Atomic Layer Deposition of Thin Film Indium Oxysulfide - a Non- Toxic Electron Transport Layer for Chalcogenide Solar Cells

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Atomic Layer Deposition of Thin Film Indium Oxysulfide - A Non- Toxic Electron Transport Layer for Chalcogenide Solar Cells

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Citation Jayaraman, Ashwin. 2018. Atomic Layer Deposition of Thin Film Indium Oxysulfide - A Non-Toxic Electron Transport Layer for Chalcogenide Solar Cells. Doctoral dissertation, Harvard University, Graduate School of Arts & Sciences.

Citable link http://nrs.harvard.edu/urn-3:HUL.InstRepos:41129213

Terms of Use This article was downloaded from Harvard University’s DASH repository, and is made available under the terms and conditions applicable to Other Posted Material, as set forth at http:// nrs.harvard.edu/urn-3:HUL.InstRepos:dash.current.terms-of- use#LAA Atomic Layer Deposition of Thin Film Indium

Oxysulfide - A Non-Toxic Electron Transport

Layer for Chalcogenide Solar Cells

A dissertation presented by

Ashwin Jayaraman

The Harvard John A. Paulson School of Engineering and Applied Sciences

in partial fulfillment of the requirements

for the degree of

Doctor of Philosophy

in the subject of

Engineering Sciences

Harvard University

Cambridge, Massachusetts

April 2018

Dissertation Advisor: Professor Roy G. Gordon Ashwin Jayaraman

Atomic Layer Deposition of Thin Film Indium

Oxysulfide - A Non-Toxic Electron Transport

Layer for Chalcogenide Solar Cells

Abstract

CZT(S,Se) (Cu2ZnSn(SxSe1-x)4) has emerged as an earth-abundant, non-toxic alternative to thin-film photovoltaic technologies based on CIG(S,Se) (Cu(In,Ga)(S,Se)2) and CdTe.

Devices employing CZT(S,Se), however, suffer from poor voltage extraction, reaching only 60

% of the Shockley-Queisser limit. The low photovoltage results primarily from interfacial recombination at the absorber-electron conductor junction, owing to mismatch in conduction band energies, lack of conformality, unpassivated defects, and elemental interdiffusion. The best existing heterojunction devices employ CdS as the electron transport layer.

We have explored In2(O,S)3 films prepared by atomic layer deposition (ALD) as an electron transport layer for CZT(S,Se). The motivation for this thesis was to tune the bands of

In2(O,S)3 by controlling the S:O ratio. This would result in higher collection of photo-generated electrons and higher open circuit voltage on coupling In2(O,S)3 with CZT(S,Se). We initially

3+ 2+ selected In2S3 on grounds that In ion possibly has lower diffusivity than Cd which should limit elemental interdiffusion. It has been previously established in literature by solution deposition that incorporation of oxygen into In2S3 allows us to obtain band positions at the p-n

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junction that should enable good minority carrier extraction. Oxygen incorporation also increases the band gap and therefore the optical transparence.

Unfortunately, solution-based methods leave a large number of uncontrollable hydroxyl groups in the film, which affords poor control over electrical properties of the film. We now report that atomic layer deposition (ALD) using alternate cycles of tris(N,N'- diisopropylformamidinato) indium(III) (indium formamidinate), water, and hydrogen sulfide at

o 200 C results in growth of a pure In2(O,S)3 film (free of halides, carbon) with close control in- situ over sulfur to oxygen ratio. We arrived at conditions for growth of In2(O,S)3 by studying

ALD of binaries In2S3 and In2O3 independently. The choice of precursor for growing In2(O,S)3 was made on the basis of kinetics of the growth process of the aforementioned binaries. As we use water as the oxygen source, there is incorporation of adventitious hydroxyl groups. We are able to control the content of hydroxyl groups in the In2(O,S)3 film though, by varying number of water containing sub-cycles.

In2(O,S)3 films exhibited an indirect band gap higher than In2S3 which minimizes absorption losses in the electron transport layer. A reasonably high mobility was obtained with wide and tunable range of carrier concentrations over 5 orders of magnitude. To limit recombination, fewer charge carriers are targeted in In2(O,S)3 close to its junction with absorber while more carriers are targeted adjacent the transparent conducting oxide in the solar cell. Band offset measurements of In2(O,S)3 with reference to CZT(S,Se) by x-ray photoelectron spectroscopy indicate that oxygen incorporation (and resultant decrease in S:O ratio) increases conduction band offset relative to the pure sulfide. We identify In2(O,S)3 with oxygen contents between 4-33 at. % as the optimal composition in combination with CZT(S,Se), to potentially boost the open circuit voltage of the solar cells.

Acknowledgements

This thesis came to fruition because of the kind help and support of many individuals. At the outset I would like to extend my sincerest thanks to my research advisor Professor Roy

Gordon for keeping me extremely motivated throughout the duration of graduate research. I like and admire his ability to perform application-oriented research no end. I deeply appreciate his help in my choice of the research topic and related projects. It was an amazing experience to learn the art of well-directed, innovative and improvised scientific research from him. This work would not have been possible without the invaluable collaboration opportunities that Professor

Roy Gordon provided with centers of learning like IBM Thomas J Watson labs and

Massachusetts institute of technology. I would like to acknowledge the generous financial support by my advisor during my time at Harvard. Simply put, I could not have had a nicer guide and I cannot thank him enough for all the academic support and help with finalizing my career going forward.

I sincerely thank my defense committee members Professor Michael Aziz and Professor

Frans Spaepen. Professor Aziz also served on my qualification committee and asked me interesting questions on my research direction. I am ever so thankful to have learnt from esteemed professors like him and express my gratitude for all his guidance. I thoroughly enjoyed the “Energy Technology” course conducted by him.

Professor Spaepen who was also on my qualification committee was an amazing mentor to me at Harvard. His course on “Kinetics of condensed phase processes”, I felt was the best class that I took at Harvard. I am simply amazed by his ability to communicate complex ideas in

very simple terms. I am very thankful for his generous contribution of time and efforts to make sure my concepts are clear.

I am also grateful to Professor David Clarke and Professor Shriram Ramanathan (Now at Purdue

University) for mentoring me in my 1st year at Harvard. I learnt immensely from the studies on thermal conductivity variations with pore geometry that I conducted with Professor Clarke. I also benefitted no end from Professor Ramanathan’s knowledge of functional oxide thin films. I would like to thank Professor Ramanathan for giving me the opportunity to serve as his teaching assistant for his course “Electrical, Optical and Magnetic Properties of Materials”

I take this opportunity to thank my undergraduate research advisor Professor Upadrasta

Ramamurty who was very influential in my choice to do a PhD. His work ethics and quest for excellence I felt were second to none and I feel blessed to have been given a chance to work with him. This work would not have been possible without the unconditional support and mentorship of 2 individuals in the Gordon group, Dr. Sang Bok Kim and Dr. Luke Davis. Dr. Sang Bok Kim synthesized precursors for growth of indium based thin films and I am very grateful for the opportunity to work with him. I sincerely feel that his knowledge in synthetic chemistry and design of novel experiments is unparalleled. I was very impressed by Dr. Luke Davis for his amazing breadth of knowledge in all fields of science and thank him for his suggestions to improve the quality of my work.

I would like to thank all Gordon group members who I have been associated with and befriended during my tenure at Harvard. I had a great time getting to know them and learning from them. I extend my humble thanks to current group members who worked with me on solar cells namely Danny Chua, Robert Gustafson, Xizhu Zhao and Lauren Hartle. I closely

collaborated with Danny Chua and Xizhu Zhao for the electrical and optical measurements respectively done in this work. I specially thank Robert Gustafson for his time and expertise in designing the valve control set up for the atomic layer deposition reactor that was used in this work for the deposition of thin films. I thank other current group members Christina Chang,

Aykut Aydin, Xian Gong, Dr. Yunlong Ji, Dr. Yan Jing, Emily Kerr, Dan Pollack, Lu Sun, Dr.

Liuchuan Tong and Dr. Marc-Antoni Goulet. I found the research of every group member equally fascinating.

I would also like to thank our lab administrator, Teri Howard for being very supportive during my stay at Harvard. I take this opportunity to thank former Gordon group members;

Prof. Jaeyeong Heo, Dr. Yeung (Billy) Au, Prof. Sang Woon Lee, Dr. Norifusa Satoh, Dr.

Eugene Beh, Dr. Xudong Chen, Prof. Xinwei Wang, Dr. Bin Xi, Dr. Prasert Sinsermsuksakul,

Dr. Jing Yang, Dr. Leizhi Sun, Michael Vogel, Dr. Kecheng Li, Dr. Jun Feng, Dr. Kaixiang Lin,

Dr. Rachel Heasley, Dr. Tamara Powers and Dr. Michael Marshak. I will fondly remember Dr.

Helen Hejin Park who I collaborated with on solar cell buffer layers. Also many thanks are due to Dr. Xiabing Lou for helping me with transmission electron microscopy studies and analysis. I would like to express my gratitude to Professor Sunghwan Lee, a visiting professor in the

Gordon lab who collaborated with us on the indium oxide project.

This work would not have been possible without the contribution of our collaborators at IBM;

Prof. David Mitzi, Dr. Richard Haight, Dr. Oki Gunawan, Dr. Wei Wang, Dr. Priscilla Antunez and Dr. Yun Seog Lee. I thank them for all their intellectual contributions and help with measurements. I sincerely thank our collaborators at MIT in the Buonassisi lab, Professor Niall

Mangan, Jeremy Poindexter, and Alex Polizzotti who I worked with during my years at Harvard.

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The staff scientists at Center for Nanoscale Systems (CNS); Dave Lange, Hao-Tu (Greg) Lin,

Jason Tresback, Arthur Mcclelland, Adam Graham and Ed Macomber were extremely helpful with training me on specific scientific equipments and sample fabrication. I am very thankful to them for their time and patience.

I reserve my biggest thanks to my family and friends for their love and moral support. My parents N.K. Jayaraman and Girija Jayaraman have put in every ounce of their energy in making sure I live a blissful life and the least I could do is thank them from the bottom of my heart for their generosity. This dissertation is dedicated to You. My sister Aishwarya Jayaraman has been a pillar of strength throughout my life. My journey leading to completion of my graduate studies would be incomplete without her and I thank her for the same. I would also like to dedicate this thesis to my late paternal grandparents Nochur Krishnasastry and Rajalakshmi Krishnasastry and my maternal grandparents T.S Ramakrishnan and Akhila Ramakrishnan. I would specially like to thank my aunt Meena Iyer and uncle T.R. Sundararaman who motivated me to work earnestly and ethically towards my goals.

Lastly, I express my gratitude to all my friends for being there with me through thick and thin. I will forever cherish the moments spent with them. I specially thank Anand Moharir,

Chaitanya Chitale, Suyash (Lara) Joshi, Rupesh Kolte, Ameya Khandekar, Chinmayee Atre, Dr.

Proloy Nandi, Katia Verteletska and Nirmal Shende. This thesis is dedicated to each one of you!

Ashwin Jayaraman April 2018

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Table of Contents

Abstract iii

Acknowledgements v

Table of Contents ix

List of Figures xii

List of Tables xviii

Chapter 1 Introduction……………..………………………………………………………….1

1.1 Need for Solar Energy and Solar Cells……………...... 1

1.2 CZT(S,Se) Solar Cells and the need for Improvement in Open Circuit Voltage (Voc)...... 5

1.3 Buffer Layers in CZT(S,Se) Solar Cells………………………………...…………………...8

1.3.1 Band offset of the buffer layer from the absorber CZT(S,Se)………………………9

1.3.2 Electrical property requirements of the buffer layer...... 11

1.4 Need for a Novel Buffer Layer for CZT(S,Se) Solar Cells…...………………...…...……..12

1.5 References…………………………………………………………………………………..13

Chapter 2 Study of Atomic Layer Deposited Indium Sulfide as Electron Transport Layer for Chalcogenide CZT(S,Se) Thin Film Solar Cells...... 15

2.1 Chapter Abstract ...... 15

2.2 Introduction ...... 16

2.3 Atomic layer deposition (ALD) of Indium Sulfide ...... 20

2.3.1 Indium Sulfide Deposition using Indium Acetylacetonate...... 21

2.3.2 Atomic layer deposition of In2S3 using tris(N,N'-

diisopropylformamidinato)indium(III)……………………………………………...27

2.4 Band offset studies by ultraviolet photoelectron spectroscopy (UPS) for In2S3 grown by

ALD using tris(N,N'-diisopropylformamidinato)-indium(III)…………………...... 43

2.5 CZT(S,Se) solar cell device studies with In2S3 buffer. …………………………………….51

2.6 Double Emitter Experiments on CZT(S,Se) solar cells…………………………………….55

2.7 Conclusions…………………………………………………………………………………59

2.8 References…………………………………………………………………………………..60

Chapter 3 Atomic Layer Deposition of Binary Indium Oxide (In2O3) Thin Films Using

Amidinate Precursors, and Studies of the Structural, Electrical, and Optical Properties...65

3.1 Chapter Abstract ...... 65

3.2 Introduction ...... 66

3.3 Experiments ...... 70

3.3.1 Atomic layer deposition of indium oxide films…………………………………..…70

3.3.2 Characterization of In2O3 films…………………………………………………...…71

3.3.3 Kinetic studies of indium oxide film growth……………………………………..…72

3.4 Results and Discussion ...... 73

3.4.1 ALD window determination……………………………………………………...…73

3.4.2 Kinetic Studies by surface saturation experiments………………………………….81

3.4.3 Scanning Electron Microscopy Studies……………………………………………..88

3.4.4 X-Ray Diffraction (XRD) studies…………………………………………………...90

3.4.5 Electrical Properties…………………………………………………………………96

3.4.6 Optical Characterization…………………………………………………………...102

3.5 Conclusions ...... 106

3.6 References ...... 108

Chapter 4 Atomic Layer Deposition of In2(O,S)3 using Indium tris(N,Nʹ- diisopropylformamidinate) and its Optical and Electrical Properties…………………..…112

4.1 Chapter Abstract ...... 112

4.2 Introduction ...... 113

4.3 Experiments ...... 116

4.4 ALD Mechanisms ...... 118

4.5 Results and Discussion ...... 123

4.5.1 Photoluminescence Experiments…………………………………………………..151

4.5.2 Band offset measurements of In2(O,S)3 with CZT(S,Se): Requirements for obtaining

a high open circuit voltage…………………………………………………………….... 153

4.6 Conclusions ...... 162

4.7 References………………………………………………………………………………... 163

Chapter 5 Conclusions and future work……...... 167

Appendix A Diagram of System used for atomic layer deposition...... 172

List of Figures

1.1.1 Comparison of Lifecycle Greenhouse Gas Emissions of Various Electricity Generation Sources, World Nuclear Association (WNA), 2011.3 (Sourced and reproduced with permission from WNA)

1.2.1 A side view of a representative CZT(S,Se) solar cell device.

1.2.2 Fraction of Shockley-Queisser limit for current and voltage achieved by record solar cells9 (Reproduced with permission from authors9).

1.3.1 Schematic energy band diagram of the TCO-Buffer layer and absorber-buffer layer junctions in a CZT(S,Se) solar cell.

1.3.1.1 Electron recombination pathway when the buffer layer conduction band is significantly below that of absorber (shown in red).

2.3.1.1 X-ray fluorescence spectrum of In2S3 deposited using indium acetylacetonate at 150°C.

2.3.1.2 (a) AFM height contrast image of In2S3 sample grown using indium acetylacetonate.

2.3.1.2 (b) Height profile obtained by AFM of In2S3 sample grown using indium acetylacetonate.

2.3.1.3 XPS survey spectrum for In2S3 deposited using indium acetylacetonate at 106 °C (with inset showing high resolution of carbon at 284.8 eV on etching).

2.3.2.1 Formula of tris(N,N'-iisopropylformamidinato)indium(III).

2.3.2.2 (a) Cross-section view scanning electron microscopy (SEM) image of rough In2S3 film obtained using universal purge time of 10 seconds, (b) Cross-section view SEM image of a slightly smoother morphology obtained by increasing purging after In precursor to 30 seconds, (c) side view SEM image of a smooth In2S3 film obtained with a high purge time (45 sec) for both co-reactants.

2.3.2.3 (a) Plan view scanning electron microscope (SEM) image of rough In2S3 film obtained using purge times of 10 seconds (b) Plan view SEM image of a slightly smoother morphology obtained by increasing purging after In precursor to 30 seconds.

2.3.2.4 (a) AFM height contrast image of an In2S3 sample deposited using indium formamidinate and hydrogen sulfide at 160 °C using optmized ALD conditions.

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2.3.2.4 (b) Height profile obtained by AFM of In2S3 sample grown using indium formamidinate and hydrogen sulfide at 160 °C.

2.3.2.5 XPS depth profile of In2S3 sample deposited using indium formamidinate.

2.3.2.6 (a) Top – Selective area diffraction pattern of an In2S3 sample grown using indium formamidinate and hydrogen sulfide (b) Bottom – Bright field TEM image of pattern of an In2S3 sample grown using indium formamidinate and hydrogen sulfide at 200 °C.

2.3.2.7 The SEM cross sections of In2S3 films grown at different temperatures (a) 140 °C, (b) 160 °C, (c) 180 °C, (d) 200 °C, (e) 230 °C, (f) 250 °C.

2.3.2.8 Growth rate as a function of temperature for In2S3 films.

2.3.2.9 High resolution C 1s scan of In2S3 deposited in the range 140 °C – 250 °C.

2.3.2.10 XRD spectra of In2S3 deposited using indium formamidinate at various temperatures.

2.3.2.11 Morphology of In2S3 films grown at different temperatures (a) 160 °C , (b) 180 °C, (c) 200 °C, (d) 230 °C.

2.3.2.12 The variation in electrical properties of In2S3 with deposition temperature in the ALD window.

2.4.1 (a) UPS energy spectrum for bare CZT(S,Se) sample showing energy levels near band edge and shift in the band edge with a 800 nm pump pulse.

2.4.1 (b) Expanded view of the valence band edge region of bare CZT(S,Se) showing Fermi level is 0.54 eV above the valence band edge.

2.4.2 Comparison of the pumped and the unpumped UPS In2O3/CZTS,Se spectra.

2.4.3 Band diagram at the CZT(S,Se)/In2O3 p-n interface.

2.4.4 Comparison of the pumped and the unpumped UPS In2S3/CZTS,Se spectra.

2.4.5 Band diagram at the CZT(S,Se)/In2S3 p-n interface.

2.5.1 Schematic of a representative CZT(S,Se) thin-film solar cell fabricated with In2S3 as the electron transport layer.

2.5.2 J–V characteristics of best solar cell tested with CZT(S,Se)-In2S3 p-n junction.

2.6.1 (a) Double emitter thin-film solar cell structure schematic with CZT(S,Se) as absorber.

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2.6.1 (b) Band bending vs anneal time of the double emitter structure at 250 °C.

2.6.2 Solar cell parameters obtained from the double emitter structure.

2.6.3 J-V curves of the double emitter devices with varying thicknesses of In2S3 in comparison to the control CdS based device.

3.3.3.1 The expected surface reactions of water with a surface saturated with 1 or 2 at 250 °C.

3.4.1.1 Formulas of precursor 1 (a) and precursor 2 (b).

3.4.1.2 Growth rate as a function of temperature showing ALD window for 1.

3.4.1.3 Growth rate as a function of temperature showing ALD window for 2.

3.4.1.4 High resolution carbon 1s scans for In2O3 films grown using 1 with carbon at. % computed by fitting shown.

3.4.1.5. High resolution carbon 1s scans for In2O3 films grown using 2 with carbon at. % computed by fitting shown.

3.4.1.6 Growth rate as a function of temperature for all reported processes for ALD of In2O3. Lines represent the carbon-free ALD window and circles denote temperatures at which the product film resistance is less than 0.01 Ω·cm. Red markers indicate water as the oxygen source while gray markers indicate other oxygen sources such as ozone, hydrogen peroxide or oxygen. The indium sources 19 15 13 are compounds 1 and 2 (our study) and TMIn (A ), Et2In(Ntms2) (B ), DADI 20 21 22-23 9 24 25 (C ), In[C(NiPr2)(NiPr)2]3 (D ), InCl3 (E ), InCp ((F ), (G )), TEIn (I ), 26 27 Me2In(EDPA) (J ), In(tmhd)3 (K ).

3.4.2.1 (a) Growth rate as a function of reaction time for precursor 1 , (b) hypothetical mechanism of attachment of 1 on to a hydroxylated surface.

3.4.2.2 (a) Growth rate as a function of doses for precursor 2, (b) hypothetical mechanism of attachment of 2 onto a hydroxylated surface.

3.4.2.3 Growth rate as a function of water reaction time for 1 at 250 °C.

3.4.2.4 Growth rate as a function of water reaction time for 2 at 250 °C.

3.4.2.5 (a) Plausible reaction of water with surface saturated with 1 or 2.

3.4.2.5 (b) Hydrophobic nature of 2 v/s hydrophilic nature of 1.

3.4.2.5 (c) Steric hindrance comparison between 1 (R=H) and 2 (R=CH3).

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3.4.3.1 (a) Top view of In2O3 films grown by ALD using tris(N,N'- diisopropylformamidinato)indium(III) and (b) Cross-sectional view of In2O3 films grown by ALD using tris(N,N'-diisopropylformamidinato)indium(III). Scale bar indicates 100 nm.

3.4.3.2 (a) Top view of In2O3 films grown by ALD using tris(N,N'- diisopropylacetamidinato)indium(III) and (b) Cross-sectional view of In2O3 films grown by ALD using tris(N,N'-diisopropylacetamidinato)indium(III). Scale bar indicates 100 nm.

3.4.4.1 XRD patterns of In2O3 films on fused quartz substrate using tris(N,N'- diisopropylformamidinato)indium(III) (1); (a) 1300 cycles grown at 150-325 °C and (b) enlarged images for the XRD of films at 150, 200, 250, and 325 °C.

3.4.4.2 XRD patterns of In2O3 films on fused quartz substrate using tris(N,N'- diisopropylacetamidinato)indium(III) (1); (a) 2000 cycles grown at 150-350 °C and (b) enlarged images for the XRD of films at 150, 225, 275, and 350 °C.

3.4.5.1 For In2O3 films grown with 1 and H2O, (a) Resistivity as a function of deposition temperature, (b) Electron mobility as a function of deposition temperature, (c) Electronic carrier concentration as a function of deposition temperature, and (d) Electron mobility as a function of electronic carrier concentration.

3.4.5.2 For In2O3 films grown with 2 and H2O, (a) Resistivity as a function of deposition temperature, (b) Electron mobility as a function of deposition temperature, (c) Electronic carrier concentration as a function of deposition temperature, and (d) Electron mobility as a function of electronic carrier concentration.

3.4.6.1 (a) Optical transmittance v/s wavelength for In2O3 deposited using 1 at different temperatures with thicknesses as shown.

3.4.6.1 (b) Optical transmittance v/s wavelength for In2O3 deposited using 2 at different temperatures with thicknesses as shown.

2 3.4.6.2 Band gap determination using plot of (ᾀhὐ) v/s photon energy for In2O3 films grown using 1 (a) and 2 (b).

3.4.6.3 Carrier concentration as a function of optical band gap for In2O3 films grown using 1.

4.3.1 Formula for indium tris(N,N’-diisopropylformamidinate).

4.4.1 (a) Indium precursor reaction with a thiolated surface (case 1).

4.4.1 (b) Reaction of indium amidinate saturated surface with incoming H2S (case 1).

4.4.1 (c) Indium precursor reaction with a single thiol group on the surface (case 2).

4.4.1 (d) Reaction of indium amidinate saturated surface with hydrogen sulfide (case 2).

4.4.1 (e) Indium precursor reaction with a hydroxylated surface (case 3).

4.4.1 (f) Reaction of indium amidinate saturated surface with incoming H2O (case 3).

4.4.1 (g) Indium precursor reaction with a single hydroxyl group on surface (case 4).

4.4.1 (h) Reaction of indium amidinate saturated surface with H2O (case 4).

4.4.1 (i) Bridging of thiol groups on an indium-thiolated surface.

4.4.1 (j) Conversion of indium oxide bonds to indium sulfide bonds by reaction with H2S.

4.5.1 CZT(S,Se) device stack used for device studies.

4.5.2 Representative XPS In 3d (a), S 2p (b) and O 1s (c) scans for an In2(O,S)3 composition deposited at 200 °C using 0.91 fraction of water containing subcycles. Positions of the In 3d5/2, In 3d3/2, O 1s, S 2p1/2 and S 2p3/2 peaks are typically seen at a binding energy of ~ 444.5 eV, ~ 452 eV, ~530 eV, 162.7 eV and ~ 161.6 eV respectively.

4.5.3 Chemical composition of the In2(O,S)3 films as a function of the fraction of water- containing subcycles.

4.5.4 Anion ratio in the films and bulk hydroxyl group content as a function of the fraction of water-containing subcycles.

- 4.5.5 (a) O 1s spectrum for the In2(O,S)3 composition with 49.1 at % O, 3.9 % OH , 9 at. % S (films obtained with a 0.91 fraction of water containing subcycles).

- 4.5.5 (b) O 1s spectrum for the In2(O,S)3 composition with 33 at % O, 1.6 % OH , 23 at. % S (films obtained with a 0.75 fraction of water containing subcycles).

- 4.5.5 (c) O 1s spectrum for the In2(O,S)3 composition with 9 at % O, 1.3 % OH , 48 at. % S (films obtained with a 0.33 fraction of water containing subcycles).

4.5.5 (d) O 1s spectrum for the In2(O,S)3 composition with 4 at. % O, below detection limit OH-, 53 at. % S (films contain a water fraction of 0.2).

4.5.6 Elemental composition inside a representative In2(O,S)3 film deposited with a 0.25 fraction of water containing subcycles, showing homogeneity of composition.

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4.5.7 Ideal growth rate/supercycle vs. real growth rate/supercycle for In2(O,S)3 grown using conditions shown. 4.5.8 Transmission electron microscopy scans of indium sulfide (a1,a2) at 200 °C, indium oxide (b1,b2) at 200 °C, and a representative intermediate In2(O,S)3 composition (18 at. % O , 38 at. % S) at 200 °C (c1,c2) deposited on Si3N4 membranes.

4.5.9 XRD analyses of In2S3, In2O3 and In2(O,S)3 thin films.

4.5.10 The top view morphology of (a) In2(O,S)3 (18 at. % Oxygen, 38 at. % Sulfur), (b) In2O3 and (c) In2S3 films grown at 200 °C. Scale bar indicates 100 nm.

4.5.11 (a) Cross section of In2(O,S)3 film with atomic composition In = 38 %, O = 49.1 %, OH- = 3.9 %, and S = 9 % and (b) Cross section of In2(O,S)3 film with atomic composition In = 42.4 %, O = 33 %, OH- = 1.6 %, and S = 23 %.

4.5.12 (a) Height contrast image under AFM for In2(O,S)3 with 49.1 at. % O, 9 at. % S.

4.5.12 (b) 3 dimensional height contrast view under AFM for In2(O,S)3 with 49.1 % O, 9 % S.

4.5.12 (c) Topographical height profile under AFM for In2(O,S)3 with 49.1 at. % O, 9 at. % S.

4.5.13 (a) Height contrast image under AFM for In2(O,S)3 with 18 at. % O, 38 at. % S.

4.5.13 (b) 3 dimensional height contrast under AFM for In2(O,S)3 with 18 at. % O, 38 at. % S.

4.5.13 (c) Topographical height profile for In2(O,S)3 with 18 at. % O, 38 at. % S.

4.5.14 (a) Carrier concentration in the In2(O,S)3 as a function of atom % sulfur.

4.5.14 (b) Electron mobility and resistivity of In2(O,S)3 films as a function of atom % sulfur.

4.5.14 (c) Carrier concentration in the In2(O,S)3 as a function of atom % oxygen.

4.5.14 (d) Electron mobility and resistivity of In2(O,S)3 films as a function of atom % oxygen.

4.5.15 % Transmittance of In2(O,S)3 thin films in the UV-Visible region of the light spectrum.

4.5.16 (a) Direct band gap approximation for determining band gap of indium based compounds.

4.5.16 (b) Indirect band gap approximation for determining band gap of indium based compounds.

4.5.17 High res. XPS scans of Cu (a) and Zn (b) after 40 s etch on In2(OS)3/CZT(S,Se).

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4.5.1.1 Photoluminescence spectra of In2O3 and In2(O,S)3 samples obtained with a 532 nm laser excitation.

4.5.2.1 (a) Core level spectra of Cu 2p for the layered sample and (b) Core level spectra of In 3d for the layered sample.

4.5.2.2 (a) Valence band edge for the CZT(S,Se) sample showing band edge extrapolation to 0 counts and (b) Cu 2p core level spectra for the CZT(S,Se) sample.

4.5.2.3 (a) Valence band edge for the In2(O,S)3 sample showing band edge extrapolation to 0 counts and (b) In 3d core level spectra for the In2(O,S)3 sample.

- 4.5.2.4 Band diagram at the interface of In2(O,S)3 (In – 42.5 %, OH - 1.5 %, S - 38 %, O - 18 %) and CZTS,Se (relative core levels and valence band levels determined by XPS, band gap by UV-Vis spectrophotometry).

4.5.2.5 Valence band offset variation with change in oxygen content in In2(O,S)3.

4.5.2.6 Conduction band offset variation with change in oxygen content in In2(O,S)3.

4.5.2.7 Variation in carrier concentration as a function of atom % oxygen in In2(O,S)3 compositions.

A.1 Diagram of system used for atomic layer deposition of In2O3, In2S3 and In2(O,S)3 thin films.

xviii

List of Tables

2.3.1.1 XRF counts of indium and sulfur in In2S3 grown using indium acetylacetonate.

2.3.1.2 Composition of In2S3 film deposited at 106 °C using indium acetylacetonate.

2.3.2.1 Thickness and growth rates of In2S3 films at different deposition temperatures.

2.3.2.2 Hall measurement results of In2S3 samples deposited in ALD window.

3.4.1 Classification of indium precursors for In2O3 ALD on the basis of ALD onset temperature and ALD window temperature range.

3.4.4.1 Crystallite size as a function of deposition temperature for In2O3 grown using 1.

3.4.4.2 Peak area as a function of deposition temperature for In2O3 grown using 1.

3.4.4.3 Crystallite size as a function of deposition temperature of In2O3 using 2.

3.4.4.4 Peak area as a function of deposition temperature of In2O3 using 2.

4.5.1 In2(O,S)3 composition as a function of dosing ratio of water containing to hydrogen sulfide containing cycles.

4.5.2 Device results we obtained from the In2(O,S)3 buffer layer incorporated solar cells.

xix

To Amma, Appa and my sister Aishwarya

Chapter 1

Introduction

1.1 Need for Solar Energy and Solar Cells

The global energy demand is growing rapidly due to population explosion and technological progress. It is thus of paramount importance to reduce reliance on fossil fuels and opt for a reliable, everlasting and cost effective energy source to meet the future energy demand.

Solar energy amongst various renewable sources of energy is most promising and most abundant. The Sun delivers energy that is 4 orders of magnitude higher than the rate at which humans use and produce energy1, thus there is enormous potential. The solar energy is mainly used in 3 ways 1) Conversion to electricity in a photovoltaic (PV) system, 2) Solar heating and cooling (SHC) and 3) Concentrated solar power (CSP). Studies performed in this thesis are focused on materials used in a particular type of thin film PV technology. Currently, PV systems made of solar cells are a growing part of electrical energy mix in Europe, United States, China, and many other countries around the world. In 2014, for example, 6-11 % of the yearly electricity generated was from PV systems in Greece, Italy and Germany2 (IEA-PVPS, 2015).

Grid connected and stand-alone applications provide power over a wide range from a tenth of a watt to hundreds of megawatts.2 At the end of 2015, the global cumulative capacity of installed

PV systems was 227 gigawatt, which is equivalent to 280 coal fired power plants.2

Previously development of PV technology was propelled by need for robust and durable electricity generation for space applications like satellites. Currently, however, the use of PV is 1

driven mainly by need to reduce carbon dioxide emissions. Amongst the various energy sources,

PV systems have extremely low CO2 emission per GWh of electricity generated, namely around

3 85 tonnes of CO2 equivalent per GWh on average as seen in Figure 1.1.1. In comparison, coal

3. fired power plants emit around 888 tonnes on average of CO2 equivalent per GWh. Thus PV systems are very important to mitigate global warming due to amplified greenhouse effect.

Figure 1.1.1. Comparison of Lifecycle Greenhouse Gas Emissions of Various Electricity Generation Sources, World Nuclear Association (WNA), 2011.3 (Sourced and reproduced with permission from WNA)

Major economies like Europe have set targets of 20 % reduction of CO2 emission by

2020 and an 80-95 % reduction in greenhouse gases by 2050 compared to 1990 emission levels2

(European Commission 2011). A mix of sustainable energy technologies including a large

volume of solar PV solar energy is thus inevitably required to meet these targets. In order to bridge the difference between the large potential of solar energy and its relatively small utilization as of now, a majority of the research in PV systems needs to be focused on improving the energy conversion efficiency while using a low cost manufacturing process with inexpensive materials. The materials used in a solar cell should preferably consist of non-toxic and earth abundant elements for large scale manufacture of environmentally friendly PV. Crystalline silicon PV cells are the most commonly used solar cells in solar panels available in the market, accounting for 85 % of the PV cell sales in 2011.4 Crystalline silicon cells have a very high efficiency up to 25 % for single crystal cells and over 20 % for multicrystalline cells on lab scale.4 Industrial silicon based solar modules achieve maximum efficiencies of 18-22 % under standard test conditions.4

Despite the high efficiencies of PV technology involving silicon, a caveat associated with the same is that the purification to single crystalline grade silicon is very expensive. Silicon is also an indirect band gap material with relatively weak absorption. Thus a thickness of crystalline silicon wafers in the range of 200 μm is employed for nearly complete absorption of sunlight.

