Perovskite Solar Cells Via Vapour Deposition Processes

Perovskite Solar Cells via Vapour Deposition Processes

by Qingshan Ma

A THESIS IN FULFILMENT OF THE REQUIREMENTS FOR THE DEGREE OF

Doctor of Philosophy

School of Photovoltaic and Renewable Energy Engineering

Faculty of Engineering

The University of New South Wales

November 2017

Acknowledgment

I would love to thank all the people who provided help and supports during my PhD adventure. I would like to acknowledge the School of Photovoltaic and renewable energy engineering for supporting my PhD studies. Without them, this thesis wouldn’t be possible.

First and foremost, I would like to express my sincere thanks to my supervisors Dr.

Shujuan Huang, Dr. Anita Ho-Baillie and Prof. Martin Green for their guidance and supports in the past 3.5 years.

I would like to thank all my perovskite group members, Dr. Sanghun Woo, Dr. Rui

Sheng, Arman Mahboubi Soufiani, Dr. Jae Yun, Dr. Yajie Jiang, Sheng Chen, Jincheol

Kim, Cho Fai Jonathan Lau, Xiaofan Deng, Adrian Shi, Dr. Meng Zhang, Jianghui Zheng,

Jueming Bing, Yongyoon Cho, Da Seul Lee and Benjamin Wilkinson, for their help in my research and the great time we had together. I also would like to thank Dr. Xiaoming

Wen for the photoluminescence characterization and Dr. Trevor Young for the proof reading of my thesis.

I thank my friend Zewen Zhang, who inspired me and always has been there with me during my time in Australia. I thank my friends Aobo Pu and Dr. Wenkai Cao for making my PhD life so enjoyable. Particularly, I will always remember the countless lunch that I have had with Aobo in the last a few years, which made me not feel so alone.

I would like to say thanks to all my friends in Australia and China!

Special thanks to Zelin Li, who encouraged me to go abroad for research studies.

Last but not least, I would love to thank my family!

This acknowledgement ends with my proposal to Xinwei Li: Marry me! 1 Abstract

Perovskite thin film solar cells have experienced astonishing efficiency improvements from 2.2% to above 22% in recent years. The high efficiency, ease and low embodied energy fabrication and the bandgap tunability make it promising as the next generation of low cost photovoltaic devices and in the application of tandem solar cells with even higher efficiencies.

This thesis focuses on developing scalable fabrication of perovskite solar cells by vapour based processes towards the commercialization and investigating inorganic lead halide perovskites with suitable bandgaps and improved thermal stability for tandem applications.

Vapour based processes which are able to deposit uniform thin films on large substrates for scalable production are developed to fabricate organic lead halide perovskite. Firstly, dual source thermal co-evaporation and sequential thermal evaporation methods are introduced to evaporate methylammonium iodide (MAI) and lead chloride onto the substrate in a vacuum chamber to form the organic lead halide perovskite. Later it is found that the evaporation of the small organic molecular MAI is not so friendly to the vacuum system. Thus a novel vapour-assisted evaporation method is proposed to fabricate MAPbIXCl3-x perovskite in which the deposition of MAI in the vacuum evaporation system is eliminated and carried out in a glass container in the nitrogen glovebox instead.

Inorganic metal halide perovskites have the advantage of better thermal stability compared to the organic counterparts and has a higher bandgap suitable for tandem application, for example, when integrated onto the silicon photovoltaic devices 2 particularly when a vapour fabrication method is employed for the perovskite deposition. This thesis explores the fabrication of inorganic CsPbIBr2 and CsPbI2Br perovskites via the dual source thermal co-evaporation method. The CsPbIBr2 and

CsPbI2Br perovskite with a bandgap of 2.05 eV and 1.9 eV respectively are suitable to work as a top cell in a 3-junction tandem cell on a 1.1 eV silicon bottom cell. The thermal stability of CsPbIBr2 and air stability of CsPbI2Br perovskite thin film are investigated as well in this thesis. The air stability of CsPbI2Br perovskite can be improved by the stoichiometry control, benefiting from the reduced crystallite size.

Table of Contents

Acknowledgment ...... 1 Abstract ...... 2 Table of Contents ...... 4 List of Publications ...... 7 Abbreviations and Symbols ...... 9 Chapter 1 Introduction ...... 12 1.1 Background ...... 12 1.1.1 Fossil Fuel Crisis and Its Problems ...... 12 1.1.2 Solar Energy and Photovoltaics (PV) ...... 12 1.1.3 Thin Film Solar Cells ...... 14 1.1.4 Thin Film Perovskite Solar Cells ...... 17 1.1.4.1 Advantages of Perovskite Solar Cells ...... 19 1.1.4.2 Use of Perovskites for Tandem Solar Cells ...... 19 1.1.4.3 Challenges for Perovskite Solar Cells ...... 20 1.2 Thesis Objective ...... 20 1.3 Thesis Outline ...... 22 References: ...... 24 Chapter 2 Literature Review ...... 26 2.1 Perovskite Solar Cells: Overview ...... 26 2.1.1 History of Perovskites ...... 26 2.1.2 Properties of Perovskite Materials ...... 26 2.1.3 History of Perovskites in Solar Cell Applications ...... 30 2.2 Perovskite Solar Cells Fabricated by Solution Processes ...... 31 2.3 Perovskite Solar Cells Fabricated by Vapour Based Processes ...... 34 2.3.1 Sequential Vapour Deposition Methods ...... 34 2.3.2 Dual Source Thermal Co-Evaporation Methods ...... 36 2.4 Inorganic Caesium Cation Perovskites...... 40 2.4.1 History and Properties ...... 40 2.4.2 Caesium Lead Halide Perovskite Solar Cells ...... 42 References: ...... 47 Chapter 3 Organic Metal Halide Perovskite Solar Cells by Vapour Deposition ...... 52

3.1 Introduction ...... 52

3.2 Vacuum Thermal Evaporation of MAPbIXCl3-x ...... 52 3.2.1 Experimental Section ...... 52 3.2.1.1 Equipment ...... 52 3.2.1.2 Calibration of the Tooling Factor (TF) and Evaporation Rate ...... 54

3.2.1.3 Fabrication and Characterization of MAPbIXCl3-x ...... 55 3.2.2 Results and Discussions...... 57

3.2.2.1 Dual Source Thermal Co-Evaporation of MAI and PbCl2 ...... 57

3.2.2.2 Sequential Evaporation of MAI and PbCl2 ...... 59

3.3 Fabrication of MAPbIXCl3-x by Vapour-Assisted Evaporation Method ...... 63 3.3.1 Methods and Experimental Section ...... 63 3.3.2 Results and Discussions...... 64

3.3.2.1 MAPbIXCl3-x Perovskite Formation Process ...... 64

3.3.2.2 I-V Characteristics of MAPbIXCl3-x Perovskite Solar Cells (Small Area Devices) ... 66

3.3.2.3 I-V Characteristics of MAPbIXCl3-x Perovskite Solar Cells (Large Area Devices) ... 67 3.4 Conclusion ...... 69 References: ...... 71

Chapter 4 Inorganic CsPbIBr2 Perovskite Solar Cells by Dual Source Thermal Co-Evaporation . 72 4.1 Introduction ...... 72 4.2 Experimental Section ...... 73 4.2.1 Device Substrate Preparation ...... 73 4.2.2 Dual Source Thermal Co-Evaporation of the Perovskite Absorber ...... 73 4.2.3 Characterization ...... 74 4.3 Results and Discussions ...... 75

4.3.1 Elemental Analysis of the Deposited CsPbIBr2 Thin Film ...... 75

4.3.2 Crystalline Structure and Morphology Analysis of the CsPbIBr2 Thin Film ...... 77

4.3.3 Optical Properties of the CsPbIBr2 Thin Film ...... 79

4.3.4 Thermal Stability of the CsPbIBr2 Thin Film ...... 83

4.3.5 Electrical Characteristics of the CsPbIBr2 Solar Cells ...... 85

4.3.6 Deposition of CsPbIBr2 on Silicon Substrate ...... 89 4.4 Conclusion ...... 90 References: ...... 91

Chapter 5 Inorganic CsPbI2Br Perovskite Solar Cells by Dual Source Thermal Co-Evaporation . 93 5.1 Introduction ...... 93 5

5.2 Experimental Section ...... 95 5.2.1 Device Substrate Preparation ...... 95 5.2.2 Dual Source Thermal Co-Evaporation of the Perovskite Absorber ...... 95 5.2.3 Characterization ...... 96 5.3 Results and Discussions ...... 97 5.3.1 Elemental and Optical Analysis ...... 97 5.3.2 The Effect of Stoichiometry on the Air Stability ...... 99 5.3.3 The Effect of Stoichiometry on the Film Morphology and Carrier Lifetime ...... 102 5.3.4 Discussion on the Improved Air Stability ...... 107

5.3.5 Electrical Characteristics of the CsPbI2Br Solar Cells with Different Stoichiometry . 108 5.4 Conclusion ...... 113 References: ...... 114 Chapter 6 Conclusion and Future Work ...... 116 6.1 Conclusion ...... 116 6.2 Future Work ...... 119 References: ...... 122

List of Publications

1. Qingshan Ma, Shujuan Huang*, Xiaoming Wen, Martin A. Green, Anita W. Y.

Ho-Baillie, Hole Transport Layer Free Inorganic CsPbIBr2 Perovskite Solar Cell by

Dual Source Thermal Evaporation，Adv. Energy Mater. 2016, 1502202.

2. Qingshan Ma, Shujuan Huang,*, Sheng Chen, Meng Zhang, Cho Fai Jonathan

Lau, Mark N. Lockrey, Hemant K. Mulmudi, Yuchao Shan, Jizhong Yao, Jianghui

Zheng, Xiaofan Deng, Kylie Catchpole, Martin A. Green, Anita W. Y. Ho-Baillie,

The Effect of Stoichiometry on the Stability of Inorganic Cesium Lead Mixed-

Halide Perovskites Solar Cells, accepted by The Journal of Physical Chemistry C.

3. Xiaoming Wen, Rui Sheng, Anita Ho-Baillie, Aleš Benda, Sanghun Woo,

Qingshan Ma, Shujuan Huang, Martin A. Green, Morphology and Carrier

Extraction Study of Organic-inorganic Metal Halide Perovskite by One- and

Two-photon Fluorescence Microscopy, J. Phys. Chem. Lett. 2014, 5, 3849−3853.

4. Cho Fai Jonathan Lau, Xiaofan Deng, Qingshan Ma, Jianghui Zheng, Jae S. Yun,

Martin A. Green, Shujuan Huang, and Anita W. Y. Ho-Baillie, CsPbIBr2

Perovskite Solar Cell by Spray Assisted Deposition, ACS Energy Lett. 2016, 1,

573−577.

5. Meng Zhang, Jae S. Yun, Qingshan Ma, Jianghui Zheng, Cho Fai Jonathan Lau,

Xiaofan Deng, Jincheol Kim, Dohyung Kim, Martin A. Green, Shujuan Huang and

Anita W. Y. Ho-Baillie, High Efficiency Rubidium Incorporated Perovskite Solar

Cells by Gas Quenching, ACS Energy Lett., 2017, 2 (2), pp 438–444.

6. Jianfeng Yang, Xiaoming Wen, Hongze Xia, Rui Sheng, Qingshan Ma, Jincheol

Kim, Patrick Tapping, Takaaki Harada, Tak W. Kee, Fuzhi Huang, Yi-Bing Cheng, 7

Martin Green, Anita Ho-Baillie, Shujuan Huang, Santosh Shrestha, Robert

Patterson and Gavin Conibeer, Acoustic-optical phonon up-conversion and hot-

phonon bottleneck in lead-halide perovskites, Nat Commun. 2017; 8: 14120.

7. Jianghui Zheng, Meng Zhang, Cho Fai Jonathan Lau, Xiaofan Deng, Jincheol Kim,

Qingshan Ma, Chao Chen, Martin A. Green, Shujuan Huang and Anita W. Y. Ho-

Baillie，Efficient Perovskite Solar Cell by Blow-drying，Solar Energy Materials

and Solar Cells 168 (2017) 165–171.

8. Cho Fai Jonathan Lau, Meng Zhang*, Xiaofan Deng, Jianghui Zheng, Jueming

Bing, Qingshan Ma, Jincheol Kim, Long Hu, Martin A. Green, Shujuan Huang,

and Anita Ho-Baillie*, Highly Efficient Strontium Doped Low Temperature

Processed CsPbI2Br Perovskite Solar Cells, submitted to ACS Energy Lett.

9. Hongjun Chen, Meng Zhang, Renheng Bo, Chog Barugkin, Jianghui Zheng,

Qingshan Ma, Shujuan Huang, Anita W. Y Ho-Baillie, Kylie R Catchpole, Antonio

Tricoli, Superior self-powered room-temperature chemical sensing with light-

activated inorganic halides perovskites, submitted to Angewandte Chemie.

