor [2.39]. The effect of Na on the formation of these compounds is not entirely clear, as this element has been reported to both suppress

[2.40] and promote [2.41] the ODC phases. With respect to the effect of sodium on the performance of the photovoltaic devices themselves, the Na is found to increase

13 the open-circuit voltage (VOC) and fill-factor (FF) [2.23, 2.27, 2.30], thereby increasing the energy conversion efficiency.

Several models describing the action of the Na in the material have been put forward in order to explain the abovementioned observations. Included in these is the

“oxygen” model [2.42-2.43], used to explain the perceived increase in acceptor

+ concentration as a decrease in the compensating donor VSe . In this model, Na acts to cause O2 molecules, present during a post-fabrication anneal in air, to dissociate into atomic O, which then bonds with In at the surface of the material, thereby settling at a location of an Se vacancy and passivating it.

The InCu model [2.44] also describes the action of the Na as leading to a decrease of compensating donors, with the donor being InCu. It is suggested that Na itself can occupy a copper site in the CuInSe2 lattice, and in doing so, prevents an indium atom from doing the same. The ability of Na occupying a place within the intact CuInSe2 crystal at a Cu site, resulting in the formation of a NaInSe2 phase, has been suggested by Granata et al [2.45] and theoretically evaluated by Wei et al [2.46].

The NaIn model [2.47] also describes the Na occupying a place within the

CuInSe2 lattice, although that place is an indium site (or a gallium site, in the case of

CIGS), rather than a copper site. In this way, the Na actively dopes the material through the creation of these acceptor-like defects. In examining material deposited on soda-lime glass using X-ray photoelectron spectroscopy (XPS), Na-Se bonds were

14 attributed to this defect, rather than to secondary Na-Se phases. The ability of Na to occupy an In site in the CuInSe2 lattice has also been theoretically proposed by Wei et al [2.46].

2.3 (b) Sodium incorporation mechanisms

As described above, the first method of incorporating sodium in the absorber layer was through the use of soda-lime glass as the substrate material. Over the course of the fabrication of the devices, the Na could diffuse from the glass through the molybdenum back contact layer, and into the CIS or CIGS. Over time, experimenters devised ways to increase the diffusion of the Na, for example, by increasing the substrate temperature during the deposition of the absorber. Shafarman et al [2.48], Wang et al [2.49], and Zhang et al [2.50] found the performance of CIGS devices made on soda-lime glass to increase with deposition temperature up to 550

°C, the maximum used in that work, despite a decrease of JSC occurring after 500 °C in the case of Wang et al [2.49].

Recently, an objective of CIGS module manufacturers is the development of high-performance devices on alternative substrates, such as foils or other flexible materials not containing Na, as well as to reduce deposition temperatures to allow for the use of a wider range of substrates. This has naturally led to efforts to develop alternative Na-incorporation methods. One such method, demonstrated by Ishizuka et al [2.51-2.52], involves depositing an “alkali-silicate glass thin layers” (ASTL) as a precursor to the Mo. In this method, a soda-lime glass thin film (SLGTF) layer in

15 sputtered onto titanium substrates, followed by a layer of molybdenum, and then

CIGS. The Na density in the films, measured using SIMS, could be controlled by increasing the thickness of the layers. In films between 50 and 460 nm, the highest performance was seen in a device made with a 120 nm SLGTF, which had an efficiency of 17.4%. Another cell made using SLGTF on a zirconia substrate achieved an efficiency of 17.7%, about 1% less than reference cells made on SLG.

Interestingly, they also observed a decrease in grain size with SLGTF compared to uncoated Ti, but an increase in efficiency, suggesting grain size is not a limiting criterion of high performance.

Aside from being contained in the substrate and diffused through the molybdenum back-contact, Na can be incorporated into the CIS or CIGS absorber layer by: i) depositing a precursor layer, prior to the deposition of the absorber layer; ii) including it during the growth of the absorber by co-evaporating this element, or a compound containing this element, together with the Cu, In, Ga and Se; iii) depositing a successor film containing Na, after the growth of the absorber; or, very recently, iv), doping the molybdenum sputtering target, used in the deposition of the back-contact, in order to create an intrinsic Mo:Na layer. These four methods are now discussed.

In the precursor method, a compound containing Na is evaporated on to the surface of the molybdenum back contact. Possible precursors include Na2S and

Na2Se, although NaF is more commonly used due to its stability in air and ability to

16 evaporate stoichiometrically [2.53]. In studies using NaF as a precursor, an increase in VOC and FF was observed [2.53-2.55]. However, it was also found that increasing the thickness of the NaF caused a decrease in JSC, although Zachmann et al [2.55] found this was only true with NaF as a co-evaporant, not as a precursor. Granath et al

[2.53] saw an improvement in performance with a small NaF layer, 30 Å thick, but noticed no change in the crystal structure, suggesting the change to be in the electrical properties. Caballero et al [2.54] found higher carrier concentrations with increased

NaF by drive-level capacitance profiling, which they ascribed to a reduction of compensating donors. In work on the sulfide material Cu(In,Ga)S2, Kaul et al [2.56] found an NaF precursor to result in increased grain size and preferred (112) orientation.

Sodium can also be incorporated in the material during growth, by co- evaporating a Na compound with the other absorber elements. Rudmann et al [2.57] found that, in films grown with NaF, the sodium contributed to a decreased grain size, as well as affected the [Ga]/[In] concentration ratios, but also an increased cell performance in devices made from the films. Lammar et al [2.58], in experimenting with co-evaporation as a means of compensation for lower fabrication temperatures, found the grain size to decrease with deposition temperature for films on SLG, but that the size did not increase with Na co-evaporation. They also noticed a drop in

VOC, JSC and FF in devices deposited at 420 °C from those deposited at 550 °C.

While VOC and FF were seen to increase in the low-temperature devices with Na co- evaporation to values comparable to those obtained from high-deposition

17 temperatures, JSC was not increased, ultimately resulting in lower conversion efficiencies. More recently, Shin et al [2.59], experimenting with Na2S as a co- evaporant, also found devices made on Na-free Corning glass performed worse than conventional CIGS cells on SLG.

Post-growth Na incorporation is similar in experimental procedure to the Na precursor film treatment, except that the Na-containing film is deposited directly onto the CIGS after growth, rather than the Mo before growth. As such, the effect of the

Na, when included post-absorber deposition, is predominately in the electrical properties of the material, rather than structural properties, as investigated by

Rudmann et al [2.60-2.61]. In depositing NaF onto grown Na-free absorbers, they found that devices made at lower temperatures on SLG performed worse than those with a post-deposition Na-treatment, but the opposite was true for devices made at temperatures exceeding 500 °C.

Finally, it was found very recently by Yun et al [2.62-2.63] that benefits in device performance associated with sodium could be achieved by incorporating the

Na directly into the molybdenum back contact through doping of the Mo-sputtering targets. Depositing CIGS on Na doped Mo layers, onto an alumina substrate, they observed increases in the (112) growth plane, as well as VOC, FF and efficiency.

These results were also found by R. Wuerz et al [2.64], and have led to the commercialization of this technology [2.65].

18 2.4 RECENT ADVANCES IN CIGS PHOTOVOLTAICS

The ratio [Ga]/[Ga+In] in the material is often used to manipulate the bandgap and, with that, improve the collection efficiency of the devices. For example, at the

2011 EU PVSEC in Hamburg, A. Chirila et al [2.66] reported that the performance of

CIGS devices on Ti foils were dramatically improved from 14% to 18% by varying the elemental co-evaporation sequence. In doing this, they were able to decrease the barrier at the CdS junction, improve the grading in the middle part of the device and increase the back-contact barrier. Such careful management of the gallium content is also largely responsible for the improvements in VOC and FF, and efficiency, seen in the record-breaking cells produced at NREL [2.67] and ZSW [2.68].

The ability of other elements to alter the bandgap of the material and, with that, improve VOC of devices, has also been explored. The choice of elements to occupy the Cu, In and Se sites within the lattice is made by choosing those in the same column of the periodic table. Researchers have also experimented with sulfur

[2.69] and tellurium [2.70] as alternates to Se, silver to the copper [2.71], and aluminum [2.72-2.73] to the indium and gallium. With these advancements, the material has progressed from the ternary of CuInSe2, to the more complicated

(Cu,Ag)(In,Ga,Al)(Se,S,Te)2 alloy. Copper-zinc-tin-sulfide (written as CZTS) is a derivative of CIS for which much interest has recently been generated. The advantage of this material over CIS is primarily in its exclusion of indium, the price of which has increased dramatically as a result of its scarcity and a growing demand

19 for ITO (indium-tin-oxide), a critical material in flat-panel displays. Contrarily, tin is much more abundant on Earth’s crust than indium, as is sulfur than selenium [2.74].

The record conversion efficiency for a CZTS cell is 10.1%, made by Barkhouse,

Gunawan, Gokmen, Todorov, and Mitzi [2.75], using a solution-deposited absorber of the form Cu2ZnSn(S,Se)4.

The traditional method of deposition of CIGS films is vacuum deposition, by co-evaporation of the Cu, In, and Ga, followed by an anneal in Se. The high manufacturing costs of this process has, in part, motivated the search for alternative deposition methods. It is also preferable to avoid the scale-up and environmental issues associated with vacuum processes. As well, the use of new alternative substrates, such as polyamide, steel and titanium foil, can place demands on deposition and fabrication temperatures, requiring that new processes be developed to accommodate these demands. One new process is paste coating [2.76], in which a precursor solution of the absorber elements in ethanol is prepared and applied to glass substrates, followed by an anneal. In the Field-Assisted Simultaneous Synthesis and

Transfer (FASST) method, developed at Heliovolt [2.77], a Cu-Se precursor ink is first synthesized, and then processed with an In-Ga film, deposited by physical vapour deposition, into CIGS. Ink-based deposition has also been experimented with as a viable method for use in conjunction with alternative, flexible substrates [2.78].

More recently, the use of spray deposition of the precursor solutions [2.79-2.80] has been investigated as further attempts to decrease the fabrication time, and increase throughput, are made.

With regard to record efficiencies, as mentioned above, the highest efficiency

CIGS solar cell was made by Jackson et al [2.68] at ZSW, Stuttgart. The efficiency of that cell was 20.3%, on an area of 0.5015 cm2, as confirmed by Fraunhofer ISE. The record module efficiency for a CIGS device is 15.7%, by MiaSole, on an area of 9703 cm2, as reported by Green et al [2.81].

Figure 2.1 Ternary phase diagram of Cu-In-Se, indicating CuInSe2 and ODC phases, including CuIn3Se5, CuIn4Se7 and CuIn5Se8, along the tie-line between Cu2Se and In2Se3 [2.37].

Figure 2.2 Psudo-binary Cu2Se-In2Se3 phase diagram of CIS phases with temperature. Note that chalcopyrite CuInSe2 (written as α-CuInSe2 in the diagram) cannot have more than 50 % of Cu2Se at room temperature [2.38].

22 Chapter 3

Transport Theory

3.1 INTRODUCTION

In this work, the effects of sodium on the transport properties of Bridgman- grown ingots of CuInSe2 were investigated by measuring the thermoelectric power

(α) and Hall coefficient (RH) on filamentary samples cut from these ingots. It was by these measurements that the change in conductivity type from p to n in stoichiometric material grown with a sufficient amount of Na was discovered, as well as the relation between the critical amount of Na required for the type change to occur and x, the amount of excess Se also present in the ampoule. This relation, the evidence for which will be given in Chapter 5, and its significance, discussed in Chapter 7, points to the creation of Se vacancies at the rate of 2:1 Na atoms added to the ampoules to

Se atoms removed from the material, and is a central discovery of this work.

Accordingly, in the present chapter, the basic theory and equations of these two quantities (α and RH) will be presented.

This chapter will first consider α and RH in single-carrier material. Following this, the equations governing α and RH in a two-carrier material will be discussed.

This discussion will include a description of the simplifying assumptions made in the derivation of the equations, and how these apply to the actual material in this work.

23 As will be shown, the single-carrier transport model of α and RH can reasonably be applied to the model, and the effect of the second carrier can be neglected.

The next section, Section 3.3, will explore how α and RH are affected by a conductivity sign change, first in near-intrinsic material, and then by doping in extrinsic material. Considering that the energy gap of CuInSe2 is in the range of 1 eV, the latter scenario is a more accurate perception of the electrical changes undergone by the ingots grown with sodium in this work. This material in the chapter is based on the standard textbooks on electron transport properties such as

“Thermoelectricity: Science and Engineering” by R.R. Heikes and R.W. Ure Jr., 1960

[3.1], and “The Hall Effect and Related Phenomena” by E.H. Putley, 1960 [3.2].

3.2 EXPLANATION OF SEEBECK AND HALL EFFECTS

3.2 (a) Seebeck effect

The thermoelectric power effect, also known as the Seebeck effect, derives from the diffusion of carriers of different energies when a temperature gradient is imposed on a material. While a detailed derivation of the equations which govern this effect is beyond the scope of this work, considering that it is the principal phenomenon by which the conclusions in this thesis are drawn, a qualitative explanation of the forces influencing this effect will be given here.

24 Considering an n-type, non-degenerate extrinsic semiconductor, the carrier concentration of electrons of energy E is expressed as:

m* 2m* (E  E )   E  E  n(E)  n n C exp  F  , (3.1) 2 3        kT 

where the term on the right consists of the density of electron states multiplied by the probability of filling at energy E (see Figure 3.1). In examining this equation, it can be understood that the Fermi distribution of carriers is broadened with increases of temperature; that is, more high-energy electrons are found at areas of the material at higher temperatures than at lower temperatures.

Therefore, a temperature gradient placed across a material will result in a non- uniform density of high and low-energy carriers, which will then diffuse in opposite directions; the high-energy carriers will diffuse in opposite direction of the temperature gradient, towards the cooler area of the material. However, the high- energy carriers diffuse faster than the low-energy carriers. In an open circuit, this will result in a buildup of electrons in the low-temperature zone, producing an electric field which opposes the diffusion. When the material reaches steady state, the diffusion current and the opposing drift current are completely balanced, resulting in no net current flow, but an electric field existing within the material. It is the generation of this electric field which is known as the Seebeck effect. The thermoelectric power, which is a characteristic property of the material, is calculated

25 as the electric potential difference ΔV across two points with a temperature difference of ΔT in a material generated per degree difference between the points, expressed as:

V    , (3.2) T

where a positive or negative sign indicates holes or electrons as the dominant carrier, respectively.

3.2(b) Hall effect

In contrast to the Seebeck effect, in which an electric potential is generated as a result of diffusion of higher-energy carriers, the Hall effect is an electric potential resulting from the Lorentz force on the drift current in the presence of a perpendicular magnetic field. It is defined as follows: consider a rectangular conductor of thickness t and width w, through which a current I is flowing, as depicted in Figure 3.2.

Suppose a magnetic induction B is applied as indicated in the direction t, then a transverse voltage VH, the Hall voltage, appears between the conductor edges, as shown. The transverse Hall electric field is given by

Ey  RH JB , (3.3) where RH is the Hall coefficient and J is the current density, given by I/wt.

The force on any particle of charge e moving with velocity v through a magnetic field of strength B is given by:

F  ev  B . (3.4)

If the moving particle is part of a stream of carriers making up a current in a material, then the Lorentz force will act on all the carriers in the current, causing them to move towards a side of the material perpendicular to both the direction of the current and the magnetic field. In a manner similarly described for the Seebeck effect, given above, in a single-carrier material, the buildup of carriers on one side will result in an electric field in the opposite direction of the Lorentz force, eventually equalizing its effect. It is the generation of this field which is known as the Hall effect.

Considering a p-type semiconductor, the Lorentz force on each hole, given in

Equation 3.4, is balanced by the electrostatic Hall field voltage VH. Therefore,

eE H  evB , (3.5a) or,

EH  VH / w  vB , (3.5b)

Given that the horizontal current I is equal to current density J multiplied by dimensions w and t, and J is given by J = pev, where p is the hole concentration, and v is the hole drift velocity, the Hall field can now be expressed as,

V JB E  H  . (3.6) H w pe

Comparing this to the Hall coefficient law, expressed in Equation 3.3, the Hall coefficient can now be expressed as,

1 tV R   H . (3.7) H pe IB

3.3 TRANSPORT FORMULAE

3.3 (a) Single carrier model

Consider a simple semiconductor with a non-degenerate concentration of holes (p), scattered by dominant acoustic modes. The Hall coefficient is given by:

3 R  , (3.8) H 8pe

where e is the electronic charge. In this equation, the factor 3π/8 arises from acoustic lattice scattering which goes to unity as degeneracy prevails.

Conversely, for a semiconductor where the electron concentration, n, is dominant over the hole concentration, the Hall coefficient is given by

3 R   . (3.9) H 8ne

Under the same non-degenerate conditions, the thermoelectric power for dominant hole and electron material is respectively given by:

k  EV  EF   p  2   , (3.10) e  kT 

and

k  EF  EC   n   2   , (3.11) e  kT 

where k is Boltzmann’s constant, T is absolute temperature and EF, EC and EV are respectively the Fermi level and energy positions of the conduction and valence band edges, as shown in Figure 3.1. The factor 2 in the brackets changes with degeneracy and different scattering modes.

3.3 (b) Two carrier model

Consider now the case when both electrons and holes are involved, as in intrinsic or near-intrinsic situations, again with non-degenerate concentrations and dominant acoustic lattice scattering. Under these circumstances, the Hall coefficient is given by:

2 2 3 p p  nn RH  2 . (3.12) 8 e( p p  nn )

Rearranging Eq. 3.12 yields:

n 1 b 2 3 p RH  , (3.13) 8pe n 2 (1 ) p

where µp and µn are respectively the mobilities of the holes and electrons, and b = µn/

µp. Equation 3.13, which is function of hole concentration and n/p, can also be expressed as a function of the electron and hole densities directly:

3 nb 2  p R   . (3.14) H 8e (nb  p)2

Under the same conditions, the thermoelectric power, with two carriers involved, is expressed as:

 E  E E  E  p (2  V F )  n (2  F C ) k  p kT n kT      . (3.15) e  p p  nn   

30 Rearranging Equation 3.15 yields:

 E  E  n  E  E  2  V F   b2  F C   kT p kT      . (3.16) k n e 1 b p

If the bandgap of this material is taken as EG, Equation 3.16 can be expressed as:

n  E  2  b2  G   p  kT  EV  EF   . (3.17) k n kT e 1 b p

It is also possible to express Equation 3.15 in terms of the effective density of conduction and valence band states, NC and NV, respectively:

k 1   n   p     2(nb  p)  nbln   pln  , (3.18) e (nb  p)   NC   NV 

where NC and NV are given by

2(2m kT)3/ 2 N  n , (3.19a) C h3

31 and

2(2m kT)3/ 2 N  p , (3.19b) V h3

where h is Planck’s constant and mn and mp are respectively the effectives masses of electrons and holes.