With the objective to reduce cost, thinner material with strong absorption in the form of amorphous silicon was developed as an alternative. The low growth rate of thin film amorphous silicon limited low cost large scale production and the maximum efficiencies reached were significantly lower than crystalline silicon. Thus there was a need for a low cost highly efficient thin film solar technology. Over the course of time, CdTe and Cu(In,Ga)(S,Se)2 (CIGS) evolved as predominant thin film technologies in the industry. CdTe has reached the highest efficiency of

22.1 % (First Solar) while CIGS has hit efficiencies of 22.6 % (ZSW) at lab scale. There are still

certain issues with employing CIGS and CdTe on a larger scale in residential and utility scale solar. Because of the limited earth abundance and/or toxicity of constituent elements in CIGS and CdTe, these are not permanent solutions to our need.5 Also, on the cost front; there has been a change over the years with module costs now being comparable for multicrystalline silicon and thin film technologies of CdTe and CIGS. On comparing the total installed system cost in dollar/W, it is seen that multicrystalline silicon costs around 1.08 $/Watt (for 10 m2 modules) while the CdTe technology (with lower module efficiency) costs around 1.3 $/Watt for the same module area.6 Thus it is imperative to look for lower cost thin film technologies which are based on earth abundant, non toxic elements. One of the other reasons for favoring thin films over crystalline silicon solar cells is the light weight nature of the thin film cells which employ around

1μm thickness of absorber material in comparison to around 200 μm for crystalline cells. Light weight PV cells are essential for example in applications involving the internet of things, where sensors deployed to track and report the state of mobile objects like vehicles and animals, need to be powered.7 Batteries do not serve this purpose because of the following. A lithium ion battery with maximum allowed footprint of 1 cm2 on a sensing device has an estimated storage capacity of only 20 mWh.7 Given estimates of power consumption for processing and data storage (100 μW), sensing (1 mW) and data transmission (5mW) and intermittent operation of 1 hr each day, such a battery will last only for 3 days.7 Thus in order for the sensing device to last for years, one needs to use light weight PV cells.

Given all the aforementioned considerations, kesterite Cu2ZnSn(SxSe1-x)4 (CZT(S,Se)) has emerged as one of the favorable thin film technologies for PV applications. At 12.6 % conversion efficiency8, it is the most efficient earth abundant/non-toxic material based thin film

absorber currently available. A thickness of 2 μm suffices to absorb most of the sunlight and thus

CZT(S,Se) cells are light-weight.

1.2 CZT(S,Se) Solar Cells and the need for Improvement in Open Circuit

Voltage (Voc)

CZT(S,Se) (chemical formula - Cu2ZnSn(SxSe1-x)4 ) is a semiconductor which contains earth abundant and non toxic elements like Cu, Zn, Sn and S and hence is environmentally friendly. CZTS (chemical formula - Cu2ZnSnS4) has a direct band gap of 1.5 eV and a large absorption coefficient of 104 cm-1 and hence is an ideal material for absorption of sunlight.

Addition of selenium results in lowering the band gap of CZTS up to around 1 eV. The resultant

CZT(S,Se) material has a higher current density and significantly enhanced efficiency compared to CZTS. A representative high efficiency CZT(S,Se) cell with all its components including front

(Ni/Al) and back (Mo) contacts is shown in Figure 1.2.1. A 2μm thick CZT(S,Se) layer usually deposited by hydrazine solution processing or vacuum deposition is used. Cadmium sulfide

(CdS) deposited by chemical bath deposition is the electron transport layer. Tin doped indium oxide (ITO) is sputtered on as transparent conducting oxide. A zinc-tin oxide (ZTO) layer is present between ITO and CdS to prevent sputtering damage on CdS layer.

Figure 1.2.1. A side view of a representative CZT(S,Se) solar cell device.

The maximum CZT(S,Se) solar cell efficiency obtained (12.6 %)8 is still far below the

Shockley Queisser (SQ) limit of 28 %. While a maximum of 81 % of the SQ limit has been extracted from a CZT(S,Se) device in terms of the short circuit current , less than 50 % of the SQ limit has been extracted as far as the open circuit voltage - fill-factor product is concerned9 (See

Figure 1.2.2). These values are poor when compared to competing thin film technologies like

CdTe for which 95 % of the current and 68 % of the Voc-fill factor product have been extracted.

The high current extraction indicates relatively complete light absorption and collection of charge carriers. The low open circuit voltage in CZT(S,Se) is because of a number of factors.

One of them is the presence of bulk point defects, primarily Cu/Zn antisite defects10 which results in recombination of photo-generated electrons. Recombination at the grain boundaries also results in lower efficiencies seen, a factor that is mitigated by air annealing and hence passivating the grain boundaries with tin oxide.11

Figure 1.2.2. Fraction of Shockley-Queisser limit for current and voltage achieved by record solar cells9 (Reproduced with permission from authors9).

Another reason for the low Voc is the interfacial recombination at the absorber-electron transport layer interface, due to mismatch in conduction band energies. Thus the choice of an ideal electron transport layer (also referred to as buffer layer) with a tunable conduction band level is very important for improving the open circuit voltage of the CZT(S,Se) cell. One of the motivations of this work was to reduce the Voc deficit (difference between absorber band gap and measured Voc) in CZT(S,Se) solar cells by optimizing the choice of the buffer layer.

Lack of conformality of the electron transport layer on the absorber, unpassivated defects and elemental interdiffusion at the interfaces also negatively affect the Voc of the CZT(S,Se) solar cell.

1.3 Buffer Layers in CZT(S,Se) Solar Cells

Cadmium sulfide (CdS) is currently used as the electron transport layer/ buffer layer for best performing CZT(S,Se) solar cells.8 Figure 1.3.1 shows a schematic energy band diagram of the TCO(transparent conducting oxide) - buffer layer and absorber - buffer layer junctions in a

CZT(S,Se) solar cell. There are a number of properties that are critical in an n-type buffer layer used in combination with p type CZT(S,Se). . The buffer layer should be free of impurities like carbon which act as recombination centers. The CZTS,Se - buffer layer junction should be pristine and free of elemental interdiffusion during growth and anneals.

Figure 1.3.1. Schematic energy band diagram of the TCO-Buffer layer and absorber-buffer layer junctions in a CZT(S,Se) solar cell.

In addition, the buffer layer should have a wide band gap to transmit a majority of the photons to the absorber as shown in Figure 1.3.1. This results in more light capture and increased short circuit current from the solar cells.

1.3.1 Band offset of the buffer layer from the absorber CZT(S,Se)

The conduction band level of the buffer should be close to that of the CZT(S,Se) absorber as shown in Figure 1.3.1. There are two possible electron pathways, 1 and 2 represented by red arrows in Figure 1.3.1. Pathway 1, through the conduction band of the buffer layer, helps in the collection of electrons from the solar cell. Pathway 2 resulting in recombination of electrons into valence band of absorber layer needs to be prevented.

If the conduction band of the buffer is very high in comparison to the absorber (> 0.4 eV spike), photo-generated electrons are not transferred to the electron transport layer resulting in loss of photocurrent and poor fill factor. On the other hand, if the conduction band of the buffer is significantly below that of the absorber layer (high cliff), the photo-generated electrons recombine into the valence band of the absorber layer as shown in Figure 1.3.1.1. This results in reduced open circuit voltage of the solar cell.

Figure 1.3.1.1. Electron recombination pathway when the buffer layer conduction band is significantly below that of absorber (shown in red).

Thus it is very important to tune the conduction band of the buffer to match that of the

CZT(S,Se) absorber layer . A small positive conduction band offset (spike) is induced by stronger band bending and thus a larger hole barrier in comparison to a small negative cliff barrier.

In theory12, a small negative cliff barrier corresponds to small interfacial band bending which leads to a large number of holes recombining with electrons on either side of the interface.

The enhanced recombination current due to this will manifest in a low Voc of the device. A small positive spike barrier on the other hand corresponds to a larger built in potential with a suppressed interface recombination due to there being fewer holes available to recombine with interface electrons. This translates into obtaining a relatively higher Voc.

1.3.2 Electrical property requirements of the buffer layer

The buffer layer should exhibit a high electron mobility to transfer photo-generated electrons from CZT(S,Se) absorber layer to the TCO. The carrier concentration of the buffer should not be significantly below that of the absorber. Under low doping of the buffer, for both the cliff and spike band offset cases, device performance goes down. This is because there is insufficient carrier collection in the absorber with a poor band bending and hence a low electric field. Thus a highly resistive buffer is not preferred. A highly resistive buffer also leads to high series resistance of the solar cell and poor fill factor.

In case of a high carrier concentration/doping in the buffer compared to the absorber, one gets most of the band bending in the absorber. The low density of interfacial holes limits forward

12 current and hence a higher Voc is obtained. Numerical simulations in the literature have shown that a spike/positive conduction band offset of 0.1 - 0.3 eV in combination with around 1x1017 -

1x1018 electrons is optimal for the buffer layer in CdTe solar cells.12

The recombination of electrons into the valence band of the absorber is accelerated if there are trap states at the absorber-buffer interface. The recombination velocity is directly proportional to the interface defect density. The interface recombination velocity SIF is given by the formula12

SIF = (Nd,IF) x (σn,p) x (vth)

-12 where Nd,IF is interface defect density, σn,p is capture cross sections of electrons and holes (~10

2 7 cm ) and vth is thermal velocity of electrons and holes (~10 cm/s). A high SIF results in low open circuit voltage.

The method of deposition of the buffer layer is thus very important to get a defect free interface and hence a high open circuit voltage from the solar cell.

1.4 Need for a Novel Buffer Layer for CZT(S,Se) Solar Cells

Cadmium sulfide (CdS), the electron transport/buffer layer in best performing CZT(S,Se) cells8 has a number of issues. CdS besides being a toxic material, has a low direct band gap of

2.4 eV and hence is not most optimal as a buffer layer as it absorbs light in the blue region. This results in absorption losses and reduced short circuit current of the solar cell. CdS is prevalently deposited by chemical bath deposition (CBD), a solution based technique in which the chemistry involved results in uncontrollable hydroxyl groups and other contaminants in the material. This leads to poor control over properties of the CZT(S,Se) – CdS interface. Also this solution based technique of growing CdS buffer layer does not provide a high yield process on coupling with vacuum based methods which are used in the industry for depositing other layers in a solar cell.

CdS also possibly does not have the most optimum band offset with CZT(S,Se) as the open circuit voltage (Voc) is significantly less by around 617 mV in comparison to the band gap of CZT(S,Se).8 Varying band positions of CZT(S,Se) surface are obtained due to change in deposition and anneal conditions and it is beneficial to use a buffer with a tunable conduction band offset. CdS does not exhibit a tunable band offset with CZT(S,Se). We also look for alternatives to CdS because the Cd2+ shows a propensity to diffuse at the absorber- buffer interface, especially when the absorber is grown after the buffer layer in the superstrate configuration of the solar cell. In2S3 with an indirect band gap of 2.1 eV emerged as one of the better alternatives as a buffer layer for CZT(S,Se). The indirect band gap of In2S3 results in lower

absorption losses. In2S3 was initially investigated because it functioned well with another quaternary semiconductor copper indium gallium sulfide selenide (CIGS-Se). Although indium is a relatively rare element, back of the envelope calculations reveal that the amount of indium needed for 20 nm buffered CZT(S,Se) devices to power the world is of the same order of magnitude as what is available in the earth’s crust. Although In2S3 is a promising material, method of growth predominantly used is CBD which as discussed above is not ideal. This thesis explores the growth of In2S3 by atomic layer deposition (ALD) using a newly developed precursor in Chapter 2. ALD has potential to grow pristine interfaces and films free of hydroxyl groups and contaminants. This work goes on to explore ALD of indium oxide in Chapter 3 on grounds that alloying indium oxide with the indium sulfide helps develop a growth recipe for

In2(O,S)3. Indium oxysulfide studied in chapter 4 is motivated because of the ability to tune conduction and valence band positions of the buffer material by changing the sulfur to oxygen ratio. This can have direct influence over improvement of Voc of the CZT(S,Se) solar cell devices as discussed in section 1.2 and 1.3.1.

1.5 References

1. Crabtree G. W. and Lewis N. S., Solar energy conversion, Physics Today, 2007,37-42.

2. Reinders A.; Verlinden P.; Sark W.V.; Freundlich A., Photovoltaic Solar Energy from Fundamentals to Applications – John Wiley & Sons, Ltd, West Sussex, United Kingdom, 2017 edition, Pg. 3.

3. Comparison of Lifecycle Greenhouse Gas Emissions of Various Electricity Generation Sources. World Nuclear Association Report Website. http://www.world- nuclear.org/uploadedFiles/org/WNA/Publications/Working_Group_Reports/comparison_of_l ifecycle.pdf (accessed 25 April 2018).

4. Crystalline Silicon Photovoltaics Research. Office of Energy Efficiency & Renewable Energy Website. https://www.energy.gov/eere/solar/crystalline-silicon-photovoltaics- research (accessed April 25 2018).

5. Andersson B.A., Materials availability for large-scale thin film photovoltaics, Prog. Photovoltaics 2000, 8(1), 61–76.

6. Horowitz K.A.W.; Fu R.; Sun X.; Silverman T.; Woodhouse M.; Alam M.A., An Analysis of the Cost and Performance of Photovoltaic Systems as a Function of Module Area - Technical Report NREL/TP-6A20-67006 April 2017.

7. Haight R.; Haensch W.; Friedman D., Solar-powering the Internet of Things, Science 2016, 353, 6295, 124-125.

8. Wang, W.; Winkler, M. T.; Gunawan, O.; Gokmen, T.; Todorov, T. K.; Zhu, Y.; Mitzi, D. B., Device Characteristics of CZTSSe Thin-Film Solar Cells with 12.6% Efficiency, Advanced Energy Materials 2014, 4 (7), 1301465.

9. Polman A.; Knight M.; Garnett E.C.; Ehrler B.; Sinke W.C., Photovoltaic materials: Present efficiencies and future challenges, Science 2016, 352, 6283

10. Scragg, . . S. Choubrac, L. Lafond, A. Ericson, T. Platzer rkman, C. A low- temperature order-disorder transition in Cu2ZnSnS4 thin films, Appl. Phys. Lett. 2014, 104, 041911.

11. Sardashti, K. ; Haight, R. , Gokmen, T.; Wang W. ;Chang L.Y. ,Mitzi D.B.; Kummel A.C. ,Impact of Nanoscale Elemental Distribution in High‐Performance Kesterite Solar Cells, Advanced Energy Materials 2015, 5 (10), 201402180.

12. Song, T.; Kanevce, A.; Sites, J. R., Emitter/absorber interface of CdTe solar cells. Journal of Applied Physics 2016, 119 (23), 233104.

Chapter 2

Study of Atomic Layer Deposited Indium Sulfide as Electron Transport Layer for Chalcogenide

CZT(S,Se) Thin Film Solar Cells

2.1 Chapter Abstract

Atomic layer deposition (ALD) of indium sulfide was studied using tris(acetylacetonato)indium(III) (also called indium(III) acetylacetonate) and tris(N,Nʹ- diisopropylformamidinato)indium(III) (also called indium(III) formamidinate) precursors. The latter was identified as a more favorable precursor to obtain smooth, carbon-free films in a desirable temperature range for application as an electron transport layer on Cu2ZnSn(S1-xSex)4

(CZT(S,Se)). X-ray photoelectron spectroscopy (XPS) and growth rate studies reveal that In2S3

(grown with indium(III) formamidinate) has an atypical carbon-free ALD window, reasons for which are highlighted. The structural characterization of the films done by X-ray diffraction

(XRD) and transmission electron microscopy (TEM) showed that the films are polycrystalline in the ALD window temperature range of 160 °C to 200 °C. Hall effect measurements indicate carrier concentrations in the right range for an electron transport layer and an increase in electron mobility with temperature. We demonstrate that the band offset of In2S3 with CZT(S,Se), as

obtained by ultraviolet photoelectron spectroscopy (UPS), is around the optimal value. Despite this fact, a representative thin-film solar cell employing our In2S3 has non-optimal device properties; the reasons for the poor performance are discussed and the need for In2(O,S)3 as electron transport layer is proposed.

2.2 Introduction

There has been extensive research for an electron transport layer that has a conduction band match with CZT(S,Se) and forms a defect-free interface with the same absorber layer.

Aforementioned criteria need to be fulfilled to obtain a high open circuit voltage from the solar cell device. CdS the currently widely used buffer for CZT(S,Se) cells has its share of issues as discussed later in this section. Alternatives to CdS like ZnO and ZnS have been investigated in the past.1 ZnS has very high spike conduction band offset of 1.1 eV and hence resulted in a very low current (0.001 mA/cm2) and open circuit voltage (0.69 mV) from the CZT(S,Se) solar cell.1

ZnO on the other hand, despite having a band alignment of ± 0.1 eV resulted in a poor solar cell efficiency of 2.46 % and low voltage of 348 mV.1 This indicated that even a small negative band offset could lead to reduced open circuit voltage. Zn(O,S) with varying S:O ratios deposited by

ALD in our lab was also tested in combination with CZT(S,Se) but high open circuit voltages compared to CdS were not obtained (unpublished work by Dr. Helen Hejin Park). A chemically bath deposited indium sulfide tested in a CZT(S,Se) solar cell, showed promising results in terms of a reasonably high open circuit voltage (435 mV) and short circuit current (29.2 mA/cm2 ).1

This motivated our work on ALD of indium sulfide for CZT(S,Se) solar cells. Indium is a relatively rare metal but buffer layers constitute only 20 to 30 nm of a device which is microns in

total thickness. Thus a relatively small quantity of indium is consumed in depositing a indium sulfide buffer layer.

2 Indium sulfide (In2S3) is a III-VI semiconductor material. It has found importance in

3 4 5 optoelectronic , photoelectric and photovoltaic applications due to its stable chemical composition6, photoconductivity7 and luminescent properties8 under ambient conditions. It is an n-type semiconductor with the optical band gap in the range of 2.1-2.3 eV. There is still controversy about whether it has indirect band gap or direct band gap, with empirical results

9-11 12-13 showing fitting consistent with both the former and the latter . The In2S3 grown in our work is empirically measured to be an indirect band gap material. In2S3 has three allotropic forms, namely, α-In2S3 (cubic structure between 420 °C and 754 °C), β-In2S3 (body-centered

14 tetragonal structure below 420 °C), and γ-In2S3 (trigonal structure above 754 °C). There is some controversy as to the existence of these phases with a later phase diagram15 claiming that there are only two phase variants of In2S3, one below 415 °C (Space group I41/amd) and the other above 750 °C (space group P m1). Of the mentioned variants, β-In2S3 has wide applications due to its very stable defective spinal structure, its photosensitivity and electrical characteristics like electron mobility and electron concentration.

β-In2S3 is evolving as one of the novel alternative buffer layers used in combination with chalcogenide absorber layers. To reiterate, cadmium sulfide (CdS), although a widely used electron transport layer in thin-film solar cells based on the chalcogenide semiconductors CuIn(1- x)GaxSe2 (CIGS) and CZT(S,Se), is a toxic material. As mentioned in chapter 1, CdS grown by chemical bath deposition has an uncontrollable amount of hydroxyl groups and other contaminants due to solution based nature of deposition. This affects the electrical properties of

CdS as well as the interface with CZT(S,Se). The thin-film solar cell world was in search of a

low-cost, contaminant-free, nontoxic and industrially applicable alternative to CdS. From an environmental and health standpoint, β-In2S3 is a good alternative to CdS, considering many countries have imposed restrictions on solar PV market share for Cd-containing solar cells. Thin film β-In2S3 also possesses other physical characteristics needed for an electron transport layer, for example, the right range of electron concentration (1016-1018, as seen in this work and previous literature16), a reasonably high electron mobility, and a wide enough indirect band gap

(around 2.1 eV) to transmit photons to the absorber. β-In2S3 buffer layer has been used

17 successfully as an effective nontoxic substitute for CdS in CIGS solar cells. Thin-film In2S3 has been deposited by various methods in the past including ALD, sputtering, ionic layer deposition, physical vapor deposition, ultrasonic spray pyrolysis, chemical spray pyrolysis, atomic layer epitaxy, metal organic chemical vapor deposition, chemical bath deposition, and electrodeposition.18 To date, the highest power conversion efficiencies for CIGS solar cells with various methods of depositing In2S3 include efficiencies of 16.4 % by ALD, 15.7 % by CBD,

19 14.8 % by evaporation, and 12.2 % by sputtering. One of the reasons why In2S3 works reasonably well as a buffer layer for CIGS is because there is no conduction band discontinuity at the interface with the CIGS absorber layer.20

Although a majority of previous work has looked at the In2S3–CIGS combination in a solar cell, there is precedence as mentioned for In2S3 electron transport layers for CZT(S,Se)

1,21 solar absorbers. In2S3 was tested in a CZT(S,Se) solar cell because CZT(S,Se) is predominantly a quaternary semiconductor like CIGS with zinc and tin replacing expensive and rare indium and gallium. The highest power conversion efficiency yet obtained for a CZT(S,Se) device with an In2S3 buffer layer (deposited by chemical bath deposition, CBD) and an MgF2 anti-reflection coating was 7.6 %.1 This is lower compared to the best efficiency obtained with

CdS as a buffer layer which is 12.6 %.43 Some of the reasons for this observation are believed to be as follows. In2S3 deposited by chemical bath deposition is not stoichiometric, and contains uncontrolled number of hydroxyl groups. These groups can result in unintentionally altering the electrical and optical properties compared with pristine In2S3 material. The current obtained from

1 the CZT(S,Se) device with In2S3 is higher than that with CdS because the indirect band gap of

In2S3 results in less buffer-layer absorption of blue photons. The device voltage on the other hand is lower for the In2S3–CZT(S,Se) device as opposed to the CdS–CZT(S,Se) one possibly due to non-optimal carrier concentration of In2S3. A too-low carrier concentration results in a high series resistance of the device, while really high carrier concentrations result in recombination of electrons with the holes in the absorber layer.

Over the years, atomic layer deposition (ALD) has evolved as a very robust technique for depositing stoichiometric thin films conformally in various applications.22 It may be a useful alternative to CBD as it should avoid unwanted hydroxyl groups in the film. ALD also produces interfaces which are relatively more defect free when compared to CBD. Also, it is a vacuum- based technique and hence when coupled with other vacuum-based techniques used in making solar cells, produces a high yield. In this chapter we explore ALD of In2S3 with 2 different precursors to replace the CBD In2S3 and CBD CdS in CZT(S,Se) solar cells.

An ALD reactor with capability to deposit conformal, stoichiometric In2S3 films (verified by Rutherford backscattering spectrometry, RBS) on CZT(S,Se) was developed. The deposition conditions were tuned to obtain carbon-free (by x-ray photoelectron spectroscopy), smooth films.

The average growth rate of In2S3 was measured as a function of temperature, and structural characterization was performed by scanning electron microscopy, x-ray diffraction, and transmission electron microscopy. Ultraviolet photoelectron spectroscopy was used to determine

the valence and the conduction band offsets with CZT(S,Se). A representative CZT(S,Se) device was fabricated with ALD In2S3 as the buffer layer deposited at 160 °C and device parameters were measured and rationalized.

2.3 Atomic layer deposition (ALD) of Indium Sulfide

There have been a number of ALD precursors previously explored for the deposition of

In2S3. Trimethylindium(III) showed no growth when pulsed sequentially with hydrogen sulfide

16 (H2S). Cyclopentadienyl indium(I) also did not produce self-saturated growth of In2S3 in

16 combination with H2S. This behavior was expected considering cyclopentadienyl ligands were difficult to remove even when mechanistic studies were done by sequential dosing with H2O and

23 16 O2 . Other precursors studied, such as tris(2,2,6,6-tetramethyl-3,5-heptanedionato)indium(III)

i and indium guanidinate (In(N Pr)2CNMe2)3 in combination with H2S failed because of low vapor pressures or evaporation rates below the decomposition temperatures of the precursors.

Tris(N,Nʹ-diisopropylacetamidinato)indium(III),16 a precursor from the amidinate class of reactive compounds, was the most promising precursor for In2S3 growth but we discounted it because the reported source temperature of 190 °C exceeded the thermal limits of our ALD valves (rated only to 150 °C). Also, the vapor pressure in a cylinder of this precursor would be relatively low at the lower temperature of 150 °C, requiring longer ALD cycle time for saturated growth and hence lower growth rate of In2S3. Additionally, In2S3 grown using tris(N,Nʹ- diisopropylacetamidinato)indium(III)16 has very poor mobility of 0.5 cm2/V·sec at temperatures near and below 200 °C, around which we are interested in growing the electron transport layer.

We focus our attention instead on two precursors, tris(2,4-pentandionato)indium(III)

(indium acetylacetonate), which has been previously studied,24-27 and tris(N,N'- iisopropylformamidinato)indium(III), recently developed for this application in our lab.

2.3.1 Indium Sulfide Deposition using Indium Acetylacetonate

The binary reaction for the indium sulfide buffer layer systhesis has been previously

24-27 explored with indium acetylacetonate. The ALD reaction of the precursor with H2S proceeds as follows:24

2 In(acac)3 + 3 H2S → In2S3 + 6 acacH

The split ALD reactions for the indium sulfide deposition are:24

InSH* + In (acac)3 → InS-In(acac)*2 + acacH (Seq.A)

In(acac)* + 3 H2S → InSH* + acacH (Seq.B)

where the * indicates surface adsorption, and acac = acetylacetonate.

We perform the above reactions at 150 °C in an ALD tube furnace, in closed-valve mode with the indium acetylacetonate precursor held at 120 °C. The precursor has a low vapor pressure and hence its flow is assisted by nitrogen flowing into the reactor. Indium acetylacetonate is dosed in with an exposure of 2 torr·s (3 doses) followed by purging with nitrogen for 30 seconds.

This is further followed by 1 dose of 4 % hydrogen sulfide in nitrogen with an exposure of 2.4

torr·s. A total of 1000 cycles were deposited on thermal oxide substrates (300 nm of wet oxidation SiO2 on Si) placed inside the tube furnace.

Film thickness measured by X-ray reflectivity experiments is ~18 nm. This implies that the growth rate is 0.18 Å/cycle, similar to the growth rate seen in previous work.24 The composition is uniform from inlet to outlet of the tube furnace. The indium and sulfur fluorescence counts obtained using X-ray fluorescence (XRF) spectroscopy are almost the same at the inlet and the outlet (see Figure. 2.3.1.1). Table 2.3.1.1 shows the XRF counts.

Figure 2.3.1.1. X-ray fluorescence spectrum of In2S3 deposited using indium acetylacetonate at 150 °C.

Table 2.3.1.1. XRF counts of indium and sulfur in In2S3 grown using indium acetylacetonate.

Sample S Kα counts In Lα counts

Outlet 152 370

Inlet 156 382

Figure 2.3.1.2 (a) and (b) show the atomic force microscopy (AFM) image and the height profile, respectively, of a sample deposited under the discussed deposition conditions. The RMS roughness of the sample is around 11.2 nm while the thickness is 18 nm. The reason for high roughness could be sparse nucleation on the thermal oxide sample.

Figure 2.3.1.2. (a) AFM height contrast image of In2S3 sample grown using indium acetylacetonate.

Figure 2.3.1.2. (b) Height profile obtained by AFM of In2S3 sample grown using indium acetylacetonate.

Films of this high a roughness are not ideal. If grown on solar absorber layers, such highly rough electron transport layers would likely separate the front transparent conducting oxide (TCO) from the absorber by a very small thickness (of the electron transport layer) in certain regions.

This could result in tunneling of electrons at such locations. If the film is not continuous, the

TCO ends up in direct contact with the absorber and shunting effects are likely.

We did not optimize the deposition conditions initially to try and reduce the roughness of the films, because there were other issues with the film growth rate and composition to be addressed. An aspect of indium acetylacetonate based ALD process for In2S3 which is deleterious for solar cells is the presence of carbon in the films obtained. Below in Figure 2.3.1.3 are the

XPS high resolution carbon 1s spectra for an In2S3 deposition on thermal oxide sample at 106 °C.

Table 2.3.1.2 shows the composition measured inside the film.

24000

Surface 22000 1 min etch 2 min etch 20000

18000

16000 Intensity (CPS) Intensity 14000

12000

10000 280 285 290 295 300 Binding Energy (eV)

Figure 2.3.1.3. XPS high resolution C 1s spectrum for In2S3 deposited using indium acetylacetonate at 106 °C (elemental carbon peak seen at 284.8 eV inside film on etching).

It is interesting to note that there was no significant silicon 2p peaks seen from the survey scans

(not shown) even after 2 min etch, which implies there is good coverage. But it is difficult to confirm on the basis of a single point small area XPS measurement if the thin film is covering the substrate uniformly at all locations. Especially with the detection limit of XPS being 1 at. % it is difficult to rule out small pinholes.

Table 2.3.1.2. Composition of In2S3 film deposited at 106 °C using indium acetylacetonate.

Element Binding energy Composition inside film

In 3d 445.10 eV 40.8 at. %

S 2p 162.2 eV 45.7 at. %

C 1s 285.02 eV 9.9 at. %

O 1s 532.10 eV 2.9 at. %

It is seen from Figure 2.3.1.3 and Table 2.3.1.2 that there exists around 10 at. % carbon in the In2S3 films deposited at 106 °C using indium acetylacetonate. This carbon is likely not due to condensation of the precursor because the sublimation temperature of the precursor is 80 °C as used in previous work.24 Thermal decomposition of the precursor is also unlikely, as this

25 precursor has been used at 125 °C previously for growing In2S3. The presence of carbon is most likely the result of incomplete ALD half reactions and the protonated ligands not being completely detached post reaction. The carbon percentage decreases with temperature but there is still around 3 at. % carbon in the film at 150 °C. Films grown at 150 °C have been previously seen to contain 5.1 at. % carbon.24 Thus it is clear that it is difficult to grow carbon-free films of

In2S3 at low temperatures in the range 106-150 °C using indium acetylacetonate. Presence of

carbon in the In2S3 films (meant to be used as electron transport layers) is deleterious because carbon potentially acts as a recombination center and thus results in loss of current in a solar cell.

Also the growth rate of the films is very low (~0.2 Å/cycle) at these temperatures. A cycle is of the order of a minute. We tried sonicating the substrates in acetone and isopropyl alcohol (to remove the organics) and performing a UV-ozone treatment (to generate surface hydroxyl groups), but the growth rates remained similar.

It is important that we observe a high growth rate by ALD for electron transport layers like In2S3 grown on absorbers such as CZT(S,Se). In the case of slow growth at moderate to high temperatures, it is possible to have interfacial diffusion between elements of the absorber and the electron transport layer in a thin-film solar cell. This affects the integrity of the p-n junction of the cell, and electron collection by the electron transport layer suffers.

Despite these drawbacks, initial solar cell devices were made with 20 nm thick In2S3 buffer layers deposited at a substrate temperature of 150 °C and indium acetylacetonate precursor source temperature of 120 °C. Solar cell devices were fabricated by Richard Haight at

IBM. The device structure of Mo/CZT(S,Se)/In2S3/ZnO/ITO/Ni/Al was employed. Results showed device efficiencies up to 2.5 %. The shunt resistance values and fill factor values were low compared to devices using the champion CdS electron transport layer. Ultraviolet photoelectron spectroscopy measurements (done by Richard Haight at IBM) indicated that there is minimal band bending at the interface between the ALD In2S3 and CZT(S,Se), leading to poor device performance.

Considering the rough nature of the films we obtained with indium acetylacetonate, the slow growth rate of the films, the presence of significant amount of carbon even at 150 °C, and the resultant poor performance of the initial devices fabricated, we opted to consider an

alternative precursor (which belongs to the amidinate class of compounds) for depositing indium sulfide as an electron transport layer for chalcogenide thin-film solar cells.

2.3.2 Atomic layer deposition of In2S3 using tris(N,N'- diisopropylformamidinato)indium(III)

Figure 2.3.2.1. Formula of tris(N,N'-iisopropylformamidinato)indium(III).

The above indium precursor (Figure 2.3.2.1), alternatively called indium formamidinate, was developed in our lab by Dr. Sang Bok Kim. It belongs to a class of precursors, amidinates, which are highly reactive compounds. The sublimation temperature was found to be around 90

°C. The first ALD experiments we did were with the precursor at 125 °C to get enough vapor pressure to saturate the surface of all substrates along the reactor tube. 4 % hydrogen sulfide in nitrogen was used as co-reactant gas. Thermal oxide (300 nm SiO2 on Si) substrates were used.