Abbreviations and Symbols

ALE atomic layer epitaxy

CBP 4,4’-bis(N-carbazolyl)-1,1’-biphenyl

CdS cadmium sulfide

CdTe cadmium telluride

CIGS copper indium gallium selenide

CIS CulnSe2 c-TiO2 compact titanium dioxide

Cs caesium

CsBr caesium bromide

CsI caesium iodide

CSS close-spaced sublimation

CZTSSe copper zinc tin sulfide selenide

DMSO dimethylsulfoxide

DSSC dye-sensitized solar cell

Ec conduction band

Eg bandgap

EV valence band

Evac vacuum level

EDS Energy-dispersive X-ray spectroscopy

EL electroluminescent

ELO epitaxial lift-off

ETL electron transporting layer 9

+ + FA or HC(NH2)2 formamidinium

FAI or HC(NH2)2I formamidinium iodide

FF fill factor

FTO fluorine doped tin oxide

GaAs gallium arsenide

HI hydroiodic acid

HCVD hybrid chemical vapour deposition

HTL hole transporting layer

HTM hole transporting material

IEP intramolecular exchange process

ITO indium tin oxide

Jsc short circuit current density

LED light emitting diode

LiTFSI lithium bis(trifluoromethane) sulfonimide

+ + MA or CH3NH3 methylammonium

MAI or CH3NH3I methylammonium iodide mp-TiO2 mesoporous titanium dioxide

NREL The National Renewable Energy Laboratory

PV photovoltaics

PCE power conversion efficiency

PFT peeled film technology

Pb lead

PEDOT:PSS poly(3,4-ethylenedioxythiophene):poly(styrene sulfonate)

PL photoluminescence 10

PbI2 lead iodide

PbCl2 lead chloride

PbBr2 lead bromide

PTAA poly[bis(4-phenyl)(2,4,6-trimethylphenyl)amine]

P3HT poly(3-hexylthiophene-2,5-diyl)

SEM scanning electron microscopy

Si silicon

Sn tin

Spiro-OMeTAD 2,2´,7,7´-tetrakis-(N,N-di-p-methoxyphenylamine)9,9´-

spirobifluorene

TCSPC time correlated single photon counting

TF tooling factor

TFT thin-film field-effect transistor

TRPL time-resolve photoluminescence t tolerance factor

UPS ultraviolet photoelectron spectroscopy

μ octahedral factor

Voc open circuit voltage

XPS X-ray photoelectron spectroscopy

XRD X-ray diffraction

4-TBP 4-tert-butylpyridine

Chapter 1 Introduction

1.1 Background

1.1.1 Fossil Fuel Crisis and Its Problems

The rate of fossil fuel formation is much slower than that of its consumption by human being. The depletion times for oil, coal and gas were estimated using various calculation models to be 35 to 40, 106 to 200 and 36 to 70 years, respectively.

Furthermore, the cost of fossil fuel will inevitably increase as resources are depleted

[1]. The pollution caused by burning fossil fuel contributes substantially to global warming, which has a catastrophic effect on the earth’s ecosystem. To accomplish the climate policy of limiting global warming to less than 2 above the pre-industrial average, the use of Earth’s fossil fuel resource has to be℃ limited [2]. There are several alternatives which create electricity without carbon by-products, including hydropower, biomass, geothermal, wind and solar energy[3]. Among these, fuel for solar energy is free and unlimited, solar energy can be widely distributed and is clean when in use.

Also, the cost of solar energy has decreased markedly as the incumbent technologies develop with dramatic reduction in manufacturing cost.

1.1.2 Solar Energy and Photovoltaics (PV)

As a renewable energy source, solar energy can be utilized in two ways: solar thermal and photovoltaics (PV). A solar thermal device harvests the heat from the sunlight directly and stores the heat energy in a transfer material for other uses, while PV absorbs the energy from the sunlight and converts it directly to electricity by solar cells.

Edmond Becquerel demonstrated the first photovoltaic device in 1839 when he generated voltage and current from a solution based structure. The first solid state photovoltaic cell was built in 1883 by coating the semiconductor selenium with a layer of gold. This device converted sunlight into electricity with power conversion efficiency

(PCE) less than 1%. A great improvement in solar cell efficiency was announced in 1954 by scientists at Bell Labs where the first practical silicon solar cell was successfully fabricated. With an efficiency of 6 percent, this was regarded as “the beginning of a new era”. At that time, however, the solar cells were expensive and thus unsuitable for commercialization from the cost point of view. Efforts were made to improve the efficiency and drive down the production cost. Interest from industry and the efforts of many researchers brought down the price of silicon solar cells from around 100 US dollar per watt in the 1970s to below 1 US dollar per watt in the 2010s. During the same period, cell efficiency increased steadily. In 1985, 20% efficient silicon solar cells were reported, which was a milestone in the development of solar cells and made photovoltaics promising for commercialisation and affordable for common use. By

2010, the efficiency of silicon solar cells was increased to more than 25%, which is close to the Shockley-Queisser limit of 30% for single p-n junction silicon solar cells[4].

Because silicon is abundant in the earth’s crust, non-toxic and produces high efficiency cells, solar cells based on silicon remain the main products in the photovoltaic market, today. After years of progress, further reduction of the cost by large scale production and increasing efficiency has become difficult for silicon solar cells. Nevertheless, cheaper absorber materials and cost effective processing technology still provides possibilities for photovoltaic cost reduction, as the high process temperature of silicon solar cell fabrication leads to the high embodied energy. Also, tandem solar cells with 13 multiple junctions can achieve PCE higher than the Shockley-Queisser limit of single junction solar cells with the aim of driving down the system cost[5] due to the less solar panels required for the same power output.

1.1.3 Thin Film Solar Cells

Thin film solar cells use a layer of semiconductor material less than 1 to a few micrometres thick as the light absorber, which is much thinner than crystallized Si based solar cells at around several hundred micrometres. Therefore thin film solar cells have the advantage of minimum material consumption compared to crystallized Si based solar cells and possibly a reduction in the production cost can be achieved.

Research into thin film solar cell technologies commenced decades ago with interest focused mainly on thin film silicon, gallium arsenide (GaAs), copper indium gallium selenide (CIGS), cadmium telluride (CdTe) and more recently copper zinc tin sulfide selenide (CZTSSe) and perovskites.

Thin film silicon solar cells are often classified as 3 types: amorphous Si, microcrystalline Si, and thin-film polycrystalline silicon (polysilicon) solar cells [6-10].

Amorphous Si, with a direct bandgap of 1.7~1.9 eV, is generally fabricated with a hydrogen passivation process, as un-hydrogenated amorphous Si is of no use in electronic devices. Thus the amorphous Si is commonly known as “a-Si:H”. The current laboratory efficiency record of single junction a-Si:H solar cell is 10.2% reported in

2013[11]. Microcrystalline Si is a mixed material of amorphous Si and crystalline Si with different percentages of these two components. As the crystallite size of the crystalline

Si component is generally in the nanometre range, microcrystalline Si is also often referred as nano-crystalline Si. Similar with the a-Si:H, hydrogen is incorporated into

14 the film during the deposition process to passivate the defects and therefore microcrystalline Si is abbreviated to μc-Si:H or nc-Si:H. The efficiency record for microcrystalline Si solar cells, 11.8%, was reported in 2015[12] using films made by a very high frequency plasma-enhanced chemical vapour deposition method. Thin film polycrystalline silicon is the case that when the micro-crystalline Si has no amorphous

Si fraction or very low concentration. Unlike the p-i-n junction of amorphous or microcrystalline Si solar cells, thin film polycrystalline silicon generally has a p-n junction.

Gallium arsenide (GaAs) is an III-V group semiconductor material that has been used in photovoltaic applications since the 1970s. GaAs is a direct bandgap (1.43 eV) material provides strong absorption at wavelengths shorter than 870 nm and high energy conversion efficiency. The cost of these solar cells, however, was very high due to the expensive GaAs material, expensive growth process, and the large amount of it used in bulk type cells. To bring down the high cost, thin film GaAs solar cells were developed by a “peeled film technology (PFT)”[13]. In 1987, based on the peeled film technology, an epitaxial lift-off (ELO) technique was introduced to form large area GaAs thin films[14], which were later broadly applied in the fabrication of GaAs thin film solar cells. The current efficiency record of 28.8% was achieved by Alta Devices in 2012 and remains the highest efficiency obtained by a single junction solar cell without a concentrator [10, 15].

Copper indium gallium selenide (CIGS) thin film solar cells were developed from

CulnSe2 (CIS) thin film solar cells. By varying the In/Ga ratio, the band gap of CIGS can be controlled from 1.04 eV for pure CuInSe2 to approximately 1.7 eV for pure CuGaSe2

[16]. The optimization of different aspects of the solar cell was carried out by 15 institutions throughout the world and rapid progress was realized, resulting in the achievement of 21.7% PCE in 2014[17].

Cadmium telluride (CdTe) is a semiconductor material with a direct band gap of 1.5 eV.

When used in thin film solar cells, cadmium sulfide (CdS) is generally necessary to form a p-n junction together with CdTe. As CdTe is a stable compound, various deposition methods has been developed for CdTe thin film solar cell fabrication including vapour deposition [18], screen printing [19], close-spaced sublimation (CSS)[20], spray technology[21] and atomic layer epitaxy (ALE) [22]. Currently, First Solar holds the record efficiency of 22.1% for CdTe thin film solar cells and is the main manufacturer of this thin film technology[10].

Chalcopyrite CIGS thin film solar cells have been manufactured due to their high efficiency but it is still difficult to compete with Si based solar cells because of the cost of the scarce indium and gallium used in these solar cells. By replacing the indium and gallium in the chalcopyrite CIGS structure with zinc and tin, copper zinc tin sulfide selenide (CZTSSe) having the kesterite structure can be formed. Chalcopyrite and kesterite have very similar crystalline structure [23] and therefore kesterite solar cells are believed to be able to perform as well as CIGS solar cells. The properties of CZTSSe solar cells can be tuned by varying the sulfide/selenide ratio in the composition. When there is no selenide, the material is referred to as CZTS and has a bandgap of 1.45eV.

Similarly, CZTSe refers to the sulfur–free homologue. The first CZTS solar cell was fabricated by a vacuum deposition method in 1997 and had an efficiency of 0.66% and an open circuit voltage of 400mV[24]. The best kesterite CZTSSe thin film solar cell was reported in 2014 and had an efficiency of 12.6% [25].

Currently, thin film solar cell products take no more than 10% of the world’s photovoltaic market share. The main issues regarding the low market share are the complex processing techniques, scarce elements consumption (gallium, indium, selenide and telluride) and the low efficiencies, resulting in the relatively higher production costs compared to silicon solar cells. Also, some of the thin film technologies involve the use of toxic elements such as selenide, arsenide and cadmium, which raise toxicity concerns during the production and disposal processes. If those issues can be addressed thin film solar cell technologies will be promising with huge potentials and applications in the photovoltaic market. The newly-emerged efficient thin film perovskite solar cells, with direct bandgaps variable from 1.4 eV to 2.3 eV, can be processed easily by various methods with low embodied energy. This type of solar cell which has the potential of using non-scarce elements and non-toxic elements becomes a prospective candidate for the next generation of low cost and flexible solar cells and in the application of tandem solar cells.

1.1.4 Thin Film Perovskite Solar Cells

Perovskites are a large group of compounds with the general formula ABX3 that crystallise with the CaTiO3 structure shown in Figure 1.1[26]. Although oxide perovskites such as CaTiO3 have divalent and tetravalent cations, most perovskites of interest for photovoltaic application are halide perovskites wherein A is a monovalent

+ cation that can be either organic (methylammonium (CH3NH3 ), formamidinium

+ + + 2+ 2+ (HC(NH2)2 ) etc.) or inorganic (Cs or Rb ). B is a divalent cation (usually Pb or Sn ) and X is a halide (I-, Br- or Cl-). Perovskites with mixed A cation and/or mixed halide

17 composition are frequently used to manipulate the properties of the material.

Perovskite thin films have emerged as efficient photovoltaic absorbers in recent years.

Figure 1.1. Cubic perovskite crystal structure[26].

Although the perovskite structure was discovered in 1839, the investigation of these materials for electronic application did not start until the 1990s[27, 28]. The first perovskite based solar cell was reported in 2006[29] and published in 2009[30]. The efficiency was only 2.2% at that time. By 2011, the efficiency reached 6.5%[31] which was still quite low compared to other commercialized solar cells. Rapid improvement started from 2012 when an efficiency of 12% was reported and continued through

2016 when the current record was set at 22.1%. Figure 1.2 shows the evolution of certified cell efficiencies including those for perovskite solar cells.

Figure 1.2. Chart of best research solar cell efficiencies from NREL.

1.1.4.1 Advantages of Perovskite Solar Cells

Perovskites solar cells are advantageous because of the diversity of fabrication methods, low embodied energy during fabrication and low cost. The perovskites absorber can be fabricated by solution deposition, vapour deposition via one-step or two-step processes, or combinations of the solution and vapour processes. Although thermal treatments are necessary during the perovskite solar cell fabrication process, the energy consumption is much less than that of the crystalline Si based solar cell fabrication processes. Also, perovskites with direct bandgaps absorb visible and NIR light strongly, which is a key factor for their high efficiency, enabling their use as thin films and consequently reducing the cost.

1.1.4.2 Use of Perovskites for Tandem Solar Cells

Perovskites have bandgaps that can be tuned easily between 1.4eV and 2.3eV by controlling their composition. By stacking wide band gap perovskite cells on top of a silicon solar cell as shown in Figure 1.3, better use of the solar spectrum can be made.

The aim of 2 junction perovskite/Si or 3 junction perovskite/perovskite/Si tandem solar cells is to achieve efficiencies higher than the Shockley-Queisser efficiency limit of single junction solar cells. As a result, the cost in dollar per watt of photovoltaic production could be reduced. The system cost can also be reducted because fewer panels are installed to achieve the same amount of power output.