3.4 SIGN REVERSAL OF α AND RH

In a material which is nominally p-type at room temperature, the sign of the

Hall coefficient and thermoelectric power from positive to negative can be observed

(a) as the temperature is increased in semiconductors where µn exceeds µp, and (b) as extra donors are added to the material at a fixed temperature. Both types of inversion are now considered.

3.4 (a) Intrinsic temperature range inversion

Consider a simple semiconductor with a small or modest energy gap, EG, so that the material can be intrinsic at a few hundred degrees centigrade but is extrinsic at room temperature. Consider also the initial hole concentration in the valence band p to be equal to NA, the acceptor concentration, while the donor concentration is zero

(ND = 0). As the temperature is increased, both electrons and hole increase due to the addition of intrinsic carriers.

32 2 As can be seen by examination of Equation 3.14, RH is positive when nb is less than p, zero when nb2 is equal to p, and negative when nb2 exceeds p. Therefore, in a material where µn > µp, it is possible for RH to be negative, indicating apparent n- type material, even though p is always larger than n. The transition of RH from positive to negative, as n approaches p, is largely dependent on the ratio of mobilities b.

For thermoelectric power, α becomes zero when

n  E  2  b2  G  E  E p kT V F    . (3.20) kT n 1 b p

For a given value of b and EG, this condition will be satisfied at a different ratio of n/p. Therefore, as the temperature of the semiconductor material is increased so that the value of ni is comparable to that of the doping NA, α and RH will be seen to change sign at different temperatures. This effect is shown in Figure 3.3, where α and

RH are plotted against temperature for two different semiconductors. In this example,

17 -3 NA was set to 5x10 cm , while ND was kept at zero. Both semiconductors had electron and hole effective masses of 0.09mo and 0.7mo, respectively, and the ratio of mobility b was 10. However, the bandgap of the first material is equal to 20kT at room temperature, half that of the second material with 40kT, the latter of which more closely resembles the actual bandgap of CuInSe2. In these plots, the bandgaps

33 of both materials were kept constant as the temperature was increased; in other words,

EG/kT decreased with temperature.

As can be seen in Figure 3.3, as the temperature is increased and the majority and minority carrier concentrations tend towards intrinsic, in materials where µn/µp >

1, there exists a temperature range in which α and RH will have opposing signs.

However, as can also be seen in the figure, for CuInSe2, with an acceptor

17 -3 concentration of NA=5x10 cm , this transition range is from about 750°C to 950°C.

While it is beneficial to understand and be aware of this phenomenon, it is seen to occur beyond the range of temperatures at which the thermoelectric power and Hall measurements on the material in this work were done.

3.4 (b) Inversion by doping at fixed temperature

Consider a non-degenerate semiconductor, held at room temperature, with a fixed doping density of acceptors NA much higher than the intrinsic carrier concentration ni, with complete ionization. Now consider that donors of concentration ND are progressively added to the material. If ND is much less than NA, the concentration of holes in the material, p, can be approximated simply as NA, while

2 the concentration of electrons n is equal to ni /NA. As ND approaches NA, the material becomes compensated, and p is approximately equal to NA-ND, while n is still given

2 14 by ni /p. In the range of doping of most semiconductors, which is between 10 -

19 -3 10 cm for NA and ND, as ND is increased past NA, an abrupt change occurs in which

2 n dramatically increases (to values approximating ND-NA), while p decreases to ni /n.

For a semiconductor such as CuInSe2, with a bandgap of 1.04 eV and effective mass ratios for electrons and holes of mn/mo = 0.09, mp/mo = 0.7,

9 -3 respectively [3.3], the intrinsic carrier concentration is of the order of 6.5 x 10 cm at

17 -3 room temperature, while doping levels are of the order of 10 cm for bulk crystals

[3.4-3.5]. At these values, the minority carrier, in most applications, can be neglected, and only the majority carrier considered. Inspection of Eq. 3.14 and Eq.

3.18 indicate that the calculation of thermoelectric power and Hall coefficient are examples of such applications. Therefore, the single-carrier RH and α expressions of

Equations 3.8 and 3.9, and Equations 3.10 and 3.11, respectively, can be used to accurately determine the carrier concentrations in extrinsic material. This is further emphasized in Figure 3.4, where the single-carrier expressions for RH and α are seen plotted as solid lines, along with the two-carrier expressions (plotted as unconnected data points), against ND-NA for CuInSe2 at room temperature, using the values for the characteristic properties of the material given in this paragraph. As can be seen, the single and two-carrier expressions are indistinguishable. Note also that, unlike the inversion by temperature, described above, in the case of increasing temperature in

CuInSe2, both α and RH are seen to change sign exactly when ND is made larger than

NA.

In this thesis, the inversions, observed and described in Chapter 5, are those by doping (obtained by adding sodium). Accordingly, the sign inversions of both α and RH should occur at exactly the same concentration of dopant, and the majority carrier concentrations can be accurately calculated by considering only the single-

35 * * carrier equations (assuming mn and mp are known for the calculations of NC and NV, respectively). As such, the considerations of Section 3.4(a) are not pertinent to the present experimental work, but are presented for clarity. Thus, α and RH can change sign at different temperatures but not at different doping levels.

conduction band EC

valence band

Figure 3.1 Energy band diagram for an n-type semiconductor material, for which EG, the energy bandgap of the material, is within the range of 1 eV.

Figure 3.2 An illustration of the Hall-effect on a p-type semiconductor material.

Figure 3.3 Two-carrier thermoelectric power (black) and Hall coefficient (red) shown plotted versus temperature for two materials of different bandgaps, the larger of which approximates the actual bandgap of CuInSe2. As can be seen, a range of temperatures exists at which α and RH are of opposing signs, despite the fact that the hole concentration exceeds the electron concentration at all temperatures. The values of the bandgaps given in the figure are true at room temperature (T = 300K), and are assumed to remain constant with increasing temperature.

Figure 3.4 Thermoelectric power (α) and Hall coefficient (RH) versus ND – NA for 17 -3 CuInSe2 at room temperature, where NA is kept at a constant value of 5x10 cm . The solid lines represent the respective single-carrier equations, while the points are from the two-carrier equations. As can be seen, both α and RH are seen to change sign exactly when ND is made larger than NA.

39 Chapter 4

Experimental Procedure

4.1 INTRODUCTION

A comprehensive description of the Bridgman-growth process was included in the master’s thesis written by the author and published in October 2008 [4.1]. It has also been described in various publications [4.2-4.3]. This process, in which the starting elements are loaded in quartz ampoules, heated past the melting point of the material and then directionally cooled, has changed little from these descriptions.

Therefore, a detailed account here would be unnecessary. However, as it is an integral part of this work, a list of steps of the process, as well as a brief description of each step, is provided in this chapter.

A major deviation in this work from previous workers in this laboratory [4.4-

4.5] is the proportion of starting elements included in the ampoules before growth.

Additionally, some variation from the heating process used by previous workers has been made. It is believed that the use of Na increases the possibility of ampoule cracking, due to reaction of this element with the quartz. The growth process has therefore been lengthened in an attempt to heat the ampoules more gradually and minimize stress. Furthermore, in some runs, the maximum temperature of the upper- zone of the Bridgman furnace has been reduced slightly, from 1100 °C to 1050 °C, so

40 as to increase the life of the furnace element. The specific growth scheme used in the growth of each ingot is described in the tables of Appendix A.

As well, in further attempts to understand the effects of Na on CuInSe2, an additional experiment was conceived in which this element was heated in a sealed ampoule with already grown CuInSe2. A separate experiment was also carried out where Na was heated with Se alone.

Following the description of the ingot growth procedures, and additional experiments, is a description of the equipment used in the characterization and analysis of the material after growth. The analysis can be divided into two categories: electrical, in which the transport properties of the material were measured, and structural, in which the elemental composition and crystal structure of the ingots and deposits were investigated. In the former category, comprising thermoelectric power, resistivity and room temperature Hall-effect measurements, the equipment used was largely custom-designed and fabricated for this work. This is described in Section

4.3. The etching experiments might conceivably be included in the electrical measurements, as the etchant used was specifically developed to cause a visible deposit to form only for p-type material, and thereby allow the type of material to be identified. For the structural investigation of the material, comprising X-ray diffraction and SEM-EDX (scanning electron microscopy – energy-dispersive X-ray spectroscopy) measurements, the equipment used in the measurements was made by

41 Bruker and Philips, respectively, as was the software used in the analysis and interpretation of the measurements.

4.2 INGOT GROWTH PROCEDURE

This section contains an overview of the procedure used in the growth of the ingots. As mentioned above, a more detailed description of the each step can be found elsewhere [4.1].

4.2 (a) Ampoule and charge preparation

Tubes of fused quartz, having an inner diameter of 12 mm and an outer diameter of 16 mm, were formed into ampoules and cleaned using an HCl / HNO3 solution known as aqua regia (3:1 HCl:HNO3, respectively), in which the ampoule was allowed to soak for 24 hours, and then in acetone, for approximately 3 hours.

The thickness of the quartz walls was necessary in order to strengthen the ampoules and prevent them from exploding as a result of the vapour pressure of the selenium at high temperatures. A coating of boron nitride was applied to the inner wall of the ampoule and flamed-in. This coating procedure was discovered in this laboratory

[4.6] to minimize voids and cracking in the grown ingots, and to prevent sticking to the ampoule walls.

42 The high-purity copper, indium and selenium pellets were etched for approximately 2 minutes, the first element in 10 % nitric acid (HNO3) and the last two in 10 % hydrochloric acid (HCl), rinsed in DI water and dried thoroughly, before being weighed in the appropriate atomic proportions and placed inside the ampoule.

The ampoule was then evacuated, to a pressure of about 10-7 torr, and sealed.

Sodium, obtained from Alfa Aesar, was kept in oil before weighing. The oil was removed using toluene only immediately before the Na was added to the ampoule so as to avoid reacting with the air. In order to compensate for the oil, 10% above the desired amount was weighed. Chemical analysis on the sodium, performed by VHG

Labs in New Hampshire, has determined the purity of the material to be in excess of

99.99%, with trace elements Fe and Al present in concentrations of 1 ppm, and Ca present at 41 ppm. The details of this analysis are included in Appendix B. One ampoule containing the charge, after evacuation and sealing, is shown in Figure 4.1.

Before undergoing the Bridgman-growth procedure, the sealed ampoule was carefully pre-heated in a Thermolyne-type 10500 brick furnace. This was to cause the Se and In to react in an enclosed environment rather than in the exposed upper- zone of the Bridgman-furnace. The selenium vapourizes during this reaction, resulting in pressures up to 40 atm to be generated within the ampoule. If a defect exists, it might cause the ampoule to break or explode; therefore, it is preferable that it occur at this stage of fabrication, which takes place in an enclosed container, rather than in the exposed Bridgman furnace, resulting in less damage to equipment.

43 As mentioned above, the specific pre-reaction heating scheme used for each ingot is given in Appendix A.

4.2 (b) Bridgman-growth

After the pre-reaction, the ampoule was attached to a quartz rod and placed in the upper-zone of the Bridgman-furnace, as pictured in Figure 4.2. The temperature was then gradually raised to a maximum 1050°C, above the melting point of CuInSe2 of 986°C, and maintained at this temperature for 24 hours. During this time, the ampoule was temporarily removed from the furnace, rocked back and forth to check liquid flow, and then quickly returned, ensuring that the solution was well mixed and homogeneous. This constituted one agitation. Often, several agitations were carried out in a growth run (see Tables A1-A5 of Appendix A).

After 24 hours, the ampoule was lowered into the lower-zone, which was kept at 700 °C. Once the entire length of the ampoule was in this zone, the temperature was lowered to ambient over 42 hours. The quartz ampoule was then broken open and an ingot was obtained, as shown in Figure 4.3.

When growing CuInSe2 ingots in this laboratory in earlier studies, researchers, including this author, used a maximum upper furnace temperature of 1100 °C.

However, as discussed in more detail in Section 4.5 of this chapter, it was noticed in this work that the furnace elements (manufactured by Thermcraft) frequently failed,

44 sometimes causing an ampoule to be abandoned. Therefore, the upper furnace temperature was lowered to 1050 °C for most of the growth runs made in this work.

4.3 ANALYSIS EQUIPMENT

This section comprises descriptions of the equipment used in the electrical and structural measurements of the ingots and deposits obtained in this study.

4.3 (a) Hot probe

The hot probe was used to determine the conductivity type of wafers cut from the completed ingots. It consisted of a modified large surface soldering iron, powered through a Variac, with a type-J thermocouple attached near the tip, as pictured in

Figure 4.4. The thermoelectric power was measured by placing the probe on the surface of the wafer while recording the voltage induced between a copper wire, attached to the soldering iron tip beside the thermocouple, and a brass base-plate, considered to be at room temperature. The voltage measured across a wafer was seen to increase linearly with increase of temperature of the probe, as shown in Figure 4.5.

One major advantage of the hot-probe method is the ability to record data points across the surface of a wafer. The final value obtained is an average of these points, and thus, perhaps, is more characteristic of that ingot in general than if only one point was taken. Furthermore, inhomogeneities within the ingot, where the

45 thermoelectric power is greatly different from the average of the other points, can be identified. These may include n-type islands in a mostly p-type ingot. The thermoelectric power can also easily be measured over the various regions of the ingot (first-zone-to-freeze or last-zone-to-freeze). However, there exists some uncertainty as to the accuracy of the measurements, particularly as the data is recorded at various locations but at only one temperature, and contact between the probe and the wafer may heat the underside of the wafer to a value above room temperature. This would result in a temperature difference that may be very different from the value indicated by the single thermocouple. For this reason, the hot probe was used in this study only as a qualifying tool to determine the conductivity type of a sample (p or n), as well as to verify the electrical homogeneity of an ingot. In order to more precisely determine the thermoelectric power of a sample, a specialized sample holder, described below, was built.

4.3 (b) Sample holder for thermoelectric power measurements

The sample holder, schematically drawn in Figure 4.6a and pictured in Figure

4.6b, allows the temperature and voltage to be simultaneously recorded at two locations of a sample. As can be seen, it consists of a Teflon holder with a rectangular notch cut out of the top face, used to support the sample, and copper blocks placed at either end. An electrical resistor (10 Ω) was embedded in one of the blocks and connected to a variable voltage source. As the voltage across the resistor was increased, the block was gradually heated and thus acted as a heat source, while the opposite block, acting as a heat sink, stayed relatively cool.

46 As can be seen in Figure 4.6, the sample of material used in this measurement is rectangular in shape. Filaments of material were cut from the ingots and abrasively lapped with 600-c grit sandpaper until the desired dimensions of approximately 2 x 2 x 8 mm3 were reached. The filaments were then soaked in warm aqua regia (3HCl +

1HNO3 heated to 70 °C) for 30 seconds, rinsed with DI water and dried. Gold was evaporated onto the ends of the filaments in the case of the p-type ingots, followed by

Wood’s metal (a low melting point solder made from bismuth, lead, tin and cadmium). In the case of n-type material, Wood’s metal was applied directly to the sample; Au and Wood’s metal were previously shown to form ohmic contacts to p- and n-type CIS, respectively [4.7]. The filaments were then placed in the groove, cut through the top of the Teflon block, which was subsequently filled with indium in order to form good thermal contacts between the ends of the filament and the copper blocks. Two type-J (iron-constantan) thermocouples were embedded in the indium, one on each end of the filament.

In order to measure and record the thermoelectric power of each filament, approximately 4 V was applied across the embedded resistor. The voltage across the

“cold-side” thermocouple was measured and recorded at 0.05 mV intervals of the

“hot-side” thermocouple as it was increased from 0.25 mV to 1.10 mV (18 data points), corresponding to a temperature range of 5 °C to 22 °C above ambient. Also, the voltage induced across the filament due to the Seebeck effect, discussed in

Chapter 3 of this thesis, was measured, with positive and negative terminals connected to the iron leads of the warm and cold-side thermocouples, respectively.

47 The data represents the “warming-up” set of values, an example of which can be seen in Figure 4.7 for a filament taken from ingot HM-B26. Once the value of the warm- side thermocouple reached 1.10 mV, the power connected to the resistor was shut off, and a second set of data points were taken as the heat-sink copper block gradually returned to room temperature.

The temperature difference across the filament was determined by interpreting the voltage reading of each thermocouple to obtain a value of temperature, in °C relative to room temperature, and then subtracting the “cold-side” temperature from the “warm-side.” This was more accurate than simply considering the cold-side to be at room temperature, as it was seen to vary by a few degrees over the course of the measurements. Finally, values of thermoelectric power were calculated at each data point by dividing the voltage induced across the filament, measured across the two iron leads, by the temperature difference. This can also be seen plotted in Figure 4.7 for sample HM-B26. By convention, positive or negative thermoelectric power values were associated to either of p or n-type samples, respectively.

Note that at low temperature differences, in both the warming-up and cooling- down plots, the apparent thermoelectric power is relatively high as compared to the values obtained at greater temperature differences. The cause of this discrepancy was not investigated and was assumed to be an effect of stray voltages, which are significant when the voltage across the filament are small but become less-so as the temperature difference is made greater, where the data points appear to descend

48 asymptotally toward a single value. Therefore, for each sample, the thermoelectric power was determined by averaging only 10 data points, in each of the warming-up and cooling-down plots, taken at the 10 greatest temperature differences for each sample. The thermoelectric power of iron, +15 mV/K [4.8], was then added to this average in order to obtain the absolute thermoelectric power value for each sample.

4.3 (c) Sample holder for electrical resistivity measurements.

The electrical resistivity (ρ) of the samples was measured by applying a constant current of 1 mA through the ends of the sample while recording the voltage at positions 0.5 mm apart along the length filament. A schematic representation of voltage probing for resistivity measurements is shown in Figure 4.8a, while a photograph of the apparatus is shown in Figure 4.8b. These voltages were then plotted versus position, as shown in Figure 4.9.

The main advantage of the voltage-probe method, used to determine the resistivity of the samples, was detecting high-resistance regions, such as at cracks and at end contacts, which could be identified in the voltage plot, allowing them to be avoided when calculating the resistivity of the sample. The resistivity was calculated using only the linear regions of a plot, using the formula shown below.

V  A   , (4.1) I l

49 where I represents current, A represents the cross-sectional area of the filament, and l represents the length of the filament in the region of voltage difference V. Voltage probing is therefore more accurate in determining the resistivity of a filament than simply measuring the voltage and current across and through the ends of the filament.

However, as the brittleness of the ingots increased, particularly with high Na additions, the resistivity measurements became increasingly difficult.

4.3 (d) Hall coefficient measurements

Room-temperature Hall-effect measurements were performed after the thermoelectric power and resistivity measurements had been completed, thereby taking further advantage of the efforts made in forming the filaments, as well as in applying the end contacts. For the Hall measurements to be carried out, two additional contacts were placed on opposite sides of the filaments, near to the middle.