We employed a standard dosing recipe for indium sulfide with 1 dose of indium precursor (at 125 °C) with nitrogen assist and 1 dose of co-reactant hydrogen sulfide passed into the reactor tube furnace at 160 °C. These reactants when dosed alternately in ALD mode should ideally saturate the surface. The initial chamber evacuation time, dosing time, exposure time and purge time used for the indium precursor and hydrogen sulfide were 15 s, 3 s, 1 s, 10 s and 15 s,

1 s, 1 s, 10 s respectively. The dosing time, purge time and chamber evacuation time (closed- valve mode ALD) were varied in order to get smooth films. Initially, with a high dosing time of

3 seconds for the indium precursor and purge time of 10 seconds after the indium precursor dosing, a needle like morphology was obtained as shown in Figure 2.3.2.2 (a). The high dosing time without adequate purge resulted in partial CVD mode reaction with the reactants not being completely evacuated and reacting in the gas phase. This resulted in non-uniform and needlelike growth. In order to alleviate this we reduced the dosing time to 1 second and compensated it with the exposure time of 3 seconds for the indium precursor. We also increased the purge time after the indium precursor to 30 seconds. On making these changes, the exposure of indium precursor is 3 torr·s and that of H2S is unchanged at 2.7 torr·s. We obtained a thin film that is smoother as evident from the cross section in Figure 2.3.2.2(b). The top view of the film also shows that morphology is smoother and the transition is seen from Figure 2.3.2.3(a) to (b). Although the film is smoother, there are some minor serrations which remain. We further went on to increase the purge time after the hydrogen sulfide cycle from 10 seconds to 45 seconds to get rid of the volatile reaction byproduct. This resulted in very smooth films as shown in Figure 2.3.2.2(c).

Thus there are quite a few advantages of using an optimally low dosing time, precursor temperature and optimally high exposure time of the precursor. They are as follows

1. Lesser CVD Component in the film preventing non-uniform growth

2. Lower Precursor consumption rate

3. Higher time for reaction meaning there is enough time for atoms to move around

in the film and hence give a smoother film

Figure 2.3.2.2. (a) Cross-section view scanning electron microscopy (SEM) image of rough In2S3 film obtained using universal purge time of 10 seconds, (b) Cross-section view SEM image of a slightly smoother morphology obtained by increasing purging after In precursor to 30 seconds, (c) side view SEM image of a smooth In2S3 film obtained with a high purge time (45 sec) for both co-reactants.

(a) (b)

200 nm 200 nm

Figure 2.3.2.3. (a) Plan view scanning electron microscope (SEM) image of rough In2S3 film obtained using purge times of 10 seconds (b) Plan view SEM image of a slightly smoother morphology obtained by increasing purging after In precursor to 30 seconds.

We measured the roughness of the films obtained in Figure. 2.3.2.2(c) by AFM. Figure 2.3.2.4(a) and 2.3.2.4(b) show the AFM height contrast image and height profile respectively of an In2S3 sample deposited at 160 °C using optimized deposition conditions as described.

Figure 2.3.2.4. (a) AFM height contrast image of an In2S3 sample deposited using indium formamidinate and hydrogen sulfide at 160 °C using optmized ALD conditions.

Figure 2.3.2.4. (b) Height profile obtained by AFM of In2S3 sample grown using indium formamidinate and hydrogen sulfide at 160 °C.

The RMS roughness of the films was found to be 1.5 nm. It is thus clear that In2S3 films deposited using indium formamidinate on thermal oxide do not have nucleation issues on thermal oxide at 160 °C. There are intermittent spikes in the height profile of the film likely indicative of periodic sharp crystal facets of In2S3. The thickness of the film obtained is ~122 nm. Thus thickness is very high in comparison to the roughness of the film. This indicates that pinholes are unlikely in the film. The growth rate per minute is 0.5 Å which is greater in comparison to that for In2S3 grown using indium acetyl acetonate which has a growth rate of

0.12 Å per minute. The growth rate per cycle of In2S3 grown with indium formamidinate is 0.81

Å/cycle while that for In2S3 grown with indium acetyl acetonate is much lower at 0.2 Å/cycle

The high growth rate of In2S3 grown using indium formamidinate is beneficial as discussed earlier.

The composition of the In2S3 films deposited at 160 °C was measured by XPS and the depth profile of elements obtained is shown in Figure 2.3.2.5. XPS was done with a Thermo

Scientific K-Alpha spectrometer with a Al Kα X-ray source (1486.6 eV), 12 kV electron beam, and argon sputtering gun. The high resolution scans were performed using a 500 eV sputtering energy with 30 s of sputtering per level.

Figure 2.3.2.5. XPS depth profile of In2S3 sample deposited using indium formamidinate.

It is seen from the depth profile that the there is some surface carbon and oxygen on the sample which disappears with argon ion etching at 500 eV. The sample is mainly composed of indium and sulfur. The correct stoichiometry 2:3 of In:S is not seen because this profile was obtained from survey scans. High resolution scans with curve fitting likely gives compositions closer to the correct stoichiometry. XPS is not the optimal technique to measure film compositions in the bulk. XRD results of these samples though show peaks representative of β

In2S3 .This is discussed later on in this section.

In order to check crystallinity, we performed transmission electron microscopy (TEM) on a 30 nm thick In2S3 sample deposited at 200 °C on 50 nm thick Si3N4 membranes supported on

0.5 × 0.5 mm Si grids purchased from Ted Pella, Inc.

(a)

(b)

Figure 2.3.2.6. (a) Top – Selective area electron diffraction pattern of an In2S3 sample grown using indium formamidinate and hydrogen sulfide (b) Bottom – Bright field TEM image of pattern of an In2S3 sample grown using indium formamidinate and hydrogen sulfide at 200 °C.

The selective area electron diffraction (SAED) pattern of an In2S3 sample grown at 200

°C is shown in Figure 2.3.2.6(a). This sample is polycrystalline, as a diffraction pattern with rings is seen. The bright field image in Figure 2.3.2.6 (b) shows lattice fringes and a darkness contrast which demarcates the grains in the sample. The SAED pattern was indexed by first

determining the interplanar d-spacing using distances of the spots on the rings from the center.

Knowing the lattice parameters (a,c) of the body centered tetragonal structure of In2S3 and the relation between the interplanar spacing and the lattice parameters (1/d2 = (h2 + k2 )/a2 + l2/c2 ), we determined the planar Miller indices.

We thus could deposit crystalline In2S3 films with indium formamidinate and hydrogen sulfide. With the above preliminary information the ALD window of In2S3 was explored by deposition at different temperatures. Depositions were performed in the temperature range 140

°C -250 °C with the source temperature at 125 °C. Basic properties like thickness and crystallinity were measured by SEM (Field-emission scanning electron microscopy, ZEISS,

Ultra-55) and X-ray diffraction (XRD, D8 DISCOVER) respectively. The composition of the films was measured by XPS and the carbon-free ALD window was determined. Van der Pauw and Hall effect measurements using a custom Hall effect rig with a sourcemeter were performed to obtain electrical properties of the films in the ALD window.

The SEM cross-sections of In2S3 films at different temperatures is shown in Figure

2.3.2.7.

(a) (b)

(e) (f)

Figure 2.3.2.7. The SEM cross sections of In2S3 films grown at different temperatures (a) 140 °C, (b) 160 °C, (c) 180 °C, (d) 200 °C, (e) 230 °C, (f) 250 °C.

The cross sections of the films show that the films are smooth only at lower temperatures (140-

180 °C). The higher temperature films show sharper serrated features perhaps representative of higher crystallinity of the films (as seen later by XRD). The thickness of the above films measured by SEM are tabulated in Table 2.3.2.1. The thickness per cycle is not constant in any

temperature range which essentially means that In2S3 has an atypical ALD window with changing growth rate.

Table 2.3.2.1. Thickness and growth rates of In2S3 films at different deposition temperatures.

Temperature Thicknesses Growth rate

(number of cycles)

140 °C (3000 cyc) 259.2 nm 0.864 Å/cycle

160 °C (1500 cyc) 122 nm 0.81 Å/cycle

180 °C (2000 cyc) 127.1nm 0.64 Å/cycle

200 °C (3260 cyc) 115.5 nm 0.35 Å/cycle

230 °C (3500 cyc) 134.1 nm 0.38 Å/cycle

250 °C (3500 cyc) 135.2 nm 0. 39 Å/cycle

0.9 Carbon free ALD window 0.8

0.7

/cycle) o

0.6

0.5 Growth Rate (A Rate Growth 0.4

0.3

120 140 160 180 200 220 240 260 280 Temperature (oC)

Figure 2.3.2.8. Growth rate as a function of temperature for In2S3 films.

It is noted that the growth rate of the films decreases with increase in deposition temperature. Figure 2.3.2.8 shows that the growth rate decreases approximately linearly in the range 160 °C – 200 °C. There could be a couple of reasons for the growth rate drop with temperature. The first being that there could be reaction of neighboring thiol groups to release

H2S and form a sulfide bridge happening at higher temperature. The sulfide bridge does not take further part in ALD growth and thus the number of sites available for reaction decreases with rise in temperature resulting in growth rate reduction. Previous work16 has shown a similar mechanism for In2S3 grown using indium acetamidinate which belongs to the amidinate family of precursors. The second hypothesis for the growth rate reduction is possible etching effect of hydrogen sulfide. Hydrogen sulfide being acidic in nature, under sufficient exposure could etch away the film and thus reduce the actual growth per cycle.

The net ALD reaction that occurs with indium formamidinate and hydrogen sulfide as co- reactants is hypothesized similar to previous work16 as

* * -(SH) x + In(amd)3 → SxIn(amd) 3-x + xH(amd) (1)

* * SxIn(amd) 3-x + (3-x)H2S → SxIn(SH) 3-x + (3-x)H(amd) (2)

The above hypothesized mechanism16 is analogous to the hydroxyl group mechanism seen in ALD processes like that involving trimethylaluminum and water, although that reaction

22 occurs only at rather high temperatures. The repetition of reactions (1) and (2) leads to In2S3 growth by atomic layer deposition. Ideally after each sequence of above cycles there should be

no residual amidinate ligands. If the reaction proceeds in the aformentioned clean manner, there is no carbon or nitrogen residue seen in the film.

To determine carbon-free ALD window, we characterized the presence of carbon in films deposited at all temperatures by XPS. The high resolution carbon spectra for In2S3 films grown at different temperatures is shown in Figure 2.3.2.9.

50000

o 40000 250 C 3 % N, 11 % C o 230 C 2 % N, 9 % C o 200 C o CPS 30000 180 C o 160 C

20000 o 4 % N, 18 % C 140 C

10000 278 280 282 284 286 288 290 292 294 296 298 Binding Energy (eV)

Figure 2.3.2.9. High resolution C 1s scan of In2S3 deposited in the range 140 °C – 250 °C.

It is seen from Figure 2.3.2.9 that In2S3 films are carbon free only in the range 160 °C to

200 °C. At 140 °C, there is 18 at. % carbon seen likely because of protonated ligands not being completely detached during the ALD reaction. Condensation is unlikely because the sublimation temperature of the precursor is well below the deposition temperature. The high amount of carbon, 9-11 at. %, seen at 230 °C and 250 °C is likely because of the surface decomposition of

the reaction product/byproduct. We did not carefully measure the transition temperature from a no-carbon film to a carbonaceous composition, but it occurs between 200 and 230 °C. The XPS results when compared with growth rate variation show that ~160 °C to 200 °C is the carbon-free

ALD window for In2S3 (using indium formamidinate), where growth rate drops approximately linearly with temperature. This ALD window is almost the same range compared to that of the analogous precursor tris(N,Nʹ-diisopropylacetamidinato)indium(III)16 which depicts an ALD window from 150-225 °C. The XRD spectra of In2S3 films of nominal thickness of 100 nm at different temperatures is shown in 2.3.2.10. It is seen that samples at 140 °C are mostly amorphous with very few reflections at high angles seen. This likely due to the presence of 18 at.

% carbon in the film which disrupts crystal growth. The main peaks corresponding to the

28 tetragonal In2S3 (103), (109) and (206) start appearing at 160 °C (Pattern JCPDS 25-0390). The tetragonal structure is expected for In3+ at a growth temperature below 415 °C15 and has been

16, 24, 29-31 noted in previous literature for In2S3 grown by ALD and other methods. The peaks become bigger and sharper at higher temperatures (200 °C onward) in comparison to 160 and

180 °C indicating larger grain size and higher crystallinity. Of the main reflections (103) dominates indicating these crystal planes are mainly aligned parallel to the surface. The films are textured with only 6 peaks predominantly visible.

4500 (103) (109) (206) (312) (318) (4012) (507)

4000

o 3500 250 C

3000

o 2500 230 C

2000 CPS

o 1500 200 C o 1000 180 C o 160 C 500 o 140 C 0

10 20 30 40 50 60 70 2

Figure 2.3.2.10. XRD spectra of In2S3 deposited using indium formamidinate at various temperatures.

The samples at 230 °C and 250 °C despite containing 9 and 11 at. % carbon respectively are highly crystalline as seen. This could mean that carbon is not impeding crystalline growth in the case of In2S3 deposition using our indium precursor. Unfortunately depositions at such high temperatures give highly resistive films (as verified by Hall measurement) and hence In2S3 grown at those temperatures cannot be employed in a solar cell.

In order to map the morphology change with temperature we performed plan view SEM of the samples in the ALD window. Figure 2.3.2.11 (a) through (d) shows the morphology of the films.

(a) 160 °C (b) 180 °C

Figure 2.3.2.11. Morphology of In2S3 films grown at different temperatures (a) 160 °C , (b) 180 °C, (c) 200 °C, (d) 230 °C.

The morphology of poorly crystallized films at lower temperatures 160 °C and 180 °C is very different from those at higher temperatures. The feature size seen at 160 °C and 180 °C is smaller compared to the well crystallized films at 200 °C. The morphology and feature size do not change significantly beyond 200 °C.

It is extremely important for the electron transport layer to have a reasonably high electron mobility. We performed Hall and Van der Pauw measurements of In2S3 samples deposited in the ALD window. A custom Hall rig with a Keithley sourcemeter was used.32 The resultant resistivity, electron concentration, and electron mobility are shown in Table 2.3.2.2.

Table 2.3.2.2. Hall measurement results of In2S3 samples deposited in ALD window.

Temperature Resistivity Electron Hall Thickness

(Ω·cm) concentration Mobility

(cm-3) (cm2/V·s)

160 °C In2S3 272 ± 11.8 1.2 ± 0.1e16 1.7 ± 0.2 122 nm

180 °C In2S3 115 ± 4.5 1.1 ± 0.1e16 4.9 ± 0.5 127 nm

200 °C In2S3 28.9 ± 1.2 1.2 ± 0.11e16 17.7 ± 1.7 115 nm

Resistivity (ohm.cm) Mobility (cm2/V.sec) 3 300 Carrier concentration (/cm ) 2x1016 20

250 18 16 200 14 12 150 10 8 100 1x1016 6

50 4 2 0 0 160 170 180 190 200 Temperature (oC)

Figure 2.3.2.12. The variation in electrical properties of In2S3 with deposition temperature in the

ALD window.

It is seen from Table 2.3.2.2 and Figure. 2.3.2.12 that the electron concentration remains mostly constant with change in temperature while the mobility increases from 1.7 cm2/V·s at 160 °C to

17.7 cm2/V·s at 200 °C.

In the substrate configuration of the thin-film solar cell, where the electron transport layer is deposited over the absorber, we would prefer lower deposition temperature to prevent interfacial diffusion. But lower temperatures such as 160 °C result in a film with lower mobility.

Thus it is a trade off between a high temperature of deposition (and corresponding high mobility) and a p-n junction with minimal interdiffusion.

It is seen that the resistivity of the sample drops with temperature, with the lowest resistivity of 28.9 Ω·cm seen at 200 °C. A high series resistance of the solar cell is obtained if the resistivity of In2S3 is high. Thus we prefer low resistivity In2S3. In the superstrate configuration of the solar cell (where the electron transport layer is deposited first), the deposition temperature of the electron transport layer is not a big concern. We can thus potentially use 200 °C for deposition which provides us with a high-mobility, low-resistivity electron transport layer.

2.4 Band offset studies by ultraviolet photoelectron spectroscopy (UPS) for

In2S3 grown by ALD using tris(N,N'-diisopropylformamidinato)-indium(III)

CZT(S,Se) samples described in this section were deposited by Dr. Wei Wang at IBM

Thomas J Watson Research Center, Yorktown Heights. The UPS measurements and analysis described in this section were performed by Dr. Richard Haight at IBM. The electron transport layer depositions and analysis of results were done by the author of this work.

The valence band edge of the CZT(S,Se) absorber material and the different electron transport layers used in combination with it can be measured by femtosecond laser pump/probe ultraviolet photoemission spectroscopy (UPS). This is described in previous works,1, 33-35 by Dr.

Richard Haight. The UPS measurements were performed at IBM Watson labs on bare

CZT(S,Se), In2S3 on CZT(S,Se), and In2O3 on CZT(S,Se) samples. 40 nm of In2S3 and In2O3

o were deposited at 160 C onto CZT(S,Se). The deposition conditions for In2S3 are described in detail in section 2.3.2, while the deposition conditions for In2O3 are described in chapter 3.

Femtosecond laser pulses (split into probe only and pump) were selectively focused onto the sample in a UHV analysis system for UPS measurements. The band bending at the CZT(S,Se) buffer heterojunction was extracted by focusing 1.55 eV (800 nm) pump pulses on the probe area. The electron-hole pairs thus created in the CZT(S,Se) screen the electric field in the depletion region and hence flattens bands resulting in an energy shift of the UPS spectrum.

. The as-grown CZT(S,Se) films were transferred from a nitrogen glove box into the

UHV analysis chamber with minimal air exposure of around 10 min. The samples are heated to

170 °C in the analysis chamber to remove moisture and other impurities resulting from air exposure. The UPS energy spectrum for the unpumped CZT(S,Se) and the 800 nm pumped

CZT(S,Se) are shown in Figure. 2.4.1(a). The Cu-3d, S-3p, and Zn-3d energy levels are seen near the band edge. It is seen that the spectrum shifts in binding energy by 170 meV on pumping

CZT(S,Se) with an 800 nm pulse. This is equivalent to a 170 meV upward band bending in

CZT(S,Se) and this upward direction of band bending is in agreement with other work.36 The surface of CZT(S,Se) thus is in slight accumulation, which is in contrast to the downward band bending seen in conventional p-type materials. The extrapolation of the band edge to the binding energy axis for the pumped CZT(S,Se) sample gives the valence band to Fermi level difference as shown in Figure. 2.4.1 (b). On extrapolating the band edge to 0 intensity it is seen that Fermi level is 0.54 eV from the valence band edge, which is mid gap for our CZT(S,Se) sample (the band gap is 1.1 eV).33

Figure 2.4.1. (a) UPS energy spectrum for bare CZT(S,Se) sample showing energy levels near band edge and shift in the band edge with a 800 nm pump pulse.

Figure 2.4.1. (b) Expanded view of the valence band edge region of bare CZT(S,Se) showing Fermi level is 0.54 eV above the valence band edge.

The results shown in Figure 2.4.1 (a) and Figure 2.4.1 (b) help us in plotting the

CZT(S,Se) side of the band diagram including band bending and valence band position with respect to Fermi level.

In regards to CZT(S,Se) samples coated with layers of indium-based buffer layers, we first discuss band offset of an indium oxide thin film deposited on CZT(S,Se), using the same precursor indium formamidinate and H2O at 160 °C (The deposition conditions are described in chapter 3). Figure 2.4.2 shows UPS spectra of In2O3/CZT(S,Se) under probe only (black spectrum) and 800 nm pump (red spectrum) conditions. The In2O3 layer thickness was 40 nm as measured by X-ray reflectivity measurements, which results in minimal band bending in the electron transport layer. Since the pump is 800 nm (1.55 eV), there is no absorption in the In2O3

(band gap 3.6 eV) expected. The energy shift seen in Figure 2.4.2 from the unpumped to the pumped spectrum is the band flattening of CZT(S,Se) giving a band bending of 200 meV. This shows that there is small rectification in the In2O3/CZT(S,Se) p-n junction. The valence edge of the pumped spectrum is extrapolated to 0 intensity (as seen in Figure 2.4.2) to obtain the valence band position of In2O3 with respect to the Fermi level in the In2O3/CZT(S,Se) film. This valence band difference from the Fermi level is 2.55 eV. We had previously determined the CZT(S,Se) has a valence band to Fermi level difference of 0.54 eV in Figure 2.4.1 (b). Using the valence band to Fermi level difference for bare pumped CZT(S,Se) and pumped In2O3/CZT(S,Se) one can extract the valence band offset between CZT(S,Se) and In2O3 at the p-n junction. For In2O3 as obtained from Figure 2.4.1 (b) and Figure 2.4.2, the valence band offset = ~2.6 eV - ~0.5 eV =

2.1 eV. Using the valence band offset and known band gaps of CZT(S,Se) (1.1 eV) and In2O3

(3.6 eV), one can construct the entire band diagram at the CZT(S,Se)-In2O3 p-n junction, which is shown in Figure 2.4.3.

Figure 2.4.2. Comparison of the pumped and the unpumped UPS In2O3/CZTS,Se spectra.

Figure 2.4.3. Band diagram at the CZT(S,Se)/In2O3 p-n interface.

It is seen from Figure 2.4.3 that the conduction band offset at the p-n interface after the bands are flattened is 0.4 eV. Conduction band offset of 0.4 eV although significant is not too high to prevent photoelectrons from being transported from the absorber layer to the electron transport layer. But the indium oxide thin film has a very high carrier concentration, >1019/cm3, and thus is not an ideally suited buffer layer. An electron transport layer with too high a carrier concentration results in recombination of electrons with holes in the absorber layer, especially when there are a high number of recombination centers like defects.

Let us now discuss the band offset of an indium sulfide thin film deposited on

CZT(S,Se), using the indium formamidinate at 160 °C . Figure 2.4.4 shows UPS spectra of

In2S3/CZT(S,Se) under probe only (black spectrum) and 800 nm pump (red spectrum) conditions. The In2S3 layer thickness was 40 nm as measured by X-ray reflectivity measurements. The band bending in the In2S3 layer, though minimal, can be extracted by a 400 nm (3.1 eV) pump scan which results in light absorption by the In2S3 buffer (band gap= 2.1 eV) and electron excitation. The bands are thus flattened and band bending can be extracted in the

In2S3 layer. We have employed a pump pulse of 800 nm (1.55 eV) and thus there is no absorption in the In2S3 (band gap 2.1 eV) expected. The energy shift seen in Figure 2.4.4 from the unpumped to the pumped spectrum is therefore the band flattening of CZT(S,Se), giving a band bending of 375 meV.

Thus there is a significant rectification in the In2S3/CZT(S,Se) p-n junction. This generates an electric field to sweep the photo-generated electrons from the CZT(S,Se) into the

In2S3 layer. The valence edge of the pumped spectrum is extrapolated to 0 intensity (as seen in

Figure 2.4.4) to obtain the valence band position of In2S3 with respect to Fermi level in the

In2S3/CZT(S,Se) film combination. This valence band difference from the Fermi level is 1.2 eV.

CZT(S,Se) as seen earlier has a valence band to Fermi level difference of 0.54 eV in Figure 2.4.1

(b).

Using the Fermi level to valence band difference for bare pumped CZT(S,Se) and pumped In2S3/CZT(S,Se) one can extract the valence band offset between CZT(S,Se) and In2S3 at the p-n junction. For In2S3 as obtained from Figure 2.4.1 (b) and Figure 2.4.4, the valence band offset = ~1.2 eV - ~0.5 eV = 0.7 eV. Using the valence band offset and known band gaps of

CZT(S,Se) (1.1 eV) and In2S3 (2.1 eV), one can construct the entire band diagram at the

CZT(S,Se)-In2S3 p-n junction which is shown in Figure 2.4.5.

Figure 2.4.4. Comparison of the pumped and the unpumped UPS In2S3/CZTS,Se spectra.

Figure 2.4.5. Band diagram at the CZT(S,Se)/In2S3 p-n interface.

Figure 2.4.5 shows that the conduction band offset at the p-n interface after the bands are flattened in CZT(S,Se) is a spike of 0.29 eV. In previous literature on CdTe solar cells37 it has been noted that a small conduction band spike offset of 0.1 to 0.3 eV helps maintain good cell efficiency despite high interfacial defect density. A spike of 0.29 eV is thus not too high to prevent photoelectron transport from CZT(S,Se) to In2S3. The carrier concentration in In2S3 at

160 °C is not as high as indium oxide, as seen in section 2.3.2, so there should be minimal recombination with holes in CZT(S,Se). Thus it should serve as a reasonable buffer layer in a solar cell. Considering the above results, we fabricated a representative solar cell with CZT(S,Se) as the absorber and In2S3 as the buffer layer. The results are discussed in the next section.

2.5 CZT(S,Se) solar cell device studies with In2S3 buffer

CZT(S,Se) samples described in this section were deposited by Dr. Wei Wang at IBM

Thomas J Watson Research Center, Yorktown Heights. The solar cell fabrication and device measurements described in this section were performed by Dr. Wei Wang and Dr. Richard

Haight at IBM. The In2S3 depositions and analysis of results were done by the author of this work.

A thin-film solar cell was fabricated in the substrate configuration as shown in Figure

2.5.1. A 2 μm CZT(S,Se) film was deposited on molybdenum coated glass by a hydrazine solution processing approach, targeting a Cu-poor, Zn-rich condition (Cu/(Sn+Zn) = 0.8,

Zn/Sn=1.1).38-40 A thermal anneal in excess of 500 °C was used to increase the grain size (to reduce recombination at grain boundaries).41 The molybdenum back contact was 500 nm thick with a thin layer of Mo(S,Se)2 forming at the interface with CZTS,Se during the thermal annealing process. The oxide layer on CZT(S,Se) was removed by dipping it in a solution with

28 wt. % ammonium hydroxide and water in ratio 1:4 for 5 minutes. After this, a thin (30 nm) layer of In2S3 was deposited by atomic layer deposition at 160 °C using process parameters as discussed in section 2.3.2. We did not choose higher temperatures for In2S3 deposition due to the poor growth rates and the increased chances of interdiffusion of elements like copper and zinc from the absorber, across the interface. The sample was exposed to air for a period of 2 days after ALD before transfer to the sputter chamber where the next layers were deposited. Thus the

In2S3 could potentially contain diffused oxygen.

Figure 2.5.1. Schematic of a representative CZT(S,Se) thin-film solar cell fabricated with In2S3 as the electron transport layer.

The solar cell was completed with a 10 nm ZnO/50 nm ITO sputtered layer as a TCO,38, 41-42 followed by a 2 μm thick Ni/Al (Ni as adhesion layer) top metal contact deposited by electron beam evaporation. The cell device area of 0.42 cm2 was defined by mechanical scribing.

Figure 2.5.2 shows the J-V characteristics of the solar cell under illumination for the best performing cell. The cell parameters obtained are as follows.

Efficiency of the device = 5.8 %,

Fill factor of the device = 36 %

Open circuit Voltage = 433.6 mV

Short circuit current = 37.2 mA/cm2

Series Resistance = 6.2 ohm.cm2

Shunt Resistance = 41.9 ohm.cm2

Figure 2.5.2. J–V characteristics of best solar cell tested with CZT(S,Se)-In2S3 p-n junction.

It is thus seen that the efficiency of the device is less than half the best performing

CZT(S,Se) devices with CdS as the buffer layer which give efficiencies of around 12.6 %.43 One of the main reasons for this is the high series resistance of the CZT(S,Se)-In2S3 solar cell device.

The series resistance is 6.2 Ω·cm2 which is high compared to the best CZT(S,Se)–CdS solar cell

2 43 (0.72 Ω·cm ). The high series resistance is likely contributed by a combination of bulk In2S3 resistance, and interfacial contact resistance of the In2S3 – CZT(S,Se) interface and the In2S3–

ITO interface. The bulk resistivity of the In2S3 layer deposited at 160 °C although high (272

Ω·cm), it is a very thin layer (30 nm). Thus its net series resistance contribution is 272x30x10-7

= 8.16 x 10-4 Ω·cm2, which is quite small. Thus we believe that the interfacial resistances mentioned (In2S3 – CZT(S,Se) and In2S3– ITO) contribute mainly to the series resistance of the solar cell. The In2S3– ITO is targeted to be an ohmic interfacial contact and thus the contact

44 resistance of ITO on In2S3 can be measured by transmission line model (TLM), where contact

pads of ITO are patterned on an In2S3 thin film layer and I-V measurements are made as a function of variable distance between pads. The contact resistance between In2S3 and CZT(S,Se) can be potentially determined in the following manner:

a) Deposit a thin film structure with the configuration Mo-CZT(S,Se)-In2S3-In and

perform I-V measurements (pass current probing the Mo, In contacts and measure the

voltage).

b) Measure resistance as a function of varying thickness of In2S3 in the aforementioned

thin film structure and determine resistance ΔR at 0 thickness of In2S3 by

extrapolation.

c) ΔR is a combination of bulk resistance of CZT(S,Se) (measurable), interfacial

resistance of CZT(S,Se)- In2S3 contact, contact resistance of In on In2S3 (measurable

by TLM patterning44) and contact resistance of Mo on CZT(S,Se) (measurable by

44 TLM patterning ). Thus CZT(S,Se)- In2S3 contact resistance is determinable.

2 The shunt resistance is also very low for the CZT(S,Se)-In2S3 device (41.9 Ω·cm ) in comparison to CZT(S,Se)–CdS cell (> 600 Ω·cm2 ).43 This implies that there are possible shunts through the In2S3 layer. The higher series resistance and lower shunt resistance in the In2S3 buffered device result in a poor fill factor of 36 %.

2 The short circuit current of the CZT(S,Se)-In2S3 cell is 37.2 mA/cm and is higher than

2 the CZT(S,Se)–CdS cell (35.2 mA/cm ) as expected. The In2S3 layer with 2.1 eV indirect band gap transmits more light into the absorber than the CdS layer with a 2.4 eV direct band gap. This results in more photons getting absorbed and converted to electricity in the In2S3-buffered device thus giving a higher short circuit current.

The open circuit voltage of the best CZT(S,Se)-In2S3 cell is 433.6 mV which is lower than the 513.4 mV obtained with a CdS buffer. Reasons for this are unclear as the band offset measured independently by UPS showed relatively good band alignment between In2S3 and

CZT(S,Se). The conduction band offset of CdS in itself is likely non optimal with CZT(S,Se) as the open circuit voltage (Voc) is significantly less by around 617 mV in comparison to the band gap of CZT(S,Se).43

As mentioned earlier there may be a small amount of oxygen diffused in the ALD In2S3 from air as the samples were exposed to air for a significant period of time before TCO and metal contact layers were deposited. The addition of oxygen might have helped in increasing the open circuit voltage. In previous work, chemical bath deposited (CBD) In2S3 in combination

1 with CZT(S,Se) gave 7.6 % efficiency devices with a Voc of 435 mV. The CBD In2S3 contains

5-10 % uncontrolled amount of oxygen due to the chemistry of this solution method. Though helpful, the effect of the presence of oxygen on electronic properties is variable and not controllable in this case. Thus we propose addition of oxygen in a controlled way by performing

ALD of indium oxysulfide. Oxygen is added in situ (during deposition) in this case. The oxygen addition helps in tuning the conduction and valence bands of In2S3 as well as tuning the carrier concentration and thus can potentially improve the open circuit voltage of the device. Further motivation for the indium oxysulfide material is discussed in chapter 3 and 4.

2.6 Double Emitter Experiments on CZT(S,Se) solar cells

CZT(S,Se) samples described in this section were deposited by Dr. Wei Wang at IBM

Thomas J Watson Research Center, Yorktown Heights. The solar cell fabrication and device

measurements described in this section were performed by Dr. Yun Seog Lee at IBM. The In2S3 depositions and analysis of results were done by the author of this work.

Open circuit voltage (Voc) deficit is defined as Eg/q -Voc (where Eg is the band gap and q is the electronic charge). While best CIG(S,Se) solar cells have Voc deficits in the range of 500 mV, current CZT(S,Se) devices with record efficiencies still suffer from a Voc deficit of 600 mV.45 Presence of antisite defect pairs involving Cu and Zn causes potential fluctuations and

45 thus limits the Voc in CZT(S,Se) devices.

Voc deficit reduction was aimed at by applying In2S3/CdS double emitters on to

CZT(S,Se) solar cells. The use of In2S3 as the second buffer with a subsequent anneal results in supply of indium into the CdS/CZT(S,Se) layer via annealing. Indium can n-type dope CdS or p- type dope CZT(S,Se) as it substitutes for Cd in CdS or for Sn in CZT(S,Se). The formation energies of In on Cd and In on Sn defects are likely low as In has a similar atomic radius compared to Cd and Sn and the ionic charges of the elements are only separated by 1.45 In terms of ionic radii, In3+ is 16 % larger than Sn4+ while it is 16 % smaller than Cd2+ .46

CZT(S,Se) films were deposited on Mo-coated soda lime glass by solution processing.43

CdS (50 nm thick) was deposited by chemical bath deposition, and In2S3 was deposited by atomic layer deposition. We used a closed-valve mode of deposition for In2S3 with a deposition temperature of 160 °C, the growth conditions are described in detail in section 2.3.2. Thicknesses of 3 and 5 nm of In2S3 were explored. To diffuse the indium into the CdS/CZT(S,Se) layer we rapidly annealed the double emitter structure at 250 °C on a hot plate with varying times up to 20 seconds. We then studied solar cells in the structure as shown below in Figure 2.6.1(a). The indium tin oxide (ITO) was deposited by sputtering and the Ni/Al contact was deposited by electron beam evaporation.

Figure 2.6.1. (a) Double emitter thin-film solar cell structure schematic with CZT(S,Se) as absorber.

450

400

350

300

In2S3 +CdS /CZT(S,Se) 250 250C anneals 200 BB vs accumulated time

150 Band Bending (meV) Bending Band 100

50 150C anneal 5s- 0eV bb 0 0 2 4 6 8 10 12 14 16 18 20 22 anneal time (s)

Figure 2.6.1. (b) Band bending vs anneal time of the double emitter structure at 250 °C.