Figure 1.3. An illustration of 2 junction and 3 junction perovskite/Si tandem solar cells with the

optimal perovskite bandgaps.

1.1.4.3 Challenges for Perovskite Solar Cells

Despite the fast progress of perovskite solar cells, the negative aspects are also significant. State-of-the-art perovskite solar cells employ lead (Pb) as the B cation in the ABX3 perovskite structure, which raises toxicity concerns during the fabrication, application and disposal processes. Another problem is that perovskites degrade rapidly in normal atmospheric conditions due to their sensitivity to moisture and high temperature environments. Also, even after the successful development of perovskite solar cells at laboratory scale, it is challenging to scale up this technology as the best performing devices were fabricated using processes which may be problematic in large scale production. 1.2 Thesis Objective

This thesis will aim to address these aspects of perovskite solar cells:

1. Develop scalable processes for fabricating large area perovskite solar cells and

perovskite tandem solar cells.

2. Develop perovskites that have bandgaps suitable for use in tandem solar cells.

3. Investigate the stability of inorganic perovskite solar cells with a focus on Cs

metal halide perovskites.

Firstly, vapour based deposition methods will be investigated as scalable processes for perovskite solar cell fabrication. Vapour deposition is suitable for fabricating large area thin film devices with good uniformity. Moreover, vapour deposition methods do not require solvents and may be used to deposit materials with poor solubility, increasing the range of material options. Therefore, vapour deposition methods are suitable for the large-scale production and possibly commercialization of perovskite solar cells.

Secondly, thin film perovskites with bandgaps of 2.05eV (CsPbIBr2) and 1.9eV

(CsPbI2Br), respectively, which are suitable for tandem solar cells, were fabricated by the vapour deposition method. Among the numerous perovskite materials, organic lead halide perovskites show the best photovoltaic performance and represent the mainstream of research interests. A major cause of the instability of photovoltaic

+ perovskite materials, however, is that the organic cation A (particularly CH3NH3 ) in the

ABX3 perovskite structure is volatile and may evaporate in time under normal operating conditions (solar panels can experience 80 during their operation), resulting in unrecoverable degradation. This makes per℃ovskite solar cells impractical for long term operation and commercialization so improvement in the perovskites stability is necessary. The inorganic Cs based perovskites developed in this thesis are more thermally stable than perovskites with organic cations. Also, this thesis proposes 21 a method to improve the structural stability of CsPbI2Br perovskite by controlling the stoichiometry. 1.3 Thesis Outline

This Introduction is followed by an in-depth review of the perovskite material and solar cell literature in Chapter 2. The history of perovskite solar cells and the different fabrication methods used to make them are reviewed with particular emphasis on thermal evaporation and vapour deposition processes. Finally, the literature of perovskites with inorganic cations is reviewed.

Chapter 3 investigates the fabrication of organic metal halide perovskites, methylammonium lead mixed iodide chloride (CH3NH3PbIXCl3-x or MAPbIXCl3-x), by different vapour deposition processes, including dual source thermal co-evaporation, sequential evaporation and vapour-assisted evaporation. It was found that methylammonium iodide (CH3NH3I or MAI) has a relatively high vapour pressure and its readily sublimation means it is very hard to control its evaporation rate during a co- evaporation process. A 6.3% MAPbIXCl3-x cell was demonstrated by a sequential evaporation process and an 11% MAPbIXCl3-x cell was made using a vapour-assisted evaporation process.

Chapter 4 explores the fabrication of inorganic CsPbIBr2 perovskites by dual source thermal co-evaporation. A series of experiments investigating the post-annealing process were performed to achieve the best film quality. The optical and electrical properties of the CsPbIBr2 perovskite films are discussed. The thermal stability of

CsPbIBr2 perovskite films is presented with significant improvement compared to the

22 organic-inorganic counterparts. Finally, a hole-transport-material-free CsPbIBr2 perovskite solar cell is fabricated with a PCE of 4.7% for the first time.

Chapter 5 builds on the work of Chapter 4 by increasing the iodine concentration in

CsPbIxBr3-x perovskite to reduce the bandgap from 2.05eV for CsPbIBr2 to 1.9eV for

CsPbI2Br via the co-evaporation of CsBr and PbI2 instead of CsI and PbBr2 thereby increasing absorption of the solar spectrum and the power conversion efficiency of Cs based perovskite solar cells. However, the CsPbI2Br material tends to transfer from the cubic phase (photovoltaic active phase) to an orthorhombic phase which is not active in the photovoltaic application at room temperature and atmospheric conditions. In this work, we found a method to delay that transition by varying the stoichiometry of

CsPbI2Br perovskites by varying the deposition rate ratio of CsBr to PbI2 during thermal co-evaporation. Interestingly, it is found that as CsBr to PbI2 ratio increases, the grain size increases and the air stability of the cubic perovskite phase (photovoltaic active phase) improves as well. Solar cells based on the different stoichiometry were fabricated and the stoichiometrically balanced CsPbI2Br perovskite solar cell has the best photovoltaic performance (PCE of 7.7% for a 0.159cm2 cell and PCE of 6.8% for a

1.2cm2 cell). The 1.2cm2 device is the largest inorganic perovskite solar cell reported to date.

Chapter 6 concludes by summarising the main results and key findings in this work and how they align with the objectives. Aspects that require further improvement, in particular, better voltages, large area cells and better stability, are identified.

References:

1. Shafiee, S. and E. Topal, When will fossil fuel reserves be diminished? Energy Policy, 2009. 37(1): p. 181-189. 2. Jakob, M. and J. Hilaire, Climate science: Unburnable fossil-fuel reserves. Nature, 2015. 517(7533): p. 150-152. 3. Schiermeier, Q., et al., ELECTRICITY WITHOUT CARBON. Nature, 2008. 454(7206): p. 816-823. 4. Shockley, W. and H.J. Queisser, Detailed Balance Limit of Efficiency of p‐n Junction Solar Cells. Journal of Applied Physics, 1961. 32(3): p. 510-519. 5. Green, M.A., Third generation photovoltaics. 2006: Springer. 6. Beaucarne, G., Silicon thin-film solar cells. Advances in OptoElectronics, 2007. 2007. 7. Chopra, K., P. Paulson, and V. Dutta, Thin‐film solar cells: an overview. Progress in Photovoltaics: Research and Applications, 2004. 12(2‐3): p. 69-92. 8. Lee, T.D. and A.U. Ebong, A review of thin film solar cell technologies and challenges. Renewable and Sustainable Energy Reviews, 2016. 9. Schropp, R.E.I., R. Carius, and G. Beaucarne, Amorphous Silicon, Microcrystalline Silicon, and Thin-Film Polycrystalline Silicon Solar Cells. MRS Bulletin, 2011. 32(3): p. 219-224. 10. Green, M.A., et al., Solar cell efficiency tables (version 49). Progress in Photovoltaics: Research and Applications, 2017. 25(1): p. 3-13. 11. Matsui, T., et al. Development of highly stable and efficient amorphous silicon based solar cells. in Proc. 28th European Photovoltaic Solar Energy Conference. 2013. 12. Sai, H., et al., High-efficiency microcrystalline silicon solar cells on honeycomb textured substrates grown with high-rate VHF plasma-enhanced chemical vapor deposition. Japanese Journal of Applied Physics, 2015. 54(8S1): p. 08KB05. 13. Konagai, M., M. Sugimoto, and K. Takahashi, High efficiency GaAs thin film solar cells by peeled film technology. Journal of Crystal Growth, 1978. 45: p. 277-280. 14. Yablonovitch, E., et al., Extreme selectivity in the lift‐off of epitaxial GaAs films. Applied Physics Letters, 1987. 51(26): p. 2222-2224. 15. Kayes, B.M., et al. 27.6% Conversion efficiency, a new record for single-junction solar cells under 1 sun illumination. in 2011 37th IEEE Photovoltaic Specialists Conference. 2011. 16. Devaney, W.E., et al., Structure and properties of high efficiency ZnO/CdZnS/CuInGaSe2 solar cells. IEEE Transactions on Electron Devices, 1990. 37(2): p. 428-433. 17. Jackson, P., et al., Properties of Cu (In, Ga) Se2 solar cells with new record efficiencies up to 21.7%. physica status solidi (RRL)-Rapid Research Letters, 2015. 9(1): p. 28-31. 18. Bonnet, D. and H. Rabenhorst. New results on the development of a thin-film p-CdTe-n- CdS heterojunction solar cell. in Photovoltaic Specialists Conference, 9 th, Silver Spring, Md. 1972. 19. Nobuo, N., et al., Ceramic Thin Film CdTe Solar Cell. Japanese Journal of Applied Physics, 1976. 15(11): p. 2281. 20. Britt, J. and C. Ferekides, Thin‐film CdS/CdTe solar cell with 15.8% efficiency. Applied Physics Letters, 1993. 62(22): p. 2851-2852. 21. Ahrenkiel, R.K., et al. Minority-carrier lifetime of polycrystalline CdTe in CdS/CdTe solar cells. in The Conference Record of the Twenty-Second IEEE Photovoltaic Specialists Conference - 1991. 1991. 22. Skarp, J., et al., Development and Evaluation of Cds/CdTe Thin Film PV Cells, in Tenth E.C. Photovoltaic Solar Energy Conference: Proceedings of the International Conference,

held at Lisbon, Portugal, 8–12 April 1991, A. Luque, et al., Editors. 1991, Springer Netherlands: Dordrecht. p. 567-569. 23. Polizzotti, A., et al., The state and future prospects of kesterite photovoltaics. Energy & Environmental Science, 2013. 6(11): p. 3171-3182. 24. Katagiri, H., et al., Preparation and evaluation of Cu2ZnSnS4 thin films by sulfurization of EB evaporated precursors. Solar Energy Materials and Solar Cells, 1997. 49(1-4): p. 407- 414. 25. Wang, W., et al., Device characteristics of CZTSSe thin‐film solar cells with 12.6% efficiency. Advanced Energy Materials, 2014. 4(7). 26. Green, M.A., A. Ho-Baillie, and H.J. Snaith, The emergence of perovskite solar cells. Nature Photonics, 2014. 8(7): p. 506-514. 27. Mitzi, D., et al., Conducting layered organic-inorganic halides containing (110)-oriented perovskite sheets. Science, 1995. 267(5203): p. 1473. 28. Mitzi, D.B., K. Chondroudis, and C.R. Kagan, Organic-inorganic electronics. IBM journal of research and development, 2001. 45(1): p. 29-45. 29. Kojima, A., et al. Novel photoelectrochemical cell with mesoscopic electrodes sensitized by lead-halide compounds (2). in Meeting Abstracts. 2006. The Electrochemical Society. 30. Kojima, A., et al., Organometal halide perovskites as visible-light sensitizers for photovoltaic cells. Journal of the American Chemical Society, 2009. 131(17): p. 6050- 6051. 31. Im, J.-H., et al., 6.5% efficient perovskite quantum-dot-sensitized solar cell. Nanoscale, 2011. 3(10): p. 4088-4093.

Chapter 2 Literature Review

2.1 Perovskite Solar Cells: Overview

2.1.1 History of Perovskites

The first perovskite, a CaTiO3 mineral, was discovered in the Ural Mountains of Russia by Gustav Rose in 1839 who named it in honour of the Russian mineralogist, Lev

Perovski. The structure of the perovskite mineral was first reported by Helen Dick

Megaw in 1945 [1] who showed that both CaTiO3 and BaTiO3 have tetragonal crystal structure at 20 . At higher temperatures, the tetragonal structure transforms to the closely related ℃cubic structure. These findings were confirmed by Forrester at almost the same time[2].

Since the initial discovery of the perovskite mineral, many synthetic compounds with the general formula ABX3 have been identified. In 1884 Topsöe discussed the organic−inorganic halide compounds that have an organic group as the A cation and halides as the X anion [3]. In 1893 Wells studied the caesium lead halides and identified the CsPbX3 compounds as well as several other double salts for the first time

+ [4]. 85 years later, Weber used the combination of methylammonium ions (CH3NH3 ) with Pb and Sn respectively to form perovskites in 1978 in which a mixed halide system was adopted as the X anion[5, 6]. These findings established the present range of synthetic perovskite compounds relevant for PV application: organic-inorganic lead-tin mixed-halide compounds having the general formula ABX3 and a crystal structure similar to the perovskite mineral, CaTiO3.

2.1.2 Properties of Perovskite Materials 26

As more perovskites became known, researchers soon recognised their broad range of properties that made them particularly suitable for optoelectronic applications.

As a continuation of Weber’s work, Poglitsch investigated the temperature-dependent structures of CH3NH3PbCl3, CH3NH3PbBr3 and CH3NH3PbI3 and their dielectric properties [7]. The methylammonium lead halides were observed to have different colour depending on the different halide ions: CH3NH3PbCl3 is colourless, CH3NH3PbBr3 is orange and CH3NH3PbI3 is black, reflecting their different bandgaps as semiconductors and the bandgap tuning possible by varying the halide composition.