The side onto which a contact was to be placed was prepared by first being rubbed with a cotton swab, soaked with aqua regia, for 30 seconds, then rinsed with DIW and air dried. Gold, in the case of the p-type samples, was then evaporated, through a mask with an opening of diameter 2 mm, onto the filament. This entire process, including the etching with aqua regia, was then repeated for the opposing side.

Wood’s metal was used to secure a thin copper wire to each of the gold spots, which were in turn connected to terminals at located at the end of a sample holder. In the case of n-type samples, Wood’s metal was applied directly to the filaments, and so the gold evaporation step was omitted.

50 The end of the sample holder containing the sample was positioned between the poles of a large, water-cooled electromagnet, powered by a large AC current source, controlled with a variable autotransformer, before being converted to DC and run through the magnet. A separate current source (Keithley 244 programmable current source) was used to supply current through the ends of the filament. The magnetic field was then increased while the Hall voltage was recorded at several different field intensities. A schematic representation of these measurements can be seen in Figure 4.10, while a plot of voltage versus magnetic field for two separate samples (one p-type, and the other n-type) can be seen in Figure 4.11. Reversal of sample current and magnetic field were necessary to separate the Hall voltage from other stray voltages.

4.3 (e) X-ray equipment

The use of X-ray diffraction served to confirm the chalcopyrite structure of the ingots, as well as to identify any extra phases which may be present in the material. Ampoule deposits were also analyzed using this method, as described below. The bulk samples were manually powdered using a porcelain mortar and pestle, while the ampoule deposits were sometimes left intact. In all cases, a Bruker

Discover D8 diffractometer was used with a VANTEC 2 dimensional area detector and a Cu-Kα source (1.54056Å wavelength), at 40 kV operating voltage and 40 mA operating current.

51 4.3 (f) SEM-EDX

Scanning Electron Microscopy – Energy Dispersive X-ray Spectroscopy

(SEM-EDX) was performed on some samples using a Philips XL30 FEG SEM with

Orientation Imaging Microscope (OIM) and equipped with an EDS Analyzer Scan

Generator EDI-2. SEM-EDX analysis was used to investigate the deposits, as well as the interior and exterior of the grown crystals. In the latter cases, wafers of material were polished for 20 minutes with 0.5 µm alumina powder, followed by a rinse in toluene, acetone and DI water. The wafers were then immediately (within 10 minutes) placed in the SEM chamber, which was subsequently evacuated. This was done quickly so as to minimize oxidation of the samples. The SEM was used to aid in the visual inspection of the samples, in order to search for microscopic patterns or markings which could differentiate various phases, or other inhomogeneities such as twin or grain boundaries. EDX served as another tool by which the elemental composition of the samples could be determined.

4.4 ADDITIONAL EXPERIMENTS

4.4 (a) Heating experiments

As will be discussed in subsequent sections, when Na was included in the growth-melt of ingots, the ingot fabrication temperatures caused the Na to react with the fused quartz of the ampoules. This reaction was seen to weaken the ampoules, sometimes causing cracks to form at some point during the growth procedure. This weakening and cracking occurred more often in runs with large amounts of Na, such

52 as in the case of the CuInSe2.05 + 11 at. % Na ingot, thereby imposing an upper limit on how much Na could be included before growth.

It was thought that some of the effects of including Na in the melts, such as the presence of Na-based compounds in the deposits, are too weak to be easily detected in the previously-described growth runs. Increasing the amount of Na in the ampoule so as to enhance these effects, as discussed above, is limited by the reaction of Na with the quartz at high temperatures. Therefore, in order to circumvent this limitation, an experiment was devised whereby the material could be only modestly heated with a relatively large amount of Na.

A piece of quartz tubing was first cleaned, and then formed into an ampoule, according to the description in Section 4.2(a). A small, solid piece of previously grown p-type CuInSe2, from a stoichiometric melt, taken from the middle region of the ingot, was abrasively polished on all sides with 600-c grit sandpaper, and then rinsed with toluene, acetone and DIW. The piece was then placed in the ampoule, along with a very large amount of elemental Na equivalent to an atomic ratio of 1:2

CuInSe2 to Na, respectively, as shown in Figure 4.12. In a regular growth run at 1050

°C, this amount of Na (i.e. 66.7 at. %) would have caused a breakage. The ampoule was immediately evacuated to a pressure of 10-7 torr, sealed and placed into the brick furnace. It was then heated according to the scheme shown in Figure 1A of Appendix

A to 300 °C. After two hours, the furnace was shut off, but the ampoule remained

53 inside until the system had cooled to room temperature. Afterwards, the ampoule was broken open and the contents quickly photographed.

It was immediately apparent when the ampoule was broken open after the heating that the original CIS material had undergone obvious structural changes, as portions of it were now very brittle. Deposits were also seen in the fragments of material and on the sides. In order to investigate these changes, portions of the brittle material were manually powdered, using a mortar and pestle, and this powder was subjected to X-ray diffraction. This was done as quickly as possible in order to avoid contamination of the components of the powder with the air, and specifically those which might contain Na. Additional to this powder from the bulk CuInSe2, some of the deposits on the quartz ampoule were analyzed by XRD.

This experiment was repeated, this time using only Se, instead of CuInSe2, with a much larger amount of Na, according to an approximately 4:1 Na to Se atomic ratio, respectively, as shown in Figure 4.13. As the ampoule in the first heating experiment did not crack, and no reaction between the quartz and the Na was apparent, the ampoule was once again heated according to Figure 1A of Appendix A, except that it was brought to the slightly higher temperature of 400 °C, which was maintained for two hours. Once again, a clear change had occurred, as the contents were visibly quite different from how they appeared before the heating. XRD was used to identify the composition of the post-heat treatment material. The results of these analyses are given in Chapter 5.

54 4.4 (b) Cold powder mixing experiments

As will be shown in the analysis of the thermoelectric power results in

Chapter 5, relationships exists between the conductivity type of a grown ingot and the starting proportion of elements placed in the ampoule before growth. One such relationship involves the proportion of elemental Na and excess Se, the amount of Se added to the ampoule above the atomic proportion of 2:1 Se to either Cu or In, respectively. This relation implies that a reaction between the Na and Se occurs in the ampoule during growth, which could result in reaction products containing Na.

However, no such compounds were detected by X-ray diffraction of powdered material after regular growth, either from the main part of the ingots, the last-zones- to-freeze, or the deposits. Na had been found in EDX analysis of the elemental compositions of some deposits using methods other than XRD, as will be shown in the next chapter, but could not be confirmed to be part of a compound containing Na and Se. It was hypothesized at the time that these reaction products, if they are in fact formed, may be at too small an amount to be detected by XRD, or that a reaction with air, after removal from the evacuated ampoule, might have caused the elements to dissociate before XRD could be performed. Therefore, an experiment was devised in order to test whether it was possible to detect a secondary Na-Se compound by XRD within a CuInSe2 powder.

A piece of CuInSe2 was taken from a stoichiometrically-prepared ingot and manually pulverized using a mortar and pestle. Separately, an amount of already- synthesized, pre-made Na2Se, obtained from the Alfa Aesar Company (powder,

55 packed in Argon, 99.8% purity) was also pulverized. A quantity of the CuInSe2 was weighed, and then an amount of Na2Se was measured that was equivalent to 10 at. % of the weighed CuInSe2. The two powders were then carefully mixed together in such a way as to avoid heating by friction, which might cause a reaction between the materials. It was observed that the Na2Se was a rust-coloured powder, but that after mixing with the CuInSe2, which was grey, the two materials were indistinguishable, and the resulting powder was completely grey. XRD was immediately conducted on the unheated mixture, within approximately 10 minutes of the pulverizing of the

Na2Se, to minimize reactions with the air.

This was then repeated with another unheated mixture comprising CuInSe2 together with 100 at. % Na2Se, comprising a 1:1 atomic ratio of CuInSe2 to Na2Se.

As this latter experiment was being conducted, the objective was again to analyze the material via XRD as quickly as possible. However, it was discovered only after the mixture was prepared that the XRD machine was out of order, and that immediate analysis was impossible. The machine was made operational four hours later, at which point the powder, which sat in a Petri dish in an ambient environment throughout this period, was analyzed. It was thought that the powder may have become contaminated during this time through a reaction with the air, and so this experiment was repeated. A second mixture with an atomic ratio of 1:1 Na2Se to

CuInSe2 was prepared, for which XRD was conducted within 10 minutes of the preparation. The results of these experiments are given in Chapter 5, where it will be seen that differences in the XRD patterns of the two mixtures (one four hours old and

56 the other ten minutes old) indicate that a reaction between Na2Se and air did indeed take place. The details of this reaction could also be ascertained through analysis of the phases within the mixture. Eventually, it will be shown that the knowledge gained by understanding this reaction contributed in identifying phases resulting from the addition of Na into the growth melts of Bridgman-grown CuInSe2 ingots and, with that, in elucidating the mechanism behind the structural and electrical effects of Na on

CuInSe2.

4.4 (c) Etching experiments

Extensive work has been carried out on the effects of various etching solutions by Tell and Bridenbaugh [4.9]. Over the course of that work, it was discovered that an etching solution combining equal volumes of hydrofluoric acid, nitric acid and water has a different effect on p and n-type bulk CuInSe2 material. Specifically, immersing a p-type sample in the solution at room temperature was shown to result in a red deposit that could be seen on the surface of the material, while etching with an n-type sample was found not to result in any deposits. They used this to determine, chemically, the precise location of a pn-junction in their material. The effects of this etch was later verified by Ahmad [4.7], in this laboratory.

In order to serve as a compliment to the hot-probe method of measuring the conductivity type, one wafer from ingot HM-B39 (p-type) and another from HM-B44

(n-type) were immersed in this etch for five minutes, followed by a quick rinse with

57 DIW. An etch of equal volumes of hydrofluoric acid, nitric acid and hydrochloric acid was also tried. The results of this experiment are given in Chapter 5.

4.5 ADDITIONAL REMARKS ON GROWTH EXPERIENCE

As mentioned above, the specifics of the fabrication procedure used for each individual ingot grown in this work are given in Appendix A. Over the course of this work, slight changes have been made to the ingot fabrication procedure, the aims of which were to increase the likelihood of success of the growth process. As this project involves the study of the effects of Na in the growth melts of the ingots, it is preferable to change the fabrication procedure as little as possible, so as to make the inclusion of Na the only variable. However, it was discovered early in the precursor to this work [4.1] that, more often than that experienced by previous workers, the ampoules sometimes cracked while in the Brick furnace, and other times cracked while in the Bridgman furnace after having survived the pre-reaction heating to 300

°C. When the cracking occurred in the Bridgman furnace, it often occurred during the initial heating stage. In order to solve some of these problems, the pre-reaction temperature was increased to 400 °C, so as to cause the elements to react more completely. It is believed that the cracking is an effect of adding the Na in the ampoule, as there were seldom problems with ampoules which did not include Na.

As will be seen in the next chapter, XRD was used to confirm that a structural change in the quartz occurred with additions of Na, which caused the quartz to weaken.

As well, the maximum temperature of the upper furnace was lowered from

1100 °C to 1050 °C, which was done to increase the lifetime of the furnace elements.

It was thought that insulating pads attached to the elements, first used for ingot HM-

B43, would allow for the soak temperature to be brought back to 1100 °C. This was done for ingots HM-B44 to HM-B66, requiring about five sets of heating elements.

However, the new furnace elements used for HM-B67 burnt out during their first use, as did a replacement set. Likely, these burnouts were due to some manufacturer error, but the soak temperature was reduced back to 1050 °C nonetheless.

In addition, the heating ramp-up scheme used in the Bridgman furnace was lengthened significantly from that used by Du [4.5] in further attempts to avoid cracking due to the Na, although the scheme was sped-up slightly for ingots HM-B60 through to HM-B63. It was thought that these runs, which did not contain sodium, posed less of a breaking risk than other runs.

Se In Cu Na

Figure 4.1 A quartz ampoule containing a charge of Cu, In, Se, and Na, shown after evacuation and sealing but before heating.

Figure 4.2 Schematic of the Bridgman-furnace with the ampoule shown in the starting position in the upper-furnace. The temperature profile of the furnace is shown at the right of the figure for an upper-zone temperature of 1100 °C.

Figure 4.3 A complete ingot, shown immediately after removal from the quartz ampoule after the completion of the Bridgman-growth process (held in rubber gloves).

Figure 4.4 The hot probe used to determine the conductivity type of the ingots prepared in this work.

61 p-type

n-type

Figure 4.5 Thermoelectric voltage versus temperature difference between the hot- probe tip and the underlying metal base plate for samples 26 (n-type) and 13 (p-type), showing essentially linear variations.

Figure 4.6 (a) Schematic of sample holder used to measure thermoelectric power and electrical conductivity on filamentary samples; (b) actual photo of the sample holder, set up in thermoelectric power measurement mode.

Figure 4.7 Voltage, shown in black, measured across a filament from ingot HM-B26 as the temperature of the copper block was increased in the sample holder. Shown in red is the thermoelectric power calculated for each data point, which is seen to converge at higher temperature differences.

a) b)

Figure 4.8 (a) A picture of a filament with a schematic depicting resistivity measurements by voltage probing; (b) photo of the sample holder set up in resistivity measurement mode.

Figure 4.9 Schematic and plot of voltage probing with distance from one end along two filamentary samples. Sample 1 shows a linear variation along most of the filament but high contact resistance at the two ends. Sample 2 shows low-resistance ohmic contacts at the ends but a large discontinuity near the middle of the filament.

Figure 4.10 Schematic representation of a filament of CuInSe2 during a Hall-effect measurement.

Figure 4.11 Plot of Hall voltage versus magnetic field for one p-type sample (right) and one n-type sample (left). Note the opposing signs of the Hall voltage between the samples at the same current direction.

Na CuInSe2

Figure 4.12 A piece of p-type CuInSe2, taken from a stoichiometrically-prepared ingot, shown in an evacuated and sealed ampoule with a large amount of Na before heating to 300 °C. The atomic ratio of CuInSe2 to Na is 1:2, respectively.

Na Se

Figure 4.13 A very large amount of Na seen in an evacuated and sealed ampoule, with a quantity of Se, equivalent to an atomic ratio of approximately 4:1 Na to Se, respectively. The ampoule was then heated to 400 °C.

66 Chapter 5

Experimental Results

5.1 INTRODUCTION

The main focus of this work is a study of the effects of Na on the electrical and structural characteristics of CuInSe2. In this chapter, the observed effects of Na on the grown material are reported when this element is added to quartz ampoules containing Cu, In and Se, followed by the Bridgman-growth procedure, as described in Chapter 4.

The electrical effects are described first. As will be shown, the grown material was nominally p-type when nominally stoichiometric proportions of the starting elements were included in the growth melts, with no Na, as determined by hot-probe analysis. A small amount of Na, below 0.3 at. %, also resulted in p-type material. However, at and above 0.3 at. % Na, the material, with otherwise nominally stoichiometric proportions, was found to be n-type. These results are consistent with the work of H.P. Wang [5.1].

Additionally, it was shown that when a small amount of excess Se above stoichiometry was also included in the ampoules before growth, the change in conductivity type from p to n by the Na was mitigated, so that more Na was required

67 for the type change to take place. Furthermore, the thermoelectric power results on filaments of material, cut from the middle part of the ingots, given in Section 5.2 (b), are used to identify a relationship between the Na and excess Se, and conductivity type. As will be explained, this relationship, discovered in this work, is the foundation of the mechanism used to describe the action of the Na on the material.

The results of the room-temperature Hall-effect measurements are then given in Section 5.2(c), and are used to reinforce the results of the thermoelectric power measurements. Finally, they are combined with the resistivity measurements of

Section 5.2(d) in order to explore the effects of Na, and excess Se, on the mobility of the material. The results of all the electrical measurements performed in this work can be found in Tables C1 and C2 of Appendix C.

Following this, the structural effects of Na are explored in Section 5.3. Here, the results of the crystal structure and compositional analysis, by XRD and SEM-

EDX, are given. Details of the results of EDX analysis on the interiors of bulk crystals are provided in Table C3 of Appendix C. As will be shown, no changes in the crystal structure of CuInSe2 were apparent with additions of up to 3 at. % Na in stoichiometric material, and no Na was detected within the bulk material, but this element was detected on the exterior surface of the ingots, and at defects. Deposits, seen in the ampoules after growth, were also examined, as described in Section 5.4, revealing, among other things, an unintended and adverse reaction between the Na and the quartz, which was most apparent in samples containing much Na and little

68 excess Se. This chapter then gives the results for the extra experiments described in

Chapter 4; the experiment where a large amount of Na was heated in an ampoule with a piece of pre-grown CuInSe2 (section 5.5(a)), the experiment where pre-made Na2Se was cold-mixed with CuInSe2 powder from a pre-grown ingot (section 5.5(b)), and the etching experiments where a mixture of HF:HNO3:H2O, and separately

HF:HNO3:HCl, was used to indicate the conductivity type (section 5.5(c)). Finally, the results contained in all the sections of this chapter are discussed (section 5.6).

5.2 ELECTRICAL MEASUREMENTS

This section contains the results of the electrical measurements done on the main parts of the ingots, including the hot-probe measurements on wafers, and the thermoelectric power, room temperature Hall effect, and resistivity measurements carried out on filaments. It has been observed by previous researchers that different parts of the ingots sometimes had slightly different, or very different, electrical properties. Specifically, the areas near to the first-zone-to-freeze, close to the lowest part of the ampoule that was the first to enter the lower temperature zone, located at the front of the ampoule, was seen by H. Du [5.2] to have slightly higher thermoelectric power values than the middle part of the ingot. The last-zone-to- freeze, located at the back of the ampoule, was considerably different, and showed very low, positive thermoelectric power values, indicating a p+ conductivity type

69 region, having a very high hole concentration, regardless of the conductivity type of the main part of the ingot, as seen in Tables A1-A5.

As the purpose of this study is to investigate the effects of Na, and other changes in stoichiometry, on the electrical and structural characteristics of the

CuInSe2, the measurements were done in order to determine if differences between the samples could be related to the starting composition of the growth melt, including additions of Na. Therefore, the measurements contained in this section were done on wafers and filaments cut from the middle portion of the respective ingot.

5.2(a) Hot-probe measurements on wafers

Hot-probe measurements were used to qualitatively determine the conductivity type of the main part of the ingots. Once this was done, thermoelectric power measurements were carried out on filamentary samples cut from the middle portion of each ingot, as described above, in order to quantify the conductivity measurements.