It is seen from Figure. 2.6.1 (b) that as the anneal time of the CdS+In2S3 double emitter structure was increased, the band bending increased from 0 meV for no annealing to 430 meV for a 20 second anneal. Thus annealing aids the diffusion of indium into the CdS layer and changes the carrier concentration in the buffer. This increases the built in potential in the p-n

junction and thus increases the rectification of the junction. The corresponding band bending due to the electric field at the junction is also amplified.

Figure 2.6.2. Solar cell parameters obtained from the double emitter structure.

As one can see from Figure 2.6.2 (solar cell device parameters), there is a 30 to 40 mV improvement in Voc on adding the In2S3 layer to the CdS based solar cell. Thus the Voc deficit is reduced by 30 - 40 mV. There is not a significant difference in Voc and Jsc on changing the thickness of the In2S3 layer from 3 nm to 5 nm. The representative J-V curves of the devices are shown in Figure 2.6.3. The average fill factor of the 5 nm In2S3 based devices at around 50 % is

better than that of the 3 nm In2S3 based devices at around 40 %. This is mainly because the series resistance of 5 nm In2S3 based devices is lower than that of the 3 nm In2S3 based devices.

Figure 2.6.3. J-V curves of the double emitter devices with varying thicknesses of In2S3 in comparison to the control CdS based device.

Indium in–diffusion and doping results in an increase in the electron concentration of the

CdS layer and the hole concentration of the CZTS,Se layer. This should result in a drop in series resistance of the device. The drop in series resistance is higher for the 5 nm In2S3 layer based device compared to the 3 nm one due to more in-diffusion and doping. Increase in the electron concentration of the emitter and the hole concentration of the absorber in tandem results in an improvement of the built in potential of the p-n junction, which in turn increases the Voc. The other parameters though, like the fill factor, current, and the efficiency, on average are lower than the control single emitter CdS buffered solar cell. This is possibly due to the up diffusion of

elements like Cu, Zn into the double emitter structure increasing the resistance of the cell.

Further analysis by SIMS depth profiling is needed to understand this better.

2.7 Conclusions

Atomic layer deposition (ALD) of indium sulfide was studied for CZT(S,Se) solar cell applications using indium(III) acetylacetonate and indium formamidinate precursors. The former results in low growth rates, and carbon-containing films at temperatures of interest for depositing buffer layers. On the other hand, indium formamidinate affords smooth, carbon-free films for application as a buffer layer on CZT(S,Se). The In2S3 grown with the indium formamidinate precursor has an atypical carbon-free ALD window with decreasing growth rate with rise in temperature, a likely reason being sulfide bridge formation. The ALD conditions were optimized by adjusting the purging time and the exposure time of the 2 co-reactants and smooth films were obtained. XRD and TEM reveal polycrystalline nature of the In2S3 films in the ALD window of

16 3 160 °C to 200 °C. The carrier concentration in In2S3 is around 1.2 x 10 /cm at all temperatures.

o 2 The In2S3 films are fairly resistive at 160 C due to low mobility of 1.7 cm /V.sec. The mobility increases with temperature and reaches 17.7 cm2/V.sec at 200 oC.

Higher mobility films are preferable in solar cells but using a high temperature of growth increases chances of interface diffusion. Considering this we decided to investigate In2S3 deposited at 160 oC for its band alignment with CZT(S,Se) and performance in a CZT(S,Se) solar cell. UPS results indicate that the band offset of In2S3 with CZT(S,Se) is 0.29 eV and hence is a good candidate buffer layer in CZT(S,Se) solar cells. A representative CZT(S,Se) thin-film solar cell employing In2S3 was fabricated and device performance measured. While the short circuit current is higher than CdS based devices, all other properties fall short. The device open

circuit voltage obtained is promising, although around 80 mV lower than the CdS based devices.

This suggests the possible presence of defect states in In2S3 which act as traps. There is also the possibility of low number of charge carriers affecting the Voc. The presence of diffused oxygen in the In2S3 buffers might be a positive influence on the Voc. These results motivate the use of

In2(O,S)3 as a buffer layer grown by in-situ incorporation of oxygen into In2S3.

2.8 References

1. Barkhouse, D. A. R.; Haight, R.; Sakai, N.; Hiroi, H.; Sugimoto, H.; Mitzi, D. B., Cd-free buffer layer materials on Cu2ZnSn(SxSe1−x)4: Band alignments with ZnO, ZnS, and In2S3. Applied Physics Letters 2012, 100 (19), 193904.

2. Mughal, M. A.; Newell, M. J.; Vangilder, J.; Thapa, S.; Wood, K.; Engelken, R.; Carroll, B. R.; Johnson, J. B., Optimization of the electrodeposition parameters to improve the stoichiometry of In2S3 films for solar applications using the taguchi method. J. Nanomaterials 2014, 2014, 4-4.

3. Cansizoglu, M. F.; Engelken, R.; Seo, H.-W.; Karabacak, T., High Optical Absorption of Indium Sulfide Nanorod Arrays Formed by Glancing Angle Deposition. ACS Nano 2010, 4 (2), 733-740.

4. Ho, C.-H., Enhanced photoelectric-conversion yield in niobium-incorporated In2S3 with intermediate band. Journal of Materials Chemistry 2011, 21 (28), 10518-10524.

5. Haleem, A. M. A.; Sugiyama, M.; Ichimura, M., Sulphurization of the Electrochemically Dep- osited Indium Sulphide Oxide Thin Film and Its Photovoltaic Applications. Materials Sciences and Applications 2012, 03 (11), 802-806.

6. Newell, M. J.; Engelken, R.; Hall, J.; Mughal, M. A.; Felizco, F.; Vangilder, J.; Thapa, S.; McNew, D.; Hill, Z. In Elemental sulfur-based electrodeposition of indium sulfide films, 2011 37th IEEE Photovoltaic Specialists Conference, 19-24 June 2011; 2011; pp 001322-001326. 7. Gilles, J. M.; Hatwell, H.; Offergeld, G.; van Cakenberghe, J., Photoconductivity in Indium Sulfide. physica status solidi (b) 1962, 2 (4), K73-K77.

8. Springford, M., The Luminescence of Some Ternary Chalcogenides and Mixed Binary Systems of Group III-VI Compounds: The Nature of Luminescence Centres in Group III-VI Compounds. Proceedings of the Physical Society 1963, 82 (6), 1029.

9. Kim, C. D.; Lim, H.; Park, H. L.; Park, H. Y.; Kim, J. E.; Kim, H. G.; Kim, Y. G.; Kim, W. T., Optical absorption of Co2+ ions in In2S3 thin films. Thin Solid Films 1993, 224 (1), 69-73.

10. Kim, W. T. Kim, C. D., Optical energy gaps of β‐In2S3 thin films grown by spray pyrolysis. Journal of Applied Physics 1986, 60 (7), 2631-2633.

11. Allsop, N. A.; Schönmann, A.; Belaidi, A.; Muffler, H. J.; Mertesacker, B.; Bohne, W.; Strub, E.; Röhrich, J.; Lux-Steiner, M. C.; Fischer, C. H., Indium sulfide thin films deposited by the spray ion layer gas reaction technique. Thin Solid Films 2006, 513 (1), 52-56.

12. Barreau, N.; Marsillac, S.; Bernède, J. C.; Ben Nasrallah, T.; Belgacem, S., Optical Properties of Wide Band Gap Indium Sulphide Thin Films Obtained by Physical Vapor Deposition. physica status solidi (a) 2001, 184 (1), 179-186.

13. Lokhande, C. D.; Ennaoui, A.; Patil, P. S.; Giersig, M.; Diesner, K.; Muller, M.; Tributsch, H., Chemical bath deposition of indium sulphide thin films: preparation and characterization. Thin Solid Films 1999, 340 (1), 18-23.

14. Diehl, R. Nitsche, R., Vapour and flux growth of γ-In2S3, a new modification of insium sesquisulphide. Journal of Crystal Growth 1973, 20 (1), 38-46.

15. Okamoto, H., In-S (Indium-Sulfur). Journal of phase equilibria and diffusion 2013, 34 (2), 149-150.

16. McCarthy, R. F.; Weimer, M. S.; Emery, J. D.; Hock, A. S.; Martinson, A. B. F., Oxygen- Free Atomic Layer Deposition of Indium Sulfide. ACS Applied Materials & Interfaces 2014, 6 (15), 12137-12145.

17. Repins, I.; Contreras, M. A.; Egaas, B.; DeHart, C.; Scharf, J.; Perkins, C. L.; To, B.; Noufi, R., 19·9%-efficient ZnO/CdS/CuInGaSe2 solar cell with 81·2% fill factor. Progress in Photovoltaics: Research and Applications 2008, 16 (3), 235-239.

18. Mughal, M. A.; Engelken, R.; Sharma, R., Progress in indium (III) sulfide (In2S3) buffer layer deposition techniques for CIS, CIGS, and CdTe-based thin film solar cells. Solar Energy 2015, 120, 131-146.

19. Hariskos, D.; Spiering, S.; Powalla, M., Buffer layers in Cu(In,Ga)Se2 solar cells and modules. Thin Solid Films 2005, 480-481, 99-109.

20. Afzaal, M.; O'Brien, P., Recent developments in II-VI and III-VI semiconductors and their applications in solar cells. Journal of Materials Chemistry 2006, 16 (17), 1597-1602.

21. Yan, C.; Liu, F.; Sun, K.; Song, N.; Stride, J. A.; Zhou, F.; Hao, X.; Green, M., Boosting the efficiency of pure sulfide CZTS solar cells using the In/Cd-based hybrid buffers. Solar Energy Materials and Solar Cells 2016, 144, 700-706.

22. George, S. M., Atomic Layer Deposition: An Overview. Chemical Reviews 2010, 110 (1), 111-131.

23. Libera, J. A.; Hryn, J. N.; Elam, J. W., Indium Oxide Atomic Layer Deposition Facilitated by the Synergy between Oxygen and Water. Chemistry of Materials 2011, 23 (8), 2150-2158.

24. Sarkar, S. K.; Kim, J. Y.; Goldstein, D. N.; Neale, N. R.; Zhu, K.; Elliott, C. M.; Frank, A. J.; George, S. M., In2S3 Atomic Layer Deposition and Its Application as a Sensitizer on TiO2 Nanotube Arrays for Solar Energy Conversion. The Journal of Physical Chemistry C 2010, 114 (17), 8032-8039.

25. Naghavi, N.; Spiering, S.; Powalla, M.; Cavana, B.; Lincot, D., High‐efficiency copper indium gallium diselenide (CIGS) solar cells with indium sulfide buffer layers deposited by atomic layer chemical vapor deposition (ALCVD). Progress in Photovoltaics: Research and Applications 2003, 11 (7), 437-443.

26. Yousfi, E. B.; Weinberger, B.; Donsanti, F.; Cowache, P.; Lincot, D., Atomic layer deposition of zinc oxide and indium sulfide layers for Cu(In,Ga)Se2 thin-film solar cells. Thin Solid Films 2001, 387 (1), 29-32.

27. Jean-François, G.; Bruno, C.; El Bekkaye, Y.; Pierre, C.; Anouk, G.; Timo, A.; Michael, P.; Dimitri, H.; Hans-Werner, S.; Daniel, L., Indium-Based Interface Chemical Engineering by Electrochemistry and Atomic Layer Deposition for Copper Indium Diselenide Solar Cells. Japanese Journal of Applied Physics 2001, 40 (10R), 6065.

28. The International Center for Diffraction Data. Powder Diffraction File (PDF) 00-025-0390.

29. Asikainen, T.; Ritala, M.; Leskelä, M., Growth of In2S3 thin films by atomic layer epitaxy. Applied Surface Science 1994, 82-83 (Supplement C), 122-125.

30. Naghavi, N.; Henriquez, R.; Laptev, V.; Lincot, D., Growth studies and characterisation of In2S3 thin films deposited by atomic layer deposition (ALD). Applied Surface Science 2004, 222 (1), 65-73.

31. Spiering, S.; Eicke, A.; Hariskos, D.; Powalla, M.; Naghavi, N.; Lincot, D., Large-area Cd- free CIGS solar modules with In2S3 buffer layer deposited by ALCVD. Thin Solid Films 2004, 451-452 (Supplement C), 562-566.

32. Gunawan, O.; Virgus, Y.; Tai, K. F., A parallel dipole line system. Applied Physics Letters 2015, 106 (6), 062407.

33. Haight, R.; Barkhouse, A.; Gunawan, O.; Shin, B.; Copel, M.; Hopstaken, M.; Mitzi, D. B., Band alignment at the Cu2ZnSn(SxSe1−x)4/CdS interface. Applied Physics Letters 2011, 98 (25), 253502.

34. Lim, D.; Haight, R., In situ photovoltage measurements using femtosecond pump-probe photoelectron spectroscopy and its application to metal–HfO2–Si structures. Journal of Vacuum Science & Technology A: Vacuum, Surfaces, and Films 2005, 23 (6), 1698-1705.

35. Lim, D.; Haight, R., Temperature dependent defect formation and charging in hafnium oxides and silicates. Journal of Vacuum Science & Technology B: Microelectronics and Nanometer Structures Processing, Measurement, and Phenomena 2005, 23 (1), 201-205.

36. Du, H.; Romero, M. J.; Repins, I.; Teeter, G.; Noufi, R.; Al-Jassim, M. M. In Nanoscale measurements of the surface photovoltage in Cu(In,Ga)Se2, Cu2ZnSnS4, and Cu2ZnSnSe4 thin films: The role of the surface electronics on the efficiency of solar cells, 2011 37th IEEE Photovoltaic Specialists Conference, 19-24 June 2011; 2011; pp 001983-001986.

37. Song, T.; Kanevce, A.; Sites, J. R., Emitter/absorber interface of CdTe solar cells. Journal of Applied Physics 2016, 119 (23), 233104.

38. Todorov, T. K.; Reuter, K. B.; Mitzi, D. B., High-Efficiency Solar Cell with Earth-Abundant Liquid-Processed Absorber. Advanced Materials 2010, 22 (20), E156-E159.

39. Barkhouse, D. A. R.; Gunawan, O.; Gokmen, T.; Todorov, T. K.; Mitzi, D. B., Device characteristics of a 10.1% hydrazine-processed Cu2ZnSn(Se,S)4 solar cell. Progress in Photovoltaics: Research and Applications 2012, 20 (1), 6-11.

40. Bag, S.; Gunawan, O.; Gokmen, T.; Zhu, Y.; Todorov, T. K.; Mitzi, D. B., Low band gap liquid-processed CZTSe solar cell with 10.1% efficiency. Energy & Environmental Science 2012, 5 (5), 7060-7065.

41. Todorov, T. K.; Tang, J.; Bag, S.; Gunawan, O.; Gokmen, T.; Zhu, Y.; Mitzi, D. B., Beyond 11% Efficiency: Characteristics of State-of-the-Art Cu2ZnSn(S,Se)4 Solar Cells. Advanced Energy Materials 2013, 3 (1), 34-38.

42. Winkler, M. T.; Wang, W.; Gunawan, O.; Hovel, H. J.; Todorov, T. K.; Mitzi, D. B., Optical designs that improve the efficiency of Cu2ZnSn(S,Se)4 solar cells. Energy & Environmental Science 2014, 7 (3), 1029-1036.

43. Wang, W.; Winkler, M. T.; Gunawan, O.; Gokmen, T.; Todorov, T. K.; Zhu, Y.; Mitzi, D. B., Device Characteristics of CZTSSe Thin-Film Solar Cells with 12.6% Efficiency. Advanced Energy Materials 2014, 4 (7), 1301465.

44. Cohen, S. S., Contact resistance and methods for its determination. Thin Solid Films 1983, 104 (3), 361-379.

45. Kim, J.; Hiroi, H.; Todorov, T. K.; Gunawan, O.; Kuwahara, M.; Gokmen, T.; Nair, D.; Hopstaken, M.; Shin, B.; Lee, Y. S., High efficiency Cu2ZnSn (S,Se) 4 solar cells by applying a double In2S3/CdS emitter. Advanced Materials 2014, 26 (44), 7427-7431.

46. Shannon, R., Revised effective ionic radii and systematic studies of interatomic distances in halides and chalcogenides. Acta Crystallographica Section A 1976, 32 (5), 751-767.

Chapter 3

Atomic Layer Deposition of Binary Indium Oxide

(In2O3) Thin Films Using Amidinate Precursors, and Studies of the Structural, Electrical, and

Optical Properties

Part of this chapter has appeared in the following paper:

Obtaining a Low and Wide Atomic Layer Deposition Window (150-275 °C) III for In2O3 Films using an In amidinate and H2O Sang Bok Kim, Ashwin Jayaraman, Danny Chua, Luke M. Davis, Shao-Liang Zheng, Xizhu Zhao, Sunghwan Lee, and Roy G. Gordon (Submitted)

3.1 Chapter Abstract

The ALD process of indium oxide was studied with the aim of coupling it with ALD of

In2S3 to obtain a ternary In2(O,S)3 . The end goal was to obtain controlled incorporation of oxygen in the ternary material. A newly synthesized amidinate tris(N,N'- diisopropylformamidinato)indium(III), compound 1, was studied, which gives a carbon-free

ALD window of reaction with H2O to form In2O3 in the temperature range 150 to 275 °C. In

comparison, another amidinate tris(N,N'-diisopropylacetamidinato)indium(III), compound 2, which can be thought of as having the H of the anionic iPrNC(H)NiPr ligands replaced with a methyl group, gives a narrower ALD window at higher temperatures 225 to 300 °C for the same reaction to form In2O3. Kinetic studies demonstrate that the faster rate of surface reactions in both parts of the ALD cycle is a potential reason for the difference in ALD windows. Electrical and optical properties were analyzed for the In2O3 obtained by ALD using both the amidinate compounds 1 and 2. Based on the ALD window, the kinetic studies of In2O3 growth, and electro- optical properties, precursor 1 was selected for growing the ternary In2(O,S)3.

3.2 Introduction

In the previous chapter we concluded that indium sulfide deposited by ALD is not the most optimal electron transport layer for the solar absorber, Cu2ZnSn(SxSe1-x)4 (CZTS,Se).

Power conversion efficiencies as high as 7.6 % have been obtained by chemical bath deposited

1 In2S3 coupled with CZT(S,Se). The band offset value between In2S3 and CZT(S,Se) in this case

2 is close to optimal (spike of 0.15). In another study of chemically bath deposited In2S3 in combination with Cu2ZnSnS4, a band offset of 0.41 spike was measured. Working devices with a relatively high Voc of 590 mV (Voc deficit of 910 mV) were obtained. These results made chemically bath deposited In2S3 one of the more sought after alternatives to CdS in CZT(S,Se) devices. The chemical bath deposited In2S3 is prepared from a solution of indium(III) chloride and thioacetamide3 and incorporates oxygen and hydroxyl groups due to the solution chemistry of the process. The oxygen content in CBD In2S3 (5-10 %) is uncontrolled though and could lead to variable electronic properties. Thus we were directed into attempting controlled incorporation of oxygen into In2S3 by ALD. Such an ALD process will involve sub-cycles where oxygen is

inducted into the material to form In-O bonds. Thus an ALD reaction of indium oxide needs to be independently studied before coupling it with the ALD indium sulfide reaction.

There are other reasons to pursue ALD of indium oxide apart from alloying the indium sulfide to obtain a well-tailored electron transport layer. Until now the most common use of indium oxide is in its tin-doped form (ITO) which is used as a transparent conducting oxide4 in flat panel displays and solar panels.5 Properties like band alignment impel the use of TCOs based on tin oxide or zinc oxide in some applications, but tin-doped In2O3 (ITO) remains the industry standard for flat panel displays. The ITO films are usually deposited by magnetron sputtering6-7 and chemical vapor deposition.8 Although these methods can give high growth rates on flat substrates, high temperatures, thickness variation, and sputtering damage are issues when applied to thin films and non-flat substrates.7-9 There is thus an increasing demand for multi-component conformal thin films (10-30 nm) containing indium and atomic layer deposition of indium oxide is a possible solution.10-11 Thus studying indium oxide deposition by ALD independently can help us come up with an optimal TCO for CZT(S,Se) devices.

To obtain a ternary compound with indium, i.e., to alloy, we should find a precursor with a growth window of the oxide beginning at a low temperature and having a wide temperature range, in which the growth rate is constant, and deposition proceeds with canonical alternating cycles of surface saturation.10-11An ALD window that begins at a very low temperature helps with sensitive substrates like polymers and organic light absorbers. In the case of multilayered devices a low ALD growth temperature means less interdiffusion between layers. Thus the elements in the ternary indium oxysulfide buffer that we will eventually grow will not diffuse into the CZT(S,Se) layer, nor will the elements in our solar absorbers diffuse out into the electron transport layer. Usually growth temperatures higher than 200 oC are not used to prevent

significant interdiffusion. A wide ALD window of growth would allow for doping/alloying In3+ into a wide range of metal oxides or other indium chalcogenides.10-12

9, 12-13 Finding a low temperature process for ALD of In2O3 has been a challenge. The only reports of ALD of In2O3 below 200 °C require an oxidant like O2 or O3. The use of an oxygen source other than water limits the compatibility of the process for alloys. Thus we should target an ALD reaction with water as a co-reactant.

It is surprising than we do not have a low temperature ALD for oxide of indium because the oxide of aluminum, which belongs to the same group as indium, is grown by a very clean ALD process in the temperature range 110-300 °C with alternating doses of trimethylaluminum and water14. Trimethylindium on the contrary has an ALD reaction with water only above 200 °C.15

Mixed alkyl-amido compounds like Et2In[N(SiMe3)2] give enhanced growth rates but there is carbon in the films below 200 °C.13 Carbon in these films is likely due to the unreacted chemisorbed Et2In[N(SiMe3)2] rather than condensation because the precursor vaporization

13 temperature used is 40 °C. The high growth rates of In2O3 ALD with Et2In[N(SiMe3)2] (~0.7

Å/cycle) compared to that with Me3In (0.39 Å/cycle) may indicate a higher reactivity of water with indium-nitrogen bonds than with indium-carbon bonds. The low growth rate of the Me3In- based ALD also means that reaction kinetics of chemisorbed Me3In is lower with water. Thus we have studied the reaction of water vapor with a newly synthesized indium amidinate precursor.

The ligand was selected in light of the excellent volatility and reactivity of the corresponding

Ca2+ amidinate.16 This study can help to define a rapid, low temperature ALD process for growing In2O3 and alloying oxygen into In2S3.

This chapter highlights the comparison of existing precursors for an indium oxide ALD process with a newly synthesized amidinate precursor, 1. The synthesis of precursors 1 and 2 and

corresponding NMR experiments were done by Dr. Sang Bok Kim. The film depositions, kinetic studies and electric and structural characterizations of the In2O3 films were done by the author of this work. Dr. Sang Bok Kim and Dr. Luke Davis helped with the analysis of results. The author thanks Xizhu Zhao, Prof. Sunghwan Lee and Danny Chua for help with UV-Vis spectrophotometry and Hall measurements respectively.

In this chapter, the basic properties of an ALD reaction such as ALD window and growth rate as a function of precursor doses are compared for 1 and its acetamidinate congener 2. The kinetics of the ALD reaction of 1 and 2 with water are compared at 250 °C (in the ALD window of both precursors). X-ray photoelectron spectroscopy (XPS) was performed to determine the carbon content in the films as a function of temperature. Band gap measurements were obtained to compare the indium oxide grown by the two precursors. Hall effect measurements were performed to identify the temperature that affords the highest electron mobility. This property is important because we aim to obtain a high-mobility indium oxysulfide film. Depositing the oxysulfide material at a temperature at which the ALD oxide and sulfide individually have high mobility would be ideal. This would potentially result in a high mobility oxysulfide film. X-ray diffraction (XRD) to analyze crystallinity and scanning electron microscopy (SEM) to view morphology were also performed. One of the amidinate precursors was eventually selected for growing oxysulfide films with controlled oxygen to sulfur ratios.

3.3 Experiments

3.3.1 Atomic layer deposition of indium oxide films

A custom-built hot wall ALD reactor17 was used for performing the reaction of indium formamidinate or acetamidinate with H2O. Please refer Appendix A for the reactor diagram.

In2O3 was deposited on thermal oxide (300 nm of SiO2 on Si) treated with isopropyl alcohol to remove organic contamination and further exposed to UV-ozone for obtaining terminal hydroxyl groups, to initiate the ALD half reactions. Precursors 1 and 2 were maintained in glass bubblers and the source temperature used was 125 °C for 1 and 140 °C for 2. H2O vapor pressure above the surface of water maintained in a glass bubbler at room temperature was used as the source of oxygen for the reaction. Distilled water was used.

5 seconds of evacuation time was provided before performing the ALD half reactions in closed valve mode. The indium amidinate precursors assisted by N2 carrier gas (10 torr) and the

H2O were injected sequentially into the reactor, which allowed for chemical reactions to occur successively on the substrate surface. 1 dose of 1 (2 doses of 2 separated by 5 s of evacuation) were used in combination with 1 dose of H2O in each cycle. The exposures of 1 and 2 were 1.4

Torr·s and 1.1 Torr·s respectively. The H2O exposure was 2.4 Torr·s. N2 purging with 600 mtorr pressure was performed after each dose of precursor or H2O to remove byproduct species and unreacted material, which in turn prevented reactions from happening in the gas phase. Different purging times were used for each precursor (10 s for 1, 40 seconds for 2, and 20 seconds for

H2O. For the ALD using 1, 1300 cycles were used and temperatures between 150 °C and 325 °C were analyzed at approximately 25 °C intervals. For ALD using 2, 2000 cycles were used between 150 °C and 350 °C. The deposition study was conducted using the closed-valve mode of

ALD reactions. In this mode the reaction chamber was evacuated and closed at one end before reactants entered the other end for reaction.

3.3.2 Characterization of In2O3 films

XRD (D8 DISCOVER with DAVINCI Design, equipped with Gobel mirror/ACC2 with

LynxEye geometry) was used to obtain the X-ray diffractogram of the In2O3 film. The thin film for XRD was 100 nm in thickness deposited on a quartz substrate. Images of top view and side view of the In2O3 film were obtained using Field Emission Scanning Electron Microscopy (FE-

SEM, ZEISS, Ultra-55). The SEM samples were thin films deposited on thermal oxide (300 nm

SiO2 on Si). X-ray photoelectron spectroscopy (XPS) was done with a ThermoScientific K-

Alpha spectrometer equipped with a monochromatized Al Kα X-ray source (1486.6 eV), 12 kV electron beam, and argon sputtering source. The high resolution scans were analyzed using a high current, 500 eV argon ion sputtering energy with 60 seconds of sputtering per level unless otherwise stated. Resistivity, carrier concentration, and Hall mobility are obtained using a custom-built Hall apparatus with a Keithley sourcemeter.18 Samples for Hall measurement were

1 cm by 1 cm with a thickness of around 100 nm deposited on thermal oxide (300 nm SiO2 on

Si). Indium metal was used as contact. A UV/visible spectrophotometer with an integrating sphere was used to measure the light transmittance and reflectance and determine the absorption coefficient (α).

3.3.3 Kinetic studies of indium oxide film growth

Kinetic studies (H2O reaction studies) were performed by varying the reaction time (sum of the dosing time and exposure time of a precursor in the reactor) of H2O with a surface saturated with chemisorbed 1 or chemisorbed 2 at 250 °C. We chose this temperature because it belongs to the ALD window of both 1 and 2 as seen in the Results section of this chapter. The reaction studied is shown in Figure 3.3.3.1

Figure 3.3.3.1. The expected surface reactions of water with a surface saturated with 1 or 2 at 250 °C.

To saturate the substrate surface with chemisorbed 1, we need a surface reaction of 1 with a surface saturated with H2O. Thus to saturate the surface, we use 1 dose of H2O (1 second) and 5 seconds of holding it in the reactor before purging. This reaction time of water is sufficient for standard ALD growth as seen from H2O reaction studies described later in this chapter. The reaction time of 1 is varied in the study to obtain conditions under which a surface saturated with

1 is seen (i.e., ALD growth rate is measured).

To obtain a surface saturated with chemisorbed 2, we used multiple doses of 2. Each dose was separated by 5 seconds of evacuation and consisted of 6 seconds of reaction time (1 second dose + 5 seconds holding in the reactor before purging). The reason why we checked for multiple doses for 2 was because 1 dose of 2 with 12 seconds of reaction time at 250 °C could not result in surface saturation. The amount of H2O used in this case for saturation was 1 second of dosing

followed by 11 seconds of exposure before purging. This amount was more than sufficient for

ALD growth as seen from the H2O reaction studies.

After this, the reaction of H2O with the surface saturated with 1 or 2 was conducted. In this study deposition conditions for saturation of chemisorbed 1 or 2 were used and reaction times of H2O were varied. The growth rate of the film was checked to obtain conditions under which saturated ALD growth rate was achieved.

3.4 Results and Discussion

3.4.1 ALD window determination

The precursor 1, also referred to as indium formamidinate, was vaporized at 125 °C while precursor 2, also referred to as indium acetamidinate, was vaporized at 140 °C. The formulas of precursors 1 and 2 are shown in Figure 3.4.1.1.

(a) (b)

Figure 3.4.1.1. Formulas of precursor 1 (a) and precursor 2 (b).

Precursor 2 should be heated to around 165 °C to equalize the vaporization rate with 1 as obtained from thermogravimetric analysis experiments done by Dr. Sang Bok Kim. But our ALD valves constrain the temperatures to 140 °C and hence we use 2 doses at that temperature to saturate the substrate surface.

The ALD sequence used is t1-t2-t3-t4 where t1 is the evacuation time to bring the reactor down to base vacuum pressure, t2 is the dosing time of the precursor or water, t3 is the holding time of the vapor in the reactor, t4 is the purging time with nitrogen flowing, with all times given in seconds. For 1, t1-t2-t3-t4 = 5-1-5-10, followed by t1-t2-t3-t4 = 5-1-5-20 for the water dose.

In the case of 2, it is 2 doses (the doses are separated by 5 seconds of vacuum) of t1-t2-t3-t4 = 5-

1-5-40 followed by t1-t2-t3-t4 = 5-1-11-20 for the water dose. Samples were deposited at different temperatures as mentioned in the Experiments section and the thicknesses of samples for growth rate determination were measured by SEM. Growth rate is plotted as a function of temperature for 1 and 2 in Figure 3.4.1.2 and Figure 3.4.1.3.

1.0 3 % Carbon

0.8 3 % Carbon ALD window using 1

0.6

1 % Carbon 0.4

0.2 Longer Purging time of 1 min for both and H2O

Growth Rate (Angstrom/cycle) Rate Growth 1 - 2 % Carbon in film

0.0 100 150 200 250 300 350 Temperature (oC)

Figure 3.4.1.2. Growth rate as a function of temperature showing ALD window for 1.

1.0 5 % Carbon 0.9

0.8

0.7 ALD window using 2

0.6

0.5

0.4 1% Carbon

0.3

0.2

Growth Rate (Angstrom/cycle) Rate Growth 0.1

0.0 150 200 250 300 350 o Temperature ( C)

Figure 3.4.1.3. Growth rate as a function of temperature showing ALD window for 2.

The ALD window for 1 was found to be from 150 °C to 275 °C with a growth rate of

0.55 Å/cycle while the ALD window for 2 was found to be from 225 °C to 300 °C with a similar growth rate of 0.52 Å/cycle. It is worth noting that the cycle time for 1 is 52 seconds while that for 2 is 93 seconds. The carbon content was measured by XPS high resolution scans and carbon to the extent of 1 - 5 at. % was seen outside the ALD window in both cases. The carbon % is computed using curve fitting by “ThermoFisher Scientific Avantage 5.957 surface chemical analysis” deconvolution software. The films deposited in the ALD window are relatively carbon free (< 1 at. % by XPS). Figure 3.4.1.4 and Figure 3.4.1.5 show the high resolution scan of carbon at different temperatures for the ALD In2O3 using 1 and 2.

24000

22000 1 3 % C 325 oC

o 20000 1 % C 300 C

275 oC 18000 250 oC

16000 225 oC

200 oC 14000 o Intensity (arb. units) (arb. Intensity 175 C

o 12000 155 C

2 % C 130 oC 10000

280 282 284 286 288 290 292 294 296 298 Binding Energy (eV) .

Figure 3.4.1.4. High resolution carbon 1s scans for In2O3 films grown using 1 with carbon at. % computed by fitting shown.

26000 2 24000 o 5 % C 350 C 22000 1 % C 330 oC 20000 300 oC

o 18000 275 C

250 oC 16000

225 oC Intensity (arb. units) (arb. Intensity 14000 200 oC

175 oC 12000 150 oC 10000 280 282 284 286 288 290 292 294 296 298 Binding Energy (eV)

Figure 3.4.1.5. High resolution carbon 1s scans for In2O3 films grown using 2 with carbon at. % computed by fitting shown.

For 1 at 130 °C we obtain a lower growth rate of 0.35 Å/cycle with a carbon content of 2 at. %. Different from the recipe at other temperatures, we used a very long purging time of 1 min 76

for the precursor and water in this case. This essentially means that the temperature of 130 °C is not sufficient for the reactions to completely saturate the surface as the growth rate is below that of the ALD window. At 140 °C with 1 and water, under normal purge conditions a high growth rate of 1 Å/cycle was obtained with high carbon content. The reason for this is that the ALD half reactions do not go to completion and the protonated ligands are not completely detached during

ALD at low temperatures. At temperatures above 275 °C, carbon is seen again in the films grown with 1 and the growth rate increases above that of the ALD window. This carbon likely results from decomposition of surface-bound ligands at such high temperatures. The precursors 1 and 2 are found not to decompose within the ALD window, at least on the TGA timescale.