Also, it was found that during the crystallisation of CH3NH3PbI3 by cooling an aqueous

. solution from 100 to room temperature, colourless CH3NH3PbI3 H20 was formed when the temperature℃ fell below 40 , indicating the instability of CH3NH3PbI3 perovskite when exposed to moisture℃ at room temperature. Additionally, all three substances undergo transition with increasing temperature through three crystalline phases: orthorhombic, tetragonal and cubic, in sequence, though the phase transition temperatures are different for the three materials, as shown in Table 2.1. The dielectric constants of each material were discontinuous at the transition from orthorhombic to tetragonal phase but there was no discontinuity at the transition from tetragonal to cubic phase. Onoda-Yamamuro also found this discontinuity of dielectric constant at the orthorhombic to tetragonal phase transition point using wider frequency and temperature range in the CH3NH3PbX3(X=Cl, Br, I) system[8], as shown in Figure 2.1. It is indicated that in these perovskites the tetragonal and cubic phase has similar dielectric property which is different to that of their orthorhombic phase.

The higher dielectric constants from the tetragonal and cubic phase imply their potential use in electronic devices. This is not the case for orthorhombic phase. 27

Table 2.1. Temperature dependent structural data of CH3NH3PbX3 (X = CI, Br, I) [7].

Figure (Text/Chart/Diagram etc.) has been removed due to copyright restrictions

Figure 2.1. Temperature dependence of the real part of the dielectric permittivity of (a) CH3NH3PbCl3,

(b) CH3NH3PbBr3 and (c) CH3NH3PbI3 taken at 1 kHz in heating (open symbols) and cooling (closed

symbols). (d) Temperature dependence of the dielectric constant of the tetragonal phase and cubic

phases of CH3NH3PbX3 (X = Cl, Br, I) taken at 100 kHz[8]. Figure (Text/Chart/Diagram etc.) has been

removed due to copyright restrictions

To understand the possible crystal structure and the formability of ABX3 halide perovskites, Li et al.[9] used Goldschmidt’s tolerance factor t and the octahedral factor

μ where t is the ratio of the distance A-X to the distance B-X and μ is ratio of the ionic radius of B and X atoms as: = ( + )/ 2( + ) and = / , where RA,

푡 푅퐴 푅푋 � 푅퐵 푅푋 � 휇 푅퐵 푅푋 RB and RX are the ionic radii of the corresponding√ ions. This study showed that both t and μ establish necessary conditions for the formation of ABX3 halide perovskites but neither condition is sufficient alone. It was concluded that generally 0.81 < t < 1.11 and

0.44 < μ < 0.90 resulted in halide perovskite (X = F, Cl, Br, I)[10]. When the tolerance factor is in the range of 0.89–1.0, a cubic structure is likely to be formed while lower t values than that give the less symmetric tetragonal or orthorhombic structures.

In 1993, Papavassiliou et al. reported the absorption and luminescence spectra of

CH3NH3Pbl3 perovskites and similar compounds with Br and/or CI instead of I as well as with Sn instead of Pb[11]. They showed that CH3NH3Pbl3 and CH3NH3Snl3 have a photoluminescence (PL) peak at 725nm and 837nm, respectively, which indicates their potential application in optoelectronic devices. Because of the sharp and strong

28 photoluminescence in perovskite semiconductors, Era et al. fabricated an electroluminescent (EL) device using a layered perovskite compound

(C6H5C2H4NH3)2PbI4 (PAPI)[12] with a working mechanism identical to a light emitting diode (LED) device. This perovskite compound has a stable exciton with a large binding energy and its EL spectrum corresponds well with its PL spectrum as shown in Figure

2.2.

Figure 2.2. EL and PL spectra of an organic-inorganic heterostructure device using a spin-coated

phenylammonium lead iodide (PAPI) film at 77 K[12]. Figure (Text/Chart/Diagram etc.) has been

removed due to copyright restrictions

In 1994, Mitzi revealed that layered organic perovskites exhibit a trend in conductivity from a semiconductor to a metal when the perovskite layer number increases[13, 14].

These findings made the perovskite materials interesting to explore as new superconducting materials and more generally to tailor conducting materials by varying the number and type of layers in the perovskite matrix.

Organic-inorganic hybrid (compound with inorganic element and organic group) perovskites have high carrier mobility similar to inorganic semiconductors because the inorganic parts in hybrid perovskites can form an extended framework bound by strong covalent or/and ionic interactions. Moreover, the organic part in hybrid perovskites enables the self-assembly during material formation and therefore simple, low cost and low temperature deposition methods as organic materials. Another advantage from the organic component is that it can be tailored to change the electronic properties of the organic-inorganic hybrid perovskites. Combining these

29 advantages, in 1999, Kagan et al. demonstrated a thin-film field-effect transistor (TFT) using an organic-inorganic hybrid material as the semiconducting channel[15].

2.1.3 History of Perovskites in Solar Cell Applications

The structure of a typical perovskite solar cell is depicted in Figure 2.3[16]. This structure shows obvious trace of dye-sensitized solar cell (DSSC) in which the mesoporous titania works as the metal oxide scaffold and perovskites work as the dye in DSSC.

Figure 2.3. Schematic diagram of a typical perovskite solar cell with the perovskite layer embedded

between electron transporting layer (ETL) and hole transporting layer (HTL)[16]. Figure

(Text/Chart/Diagram etc.) has been removed due to copyright restrictions

The development of perovskite solar cells correspond well with that of DSSC from liquid-electrolyte hole transporting material (HTM) to solid-state HTM. The first perovskite solar cell reported in 2006 by Kojima et al.[17] used an electrolyte solution as the HTM. CH3NH3PbBr3 perovskite was spin-coated onto the nano-porous TiO2 and

2.2% efficiency was achieved. The first solid HTM perovskite solar cells were reported by the same research group 2 years later[18] but the results of only 0.21% efficiency in the CH3NH3PbBr3 solar cell and 0.37% efficiency in the CH3NH3PbI3 solar cell were not very encouraging at the time.

Despite the similarities with DSSC, perovskites work not only as light absorbers in the solar cell but can also transport both electrons and holes, which dyes are not capable of. This was proved by the use of insulating Al2O3 instead of conductive TiO2 as the metal oxide scaffold and also by a planar junction cell structure. In 2012, Lee et al. [19] used an insulating Al2O3 scaffold in a perovskite solar cell that confined the generated

30 carriers to travel in the perovskite material itself until reaching the n-type compact

TiO2 and the HTM. This perovskite solar cell using 2,2´,7,7´-tetrakis-(N,N-di-p- methoxyphenylamine)9,9´-spirobifluorene (spiro-OMeTAD) as the HTM produced a power conversion efficiency as high as 10.9%, confirming the bipolar charge transporting property of perovskite materials. Since then, various deposition methods were developed and the formation of dense perovskite thin films with less defects were investigated, resulting in the boost of PCE to above 22% in a few years[20]. 2.2 Perovskite Solar Cells Fabricated by Solution

Processes

At early stages, perovskite solar cells commonly employ spin-coating as the deposition method for the light-absorbing materials. In the so-called “one-step” spin coating processes, all the precursor ions such as MAI and PbI2 are dissolved in a single solution and applied to the substrate to form MAPbI3 perovskite, at once. By this method, the efficiency of perovskite solar cell was improved from a few percent to above 15% [17,

19, 21-25].

To further improve the photovoltaic performance, different techniques were developed to control the crystallisation and morphology of the perovskite films. Huang et al. reported a gas-assisted solution process for fabricating 17% efficient perovskite solar cells[26]. In this method, a high pressure stream of inert gas was blown on the substrate during the spin-coating process to induce rapid nucleation, as depicted in

Figure 2.4.

Figure 2.4. Schematic procedure for the gas-assisted spin-coating method progressing from left to

right[26]. Figure (Text/Chart/Diagram etc.) has been removed due to copyright restrictions

Later, Zhang et al. eliminated the need to spin the substrate and just used compressed gas to blow-dry[27] the perovskite layer. Based on this process, Zheng et al. fabricated an entirely spin-coating free perovskite solar cell with a PCE of over 17%[28]. In this work, the mesoporous TiO2, methylammonium lead iodide (CH3NH3PbI3 or MAPbI3) perovskite and spiro-MeOTAD layers were all deposited by the blow-drying technique.

Similarly to the gas-assisted techniques that rapidly dry the perovskite precursor solution, Jeon et al. developed a solvent engineering process to fabricate perovskite solar cells in which an anti-solvent (toluene) was drop-cast onto the substrates during the spin-coating process. The anti-solvent displaced the perovskite precursor solvent, resulting in rapid nucleation and drying and an extremely dense and uniform perovskite layer[29], as illustrated in Figure 2.5. Later Li et al. published their work on a vacuum flash–assisted solution process to fabricate efficient perovskite solar cells.

After spin-coating the perovskite precursor solution, the substrates were transferred into a chamber and a vacuum was applied to quickly remove the solvent to facilitate nucleation and perovskite crystallization[30]. This technique produced a 20.5% efficient perovskite solar cell and potentially it can be applied in the up-scaling production of perovskite solar cells.

Figure 2.5. Solvent engineering procedure for preparing the uniform and dense perovskite film[29].

Figure (Text/Chart/Diagram etc.) has been removed due to copyright restrictions

In addition to these one-step deposition techniques, two-step sequential solution deposition processes were also developed. Burschka et al. reported a sequential

32 deposition method and fabricated a 15% efficient perovskite solar cell[31]. In this work,

PbI2 solution was cast onto and infiltrated into a mesoporous titanium dioxide (mp-

TiO2) layer by the spin-coating method. After drying on a hot plate, the PbI2 film was dipped into a solution of methylammonium iodide to form the perovskite, MAPbI3.

Using the same sequential technique, Liu and Kelly demonstrated a planar heterojunction perovskite solar cell with a PCE of 15.7% in which the mesoporous metal oxide layer was eliminated[32]. Yang et al. proposed a modified sequential deposition technique called the intramolecular exchange process (IEP). In this process, a PbI2–dimethylsulfoxide (DMSO) complex was prepared initially and spin-coated onto a mp-TiO2 substrate. Subsequently, a formamidinium iodide (FAI) solution was spin- cast onto the PbI2(DMSO) film to form FAPbI3 perovskite by intramolecular exchange.

By replacing the DMSO molecules with FAI molecules, the unfavourable volume expansion that normally occurs when FAI infiltrates PbI2 was avoided and FAPbI3 based perovskite solar cells with efficiency over 20% were fabricated[33] with a small active area of 0.096 cm2. The efficiencies listed above from solution processes are all based on small area devices measured by a mask with area around 0.1cm2.

At the same time, the research community and industry have been developing scalable solution techniques to fabricate large area perovskite solar cells for commercialization, including hot-casting[34], printing[35], spray-coating[36, 37], blade-coating[38-40], inkjet printing[41-43] and slot-die coating[44-47]. The efficiencies of perovskite solar cells produced by these methods, however, were normally in the range of a few percent to around 12%, substantially less than those produced on small substrates by the solution techniques mentioned above.

To conclude, solution deposition techniques have been at the forefront of improving the photovoltaic performance of perovskite solar cells and attempts have been made to develop scalable solution processes. It is still difficult, however, to fabricate large efficient perovskite solar devices by solution deposition techniques because the uniform deposition of solution onto large substrate is challenging and uniform removal of the solvents after deposition, which is crucial to nucleation and perovskite crystallization, is difficult to control. Therefore, vapour-based deposition methods which involve no or less solvents (with an additional advantage of cost reduction) may be a worthwhile option to consider for the scale-up production of perovskite solar cells. 2.3 Perovskite Solar Cells Fabricated by Vapour

Based Processes

Compared to the solution based fabrication processes, vapour based deposition methods are suitable for depositing uniform thin films on large substrates, involving no solvents and the deposition of insoluble materials.

2.3.1 Sequential Vapour Deposition Methods

Chen et al. [48]reported a 12.1% efficient CH3NH3PbI3 perovskite solar cell with an

2 active area of about 0.11cm via a sequential vapour-assisted method, in which a PbI2 film was firstly deposited by a solution process followed by a CH3NH3I vapour treatment to form the perovskite layer, as shown in Figure 2.6. This method was also used to fabricate methylammonium lead bromide (CH3NH3PbBr3 or MAPbBr3) perovskite solar cells by Sheng et al.[49] and a PCE of 8.7% was achieved.

Figure 2.6. Schematic illustration of perovskite film formation through vapour-assisted process[48].

Figure (Text/Chart/Diagram etc.) has been removed due to copyright restrictions

Leyden et al. proposed a two-step hybrid chemical vapour deposition (HCVD) method to fabricate a perovskite solar cell with a PCE of 11.8% from an active area of around

2 0.1cm [50]. A thin film of PbCl2 was deposited by thermal evaporation in a vacuum system and then the substrate was transferred into a multi-zone tube furnace for the chemical vapour deposition of MAI, as illustrated in Figure 2.7.

Figure 2.7. Diagram of the HCVD furnace and MAI deposition onto metal halide seeded substrates[50].

Figure (Text/Chart/Diagram etc.) has been removed due to copyright restrictions

Replacing the tube furnace for the MAI deposition with a low pressure vacuum test tube, Li et al. reported a low-pressure vapour annealing process to fabricate a 16.8% efficient perovskite solar cell with an active area of 0.062cm2 in which the MAI is deposited onto the PbCl2/PbI2 substrate via vapour state in a test tube with a pressure of ~0.3 Torr [51]. Using a similar method, Luo et al. [52]reported a low-pressure chemical vapour deposition method to carry out the MAI deposition in which the MAI powder and the PbI2 coated substrate were place in a tubular furnace with a pump to achieve a low pressure, as shown in Figure 2.8. By this technique, a 12.7% efficient perovskite solar cell with an active area of 0.12cm2 was fabricated.