A summary of the conductivity type of all ingots grown in this study, obtained by hot-probe measurements on wafers, is given in Table 5.1. In this table, each column contains a series of samples grown from the same starting proportions of Cu,

In, and Se corresponding to the chemical formula CuInSe2+x, where x is the amount of excess Se above the stoichiometric amount. As can be seen, x is increasing from left to right, starting at x = 0 (CuInSe2), and finishing at x = 0.4 (CuInSe2.4). The rows of

70 this table are used to indicate the atomic percent of sodium included in the ampoule, along with the Cu, In and Se, before growth. The samples represented on the first row of the table, having an atomic percent of zero, were grown without Na.

As can be seen in this table, the conductivity type of the ingots grown with no excess of Se, x = 0, are p-type when no Na is added to the melt. These ingots remained p-type for small additions of Na, 0.1 and 0.2 at. %, but were n-type at higher additions of 0.3 at. % or more. This is consistent with the work of H.P. Wang

[5.1], in which ingots grown with 0.25 at. % Na or more were n-type. It can also be seen, from examination of Table 1, that as excess Se is increased, more Na is required for the type conversion from p to n to occur. This relation, first discovered in the present work, is better quantified in the results of the thermoelectric power measurements conducted on filamentary samples, given in Section 5.2(b).

While the results obtained from the filamentary samples are considered more accurate than the hot-probe measurements, they are nevertheless one measurement, and questions could arise as whether the result is characteristic of the ingot as a whole. In order to address these questions, the measurements depicted in Figure 5.1 were made. These are the results of hot-probe measurements taken at multiple points along the front and back surfaces of two wafers cut from two different ingots, HM-

B39 (p-type) and HM-B44 (n-type). As can be seen from the figures, the thermoelectric power results are largely consistent across the surfaces, indicating that the samples are relatively uniform. These results indicate that a result taken from a

71 filament can safely be considered characteristic of the main part of the ingot.

However, hot-probe also confirmed that this is not true for the last-zone-to-freeze, which is always p+-type, regardless of the starting proportions of the melt or the conductivity of the main part of the ingot, as mentioned above. Therefore, after cutting wafers from a grown ingot, the hot-probe was used to determine that the thermoelectric power of a specific wafer, or part of a wafer, was essentially equivalent to the majority of the rest of the ingot, apart from the last-zone-to-freeze.

A filament was then made from that specific wafer, thereby ensuring the thermoelectric power measurements were representative of the ingot.

5.2(b) Thermoelectric power measurements on filamentary samples

It can also be seen, from examination of Table 1, that as excess Se is increased, more Na is required for the type conversion from p to n to occur. This can also be seen in the plot of Figure 5.2, which shows the thermoelectric power (α) versus amount of Na in the melt for each sample. The α of a series of ingots having the same x (excess Se) value, and thus containing equivalent excess Se above stoichiometry, are indicated by the same symbol and are connected by a line. As x was increased, the critical amount of Na at which the ingot was seen to change type, designated as [Na]crit, was also increased. These critical Na values are seen plotted, against x, in Figure 5.3. Note that in this figure, the dashed red line represents an atomic slope of 2, corresponding to 2 atoms of Na to 1 atom of Se. As can be seen, the line crosses within the error bars for [Na]crit. Therefore, it is proposed that the p to n transition evidently involves the binary compound Na2Se, where the ratio of [Na] to

72 [Se] is 2:1. The uncertainty in the exact in the value of [Na]crit is limited by number of samples in each run.

Considering that samples grown with a deficiency of Se are n-type [5.3], theoretical work performed elsewhere [5.4] that indicates vacancies of Se (VSe) in

CuInSe2 to act as donors, and the favorability for Na and Se to react together, as pointed out by Braunger et al [5.5], it is suggested here that the change of conductivity type is caused not by the doping of the material with Na, but rather through a reaction between this element and the Se, thus preventing it from occupying its normal place in the CuInSe2 lattice. The thermoelectric power results presented here indicate, for the first time, that this reaction does indeed take place in a growth chamber containing both CuInSe2 and Na. More specifically, these results suggest that the binary formed between the Na and the Se is basically Na2Se.

Note in Figure 5.2 that the data point corresponding to ingot HM-B54

(CuInSe2.05 + 11 at. % Na) is connected to its adjacent data point by a broken line.

This is to emphasize that there is some doubt as to the validity of this point, as the ampoule was found, after removal from the furnace, to have cracked at some point during the growth. As is explained below, this cracking is due to a weakening of the quartz through a reaction with the Na. The ingot itself was found, by hot probe, to contain regions of p-type conductivity, although the majority of the ingot was n-type.

73 5.2(c) Hall effect measurements on filamentary samples

One major complication that was frequently encountered when carrying out the Hall effect measurements was voltage drift. This was not as much of a problem when taking measurements from filaments cut from n-type ingots; the high mobility of electrons, as opposed to that of holes, resulted in greater Hall voltages for these samples, which meant a higher voltage drift could be tolerated. For the p-type samples, the measured change in Hall voltage with increased magnetic field was of the order of a few micro-Volts, and so it was required that the voltage remained stable enough to detect this change. In some cases, a sample was allowed to sit for a length of time (approximately 20 minutes) after current had been applied to ensure voltage stabilization; this only worked for some samples. Another attempt to stabilize the voltage drift involved re-positioning the gold spots, so that the voltage was measured at a different location on the sample. Again, this worked for some samples.

However, the voltage reading for many samples did not stabilize enough for meaningful Hall measurements to be recorded. Therefore, the values of Hall coefficient are not given for all samples for which thermoelectric power measurements are provided. However, worth noting is a consistency in the samples for which Hall measurements were not recorded; they were all p-type samples of the

CuInSe2.2-set of data with an added concentration of sodium above 0.4 at. %. A comparison of the Hall coefficient and thermoelectric power values obtained for the stoichiometric ingots grown with Na is shown in Figure 5.4. Note that the conductivity type of the material is consistent for RH and α.

74 5.2(d) Electrical resistivity measurements on filamentary samples

The problems in cohesivity and brittleness which plagued the Hall-effect measurements also made the resistivity measurements quite difficult. This was especially true for the p-type samples grown with larger quantities of Na. Shown in

Figure 5.5 is a plot of resistivity (ρ) versus sodium addition for the stoichiometric and

CuInSe2.2 ingots. As can be seen, the resistivities of the samples with lower quantities of Na is of the order of 1 Ω∙cm, although, for ingot HM-B20 (CuInSe2.2 + 0.1 at. %

Na) the value increases to 10 Ω∙cm before dropping back down for ingot HM-B40

(CuInSe2.2 + 0.2 at. % Na). As the amount of Na in the growth melts increase, so too does ρ. However, it is unclear whether this increase is caused by the presence of microcracks, at which large voltage drops would be localized, rather than an actual change in the electrical characteristics of the bulk material.

5.2(e) Carrier concentration and mobility

It is possible to calculate the carrier concentration for each of the samples from their respective Hall coefficients, according to Eq. 3.8 and 3.9. These are shown plotted against sodium addition for the stoichiometric samples in Figure 5.6.

Additionally, the carrier concentration for the CuInSe2.2 ingots, with lower Na additions, can also be seen on this plot. As shown, there is no clear change in carrier concentration with additions of Na.

Carrier concentrations for the n- and p-type samples, as determined by analysis of thermoelectric power, differ slightly from those determined by Hall

75 coefficient. For the p-type samples,  and RH-derived carrier concentrations are of the order of 1x1018 cm-3 and 3x1017 cm-3, respectively, as can be seen in Figure 5.7.

It is believed that this difference is partly due to a discrepancy in the hole effective mass used, which factors heavily in the calculation of -derived carrier concentrations. The electron concentrations for the n-type samples were calculated to

16 15 -3 be about 1 x 10 and 6 x 10 cm for  and RH, respectively. There was no discernable trend in  or RH in correlation to the amount of Na included in the ampoule, with the exception of the change in conductivity type. However, the

2 -1 -1 mobility of p-type samples made with excess Se (CuInSe2.2), of about 20 cm V s , determined from RH and ρ, was consistently higher than that for p-type stoichiometric samples at concentrations of Na below 0.3 at. %, as can be seen in Figure 5.8. At and above 0.3 at. % Na, the stoichiometric material becomes n-type, for which Hall mobility ranged from approximately 200 to 400 cm2 V-1s-1.

A range of mobility values has also been reported for bulk crystals of between

15 and 150 cm2V-1s-1 for majority carrier holes, and between 90 and 900 cm2V-1s-1 for electrons in n-type material [3.7], as stated by Shafarman and Stolt [3.8] (see that work for thorough review of the electrical and physical properties of CuInSe2 and

CIGS).

76 5.3 STRUCTURAL AND COMPOSITIONAL INVESTIGATIONS

OF CuInSe2 AFTER GROWTH

It was also noticed that ingots grown with starting compositions containing little or no Na were relatively solid and were comprised of cm-sized monocrystals.

As the concentration of Na was increased, the ingots were seen to become brittle, with large cracks throughout the material. This increased cracking and brittleness was more apparent in the case of the p-type samples grown with large amounts of both Na and excess Se. This contributed to the already difficult task of forming filaments, and rendering the Hall and resistivity measurements in particular difficult or impossible. However, these problems were less apparent in the n-type samples containing only a large quantity of Na and little or no excess Se.

These structural changes to the material when Na is added to the growth melt, combined with the obvious change of conductivity from p to n, are clear indications that something is happening to the bulk CuInSe2. Additionally, the presence of growth deposits, discussed in Section 5.4, are further indications of reactions taking place between the constituents of the growth melts, or with the CuInSe2 in liquid or solid form, or with the quartz. Therefore, structural and compositional analysis was undergone in order to investigate whether changes in the crystal structure of the

CuInSe2 lattice accompany the changes in conductivity type, as well as to identify the deposits in further efforts to elucidate what is happening within the crystal.

77 5.3(a) X-ray diffraction on bulk material

X-ray diffraction (XRD) was performed on powdered samples from the main part of the ingots grown with stoichiometric proportions of the starting growth constituents (CuInSe2) plus increasing quantities of Na ranging from 0 to 3 at. %.

This is an extension of the previous work [5.8] in which only sample HM-B26 was analyzed. This sample, which was grown with the highest quantity of Na of any stoichiometric sample, clearly showed all the peaks corresponding to the chalcopyrite.

This confirmed that the material was indeed chalcopyrite CuInSe2, despite having undergone a conductivity type change from p to n as a result of the sodium.

However, there were traces of small unidentified peaks in the vicinity of 39.5°, 58° and 75.5°.

More recently, XRD was done on all ingots from this series (stoichiometric plus Na0). The results are given in Figure 5.9. As can be seen, all of the peaks belonging to the chalcopyrite are visible in all of the diffraction patterns, even as the material becomes n-type at 0.3 at. % Na. It is thought that slight changes in the peak shape, or differences in the relative intensities of the peaks of each sample, may be due to inconsistencies in the powdering of the material, as they were done manually, and it is possible for the effects of preferred orientation to play a role in the shape of a specific peak if there is an atypically large crystal present in the powder. This may overwhelm any effects of minor changes of stoichiometry within the samples. As well, the unidentified peaks at 39.5°, 58° and 75.5° seen in the scan of HM-B26 could also be seen in the scans of some of the other samples, particularly for HM-B24

78 (0.1% Na). Whatever the cause of the extra peaks, they do not match those of the β- phase, and no β-phase peaks were found in any of the scans (see Figure 5.10). As well, their intensities appear independent of Na concentration and conductivity type.

Therefore, it can be concluded that no change in crystal structure is apparent with increases of Na, notwithstanding the change in the electrical characteristics of the materials.

5.3(b) SEM-EDX

SEM-EDX analysis on areas of material completely in the interior of a grown ingot has indicated an absence of Na, as shown in the area analysis depicted in Figure

5.11 (a). However, a fair amount of this element has been found on analysis of the exterior surface of this ingot, as shown in Figure 5.11 (b). This is the first time that

Na, having been included in the melt prior to compound synthesis, was found to have been segregated from the material and deposited on the surface of a CIS crystal.

It can also be seen in Figure 5.11 (a) that the grown material within the bulk is

Cu-deficient ([Cu]/[In] < 1) and Se-rich ([Se]/{[Cu]+[In]} > 1). This was found to be the case regardless of starting composition of the melt, and regardless of conductivity type, as can be seen in Figure 5.12 (a), which depicts a similar analysis done on an n- type sample. A similar phenomenon was observed by Ziad Shukri [5.9], in which the grown material tended towards an equilibrium composition, even as the composition of the melt was changed. However, the analysis done here indicates that the samples are in the range of atomic compositions where mixed phases, α + β, could be present.

79 If these extra phases are in fact present in the material, they are there in concentrations too small to be detectable by XRD.

The last-zone-to-freeze of our material, in contrast to the main part of the ingot, was Cu-rich and slightly Se-deficient, as indicated in Figure 5.12 (b).

Additionally, a small amount of Na was detected there. Far from an intact, defect- free crystal, the last-zone-to-freeze contains many cracks and is often very brittle

[5.10]. With this in mind, it is suggested that the Na detected in this region is residing on cracks, and other defects, rather than within the crystal. Therefore, the appearance of Na in this region, rather than the bulk material, is further evidence of Na- segregation from the melt during growth. The phase diagram would suggest that

Cu2Se exists as a separate phase. However, this was not detected by XRD. Further, it was not detected by Shukri [5.11] using EPMA in the last-zone-to-freeze. It should also be noted that precise compositional measurements require a standard sample, and since these measurements were analyzed without the use of a standard, their accuracy could be compromised.

Figure 5.13 shows concentration determinations across a 3 micron-wide imperfection on the surface of wafer cut from ingot HM-B31, grown with large excess of Se, corresponding to the chemical formula CuInSe2.4, and a relatively large concentration of Na (10 at. %). An abrupt increase in Na, accompanied with an increase in Se and a decrease of In and Cu is noted at this region (note that the solid lines of the SEM-EDX data are used only to connect adjacent data points, and do not

80 represent a trend of the atomic concentration of the respective element between the points). This figure is further evidence for Na to collect only on grain boundaries and in defects, but not within the intact crystal. Additionally, it is noticed that composition of the bulk material is similar to that seen in the area-analysis measurements described above, i.e. slightly Cu-deficient and slightly Se-rich. While the increase of Na and Se at the imperfection is accompanied by decreases of Cu and

In, the value of [Cu] / [In] was seen to increase above unity at that point.

The presence of additional phases, aside that from the chalcopyrite CuInSe2, was noted in SEM images, taken using a backscattering detector, on a polished sample from the last-zone-to-freeze (LZTF) of ingot HM-B70 (selenium deficient), as shown in Figure 5.14. While these extra phases were not identified, it is presumed that they include Cu7In4, Cu9In4 and InSe, as were identified in the LZTF of a Se- deficient sample by Shukri [5.11].

5.4 CHARACTERIZATION OF DEPOSITS

This section describes the differences seen in ampoules grown with varying charge compositions. A visual inspection of the ampoules, after their removal from the Bridgman furnace, revealed striking differences in appearance corresponding to increases in both sodium and excess selenium content. A white deposit, which formed on the inner walls of ampoules that contained sodium, was seen

81 predominately when little excess Se was also included, resulting in n-type material.

As described, SEM-EDX analysis of this deposit determined it to be comprised of mostly silicon and oxygen, presumably from the quartz, as well as some Na, while

XRD found only peaks corresponding to crystalline SiO2. Less of this deposit was seen to have been formed when the amount of excess Se above stoichiometry was increased, resulting in p-type material. In this case, another deposit, comprising of a reddish-brown layer of material between the exterior of the ingot and the interior of the quartz, was noticed. SEM-EDX measurements were taken in order to determine the atomic compositions of this deposit as well.

5.4(a) Red deposit

Shown in Figure 5.15 are two ampoules, pictured immediately after being removed from the Bridgman-furnace. As can be seen, various deposits were found within the ampoules after growth, and the presence and proportions of these deposits were related to the starting composition of the melts. In ampoules containing p-type material (see Figure 5.15 (b)), grown with Na and excess Se, a reddish-brown deposit was found between the exterior surface of the ingot and the interior wall of the ampoule. Analysis, by SEM-EDX of the interior surface of an ampoule, where this deposit was seen, indicated significant quantities of Na and Se, in the approximate proportion of 3 to 1, respectively, as well as traces of the other starting materials, as shown in Figure 5.16. The suggestion here is that Na is deposited on the surface of the ingot in the form of a selenide. However, XRD analysis of this deposit failed to uncover any Na compounds within the scans. It is possible for Na compounds to

82 exist in this deposit in an amorphous form, thereby rendering them undetectable by

XRD, or that they are found in too small concentrations after growth. There is also some evidence for Cu-Se phases to be present in this compound; Cu3Se2 has been found previously in deposits taken from ampoules which contained p-type samples that had grown from melts containing Na2Se [5.8].

5.4(b) White deposit

In the case of ampoules grown with Na and little or no excess Se, a white deposit was present, as shown in Figure 5.15 (a), the quantity of which was increased in samples grown with greater additions of Na. EDX analysis of this deposit showed it to contain Na at higher concentrations than the other elements (Cu, In or Se) as well as Si and O, presumably from the quartz. More recent attempts to determine the structural composition of this material were made using XRD, which uncovered crystalline SiO2 in the diffraction pattern, shown in Figure 5.17 (a). XRD on an area of quartz where the white deposit had been scratched away resulted in a pattern indicative of an amorphous material. The ampoule itself is made from fused quartz, and so it was a surprise to find unmistakably a crystalline form of this compound.

However, it had been suspected that Na weakens the quartz, causing it to crack when high concentrations are added to the ampoules. The evidence collected from XRD, as well as the knowledge that this deposit only appears when Na is included in the melts, leads to the conclusion that this element causes the quartz to nucleate, which explains the weakening as well as the white colour of the deposit.

83 5.4(c) Copper nodules

Another type of deposit, shown Figure 5.18, was found only in ampoules into which Na had been included and when the resulting ingot material was n-type, with the exception of one sample grown without Na, described below. SEM-EDX analysis determined the copper-coloured precipitates, which in this work have been referred to as nodules, to be very copper rich, as is shown in Figure 5.19. Here, EDX indicated the area around the nodules to be 75 at. % copper, and the globule on which it appeared could be described as CuInSe2 with excess Cu and a deficiency of In. Note that the nodule area where the scan was taken is not uniform; scans focusing on the nodules revealed them to be 95% Cu.

Copper nodules were found in all samples in which Na had been included and the resulting bulk material was n-type. It is thought that the amount of elemental Cu found in these ampoules increased with Na, however the nodules were never quantified and therefore this is merely a general perception rather than a statement of fact. Aside from the ampoules grown with Na, precipitates of what appear to be Cu were found on a piece of material taken from ingot HM-B70, grown with no Na but with a deficiency of selenium such that the starting melt composition was CuInSe1.7.