For ALD reaction involving 2 and water, the temperature range between 150 °C and 225

°C produces carbon-free films but the growth rate is below that of the ALD window. This points to the fact that the surface saturation is not observed even though 12 seconds were given for reaction completion.

Figure 3.4.1.6. Growth rate as a function of temperature for all reported processes for ALD of In2O3. Lines represent the carbon-free ALD window and circles denote temperatures at which the product film resistance is less than 0.01 Ω·cm. Red markers indicate water as the oxygen source while gray markers indicate other oxygen sources such as ozone, hydrogen peroxide or oxygen. 19 15 13 The indium sources are compounds 1 and 2 (our study) and TMIn (A ), Et2In(Ntms2) (B ), 20 21 22-23 9 24 25 DADI (C ), In[C(NiPr2)(NiPr)2]3 (D ), InCl3 (E ), InCp ((F ), (G )), TEIn (I ), 26 27 Me2In(EDPA) (J ), In(tmhd)3 (K ).

Figure 3.4.1.6 shows growth rate as a function of temperature for all ALD In2O3 processes in the literature. It is seen here that there are no other precursors apart from 1 that deposit ALD In2O3 with just water as the oxidant at temperatures below 200 °C. The amidinate precursors 1 and 2 result in In2O3 with resistivity values below 0.01 Ω·cm over a wide range of temperatures. The other precursors are able to produce conductive films only in one portion of their ALD window, typically either at high temperature or at low temperature. Compound 1 is neither oxygen sensitive like InCp nor pyrophoric like the indium(III) alkyls. Among all In2O3

ALD reactions with water, 1 affords the lowest ALD onset temperature and the widest ALD window. Many ALD windows of metal oxides overlap with the ALD window of reaction between compound 1 and water.11

We have classified the different indium precursors in terms of their ALD onset temperature and ALD window width (See Table 3.4.1.). Note that the ALD cycle time is

different for different precursors. It is seen that while 1 belongs to class 1 of precursors, 2 belongs to class 4. It is seen that very few precursors have reactions with just water as the oxygen source.

Table 3.4.1. Classification of indium precursors for In2O3 ALD on the basis of ALD onset temperature and ALD window temperature range.

Classification ALD of In2O3

9 Class 1 (window temperature In(I)Cp with simultaneous H20 and O2 range ≥ 100 °C, min temp < 200 140-250 °C, 1.6 Å/cycle, ~ 1.1 × 10-3 Ω·cm

25 °C) TEIn(III) with O3

100-200 °C, 0.6 Å/cycle; 2×10-3 -8.58 × 10-5 Ω·cm

27 In(tmhd)3 with O2 plasma

100-400 °C, 0.14 Å/cycle; 1.8 × 10-2 at 100 °C, 1.7 × 10-2

at 150 °C, 2.5 × 10-3 at 300 °C, 2.6 × 10-3 Ω·cm at 400 °C

Class 2 (window temperature In(I)Cp with ozone24 range ≥ 100 °C, min temp ≥ 200 200-450 °C, 1.3 A/cycle; 1.6 × 10 -2 Ω·cm at 275 °C.

°C Resistivity at other temperatures not reported in paper

22 InCl3 with aqueous H2O2

Possibly 300-500 °C, 0.4 Å/cycle; 1.0 × 10-2 Ω·cm at 300

°C, 1.5 × 10-3 Ω·cm at 500 °C

Table 3.4.1. (Continued) Classification of indium precursors for In2O3 ALD on the basis of ALD onset temperature and ALD window temperature range.

28 Class 3 (window In(acac)3 with H20 temperature < 100 °C, min 165-200 °C, 0.2 Å/cycle 60 Ω·cm at 175 °C, 7 ×10-2 at 200 °C

28 temp < 200 °C) In(acac)3 with O3

165-225 °C, 0.12 Å/cycle > 1000 Ω·cm

13 Et2InN(TMS)2 with H20

175-250 °C, 0.7 Å/cycle; 10-1 Ω·cm at 175 °C to 10-5 at 250 °C

in the range 10-3 to 10-5 only between 210 and 250 °C

26 Me2In(EDPA) with O2 plasma

90-180 °C, 0.53 Å/cycle; 1.5×10-1 Ω·cm

15 Class 4 (window TMIn with H2O temperature < 100 °C, min 200-250 °C, ~ 0.4 Å/cycle; 2.8 ×10-3– 5 ×10-3 Ω·cm at 200-250 temp ≥ 200 °C) °C

21 Indium Guanidinate with H2O

230-300 °C, 0.4-0.45 Å/cycle; resistivity data not reported

22, 29 Etc. (No reported window) InCl3 with water

0.23 Å/cycle (400 °C), 0.27 Å/cycle (500 °C); 3 ×10-3 Ω·cm

20 [3-(dimethylamino)propyldimethylIn(3) with H2O

ALD only at 275 °C, 0.6 Å/cycle; 9.2 ×10-5 Ω·cm at 275 °C

(deposition over 275 °C was not reported)

3.4.2 Kinetic Studies by surface saturation experiments

Kinetic studies at 250 °C as described in the Experimental section were performed to understand why the lowest ALD temperature for 1 with water is lower compared to that of 2 with water. 250 °C falls within the ALD window of both precursors. To compare reactivity of the two precursors with water, we saturated the surface of the substrate with the precursor before going on to study the exposure time necessary for water to completely saturate the substrate surface and achieve the ALD growth rate. As the temperature of 250 °C is within both ALD windows we are confident about the surface saturation with the indium source.

Surface saturation with precursor 1 or 2 was obtained by experiments as discussed in section

3.3.3. As shown in Figure 3.4.2.1 (a), a single dose of precursor 1 was used in combination with sufficient water (1 second dose, 5 seconds exposure) at around 150 °C and growth rate variation was studied. It is seen in Figure 3.4.2.1 (a) that 4 seconds of reaction time (1 second dosage + 3 seconds exposure) of precursor 1 saturates the surface and close to the ALD growth rate (0.55

Å/cycle) is achieved. It is thus expected that it takes less time than 4 seconds for 1 to saturate the surface at 250 °C. ALD growth was indeed obtained on dosing precursor 1 for 1 second and holding it for 1 second in the reactor at higher temperatures. Figure 3.4.2.1 (b) shows the hypothetical mechanism of attachment of 1 on to a hydroxylated surface.

0.60

0.55

0.50

0.45

0.40

0.35

0.30

Growth(Angstrom/cycle) Rate 0.25 1 dose 150 oC 1 0.20 2 3 4 5 6 Reaction time (sec)

(a) (b)

Figure 3.4.2.1. (a) Growth rate as a function of reaction time for precursor 1 , (b) hypothetical mechanism of attachment of 1 on to a hydroxylated surface.

In the case of precursor 2, because the vaporization temperature was not sufficient (as discussed above), multiple doses are used for saturation. As seen in Figure 3.4.2.2 (a), 1 dose (6 seconds) is not enough for saturation as it gives a growth rate of only 0.2 Å/cycle. The maximum growth rate is achieved by the 2nd dose after 12 seconds and the 3rd dose (exposure time beyond 12 seconds) does not offer any more growth. Each dose thus as discussed earlier has 6 seconds of reaction time and doses are separated by 5 seconds of vacuum. Figure 3.4.2.2 (b) shows the hypothetical mechanism of attachment of 2 onto a hydroxylated surface.

0.55

0.50

0.45

0.40

0.35

0.30 o 0.25 250 C

0.20

GrowthRate (Angstrom/cycle) 0.15 6 8 10 12 14 16 18 Reaction time

(a) (b)

Figure 3.4.2.2. (a) Growth rate as a function of doses for precursor 2, (b) hypothetical mechanism of attachment of 2 onto a hydroxylated surface.

Since we use the reaction rate between the above saturated surfaces and H2O vapor, we avoid confounding the reaction rate with the difference in vapor pressure between 1 and 2 at 250

°C. The similarity of ALD growth rate between the two precursors 1 (0.55 Å/cycle) and 2 (0.52

Å/cycle) rules out confounds resulting from different extents of surface coverage. In order to obtain growth rates of 0.55 and 0.52 Å/cycle the coverage in case of 1-saturated and 2-saturated surfaces should be comparable. The saturated surfaces are treated with water vapor and the results of the kinetics experiments are shown in Figure 3.4.2.3 and Figure 3.4.2.4.

0.56

0.54

0.52

/cycle) o 0.50

0.48

0.46 1 Indium Formamidinate GrowthRate (A 0.44

0.42 1 2 3 4 5 6 Water Reaction time with Surface (s)

Figure 3.4.2.3. Growth rate as a function of water reaction time for 1 at 250 °C.

0.54

0.52

0.50

0.48 /cycle)

o 0.46

0.44

0.42 0.40 2 0.38 Indium

GrowthRate (A 0.36 Acetamidinate 0.34

0 2 4 6 8 10 12 Water Reaction time with Surface (s)

Figure 3.4.2.4. Growth rate as a function of water reaction time for 2 at 250 °C.

Reaction time is defined again as the sum of the dosing time and holding time inside the reactor. It is seen that for precursor 1 the reaction time for water to achieve the ALD growth rate of 0.55 Å/cycle is just 2 seconds (1 second of dosing and 1 second of holding in the reactor). In contrast, the reaction of a surface saturated with precursor 2 with water is very slow and growth rate increase is very gradual with reaction time and reaches saturation to ALD growth rate after 7 seconds. This observation that water reacts readily with a surface saturated with 1 as opposed to a surface saturated with 2 at the same temperature helps us rationalize the earlier ALD growth

onset in the case of 1. The surface reactions seemingly go easily to completion in the case of 1 and hence lower growth temperatures are possible. The plausible reaction of precursor 1 or 2 with water is shown in Figure 3.4.2.5(a)

Figure 3.4.2.5. (a) Plausible reaction of water with surface saturated with 1 or 2.

Figure 3.4.2.5. (b) Hydrophobic nature of 2 v/s hydrophilic nature of 1.

In Figure 3.4.2.5. (b) it is seen that 2 with its backbone methyl group is sterically hindred compared to the more hydrophilic nature of 1 which has a backbone hydrogen attached.

Figure 3.4.2.5. (c) Steric hindrance comparison between 1 (R=H) and 2 (R=CH3).

The saturated surface with chemisorbed 2 may exhibit more steric hindrance effect (see Figure

3.4.2.5(c)) on the formation of its activated complex. During the ALD reaction incoming water needs to get to the In-N bonds of the precursor to break them and form In-OH and H-N bonds.

Also incoming indium precursor needs to reach the OH groups on the surface (again to make In-

O bonds and H-N bonds). In both the aforementioned cases, the acetamidinate ligand takes up more space compared to the formamidinate and is less flexible, making it harder for the water/hydroxide to reach the In-N bonds.

A possible mechanism determining the ALD window is the hydrophobicity or hydrophilicity of the precursor saturated surface. 1 is likely more hydrophilic by nature compared to 2. The rate of ALD reaction of 1 with a hydroxyl terminated surface depends on temperature. While compound of 1 requires 4 seconds of reaction time for saturation at 150 °C, 2 seconds suffice for a complete reaction of 1 with a hydroxyl-terminated surface at 250 °C.

No matter the detailed mechanism, we find that 1 has higher reactivity with water as compared to 2. This allows ALD at lower temperatures and implies we can deposit films more

quickly; although 1 and 2 have similar growth per cycle, the cycles employing 1 can be significantly shorter. Another advantage we find for 1 over 2 is that the ALD window of 1 with

H2O begins at a lower temperature. We know that the ALD window of indium sulfide using 1 is from 160-200 °C (see chapter 2). In order to dope or alloy the indium sulfide with indium oxide in ALD mode, the temperature chosen for deposition must ideally fall in ALD window of both the sulfide and the oxide. The In2O3 ALD window of 2 (225-300 °C) barely overlaps with that of

In2S3 from 1, while the ALD window of 1 (150-275 °C) shares a wide range with that of In2S3 from 1. Together, these features mean we can use 1 with water for doping the indium sulfide in a ternary ALD process, and can deposit this electron transport layer at a fast rate. The higher rate and lower temperature allowed by the enhanced reactivity of 1 should result in less interdiffusion of elements across the junction during film growth.

3.4.3 Scanning Electron Microscopy Studies

a) b)

cycles

3300 3300

1300 cycles 1300 ALD WINDOW ALD

Figure 3.4.3.1. (a) Top view of In2O3 films grown by ALD using tris(N,N'- diisopropylformamidinato)indium(III) and (b) Cross-sectional view of In2O3 films grown by ALD using tris(N,N'-diisopropylformamidinato)indium(III). Scale bar indicates 100 nm.

Figure 3.4.3.1 (a) and (b) show top view and cross-sections respectively of the In2O3 films grown with 1 and H2O under ALD conditions. The growth consisted of 1300 ALD cycles.

It is seen from the top view that the morphology of In2O3 consists of tetrahedral crystallites at most temperatures. The feature size increases with temperature in the ALD window and above the ALD window starts decreasing. This is consistent with the fact that there is carbon in the

films above the ALD window and this carbon is likely interfering with the crystal growth. The crystals seem to be in the growth phase at temperatures 130-175 °C. From 200 °C onward clear crystalline facets are seen. The cross-sectional view shows that all films in the ALD window have grown to similar thickness ~ 72 nm. The film at 130 °C is thicker because it is grown with

3300 cycles to account for its lower growth rate.

Figure 3.4.3.2 shows the top view (a) and cross-sections (b) of ALD In2O3 films grown by 2. The growth consisted of 2000 ALD cycles. It is seen from the morphology that the crystal facets of In2O3 do not clearly emerge until the ALD window onset at 225 °C. At lower temperatures (150-200 °C) the feature size is small. The feature size increases with temperature in the ALD window range between 225 and 300 °C. Beyond 300 °C, the size of the facets seen, starts decreasing, most likely due to carbon in the films interfering with crystal growth. The morphology itself appears to change at 330 °C, outside the ALD window range. The cross- sectional view shows that the growth rate is lower in the temperature range 150-200 °C and thickness attained for 2000 cycles is around 60 nm. In the ALD window range the thicknesses attained are around 105 nm at a growth rate of 0.55 Å/cycle. At 350 °C, much beyond the ALD window the thickness increases to 169 nm. This is likely due to precursor decomposition on the surface of the substrate at high temperatures. The growth mechanism thus changes at higher temperatures outside the ALD window.

a) b)

2000 cycles 2000 ALD WINDOW ALD

Figure 3.4.3.2 (a) Top view of In2O3 films grown by ALD using tris(N,N'- diisopropylacetamidinato)indium(III) and (b) Cross-sectional view of In2O3 films grown by ALD using tris(N,N'-diisopropylacetamidinato)indium(III). Scale bar indicates 100 nm.

3.4.4 X-Ray Diffraction (XRD) studies

Figure 3.4.4.1 shows the XRD patterns obtained from In2O3 deposited on quartz at different temperatures. XRD data is collected from 15° to 75° 2θ with an increment of 0.05° after each point and a collection time of 5 sec per point. The same data collection parameters are used at all temperatures. An offset is used to avoid overlapping of the patterns. It is seen that at temperatures 130 and 150 °C, the material is likely mainly amorphous with tiny diffraction peaks pointing to early nucleation of first crystallites. At higher deposition temperatures, the main crystalline orientations of In2O3 appear namely (211), (222), (400), (440) and (622) with

diffraction peaks around 21o, 30o, 35o, 51o and 60o respectively (as seen in Figure 3.4.4.1 (b)). To determine the preferred orientation XRD peak intensities need to be normalized to the XRD peaks in the powder spectrum. (222) seems to be the preferred orientation at 200 °C and above.

The other orientations have minute diffraction peaks as seen from Figure 3.4.4.1 (a) and (b).

a) b) 20 25 30 35 40 45 50 55 60 65 70 20 25 30 35 40 45 50 55 60 65 70 2000 2000 1500 1000 325 C 325 C 1500

500 0 2000 1000

1500 300 C 1000 500

500 0 0 2000 1500 275 C 1000 2000 (211) (222)

500 250 C 0 1500 2000 (400) (431) (440) (622) 1500 250 C 1000

1000

500 0 500 2000 1500 225 C 0 1000

500 2000 0 200 C 2000 1500 1500 200 C 1000 1000

500

0 2000 500

Intensity (arb. units) 1500 175 C

1000 Intensity (arb. units) 0

500 0 2000 2000 1500 150 C 150 C 1000 1500

500 0 1000

2000 1500 130 C 500 1000

500 0 0 20 25 30 35 40 45 50 55 60 65 70 20 25 30 35 40 45 50 55 60 65 70 2theta (degree) 2theta (degree)

Figure 3.4.4.1. XRD patterns of In2O3 films on fused quartz substrate using tris(N,N'- diisopropylformamidinato)indium(III) (1); (a) 1300 cycles grown at 150-325 °C and (b) enlarged images for the XRD of films at 150, 200, 250, and 325 °C.

The crystallite size can be determined by the Scherrer equation30 and is tabulated in Table

3.4.4.1. The crystallites range in size from 22 nm to 36 nm. There is no particular trend seen in crystallite size change with temperature although the maximum crystallite size is obtained at 200

°C. Crystallinity of a sample is measured by area under the peaks31. For films of similar

thickness (~ 72 nm) grown from 150-275 °C, the product of peak width and the peak height goes up which results in the peak area going up. The peak area for different temperatures is tabulated in Table 3.4.4.2. Thus the total crystallinity of the samples increases from 150-275 °C (it is noted that samples are of similar thickness of 72 nm). The peak area at 300 and 325 °C is the least and despite the samples being 5-10 nm thicker than other temperatures, the crystallinity is poor. This is consistent with the presence of carbon disrupting crystalline growth at 300 °C and above.

Table 3.4.4.1. Crystallite size as a function of deposition temperature for In2O3 grown using 1.

Deposition Temperature of In2O3 using 1 Crystallite Size by Scherrer Equation

130 °C 24.4 nm

150 °C 22.5 nm

175 °C 32.6 nm

200 °C 36.2 nm

225 °C 30 nm

250 °C 26.2 nm

275 °C 30.5 nm

300 °C 30.3 nm

325 °C 28.5 nm

The crystalline nature of In2O3 samples deposited at low temperatures is also seen in the case of InCp.9 At temperatures of 150 °C highly crystalline films were obtained. Obtaining polycrystalline films in the case of binary In2O3 and In2S3 (as see in chapter 2) does not mean

that we will necessarily obtain crystalline films on alloying the indium oxide with the indium sulfide. This is because the surface reactions occurring in the case of the ternary material will likely be different. Independent of this, since In2O3 and In2S3 have different crystal lattices alloying them to a crystalline material is challenging by itself.

Table 3.4.4.2. Peak area as a function of deposition temperature for In2O3 grown using 1.

Deposition Temperature of In2O3 using 1 Peak Area (CPS x deg)

130 °C 108.2

150 °C 158.1

175 °C 176.3

200 °C 239.3

225 °C 349.4

250 °C 377.1

275 °C 400.3

300 °C 216

325 °C 201.4

o The XRD spectra of In2O3 films grown using 2 are shown in Figure 3.4.4.2. (222) at 30 is seen as the major peak at all temperatures. The other orientations and their diffractions appear as minority peaks at temperatures in the ALD window (225-300 °C), and at 330, 350 °C. The

In2O3 film at 150 °C is likely more crystalline than its counterpart deposited using 1 as the (222) peak intensity is higher for a comparatively smaller thickness of film. Also the peak width-peak

height product is higher for film grown using 2 at 150 °C in comparison to film grown using 1 at the same temperature (see Table 3.4.4.4). The deposition time for the film grown using 2 at 150

°C, is 2000 cycles × 93 seconds in comparison to 1300 cycles × 52 seconds for film grown using

1. Thus we have obtained higher crystallinity by deposition for more time at the same temperature. The Scherrer equation30 can be used to estimate crystallite size and this estimate is tabulated for all temperatures in Table 3.4.4.3. The maximum crystallite size of 43.5 nm is achieved at 300 °C and the crystallite size is mostly constant around 38 nm for majority of temperatures in the ALD window. From Figure 3.4.4.2 (b) it is seen that peak height initially goes up from low temperatures (150 °C) to 275 °C and then goes down at temperatures outside the ALD window like 350 °C.

a) b) 25 30 35 40 45 50 55 60 65 70 20 25 30 35 40 45 50 55 60 65 70 5000 4000 3000 2000 2000 350 C 350 C

1000 1500 0 5000 4000 1000 3000 330 C 2000 500

1000 0 5000 0 4000 3000 300 C (211) (222) 2000 2000 275 C

1000 0 5000 1500 (400) (431) (440) (622) 4000 3000 275 C 2000 1000

1000 0 500 5000 4000 3000 250 C 0 2000

1000 0 2000 225 C 5000 4000 3000 225 C 1500 (411) 2000 1000 1000 0 5000

4000 500 Intensity (arb. units)

Intensity (arb. units) 3000 200 C 2000 0

1000 0 5000 150 C 4000 2000 3000 175 C 2000 1500

1000 0 1000 5000 4000 3000 150 C 500 2000

1000 0 0 25 30 35 40 45 50 55 60 65 70 20 25 30 35 40 45 50 55 60 65 70 2theta(degree) 2theta(degree)

Figure 3.4.4.2. XRD patterns of In2O3 films on fused quartz substrate using tris(N,N'- diisopropylacetamidinato)indium(III) (1); (a) 2000 cycles grown at 150-350 °C and (b) enlarged images for the XRD of films at 150, 225, 275, and 350 °C.

The product of peak height and peak width goes up in the temperature range 150 °C to 300 °C

(see Table 3.4.4.4.) and then decreases at higher temperatures. We can thus make the argument that % crystallinity of sample increases with temperature in the ALD window considering that the thickness of the films is similar in the ALD window of 225 °C – 300 °C. The crystallinity is very poor at high temperatures of 330 °C and above, despite the thicker films. This is likely due to existence of 1-5 % carbon in the film at 330 to 350 °C. The peak area at different deposition temperatures of In2O3 using 2 is shown in Table 3.4.4.4. Thus indium oxide films grown using 1 and 2 show similar crystallinity trends. The indium oxide films obtained using 1 and 2 indicate that (222) is the favored growth plane. Thermodynamic calculations by Walsh32 show that (111) is the lowest energy cleavage plane in In2O3 and so texturing along (222) seen is not surprising.

Table 3.4.4.3. Crystallite size as a function of deposition temperature of In2O3 using 2.

Deposition Temperature of In2O3 using 2 Crystallite Size by Scherrer Equation

150 °C 30.9 nm

175 °C 23 nm

200 °C 27.8 nm

225 °C 38.7 nm

250 °C 38.5 nm

275 °C 38.3 nm

300 °C 43.5 nm

330 °C 37 nm

350 °C 38.3 nm

Table 3.4.4.4. Peak area as a function of deposition temperature of In2O3 using 2.

Deposition Temperature of In2O3 using 2 Peak Area (CPS x deg)

150 °C 252.7

175 °C 275.1

200 °C 278.6

225 °C 293.3

250 °C 326.2

275 °C 326.5

300 °C 326.9

330 °C 211.9

350 °C 185.7

3.4.5 Electrical Properties

The electrical properties shown by Van der Pauw 4 point probe (resistivity) and Hall effect measurements (carrier concentration and mobility) for In2O3 films (using 1) are shown in

Figure 3.4.5.1. Electrons are found to be majority charge carriers for these thin films. Figure

-3 -3 3.4.5.1 (a) shows that resistivity of In2O3 (using 1) ranges from 1×10 to 3.9×10 Ω·cm for temperatures between 150 °C and 325 °C. The mobility of In2O3 films grown using 1 and water

(deposited at 130-325 °C, ALD window: 150-275 °C), increases from 2.4 cm2/V·s to 17.5 cm2/V·s when temperature increases from 130 to 150 °C (see Figure 3.4.5.1 (b)). The carrier concentration increases from 2×1018 to 9.6×1019 cm-3 (Figure 3.4.5.1 (c)) in this same interval.

To rationalize these trends, we consider literature discussions of mobility change of indium oxide with variables like temperature or carrier concentration. Nomura et al. said that post transition metal cation (e.g., In, Sn, Zn, Ga, Pd, Tl, etc.) based amorphous oxide semiconductors which have metal-oxygen-metal bonds show Hall mobilities similar to their crystalline counterpart phase.33 They attribute this observation to metal s orbitals with isotropic shape directly overlapping with neighboring s orbitals.33 Paine et al.34 found that there is no significant difference in electron mobility between amorphous indium oxide (44.5 cm2/V·s) and crystalline indium oxide (heated at 350 °C) when deposited by DC magnetron sputtering. Under

ALD conditions, perfectly amorphous or crystalline films are not produced. Films with crystallinity spanning from mostly amorphous to mostly crystalline are produced according to the kind of indium precursor, oxygen source, reaction time, purging time and deposition temperature. Chang et. al.35 in their study of amorphous indium oxide showed Hall mobility change from perfectly amorphous phase to perfectly crystalline In2O3 phase grown by pulsed laser deposition using variable deposition temperatures of 100 to 600 °C. They proposed that initial crystallization in a completely amorphous phase results in the nano-crystallites acting as scattering centers and thus lowering the Hall mobility with increasing temperature. With further increase in deposition temperature, they report that the crystallites/scattering centers grow, decreasing the mobility until the crystalline phase becomes the major percolation pathway through the material. Once this stage is reached, the amorphous regions become the scattering phase and thus with increase in temperature and decrease of the amorphous phase, Hall mobility goes up. Their XRD studies35 show that the (222) peak is asymmetric and is a combination of the

(222) crystalline peak and an amorphous hump in the temperature range in which mobility decreases. The amorphous hump disappears and symmetric sharp (222) peaks appear when the

temperature of deposition is above a particular value (where the crystalline fraction is 0.82 and above), and the mobility increases in this range.35 It is noted that the maximum Hall mobility of

2 the fully crystalline In2O3 phase (65 cm /V·s) is very similar to its fully amorphous counterpart

(~ 60 cm2/V·s).35 Also, the Hall mobility does not change drastically once a certain level of crystallinity is achieved (600 °C has a mobility of 70 cm2/V·s, vs. 400 °C, with 65 cm2/V·s).35

This is in contrast to 25 cm2/V·s at 200 °C.35 The above discussion shows that crystalline fraction is a major contributor to trends of mobility change. Another important factor which governs change in mobility is ionized impurity scattering36 and the scattering amongst electrons/holes. The 3rd important rationalization of change in mobility is grain boundary scattering.15 The larger the size of the grains, the lower the scattering from the boundaries is and the higher the mobility.

In reference to the change in mobility we observe for In2O3 grown using 1 and water, the mobility is seen to increase from 2.4 cm2/V·s at 130 °C to 48.4 cm2/V·s at 200 °C (see Figure

3.4.5.1(b)). This comes despite the carrier concentration increasing from 2×1018 to 5×1019 cm-3

(3.4.5.1(c)) at these temperatures. This means that ionized impurity scattering or scattering amongst electrons is not the major contributor to the change of mobility. The rationale for this observation is the reduction of the scattering amorphous phase in the already crystallized In2O3 which results in increase of mobility with rise in deposition temperature.35 Beyond 200 °C in the

ALD window the mobility remains constant as a function of temperature (Figure 3.4.5.1 (b) and carrier concentration (seen in Figure 3.4.5.1 (d)). Eventually at the high temperatures of 300 °C and 325 °C outside the ALD window the mobility decreases with increase in carrier concentration likely due to scattering amongst electrons or by carbon impurities.

1.48 1.47 50 1.46 1.45

1.44 40 /V.sec) 1.43 2 1.42 1.41 30 1.40

0.004 20 0.003 10 Resistivity( ohm.cm) Resistivity( 0.002

0.001 ElectronMobility (cm 0 0.000 100 150 200 250 300 350 100 150 200 250 300 350 Temperature (oC) Temperature (oC)

(a) (b)

3.00E+020 50 ) 200 oC

-3 2.50E+020

o cm 40

( 275 C

/V.sec) o

2.00E+020 2 175 C

30 250 oC 1.50E+020 225 oC

1.00E+020 20 o 150 oC 300 C 325 oC 5.00E+019 10

0.00E+000 Electron(cm Mobility

0 Carrier Concentration Concentration Carrier

-5.00E+019 0 1x1020 2x1020 3x1020 100 150 200 250 300 350 -3 Temperature (oC) Carrier Concentration (cm )

Figure 3.4.5.1. For In2O3 films grown with 1 and H2O, (a) Resistivity as a function of deposition temperature, (b) Electron mobility as a function of deposition temperature, (c) Electronic carrier concentration as a function of deposition temperature, and (d) Electron mobility as a function of electronic carrier concentration.

1.0x10-2

60 ) 8.0x10-3 55

/V.sec 2

6.0x10-3 45

cm ( 40

-3 4.0x10 35

Resistivity(ohm.cm) -3 2.0x10 25

20 0.0 Mobility Electron 150 200 250 300 350 15 150 200 250 300 350 Temperature (oC) Temperature (oC)

(a) (b)

2.00x1020 60 150 oC 20

) 1.75x10

55 -3 o 1.50x1020 50 175 C /V.sec) o 2 225 C 20 45 1.25x10 250 oC o 40 200 C 1.00x1020 35 275 oC o 7.50x1019 330 C 30 5.00x1019 o 25 300 C

350 oC Carrier concentration Carrier (cm 2.50x1019 Electron(cm Mobility 20

15 150 200 250 300 350 4.0x1019 6.0x1019 8.0x1019 1.0x1020 o Temperature ( C) Carrier concentration (cm-3)

Figure 3.4.5.2. For In2O3 films grown with 2 and H2O, (a) Resistivity as a function of deposition temperature, (b) Electron mobility as a function of deposition temperature, (c) Electronic carrier concentration as a function of deposition temperature, and (d) Electron mobility as a function of electronic carrier concentration.

-3 The resistivity of In2O3 films grown using 2 and H2O is of the order of 2 x 10 Ω·cm –

-3 3.8 x 10 Ω·cm as shown in Figure 3.4.5.2 (a). For ALD of 2 with H2O (ALD window:225-300

°C), the Hall mobility decreases from 58 to 19.7 cm2/V·s (Figure 3.4.5.2 (b)) as the temperature goes up from 150 °C to 350 °C, and the carrier concentration increases from 4×1019 to 1×1020

100

cm-3 (Figure 3.4.5.2 (c)) . This trend seen in Figure 3.4.5.2 (d) (where the electron mobility drops with increase in carrier concentration except for a couple of outliers), can likely be attributed to ionized impurity scattering. There is also the possibility that scattering by carbon impurities at temperatures above ALD window reduces the mobility. It is hard to confirm that the crystallinity at the lower end of these temperatures (150-200 °C) is high enough that nanocrystallites are not acting as scattering centres and reducing the mobility. The XRD pattern in the case of films grown using 2 does not show asymmetric amorphous peaks like Chang et. al.35 Thus ideally there is a possibility that we are in the regime where mobility increases as a function of crystallinity and temperature. But our observation is contrary to this.

With increase in temperature (especially at higher temperatures, 225 to 300 °C), the crystallinity increases (increase in (222) peak intensity for similar thickness and FWHM), but the mobility decreases. We conclude that mobility change cannot be explained by crystallinity variation for In2O3 grown using 2. Ionized impurity scattering is a plausible explanation for the observed mobility change. There might be other factors like grain boundary scattering or carbon impurity in the film which affect the mobility too but these are not active in the ALD window.

This is because carbon is negligible in the ALD window; also the grain boundary scattering should have reduced at higher temperatures (due to increase in grain size) and resulted in higher mobility. At the temperatures around which there is carbon, 330-350 °C, the mobility drops down possibly because carbon disrupts the crystalline growth.

101

3.4.6 Optical Characterization

Optical transmittance and band gap studies were done on In2O3 samples using a spectrophotometer with an integrating sphere. Figure 3.4.6.1 (a) and (b) show transmittance as a function of wavelength for In2O3 films deposited using 1 and 2 respectively on quartz.

Figure 3.4.6.1. (a) Optical transmittance v/s wavelength for In2O3 deposited using 1 at different temperatures with thicknesses as shown.

102

a) b) 100 130 C 90 90 1 1 80 85 70 130 C 150 C 80 60 175 C

200 C 50 225 C 75

40 250 C 150 C ALD Window ALD 30 275 C 70 225 C ansmittance (%) 130 C: 112 nm (2% C, 3300 cycles) 300 C ransmittance (%) 250 C

Tr ALD win.: ca.73 nm 20 T 1300 cycles 325 C 300 C: 72 nm (1% C) 65 275 C 10 325 C: 83 nm (3% C) 300 400 500 600 700 800 300 400 500 600 700 800

Wavelength (nm) Wavelength (nm) c) d) 100 100 90 2 95 2 80 70 90 60 150 C 175 C 85

50 350 C 200 C

80 40 225 C 250 C 150 C 30 150 C: 58 nm 75 275 C 200 C

20 175 C: 52 nm ansmittance (%) 200 C: 63 nm 300 C ansmittance (%) 225 C

ALD Window ALD 70

ALD win.: ca.100 nm Tr 10 330 C Tr 275 C 330 C: 100 nm (1% C) 2000 cycles 0 350 C: 170 nm (5% C) 350 C 65 300 C

300 400 500 600 700 800 300 400 500 600 700 800 Wavelength (nm) Wavelength (nm)

Figure 3.4.6.1. (b) Optical transmittance v/s wavelength for In2O3 deposited using 2 at different temperatures with thicknesses as shown.