Figure 2.8. Schematic diagram of the low-pressure chemical vapour deposition apparatus[52]. Figure

(Text/Chart/Diagram etc.) has been removed due to copyright restrictions

Hu et al. proposed a sequential thermal evaporation method to fabricate the MAPbI3 perovskite layer in a vacuum chamber[53]. In this work, once a base pressure of 5×10-3

Torr was achieved in the chamber, PbI2 and MAI were evaporated in sequence onto the substrates with a fixed source-substrate distance of 20 cm. Subsequently, the as- 35 deposited film was taken out from the chamber and annealed at 100˚C for 30 min to complete the perovskite crystallization. A perovskite solar cell with an efficiency of 5.4% from an active area of 0.09cm2 was produced without an HTM. Similarly, Chen et al. reported a high vacuum sequential thermal evaporation method to fabricate perovskite layers in a vacuum chamber with a base pressure of 1×10-6 Torr [54].

Perovskite solar cells were fabricated by evaporating PbCl2 and MAI to form MAPbIXCl3- x perovskite, fullerene C60 and bathophenanthroline (Bphen) to form the ETL and Ca and Ag to form the metal contact in sequence on a poly(3,4- ethylenedioxythiophene):poly(styrene sulfonate) (PEDOT:PSS) coated indium tin oxide

(ITO) glass, as shown in Figure 2.9. This solar cell produced a PCE as high as 15.4% from an active area of 0.05cm2.

Figure 2.9. Schematic illustration of perovskite solar cells fabricated by sequential thermal

evaporation method of Chen et al. [54] Figure (Text/Chart/Diagram etc.) has been removed due to

2.3.2 Dual Source Thermal Co-Evaporation Methods

A dual source thermal co-evaporation method for perovskite film formation was proposed by Liu et al. in which the MAI and PbCl2 were evaporated simultaneously onto a rotating substrate in a vacuum chamber [55], as shown in Figure 2.10. In comparative studies, it was shown that the dual source evaporation method can produce a more uniform thin film than the solution process. A planar heterojunction

MAPbIXCl3-x perovskite solar cell with a PCE over 15% was demonstrated by this method. Liu et al. concluded that this deposition technique and the perovskite films produced are compatible with first- and second-generation photovoltaic technologies

36 and hence are suitable for the fabrication of efficient tandem solar cells. Later, Gao et al. fabricated perovskite solar cells with this dual source thermal co-evaporation method and did a study of the effects of PbCl2 to CH3NH3I evaporation rate ratio on the structural, optical and electrical properties of the deposited perovskite films[56]. It was found that with additional MAI, the PbCl2 transitioned initially to PbI2 and then to perovskites. After optimization, a 10.5% efficient perovskite solar cell with negligible hysteresis was fabricated. Following this co-evaporation methodology, Kim et al.[57] demonstrated a fully vacuum–processed perovskite solar cell with a PCE of 13.7% wherein all function layers including ETL, HTM and back metal electrode were deposited by thermal evaporation and the perovskite layer was deposited by co- evaporation of PbI2 and MAI.

Figure 2.10. a) Dual source thermal evaporation system for depositing perovskite films. b) Cross-

sectional scanning electron microscopy (SEM) images of completed solar cells constructed from a

vapour deposited perovskite film and a solution-processed perovskite film[55]. Figure

(Text/Chart/Diagram etc.) has been removed due to copyright restrictions

During a dual source evaporation process, however, it is difficult to control the deposition of MAI and the composition of deposited films because the MAI vapour increases the vacuum pressure significantly and its accurate detection by the sensor becomes difficult. Besides, the evaporation of MAI is vapour-induced and non- directional, i.e. not a line-of-sight evaporation. To address this problem, a method to control the MAI evaporation rate through the MAI vapour partial pressure inside the vacuum chamber was proposed by Ono et al.[58]. Unlike the co-evaporation chamber set-up in Figure 2.10 where the two rate sensors are located close to the sources, in

Ono’s work the rate sensors are positioned near the substrate with the MAI sensor facing upwards, away from the MAI source and the PbCl2 sensor facing downwards, towards the PbCl2 source, as shown in Figure 2.11. When MAI evaporates, the chamber fills with MAI vapour. The sensor facing up detects MAI only while the sensor facing down detects the total rate of PbCl2 and MAI deposition. The rate detected by the sensor facing upwards is not the accurate rate but is an indication of the MAI vapour pressure, which reflects the possible MAI deposition rate. Combining the rates detected by the two sensors and the characterization of the deposited films, a co- evaporation system with better deposition rate control was established. Using this method, a MAPbIXCl3-x perovskite solar cell with a PCE of 9.9% was fabricated.

Figure 2.11. Side view of the co-evaporation system: (1) main vacuum chamber; (2) pumping system;

(3) substrate holder stage; (4) substrate; (5) substrate shutter; (6) rate sensor facing downwards; (7)

rate sensor facing upwards; (8) MAI evaporation source; (9) widely opened dish-shaped crucible for

the evaporation of lead halide compounds; (10) spiral-shaped tungsten wire; (11) electric feedthroughs; (12) lead halide shutter; (13) pressure gauge[58]. Figure (Text/Chart/Diagram etc.) has

been removed due to copyright restrictions

As well as detecting MAI deposition by a rate sensor, the deposition rate of MAI can be inferred from the pressure in the vacuum chamber during the thermal co-evaporation process because MAI evaporation increases the chamber pressure significantly while the lead halide compounds’ evaporation do not, as reported by Kim et al. [59]. In their work, the deposition rate of PbI2 was fixed and the vacuum chamber pressure, controlled by adjusting the MAI source power, was the MAI deposition parameter. X- ray diffraction (XRD) patterns of films deposited at different chamber working pressures were obtained to assess the MAPbI3 film quality, as shown in Figure 2.12. 38

-5 After optimizing the deposition pressure to 5.1×10 Torr, a MAPbI3 perovskite solar cell with a PCE of 14.5% was successfully demonstrated. Based on this technique, Zhao et al.[60] reported a MAPbI3 perovskite solar cell with an efficiency of 15.7% by employing vacuum-deposited fullerenes C60 and C70 as the ETL, in which the MAPbI3 perovskite layer was deposited by co-evaporation of PbI2 and MAI at a chamber pressure of 5×10-5 Torr.

Figure 2.12. a) Schematic of the vacuum thermal co-evaporation system for the perovskite layer. b)

XRD patterns of the perovskite films under different pressures on the ITO substrate and a magnified

image of the XRD patterns in the red square region[59]. Figure (Text/Chart/Diagram etc.) has been

removed due to copyright restrictions

In conclusion, various vapour based deposition methods suitable for large scale production have been developed to fabricate organic metal halide perovskite solar cells. These include sequential vapour deposition and dual source thermal co- evaporation methods, in which dual source thermal co-evaporation methods have attracted particular interest because of the simple and efficient one-step deposition process. Efficiencies over 15% have been achieved by these vapour deposition techniques while the large scale solution deposition processes generally produced perovskite solar cells with efficiencies below 12%. Accurate control of the MAI flux during the vapour deposition, however, is challenging due to the volatility of the small organic MAI molecule, requiring additional engineering processes to indicate the deposition rate. If inorganic cations such as Cs+ and Rb+ are employed to replace the

+ organic cation CH3NH3 , more precise control of the deposition rate could be realised in the vapour based fabrication processes for perovskite solar cells.

2.4 Inorganic Caesium Cation Perovskites

Due to the rapid progress of perovskite solar cells, the attention of the research community has been drawn to the issues of stability and the commercialization of this technology. Although organic cations in perovskites provide suitable bandgap for single junction solar cell application and high power conversion efficiency, volatility of the organic component also contributes to the instability of the corresponding perovskite compounds. Therefore, inorganic cations such as Cs+ and Rb+ emerge as possible substitutes for the organic cations MA+ and FA+ in perovskites. According to the tolerance factor for indicating the formability of ABX3 halide perovskites discussed in section 2.1.2, however, Rb cannot form a perovskite with lead and iodine as the tolerance factor is less than 0.81[61]. Therefore, inorganic caesium lead halide perovskites with tolerance factors above 0.81 are more attractive in solar cell applications.

2.4.1 History and Properties

The first Cs perovskite was described in 1893 in the compound CsPbX3 (X=Cl, Br or I)[4], which was the first description of lead halide perovskites as well. In 1958[62], Moller investigated the crystal structure of CsPbX3 (X=Cl, Br or I) compounds. It was found that

CsPbCl3 and CsPbBr3 have a tetragonal perovskite structure at room temperature and undergo a transition to the cubic perovskite phase at 47 and 130 respectively, while CsPbI3 has an orthorhombic phase at room temperature℃ and ℃transitions to the black perovskite structure at 305-308 . Moller also noted that the CsPbX3 crystals are intensely coloured; CsPbCl3 is pale yellow,℃ CsPbBr3 is orange and CsPbI3 is black, implying that if this kind of material was used in semiconductor applications the 40 bandgap could be easily tuned by varying the halide composition. Another interesting finding of this work was that the coloured CsPbX3 crystals are photoconductive, an indication of their potential application in photo-electronic devices such as solar cells.

Later in 1974 [63], Hirotsu et al. confirmed the phase transition in CsPbBr3 from cubic perovskite phase to tetragonal phase with cooling through 130 and also identified a transition to orthorhombic phase at 88 . It can be concluded tha℃ t CsPbX3 crystals show a trend of decreasing symmetry from℃ Cl to Br and then to I. This is particularly true for CsPbI3 which is non-perovskite orthorhombic phase at room temperature [64] and has no practical value in semiconductor applications.

Therefore, researchers tend to investigate the more structurally stable Br-containing

CsPbX3 perovskites. In 2013[65], Stoumpos et al. synthesized a single crystal of CsPbBr3 perovskite semiconductor with a strong PL emission, as shown in Figure 2.13. They found that the CsPbBr3 perovskite has the properties of high attenuation, high resistivity and significant photoconductivity response, making it suitable in the detector applications.

Figure 2.13. a) Photograph of CsPbBr3 single crystals. b) A photograph of the strong PL emission of a

CsPbBr3 specimen at 46 K [65]. Figure (Text/Chart/Diagram etc.) has been removed due to copyright

restrictions

Protesescu et al. synthesized nanocrystals of CsPbX3 (X=Cl, Br or I) perovskites as well as mixed halide CsPbX3 perovskites such as CsPb(I/Br)3 and CsPb(Br/Cl)3[66]. The properties of the CsPbX3 perovskite nanoparticles such as their crystal colour, XRD patterns, absorption spectra and PL spectra trend consistently as the halide composition was varied from Cl to I, as shown in Figure 2.14, indicating that the

41 bandgaps of CsPbX3 (X=Cl, Br or I) perovskites can be readily tuned for various optoelectronic applications, like solar cells, by engineering the halide composition.

Figure 2.14. a) Nanocrystal perovskites solutions under UV lamp (λ = 365 nm). b) XRD patterns for typical ternary and mixed-halide nanocrystal perovskites. c) Typical optical absorption and PL spectra

of the nanocrystal perovskites. [66] Figure (Text/Chart/Diagram etc.) has been removed due to

2.4.2 Caesium Lead Halide Perovskite Solar Cells

In 2015, Kulbak et al.[67] fabricated the first solar cells using CsPbBr3 perovskite thin films as the light absorbing material. With a direct bandgap of 2.32 eV, these cells produced efficiencies around 5%. In this work, a two-step solution process was employed to fabricate the CsPbBr3 as the authors were unable to find a solvent to dissolve both the CsBr and the PbBr2, indicating that it may be problematic to fabricate

CsPbX3 thin films by one-step solution processes due to the poor solubility of the CsX and PbX2 (X=Cl, Br or I) materials. They also measured the valence band (EV) of the

CsPbBr3 perovskite on mp-TiO2 to be at 5.97 eV versus the vacuum level (Evac) by ultraviolet photoelectron spectroscopy (UPS), which is similar to the 5.9 eV for the EV of MAPbBr3 they had measured previously. Based on the EV of the CsPbBr3 perovskite, different HTMs, i.e., spiro-OMeTAD, poly[bis(4-phenyl)(2,4,6-trimethylphenyl)amine]

(PTAA), and 4,4’-bis(N-carbazolyl)-1,1’-biphenyl (CBP) were used in demonstrated devices, their photovoltaic performances as well as those for HTM-free devices were compared. Since then, many researchers have worked on solar cells employing caesium lead halide thin films as the light absorber, particularly Br-containing CsPbX3 perovskites.