This deposit can be seen in Figure 5.20. While it appears in a different form than the nodules, seen in Figure 5.18, it nonetheless is another similarity shared with the ingots grown with sufficient Na to be made n-type. Therefore, the appearance of the elemental Cu in samples grown with Na is consistent with what occurs in samples

84 deficient in Se. This is further evidence for an interaction between the Na and the Se to occur when Na is added to the melt, resulting in Se-deficient material.

5.5 RESULTS OF EXTRA EXPERIMENTS

The main advancement of knowledge of this thesis is the discovery and characterization of the change of conductivity type with addition of Na. Analysis of the relationship between Na, excess Se, and conductivity type has led to the development of a model to describe the action of the Na and Se on the CuInSe2 crystal, which will be given in Chapter 7. The measurements and experiments described in this section (Section 5.5) are meant as a support to the thermoelectric power measurements, described above, and to aid in the development of the model used to explain the results of those measurements.

5.5(a) Heating experiments

The XRD results, described above in section 5.3(a), in which no Na, Na2Se, or any phases were detected apart from the chalcopyrite in the bulk region of stoichiometric ingots, grown with up to 3 at. % Na, is in some ways inconsistent with the thermoelectric power results described in section 5.2(b). Those results point clearly at a 2:1 relation between the Na and excess Se, respectively, and the conductivity type of the ingots, leading to the conclusion that Na2Se is formed in the ampoule at some stage during growth. Additionally, as described below in section

85 5.5(b), Na2Se was detected by XRD in a mixture of pre-grown stoichiometric

CuInSe2, but even in an unheated 1:1 molar mixture of Na2Se to CuInSe2, the peaks corresponding to the binary were quite weak. Therefore, a solid piece of CuInSe2 was heated modestly in an evacuated ampoule with a large amount of elemental Na (2:1 molar ratio of Na to CuInSe2, respectively), which would have cracked the ampoule in a normal growth run at 1050 °C. The resulting material was mostly brittle, as shown in Figure 5.21, thus confirming the effect of Na in increasing the brittleness of ingots grown with increasing quantities of this element, described in section 5.3.

A part of the brittle material was powdered manually using a mortar and pestle, and analyzed by XRD. The resulting diffraction pattern is shown in Figure

5.22. As can be seen, lines of Na2Se are present in the powder, as well as the hydrate

Na2Se(H2O)5. This is significant, as it is the first time selenium, originally part of the

CuInSe2 lattice, was shown to be removed from the crystal in order to react with Na in a 2:1 atomic ratio. Additionally, the originally p-type piece of ingot HM-B50 was found, by hot-probe, to have become n-type after the heating with Na. An orange- coloured deposit seen on the piece after the heating, shown in Figure 5.23, was confirmed by XRD to contain elemental Cu, thereby reinforcing the results of the growth experiments in which elemental Cu was also found in n-type samples grown with Na, as well as in an n-type sample grown with an intentional deficiency of Se.

The strong affinity of Na and Se was also evident in another heating experiment in which only these two elements were heated together. The resulting

86 material, found after the heating, was also analyzed by XRD and was found to contain compounds of Na and Se in a 2:1 atomic ratio, respectively, as shown in Figure 5.24.

Therefore, it is clear that compounds containing Na and Se are formed preferentially to those containing Na with Cu and In, and that they form in binaries, and associated hydrates, containing 2 atoms of Na for every one atom of Se.

5.5(b) Mixing experiments

As described in Section 4.4(b) of this thesis, the mixing experiments were conducted in order to determine whether XRD could be used to detect a Na2Se-phase in a powder containing predominately CuInSe2, obtained from ingot HM-B50. As a precursor to this experiment, the original Na2Se, as supplied by the manufacturer, was analyzed. The results, shown in Figure 5.25, indicate that Na2Se alone is indeed visible to XRD.

A total of three powders were made, one containing 10 at. % Na2Se in

CuInSe2, and two with 100 at. % Na2Se in CuInSe2. As mentioned in Chapter 4, while conducting this experiment, the intention was to perform the XRD analysis as soon as possible after mixing in order to minimize any reaction that might take place between the Na-compound and the air. In the case of the 10 at. % mixture, this was achieved, and the sample was analyzed within 10 minutes of the Na2Se being removed from the sealed bottle. However, there was a delay of four hours between the preparation and analysis of the first 100 at. % Na2Se to CuInSe2 mixture prepared.

For this reason, a second 100 % mixture was prepared.

87 The XRD scan of the 10 at. % mixture is shown in Figure 5.26. As can be seen, perhaps surprisingly, no peaks from the extra phases are present, and only the chalcopyrite lines from the CuInSe2 are visible, with the exception of a small peak at approximately 29.5°, identified as the (101) peak of elemental selenium. It was this result that prompted the experiment to be repeated with much more Na2Se. The first time this was done, the scan completed after the delay also revealed no Na2Se, as can be seen in Figure 5.27. However, Na2(SeO3), having a 2:1 [Na]:[Se] ratio, as well as elemental Se, was identified. It was speculated that the time delay may have caused the Na2Se to react with air and change form. Therefore, the experiment was repeated, only this time the new 100 at. % Na2Se to CuInSe2 mixture was analyzed within 10 minutes of preparation. As shown in Figure 5.28, this new sample reveals clear, though weak, Na2Se peaks, along with those of the chalcopyrite. Peaks belonging to the selenite phase can also be seen.

With these results, one might initially assume that the elemental Se was somehow dissociated from the Na2Se during oxidation. However, as both sodium selenide and sodium selenite have equivalent Na to Se ratios, and that all phases in the mixture were identified, it is more likely that the elemental Se was somehow dissociated from the CuInSe2 as a result of mixing with the Na2Se.

5.5(c) Etching experiments

The sign change is also reflected in chemical etching behaviour. Shown in

Figure 5.29 are two pieces of Bridgman-grown material before and after abrasive

88 polishing and etching in room temperature HF:HNO3:HCl (1:1:1), first used in the present work. It was found that the etch reacted with a p-type wafer of CuInSe2 to form a red deposit, after soaking in the mixture for approximately five minutes, but this did not occur with an n-type wafer. As can be seen in the figure, a reddish-brown deposit is apparent on the p-type sample from HM-B39, which was grown from a melt containing CuInSe2 with 0.2 at. % Na. The deposit appears more predominately in the center of the sample, as compared to the n-type wafer cut from ingot HM-B44

(CuInSe2 + 0.3 at. % Na). However, some red deposit can still be seen around the side near the top and right of the n-type sample. Nevertheless, the red deposit on the p-type sample is much more apparent.

Ahmad [5.12] in his work on etching, described the reaction between CuInSe2 and HF:HNO3:H2O as “violent,” and stated that a 30-second etch time was enough for a clear deposit to be visible. However, when this etchant was used in the present work, no visible reaction took place as the wafers were placed in the beaker containing the acid mixture, and that no deposit appeared to form on either wafer at the 30-second mark. The wafers were therefore left to soak in the acid for 5 minutes, but still no reaction or deposit became apparent. Note that the wafers used by Ahmad were cut from ingots grown without Na. The role that Na, added to ampoules before growth, plays in this reaction is unclear. However, with the etchant contained HCl, it was noted that the small amount of red deposit on the wafer of HM-B44 appeared to be localized to individual grains. As can be seen from the before-and-after pictures in

Figure 5.29, the sides at the top and right of the wafer are part of a different crystal

89 than the middle, which also means that different crystal planes were exposed to the acid. This is in contrast to the p-type wafer of HM-B39, which is largely from one crystal. Of course, the role of crystal plane orientation on the reaction between

CuInSe2 and the etchants is also unclear.

5.6 DISCUSSION

It is possible to use the XRD results of the cold mixing experiment, described in section 5.5(a), to quantify the detection limit of Na2Se in the scans of Figure 5.9. It is a fundamental principal of XRD that the normalized area under the Bragg peak, belonging to a specific phase, is proportional to the volume fraction of that phase within the greater material, assuming the scans were made using identical spectrometer settings and times. In the scans of Figure 5.9, the counts were taken over a period of 300s. Now, assuming equal grain sizes in each powder, the ratio of the height of a certain peak in two different scans would give an indication of the ratio of the phase present in the material between the scans. As can be seen in Figure

5.22, the (111) peak of Na2Se, which occurs at a 2θ-value of approximately 22.6°, is roughly 12000 counts high, although this number represents a scale-up factor of 100.

No corresponding peak is present in any of the scans of powdered bulk material grown in the presence of Na in Figure 5.9. In that figure, the background noise for any scan is of the order of 10 counts, meaning any peak must be greater than 10 counts high in order to be visible. Therefore, the ratio of the background noise in

90 Figure 5.9 to the peak height of a phase in Figure 5.22 should be equal to the ratio of the detection limit of that phase in Figure 5.9 versus the amount in the beam in Figure

5.22. Assuming a volume fraction of 25% Na2Se in Figure 5.22, the detection limit of Na2Se in the bulk material is 4 at. %. Therefore, if Na2Se is present in the bulk material of the stoichiometric ingots grown with Na, it is estimated to be less than 4 at. % of the material in the beam. Considering the hydrate Na2Se(H2O)5, of which the

(111) peak height in Figure 5.22 is roughly 6000 counts, also scaled up by 100, the detection limit in the scans of Figure 5.9 becomes 2 at. %. Note that these are merely order of magnitude estimates. The volume fraction of 25 at. % for the Na compounds in Figure 5.22 is purely a guess, and the actual number is probably somewhere between 10% and 60%. As well, if the grain sizes of the Na compounds are much smaller in the samples of Figure 5.9 than those of Figure 5.22, the detection limit of those compounds in Figure 5.9 could be higher. Of course, the cold-powder mixture of 10 % Na2Se in CuInSe2 did not reveal any Na2Se diffraction peaks, as seen in

Figure 5.26, which would indicate that the detecton limit may be higher than that estimated in this discussion.

Figure 5.1 Back and front of slices cut from ingots grown with (a) 0.2 and (b) 0.3 at. % Na. Also shown are the α-values found by hot probe at the locations indicated on the slices (note the polarities).

Figure 5.2 Thermoelectric power versus sodium addition of ingots grown with a starting composition of CuInSe2+x; positive values, in red, indicate p-type material, while negative values, in blue, indicate n-type material.

Figure 5.3 Critical amount of Na required for the conductivity type to change from p to n versus x, the amount of excess Se present in the ampoules before growth. The dashed red line is used to mark the Na:Se molar ratio of 2:1, corresponding to the formula Na2Se.

Figure 5.4 Hall coefficient (RH) and thermoelectric power (α) measured on filaments obtained from ingots grown from melts containing copper, indium and selenium in the atomic proportions of 1:1:2, plus varying amounts of added sodium. An abrupt change in conductivity type was seen for Na additions of 0.3 at. % and above.

Figure 5.5 Electrical resistivity measured on filamentary samples plotted against added sodium content. The red symbols indicate p-type conductivity and the black symbols n-type conductivity.

Figure 5.6 Carrier concentrations, determined by room-temperature Hall-effect measurements, on stoichiometric samples grown with Na (p-type samples in red and n-type samples in black), as well as those on samples made from a melt composition of CuInSe2.2 (p-type, in blue).

Figure 5.7 Majority carrier concentration, calculated from both thermoelectric power and Hall measurements, for CuInSe2.2 versus atomic percent of sodium present during growth.

Figure 5.8 Hall-derived carrier mobility versus sodium addition for CuInSe2 (x = 0) and CuInSe2.2 (x = 0.2), showing the excess Se samples to be consistently higher at low sodium additions.

Figure 5.10 Semilog plot of the XRD of sample HM-B26 (stoichiometric + 3% Na) in the critical region where β-phase peaks (marked by black arrows) would appear near the chalcopyrite peaks (marked by red lines). No β-phase peaks are discernable above the background noise of the scan.

Figure 5.11 Compositional EDX analysis of the interior surface (a) and exterior (b) of an ingot grown from a melt with a starting composition of CuInSe2.05 (x = 0.05) with 5 at. % Na.

Figure 5.12 Compositional EDX analysis of the main part (a) and last zone to freeze (b) of an ingot grown from a stoichiometric melt with 0.3 at. % Na.

100

Figure 5.13 SEM image of an imperfection on a polished wafer of an ingot grown from a melt with the composition CuInSe2.4 + 10 at. % Na. EDX measurements below indicate Na at the imperfection, as well as increased Se but decreased Cu and In.

101

Figure 5.14 SEM images, taken using a backscattering detector, on the polished last- zone-to-freeze of ingot HM-B70 (CuInSe1.7, Se-deficient). The presence of several phases of material is apparent in the images, possibly including Cu7In4,Cu9In4 and InSe [5.11].

102

Figure 5.15 Ampoules shown after the Bridgman-growth process, both containing 3 at. % Na; (a) an n-type sample, grown with less excess Se, CuInSe2.005, showing a large amount of white deposit. (b) A p-type sample, grown from a melt with composition CuInSe2.05, showing some reddish-brown deposit, adjacent to the ingot (left of picture), and only some white deposit.

Figure 5.16 SEM image of the red-brown deposit on the inner surface of quartz from an ampoule containing CuInSe2.05 + 5 at. % Na after growth. Shown also is the elemental composition of the material within the white square, as determined from EDX analysis.

103

Figure 5.17 (a) XRD of the white deposit seen in an ampoule in which Na was included, indicating peaks of crystalline SiO2; (b) XRD on the ampoule after the white deposit was scratched off, indicating the material under the white deposit to be amorphous.

104

Figure 5.18 Nodules of copper-coloured material found on dots of material adhered to the sides of an ampoule containing Cu, In, Se and Na after growth. These deposits, later identified as elemental Cu, were seen only when Na was included in the melt and the resulting ingot was n-type, as in this case.

Figure 5.19 SEM image of a globule of material, with copper-coloured precipitate (left) taken from the inside of an ampoule containing CuInSe2.005 + 3 at. % Na after growth. Shown also is the average elemental composition of the globule and precipitate (areas inside respective red rectangles), as determined by EDX. On right is a photograph of the globule, taken through a microscope.

Figure 5.20 A piece of material that had broken off an ingot grown with a starting composition of CuInSe1.7 (Se-deficient), showing what appears to be precipitated copper; this has yet to be confirmed by compositional analysis.

105

Figure 5.21 Brittle material seen in an ampoule originally containing a solid piece of 0 CuInSe2 and a large piece of Na , after heating to 300 °C for 1 hour.

Figure 5.22 XRD of material resulting from heating p-type CuInSe2 with Na, in a molecular ratio of 1:2, to 300 °C, where the peaks from CuInSe2 have been removed, along with the noise below these peaks, revealing only peaks from extra phases. The , blue and red lines represent the diffraction patterns of Na2Se and Na2Se(H2O)5 respectively, while some peaks remain unidentified.

106

Figure 5.23 XRD on a solid piece of material taken from the ampoule of Fig. 5.21, originally containing CuInSe2 and elemental Na in a 1:2 atomic ratio, after heating to 300 °C, indicating Cu, Cu7In3, and possibly Na2(SeO3), along with some CuInSe2 peaks. The beam was aimed at the orange part on the surface of the sample. Note that the sample was not rotated during the scan, and so the effects of preferential orientation may be apparent in the absence of some peaks, notably the (112) peak of CuInSe2.

Figure 5.24 XRD of material resulting from heating Se with Na to 400 °C, indicating compounds primarily consisting of 2 Na atoms with 1 Se atom.

107

Figure 5.25 XRD pattern of commercially-fabricated, unheated Na2Se, from Alfa Aesar, of nominally 99.8% purity. As can be seen, all the standard Na2Se peaks, shown in blue, are present in the experimental sample. The shifting is likely due to stress in the material.

Figure 5.26 XRD pattern of unheated, powdered CuInSe2, from ingot HM-B50, mixed with Na2Se in a 10:1 atomic ratio, respectively. The pink lines represent the standard CuInSe2 peaks. Note that no peaks from Na2Se are present, while a peak just before 30° was identified as Se0.

108

Figure 5.27 XRD pattern of unheated, powdered CuInSe2 mixed with Na2Se in a 1:1 atomic ratio, after having sat on a counter at room temperature for 4 hours. The pink lines indicate the standard peaks of CuInSe2. Some peaks identified as Na2(SeO3) are marked with black arrows, and the large peak just before 30° is from elemental Se. The blue lines show the standard peaks of Na2Se, which is clearly absent from the material.

Figure 5.28 XRD pattern of unheated, powdered CuInSe2 with Na2Se in a 1:1 atomic ratio, this time taken 10 minutes after the preparation of the powders. Now, the peaks of Na2Se, marked in blue, are clearly present.

109

Figure 5.29 Photographs of p-type (left) and n-type (right) sample before (top) and after a 5-min etch in HF:HCl:HNO3. A red deposit is more prominent on the p-type sample after the etching.

110 Chapter 6

Photovoltaic Cells made from mono-CIS

6.1 INTRODUCTION

The primary objective of fabricating photovoltaic cells from the ingots grown over the course of this work was to establish a routine for fabricating monocrystalline cells with consistent, reproducible characteristics above around 7 % efficiency.

Ultimately, it was hoped that such a process would allow for comparisons between devices made from ingots grown with and without sodium.

In total, more than 50 cells were fabricated, but the performance of these cells varied greatly. As is explained in this chapter, the source of some of these discrepancies in performance was determined, which gained insight into the loss mechanisms of the cells, thus allowing the higher performing devices to be made.

However, the performance and characteristics of devices made following the same procedures, and even, in some cases, in the same batch, remained inconsistent. The highest-performing cell made in this work, called cell 62-1, had an energy conversion efficiency of 8.8 %, achieved with no anti-reflection coating, and had a total area of

16 mm2. In comparison, the highest-performing single crystal cells made in this laboratory by previous workers L.S. Yip [6.1], Z.A. Shukri [6.2], H. Wang [6.3] and

H. Du [6.4] had efficiencies of 11.5 % (4 mm2 active area), 6.3 % (5.3 mm2 active

111 area), 8.1 % (active area not stated) and 12.5 % (12.6 mm2 active area), respectively, the last result being the highest ever reported for a monocrystalline CuInSe2 device.

However, these workers also noted inconsistencies in the performance of their cells, the sources of which were never fully explained. The performance of a select few cells made in the present work, along with the fabrication procedure, is given in

Appendix D. The majority of the other approximately 45 cells had efficiencies below or very close to 1 %, including all of those made from wafers grown with sodium. As these values are too low for meaningful comparative measurements to be made, the effect of Na on the performance and characteristics of the cells cannot be determined.

Nevertheless, valuable experience and insight was gained into the functionality, and performance losses, of photovoltaic cells made from monocrystalline CuInSe2.

Section 6.2 of this chapter discusses the fabrication and testing procedure.

Following this, the performance of a few cells are reviewed and discussed in Section

6.3. The chapter then concludes with a discussion of the determination of diffusion length made on cell 62-1 by photocurrent-capacitance measurements.