It is seen from Figure 3.4.6.1 (a) and (b) that the transmittance of In2O3 deposited using both 1 and 2 in the ALD window is above 70 % throughout the wavelength range from 400 nm to 850 nm. Thus these films can be used as TCOs in solar cells considering they also have a high carrier concentration at many deposition temperatures in the ALD window, as studied in section

3.4.5 of this chapter. The film deposited at 350 °C with 2 has a transmittance as low as 60 % at certain wavelengths which may be due to the presence of 5 % C in the films.

103

a) b)

0.8 Films in ALD window 0.6

0.7 1 band gap )

) 1 2 2 150 C 3.79 eV 0.6 130 C 0.5 175 C 150 C 3.79 eV 200 C 0.5 175 C 0.4 3.73 eV

225 C 3.78 eV eV/cm

eV/cm 0.4 200 C 12

12 250 C 3.75 eV

225 C 0.3 275 C

0.3 250 C 3.74 eV

x10 x10 (

( 275 C

0.2 2

2 0.2 )

) 300 C 

 325 C h

h 0.1 0.1

 

( ( 0.0 0.0

1.5 2.0 2.5 3.0 3.5 4.0 3.5 3.6 3.7 3.8 3.9 4.0 4.1 Energy (eV) c) Energy (eV)

0.7 (a) (b)

) 1 2 0.6 band gap 130 C 3.71 eV 0.5 2 300 C 3.77 eV Figure 3.4.6.2. Band gap determination using plot of (ᾀhὐ) v/s photon energy for In2O3 films

eV/cm 0.4 325 C 3.80 eV grown using 1 (a) and 2 (b).

0.3

x10

(

2 0.2 )

 In2O3 is known to be a direct band gap material. In order to determine the band gaps of

h 0.1  ( 2 0.0 the In2O3 films, the square of the absorption coefficient times the photon energy ((αhν) ) was 3.5 3.6 3.7 3.8 3.9 4.0 4.1 Energy (eV) plotted against the photon energy. This is the standard Tauc method for direct band gap

materials.37 The absorption coefficient (α) was estimated as –ln(T)/d where d is the thickness of

the film and T is the transmittance.37-38 The band gap values of films grown at different

temperatures were determined from the intercept on the x-axis (photon energy) in Fig. 3.4.6.2.

The band gap values for films grown using 1 at all temperatures (130-325 °C) were found to be

in the range 3.7 to 3.8 eV. Films made using 2 at all temperatures (150-350 °C) also exhibited

band gaps of 3.6 eV to 3.8 eV. Optical band gap values of films grown using 1 and 2 are in the

range of direct band gaps reported in the literature for indium oxides (3.6 eV to 3.8 eV).25, 37, 39

Carrier concentration as obtained in section 3.4.5 is plotted against respective band gaps of In2O3

films (grown with 1) in Figure 3.4.6.3. The optical band gap increases with the carrier

concentration as shown. The band gap increase because of increase in carrier concentration in

104

semiconductors is explained partly by the Burstein-Moss effect.13, 40 According to the method of deposition of the In2O3 film, a particular carrier concentration of a film corresponds to a particular band gap (i.e., the slope of carrier concentration vs. optical band gap curve is different).13, 25, 37, 41 In our case, although the carrier concentration of films grown using 1 is similar to that using 2 at certain temperatures, very small difference in the band gaps are observed (~0.05 eV) likely because of different precursors being used and the deposition conditions being different.

1.30E+020

1.20E+020

) -3

1.10E+020

cm ( 1.00E+020

9.00E+019

8.00E+019

7.00E+019

6.00E+019 CarrierConcentration

5.00E+019

3.73 3.74 3.75 3.76 3.77 3.78 3.79 Optical band gap (eV)

Figure 3.4.6.3. Carrier concentration as a function of optical band gap for In2O3 films grown using 1.

41 Hamberg et.al. have done a detailed study of the optical properties of Sn-doped In2O3 films and show the band gap shift contributions from Burstein-Moss shift, electron impurity scattering, and electron-electron scattering. In our case too band gap shifts are likely due to all 3 contributions but this deconvolution is not trivial.

105

3.5 Conclusions

ALD In2O3 growth with 1 and 2 was studied in the temperature range 130-350 °C. Of the known In2O3 growth processes performed with H2O as the oxygen source, the ALD growth using

1 showed the lowest ALD onset temperature (150 °C) and widest ALD temperature window

(150-275 °C). The ALD process of another amidinate precursor, 2, was found to have an ALD onset of 225 °C and an ALD window from 225 °C to 300 °C. ALD using 1 is proven to have faster kinetics to that using 2. The comparison of ALD windows between In2O3 deposited using

1 and 2 and the kinetic studies show that 1 is more reactive with H2O compared to 2. It is found that a surface saturated with chemisorbed 1 completes surface reactions with H2O 3 to 4 times more quickly compared to a surface saturated with 2 at 250 °C. At lower temperatures of 150 °C, surface saturation with chemisorbed H2O is completed in 4 seconds for 1 while the surface saturation with chemisorbed H2O under ALD mode using 2 is incomplete even after 12 seconds of reaction time (0.3 Å/cycle as opposed to ALD growth rate of 0.52 Å/cycle).

The high rate of surface reactions for 1 likely results in a low temperature onset wide

ALD window. The enhanced reactivity of compound 1 bearing the N,N'- diisopropylformamidinato ligand is potentially very useful in developing precursors for low temperature ALD. There are two advantages to the enhanced reactivity of 1 with H2O. 1) The resultant wide ALD window means that this ALD In2O3 can be coupled with various other metal oxide and metal chalcogenide processes for obtaining ternary materials. We selected the ALD process of 1 with water (window 150-275 °C) to obtain In2O3 to be alloyed with indium sulfide

(ALD window 160-200 °C) (see chapter 2). This combined process can result in oxysulfide materials with controllable composition as desired (see chapter 4). 2) The faster process of ALD between 1 and H2O results in higher growth rate and quicker deposition of an electron transport

106

layer or a TCO. This in addition to a lower temperature process would result in less interfacial interdiffusion of elements in the layers of a thin-film solar cell.

The morphologies and side views of In2O3 grown using 1 and 2 were also compared. The trends seen are similar with feature size going up with temperature in the ALD window and then decreasing at still higher temperatures. The XRD patterns of In2O3 grown using 1 and 2 show that the crystallinity increases with temperature in the ALD window and eventually decreases at the highest temperatures studied. The similar behavior of In2O3 grown by 1 and 2 is likely because they are precursors based on similar amidinate ligands. The electrical characterization shows the difference between the precursors in terms of change in resistivity, carrier concentration, and mobility. The trends seen are explained in section 3.4.5. The important point to note is that between the two precursors the highest mobility seen at the lowest temperature in

2 the ALD window is 48.4 cm /V·s with 1 (and H2O) at 200 °C. This is critical because we know that in order to grow a high mobility oxysulfide material (as an electron transport layer in a solar cell), we can now use a process with interspersed In2O3 and In2S3 (both grown using 1) cycles grown at 200 °C. The mobilities of both the individual binary materials In2O3 and In2S3 is measured to be highest at 200 °C. The ALD process with 1 also provides films with resistivity ~

1×10-3 Ω·cm at all temperatures in the ALD window. Thus films grown with 1 can be used as a

TCO in the widest temperature range. In general, TCOs have resistivity less than 1×10-3 Ω·cm and transmittance greater than 80 % over a wide wavelength range.6 There is not much difference in optical band gaps (~0.05 eV) between In2O3 films grown using 1 and 2. This is not a criterion of selection between the two precursors for alloying indium sulfide with oxygen. The band gap tends to increase with carrier concentration and is rationalized predominantly by the

Burstein-Moss shift.

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3.6 References

2. Yan, C.; Liu, F.; Song, N.; Ng, B. K.; Stride, J. A.; Tadich, A.; Hao, X., Band alignments of different buffer layers (CdS, Zn(O,S), and In2S3) on Cu2ZnSnS4. Applied Physics Letters 2014, 104 (17), 173901.

3. Hariskos, D.; Spiering, S.; Powalla, M., Buffer layers in Cu(In,Ga)Se2 solar cells and modules. Thin Solid Films 2005, 480–481, 99-109.

4. Ellmer, K., Past achievements and future challenges in the development of optically transparent electrodes. Nat. Photon. 2012, 6 (12), 809-817.

5. Kumar, A.; Zhou, C., The Race To Replace Tin-Doped Indium Oxide: Which Material Will Win? ACS Nano 2010, 4 (1), 11-14.

6. Minami, T., Transparent conducting oxide semiconductors for transparent electrodes. Semicond. Sci. Technol. 2005, 20 (4), 35-44.

7. Xirouchaki, C.; Kiriakidis, G.; Pedersen, T. F.; Fritzsche, H., Photoreduction and oxidation of as‐deposited microcrystalline indium oxide. Journal of Applied Physics 1996, 79 (12), 9349- 9352.

8. Kim, N. H.; Myung, J. H.; Kim, H. W.; Lee, C., Growth of In2O3 thin films on silicon by the metalorganic chemical vapor deposition method. physica status solidi (a) 2005, 202 (1), 108- 112.

9. Libera, J. A.; Hryn, J. N.; Elam, J. W., Indium Oxide Atomic Layer Deposition Facilitated by the Synergy between Oxygen and Water. Chem. Mater. 2011, 23 (8), 2150-2158.

10. George, S. M., Atomic Layer Deposition: An Overview. Chemical Reviews 2010, 110 (1), 111-131.

11. Miikkulainen, V.; Leskelä, M.; Ritala, M.; Puurunen, R. L., Crystallinity of inorganic films grown by atomic layer deposition: Overview and general trends. Journal of Applied Physics 2013, 113 (2), 021301.

12. Bugot, C.; Schneider, N.; Lincot, D.; Donsanti, F., Synthesis of indium oxi-sulfide films by atomic layer deposition: The essential role of plasma enhancement. Beilstein J. Nanotechnol. 2013, 4, 750-757.

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13. Maeng, W. J.; Choi, D.-w.; Chung, K.-B.; Koh, W.; Kim, G.-Y.; Choi, S.-Y.; Park, J.-S., Highly Conducting, Transparent, and Flexible Indium Oxide Thin Film Prepared by Atomic Layer Deposition Using a New Liquid Precursor Et2InN(SiMe3)2. ACS Appl. Mater. Interfaces 2014, 6 (20), 17481-17488.

14. Ylivaara, O. M. E.; Liu, X.; Kilpi, L.; Lyytinen, J.; Schneider, D.; Laitinen, M.; Julin, J.; Ali, S.; Sintonen, S.; Berdova, M.; Haimi, E.; Sajavaara, T.; Ronkainen, H.; Lipsanen, H.; Koskinen, J.; Hannula, S.-P.; Puurunen, R. L., Aluminum oxide from trimethylaluminum and water by atomic layer deposition: The temperature dependence of residual stress, elastic modulus, hardness and adhesion. Thin Solid Films 2014, 552 (Supplement C), 124-135.

15. Lee, D.-J.; Kwon, J.-Y.; Lee, J. I.; Kim, K.-B., Self-Limiting Film Growth of Transparent Conducting In2O3 by Atomic Layer Deposition using Trimethylindium and Water Vapor. The Journal of Physical Chemistry C 2011, 115 (31), 15384-15389.

16. Kim, S. B.; Yang, C.; Powers, T.; Davis, L. M.; Lou, X.; Gordon, R. G., Synthesis of Calcium(II) Amidinate Precursors for Atomic Layer Deposition through a Redox Reaction between Calcium and Amidines. Angewandte Chemie International Edition 2016, 55 (35), 10228-10233.

17. Heo, J.; Hock, A. S.; Gordon, R. G., Low Temperature Atomic Layer Deposition of Tin Oxide. Chemistry of Materials 2010, 22 (17), 4964-4973.

18. Gunawan, O.; Virgus, Y.; Tai, K. F., A parallel dipole line system. Applied Physics Letters 2015, 106 (6), 062407.

19. McCarthy, R. F.; Weimer, M. S.; Emery, J. D.; Hock, A. S.; Martinson, A. B. F., Oxygen- Free Atomic Layer Deposition of Indium Sulfide. ACS Applied Materials & Interfaces 2014, 6 (15), 12137-12145.

20. Maeng, W. J.; Choi, D.-w.; Park, J.; Park, J.-S., Atomic layer deposition of highly conductive indium oxide using a liquid precursor and water oxidant. Ceramics International 2015, 41 (9, Part A), 10782-10787.

21. Gebhard, M.; Hellwig, M.; Parala, H.; Xu, K.; Winter, M.; Devi, A., Indium-tris- guanidinates: a promising class of precursors for water assisted atomic layer deposition of In2O3 thin films. Dalton Trans. 2014, 43 (3), 937-940.

22. Ritala, M.; Asikainen, T.; Leskelä, M., Enhanced Growth Rate in Atomic Layer Epitaxy of Indium Oxide and Indium‐Tin Oxide Thin Films. Electrochemical and Solid-State Letters 1998, 1 (3), 156-157.

23. Timo Asikainen; and, M. R.; Leskelä, M., Growth of In2O3 Thin Films by Atomic Layer Epitaxy. J. Electrochem. Soc. 1994, 141 (11), 3210-3213.

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24. Elam, J. W.; Martinson, A. B. F.; Pellin, M. J.; Hupp, J. T., Atomic Layer Deposition of In2O3 Using Cyclopentadienyl Indium: A New Synthetic Route to Transparent Conducting Oxide Films. Chem. Mater. 2006, 18 (15), 3571-3578.

25. Maeng, W. J.; Choi, D.-W.; Park, J.; Park, J.-S., Indium oxide thin film prepared by low temperature atomic layer deposition using liquid precursors and ozone oxidant. Journal of Alloys and Compounds 2015, 649, 216-221.

26. Kim, H. Y.; Jung, E. A.; Mun, G.; Agbenyeke, R. E.; Park, B. K.; Park, J.-S.; Son, S. U.; Jeon, D. J.; Park, S.-H. K.; Chung, T.-M.; Han, J. H., Low-Temperature Growth of Indium Oxide Thin Film by Plasma-Enhanced Atomic Layer Deposition Using Liquid Dimethyl(N-ethoxy-2,2- dimethylpropanamido)indium for High-Mobility Thin Film Transistor Application. ACS Applied Materials & Interfaces 2016, 8 (40), 26924-26931.

27. Ramachandran, R. K.; Dendooven, J.; Poelman, H.; Detavernier, C., Low Temperature Atomic Layer Deposition of Crystalline In2O3 Films. The Journal of Physical Chemistry C 2015, 119 (21), 11786-11791.

28. Nilsen, O.; Balasundaraprabhu, R.; Monakhov, E. V.; Muthukumarasamy, N.; Fjellvåg, H.; Svensson, B. G., Thin films of In2O3 by atomic layer deposition using In(acac)3. Thin Solid Films 2009, 517 (23), 6320-6322.

29. Asikainen, T.; Ritala, M.; Leskelä, M., Growth of In2O3 Thin Films by Atomic Layer Epitaxy. Journal of The Electrochemical Society 1994, 141 (11), 3210-3213.

30. Patterson, A. L., The Scherrer Formula for X-Ray Particle Size Determination. Physical Review 1939, 56 (10), 978-982.

31. Hindeleh, A. M.; Johnson, D. J., Crystallinity and crystallite size measurement in polyamide and polyester fibres. Polymer 1978, 19 (1), 27-32.

32. Walsh, A., Electronic and structural properties of the surfaces and interfaces of indium oxide. 2011; Vol. 1399, p 189-190.

33. Nomura, K.; Ohta, H.; Takagi, A.; Kamiya, T.; Hirano, M.; Hosono, H., Room-temperature fabrication of transparent flexible thin-film transistors using amorphous oxide semiconductors. Nature 2004, 432 (7016), 488-492.

34. Paine, D. C.; Yaglioglu, B.; Beiley, Z.; Lee, S., Amorphous IZO-based transparent thin film transistors. Thin Solid Films 2008, 516 (17), 5894-5898.

35. Buchholz, D. B.; Ma, Q.; Alducin, D.; Ponce, A.; Jose-Yacaman, M.; Khanal, R.; Medvedeva, J. E.; Chang, R. P. H., The Structure and Properties of Amorphous Indium Oxide. Chemistry of Materials 2014, 26 (18), 5401-5411.

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36. Ellmer, K.; Mientus, R., Carrier transport in polycrystalline transparent conductive oxides: A comparative study of zinc oxide and indium oxide. Thin Solid Films 2008, 516 (14), 4620-4627.

37. Weiher, R. L.; Ley, R. P., Optical Properties of Indium Oxide. Journal of Applied Physics 1966, 37 (1), 299-302.

38. Haarindraprasad, R.; Hashim, U.; Gopinath, S. C. B.; Kashif, M.; Veeradasan, P.; Balakrishnan, S. R.; Foo, K. L.; Poopalan, P., Low Temperature Annealed Zinc Oxide Nanostructured Thin Film-Based Transducers: Characterization for Sensing Applications. PLOS ONE 2015, 10 (7), e0132755.

39. Erhart, P.; Klein, A.; Egdell, R. G.; Albe, K., Band structure of indium oxide: Indirect versus direct band gap. Physical Review B 2007, 75 (15), 153205.

40. Liang, S.; Bi, X., Structure, conductivity, and transparency of Ga-doped ZnO thin films arising from thickness contributions. Journal of Applied Physics 2008, 104 (11), 113533.

41. Hamberg, I.; Granqvist, C. G., Evaporated Sn‐doped In2O3 films: Basic optical properties and applications to energy‐efficient windows. Journal of Applied Physics 1986, 60 (11), R123-R160.

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Chapter 4

Atomic Layer Deposition of In2(O,S)3 using

Indium tris(N,Nʹ-diisopropylformamidinate) and its Optical and Electrical Properties

4.1 Chapter Abstract

Indium oxysulfide (In2(O,S)3) thin films, useful as electron transport layers in chalcogenide thin film solar cells, were deposited by atomic layer deposition using the newly developed precursor indium tris(N,Nʹ-diisopropylformamidinate) and co-reactants water and hydrogen sulfide. Controlled incorporation of oxygen in the oxysulfide films was obtained. X- ray photoelectron spectroscopy was used to estimate the content of hydroxyl groups in the thin films. Structural characterization of the films was performed using scanning electron microscopy

(SEM), transmission electron microscopy (TEM) and X-Ray diffraction (XRD). The thin films were electrically and optically studied using Hall effect measurements and UV-Visible spectrophotometry, respectively. A wide range of carrier concentrations from 6×1015 cm-3 to

4×1020 cm-3 was obtained in the films by varying the oxygen to sulfur ratio, thus demonstrating the possibility of use as an electrically graded buffer layer. An increase in the indirect band gap was noted on adding oxygen to indium sulfide thus allowing for lower solar cell absorption loss, relative to In2S3. In this chapter we demonstrate tuning of the conduction band position of indium oxysulfide with respect to CZT(S,Se), by adjusting the S:O ratio.

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4.2 Introduction

Chalcogenide thin film solar cells like (Cu)2ZnSn(S,Se)4 (CZT(S,Se)) and Cu(In,Ga)Se2

(CIGS) have used CdS as a conductive electron transport layer in some of the best devices developed so far.1,2 The toxicity of cadmium and collection losses in the blue and short- wavelength UV regions, owing to a direct band gap of 2.4 eV, remain issues with CdS.

As noted in previous chapters, indium sulfide (In2S3) has recently been studied as a possible nontoxic electrical buffer for chalcogenide absorbers because, despite its smaller bandgap of 2.1 eV, the gap is indirect, resulting in a substantially lower absorption coefficient

3-4 when compared to direct bandgap CdS, thus minimizing the collection losses. In2S3 films have the added possible advantage of minimal interfacial diffusion of the In3+ cation during deposition, and annealing of the solar cells for increasing the grain size and minimizing recombination. A method of improving the photon transmittance of the indium sulfide material is

5-8 known to be alloying it with oxygen to form indium oxysulfide (In2(O,S)3). With the same electron valence and close enough ionic radii, O2− (r = 1.32 Å) compared to S2− (r = 1.82 Å), they are candidates for substituting each other9. The electronic shells of S2−and O2−are also structurally similar, resulting in similar chemical properties of their compounds.10 The substitution of just 5 % of sulfur atoms by oxygen in In2S3 has been previously shown to increase the optical transmission threshold from 2.06 eV to 3.09 eV.11 Ab initio calculations showed that substitutional oxygen impurities produce compressive effects in the lattice11 and there is a clear trend of increasing the band gap as the system undergoes a compression. Of course the bandgap of pure indium oxide is also larger than that of indium sulfide. All of the above reasons are strong pointers to In2(O,S)3 being a good candidate buffer material for thin film solar cells. The ability to tune the conduction and valence bands of the material by changing the S:O ratio in

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In2(O,S)3, in order to match the band positions of CZT(S,Se), CIGS, or SnS is the major reason why In2(O,S)3 is targeted. This adjustment of band positions to optimal values should result in maximum carrier extraction from the solar cell absorber to the buffer. The idea is analogous to previous studies where the energy band alignment between CIGS or SnS absorbers and Zn(O,S) buffers is improved by optimizing the amount of oxygen in Zn(O,S) buffers.12-13

A pinhole-free, uniformly thick In2(O,S)3 buffer is preferred to prevent shunting between the absorber layer and transparent conducting oxide of the solar cells. Solution-based methods like the chemical bath deposition (CBD) used previously to deposit similar buffers like indium hydroxysulfide,14 can have poor control over deposition conditions leading to difficulty tuning

O:S ratios and poor reproducibility. Also, as discussed in chapter 1, CBD is incompatible with the vacuum processes used for depositing other layers in the solar cell, decreasing the production

15 throughput. Sputtering is another method which has been recently used to make In2(O,S)3 compositions,16 but it is generally not preferable as the high energy particles cause defects at the absorber-buffer interface that reduce the integrity of the p-n junction. Atomic layer deposition

(ALD), with its good step coverage and compositional uniformity, has been shown in recent years to be a robust alternative for depositing electron transport layers in thin film solar cells.17-18

ALD uses a sequence of self-saturating surface reactions to grow uniform and conformal thin films with high control over composition.19-21 ALD of indium oxysulfide films is relatively unstudied, with only two recent studies22-23 exploring the use of oxygen plasma and hydrogen sulfide with indium acetylacetonate to synthesize this material. Conventional oxygenation by water or oxygen did not work, so the authors resorted to the use of plasma. It is interesting to note that oxygen plasma helped in promoting the exchange reactions between S and O atoms.

However, plasma could cause the surface damage of the solar absorber layer on which the buffer

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is being deposited; hence it is not an optimal technique for growing films in our application.

Also, the use of a less reactive indium acetylacetonate precursor led to incomplete ALD half reactions leaving 6 % carbon in the film.22-23

In the previous chapter we selected indium tris(N,N’-diisopropylformamidinate) (hereafter referred to as indium formamidinate) for growing a ternary indium oxysulfide composition by

ALD. This choice was made on the basis of fast rate of its ALD reactions with water to form

In2O3, (chapter 3) as well as its fast rate of ALD reactions with hydrogen sulfide to form In2S3

(chapter 2). In this chapter we explore the coupling of these facile ALD reactions into a ternary sequence involving alternate cycles of the highly reactive indium precursor, water (as the oxygen source), and hydrogen sulfide (as the sulfur source) on the surface of a substrate at 200 °C. This ternary process, resulting in an In2(O,S)3 composition with relatively few hydroxyl groups in the film, gives us precise control over the S to O ratio in the films. As discussed in chapter 2, carbon impurities are likely deleterious to the performance of an electron transport layer, as carbon atoms could act as recombination centers resulting in loss of current in solar cells. Our ALD method results in pure, carbon-free films of In2(O,S)3 due to the ALD reactions completely removing the ligands. The adventitious presence of some hydroxyl groups helps obtain a wide range of carrier concentrations in the thin films. The hydroxyl group content in the films is studied by X-ray photoelectron spectroscopy (XPS). Compositions with hydroxyl group below the detection limit under XPS are identified. The electrical and optical properties of the In2(O,S)3 films are studied for application as electron transport layers in solar cells. Structural properties of the In2(O,S)3 are revealed by XRD and TEM. Band offset measurements are performed using

XPS.

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4.3 Experiments

Figure 4.3.1. Formula for indium tris(N,N’-diisopropylformamidinate).

ALD was performed in a custom-built reactor with an oven for heating the indium precursor shown in Figure 4.3.1. The precursor is extremely reactive to moisture and hence was loaded into a glass bubbler in the nitrogen atmosphere of a glovebox. A nitrogen assist gas pressure of 8 torr was used along the precursor line to aid the flow of the precursor. Deionized water (degassed by evacuating air above it and bubbling nitrogen through it) and 4 wt. % hydrogen sulfide in nitrogen (Airgas) were used as co-reactants to grow the In2(O,S)3 films while purified nitrogen was used as the purge gas. No carrier gas assist was used for water. The exposures for indium precursor, water, and hydrogen sulfide used were 1.3 torr·s, 2.4 torr·s, and

1.8 torr·s, respectively. Just 1 s of exposure time was used in all cases. To control the stoichiometry and thus the S:O ratio of the films, a dosing sequence of indium precursor/N2/H2O/N2 (water-containing subcycle) followed by indium precursor/N2/H2S/N2

(hydrogen sulfide-containing subcycle) was used. The reactants were dosed in a closed-valve mode,which involves percursors being trapped in an evacuated chamber for reaction on the substrate, followed by a purge with N2 gas flow. The indium precursor was sublimed at a temperature of 120 °C while the coreactants, H2O vapor and H2S gas, were kept at room temperature. A tube furnace contained the substrates used for the ALD reaction. Substrate temperatures between 120 °C and 200 °C were used for ALD. Variable autotransformers

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regulating an omegalux heating tape were used to maintain the same temperature of the reactor tube at its inlet and outlet.

The stochiometry of the deposited films was measured by x-ray photoelectron spectroscopy (XPS) high-resolution scans. XPS was done with a Thermo Scientific K-Alpha spectrometer equipped with a monochromatized Al Kα X-ray source (1486.6 eV), 12 kV electron beam, and argon ion sputtering gun. Fresh surfaces were prepared for the high-resolution scans using a high current, 500 eV sputtering energy with 30 s of sputtering per level unless otherwise stated. All XPS spectra were recorded using a pass energy of 26 eV and a step size of 0.025 eV.

Curve fitting was done by the “ThermoFisher Scientific Avantage 5.957 surface chemical analysis” deconvolution software. A Lorentzian convolution with a Shirley-type background subtraction was used. The high resolution scans of oxygen are used to detect hydroxyl group content.24-25 Transmission electron microscopy was performed on samples deposited on 50 nm thick Si3N4 membranes supported on 0.5 × 0.5 mm Si grids purchased from Ted Pella, Inc. Van der Pauw and Hall measurement techniques were used to detect carrier type and carrier density.

A custom-built Hall apparatus with a Keithley sourcemeter was used.26 The sample size was 1 cm × 1 cm with a film thickness of around 100 nm. A UV/visible spectrophotometer with an integrating sphere was used to measure the light transmittance and reflectance, to extract the absorption coefficient (α). Films for Hall measurements were deposited on substrates made of silicon covered with 300 nm of thermal oxide while those for optical studies were deposited on quartz substrates. The substrates were cleaned with semiconductor grade acetone and isopropyl alcohol before deposition to remove organic contamination and further treated with UV-ozone to obtain terminal hydroxyl groups for initiating the ALD surface reactions. The film crystallinity is subject to change depending on substrate27 which could affect the properties of the film. The

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substrates used for characterization here, although different from ones used in solar cell applications, should still be representative of the electrical and optical behavior of the films in solar cells or other applications.

4.4 ALD Mechanisms

Based on the structure of the amidinate precursor and the existence of surface thiol and hydroxyl groups, we formulated hypothetical reaction mechanisms that are likely to deposit an

In2(O,S)3 composition. The indium formamidinate precursor has 3 amidinate ligands attached to it. Thus as seen in Figure 4.4.1(a) when the precursor molecule impinges on a surface terminated with multiple thiol groups there is a possibility that 2 of the ligands are protonated and one ligand with the indium ion attaches itself to the sulfur atoms. In the following ALD cycle when hydrogen sulfide arrives at the surface, as seen in Figure 4.4.1 (b), the last ligand attached to indium is protonated and one thiol group is restored. The aforementioned cycle then repeats. This is one of the likely reaction pathways in the indium sulfide subcycle of our In2(O,S)3 ALD process. This is case 1.

Figure 4.4.1. (a) Indium precursor reaction with a thiolated surface (case 1).

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Figure 4.4.1. (b) Reaction of indium amidinate saturated surface with incoming H2S (case 1).

There is another possible reaction pathway (case 2) in the indium sulfide subcycle. Here, one ligand is protonated by a surface thiol, and indium with two amidinate ligands attaches itself to the surface sulfur (seen in Figure 4.4.1 (c)). In this case, in the following step hydrogen sulfide attacks both the existing ligand groups resulting in their protonation and two thiol groups are restored at the ligand positions (see Figure 4.4.1 (d)). This ALD cycle then repeats. It is highly unlikely that all three ligands are protonated in 1st step.

Figure 4.4.1. (c) Indium precursor reaction with a single thiol group on the surface (case 2).

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Figure 4.4.1. (d) Reaction of indium amidinate saturated surface with hydrogen sulfide (case 2).

Similarly in the indium oxide subcycle of the In2(O,S)3 supercycle we have 2 possible reaction pathways. In case 3, when the precursor impinges the surface with hydroxyl groups, two of the ligands are protonated and the indium and its remaining ligand attach to the surface oxygen (See Figure 4.4.1 (e)). In the following half cycle when water arrives the last ligand attached to indium is protonated and one hydroxyl group is restored (See Figure 4.4.1 (f)). This

ALD cycle then repeats. This is one of the likely pathways in the indium oxide subcycle of our

In2(O,S)3 ALD process.

Figure 4.4.1. (e) Indium precursor reaction with a hydroxylated surface (case 3).

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Figure 4.4.1. (f) Reaction of indium amidinate saturated surface with incoming H2O (case 3).

There is another pathway (case 4) in the indium oxide subcycle also. Here, upon arriving at a hydroxyl-terminated surface, one ligand is protonated and indium with two amidinates attaches itself to the oxygen of the suface In-O bond. (seen in Figure 4.4.1 (g)). In this case, in the following step water attacks both the remaining ligand groups resulting in their protonation and two hydroxyl groups are restored at the ligand positions (see Figure 4.4.1 (h)). This ALD cycle then repeats. Similar to the indium sulfide subcycle, it is highly unlikely that all the ligands are protonated at once.

Figure 4.4.1. (g) Indium precursor reaction with a single hydroxyl group on surface (case 4).

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Figure 4.4.1. (h) Reaction of indium amidinate saturated surface with H2O (case 4).

The indium oxysulfide deposition proceeds in cycles of the reactions described in cases 1 through 4, owing to the similarity of the reaction mechanism of the indium precursor with surface thiols and surface hydroxyls. Additional complications can arise, however.

There are certain side reactions which can be expected at high temperatures while depositing In2(O,S)3. The thiol groups on the surface can bridge with a release of H2S as shown in Figure 4.4.1 (i). Also, saturated In-O bonds on the surface can convert to In-S bonds on dosing hydrogen sulfide (Figure 4.4.1 (j)). This is because converting indium oxide to indium sulfide is an exothermic reaction at 200 °C:28

In2O3(s) + 3 H2S (g) → In2S3 (s) + 3 H2O ΔG = -98.4 kJ/mole at 200 °C

Figure 4.4.1. (i) Bridging of thiol groups on an indium-thiolated surface.

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Figure 4.4.1. (j) Conversion of indium oxide bonds to indium sulfide bonds by reaction with

H2S.

It is possible that indium hydroxyls can be converted to indium thiols by a similar process, although the free energy of this reaction cannot be calculated from current thermodynamic tables.

4.5 Results and Discussion

We started with In2(O,S)3 depositions at a low temperature of 120 ° C. This is mainly meant for the substrate configuration of the solar cell, where the electron transport layer is deposited on top of the absorber layer. The lower the temperature of deposition of the n-type

In2(O,S)3 electron transport layer, the better we expect the integrity of the p-n junction with the p-type absorber layer to be. Table 4.5.1 shows a control In2S3 and In2(O,S)3 compositions at 120

°C, as characterized by XPS high resolution scans.

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Table 4.5.1. In2(O,S)3 composition as a function of dosing ratio of water containing to hydrogen sulfide containing cycles.

Sample Substrate Dosing ratio In S O C N

Type temp (In/H2O:In/H2S) (at. %) (at. %) (at. %) (at. %) (at. %)

2 120 °C 2:1 42.4 34 15.6 6 2

1 120 °C 1:1 44.4 40.8 8.5 5.5 0.8

Control 120 °C 0:1 44.5 48 0 6 1.5

In2S3

The In2(O,S)3 compositions at 120 °C contained ~6 at. % carbon and ~1-2 at. % nitrogen.

This indicates that the ALD half reactions do not go to completion and some of the protonated ligands are not completely detached post ALD at such low temperatures. This result is not surprising, as 120 °C is a temperature outside the ALD window overlap of In2O3 and In2S3 which means that In2(O,S)3 is not grown under perfect ALD conditions.