Later, also using the two-step solution process, an entirely inorganic perovskite solar cell with the structure of glass/FTO/ c-TiO2/mp-TiO2/ CsPbBr3/carbon was demonstrated by Liang etc. without the use of expensive HTMs and metal electrode, as shown in Figure 2.15, producing a PCE of 6.7%[68]. The thermal stability of CsPbBr3 solar cells was compared with that of similarly structured MAPbI3 solar cells. The results showed that cells based on CsPbBr3 were much more stable at 100 and high- humidity ambient than those made from MAPbI3. ℃

Figure 2.15. a) Schematic cross-sectional view of the CsPbBr3 all-inorganic perovskite solar cell. b)

Energy level diagram of the device. c) Normalized PCEs of CsPbBr3/carbon-based all-inorganic solar

cells and MAPbI3/carbon-based hybrid solar cells as a function of time heated at high temperature

(100 °C) in a high-humidity ambient environment (90−95% relative humidity (RH), 25 °C) without

encapsulation. [68] Figure (Text/Chart/Diagram etc.) has been removed due to copyright restrictions

As mentioned previously, the bandgap of caesium lead halide perovskites can be reduced and the light absorption spectrum can be widened by incorporating iodine into the halide composition. Therefore, caesium lead halide solar cells with higher iodide composition such as CsPbIBr2, CsPbI2Br and CsPbI3 were explored, even though

CsPbI3 does not form a stable perovskite at room temperature. Lau et al. demonstrated a mix halide CsPbIBr2 solar cell with a stabilized efficiency of 6.3% using the spray-assisted deposition technique[69], in which the PbBr2 film was introduced onto the substrate by spin-coating and then the CsI was sprayed on under ambient conditions to form the perovskite. The deposited CsPbIBr2 film has a reduced optical bandgap of around 2.05 eV compared to CsPbBr3 (2.3 eV) and the film was proved to be thermally stable by heating at 300 for 1.5 hours.

℃ 43

Sutton et al. reported producing mixed halide CsPbI2Br perovskite solar cells by solution processes with a stabilized efficiency of 5.6%[70]. They found that the two- step solution method can be used to form the full series of caesium lead halide perovskite thin ﬁlms from CsPbBr3 to CsPbI3, while the one-step solution process can only create uniform thin ﬁlms of CsPb(IxBr1− x)3 with iodide fraction x between 0.67 and

1. The absorption onsets and PL peak positions varied linearly with x between 0.67 and

1, as shown in Figure 2.16a. The thermal stability of CsPbI2Br and MAPbI2Br was assessed by comparing the XRD patterns of the films before and after heating at 85°C for 270 min. The results in Figure 2.16b show that CsPbI2Br is more thermally stable than MAPbI2Br. Almost at the same time, Beal et al. reported the synthesis of solution processed mixed halide CsPb(IxBr1− x)3 perovskites. A solar cell based on CsPbI2Br perovskite with a bandgap of 1.9 eV was fabricated, producing a stabilized PCE of 6.5%

[71]. It was predicted that Cs-based mixed-halide perovskites can be optimized to serve as absorbers in tandem top-cells and other optoelectronic applications such as

LEDs.

Figure 2.16. a) Absorption onsets (black squares), determined by Tauc plots of absorbance spectra

taken with an integrating sphere, and PL peak positions (red open squares); solid line is the linear ﬁt

to the absorbance. Error bars indicate the range of values obtained. b) XRD spectra before and after

heating CsPbI2Br and MAPbI2Br films at 85 °C in 20–25% RH for270 min. The PbI2−xBrx peak position is

indicated by a diamond. [70] Figure (Text/Chart/Diagram etc.) has been removed due to copyright

restrictions

Later, Niezgoda et al. reported a CsPbI2Br perovskite solar cell with a PCE of 10.3% produced by solution processes[72]. The improved efficiency was ascribed to improved charge collection resulting from light-induced dealloying of CsPbI2Br. Nam et al. 44 studied the crystal formation process of CsPbI2Br thin films at different annealing temperatures[73]. It was found that the film has the best, pin-hole-free surface morphology when annealed at 280 °C, as shown in the SEM images in Figure 2.17. By optimizing the fabrication processes, a 10.7% efficient solar cell was produced.

Figure 2.17. SEM images of CsPbI2Br films annealed at temperatures from 100 to 350 °C for 2 min.

Scale bar: 1 μm. [73] Figure (Text/Chart/Diagram etc.) has been removed due to copyright restrictions

A CsPbI3 perovskite solar cell was produced by Eperon et al. using solution processes in which hydroiodic acid (HI) was added to the precursor solution to stabilize the cubic phase of the annealed film with a bandgap of 1.73 eV in ambient conditions for hours[74]. The solar cell, however, only had a PCE of 2.9%. Later, the efficiency of

CsPbI3 perovskite solar cells was increased to 4.68% by Ripolles et al. [75] through the addition of an electron blocking MoO3 layer between the poly(3-hexylthiophene-2,5- diyl) (P3HT) HTM and the Au electrode, as shown in Figure 2.18.

Figure 2.18. Schematic architecture of the caesium perovskite solar cell as glass/FTO/blocking TiO2(bl-

TiO2)/mp-TiO2/CsPbI3/P3HT/MoO3/Au. [75] Figure (Text/Chart/Diagram etc.) has been removed due

to copyright restrictions

As the solution processes limit the control of the film thickness, uniformity, purity and crystallinity of caesium lead halide perovskite thin films and the solution concentration is limited by the poor solubility of the inorganic materials, vapour based processes may be preferable for depositing high quality CsPbX3 thin films. Frolova et al. reported fabricating CsPbI3 perovskite solar cells via thermal co-evaporation of CsI and PbI2 and planar junction solar cells with efficiency up to 10.5% were demonstrated[76].

Later, Chen et al. fabricated CsPbX3 perovskite solar cells in which the ETL, perovskite layer and HTM were all deposited by thermal evaporation. For comparison, CsPbI2Br and CsPbI3 perovskites were employed as absorber materials. CsPbI2Br and CsPbI3 were deposited by co-evaporating CsBr with PbI2 and CsI with PbI2, respectively[77].

Efficiencies of 11.8% and 9.4% were achieved for the CsPbI2Br and CsPbI3 based solar cell respectively, illustrating that CsPbI2Br cells have better photovoltaic performance, particularly higher open circuit voltage (Voc) and fill factor (FF), than CsPbI3 based cells as shown in Table 2.2, although the CsPbI3 perovskite has a more favourable band gap and wider light absorption spectrum.

Table 2.2. Performance of CsPbI3 and CsPbI2Br solar cells[77].

Voc Jsc PCE Device FF [V] [mA cm-2] [%]

CsPbI3 0.97 17.4 0.56 9.4

CsPbI2Br 1.15 15.2 0.67 11.8

In conclusion, organic lead halide perovskites and inorganic caesium lead halide perovskites are being developed as semiconductor materials for photovoltaic applications. Their bandgaps are tuneable by engineering the halide ion composition.

The organic component of organic lead halide perovskites can be evaporated easily resulting in thermal instability, while the inorganic caesium lead halide perovskites are more thermally stable as indicated by their much higher annealing temperature, 250 to 350°C compared to around 100 °C for organic perovskites. Various fabrication processes have been developed for perovskite thin films. Among these processes, vapour based deposition methods are advantageous for fabricating uniform thin films 46 on large area substrates and are suitable for large scale production and commercialization while solution processes are not ideal for large scale production of highly efficient perovskite solar cells. However, it has been found that the evaporation of MAI is not so friendly to the evaporation vacuum system. Therefore, this thesis proposes some modified methods based on evaporation to fabricate organic-inorganic perovskite solar cells. In addition, dual source thermal co-evaporation for the deposition of inorganic caesium lead halide perovskites which have poor precursor solubility and therefore are not amenable to solution fabrication processes is proposed as well.

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Chapter 3 Organic Metal Halide Perovskite

Solar Cells by Vapour Deposition

3.1 Introduction

This chapter focuses on the fabrication of perovskite thin films with organic cations, specifically MAPbIXCl3-x, by vapour-based methods. Firstly, an attempt was made to establish a baseline process for MAPbIXCl3-x by dual source thermal co-evaporation[1], though some problems with this technique were recognized very soon. Then, a modified sequential thermal evaporation method[2, 3] was used for the deposition of this material and a solar cell with a PCE of 6.3% by reverse scan (from Voc to Jsc (short circuit current density)) and 2.44% by forward scan (from Jsc to Voc) was fabricated. A serious problem was detected that MAI vapour is damaging to the evaporation system.

Finally, a vapour-assisted evaporation (atmospheric pressure MAI vapour deposition process) was developed to fabricate MAPbIXCl3-x thin films in which the detrimental evaporation of MAI in the vacuum system was eliminated. Solar cells using this method had an efficiency of 11.0% by reverse scan which shows the potential of this technique.

3.2 Vacuum Thermal Evaporation of MAPbIXCl3-x

3.2.1 Experimental Section

3.2.1.1 Equipment

The vacuum thermal evaporation process was carried out in a high vacuum chamber

(Kurt J. Lesker Mini Spectros) integrated in a nitrogen filled glove box. A schematic

52 drawing of the dual source thermal evaporation system is shown in Figure 3.1. Two evaporation sources were positioned at the bottom of the evaporator chamber. Each source consisted of a crucible to hold the evaporation material, a heating element around the crucible, a temperature sensor below the crucible, heat insulating layers outside the heating element and a shutter. A deposition rate sensor with a gold coated

6 MHZ crystal oscillator was placed closely above each evaporation source and the rate sensors are cooled by external cooling water. A 20 cm diameter substrate holder at the top of the evaporator chamber faced down, towards the sources. The substrate holder was rotated during the deposition at adjustable rotation speeds and its temperature was controlled in the range of -20 to +85 by flowing coolant from an external system. A main shutter close to the substrate℃ holder controlled the beginning and end of the deposition.

A turbo pump and a roughing pump were connected to the evaporation chamber in series to achieve the vacuum condition for the thermal evaporation. Once the required vacuum pressure was reached, the shutter for the evaporation source was opened and the source was heated to vaporize the material inside the crucible. The heating power was adjusted to control the evaporation rate which was detected by the sensor above.

Once the desired evaporation rate was established, the main shutter for the substrate was opened to start the deposition. After depositing the required thickness, the main shutter was closed to stop the deposition process.

Figure 3.1. An illustration of the evaporation chamber set up for dual source thermal evaporation

processes

3.2.1.2 Calibration of the Tooling Factor (TF) and Evaporation Rate

As can be seen from Figure 3.1, the rate sensor detects the evaporation rate near the heating source whereas the real deposition rate at the substrate holder is different.

Therefore a tooling factor (TF), the ratio of the deposition rate at the substrate holder to the deposition rate detected by the sensor, is necessary to indicate the real deposition rate for different materials. To calculate the TF, firstly, the TF was set at 1 in the evaporation programme, and then a deposition of some material was carried out.

The thickness deposited on the substrate was measured by a surface profilometer. By dividing the thickness measured on the substrate by the thickness measured by the sensor the TF was estimated. Several more TF calibrations were repeated to confirm the results. Then, the real deposition rate/thickness on the substrate can be estimated by the calculated TF and sensor reading for this kind of material. Note that different materials have different TFs and even with the same material when the deposition

54 conditions vary significantly, the TF may be different. Also, the TF only provides an estimation of the deposition rate and 10% to 20% differences from the real deposition rate may apply depending on the materials deposited. The TF for MAI and PbCl2 evaporation at a deposition rate range of 0.1 to 0.3 Å and 0.1 to 0.5 Å were calibrated at 0.13 and 0.31, respectively.

3.2.1.3 Fabrication and Characterization of MAPbIXCl3-x

For dual source thermal co-evaporation of MAPbIXCl3-x, MAI and PbCl2 were loaded in source 1 and 2 respectively and clean glass substrates were fixed onto the substrate holder, facing towards the sources. After the pressure of the evaporator chamber was

−6 pumped down to 5x10 mbar, MAI and PbCl2 were heated to the set temperature of

70 °C and 240 °C, respectively. Once these temperatures were reached, the shutter for each source was opened simultaneously and the deposition rate of each material was increased to the level required by fine adjustment of the source heating power. After the deposition rates stabilized at the required values, the main shutter was opened to commence deposition on the substrates.

For sequential evaporation of MAPbIXCl3-x, MAI and PbCl2 were also loaded in source 1 and source 2, respectively and the clean substrates were fixed onto the substrate holder facing towards the sources. When a chamber pressure of 5x10−6 mbar was achieved, PbCl2 was heated to the set temperature at 240 °C before opening shutter 2.

After the desired deposition rate stabilized, the main shutter was opened to start the

PbCl2 deposition. The main shutter and shutter 2 were closed when the PbCl2 deposition was finished and source 2 was cooled down to room temperature. Then the

MAI source was heated up and shutter 1 was opened after the source temperature of

70 °C was reached. Again, the MAI evaporation rate was adjusted and stabilized to a 55 set value before the main shutter was opened to start the MAI deposition. When the

MAI deposition was completed the shutters were closed and the source cooled down to room temperature before venting the chamber. After removing the substrates from the evaporation chamber, they were annealed at 110 °C on a hotplate in the glovebox.

MAI was synthesized by reacting 24 mL of methylamine (33 wt% in ethanol, Sigma-

Aldrich), 10 mL of hydroiodic acid (57 wt% in water, Sigma-Aldrich), and 100 mL ethanol in a 250 mL round-bottom flask at 0 °C for 2 hours with stirring. After the reaction, the precipitate was recovered by rotary evaporation at 60 °C and then dissolved in ethanol followed by sedimentation in diethyl ether by stirring, which was repeated until the white MAI powder appeared. The final product was collected and dried at 60 °C in an oven and dehydrated in a vacuum chamber. PbCl2 (puratronic,

99.999%) was purchased from Sigma-Aldrich and used without any further purification.