6.2 CELL FABRICATION PROCEDURE

Shown in Figure 6.1 (a) is a cross-sectional drawing of the cell constructed in this work, having a structure of Au - p-CIS – n-CdS – ZnO – In. Cells were fabricated by first cutting a 3 mm thick wafer of CuInSe2 using a diamond saw. The

112 hot-probe was used to confirm the p-type conductivity of the wafer. One side of the wafer was then manually lapped with 600-c, 800-c, and 1200-c grit paper for 5 minutes each, followed by polishing with a suspension of 0.05 µm alumina powder in de-ionized water (DIW) for 2 hours. Powders with different diameters, ranging from

10 µm to 0.05 µm, were tested, but no effect on cell performance was observed. The wafer was then cleaned successively with toluene in an ultrasonic bath for 8 minutes, followed by acetone and DIW. In most cases, this cleaning was followed by an anneal at 350 °C in argon for 2 hours, followed by CdS deposition by chemical bath.

However, in the case of cell 62-1, the annealing step was skipped, and CdS was deposited immediately after the cleaning. A circular active region 4.5 mm in diameter was then defined, onto which ZnO was deposited by sputtering. Top and bottom contacts of In and Au, respectively, were then deposited by evaporation.

Finally, the cell was mounted onto an aluminum stud using silver epoxy, and a thin copper wire was attached to the indium using Wood’s metal. A finished cell can be seen in Figure 6.1 (b). For a detailed description of the exact process followed for each individual cell, please refer to Table D of Appendix D.

6.2 (a) Wafer cutting, polishing, cleaning and surface etching

The first step in the fabrication of photovoltaic cells in this work was the choice of ingot from which the absorber layer would be made. There were several requirements, the first one being that the ingot had to be uniformly p-type. It was noticed by previous researchers [6.5], as well as by the author in the precursor to this work [6.6], that some ingots had small n-type regions in what was otherwise p-type

113 material. Furthermore, the ingot had to be relatively solid in order to withstand the fabrication process, and be free of cracks over a region large enough to define an active area. As all of the steps were done manually, a wafer must have a defect-free area of 1 cm2 and be 2 mm thick, before polishing.

Therefore, an ingot was chosen from a melt composition known to result in p- type material and appeared solid, and was then cut into 3 mm-thick wafers using a diamond saw. This thickness was necessary to allow the wafer to be handled without breaking, as subsequent polishing steps remove approximately 1 mm of material.

Following this, the surface of the wafer was abraded using 600-c grit sandpaper, with

DIW as a lubricant. After this step in the polishing, the surface of the wafer was tested using the hot probe to ensure it was entirely p-type. The wafer was then abraded using 800-c grit sandpaper, and then 1200-c grit sandpaper, again with DIW.

Finally, the surface was finely polished with 0.05 µm alumina powder in DIW for between one and two hours.

The wafers were then cleaned by emersion in toluene in an ultrasonic bath for

48 minutes, with the toluene replaced every eight minutes (TCE was used in previous work on cells in this laboratory). This was followed by acetone, again in the ultrasonic bath for eight minutes, and finally eight more minutes in DIW.

Previous workers had experimented with various etchants, including HCl, aqua regia and, slightly more recently, bromine methanol. As the cell with the

114 highest conversion efficiency had undergone an etching in bromine methanol (BM) prior to the annealing step, this etch was the one used in this work. A 0.5% v/v BM solution was prepared after the cleaning, and the wafer was immersed in the solution for 30 seconds; a longer etch was seen to result in etch-pits. After the 30 seconds had passed, the wafer was soaked in DIW to stop the etching process, and then air-dried.

6.2 (b) Wafer annealing

Previous work in this lab performed by L.S. Yip [6.1] has shown annealing to be an important step in achieving higher-efficiency solar cells. In that work, the effects of annealing at different temperatures for different times and with different gases was explored, including argon, hydrogen and air, as well as in a vacuum. It was found that the condition resulting in the best cell performance was 350 °C for 2 hours in argon.

Annealing took place immediately after polishing and etching so as to minimize oxidation of the surface. The wafers were first loaded into a Pyrex tube about 6 cm in diameter, which was flooded with argon gas. After 20 minutes, a controller was turned on that passed electricity through a heating coil wrapped around the center part of the tube. The controller was set to maintain a temperature of 350 °C inside the tube. It took about 15 minutes for the furnace to reach the desired temperature, at which point the annealing was considered to have begun. Exactly two hours later, the controller was shut off, allowing the system to return to room temperature, which took about 30 minutes, after which the wafers were removed. A

115 constant flow of argon gas was maintained throughout the entire process. Note that cell 62-1, the best cell made in this study, did not undergo an annealing process. In the case of this cell, the annealing step was skipped, and the CdS deposition was carried out immediately after the etching.

6.2 (c) Cadmium sulfide layer formation by chemical bath deposition

Cadmium sulfide (CdS) was used as the n-type buffer layer between the p- type CuInSe2 and transparent conductive oxide layer, described in Section 62(e). The

CdS layer was applied by chemical bath deposition (CBD) immediately after the annealing step so as to avoid any effects of oxidizing of the absorber surface.

Beforehand, two solutions were made and kept in separate containers. The first solution, called solution A, was comprised of CdCl2, NH4Cl and (NH2)2CS (thiourea, as the source of sulfide ions) in the proportions of 2 mM, 20 mM and 20 mM, respectively. The second solution, called Solution B, was 200 mM NaOH, used to control the pH. During the deposition process, equal parts of each solution (usually

25 ml) were poured into a beaker, which was then placed in a warm bath of DIW within a larger container on a hot plate. When mixed together at room temperature, the resulting solution was completely clear. After one minute, the substrate was placed into the beaker containing the salts mixture, and allowed to sit for 12 minutes, during which time the liquid was seen to turn yellow and develop a milky consistency. The samples were then removed and immediately rinsed with DIW. In the case of cell 60-1, this procedure was performed only once, but for the other cells it was repeated twice. It had been found by Wu [6.7], while studying the properties of

116 CdS for use in electronic resonators, that repetitive depositions were necessary in order to increase the thickness of the layer; increasing the deposition time resulted in rough, uneven surfaces, which would be undesirable for photovoltaic applications.

Using this process, the layer of CdS was seen to be approximately 40 nm per deposition, as verified by SEM in Figure 6.2 (a), with a resistivity of approximately

1.5 x 105 Ω∙cm, as measured by four-point-probe. This is similar to the thickness and resistivity measured on CdS by previous workers in this laboratory [6.1, 6.5].

However, the layer was also seen by SEM to be uneven, with large clumps of material seen scattered around the surface, as shown in Figure 6.2 (b).

In this work, some time was spent investigating the effects of the deposition parameters on the film properties, and specifically the results of increasing the deposition temperature. Shown in Figure 6.3 are four glass slides, each with five layers of CdS deposited at different temperatures. As can be seen, as the deposition temperature is increased, the colour of the CdS becomes darker, and eventually appears non-uniform and patchy at 80 °C. The CdS layer of the best cell in this work, cell 62-1, was deposited at 70 °C, however more work is needed in order to isolate the effects of the CdS layer on cell performance and to determine the optimal deposition conditions.

6.2 (d) Back contact deposition

Gold has been traditionally used in this laboratory as the back-contact material for CuInSe2 solar cells, as it is easy to deposit by evaporation and, being a high-

117 work-function material, forms an ohmic junction with p-type material. An Edwards model E12-E3 evaporation chamber was used to deposit the gold to the back of the

CuInSe2 wafer. Approximately 2 cm of gold cut from a wire of 99.8% purity, manufactured by Alfa Aesar, was placed in a molybdenum boat. The wafer was affixed approximately 5 cm above the boat. The chamber was then evacuated to a pressure of 10-7 torr, at which time 50 A of current was passed through the boat, which caused the molybdenum to heat up and the gold to liquefy. After about 30 seconds, a mask between the boat and substrate was removed, and the current was gradually increased to 80 A for 2 minutes, at which point it was shut off. The system was then allowed to cool for 1 hour before the chamber was flooded with air, thus returning to atmospheric pressure, and the sample removed.

Early cells fabricated in this work were seen to have low fill factors, which was attributed to high bottom-contact resistance, resulting in high series resistance.

In order to study the resistance of the gold contact, two wafers were taken from a stoichiometric ingot (HM-B36) and manually abraded with 600-c grit sandpaper.

One of the wafers was then annealed at 350 °C in argon for two hours. Two gold dots about 1 mm in diameter were then deposited on the same surface of each wafer. The resistance measured between the dots was considerably higher for the annealed wafer than the un-annealed wafer. The surface of the annealed wafer was then re-abraded and the dots re-deposited, this time showing a low resistance, ohmic profile similar to the second wafer. This experiment, visualized in Figure 6.4, showed the necessity of

118 abrading the bottom of the wafer before metallization. As such, this was done for all of the cells made after this experiment, including those listed in Appendix D.

6.2 (e) ZnO deposition by sputtering

A low-resistive zinc oxide (ZnO) layer served as the transparent conductive oxide (TCO) for the cells made in this work. The ZnO was deposited by rf magnetron sputtering from a target containing 2 % Al2O3, which served to dope the layer with aluminum in order to decrease the sheet resistance. A metal mask was used to confine the ZnO to a defect-free region on the wafer, and thereby defining the cell area. The sputtering was done using an Edwards E306A coating system, which was first evacuated to a pressure of 10-6 torr before argon gas was pumped in until the pressure was raised to around 0.2 torr, at which point the rf power supply was switched on, forming a plasma. The pressure within the chamber was then adjusted to 0.06 torr and maintained for 8 hours, resulting in a layer of ZnO:Al approximately

700 nm thick. This step was carried out each time by Dr. Ishiang Shih.

Note that only one layer of conducting ZnO, of resistivity approximating 7 x

10-3 Ω·cm [6.5], was used in the fabrication of the cells in this work. Du had previously experimented with a double layer of ZnO: a highly-resistive layer adjacent to the CdS, followed by a low-resistance layer, as is typically used in the fabrication of polycrystalline, thin film CIGS devices. However, he observed no benefit in performance over a single, Al-doped layer [6.5].

119 6.2(f) Indium top contact deposition and cell mounting

Indium was used as the top-contact material for the cells made in the present work, and was deposited after the deposition of the ZnO. This was done using the same metallization process described for the back contact in Section 6.2 (d) of this work, except that the lower melting point of In meant the material liquefied at about

40 A of current through the molybdenum boat and vaporized at about 60 A. A mask was also used to define the contacts. Different mask shapes were sometimes used, but, as the performance of the devices made in this work was largely inconsistent, as described in Section 6.1, the influence of the top contact pattern on cell performance could not be determined. Note that the highest cell in this work, cell 62-1, was made using a dot of In, rather than two stripes, but that subsequent cells also made with a dot of In failed to achieve the same conversion efficiencies as this cell.

After the top-contact deposition, the cell was mounted onto an aluminum stud using silver epoxy, which was followed by heating at 80 °C in air for 10 minutes to cause the epoxy to set. Wood’s metal was used to attach fine copper wires to the indium top contacts. At this point, the cell was ready for testing.

6.2(g) Testing

The illuminated IV characteristics of the cells discussed in this work were measured on the roof of the McConnell Engineering building, located at the McGill

University downtown campus in Montreal, in direct sunlight on November 11, 2010, at around noon, except cell 62-1, which was measured on November 24, 2010. The

120 solar irradiance on both these days was 96 mW/cm2, measured using an Eppley

Precision Spectral Pyranometer model PSP with a model 3478A multimeter manufactured by Hewlett Packard. During testing, the devices were positioned at an angle so that the plane of the cell was perpendicular to the rays of the sun, and the open-circuit voltage was at a maximum. A Keithley programmable current source was then used to vary the current through the device as the voltage was read using the multimeter.

In one case, data is reported on a cell which was measured using a xenon arc lamp solar simulator, LH 150/1 manufactured by Kratos, as an illumination source.

In this case, the filtered lamp was calibrated using a silicon reference cell, manufactured by Solarex Corporation, for which an open-circuit voltage of 3.1 mV was generated under AM1.5G conditions. A comparison of the solar simulator to the actual sun can be seen in Figure 6.5 for an early cell made in this work. As indicated by the figure, the performance of the cell measured under the xenon lamp is a slight underestimate of the performance under actual sunlight, although this depends on the adjustment of the lamp, and, of course, the weather conditions at the time the measurements are taken. Note that the measurements under the sun in Figure 6.5 were taken in Montreal on February 24, 2009, at 11:45AM, under clear skies.

121 6.3 RESULTS ON CELLS

Shown in Figure 6.6 are the J-V characteristic curves for four cells made from two different ingots, each with a proportion of the starting elements corresponding to the chemical formula CuInSe2.05. This is in contrast with cells made in previous studies in this laboratory, which were always fabricated from stoichiometric ingots, with the exception of the work H. Du, in which two cells were fabricated from two ingots with starting compositions of CuInSe2.2 and CuInSe2.4 [6.8]. The cells in

Figure 6.6 were all measured on the same day (November 11 2010) under sunlight.

As can be seen, the efficiencies achieved by these cells were between 1.4% and 5.8%.

The J-V characteristic for another cell, 62-1, also measured under sunlight, is shown in Figure 6.7. As can be seen, this cell, on this particular day (November 24

2010), was demonstrated to have a conversion efficiency of approximately 9%, with a fill factor (FF) of 52%. A few days later, on December 16, an efficiency of over 10% was measured using the xenon lamp, as shown in Figure 6.9. One interesting feature about this cell, apart from its relatively high conversion efficiency, is that it was not annealed, in comparison with the four cells described above.

The dark IV characteristic for cell 62-1, measured on November 24 2010, is shown in Figure 6.8. Plotted also is cell 35, fabricated by H. Du [6.4], and a theoretical cell slope with an ideality factor of 2. As can be seen, the ideality factor of 62-1 is lower than that of cell 35.

122 6.4 DIFFUSION LENGTH ESTIMATE BY PHOTOCURRENT-

CAPACITANCE MEASUREMENTS

The relatively high performance of cell 62-1 allowed for photocurrent- capacitance measurements, which are used to estimate the minority carrier diffusion length Ln [6.9]. The capacitance measurements were performed using a 4192A

Impedance Analyzer manufactured by Hewlett Packard.

For a one-sided, reverse-biased junction, the depletion capacitance is given by the following formula:

  A C  r 0 , 6.1 p W

where Cp is the depletion capacitance, εr is the relative dielectric constant of the material, ε0 is the permittivity of a vacuum, A is the junction area and W is the depletion width across the junction. Assuming W to be much smaller than the thickness of the absorber material, and assuming the penetration depth of the incoming radiation to be larger than the sum of W and diffusion length Ln, then the photocurrent Iph is given by:

I ph  K(W  Ln ) , 6.2

123 where K is a coefficient independent of voltage bias. Therefore, combining

Equations 6.1 and 6.2:

 1 L  I  K  A  n  . 6.3 ph r 0    C p  r 0 A

Figure 6.10 shows a ΔI versus 1/Cp plot for cell 62-1, where ΔI is the illuminated-to-dark current, equal to Iph for an ideal cell, measured under 1.3 μm light with reverse bias. Assuming a relative dielectric constant of εr = 10 and a cell area of

3 -1 A = 0.16 cm , the extrapolated intercept of 1/ Ci = -4.5 nF on the negative abscissa yields an estimated electron diffusion length Ln of about 6 μm. This is somewhat in agreement with the values for diffusion length estimated for annealed CdO-CuInSe2 cells by Z. Shukri [6.2], which were between 2.6-3.4 µm, but is an order of magnitude larger than that measured on the high-performing 12.5 % cell by H. Du et al [6.4].

6.5 DISCUSSION

While this work would certainly benefit from a more extensive study of photovoltaic cells, a consistent fabrication process is a requirement if the effect of the addition of Na in the melts, on cell performance is to be isolated. This is something that was not required of previous studies, where the objective was simply to make an increasingly higher-performing cell, and so then consistency was not a primary

124 objective. While a relatively high-performance cell has been made in this work, considerable effort is still necessary if meaningful comparisons between ingots are to be made. With respect to ingots made with Na, brittleness is a major impediment to the fabrication of good cells at higher Na concentrations. Therefore, it was not possible to study the characteristics of monocrystalline CuInSe2 grown under different conditions through fabrication and analysis of their respective photovoltaic devices. In order for this to be accomplished, it is necessary to first identify all sources of performance losses.

As explained in this chapter, some sources of performance loss were identified over the course of this work. However, many variables were left uncontrolled, and it is not clear what effect they might have on performance. Considering that several medium-performance cells and one higher-performance cell were made in the present work, and in preceding works, the absorber material itself, and the material growth process, is not thought to be a source of performance loss. However, the presence of micro-cracks in the CuInSe2 wafers, while avoided when visible, could not guarantee that no performance-inhibiting defects existed within the active areas chosen on the wafer surfaces. The metallization process, while initially considered a possible source of high series resistance, was also standardized during this work so that low- resistance, ohmic contacts could be made. Therefore, possible sources of performance loss include the polishing procedure, which was manual and therefore prone to inconsistency, the CdS layer, and the ZnO layer. However, despite this

125 uncertainty, both insight and experience was gained by the fabrication of photovoltaic cells over the course of this work.

126 a

copper thread b

indium with Wood’s metal

ZnO:Al

CdS on top of CuInSe2

aluminum stud

Figure 6.1 (a) Cross-sectional drawing of the cells fabricated in this work; (b) top photograph of a finished cell, in this case, cell 62-1.

127 a

Figure 6.2 (a) Cross-sectional SEM image of a layer of CdS deposited on a silicon wafer by the chemical bath method, using three depositions of 10 minutes each at 70 °C, showing the final layer to be 122 nm thick. Note also the uneven texture of the CdS layer, with large clumps of material at spots. (b) Top view, as seen by SEM, of a silicon wafer with two CdS depositions, also showing clumps of material across the surface of the wafer.

128

Figure 6.3 CdS on glass slides, showing a range of colours and textures, from light and smooth (left) to dark and patchy (right). All were deposited on glass (cleaned with toluene and rinsed with water before deposition), 5 repetitions of 8 minute deposition time, 50 °C, 60 °C, 70 °C, 80 °C (from left to right).

129

Figure 6.4 Current-voltage characteristics made between two evaporated gold contacts on an annealed surface (1) without lapping, and on annealed and then lapped surfaces (2 and 3), showing considerably lower resistances in the later case.

130

Figure 6.5 Illuminated IV characteristics for an early cell made in this work, using a filtered xenon arc lamp solar simulator and the actual sun. As indicated, the xenon lamp is an underestimate of the actual sunlight.

Figure 6.6 J-V characteristics for four cells made during this study, measured under sunlight in Montreal on November 11 2010.

131

Figure 6.7 J-V characteristic of cell 62-1 measured under sunlight in Montreal on November 24 2010, demonstrating an efficiency of approximately 9% and a FF of 52%.