The S:O ratio is approximately 5:1 when a 1:1 dosing ratio of In/H2O:In/H2S is employed. X-ray diffraction of 100 nm thick films of In2(O,S)3 deposited at 120 °C shows the amorphous nature of the films. Thin film solar cells were made with In2(O,S)3 layer deposited at

120 ° C on CZT(S,Se). The entire device stack is shown in the Figure 4.5.1. Device results are shown in Table 4.5.2.

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Figure 4.5.1. CZT(S,Se) device stack used for device studies.

Table 4.5.2. Device results we obtained from the In2(O,S)3 buffer layer incorporated solar cells.

Substrate Dosing In S O C N Eff FF Voc Jsc

temp, ratio (at. (at. (at. (at. (at. (%) (%) (mV) (mA/

2 sample (In/H2O %) %) %) %) %) cm )

Type :

In/H2S)

120 °C 2:1 42.4 34 15.6 6 2 1.11 27.80 298.6 13.29

In2(O,S)3

120 °C 1:1 44.4 40.8 8.5 5.5 0.8 1.14 27.66 289.3 14.30

In2(O,S)3

120 °C In2S3 0:1 44.5 48 0 6 1.5 1.75 27.88 347.8 18.00

160 °C In2S3 0:1 44 54.1 <1 <1 <1 5.83 36.08 433.6 37.27

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We can observe from Table 4.5.2 that addition of oxygen to the control indium sulfide worsens the properties (Jsc and Voc) of the solar cell. This is in stark contrast to our expectation that oxygen addition will tune the energy bands to better align with CZT(S,Se) and hence will give higher open circuit voltage. The open circuit voltage in fact reduces by close to 50 mV while the short circuit current goes down by around 5 mA/cm2 and the net efficiency takes a hit of 4 %. We suspect that presence of 6 at. % carbon and 2 at. % nitrogen in the In2(O,S)3 is the primary reason for the poor performance of these solar cells. Thus we decided to explore higher temperatures for deposition of In2(O,S)3.

The ALD window of In2O3 using indium formamidinate is 150 °C to 275 °C as seen in

Chapter 3. Also, the ALD window of In2S3 using the same precursor is 160 °C to 200 °C as seen in Chapter 2. Studies in Chapter 2 have shown a high electron mobility of 17.7 cm2/V·s for carbon-free In2S3 grown with the same indium precursor at 200 °C. Studies in Chapter 3 have

2 shown a high electron mobility of 48.4 cm /V·s for carbon-free In2O3 grown with the same indium precursor at 200 °C. Targeting a high electron mobility in our In2(O,S)3 films formed using interspersed In2O3 subcycles, we performed our depositions at 200 °C. Interestingly, given the lower ALD onset temperature for each binary material, this was lowest carbon-free deposition temperature for In2(O,S)3. At 160 °C and 180 °C, we obtained small amounts of carbon in the film. 200 °C also lies in the ALD window overlap of In2O3 and In2S3.

The representative high resolution In 3d, O 1s, and S 2p scans (by XPS) for an In2(O,S)3 composition deposited at 200 °C using 0.91 fraction of water containing subcycles (10:1 subcycle ratio of In2O3 to In2S3) is shown in Figure 4.5.2. The fraction of water containing subcycles is defined as the ratio of the subcycles of In2O3 to the total number of subcycles dosed including those corresponding to In2O3 and In2S3. The different compositions of In2(O,S)3 at 200

126

°C on variable dosing of water and hydrogen sulfide containing subcycles, as obtained by XPS high resolution scans, are shown in Figure 4.5.3. It is seen that as the ratio of the number of

In2O3 to In2S3 subcycles in a supercycle varies, the oxide, hydroxyl group and sulfur contents vary while the indium content remains more or less constant. We attempt to study the change in composition from sulfur-rich to sulfur-poor In2(O,S)3 and XPS shows us the correct trend.

The (oxide + hydroxyl) content as obtained from XPS high-resolution scans is seen to increase and sulfur content is seen to decrease proportionally, as the number of water contaning subcycles in a supercycle increases. This is a natural occurance as water is the source of hydroxyl groups and oxygen, and increasing the number of indium oxide subcycles (containing water) increases those species. Figure 4.5.4 shows the ratio of concentrations (OH- +O2-)/S2- changing from from values as low as 0.08 for a film deposited with a low fraction of 0.2 of water-contaning subcyles (which corresponds to 1:4 subcycle ratio) to values as high as 5.4 for a film made with a 0.91 fraction of water containing subcycles (which corresponds to 10:1 subcycle ratio). The (OH- +O2-)/S2- ratio is 0.5 for equal dosage of water and hydrogen sulfide containing subcycles. One of the reasons for higher sulfur incorporation is possibly that H2S is converting either hydroxyls or oxides to thiols or sulfides, as proposed in Figure 4.4.1 (j). Thus we obtain tunability of the composition of the film to heavily oxygen poor and sulfur rich, and vice versa. Simultaneously, the hydroxyl groups in the bulk are higher in concentration when the number of water-containing subcycles is larger. The deconvolution of the O1s scans as seen in

Figures 4.5.5 (a) to (d) show 2 main peaks at binding energies of 531.6 eV and at around 529.9 eV for the O1s scans of different compositions of In2(O,S)3. While the peak at 529.9 eV is assignable to oxygen-indium bond as seen in previous studies, the peak at 531.6 eV could be

- 2- 14 assignable to OH , SO4 or C=O bonds. Considering that the films are relatively free of carbon,

127

2- the peaks are not assigned to C=O. The SO4 peak also occurs at a binding energy of 169 eV as seen in the sulfur 2p scan, but we did not observe its signature in our high resolution S 2p scans.

These facts corroborate the presence of the hydroxyl species in our films which is assigned to the peak at 531.6 eV. On performing the peak area fitting by using the deconvolution software, it is noted that bulk OH- is as high as 3.9 at. % when the fraction of water contaning subcycles is 0.91 while it goes down to relatively low values of slightly above 1 at. % for fewer water containing subcycles (refer Figure 4.5.4).

Eventually when the fraction of water containing subcycles reaches 0.2 there are no hydroxyl groups detected in the samples by XPS. The detection limit of XPS being 1 at. %, we identify 1 at. % to be the upper bound of the hydroxyl group content when we employ a very low fraction (< 0.2) of water containing subcycles to grow In2(O,S)3. This implies that when the fraction of water containing subcycles is lower than 0.2, the samples are nearly free of hydroxyl groups. Hydroxyl groups are usually electron donors in indium oxide and hence it is important to quantify them to know the extent to which they contribute to the number of charge carriers in the system. The hydroxyl groups thus likely control the electrical conductivity of the buffer layers.

Also there is the possibility of hydrogen acting as an extrinsic n-type dopant, as it often does in metal oxides.

128

1000000 (a) Indium Scan 40000 (b) Sulfur scan CPS CPS S2p3/2 800000 In3d5/2 35000 In3d S2p'1/2 3/2 Bck 600000 Bck ENV

ENV 30000

CPS CPS 400000 25000

200000 20000

0 15000 156 158 160 162 164 166 168 170 172 174 176 435 440 445 450 455 460 Binding Energy (eV) Binding Energy (eV)

220000 (c) Oxygen Scan 200000 CPS 180000 O1s OH - 160000 Bck ENV CPS 140000

120000

100000

80000

60000 525 530 535 540 545 Binding Energy (eV)

Figure 4.5.2. Representative XPS In 3d (a), S 2p (b) and O 1s (c) scans for an In2(O,S)3 composition deposited at 200 °C using 0.91 fraction of water containing subcycles. Positions of the In 3d5/2, In 3d3/2, O 1s, S 2p1/2 and S 2p3/2 peaks are typically seen at a binding energy of ~

444.5 eV, ~ 452 eV, ~530 eV, 162.7 eV and ~ 161.6 eV respectively.

129

70 % In 2- - 60 % O +OH % S 50

30 Atomic% 20

0.0 0.2 0.4 0.6 0.8 1.0 Fraction of water containing subcycles

Figure 4.5.3. Chemical composition of the In2(O,S)3 films as a function of the fraction of water- containing subcycles.

6 4 5

4 3

/S ) - - 3

groups bulk in

- +OH

2- 2 O ( 1

at. % OH at. %

0 0

0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1.0 Fraction of water containing subcycles

Figure 4.5.4. Anion ratio in the films and bulk hydroxyl group content as a function of the fraction of water-containing subcycles.

130

140000 529.9eV 531.6eV 135000 CPS 130000 O1s OH - 125000 Bck ENV CPS 120000

115000

110000

105000

100000 525 530 535 540 545 Binding energy (eV)

- Figure 4.5.5. (a) O 1s spectrum for the In2(O,S)3 composition with 49.1 at % O, 3.9 at. % OH , 9 at. % S (films obtained with a 0.91 fraction of water containing subcycles).

110000 531.6 eV 529.9 eV CPS O1s 100000 OH - Bck ENV

90000 CPS

80000

70000

60000 525 530 535 540 545 Binding Energy (eV)

- Figure 4.5.5. (b) O 1s spectrum for the In2(O,S)3 composition with 33 at % O, 1.6 at. % OH , 23 at. % S (films obtained with a 0.75 fraction of water containing subcycles).

131

150000 529.9 eV 531.6 eV 140000 CPS 130000 O1S OH - 120000 Bck Env

110000 CPS 100000

90000

80000

70000

60000 525 530 535 540 545 Binding Energy (eV)

- Figure 4.5.5. (c) O 1s spectrum for the In2(O,S)3 composition with 9 at. % O, 1.3 at. % OH , 48 at. % S (films obtained with a 0.33 fraction of water containing subcycles).

114000 CPS Bck 112000 ENV)

110000

108000 CPS

106000

104000

102000

525 530 535 540 545 Binding Energy (eV)

Figure 4.5.5. (d) O 1s spectrum for the In2(O,S)3 composition with 4 at. % O, below detection limit OH-, 53 at. % S (films obtained with a 0.2 fraction of water containing subcycles).

132

Si S 60 OH- + O2- In 50 C

atomof % element 10

-200 0 200 400 600 800 1000 1200 1400 Etch time (Sec)

Figure 4.5.6. Elemental composition inside a representative In2(O,S)3 film deposited with a 0.25 fraction of water containing subcycles, showing homogeneity of composition.

Figure 4.5.6 shows that the composition of the films remains constant throughout the film. One of the representative compositions of In2(O,S)3 with a water containing subcycle fraction of 0.25 is scanned by high resolution XPS and shows uniformity of elemental content throughout the film. Such uniformity of composition is critical for homogeneous electrical and optical properties of the film.The carbon concentration is below detection limit of XPS, which is about 1 atomic percent. The real growth rate per supercycle of the characterized In2(O,S)3 films with varying atomic % sulfur is different from the ideal growth rate per supercycle of the films calculated, ideal growth rate being the sum of the indium oxide and indium sulfide growth rates at that temperature (See Figure 4.5.7 ).

133

7 7 ) 10:1 )

6 6

/cycle

(

( )

5 5 )

C precursor C

C precursor C o

4 4 o

125 125

120 120

(

S exposure S

S exposure S 2 3 3 2 1:4 2 3:1 2 1:2

1:1 Ideal growth rate rate growth Ideal

1 1 Real rate growth

and 2.7 torr.sec H torr.sec 2.7 and and 1.7 torr.sec H torr.sec 1.7 and 0 0 10 20 30 40 50 60 Atomic % Sulfur

Figure 4.5.7. Ideal growth rate/supercycle vs. real growth rate/supercycle for In2(O,S)3 grown using conditions shown.

Ideal growth rate assumes 125 °C temperature of precursor and 2.7 torr·s exposure of H2S while we grew the films using a precursor temperature of 120 °C and an H2S exposure of 1.7 torr·s. We used undersaturated conditions for both the precursor and hydrogen sulfide so as to obtain conductive In2(O,S)3 films. A film with sulfur vacancies is relatively more conductive than one with fully saturated In-S bonds. Figure 4.5.7 also indicates that growth of In2(O,S)3 is suppressed most at low sulfur content. It is possible that the surface hydroxyls are replaced by surface thiols

o on introducing H2S and the surface thiols further bridge at 200 C (as discussed in section 4.4) to supress the growth of the film.

134

Figure 4.5.8. Transmission electron microscopy scans of indium sulfide (a1,a2) at 200 °C, indium oxide (b1,b2) at 200 °C, and a representative intermediate In2(O,S)3 composition (18 at. % O , 38 at. % S) at 200 °C (c1,c2) deposited on Si3N4 membranes.

It is clearly seen from Figure 4.5.8 that the pure indium sulfide and indium oxide samples made with the same precursor are well crystallized showing ring electron diffraction patterns dispersed with spots representative of a polycrystalline sample. The TEM bright field images of these samples also show lattice fringes and darkness contrast indicative of grains in the sample.

In contrast, the indium oxysulfide sample of an intermediate composition (18 at. % O, 38 at. %

S) shows a very hazy selective area diffraction pattern indicative of poor crystallization. The bright field image of the In2(O,S)3 sample shows very few lattice fringes pointing to its nanocrystalline/mostly amorphous nature. These results are corroborated by XRD on these samples (see Figure 4.5.9). The thickness of all the XRD samples is ~ 100 nm. The XRD of the

o pure In2O3 sample grown at 200 C shows a cubic structure primarily oriented along (222) and the pure In2S3 shows a tetragonal structure oriented along (103). However, the intermediate

In2(O,S)3 compositions with different oxygen contents do not show any XRD peaks and hence appear to be amorphous. The broad peaks seen at 22° correspond to the fused quartz substrate

135

used for the deposition. This amorphous nature of In2(O,S)3 films can be beneficial in solar cells as its conformal deposition on crystalline absorbers prevents direct contact between the absorber and front electrode layers.

16000 (1 0 3) 14000 (2 0 6) (3 0 9) 12000 (4 0 12) In S 10000 2 3

8000 In2(O,S)3 - 18 % O

CPS 6000 In2(O,S)3 - 33 % O 4000 In2(O,S)3 - 49 % O (2 2 2) 2000 (4 0 0) In2O3 0

10 20 30 40 50 60 70 80 

Figure 4.5.9. XRD analyses of In2S3, In2O3 and In2(O,S)3 thin films.

The morphology of a representative indium oxysulfide (18 at. % O, 38 at. % S) sample is shown in the Figure 4.5.10(a) indicating nanocrystalline nature of the films. The top view shows very small nano-sized features which are hardly distinguishable in the cross section (not shown here) revealing a very smooth film. Such smoothness and complete coverage as seen is very critical in solar cells to prevent shunting effects. The morphology of the indium oxide and sulfide samples are different with larger grain sizes of 25 nm and 100 nm (see Figure 4.5.10(b) and

Figure 4.5.10(c) respectively). The XRD, TEM and SEM results thus show that, although the

136

binary ALD reactions separately provide polycrystalline films, the alternation of the H2O and

H2S results in an amorphous In2(O,S)3 structure.

Figure 4.5.10. The top view morphology of (a) In2(O,S)3 (18 at. % Oxygen, 38 at. % Sulfur), (b) In2O3 and (c) In2S3 films grown at 200 °C. Scale bar indicates 100 nm.

The cross-sections of the In2(O,S)3 films also show that these are smooth and conformal films. Figure 4.5.11 (a) shows a representative cross section of an In2(O,S)3 film with the atomic composition of In = 38 %, O = 49.1 %, OH- = 3.9 %, and S = 9 %. Figure 4.5.11 (b) shows the representative cross section of an In2(O,S)3 film with the atomic composition of In = 42.4 %, O =

33 %, OH- =1.6 %, and S = 23 %. The higher sulfur composition films were charging a lot due to their lower electron mobility and electronic conductivity and hence could not be imaged.

137

(a) (b)

Figure 4.5.11. (a) Cross section of In2(O,S)3 film with atomic composition In = 38 %, O = 49.1 %, OH- = 3.9 %, and S = 9 % and (b) Cross section of In2(O,S)3 film with atomic composition In = 42.4 %, O = 33 %, OH- = 1.6 %, and S = 23 %.

The roughness of sulfur-rich and sulfur-poor In2(O,S)3 was measured using atomic force microscopy (AFM). The RMS roughness of In2(O,S)3 films with oxygen-rich composition 49.1 at. % O, 9 at. % S was found to be 0.2 nm. Figure 4.5.12(a) shows the AFM height contrast image of the aforementioned In2(O,S)3 composition. Figure 4.5.12(b) shows the 3-dimensional height contrast view. It is seen that the film is fairly uniform and smooth. Figure 4.5.12(c) shows the height profile across the surface of the film.

Figure 4.5.12. (a) Height contrast image under AFM for In2(O,S)3 with 49.1 at. % O, 9 at. % S. 138

Figure 4.5.12. (b) 3 dimensional height contrast view under AFM for In2(O,S)3 with 49.1 % O, 9 % S.

Figure 4.5.12. (c) Topographical height profile under AFM for In2(O,S)3 with 49.1 at. % O, 9 at. % S.

The RMS roughness of In2(O,S)3 films with sulfur-rich composition 18 at. % O, 38 at. %

S was found to be 0.4 nm. Figure 4.5.13 (a) shows the AFM height contrast image of the aforementioned In2(O,S)3 composition. Figure 4.5.13 (b) shows the 3-dimensional height contrast view. It is again seen that the film is fairly uniform and smooth. Figure 4.5.13 (c) shows the height profile across the surface of the film.

139

Figure 4.5.13. (a) Height contrast image under AFM for In2(O,S)3 with 18 at. % O, 38 at. % S.

Figure 4.5.13. (b) 3-dimensional height contrast under AFM for In2(O,S)3 with 18 at. % O, 38 at. % S.

140

Figure 4.5.13. (c) Topographical height profile for In2(O,S)3 with 18 at. % O, 38 at. % S.

It is thus seen that the oxygen-rich and the sulfur-rich In2(O,S)3 compositions are smoother compared to the pure In2S3 films (RMS = 1.5 nm as seen in Chapter 2).

) 10 3 20

/cm 10 ( 1019 1018 1017

16 Concentration 10

1015 arrier

C 1014 0 10 20 30 40 50 60 In O In S 2 3 at. % Sulfur in In (O,S) 2 3 2 3

Figure 4.5.14. (a) Carrier concentration in the In2(O,S)3 as a function of atom % sulfur.

141

100 50

) 10 40

/V.sec 1

2 30 cm ( 0.1 20

0.01 Resistivity(ohm.cm)

10 Mobility Mobility 0.001 0 0 10 20 30 40 50 60 In2O3 In2S3 at. % Sulfur in In2(O,S)3

Figure 4.5.14. (b) Electron mobility and resistivity of In2(O,S)3 films as a function of atom % sulfur.

1021

)

3 1020 /cm ( 1019 1018 1017 1016

arrier Concentration Concentration arrier 10 C 1014 0 10 20 30 40 50 60 In S In O 2 3 at. % Oxygen in In (O,S) 2 3 2 3

Figure 4.5.14. (c) Carrier concentration in the In2(O,S)3 as a function of atom % oxygen.

142

100 50

) 10 40

/V.sec 1

2 30 cm ( 0.1 20

0.01 Resistivity(ohm.cm)

10 Mobility Mobility 0.001 0 0 10 20 30 40 50 60 In2S3 In2O3 at. % Oxygen in In2(O,S)3

Figure 4.5.14. (d) Electron mobility and resistivity of In2(O,S)3 films as a function of atom % oxygen.

The electrical properties as revealed by Van der Pauw and Hall measurements are shown in Figure 4.5.14. Electrons are determined as the charge carriers for these thin films and a wide range of tunability of electron concentration is seen. Figure 4.5.14 (a) shows that the number of

19 -3 electrons are in the mid 10 cm range for the pure In2O3 films. The number of carriers goes

20 -3 above 10 cm for oxygen-rich In2(O,S)3 thin films (49.1 at. % O, 9 at. % S) . This composition is rich in hydroxyl groups (3.9 %) which likely act as electron donors that generate the large number of carriers. Indium sulfide films are less conductive than indium oxide films due to the deeper sulfur vacancy donors compared to the oxygen vacancies. Hence addition of sulfur to indium oxide reduces the conductivity, but we notice the opposite trend initially due to a high content of hydroxyl groups in films with low sulfur content. With further addition of sulfur, as seen from the measurements for the other In2(O,S)3 films, the carrier concentration reduces to

2×1019 cm-3 for 23 at. % sulfur containing films and further to 5×1016 cm-3 for the 38 at. % sulfur containing films. This trend is reaffirmed by the fact that the content of hydroxyl groups goes

143

down from 3.9 % for the 9 at. % S containing films to 1.6 % for the 23 at. % S containing films and further on to 1.35 % for the 38 at. % S containing films. The decrease in % hydroxyl groups is thus representative of fewer electron donating groups in the thin films. At the other end of the

16 -3 spectrum, the carrier concentration is 1×10 cm for the pure In2S3 in which the carriers are contributed by indium interstitials along with the sulfur vacancies.29-30 Thus we are able to control the number of carriers in situ by tuning the S2-:(O2-+OH-) ratio in the material. This is very critical in solar cells because a lower number of carriers is usually preferred in the buffer layer near the p-n junction interface especially when there are many recombination sites. Away from the interface, closer to the transparent conducting oxide (TCO), a high carrier concentration is preferable to have minimal contact resistance with the TCO. One could now imagine growing a film of indium oxysulfide with a recipe changing over the course of the deposition, starting with sulfur-rich at the absorber side and ending with sulfur-poor at the TCO side, meeting both of these carrier concentration needs.

Pure indium sulfide grown with indium acetamidinate and indium acetylacetonate have been previously reported to have a carrier concentration of 2×1018 cm-3 and 1×1017 cm-3 respectively.31-32 Thus, in comparison to these precursors, the indium formamidinate precursor possibly produces fewer indium interstitials and sulfur vacancies. It is plausible that the number of sulfur vacancy donors goes down with oxygen incorporation and hence the number of carriers

16 -3 15 -3 initially goes down from 10 cm for pure In2S3 (~ 58 at. % S) films to 6×10 cm for ~ 48 at.

% sulfur containing In2(O,S)3 film compositions. Figure 4.5.14 (b) shows the electron mobility and resistivity variation as a function of atom % sulfur in the In2(O,S)3 films. The indium oxide

2 films have a high mobility of around 48.4 cm /V·s. The In2(O,S)3 films show a lower mobility of around 10 cm2/V·s for all the sulfur compositions tested. This lower number is expected because

144

of the disordered nature of the films as evidenced by XRD and TEM studies. Also this stands reasonably in comparison to some of the other ternary buffer compositions like Zn(O,S) tested

17-18 under ALD conditions. Unlike Zn(O,S), all the mixed In2(O,S)3 compositions show similar

2 mobility values. The electron mobility of pure In2S3 polycrystalline films is around 17 cm /V·s.

This is a much lower value than bulk single crystalline In2S3, which has mobility around 200 cm2/V·s.30

The carrier concentration can also be plotted as a function of oxygen content in the

In2(O,S)3 . This is shown in Figure 4.5.14 (c). The oxygen substitution for sulfur results in increase in carrier concentration for the In2(O,S)3 studied. The electron mobility and resistivity variation can also be represented with respect to change in oxygen content of In2(O,S)3 . Figure

4.5.14 (d) depicts the electron mobility and resistivity variation as a function of atom % oxygen in the In2(O,S)3 films. The mobility of the In2(O,S)3 remains mostly constant with addition of oxygen while the resistivity increases as shown in Figure 4.5.14 (d)

The transmittance of the In2(O,S)3 films in the UV-Vis region is shown in Figure 4.5.15, in comparison to In2O3 and In2S3 films. All the films are around 100 nm in thickness.

145

100

o In O at 200 C 40 2 3 o In2(O,S)3 ~49 at. % O at 200 C o In (O,S) ~30 at. % O at 200 C % Transmittance % 20 2 3 o In2(O,S)3 ~6 at. % O at 200 C o In S at 200 C 0 2 3

200 300 400 500 600 700 800 900 1000 Wavelength (nm)

Figure 4.5.15. % Transmittance of In2(O,S)3 thin films in the UV-Visible region of the light spectrum.

It is noted that while the indium sulfide films grown at 200 oC transmit the least in the UV-vis region, the indium oxide films are most transmitting (around 80 %). The mixed In2(O,S)3 compositions show that higher the oxygen content in the film, the higher the transmittance in the

UV region. This is one of the reasons why oxygen addition is preferable in these buffer materials as it increases the light transmitted to the absorber.

We have plotted the absorption coefficient against the photon energy for various

33 stoichiometries of In2(O,S)3. Tauc’s band gap relation for both direct and indirect transitions were plotted as shown in Figure 4.5.16 (a) and 4.5.16 (b), respectively. In the case of indium oxide, which is known to be a direct band gap material, only a direct transition corelation was plotted.

146

0.5 α(hν) (hν Eg) - For direct transitions

2 α(hν) (hν Eg) - For indirect transitions

where α is the absorption coefficient, Eg is the optical bandgap, and hν is the photon energy. The effective mass of electrons and holes is assumed to be constant. The optical band gap is obtained by linear extrapolation of the square of the absorption coefficient (or the square root of the absorption coefficient) plotted against photon energy, near the absorption edge. As seen in

Figure 4.5.16 (a), the direct band transition does not give a distinct linear extrapolation near the absorption edge for the In2(O,S)3 films while the extrapolation for the direct band gap In2O3 gives a band gap of around 3.7 eV. It is seen that the extrapolation near the absorption edge is more nearly linear when we assume an indirect band transition for the In2(O,S)3 films (See

Figure 4.5.16 (b). A small increase in the band gap from around 2.2 eV for the In2S3 to around

2.4 eV for the In2(O,S)3 compositions is observed. Although the band gap is rather low, similar to the CdS buffer layers (2.42 eV direct band gap34), the absorption in the blue region is still comparatively lower for In2(O,S)3 due to the indirect nature of the band gap. The increase in band gap on alloying with oxygen is smaller than the amount found in a previous study where a band gap increase of 1.1 eV was observed.22 However, this larger increase could be due to the

22 higher oxygen content in all the In2(O,S)3 compositions measured previously. This oxygen concentration is higher than the oxygen contents in In2(O,S)3 evaluated in our study. The presence of 6 atomic % carbon in the films22 could also have affected the band gap.

147

12 o 4x10 In2O3 at 200 C o

In2(O,S)3 - 49 at. % O at 200 C

2 o ) In2(O,S)3 - 33 at. % O at 200 C 3x1012 o

eV In2(O,S)3 - 18 at. % O at 200 C

-1 o In2S3 at 200 C

cm Linear fit for In O , x=3.7 eV ( 12 2 3

2 2 2x10

)



(  ( 1x1012

0 0.5 1.0 1.5 2.0 2.5 3.0 3.5 4.0 4.5

h (eV)

Figure 4.5.16. (a) Direct band gap approximation for determining band gap of indium based compounds.

In2S3 3 o 2.0x10 In2(O,S)3 - 18 at % O at 200 C

 o

In2(O,S)3 - 33 at % O at 200 C 0.5 o

1.5x103 In2(O,S)3 - 49 at % O at 200 C

eV)

-1 (cm

 3

 1.0x10



)

) 

h 2

( 5.0x10

 (

0.0

0.5 1.0 1.5 2.0 2.5 3.0 3.5 4.0 4.5 5.0 h (eV) Figure 4.5.16. (b) Indirect band gap approximation for determining band gap of indium based compounds.

148

The band gap increase obtained in our In2(O,S)3 compositions in comparison to pure

35 In2S3 may be partially produced by the Burstein-Moss effect, as the carrier concentrations in the higher oxygen containing In2(O,S)3 films (49.1 at. % O) are up to degenerate levels (above

20 3 10 /cm ). In the lower oxygen containing 18 at. % O In2(O,S)3 films the band gap increase cannot be explained by the same effect because the carrier concentrations are fairly low, on the order of 1016 electrons/cm3. Previous studies by Matar et al.36 have shown that oxygen orbitals

37 contribute to both the conduction and valence bands of In2O3. Similarly Robles et. al. have concluded that sulfur orbitals contribute to both conduction and valence bands of In2S3 . There is thus an electronic effect leading to an increase in optical band gap as both conduction and valence bands shift in the mixed composition materials as a function of oxygen content. Other studies11 have argued that lattice compression leads to optical band gap increase. However in our case of an amorphous structure, this mechanism for the band gap increase cannot be tested.

The In2(O,S)3 compositions we studied are moisture resistant and also largely resistant to air oxidation. We conducted tests where we exposed the films to air for 2 weeks and remeasured the composition by XPS. The composition remained mostly constant with variability of ±3 %.

Thus exposing the electron transport layer to air before depositing the absorber (in the superstrate configuration of the solar cell) is not problematic. This finding is consistent with the thermodynamics presented previously, which suggest that substitution of sulfur for oxygen is favorable in indium oxysulfide (and therefore, substitution of oxygen for sulfur is not).

In order to check whether In2(O,S)3 provides pinhole-free coverage over CZT(S,Se), we deposited 30 nm (nominal buffer layer thickness) of In2(O,S)3 at 200 ºC on CZT(S,Se) substrates and performed high resolution (high res.) XPS scans of substrate elements like Cu and Zn which are prone to diffusion during growth of buffer. It was interesting to note that there was no

149

signature of Cu or Zn at the surface of In2(O,S)3. This confirmed that the respective elements had not diffused all the way from the substrate to the surface of In2(O,S)3. But this result did not preclude the possibility of interdiffusion at the interface between In2(O,S)3 and CZT(S,Se). Thus we etched the surface of the sample with argon ions to check the bulk composition of In2(O,S)3.

Despite etching for 40 seconds, we did not see resolvable signature of Cu 2p and Zn 2p peaks as shown in Figure 4.5.17 (a) and (b) respectively.

(a)

(b)

Figure 4.5.17. High res. XPS scans of Cu (a) and Zn (b) after 40 s etch on In2(OS)3/CZT(S,Se).

150

XPS results prove that within the detection limit (1 at. %), Cu and Zn have not diffused at least to the depth inside In2(O,S)3 that 40 seconds of etching results in. Thus coverage of In2(O,S)3 is good but there is still the possibility that Cu and Zn have diffused in miniscule amounts near the interface. Studies by atom probe tomography might help resolve if there is diffusion right at the interface.

Although ALD In2(O,S)3 is a very uniform film as seen from SEM images shown in

Figure 4.5.11, it is useful to do high resolution XPS scans for substrate elements at multiple locations to confirm complete coverage within the detection limit of XPS. Extremely small pinholes are likely not detected by XPS because of the photoelectron signal to noise ratio coming from the holes being very low.

4.5.1 Photoluminescence Experiments

Photoluminescence experiments were performed on the In2O3 and In2(O,S)3 samples by using a Horiba multiline Raman spectrometer. A 532 nm laser excitation was used. The tool has an 800 mm spectrometer with 600 blaze grating and synapse CCD detector. The samples were placed on an XYZ motorized stage and focused and calibrated with respect to silicon.

Figure 4.5.1.1 shows the photoluminescence spectra of In2O3 and In2(O,S)3 samples obtained with a 532 nm laser excitation. The Si-O Raman stretch is seen on all samples at around

561 nm and this is because of thermal oxide substrates used for deposition. The intensity of the peak at 561 nm is different for different compositions of In2(O,S)3 and In2O3. This is likely due to the different absorptions in the respective compositions preceding the Si-O bond stretch.

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2.34 eV In2O3

In2(O,S)3 - 9 % S 6000 2.29 eV Si-Si Raman stretch In2(O,S)3 - 23 % S

In2(O,S)3 - 38 % S 5000

4000

3000 Counts per sec per Counts 2000

1000 Si-O Raman stretch

0 530 540 550 560 570 580 590 600 Wavelength (nm)

Figure 4.5.1.1. Photoluminescence spectra of In2O3 and In2(O,S)3 samples obtained with a 532 nm laser excitation.

The sharp high intensity Si-Si Raman stretch is seen clearly at 547 nm and is also indicative of the thermal oxide substrate (300 nm SiO2 on Si) used. Photoluminescence peaks are seen at 2.29 eV and 2.34 eV for all the different indium oxysulfide compositions with the intensity of the peaks in CPS decreasing with the decrease in sulfur content of In2(O,S)3 . The pure In2O3 composition shows a very small peak at 2.29 eV. These results essentially mean that, with the decrease in sulfur content, the samples luminesce less from recombination of electrons with holes at defect level energies separated from the conduction or valence band by 2.29 eV and

2.34 eV. The higher the sulfur content, the greater the photoluminescence from these 2.29 eV and 2.34 eV defect levels. Interestingly it has been noted in photoluminescence studies of In2S3 that defect levels of sulfur vacancy (Vs) and indium vacancy (Vin) are separated by 2.19 eV while defects like In interstititals and OVs (oxygen on a sulfur vacancy) are separated by 1.88 eV

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38 within the band gap of In2S3. The In2O3 sample is almost free of defect related photoluminescence peaks. Above results alternatively mean that with more oxygen addition to

In2(O,S)3, we obtain compositions with fewer of the aforesaid defect levels. This could be another beneficial aspect of adding oxygen in situ by ALD to In2S3 . Thus adding oxygen to

In2S3 results in buffer compositions with lesser probability of recombination at certain defect levels. It can be argued however that the oxide could provide non-emissive defect levels which allow for faster, non-radiative recombination of the electron and hole. Further measurements like carrier lifetime are needed to verify the same.