To fabricate perovskite solar cells, fluorine-doped tin oxide (FTO)-coated glass substrates (Pilkington, 15 Ω/sq) were patterned with 2 M HCl and zinc powder and then cleaned in 2% Hellmanex detergent, acetone and isopropanol sequentially for 10 minutes each in an ultrasonic bath. This was followed by a UV-ozone (UVO) treatment for 15 min. The n-type compact TiO2 (ETL) layer was deposited by spin-coating a mildly acidic solution of titanium isopropoxide in ethanol onto the substrates at 2500 rpm for

60 s followed by annealing at 500 °C for 30 min. Then the MAPbIXCl3-x perovskites were deposited onto the compact TiO2 by a thermal evaporation method described above. A

HTM was deposited by spin-coating a spiro-MeOTAD solution onto the substrates at

2000 rpm for 60 s. The HTM solution was prepared by dissolving 72.3 mg spiro-

MeOTAD, 28.8 mL of 4-tert-butylpyridine (4-TBP) and 17.5 mL of a stock solution of

520 mg/mL lithium bis(trifluoromethane) sulfonimide (LiTFSI) in acetonitrile in 1 mL of 56 chlorobenzene. The solar cells were completed by depositing 100 nm of gold by thermal evaporation through a shadow mask that defined the active area. The device structure is shown in Figure 3.2.

Figure 3.2 Cross sectional schematic of the solar cell structure.

XRD patterns were measured using a PANalytical Xpert Materials Research diffractometer system with a Cu Kα radiation source (λ=0.1541 nm) at 45 kV and 40mA.

The current density–voltage (J–V) measurements were performed using an IV5 solar cell I–V testing system from PV measurements, Inc. (using a Keithley 2400 source meter) under illumination power of 100 mW/cm2 from an AM1.5G solar simulator

(Oriel model 94023A) with an 0.159 cm2 aperture and a scan rate of 1.2 V/s.

3.2.2 Results and Discussions

3.2.2.1 Dual Source Thermal Co-Evaporation of MAI and PbCl2

The dual source thermal co-evaporation of MAI and PbCl2 to form MAPbIXCl3-x was carried out as described in the literature[1]. First, the deposition rate ratio of MAI to

PbCl2 required to form stoichiometric MAPbIXCl3-x was calculated. Given that the deposition rate is proportional to the thickness, volume, mass and molar quantity of the deposited material, the required 1:1 molar ratio, the molar masses of MAI and 57

3 PbCl2 (158.9 and 278.1 g/mol) and the densities of MAI and PbCl2 (1.1 and 5.85 g/cm ) the required deposition rate ratio of MAI to PbCl2 is

× / = = = = 3.09 푀퐴퐼× 푀퐴퐼 / 푀퐴퐼 푅푀퐴퐼 푇푀퐴퐼 푉푀퐴퐼 푀 푛 휌 � 푃푏퐶푙2 � 푃푏퐶푙2 � 푃푏퐶푙2 2 2 2 푅 푇 푉 푀푃푏퐶푙 푛푃푏퐶푙 휌푃푏퐶푙 where R, T, V, ρ, M and n represent the deposition rate, thickness, volume, density, molar mass and amount of substance in mol for the corresponding materials, respectively. Therefore, the deposition rates of MAI and PbCl2 were set to 0.27 Å/s and

0.09 Å/s respectively.

Controlling the deposition rates during the co-evaporation was difficult, however, and the MAPbIXCl3-x perovskite thin film could not be produced. It was noted that the evaporation of MAI increased the pressure in the chamber from 10-6 mbar to >10-4 mbar and also affected the rate reading on the PbCl2 sensor. An experiment was performed to observe this more closely. During a 3-hour MAI evaporation from source

1, the PbCl2 source remained off and shutter 2 remained closed at all times. The deposition rates on sensor 1 for MAI and sensor 2 for PbCl2 were recorded, as shown in Figure 3.3. It can be seen that the evaporation rate of MAI initially ramped up to 0.3

Å/s and then stabilized at around 0.27 Å/s for about 2 hours. Even though PbCl2 was not evaporated, sensor 2 gave a high reading during the whole process. The rate increased initially to 1 Å/s and then deceased slowly until the end of the MAI evaporation (the rate reading on sensor 2 is not accurate as the TF used is for PbCl2 deposition instead of MAI). It is concluded that the evaporation of MAI is not line-of- sight and that it affects the rate reading for the evaporation of PbCl2 during the co- evaporation process. To address this problem, a modified sequential evaporation of

MAI and PbCl2 was proposed for the vacuum thermal evaporation of MAPbIXCl3-x.

Figure 3.3. Observed deposition rates on sensors 1 and 2 during the evaporation of pure MAI.

3.2.2.2 Sequential Evaporation of MAI and PbCl2

Sequential evaporation of MAI and PbCl2 was performed in the vacuum evaporator.

Using this method, only one material at a time was evaporated. Also, the final thickness of each layer was more important than the deposition rate so the process was much easier to control. First, PbCl2 thin films with thicknesses of 50, 100, 150 and

200 nm were deposited on the substrates. Then, without breaking vacuum, 200 nm of

MAI was deposited on the PbCl2-coated substrates. After the evaporations, the samples were annealed for 60 minutes at 110 °C on a hotplate in the glovebox. XRD analysis was performed to check the formation of MAPbIXCl3-x perovskite with the results shown in Figure 3.4. The main peaks at 14.15°, 28.45° and 31.95° correspond to diffractions from the (110), (220) and (310) planes of the tetragonal phase of

MAPbIXCl3-x perovskite, respectively. It can be seen that 100 nm PbCl2 and 200 nm MAI produced the best crystalline MAPbIXCl3-x perovskite thin film with minimum impurities.

Figure 3.4. XRD patterns of films deposited by sequential evaporation followed by a post-deposition

annealing at 110 °C for 60 min. “*” indicates the peaks for the tetragonal phase of MAPbIXCl3-x

perovskite, “#” indicates the peaks for MAI and “+” indicates the peaks for PbI2.

The effect of post-deposition annealing time on film quality was investigated further.

Films deposited by sequential evaporation of 100 nm PbCl2 and 200nm MAI were annealed on a 110 °C hotplate for 40, 60, 80 and 100 min and their XRD patterns were compared, as shown in Figure 3.5. It can be seen that films annealed for different time all show good crystallinity of MAPbIXCl3-x perovskite but a small peak corresponding to

PbI2 appears at around 12.8° as the annealing time increases.

Figure 3.5. XRD patterns of films deposited by sequential evaporation followed by post-deposition

annealing on a hotplate at 110 °C for 40, 60, 80 and 100 min in the glovebox. “*” indicates the peaks

for the tetragonal phase from MAPbIXCl3-x perovskite.

Twenty solar cells were fabricated from films with four different post-deposition annealing times. Their efficiencies measured by reverse scan are compared in Figure

3.6a. It can be seen that devices annealed for 40 and 60 min had efficiencies below 0.5% while longer annealing time resulted in improved performance. The best device with

80 min post-deposition annealing had a PCE of 3.7% and the best device with 100 min post-annealing had a PCE of 6.3%. The J-V curves of the best performing device are presented in Figure 3.6b.

Figure 3.6. a) Efficiencies of 20 cells processed with different post-deposition annealing times from 40

to 100 minutes. b) J-V curves of the best performing device that was annealed for 100 minutes.

The performance might be improved further by optimizing the film thickness, PbCl2:

MAI ratio, substrate temperature during the evaporation processes, the post- annealing temperature/time or by using multilayer stacks. Unfortunately, a serious fault occurred and the turbomolecular pump of the evaporator was damaged.

Investigation revealed that MAI powder had collected in the pump and corroded the turbine blades. Consequently, the evaporation of MAI was no longer carried out in the vacuum evaporation system and a modified vapour-assisted evaporation (atmospheric pressure MAI vapour deposition process) was proposed to fabricate MAPbIXCl3-x perovskite solar cells, as discussed in the following section.

3.3 Fabrication of MAPbIXCl3-x by Vapour-Assisted

Evaporation Method

3.3.1 Methods and Experimental Section

The vapour-assisted evaporation process (atmospheric pressure MAI vapour-assisted process) is illustrated in Figure 3.7. First, 150 nm of PbCl2 was deposited on the substrate by thermal evaporation under vacuum, as discussed previously. Then, instead of depositing MAI in the vacuum evaporator, the samples were exposed to

MAI vapour at atmospheric pressure in a glovebox. Before the MAI vapour deposition, a closed glass petri dish with MAI powder at the bottom was pre-heated at 130 for

30 min. Then the PbCl2 coated substrate was transferred into the container and ℃ exposed to the MAI vapour. The PbCl2-coated substrate was placed approximately 2 mm away from the MAI powder, facing towards it. Because the substrate was very close to the container bottom and the hotplate, the MAI deposition and the annealing that crystallizes MAPbIXCl3-x from MAI and PbCl2 took place almost simultaneously.

Therefore, a post-deposition annealing process was not necessary for this method.

After removing the sample from the MAI vapour, it was blown by a nitrogen gun to remove excess MAI powder from the surface.

Figure 3.7. An illustration of the MAI vapour-deposition process.

Solar cells were fabricated and characterised using the procedures described previously in section 3.2.1.3, except for the deposition of the MAPbIXCl3-x perovskite layer. Top view scanning electron microscopic (SEM) images were obtained using a field emission SEM (NanoSEM 230).

3.3.2 Results and Discussions

3.3.2.1 MAPbIXCl3-x Perovskite Formation Process

The MAPbIXCl3-x perovskite formation process is recorded in Figure 3.8. The top row shows how the colour of the films change during the MAI vapour deposition process and the bottom row shows the corresponding SEM image of the surfaces. It can be seen that a light brown coloured film formed after the PbCl2-coated substrate was exposed to MAI vapour for 45 min and the film became dark brown after 80 min exposure. From the SEM images it can be seen that films with 45 or 80 min exposure to

MAI vapour have different morphologies. 80 min MAI vapour treatment produced larger and more compact grains than 45 min. These compact grains grew even larger after 110min of MAI exposure. The dark brown film started to discolour from the centre after 120 min of MAI vapour treatment. At the same time, the MAPbIXCl3-x perovskite started to degrade, the grain shape became less clear and pin holes

64 appeared. It appears likely that the best MAI vapour exposure time is in the range 80-

110 min.

Figure 3.8. Photographs of PbCl2 films after exposure to MAI vapour for different times and the

corresponding top view SEM images of the films. The scale bars indicate a length of 500 nm.

3.3.2.2 I-V Characteristics of MAPbIXCl3-x Perovskite Solar Cells (Small Area Devices)

Solar cells were made from MAPbIXCl3-x perovskite films that had been exposed to MAI vapour for different times from 80 to 120 min. The Au top electrodes were deposited by thermal evaporation through a shadow mask defining solar cell active area at

0.4cm2 and the efficiencies were measured by an aperture with a window area of

0.159cm2. Five cells were made using each exposure time and the results are shown in

Figure 3.9a. When the MAPbIXCl3-x perovskite was just formed after 80 min MAI vapour deposition, the devices had efficiencies below 2%. When the MAI vapour deposition time was increased to 90 and 100 min, the efficiencies improved to around 5% to 6%.

110 min of exposure to MAI vapour produced the best performing devices with efficiencies up to 11%, and a Voc of 0.87V, a Jsc of 25.08 mA/cm2 and a FF of 50.6%

(there was some overestimation of the Jsc and efficiency probably due to the improper calibration of the solar simulator). As expected, 120 min exposure to MAI vapour degraded the MAPbIXCl3-x perovskite and produced solar cells with efficiencies <2%.

The J-V curves of the device with the highest efficiency are shown in Figure 3.9b. It can be seen that the reverse and forward scans give very different results. The 11.0% PCE obtained by the reverse scan dropped to only 2% in the forward scan. During the MAI vapour deposition process, MAPbIXCl3-x perovskite formation starts from the surface of the PbCl2 film. The severe hysteresis observed in Figure 3.9b may be caused by unreacted PbCl2 remaining at the bottom of the film. Further optimization is necessary to address this hysteresis problem.

Figure 3.9 a) PCE of MAPbIXCl3-x perovskite solar cells made from films with different exposure times

to MAI vapour, as measured in reverse scan. b) The J-V curves of the best performing device which

had 110 min exposure to MAI vapour.

3.3.2.3 I-V Characteristics of MAPbIXCl3-x Perovskite Solar Cells (Large Area Devices)

The vapour-assisted evaporation process (atmospheric pressure MAI vapour-assisted method) was used to fabricate larger devices on 6.25 cm2 substrates with an active area of around 3cm2, instead of on a small 1cm2 substrate with an active area of only ~

0.4 cm2. Gold back contact is evaporated onto almost the whole area of the perovskite

67 film, in which the active area of these cells was ~3 cm2, as defined by the area of the perovskite film covered by the Au contact and shown in Figure 3.10a. To check the efficiency uniformity within the large area device and to keep the results comparable to those obtained previously using small cells, J-V measurements were made at 3 points using a 0.159 cm2 (4.5 mm diameter) aperture. The middle of the cell gave the best performance with a PCE of 8.0%, a Voc of 0.73V, a Jsc of 21.8 and a FF of 50.38% under reverse scan and a PCE of 2.7% under forward scan. The J-V curves are presented in Figure 3.10b. As can be seen, regions to the left and right only produced about half the performance of the central part. This is probably caused by the temperature variation of the hotplate surface during the MAI vapour process whereby different regions were annealed at different temperatures. As a result, different areas of the large perovskite film may have different crystal quality. For the hotplates used in our laboratory, the centre always has higher temperature than the sides. Therefore, the side area of the perovskite film with a lower annealing temperature may not be fully crystallized, resulting in lower efficiencies. This non-uniformity problem can be addressed simply by using a hotplate with a uniform surface temperature.