Figure 6.8 Forward and reverse dynamic IV dark characteristics of cell 62-1, measured on November 24 2010. Also shown is that for cell 35, the best cell ever made in this lab, fabricated by H. Du, measured on January 23 2003. For comparative purposes, a theoretical cell slope with an ideality factor of 2 is also plotted.

132

Figure 6.9 J-V characteristic of cell 62-1, measured on December 16 2010 using a xenon lamp, calibrated to AM 1.5 intensity, demonstrating an efficiency of more than 10 % and a FF of 53 %.

Figure 6.10 Photocurrent-reciprocal capacitance plot for a cell made from Bridgman- grown CuInSe2.

133 Chapter 7

Discussion and Conclusions

7.1 INTRODUCTION

This chapter contains a summary and review of results described in Chapter 5.

As well, the results are interpreted in order to explain the action of the sodium in influencing the electrical changes seen in CuInSe2 when grown in the presence of this element. The chapter then concludes with a list of the original contributions to knowledge made in this work.

7.2 MECHANISM OF SODIUM ACTION IN BRIDGMAN-

GROWN CuInSe2

It has long been established in this laboratory [7.1-7.5] that stoichiometric melts containing copper, indium and selenium, corresponding to the chemical formula

CuInSe2, when solidified into a monocrystalline ingot by the one-ampoule Bridgman- method, results in chalcopyrite material that is always p-type. Furthermore, it had been shown previously, in the work of Wang et al [7.6], that sodium, when added in

134 sufficient quantities to the stoichiometric melts containing copper, indium and selenium corresponding to the chemical formula CuInSe2, results in n-type material.

The present work, for the first time, has demonstrated a unique relation between the conductivity type change from p to n with sodium, and the amount of excess selenium, above stoichiometry, present in the melt before compound synthesis and growth. Specifically, it has been shown that the presence of excess Se inhibits the type change, so that the critical amount of Na necessary for the change to occur is increased. In examining the critical amounts of Na with respect to excess Se, determined by fabricating material with increasing quantities of Na and Se until a pattern became apparent, it was determined, as seen in Figure 5.3, that the conductivity type of CuInSe2+x is governed by a 2:1 Na to x atomic ratio. In other words, if the value of [Na]/[x] is less than or equal to 2, the resulting material is p- type, whereas if it is greater than 2, the material is n-type. Here, x represents the amount of excess Se included in the melt above stoichiometry, so that [x] is the number of moles of excess Se. Initial attempts to find Na2Se-related phases by XRD in the residues of Bridgman-growth runs containing sodium were unsuccessful, perhaps due to the insensitivity of XRD to Na2Se, as demonstrated in Figures 5.26 to

5.28. However, Na2Se, and the hydrated compound Na2Se(H2O)5, both containing Na and Se with an [Na]/[Se] ratio of two, were detected by XRD when a piece of

CuInSe2 was heated to lower temperatures with more initial Na than would be possible to have included in the regular Bridgman-growth process.

135 With the verification of this relation between the Na, excess Se and conductivity type of the material, a model of the action of the Na which results in the p- to n-type conversion can be developed: the Na reacts with the Se in the growth melt, owing to the strong chemical affinity of Na for Se [7.7], to form the binary compound Na2Se. Once all of the Na present in the ampoule is exhausted, the remainder of the un-reacted Se is available to become part of the chalcopyrite. If this remainder is less than the required stoichiometric quantity, it leaves the ternary compound Se-deficient, and therefore n-type through the creation of donor-like Se- vacancies (VSe) [7.8]. However, if sufficient excess Se is added to the ampoule, above the stoichiometric quantity, then the Na will bond with that element, leaving the remaining Se to take its place within the CuInSe2 lattice so that the number of donor VSe sites does not exceed the number of acceptors already present in the material. The formed Na2Se is not able to reside within the interior of the intact

CuInSe2, and is thus deposited on the outside surface of the ingot, and in cracks and defects, such as those found in the last-zone-to-freeze.

Note that in stoichiometrically-prepared ingots, which are normally p-type, there is no apparent excess of Se and the ampoule, apart from the ingot itself, is clear of any deposits. Here, since x = 0, it was an amount of added Na greater than a certain critical level that resulted in a conductivity type change in the grown material.

This can be seen in the deviation from the origin of the left-most data point in Figure

5.3, which corresponds to an excess Se value of zero, This critical amount was identified in this work as greater than 0.2 at. %, as shown in Table 5.1, which is in

136 agreement with that identified in the work of Wang et al [7.6]. Therefore, intrinsic acceptors are present in some form in the material as grown. These acceptors are not specifically identified in this work, but could arise from the small deficiency of copper detected by the EDX analysis of the main part of the ingots.

These findings are also similar to those of previous workers in this laboratory in experiments on Bridgman ingots of CuInSe2 grown from melts intentionally-made selenium deficient [7.4, 7.9-7.10]. In these cases, the resulting material was n-type, with multi-phases in the last-zone-to-freeze. This experiment was repeated here for an ingot grown from a melt corresponding to the chemical formula CuInSe1.7, which was also found to be n-type. However, another similarity exists between the results of the experiment with intentional Se-deficiency and the experiments with Na, and that is the presence of precipitated copper, found on deposits within the sealed ampoules, exterior to the ingots. Z. Shukri et al [7.9] found that material grown from non-stoichiometric melts tended towards stoichiometry. So, with Se missing, the Cu and In were precipitated out; note that CuInSe1.7 could also be written as

Cu1.18In1.18Se2, that is, the chalcopyrite with excesses of Cu and In. The copper nodules seen in the present work may therefore be explained as being expelled from the material as it tended towards stoichiometry in order to compensate for the deficiency of selenium. Indium, however, was not detected as the element, but is perhaps present in extra phases in the last-zone-to-freeze as Cu7In4 and InSe [7.2].

However, Cu is more mobile and diffusive in CuInSe2 than the other elements [7.11],

137 and so is more easily removed from the material. As well, its colour makes it more visible and easier to detect than indium.

The white deposit, seen predominately in ampoules containing Na with little or no excess Se, was found by XRD to be a crystalline form of SiO2 called cristobalite, while underneath, the quartz was found to still be amorphous. SEM-

EDX analysis of the deposit uncovered mostly Si, O but also Na, with only trace amounts of the other elements. Therefore, it was concluded that the Na reacts with the fused quartz of the ampoule, causing it to nucleate and thereby weaken. This weakening, which can result in the ampoule cracking and breaking, as in the case of ingot HM-B54 (CuInSe2.05 + 11 at. % Na), imposes a limit on how much Na can be added to the melts. However, less of this white deposit was found in ampoules with increased amounts of excess Se. In these cases, a reddish-brown deposit, adhered to the quartz immediately adjacent to the ingot, was predominant. Analysis of this deposit showed it to contain all of the starting elements but primarily Na and Se. This has led to the belief that the Na acts with the Se to form the red deposit, but the quartz wall of the ampoule itself also competes to react with the Na. When less Se is present, the Na is then free to react with the quartz, weakening it and forming the white deposit. These colour and analysis observations are consistent with the selenium “starvation” model described earlier.

Investigation into the electrical properties of the material found no apparent clear change in Hall mobility with increase of Na alone, but an increase with excess

138 Se, from CuInSe2 to CuInSe2.2 was observed with Na also present. It is suggested that this increase is due to the filling of Se-vacancies by the excess Se, leading to a reduction of compensating donors and, therefore, a decrease of scattering centers.

However, SEM-EDX measurements did not indicate the atomic proportion of Se within the bulk material to increase, on average, with increases of excess Se in the growth melt. Similarly, analysis of the results obtained by SEM-EDX on material grown with Na found no apparent change in the atomic proportions, despite the change in conductivity type. Equally surprising is the results of the analysis by XRD, in which no change in crystal structure was found, even with additions of sodium as high as 3 at. % in stoichiometric material. It was imagined that structural changes would naturally accompany the obvious electrical changes seen in the material, as the conductivity type of the ingots was seen to be dependent on the composition of the growth melts, but this proved not the case. However, the findings are consistent with the work of Shukri et al [7.9], where the conductivity type of the material was found to be a function of the starting composition of the melt, but apparently largely independent of the finished composition of the ingot, as determined by EPMA. Note also that no evidence for the formation of ordered defect compounds (ODC) with Na was observed in this work; there were no traces of any of peaks belonging to an ODC in any of the XRD scans. Since such ODC forms only at the surface of thin films

[7.12], if it occurred also on the present Bridgman material before powdering, its presence would have been at too low a level to be detected by XRD.

139 With regard to the specifics of the compositional analysis of the main part of each ingot by SEM-EDX, it was found that the material was slightly copper deficient, having an atomic percentage of Cu less than 25 %, and slightly selenium rich, with atomic percentages over 50 %. This was true for all material for which SEM-EDX analysis was performed, regardless of conductivity type or the starting composition of the melts. However, the opposite was true for the material in the last-zone-to-freeze of all ingots examined, where SEM-EDX showed excess Cu but reduced Se. This confirms that the stable form of the compound from the one-ampoule Bridgman- process is slightly Cu-deficient and Se-rich material.

Note also that the material became increasingly brittle with Na. It is thought that the increase in brittleness results from this element removing Se from the compound and settling on microcracks and defects. This was especially apparent in the experiment in which a solid piece of pre-grown, p-type CuInSe2, after being heated with a large amount of Na, crumbled into fragments, which were seen to be n- type. This reduction of cohesiveness with Na hampered resistivity measurements, and so the above comments on mobilities are made only for material where Na content was low.

To be clear, this work does not indicate Na to be a donor in CuInSe2. Rather, it is believed that the action of the Na is the removal of Se, resulting in the creation of donors in the material. Thus, it is the Se vacancies which are the donors. In earlier work [7.13], when Na2Se was added to the melts, rather than elemental Na, the

140 material remained p-type (see Figure 7.1). There, the thermoelectric values were seen to decrease, indicating that Na2Se has an acceptor action in the material. This is in agreement with research done elsewhere on thin films [7.14-7.16], which are always grown with excesses of Se, although it is not clear what the acceptor action of the Na is in this case. Nevertheless, the present work may be viewed as somewhat controversial, considering that many researchers regard Na as an acceptor, while the results here indicate that Na apparently has donor action under some circumstances.

Finally, it was shown in this work that it is possible to fabricate reasonably good photovoltaic devices from the Bridgman-grown material. Difficulties associated with the fabrication process prevented comparisons to be made between monocrystalline cells with and without sodium. Therefore, the effect of the presence of Na on monocrystalline cells made from Bridgman-ingots, as well as its relevance to the increase in performance seen in thin-film polycrystalline cells, can only be speculated upon. The observations made in this thesis on the electrical characteristics of the Bridgman-grown material, however, may be used to help explain the performance of thin-film photovoltaic devices.

One such electrical characteristic is the conductivity type change from p- to n- type. It has been demonstrated here, and in previous work [7.6] that Na has, under the right circumstances, a net donor action on CuInSe2, even though it is not a true donor. If these circumstances exist at specific locations in a thin-film of CuInSe2, the result may be the formation of n-type regions, or perhaps even an n-type layer within,

141 or at a surface of, the film. CIGS photovoltaic devices, which were traditionally believed to consist of a pn-junction heterojunction between the p-type CIGS and the n-type CdS, are now believed to actually consist of a buried homojunction beneath the surface of the CIGS film [7.17-7.18], or a heterojunction of p-type CuInSe2 and n- type CuIn3Se5 [7.12]. If Na plays a role in the creation of this junction, it may act to enhance this layer and allow it to be created farther away from physical CIGS-CdS boundary, thereby lessening the negative effects of interface scattering centers, which reduce lifetime and minority diffusion length. This idea, of course, has not been tested.

As well, the observed increased majority carrier mobility, from Hall measurements, in material grown with excess Se at low Na concentrations, may lead to speculation that the minority carrier mobility also increases. If this were true, it would mean that the minority diffusion length also increases, and with it, the capture distance in a photovoltaic cell. However, as with the hypothesis above, this is merely conjecture, and more work would be needed to confirm its validity.

Lastly, it is possible that sodium will have no beneficial effect at all on single crystal photovoltaic cells. Most researchers agree that there is an improvement in

{112} orientation on thin-film polycrystalline cells grown with Na. However, single crystals are, in principal, perfectly oriented, so the effect of preferred {112} orientation in polycrystalline devices would, in monocrystalline devices, be

142 dependent on wafer cutting and polishing procedures, rather than directly on growth procedures.

Considering that such hypotheses are based on results of experiments done on single crystals, it is hoped that work on monocrystalline CuInSe2 can continue.

7.3 ORIGINAL CONTRIBUTIONS TO KNOWLEDGE

1. A conductivity type change from p- to n-type, which occurs in CuInSe2 grown

from stoichiometric Bridgman-grown melts to which Na is added, occurs

abruptly with Na additions between 0.2 - 0.3 at. %.

2. The critical amount of Na required for the type conversion to occur when

excess Se is present is approximately proportional to the amount of excess Se

also included in the ampoule, and corresponds to 2 atoms of Na to 1 atom of

excess Se, consistent with the chemical formula for sodium selenide (Na2Se).

3. Within residues from heating pre-grown, p-type Bridgman CuInSe2 with a

very large amount of Na, where p- to n-type conversion had taken place,

Na2Se, Na2Se(H2O)5 and possibly Na2SeO3 were detected by X-ray

diffraction, all having a ratio of two atoms of Na to one of Se, in accordance

with Contribution 2, above.

4. The chalcopyrite structure of CuInSe2 in the bulk was maintained with

additions of Na into the melt, up to at least 3 at. %, as confirmed by X-ray

143 diffraction, even as the material is converted from p- to n-type as a result of

the addition.

5. Analysis by SEM-EDX on the main part of ingots found [Cu]/[In] < 1, while

[Se]/{[Cu]+[In]} > 1. This was true for all samples for which measurements

were taken, and no correlation was apparent between these values and the

addition of Na or excess Se, or the final conductivity type of the ingots.

6. By contrast, the last-zones-to-freeze of samples measured by SEM-EDX were

all found to be copper-rich and selenium-deficient, also irrespective of starting

composition. This region was always p+-type, regardless of the conductivity

type of the main part of the ingot.

7. Na was found in deposits within the ampoules after growth, and in the last-

zone-to-freeze, but none was detected by SEM-EDX within the intact bulk

material itself.

8. Elemental copper was detected with n-type samples grown with Na, and in a

deposit found on a previously-grown CuInSe2 piece heated with Na. It was

also found in deposits exterior to an ingot grown without Na but with an

intentional deficiency of selenium, corresponding to the chemical formula

CuInSe1.7.

9. The cohesivity of grown ingots was seen to decrease with increases of Na,

particularly when excess Se was also added to the melt.

10. The presence of Na within the ampoules was found to weaken the fused-

quartz ampoules by causing the amorphous quartz to crystallize.

144 11. An increase in mobility for p-type samples as excess Se was increased in the

melt, at least for low Na additions, was apparent.

12. The majority carrier concentrations in the grown ingots material were

determined, by Hall measurements, to be ~1017 cm-3 for p-type material, and

at least an order of magnitude less for n-type material. These trends are also

confirmed by the thermoelectric power-derived concentration values.

13. A selenium “starvation” model to describe the mechanism by which the Na

acts to convert the material from p- to n-type was developed.

145

Figure 7.1 Variation of thermoelectric power, measured by hot-probe, on Bridgman- grown stoichiometric CuInSe2 with added Na2Se [7.13].

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156 Appendix A

Ingot Growth Conditions

This appendix contains the details of the growth conditions for every ingot grown in this study, starting September 1, 2008. Some of the ingots mentioned in the text of this work were grown in the precursor to this study, which was completed in

August of 2008. For the details of the growth conditions of those ingots, consult

“Studies on the effect of sodium in Bridgman-grown CuInSe2” by H.F. Myers,

McGill University, Department of Electrical and Computer Engineering Master’s

Thesis, 2008.

157

Figure 1A Pre-reaction heating ramp-up scheme used for all ingots grown in this work.

Figure 2A Temperature ramp-up schemes used to bring the ampoules to the maximum soaking temperature.

158 Table A1 Ingot Growth Conditions Ingot Number HM-B33 HM-B34 HM-B35 HM-B36 HM-B37 HM-B38 HM-B39 HM-B40 Composition (g) 5g In; 5g In; 5g In; 5g In; 5g In; 5g In; 5g In ; 5g In ; 6.8770g Se; 7.565g Se; 6.8770g Se; 6.8770g Se; 6.8770g Se; 7.565g Se; 6.8770g Se; 7.565g Se ; 2.7672g Cu; 2.767g Cu; 2.7672g Cu; 2.7672g Cu; 2.7672g Cu; 2.7672g Cu; 2.7672g Cu; 2.7672g Cu ; 0.2864g Na2Se 0.2864g Na2Se 0.011g Na 0.0011g Na 0.0022g Na 0.0022g Na Composition (mol) (CuInSe2)0.95 (CuInSe2.2)0.95 CuInSe2 CuInSe2 (CuInSe2)0.99 (CuInSe2.2)0.999 (CuInSe2)0.998 (CuInSe2.2)0.998 + (Na2Se)0.05 + (Na2Se)0.05 + Na0.01 + Na0.001 + Na0.002 + Na0.002 Evacuation Time 3:00 3:20 5:15 4:55 6:12 5:55 4:25 4:15 (h:m) Pre-reaction time 4:23 4:10 3:50 7:18 4:45 5:40 4:55 5:05 (h:m)* Peak Upper Furnace 1050 1050 1050 1050 1050 1050 1050 1050 Temperature (°C) Reaction Time (h) 23:50 24:05 24:00 24:05 24:00 24:00 24:05 24:05

15 Number of Agitations 1 1 2 1 1 1 1 1

9 Lowering Rate 4.14 4.20 4.24 4.12 2.16 2.16 3.63 3.63 (mm/h) Peak Lower Furnace 700 700 700 700 700 700 700 700 Temperature (°C) Heating scheme Fig. 2A a Fig. 2A a Fig. 2A a Fig. 2A a Fig. 2A a Fig. 2A b Fig. 2A a Fig. 2A a Cooling Rate (°C/h) 15.71 15.71 7.86 15.71 15.71 15.71 15.71 15.71 Conductivity type p p n p n p p p Conduc. Type LZTF p+ p+ p+ p+ p+ p+ p+ p+ Comments A crack formed This run is Meant to Ampoule in the ampoule considered to be include only brought up to sometime after an error, 0.1 % 1100 °C in the lowering had although the sodium, but brick furnace been complete. cause of the n- more was prior to The temperature type inadvertently Bridgman- at the time was conductivity included as growth. between 500 and was never determined 700 °C. The run determined. by the was not amount of interrupted. white deposit on the ampoule.