4.5.2 Band offset measurements of In2(O,S)3 with CZT(S,Se): Requirements for obtaining a high open circuit voltage

As discussed in chapter 1, In2(O,S)3 should ideally have a small spike type conduction band offset with CZT(S,Se) to obtain a high open circuit voltage. The band offset of In2(O,S)3 compositions with respect to CZT(S,Se) was measured by using the Kraut method.38 CZT(S,Se) samples were deposited by Dr. Priscilla Antunez at IBM Thomas J Watson research labs. The band gap of CZT(S,Se) was measured by UV-Vis spectrophotometry to be 1.1 eV. 3 different samples were used for the band offset measurement namely a 100 nm bulk sample of In2(O,S)3

(of which composition was varied) deposited at 200 oC, a 2 μm thick bulk CZT(S,Se) sample

o and 3 nm In2(O,S)3 of the same composition deposited at 200 C on the 2 μm thick CZT(S,Se) sample. XPS measurement is performed on the aforementioned 3 samples to obtain core levels of the indium and copper cations respectively. Valence band edge scans from XPS are obtained to determine the valence band maximum of In2(O,S)3 and CZT(S,Se) respectively.

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- Let us consider the example of In2(O,S)3 with composition of In – 42.5 %, OH – 1.5 %, S –

38 %, O – 18 % (measured by XPS depth profile) . 3 nm of this composition was deposited at

200 oC on CZT(S,Se); this sample shall be referred to as the layered sample from here on.

Besides this bulk samples of 100 nm In2(O,S)3 and 2 μm CZT(S,Se) were used for the measurement. The core levels of indium and copper were measured by XPS on the layered sample. The difference in these core levels at the interface is used for the band offset measurement. The interfacial core level spectra of indium and copper from XPS are shown in

Figure 4.5.2.1(a) and 4.5.2.1 (b) respectively. The difference between the valence band maximum and the core level in bulk CZT(S,Se) and bulk In2(O,S)3 are extracted from XPS from valence band edge and copper and indium core level scans for the respective samples.

The valence band edge scan and copper core level scan for CZT(S,Se) are shown in Figure

4.5.2.2 (a) and Figure 4.5.2.2 (b) respectively. The valence band edge scan and indium core level scan for In2(O,S)3 are shown in Figure 4.5.2.3 (a) and Figure 4.5.2.3 (b) respectively

Using all the aforementioned parameters, the valence band offset (∆EV) between CZT(S,Se)

39 and In2(O,S)3 can be extracted using the following formula

CZT(S,Se) CZT(S,Se) In (O,S) In (O,S) ∆EV = (EVBM – ECu2p ) + ∆ECL - (EVBM 2 3 – EIn3d 2 3)

In (O,S) /CZT(S,Se) In (O,S) /(CZTS) where ∆ECL = ECu2p 2 3 –EIn3d 2 3

= (0.31-931.88) + (932.38-444.98) – (1.37-444.48) = - 1.06 eV

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130000 400000 In3d5/2 Cu2p3/2

120000 350000

300000 In3d3/2 110000 250000 100000 Cu2p1/2 200000

90000 Counts/second 150000 Counts/second

80000 100000

50000 70000 444.98 eV 932.38 eV 435 440 445 450 455 460 920 930 940 950 960 970 Binding energy (eV) Binding Energy (eV)

(a) (b)

Figure 4.5.2.1. (a) Core level spectra of Cu 2p for the layered sample and (b) Core level spectra of In 3d for the layered sample.

500000 25000 Equation y = a + b*x Cu 2p3/2 450000 Plot B 20000 Weight No Weighting 400000 -2214.34286 ± Intercept 265.99061 350000 15000 7922.85714 ± Slope 125.57909 300000 Cu 2p1/2 10000 Residual Sum 88312.59714

of Squares 250000 Counts/second Counts/second Pearson's r 0.99937 200000 5000 R-Square(CO 0.99875 Adj. R-Square 0.99849 150000 0 931.88 VB edge intercept = 0.31 eV 100000 920 930 940 950 960 970 -6 -4 -2 0 2 4 6 8 10 12 Binding energy (eV) Binding energy (eV)

(a) (b)

Figure 4.5.2.2. (a) Valence band edge for the CZT(S,Se) sample showing band edge extrapolation to 0 counts and (b) Cu 2p core level spectra for the CZT(S,Se) sample.

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2500

Equation y = a + b*x 500000 In3d5/2 Plot B 2000 Weight No Weighting -1631.02144 ± 354. 400000 Intercept 90797 In3d3/2 1500 1190.12783 ± 116.0 Slope 2348 300000 Residual Sum of Sq 769994.82037 1000 uares 200000

Pearson's r 0.95562 Counts/second

R-Square(COD) 0.91321 Counts/second 500 Adj. R-Square 0.90453 100000

VB Edge intercept = 1.37 eV 0 0 444.48 eV -10 0 10 20 30 40 435 440 445 450 455 460 Binding Energy (eV) Binding Energy (eV)

(a) (b)

Figure 4.5.2.3. (a) Valence band edge for the In2(O,S)3 sample showing band edge extrapolation to 0 counts and (b) In 3d core level spectra for the In2(O,S)3 sample.

To elucidate, the valence band to core level difference in In2(O,S)3 when subtracted from the sum of interfacial core level difference and valence band to core level difference in

CZT(S,Se) , gives the valence band offset. Thus, for the In2(O,S)3 under consideration we obtain a valence band offset of - 1.06 eV. 3 such measurements using 3 sets of bulk In2(O,S)3 , bulk

CZT(S,Se) and layered samples at each composition of In2(O,S)3 are made to determine mean and error bars. The valence band offset of - 1.06 eV is sufficiently high for a hole blocking layer which prevents holes from crossing over from absorber to buffer.

One can use the valence band offset and the band gaps of CZT(S,Se) and In2(O,S)3 to determine the conduction band offset (∆Ec ) as follows

CZT(S,Se) In (O,S) ∆Ec = ∆EV + (-Eg +Eg 2 3 )

∆Ec =-1.06 eV + (2.4-1.1) eV = 0.24 eV spike (The band gap of the In2O,S3 determined by UV-

Vis spectrophotometry, as discussed before in this chapter, is 2.4 eV.)

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Thus we obtain a positive/spike conduction band offset of 0.24 eV for In2(O,S)3 (composition of

In – 42.5 %, OH- – 1.5 %, S – 38 %, O – 18 %) with respect to CZT(S,Se). The band diagram for the CZT(S,Se) - In2(O,S)3 interface can be plotted using the values of conduction and valence band offsets and the band gaps. It is shown in Figure 4.5.2.4.

- Figure 4.5.2.4. Band diagram at the interface of In2(O,S)3 (In – 42.5 %, OH - 1.5 %, S - 38 %, O - 18 %) and CZTS,Se (relative core levels and valence band levels determined by XPS, band gap by UV-Vis spectrophotometry).

Analogous to what was shown above, the valence and conduction band offsets of all the different compositions of In2(O,S)3 deposited in this study is determined with respect to

CZT(S,Se). The relative difference between valence band maximum and core level was used to obtain the valence band offsets of the different In2(O,S)3 compositions. The variation in valence band offset with respect to change in oxygen content of In2(O,S)3 is shown in Figure 4.5.2.5. It is noted that while 0 atom % oxygen corresponds to pure In2S3, 60 atom % oxygen corresponds to

157

pure In2O3. It is seen that In2S3 has a pretty high absolute value of valence band offset of 0.97 eV.

On adding 4 atom % oxygen, the resultant In2(O,S)3 composition shows an increased valence band offset (VBO) of 1.15 eV with CZT(S,Se). Further addition in oxygen in In2(O,S)3 results in decrease in valence band offset up to 0.8 eV for In2(O,S)3 containing 49 % oxygen. The highest value of valence band offset of 2.27 eV is obtained for the In2O3 composition with 60 % oxygen.

The reduction in valence band offset across In2(O,S)3 compositions as seen in Figure 4.5.2.5 is likely because of increase in charge carriers as seen in Figure 4.5.2.7. The increase in carrier concentration in the buffer layer (with oxygen addition) results in relatively increased band bending in the absorber and thus the valence band shifts further down. This results in decrease of valence band offset.

2.4

2.2

2.0

1.8

1.6

1.4

1.2 Valence Band offset Band Valence 1.0

0.8

0.6 0 10 20 30 40 50 60 In2S3 In2O3 at. % Oxygen in In2(O,S)3

Figure 4.5.2.5. Valence band offset variation with change in oxygen content in In2(O,S)3.

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The VBO results indicate that all the In2(O,S)3 compositions and the In2O3 and In2S3 compositions have high enough a valence band offset with CZT(S,Se) to prevent hole flow from the absorber into the respective buffer layers. It has been noted that valence band offsets of

40 above 0.8 eV are sufficiently high to block holes. Thus all the In2(O,S)3 compositions tested are good hole blocking layers.

Using the band gaps and the valence band offsets of the various indium compounds, one can obtain the conduction band offset as discussed earlier in this section. The variation in conduction band offset with change in atom % oxygen is shown in Figure 4.5.2.6. It is noted that while 0 atom % oxygen corresponds to pure In2S3, 60 atom % oxygen corresponds to pure In2O3 .

It is seen from Figure 4.5.2.6 that the conduction band offset of indium compounds with

CZT(S,Se) increases monotonically from 0.14 eV spike for pure In2S3 to 0.5 eV on average for an In2(O,S)3 composition containing 49 atom % oxygen. The conduction band offset drops down though to a mean of 0.36 eV for pure In2O3. The rising trend seen in conduction band offset across In2(O,S)3 compositions with addition of oxygen is a result of the increased band bending in the CZT(S,Se) layer due to increase in carrier concentration of the buffer. In recent literature on device simulations41,42 of chalcogenide CIGS solar cells it was determined that a small conduction band spike offset in the range 0 to 0.4 eV (of CdS buffer with respect to CIGS ) results in good cell efficiency despite high interfacial defect density.43

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0.80 0.75 0.70 0.65 Tunability from +0.1 to +0.54 eV spike 0.60 0.55 0.50 0.45 0.40 0.35 0.30

0.25 Conduction Band offset Band Conduction 0.20 0.15 0.10 0.05 0 10 20 30 40 50 60 In S In O 2 3 at. % Oxygen in In (O,S) 2 3 2 3

Figure 4.5.2.6. Conduction band offset variation with change in oxygen content in In2(O,S)3.

CZT(S,Se) being very similar to CIG(S,Se) in terms of its material property variations, we hypothesize that a small spike band offset of 0.1 to 0.4 eV corresponds to sufficient band bending leading to rectification in the device and is not too high to prevent transfer of photogenerated electrons from the CZT(S,Se) to the buffer layer. Although we get tunability of conduction band offset from 0.1 eV to 0.54 eV by changing the atom % oxygen in the indium based compound, only compositions that give a small offset of 0 to 0.4 eV are preferred for the buffer layer (this range is demarcated in Figure 4.5.2.6). Going by this rule of thumb, In2S3,

In2O3 and In2(O,S)3 compositions with 4 to 33 % oxygen should be targeted as buffer layers for

CZT(S,Se) to obtain high efficiency. But it is seen from Figure 4.5.2.7 that the carrier

16 concentration of In2S3 and the In2(O,S)3 with 4 % oxygen is very low (~1x10 or less). This implies that these compositions will likely result in high series resistance and low fill factor of

19 the devices. The high carrier concentrations above 1x 10 of In2 (O,S)3 with 33 to 49 % oxygen

160

and In2O3 indicate that these compositions are not suitable at the absorber – buffer interface because this would result in heavy recombination with absorber holes especially if there are recombination pathways like defects. Considering all the aforementioned points, In2(O,S)3 buffers in the composition range in between 4 and 33 % oxygen would be optimal in combination with CZT(S,Se), to obtain a high open circuit voltage. Of the allowed compositions, one giving a higher or a lower conduction band position can be chosen depending on the band positions of the CZT(S,Se).

1021

)

3 1020 /cm ( 1019 1018 1017 1016

arrier Concentration Concentration arrier 10 C 1014 0 10 20 30 40 50 60 In S In O 2 3 at. % Oxygen in In (O,S) 2 3 2 3

Figure 4.5.2.7. Variation in carrier concentration as a function of atom % oxygen in In2(O,S)3 compositions.

Thus it is seen that we are able to obtain significant tunability in terms of conduction band, valence band positions and conductivity of the various In2 (O,S)3 compositions and we can employ a composition that has the best match with CZT(S,Se) of a given band gap to obtain high open circuit voltage and solar cell efficiency. It is noted that within the allowable range, an

In2(O,S)3 composition with a relatively lower conductivity and a spike offset of 0.1 to 0.4 eV is

161

suitable near the absorber-buffer interface to provide significant electron collection from the absorber with minimal recombination. In contrast, compositions with higher conductivity are preferable near the TCO-buffer interface to minimize contact resistance. The different composition requirements at the 2 interfaces (absorber- buffer and buffer-TCO) can be obtained by using atomic layer deposition to control the S:O ratio during growth.

4.6 Conclusions

The growth of carbon-free In2(O,S)3 films by ALD was demonstrated. The composition of the films can be tuned from sulfur-rich to sulfur-poor by adjusting the ratio of water- containing cycles to hydrogen sulfide-containing cycles . The variable S:O ratio in films allows us to tune the conduction and valence band offsets with absorbers like CZT(S,Se).

The content of hydroxyl groups can be reduced to negligible levels in the films deposited with less than 20 % water-containing subcycles. The In2(O,S)3 films were found to be amorphous with an increased band gap compared to pure indium sulfide films. The increased band gap correlates with the reduced absorption in the UV-Vis region compared to CdS as a buffer. A reasonably high mobility of around 10 cm2/V·s was obtained for the amorphous films.

Tunability over a wide range of carrier concentrations in the films, ranging from 4×1020 cm-3 to

6×1015 cm-3 was obtained. Such electrical gradation of electron transport layers is critical in combination with photovoltaic absorbers to improve the performance of chalcogenide thin film solar cells.

It is empirically found that the conduction and valence band positions can be tuned by changing the sulfur to oxygen ratio in In2(O,S)3. The valence band offset of all In2(O,S)3 compositions and the pure In2O3 and In2S3 is sufficiently high (> 0.8 eV) to produce a high

162

barrier to holes. But in terms of conduction band positions, only In2(O,S)3 in the composition range between 4 and 33 at. % oxygen provides a small enough spike offset (for enhanced collection of photogenerated electrons from the absorber) in combination with optimal carrier concentration to afford rectification and prevent recombination.The ability to tune the absorber- buffer conduction band offset and carrier concentration of In2(O,S)3 by changing the S:O ratio allows for choice of an In2(O,S)3 composition with the optimal band positions and conductivity to extract the maximum number of photogenerated electrons from an absorber. Such an In2(O,S)3 composition will possibly result in minimal interfacial recombination thus leading to high open circuit voltage and high efficiency CZT(S,Se) solar cell devices.

4.7 References

1. Repins, I.; Contreras, M. A.; Egaas, B.; DeHart, C.; Scharf, J.; Perkins, C. L.; To, B.; Noufi, R., 19·9%-efficient ZnO/CdS/CuInGaSe2 solar cell with 81·2% fill factor. Progress in Photovoltaics: Research and Applications 2008, 16 (3), 235-239.

2. Wang, W.; Winkler, M. T.; Gunawan, O.; Gokmen, T.; Todorov, T. K.; Zhu, Y.; Mitzi, D. B., Device Characteristics of CZTSSe Thin-Film Solar Cells with 12.6% Efficiency. Advanced Energy Materials 2014, 4 (7), 1301465-n/a.

3. Barkhouse, D. A. R.; Haight, R.; Sakai, N.; Hiroi, H.; Sugimoto, H.; Mitzi, D. B., Cd-free buffer layer materials on Cu2ZnSn(SxSe1−x)4: Band alignments with ZnO, ZnS, and In2S3. Applied Physics Letters 2012, 100 (19), 193904.

4. Yan, C.; Liu, F.; Song, N.; Ng, B. K.; Stride, J. A.; Tadich, A.; Hao, X., Band alignments of different buffer layers (CdS, Zn(O,S), and In2S3) on Cu2ZnSnS4. Applied Physics Letters 2014, 104 (17), 173901.

5. Barreau, N.; Bernède, J. C.; Marsillac, S.; Mokrani, A., Study of low temperature elaborated tailored optical band gap β-In2S3−3xO3x thin films. Journal of Crystal Growth 2002, 235 (1–4), 439-449.

6. arreau, N. Marsillac, S. ernède, . C. Assmann, L., Evolution of the band structure of β- In2S3−3xO3x buffer layer with its oxygen content. Journal of Applied Physics 2003, 93 (9), 5456- 5459.

163

7. Maha, M. H. Z.; Bagheri-Mohagheghi, M. M.; Azimi-Juybari, H.; Shokooh-Saremi, M., The structural, thermoelectric and photoconductive properties of sulfur doped In2O3 thin films prepared by spray pyrolysis. Physica Scripta 2012, 86 (5), 055701.

8. Norio, T.; Hideki, M.; Kosuke, C.; Sho, Y.; Masahiro, M.; Shogo, I.; Hajime, S.; Akimasa, Y.; Koji, M.; Shigeru, N., Characterization of electronic structure of oxysulfide buffers and band alignment at buffer/absorber interfaces in Cu(In,Ga)Se2 -based solar cells. Japanese Journal of Applied Physics 2014, 53 (5S1), 05FW09.

9. (ed), W. R. C., CRC Handbook of Chemistry and Physics. (Boca Raton, FL: CRC Press) ( 69th edn), p C-375. 10. Geng, B. Y.; Wang, G. Z.; Jiang, Z.; Xie, T.; Sun, S. H.; Meng, G. W.; Zhang, L. D., Synthesis and optical properties of S-doped ZnO nanowires. Applied Physics Letters 2003, 82 (26), 4791-4793.

11. Robles, R.; Barreau, N.; Vega, A.; Marsillac, S.; Bernède, J. C.; Mokrani, A., Optical properties of large band gap β-In2S3−3xO3x compounds obtained by physical vapour deposition. Optical Materials 2005, 27 (4), 647-653.

12. Platzer-Björkman, C.; Törndahl, T.; Abou-Ras, D.; Malmström, J.; Kessler, J.; Stolt, L., Zn(O,S) buffer layers by atomic layer deposition in Cu(In,Ga)Se2 based thin film solar cells: Band alignment and sulfur gradient. Journal of Applied Physics 2006, 100 (4), 044506.

13. Sun, L.; Haight, R.; Sinsermsuksakul, P.; Bok Kim, S.; Park, H. H.; Gordon, R. G., Band alignment of SnS/Zn(O,S) heterojunctions in SnS thin film solar cells. Applied Physics Letters 2013, 103 (18), 181904.

14. Bayón, R.; Maffiotte, C.; Herrero, J., Chemical bath deposition of indium hydroxy sulphide thin films: process and XPS characterization. Thin Solid Films 1999, 353 (1–2), 100-107.

15. Hariskos, D.; Spiering, S.; Powalla, M., Buffer layers in Cu(In,Ga)Se2 solar cells and modules. Thin Solid Films 2005, 480–481, 99-109.

16. Ho, W.-H.; Hsu, C.-H.; Wei, S.-Y.; Cai, C.-H.; Huang, W.-C.; Lai, C.-H., Sputtered Inx(O,S)y Buffer Layers for Cu(In,Ga)Se2 Thin-Film Solar Cells: Engineering of Band Alignment and Interface Properties. ACS Applied Materials & Interfaces 2017, 9 (20), 17586- 17594.

17. Hejin Park, H.; Heasley, R.; Gordon, R. G., Atomic layer deposition of Zn(O,S) thin films with tunable electrical properties by oxygen annealing. Applied Physics Letters 2013, 102 (13), 132110.

18. Park, H. H.; Jayaraman, A.; Heasley, R.; Yang, C.; Hartle, L.; Mankad, R.; Haight, R.; Mitzi, D. B.; Gunawan, O.; Gordon, R. G., Atomic layer deposition of Al-incorporated Zn(O,S) thin films with tunable electrical properties. Applied Physics Letters 2014, 105 (20), 202101.

164

19. Delft, J. A. v.; Garcia-Alonso, D.; Kessels, W. M. M., Atomic layer deposition for photovoltaics: applications and prospects for solar cell manufacturing. Semiconductor Science and Technology 2012, 27 (7), 074002.

20. Miikkulainen, V.; Leskelä, M.; Ritala, M.; Puurunen, R. L., Crystallinity of inorganic films grown by atomic layer deposition: Overview and general trends. Journal of Applied Physics 2013, 113 (2), 021301.

21. Bakke, J. R.; Pickrahn, K. L.; Brennan, T. P.; Bent, S. F., Nanoengineering and interfacial engineering of photovoltaics by atomic layer deposition. Nanoscale 2011, 3 (9), 3482-3508. 22. Bugot, C.; Schneider, N.; Lincot, D.; Donsanti, F., Synthesis of indium oxi-sulfide films by atomic layer deposition: The essential role of plasma enhancement. Beilstein J Nanotechnol 2013, 4, 750-7.

23. Bugot, C.; Schneider, N.; Bouttemy, M.; Etcheberry, A.; Lincot, D.; Donsanti, F., Study of atomic layer deposition of indium oxy-sulfide films for Cu(In,Ga)Se2 solar cells. Thin Solid Films 2015, 582, 340-344.

24. P. Liu, T. P. C., X.D. Li, Z. Liu, J.I. Wong, Y. Liu and K.C. Leong, Effect of Exposure to Ultraviolet- Activated Oxygen on the Electrical Characteristics of Amorphous Indium Gallium Zinc Oxide Thin Film Transistors. ECS Solid State Letters 2013, 2 (4), Q21-Q24.

25. Purvis, K. L.; Lu, G.; Schwartz, J.; Bernasek, S. L., Surface Characterization and Modification of Indium Tin Oxide in Ultrahigh Vacuum. Journal of the American Chemical Society 2000, 122 (8), 1808-1809.

26. Gunawan, O.; Virgus, Y.; Tai, K. F., A parallel dipole line system. Applied Physics Letters 2015, 106 (6), 062407.

27. Wada, T., Microstructural characterization of high-efficiency Cu(In,Ga)Se2 solar cells. Solar Energy Materials and Solar Cells 1997, 49 (1), 249-260.

28. Bugot, C.; Schneider, N.; Lincot, D.; Donsanti, F., Synthesis of indium oxi-sulfide films by atomic layer deposition: The essential role of plasma enhancement. Beilstein Journal of Nanotechnology 2013, 4, 750-757.

29. Jayakrishnan, R.; Teny Theresa, J.; Kartha, C. S.; Vijayakumar, K. P.; Abe, T.; Kashiwaba, Y., Defect analysis of sprayed β-In2S3 thin films using photoluminescence studies. Semiconductor Science and Technology 2005, 20 (12), 1162.

30. Rehwald, W. Harbeke, G., On the conduction mechanism in single crystal β-indium sulfide In2S3. Journal of Physics and Chemistry of Solids 1965, 26 (8), 1309-1324.

31. Naghavi, N.; Henriquez, R.; Laptev, V.; Lincot, D., Growth studies and characterisation of In2S3 thin films deposited by atomic layer deposition (ALD). Applied Surface Science 2004, 222 (1), 65-73.

165

32. McCarthy, R. F.; Weimer, M. S.; Emery, J. D.; Hock, A. S.; Martinson, A. B. F., Oxygen- Free Atomic Layer Deposition of Indium Sulfide. ACS Applied Materials & Interfaces 2014, 6 (15), 12137-12145.

33. Tauc, J., Absorption edge and internal electric fields in amorphous semiconductors. Materials Research Bulletin 1970, 5 (8), 721-729.

34. Caballero, R.; Kaufmann, C. A.; Cwil, M.; Kelch, C.; Schweigert, D.; Unold, T.; Rusu, M.; Schock, H. W.; Siebentritt, S., The role of the CdS buffer layer in CuGaSe2 -based solar cells. Journal of Physics: Condensed Matter 2007, 19 (35), 356222.

35. Burstein, E., Anomalous Optical Absorption Limit in InSb. Physical Review 1954, 93 (3), 632-633.

36. Matar, S. F.; Villesuzanne, A.; Campet, G.; Portier, J.; Saikali, Y., Étude des structures électroniques de In2O3 pur et dopé avec l’étain. Comptes Rendus de l'Académie des Sciences - Series IIC - Chemistry 2001, 4 (5), 367-373.

37. Robles, R.; Vega, A.; Mokrani, A., Theoretical study of the gap evolution of In2X3 (X=O, S, Se, Te) with lattice compression. Optical Materials 2001, 17 (4), 497-499.

38. Jayakrishnan, R.; Teny Theresa, J.; Kartha, C. S.; Vijayakumar, K. P.; Abe, T.; Kashiwaba, Y., Defect analysis of sprayed β-In 2 S 3 thin films using photoluminescence studies. Semiconductor Science and Technology 2005, 20 (12), 1162.

39. Kraut, E. A.; Grant, R. W.; Waldrop, J. R.; Kowalczyk, S. P., Semiconductor core-level to valence-band maximum binding-energy differences: Precise determination by x-ray photoelectron spectroscopy. Physical Review B 1983, 28 (4), 1965-1977.

40. Zelentsov, K. S.; Gudovskikh, A. S., GaP/Si anisotype heterojunction solar cells. Journal of Physics: Conference Series 2016, 741 (1), 012096.

41. Teimouri, R., Optimization of Cu (In, Ga) Se₂ Based Thin Film Solar Cells: Simulation. World Academy of Science, Engineering and Technology, International Journal of Energy and Power Engineering 2018, 5 (4).

42. Xiao, H.; III, W. A. G., Predicted roles of defects on band offsets and energetics at CIGS (Cu(In,Ga)Se2/CdS) solar cell interfaces and implications for improving performance. The Journal of Chemical Physics 2014, 141 (9), 094701.

43. Minemoto, T.; Matsui, T.; Takakura, H.; Hamakawa, Y.; Negami, T.; Hashimoto, Y.; Uenoyama, T.; Kitagawa, M., Theoretical analysis of the effect of conduction band offset of window/CIS layers on performance of CIS solar cells using device simulation. Solar Energy Materials and Solar Cells 2001, 67 (1), 83-88.

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Chapter 5

Conclusions and future work

This thesis explores a novel non-toxic electron transport material, In2(O,S)3 grown by

ALD for solar cells based on chalcogenide absorbers such as CZT(S,Se). Since CZT(S,Se) contains earth abundant and non-toxic elements like Cu, Zn, Sn and S, it is a scalable, potentially low-cost technology. Atomic layer deposition (ALD) is identified as a preferred technique for growing In2(O,S)3 films due to its ability to produce a well-controlled composition. Thus an optimal S:O ratio in the films can be deposited in different parts of the layer. Consequently, an optimal conduction band offset can be formed at the CZT(S,Se)-buffer interface. This in combination with optimal charge carrier density in In2(O,S)3 should result in efficient minority charge carrier extraction from the absorber with minimal recombination. This should lead to high efficiency of the solar cell devices.

We initially worked on ALD of binary materials, namely In2S3 and In2O3, before studying the deposition of In2(O,S)3 in detail. In chapter 1, we performed ALD of In2S3 using 2 precursors, indium (III) acetylacetonate and indium formamidinate. The latter was demonstrated to give carbon-free, smooth films with relatively high growth rates. The growth rate of In2S3 was seen to decrease with increasing temperature, a possible reason being sulfide bridge formation at higher temperatures. We found that while the charge carrier concentration of In2S3 films is relatively constant within the ALD window, the mobility goes up by an order of magnitude from 160 oC to

200 oC. Despite higher mobilities, one prefers lower temperature of deposition of the buffer (in

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the substrate configuration of the solar cell) to minimize diffusion across the interface. The

o conduction band offset of In2S3 films (deposited at 160 C) with CZT(S,Se) was measured to be

0.29 eV. Despite the promising band offset, solar cells fabricated with CZT(S,Se) and In2S3

o deposited at 160 C showed a Voc which is 80 mV short of CdS based devices. The low Voc is potentially because of many reasons, like low density of charge carriers in In2S3 and existence of defect states in In2S3. Jsc was higher for our In2S3 buffered cells in comparison to best performing

CdS based devices, indicating lower absorption losses due to the indirect nature of the band gap of In2S3. The best-performing CBD In2S3 buffers contain oxygen in them. Thus the likely presence of diffused oxygen in the air exposed In2S3 used in our devices could have boosted the device performance. This motivated us to incorporate oxygen into our sulfide films in situ and deposit controlled compositions of In2(O,S)3. Oxygen also increases the band gap and thus reduces absorption losses. Besides we achieve conduction band and electrical conductivity tuning of the buffer by adjusting the S:O ratio, as discussed in chapter 4.

During the quest to develop a recipe for In2(O,S)3, ALD of binary In2O3 was studied using indium formamidinate and indium acetamidinate precursors over a wide temperature range of 130-350 oC. This work is described in chapter 3. Deionized water was used as a co-reactant.

Indium formamidinate showed a low onset of ALD at 150 oC and a wider ALD window of 150 -

275 oC. Indium acetamidinate had a higher onset of ALD growth at 225 oC with a comparatively narrower ALD window of 225 oC - 300 oC. We show using kinetic studies that a surface saturated with indium formamidinate reacts at a 3 to 4 times faster rate with water compared to a surface saturated with indium acetamidinate. While it takes just 2 seconds for water to react with and saturate a surface with chemisorbed indium formamidinate, 7-8 seconds are needed for the same process using indium acetamidinate. The earlier onset of ALD growth for the

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formamidinate precursor is likely due to faster rate of reaction with water even at lower temperatures. ALD with the formamidinate precursor over a wide temperature range is advantageous because it can be coupled with various other metal oxide and metal chalcogenide processes for obtaining ternary materials. This capability is exploited by developing an ALD process using indium formamidinate, water and hydrogen sulfide to make indium oxysulfide, as described in chapter 4. The higher reactivity of the formamidinate precursor translates into faster growth of electron transport layers and TCOs at lower temperatures. This is beneficial to prevent interfacial diffusion in the solar cells and resultant recombination of electrons. The choice of precursor to grow indium oxysulfide is also facilitated by its shorter growth time, and thus indium formamidinate was chosen as our preferred precursor. In chapter 3 we found that In2O3 grown using indium acetamidinate and indium formamidinate show similar increasing crystallinity trends with temperature in the ALD window because of the similar amidinate scaffold of the 2 precursors. Indium formamidinate affords polycrystalline In2O3 films with high

2 o mobility (48.4 cm /V·s) at reasonably low temperatures of 200 C. In2S3 grown with the same precursor produces films with reasonably high mobility of 17.7 cm2/V·s at the same temperature.

These results indicated that using a combined process for indium oxysulfide with interspersed

In2O3 and In2S3 grown using indium formamidinate, we are likely to attain a reasonably high mobility at 200 oC, even though the mixed films turned out to be amorphous.

Chapter 4 describes the ALD of carbon-free In2(O,S)3 films using indium formamidinate. We successfully demonstrate In2(O,S)3 depositions with capability to adjust the S:O ratio by varying the number of water containing sub-cycles and hydrogen-sulfide containing sub-cycles used in combination with indium formamidinate. The films contain adventitious hydroxyl groups but their content can be controlled to negligible levels by reducing the fraction of water containing

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subcycles to less than 20 %. The hydroxyl groups act as electron donors and we obtain higher

- - conductivity for In2(O,S)3 films with higher OH content. The varying OH content and S:O ratio results in films with a wide range of electron concentrations from 4×1020 cm-3 to 6×1015 cm-3.

This opens the possibility of using an electrically graded buffer with in-situ control over the S:O ratio so that we obtain lower carrier concentrations near the absorber-buffer interface and higher carrier concentrations near the buffer-TCO interface. This helps in minimizing recombination and obtaining a higher open circuit voltage of the solar cell. The In2(O,S)3 films are found be amorphous with a reasonably high electron mobility of around 10 cm2/V.sec independent of S:O ratio. The indium oxysulfide films also exhibit a wider indirect band gap (2.4 eV) compared to their sulfide counterparts (2.1-2.2 eV) and thus are preferred in CZT(S,Se) solar cells to minimize absorption losses. Due to the indirect nature of the band gap they also likely exhibit lower absorption compared to the direct band gap CdS films (~2.4 eV).

Band offset measurements by XPS and calculations in Chapter 4 show that the conduction and valence band positions of In2(O,S)3 can be tuned by changing the sulfur to oxygen ratio. The valence band offset is sufficiently high (> 0.8 eV) for all In2(O,S)3 compositions studied and thus they act as good hole blockers. In2(O,S)3 in the composition range between 4 and 33 % oxygen affords a small spike offset (0.1-0.4 eV) in combination with optimal carrier concentration to afford rectification and reduce recombination. This should ideally result in a high number of minority charge carriers collected from the absorber, a high Voc and hence a high efficiency of the solar cell.

In terms of future work, the compositions of In2(O,S)3 with optimal band positions can be coupled with CZT(S,Se) to check if their junction is rectifying and solar cells can be fabricated and tested. Also, the carrier concentration in In2(O,S)3 can be controlled independent of its band

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position by extrinsic doping with appropriate elements. This helps us to have small spike type offsets of In2(O,S)3 with CZT(S,Se) in combination with a high or a low carrier concentration depending on requirements. Future work in this field could also involve testing In2(O,S)3 in combination with other earth abundant non-toxic chalcogenide absorbers like SnS in solar cells.

Fabrication of solar cells can be done in the superstrate configuration where the absorber material is deposited after the electron transport layer. In such a case, the stability of In2(O,S)3 to high temperature anneals performed after deposition of the absorber will be one of the factors dictating the performance of the solar cells.

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Appendix A

Diagram of System used for atomic layer deposition

Figure A.1 Diagram of system used for atomic layer deposition of In2O3, In2S3 and In2(O,S)3 thin films.

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