Table A2 Ingot Growth Conditions Ingot Number HM-B41 HM-B42 HM-B43 HM-B44 HM-B45 HM-B46 HM-B47 HM-B48 Composition (g) 5g In; 5g In; 5g In; 5g In; 5g In; 5g In; 5g In; 5g In ; 6.8770g Se; 7.565g Se; 7.565g Se; 6.8770g Se 6.8770g Se; 7.565g Se; 7.9085g Se; 7.2208g Se ; 2.7672g Cu; 2.767g Cu; 2.767g Cu; 2.767g Cu; 2.767g Cu; 2.767g Cu 2.767g Cu 2.767g Cu 0.003g Na. 0.004g Na 0.003g Na Composition (mol) CuInSe2 (CuInSe2.2)0.997 + (CuInSe2.2)0.996 (CuInSe2)0.997 + CuInSe2 CuInSe2.2 CuInSe2.3 CuInSe2.1 Na0.003 + Na0.004 Na0.003 Evacuation Time (h:m) 5:16 3:10 3:10 24:15 3:15 3:10 4:55 3:25 Pre-reaction time 5:55 4:45 3:35 2:30 5:35 4:00 5:30 3:30 (h:m)* Peak Upper Furnace 1050* 1050 1100 1100 1100 1100 1100 1100 Temperature (°C) Reaction Time (h) 24:00 24:00 24:00 24:00 24:00 24:00 24:00 24:00 Number of Agitations 8 1 4 3 2 2 2 3 Lowering Rate (mm/h) 4.86 4.66 4.12 4.28 3.45 3.96 4.05 3.34

1 Peak Lower Furnace 700 700 700 700 700 700 700 700 60 Temperature (°C) Pre-reaction scheme Fig. 1A Fig. 1A (up to Fig. 1A (up to Fig. 1A (up to Fig. 1A (up to Fig. 1A (up Fig. 1A (up to Fig. 1A (up to 400 °C) 400 °C) 400 °C) 400 °C) to 400 °C) 400 °C) 400 °C) Heating scheme Fig. 2A a Fig. 2A a Fig. 2A a Fig. 2A a Fig. 2A a Fig. 2A a Fig. 2A a Fig. 2A a Cooling Rate (°C/h) 12.69 15.71 15.71 15.71 15.71 15.71 15.71 15.71 Conductivity type p p p n p p p p Conduc. Type LZTF p+ p+ p+ p+ p+ p+ p+ p+ Comments *Top furnace malfunctioned during run. Temperature dropped during lowering to 850°C before being brought up to 1030°C.

Table A3 Ingot Growth Conditions Ingot Number HM-B49 HM-B50 HM-B51 HM-B52 HM-B53 HM-B54 HM-B55 HM-B56 Composition (g) 0.611g In; 5g In; 5g In; 5g In; 5g In; 5g In; 5g In; 5g In; 0.924g Se; 2.767g Cu; 2.767g Cu; 2.767g Cu; 2.767g Cu; 2.767g Cu; 2.767g Cu; 2.767g Cu; 0.338g Cu; 6.8770g Se 7.0489g Se 6.894g Se; 7.0489g Se 7.0489g Se ; 6.9457g Se; 6.9457g Se; 0.1g Na 0.033g Na 0.033g Na. 0.055g Na 0.13g Na 0.056g Na. 0.033g Na.

Composition (mol) (CuInSe2.2)0.55 + CuInSe2 (CuInSe2.05)0.97 (CuInSe2.005)0.97 (CuInSe2.05)0.95 (CuInSe2.05)0.89 (CuInSe2.02)0.95 (CuInSe2.02)0.97 Na0.45 + Na0.03 + Na0.03 + Na0.05 + Na0.11 + Na0.05 + Na0.03 Evacuation Time 3:30 3:25 3:57 3:15 3:30 3:40 3:40 3:35 (h:m) Pre-reaction time 4:25 4:15 4:40 5:40 4:05 3:25 5:40 4:15 (h:m)* Peak Upper Furnace 1100 1100 1100 1100* 11:00 1100 1100 1100 Temperature (°C) Reaction Time (h) 24:00 24:00 24:10 24:00 24:00 21:00* 24:00 24:00

61 Number of Agitations 3 2 3 2 3 3 3 2

Lowering Rate 3.91 3.91 3.91 3.91 3.91 5.33 3.58 3.58 (mm/h) Peak Lower Furnace 700 700 700 700 700 700 700 700 Temperature (°C) Pre-reaction scheme Fig. 1A (up to Fig. 1A (up to Fig. 1A (up to Fig. 1A (up to Fig. 1A (up to Fig. 1A (up to Fig. 1A (up to Fig. 1A (up to 400°C) 400°C) 400°C) 400°C) 400°C) 400°C) 400°C) 400°C) Heating scheme Fig. 2A a Fig. 2A a Fig. 2A a Fig. 2A a Fig. 2A a Fig. 2A a Fig. 2A a Fig. 2A a Cooling Rate (°C/h) 15.71 15.71 15.71 15.71 15.71 15.71 15.71 15.71 Conductivity type undetermined p p n p n/p n p Conduc. Type LZTF undetermined p+ p+ p+ p+ p+ p+ p+ Comments Thick, flaky *Top furnace Shortened soak white deposit on malfunctioned, time to try to outside of causing avoid reaction ampoule; temperature to between perhaps sodium drop to 935°C. sodium and diffused through Aluminum foil quartz quartz. Lower used to insulate ampoule; part broke off in furnace, raising cracked during furnace; no solid temperature to cool down ingot recovered. around 1050°C. phase.

Table A4 Ingot Growth Conditions Ingot Number HM-B57 HM-B58 HM-B59 HM-B60 HM-B61 HM-B62 HM-B63 HM-B64 Composition (g) 5g In; 5g In; 5g In; 5g In; 5g In; 5g In; 5g In; 5g In; 6.8942g Se; 2.767g Cu; 2.767g Cu; 2.767g Cu; 2.767g Cu; 2.767g Cu; 2.767g Cu; 2.767g Cu; 2.767g Cu; 6.9457g Se; 6.8770g Se 7.0489g Se; 7.0489g Se; 7.0489g Se; 6.8770g Se; 6.8942g Se; 0.0055g Na 0.022g Na; 0.17g Na2Se 0.01g Na ;

Composition (mol) (CuInSe2.005)0.995 (CuInSe2.02)0.98 + (CuInSe2)0.97 + CuInSe2.05 CuInSe2.05 CuInSe2.05 CuInSe2 (CuInSe2.005)0.99 + Na0.005 Na0.02 (Na2Se)0.03 + Na0.01 Evacuation Time (h:m) 3:15 4:10 4:05 6:50 3:05 3.05 3:45 3:25 Pre-reaction time 3:20 5:30 6:15 3:00 3:55 5:20 4:05 3:30 (h:m)* Peak Upper Furnace 1100 1100* 1100 1100 1100 1100 1100 1100 Temperature (°C) Reaction Time (h) 24:00 24:00 24:00 24:00 24:00 24:00 24:00 24:00 Number of Agitations 2 1 2 2 2 1 2 2

1 Lowering Rate (mm/h) 3.58 4.66 3.85 3.94 3.81 3.80 3.70 3.60 62 Peak Lower Furnace 700 700 700 700 700 700 700 700 Temperature (°C) Pre-reaction scheme Fig. 1A (up to Fig. 1A (up to Fig. 1A (up to Fig. 1A (up Fig. 1A (up to Fig. 1A (up to Fig. 1A (up to Fig. 1A (up to 400°C) 400°C) 300°C) to 400°C) 400°C) 400°C) 400°C) 400°C) Heating scheme Fig. 2A a Fig. 2A a Fig. 2A a Fig. 2A b Fig. 2A b Fig. 2A b Fig. 2A b Fig. 2A a Cooling Rate (°C/h) 15.71 15.71 15.71 15.71 15.71 15.71 15.71* 15.71* Conductivity type p p undetermined p p p p p Conduc. Type LZTF p+ p+ undetermined p+ p+ p+ p+ p+ Comments *Top furnace Some cracks Good run Good run Good run *Uncertainty *Uncertainty as malfunctioned. apparent near as to exact to exact cooling Ampoule was top of ampoule cooling rate rate (controller cooled to room from sodium. (top for bottom temperature after thermocouple furnace reaching 1100 sensor malfunction). °C and then malfunction). reheated with new furnace element.

Table A5 Ingot Growth Conditions

Ingot Number HM-B65 HM-B66 HM-B67 HM-B68 HM-B69 HM-B70 HM-B71 Composition (g) 5g In; 5g In; 10g In; 5g In; 5g In; 5g In ; 5g In ; 6.8942g Se; 6.9457g Se; 5.5345g Cu; 6.9457g Se; 7.565g Se; 5.845g Se ; 6.8770g Se 2.767g Cu; 2.767g Cu; 13.8915g Se; 2.767g Cu 2.767g Cu 2.767g Cu 2.767g Cu ; 0.022g Na 0.044g Na 0.066g Na. 0.066g Na. 0.1683g Na2Se

Composition (mol) (CuInSe2.005)0.98 (CuInSe2.02)0.96 + CuInSe2.02 (CuInSe2.02)0.94 + (CuInSe2.2)0.94 + CuInSe1.7 (CuInSe2)0.97 + + Na0.02 Na0.04 Na0.06 Na0.06 (Na2Se)0.03 Evacuation Time 7:50 3:40 4:40 3:20 4:45 3:20 3:20 (h:m) Pre-reaction time 5:15 3:20 3:15 3:30 5:25 4:45 4:00 (h:m)* Peak Upper Furnace 1100 1100 1050 1050 1050 1050 1050

1 Temperature (°C) 63 Reaction Time (h) 24 24 24 24 24 24 24

Number of Agitations 2 2 1 2 3 4 2 Lowering Rate 3.81 3.90 3.98 3.52 3.52 3.41 3.53 (mm/h) Peak Lower Furnace 700 700 700 700 700 700 700 Temperature (°C) Pre-reaction scheme Fig. 1A (up to Fig. 1A (up to Fig. 1A (up to Fig. 1A (up to Fig. 1A (up to Fig. 1A (up to Fig. 1A (up to 400°C) 400°C) 400°C) 400°C) 300°C) 400°C) 400°C) Heating scheme Fig. 2A a Fig. 2A a Fig. 2A a Fig. 2A a Fig. 2A a Fig. 2A a Fig. 2A a Cooling Rate (°C/h) 15.71 15.71 15.71 15.71 15.71 15.71 15.71 Conductivity type n p p n p n p Conduc. Type LZTF p+ p+ p+ p+ p+ p+ p+ Comments Heated twice to Indium boat 1100 °C and method to agitated 4 times; include Na2Se furnace in ampoule. malfunctions forced run to be halted; third time to 1050 °C successful.

Appendix B

Results of Assay on Elemental Sodium

A sample of the metallic sodium used in this work was sent for analysis to

VHG Labs, in New Hampshire. This was to determine the purity of the material, and to identify any impurities which may be present. The results of this assay, given in this appendix, indicate the sodium to be in excess of 99.99% pure. Iron, aluminum and calcium were the only trace impurities detected in concentrations of 1 ppm, 1 ppm and 41 ppm, respectively.

164

165

Appendix C

Results of Electrical and EDX Analysis of Ingots

These results are offered as an extension to those given in Chapter 5, in which only selected data are shown. Table C1 and C2 contain the all the results of the electrical measurements on filaments made in this work. Table C3 gives the composition results measured on wafers cut from the interior of the ingots. This complete accounting is meant to enable a critical review of the results and, with it, the conclusions presented in this thesis.

166

Table C1 Results of Electrical Measurements on all Samples

Majority carrier Majority Majority carrier mobility carrier mobility concentration Majority carrier interpreted from interpreted Hall interpreted from concentration Seebeck from Hall Seebeck coefficient Electrical Seebeck interpreted from coefficient5 coefficient6 2,3 4 % Na in Conductivity coefficient RH resistivity coefficient Hall coefficient µpα, µnα µpH, µnH 1 3 -3 -3 2 -1 -1 2 -1 -1 x Run melt type α (µV/K) (cm /Coul) ρ (Ω∙cm) pα, nα (cm ) pH, nH (cm ) (cm ∙V ∙s ) (cm ∙V ∙s ) HM-B50 0 p 387 13.8 0.8 1.2E+18 5.3E+17 6.3 14.1 HM-B24 0.1 p 314 62.3 8.0 2.8E+18 1.2E+17 0.3 6.6 HM-B39 0.2 p 371 20.1 1.3 1.5E+18 3.7E+17 3.4 13.3

HM-B44 0.3 n -334 -112.0 0.3 1.0E+17 6.6E+16 189.8 296

0 0

16 HM-B27 0.5 n -421 -1010.0 5.0 3.7E+16 7.3E+15 33.4 171

7 HM-B28 1 n -562 -1470.0 7.1 7.3E+15 5.0E+15 122.1 177

HM-B29 2 n -416 -1200.0 3.4 4.0E+16 6.2E+15 46.6 301 HM-B26 3 n -553 -1510.0 3.5 8.1E+15 4.9E+15 221.4 364

HM-B57 0.5 p 526 - - 2.4E+17 - - -

HM-B64 1 p 352 - - 1.8E+18 - - -

0.005 - - - HM-B65 2 n -389 - 5.4E+16 - HM-B52 3 n -256 - - 2.6E+17 - - - HM-B58 2 p 600 - - 1.0E+17 ------

HM-B56 3 p 468 - 4.7E+17 -

HM-B66 4 p 473 - - 4.4E+17 - - - 0.02 HM-B55 5 n -454 - - 2.6E+16 - - - HM-B68 6 n -255 - - 2.6E+17 - - - 1 Corresponding to formula CuInSe2+x 2 mp/m0 = 0.7, mn/m0 = 0.09 3 * * αp = (k/e)∙(2-ηp ), αn = (k/e)∙(2-ηn ), assuming dominant acoustic lattice scattering. 4 pH, nH = (3π/8)∙(1/eRH), assuming dominant acoustic lattice scattering 5 µpα = 1/(pαeρ), µnα = 1/(nαeρ). 6 µpH = 1/(pHeρ), µnH = 1/(nHeρ).

Table C2 Results of Electrical Measurements on all Samples

Majority carrier Majority carrier Majority carrier concentration Majority carrier mobility interpreted mobility Hall interpreted from concentration from Seebeck interpreted from Seebeck coefficient Electrical Seebeck interpreted from coefficient Hall coefficient % Na in Conductivity coefficient RH resistivity coefficient Hall coefficient µpα, µnα µpH, µnH 3 -3 -3 2 -1 -1 2 -1 -1 x Run melt type α (µV/K) (cm /Coul) ρ (Ω∙cm) pα,nα (cm ) pH,nH (cm ) (cm ∙V ∙s ) (cm ∙V ∙s )

HM-B51 3 p 550 368.0 114 1.8E+17 2.0E+16 0.3 2.7 HM-B53 5 p 446 - 1675 6.1E+17 - 0.0 - 0.05 HM-B54 11 n/p1 -445 - - 2.8E+16 - - - 0.1 HM-B48 0 p 445 17.9 1.8 6.1E+17 4.1E+17 5.7 8.5 HM-B13 0 p 321 19.3 0.9 2.6E+18 3.8E+17 2.8 19.0 HM-B20 0.1 p 429 18.5 1.0 7.4E+17 4.0E+17 8.1 15.1 HM-B40 0.2 p 356 24.3 0.9 1.7E+18 3.0E+17 4.2 23.7

HM-B42 0.3 p 348 27.2 3.7 1.9E+18 2.7E+17 0.9 6.2

16 HM-B43 0.4 p 430 - 2.0 7.3E+17 - 4.3 - 0.2

8 HM-B21 0.5 p 335 - 0.6 2.2E+18 - 4.4 - HM-B23 1 p 275 - 23.7 4.4E+18 - 0.1 - HM-B19 2 p 400 - 97 1.0E+18 - 0.1 - HM-B22 3 p 377 - 2.2 1.4E+18 - 2.1 - HM-B69 6 p 265 - - 5.0E+18 - - - 0.3 HM-B47 0 p 275 15.5 0.8 4.4E+18 4.8E+17 1.7 15.8

HM-B32 0 p 426 - 0.6 7.7E+17 - 14.0 -

HM-B25 3 p 412 - 6.1 9.0E+17 - 1.1 - 0.4 HM-B31 10 p 332 - 22.4 2.3E+18 - 0.1 - 1Sample contained both n-type and p-type regions; filament cut from n-type region. Table C3 Results of SEM/EDX analysis on bulk material from center of ingot Conductivity Run x % Na in melt type [Cu]2 [In] [Se] [Na] HM-B19 0.2 2 p 21.6 26.1 52.3 0 HM-B22 0.2 3 p 22.0 25.5 52.4 0 HM-B44 0 0.3 n 22.0 24.7 52.7 0.53 HM-B45 0 0 p 22.4 25.0 52.6 0 HM-B29 0 2 n 22.4 25.0 52.5 0 2Designates percent concentration of element in the bulk; average of three measurements. 3Possibly due to deposition of Na on imperfections.

Appendix D

Photovoltaic Cell Fabrication Procedures

This appendix contains the fabrication procedures for five cells made during this work. In total, approximately 50 cells were made, although the performance of the majority was not sufficient for meaningful measurements or comparisons to be made. Table D includes some of the higher-performing cells made in this work, including cell 62-1, the highest-performing cell. Note that the identifying cell numbers for each cell includes the ingot number, followed by the cell number from that ingot in order of their fabrication; e.g. cell 62-1 was the first cell made from ingot HM-B62.

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Table D Cell Fabrication Procedures

Cell number 60-1 61-1 61-2 61-3 62-1 600-c 5 5 5 5 5 (min) 800-c 5 5 5 5 5 (min) 1200-c 5 5 5 5 5 (min) 0.05µm 1 1.5 1.5 2 1.5

Polishing (hour)

Etching 0.5% BM 0.5% BM (30) 0.5% BM (30) 0.5% BM 0.5% BM (30) (30) (30) Annealing 350°C for 350°C for 2hrs 350°C for 2hrs 350°C for No anneal 2hrs 2hrs Back Contact Gold Gold Gold Gold Gold CdS 1 x 60°C 2 x 60°C 12 2 x 60°C 12 2 x 60°C 12 2 x 70°C 12 12 min min min min min ZnO 7 hrs 7 hrs 7 hrs 7 hrs 8 hrs Top Contact Indium Indium stripes Indium stripes Indium Indium dot stripes stripes Cell area 16 16 16 16 16 (mm2) Efficiency (%) 3.6 5.8 3.9 1.4 8.8 VOC (V) 0.42 0.4 0.37 0.26 0.43 JSC (mA) 24 31.5 24 16.5 37.5 Fill Factor (%) 35 43 42 30 52 Ideality factor 1.9 1.8 1.5 2.5 1.4 Comments ZnO peeled Best cell off. made in this work.

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