Figure 3.10 a) Images of devices fabricated on 2.5 cm square substrates showing their relatively large

active area. b) The efficiency distribution across the active area of a large area device and the J-V

curves of the middle region that produced the highest efficiency. Efficiencies were measured using a

0.159 cm2 aperture. 3.4 Conclusion

This chapter described the fabrication of MAPbIXCl3-x perovskite solar cells by physical vapour based deposition methods. Firstly, dual source thermal co-evaporation of MAI and PbCl2 based on a literature method was used to form MAPbIXCl3-x perovskite thin films. Controlling the rates of MAI and PbCl2 deposition at the same time was difficult, as the evaporation of MAI is not line-of-sight and the MAI vapour affects the deposition rate observed at the PbCl2 sensor. To address this, a process using sequential evaporation of MAI and PbCl2 was proposed for the deposition of

MAPbIXCl3-x perovskite thin films and a solar cell with a PCE of 6.3% by reverse scan was fabricated by this method. Before any further optimization was possible, however,

MAI vapour corroded the vacuum pump, indicating that thermal evaporation of MAI is 69 not amenable to vacuum evaporation systems. Therefore, a vapour-assisted evaporation process (atmospheric pressure evaporation method) was proposed to fabricate MAPbIXCl3-x perovskite thin films in which the PbCl2 was first coated onto the substrate by thermal vacuum evaporation and then the substrate was removed from the vacuum evaporator and transferred into a heated container with MAI powder. MAI vapour formed by heating the container on a hotplate at 130 was deposited on and reacted with the PbCl2 substrate, forming MAPbIXCl3-x perovskite.℃ A PCE of reverse scan as high as 11.0% was achieved by this method even though hysteresis was severe.

2 MAPbIXCl3-x perovskite solar cells with active area around 3 cm were also demonstrated by this vapour-assisted evaporation process (atmospheric pressure thermal evaporation method). For the fabrication methods mention above, the corresponding devices all showed large hysteresis. Hysteresis in perovskite solar cells fundamentally results from ion migration during solar cell operation and barriers against carrier extraction/transportation. In our cases, the hysteresis were mainly from the barriers against carrier extraction/transportation and can be explained in several aspects detailedly: one is that in these method a compact and well crystalized PbCl2 layer was firstly evaporated onto the substrates and the compact PbCl2 was difficult to convert to perovskite after the MAI was evaporated onto the PbCl2 surface (perovskite formed from this surface) upon annealing, probably leaving a thin unreacted PbCl2 layer between perovskite and the TiO2, acting as a barrier and hindering the carrier transfer; another reason is the nature of planar device architecture where the interface between TiO2 and perovskite is not sufficient enough for the carrier extraction; also, the cause of the hysteresis may be the imperfect quality of the perovskite layer which may contains defects that trap the carriers, acting as barriers. 70

Further optimization is necessary to reduce the hysteresis and improve the photovoltaic performance of the perovskite solar cells fabricated by these methods.

References:

1. Liu, M., M.B. Johnston, and H.J. Snaith, Efficient planar heterojunction perovskite solar cells by vapour deposition. Nature, 2013. 501(7467): p. 395-398. 2. Hu, H., et al., Vapour-based processing of hole-conductor-free CH3NH3PbI3 perovskite/C60 fullerene planar solar cells. RSC Advances, 2014. 4(55): p. 28964-28967. 3. Chen, C.W., et al., Efficient and uniform planar‐type perovskite solar cells by simple sequential vacuum deposition. Advanced Materials, 2014. 26(38): p. 6647-6652.

Chapter 4 Inorganic CsPbIBr2 Perovskite

Solar Cells by Dual Source Thermal Co-

Evaporation

4.1 Introduction

It has been proved that the dual source thermal co-evaporation method is a simple and efficient method to fabricate uniform perovskite thin films and is promising for scale-up production of perovskite solar cells[1-6]. However, the evaporation of organic

MAI is not line-of-sight, i.e., non-directional due to its relatively high vapour partial pressure. This makes the accurate control of the rate of PbCl2 deposition difficult as its rate reading is inevitably affected by the MAI sublimation during co-deposition as discussed in Chapter 3. Additionally, the organic molecule MAI is found to be detrimental to the vacuum evaporation system. Therefore, this dual source thermal co-evaporation method is more suitable for the deposition of inorganic materials, like inorganic Cs cation based perovskites.

The inorganic caesium (Cs) perovskites have the potential of better thermal and light stability compared to MA or FA-based perovskites. It has been shown that these inorganic Cs perovskites are stable at 250°C and 330°C during processing while the organic containing perovskites degrade[7, 8].

This chapter explores the fabrication of inorganic caesium lead bromide iodide perovskite (CsPbIBr2) perovskite solar cells by dual source thermal co-evaporation of

72 inorganic caesium iodide (CsI) and lead bromide (PbBr2). A series of experiments investigating the effect of post-anneal conditions on the perovskite crystalline structure is conducted. Films achieve best quality in terms of crystallinity, thickness uniformity, and grain size uniformity when the samples are annealed at 250 for

10min. Finally, a hole-transport-material-free planar solar cell of CsPbIBr2 is ℃ demonstrated for the first time, achieving an efficiency of 4.7% by reverse scan. Also this is the first demonstration of caesium lead halide perovskite solar cell by co- evaporation method. Compared to the perovskites with organic cations, CsPbIBr2 demonstrates better thermal stability. The mixed halide CsPbIBr2 perovskite absorber with a bandgap of 2.05 eV has a potential to be used in a 3-junction tandem (Figure 1.3) as the quality of the material and therefore voltage of the device further improves. 4.2 Experimental Section

4.2.1 Device Substrate Preparation

FTO-coated glass (TEC15, 15 Ω/☐sheet resistivity) is patterned by laser scribing, followed by cleaning by sonication in solutions of 2% Hellmanex in deionized water, acetone and isopropanol for 15 minutes. After drying, the substrate is treated by UV ozone cleaner for 10min. To form the compact TiO2 layer, a solution of titanium diisopropoxide bis(acetylacetonate) in ethanol is deposited on the clean substrates by spray pyrolysis at 450°C and the substrate is subsequently annealed on a hot plate at

450°C for 40 min.

4.2.2 Dual Source Thermal Co-Evaporation of the Perovskite

Absorber

The dual source thermal co-evaporation of the two precursors caesium iodide (CsI) and lead bromide (PbBr2), are carried out in a thermal evaporation system (Kurt J. Lesker

Mini Spectros) integrated in a glove box. CsI and PbBr2 are loaded in separate crucible heaters and the sample substrates are fixed on a rotatable substrate holder with the compact TiO2 side facing towards the precursor sources. After the pressure of the

-6 evaporator chamber is pumped down to 10 mbar, CsI and PbBr2 are then heated to the set temperature of 350°C and 180°C, respectively. Once the temperatures are reached, the shutter for each source is opened to commence deposition. The temperature of the substrate holder is kept at 20°C or 75°C during the deposition. The deposition rates of CsI and PbBr2 are set at 0.21 Å/s and 0.2 Å/s respectively to achieve a molar ratio of 1:1 for the two materials. After the evaporation, the samples are annealed on a hot plate at 100°C or 250°C for 10 min in the glove box. To complete device fabrication, 100 nm of gold is thermally evaporated on the samples to form rear contacts.

4.2.3 Characterization

X-ray photoelectron spectroscopy (XPS) measurement is performed on a Thermo

Scientific ESCALAB250Xi, with a monochromatic Al X-ray source, to confirm the chemical composition of the mixed halide Cs perovskite samples.

X-ray diffraction (XRD) patterns are measured using a PANalytical Xpert Materials

Research diffractometer system with a Cu Kα radiation source (λ=0.1541 nm) at 45 kV and 40mA.

Top view and cross-sectional scanning electron microscopic (SEM) images are obtained using a field emission SEM (NanoSEM 230). The EDS measurements are also carried out by the NanoSEM 230 using a Bruker SDD-EDS detector.

The optical reflection and transmission spectra are measured using Perkin Elmer

Lambda1050 UV/Vis/NIR spectrophotometer.

PL spectra are measured using a back scattering geometry with 405 nm laser excitation and a thermo-electrically cooled Si-CCD detector. PL imaging, fluorescence lifetime imaging microscopy (FLIM) and the PL decay traces are measured by Microtime200 microscope (Picoquant) using time correlated single photon counting (TCSPC) technique with excitation of 470 nm laser at 5 MHz repetition rate and detection through 620/40 nm band-pass filter.

The current density–voltage (J–V) measurements are performed using an IV5 solar cell

I–V testing system from PV measurements, Inc. (using a Keithley 2400 source meter) under illumination power of 100 mW/cm2 by an AM1.5G solar simulator (Oriel model

94023A) with an 0.159 cm2 aperture and a scan rate of 1.2 V/s.

All measurements are undertaken at room temperature in ambient condition. 4.3 Results and Discussions

4.3.1 Elemental Analysis of the Deposited CsPbIBr2 Thin Film

As described in the Experimental Section, the CsPbIBr2 samples are prepared by evaporating the same molar quantity of CsI and PbBr2 onto the substrates. The chemical composition of the samples is evaluated by X-ray photoelectron spectroscopy

(XPS). The atomic ratio of Pb/Cs and Br/I is estimated to be 1.1 and 2.3 respectively, which is in good agreement with the CsPbIBr2 composition. The XPS spectra are shown 75 in Figure 4.1. An Energy-dispersive X-ray spectroscopy (EDS) measurement at 15kV is also carried out by a 20μm line scan of the CsPbIBr2 film showing the atomic ratios of

Pb/Cs and Br/I to be 1.2 and 1.94 respectively. The EDS spectra are shown in Figure 4.2.

One reason for the deviation between EDS and XPS results is the difference in accuracy between the measurements (XPS, ±5%, EDS, ±15%). It is also noted that the XPS carried out measures the elemental composition of ~10 nm in depth from the surface.

For bulk measurement Ar ion etching will be required causing damage to the CsPbIBr2 film. Given the uniform column grains formed from the bottom to the surface of the

CsPbIBr2 film as shown in Figure 4.11(a), the atomic ratios of Pb/Cs and Br/I measured by XPS will be a reasonable representation for the entire film. On the other hand, the depth of EDS measurement can be ~1000nm, providing good bulk information of the elements.

Figure 4.1. XPS spectra of CsPbIB2 film deposited at 75℃ and annealed at 250℃ for 10 min: (a) Pb 4f,

(b) Cs 3d, (c) Br 3d and (d) I 3d.

Figure 4.2. EDS spectra on CsPbIBr2 film by a 20μm line scan. The atomic ratios 1.2 of Pb/Cs and 1.94

of Br/I are calculated by averaging the atomic ratios of the 100 lines-canning points respectively.

4.3.2 Crystalline Structure and Morphology Analysis of the

CsPbIBr2 Thin Film

Figure 4.3 shows the X-ray diffraction (XRD) patterns of the CsPbIBr2 films deposited via dual source thermal co-evaporation on c-TiO2/FTO glass substrates at 20°C or 75°C followed by a post anneal at 100°C or 250°C for 10min. The main diffraction peaks at

15.05°，21.35° and 30.25° correspond to the (100), (110) and (200) planes of the

CsPbIBr2 perovskite cubic phase. The perovskite crystallinity improves with the increase of substrate temperature during deposition as well as the post annealing temperature.

Figure 4.3 XRD patterns of the CsPbIBr2 films on c-TiO2/FTO glass substrates followed by post

annealing at different temperatures for 10min. Ts represents the substrate temperature during

deposition and Tp represents the temperature of the post-anneal.

The morphology of the perovskite films are studied by top view scanning electron microscopy (SEM), as shown in Figure 4.4. The films without post-annealing show small crystal grains in Figure 2(a) and (c). After post annealing, the crystal grains grow larger

78 and are more compact. The film deposited at 75°C substrate temperature gives larger crystalline grains than that deposited at 20°C. When the 75°C deposited film is annealed at 250°C, its grains grow as large as 500-1000nm in size.

Figure 4.4. Top view SEM images of the CsPbIBr2 perovskite deposited via dual source thermal co-

evaporation on c-TiO2/FTO glass substrate at different conditions: a) Ts: 20°C without post annealing;

b) Ts: 20°C, annealed at 250°C for 10 min, c) Ts: 75°C without post annealing, and d) Ts: 75°C,

annealed at 250°C for 10 min. The scale bar is 1 μm.

4.3.3 Optical Properties of the CsPbIBr2 Thin Film

The optical transmission and reflection spectra of the CsPbIBr2 films were measured using a Perkin Elmer LAMBDA 1050 UV-VIS-NIR spectrophotometer. The absorption coefficient of the samples was calculated using = ln (1 )/ , where R and T −1 훼 푡 � − 푅 푇� 79 are the reflection and transmission and t is the thickness of the samples. All of the samples annealed at different temperatures show similar absorption coefficients, over

5x104 cm-1 in the absorption range of 250-580 nm, which are very similar to that of

CH3NH3PbI3. Figure 4.5(a) shows the absorption coefficient of the sample deposited at

75°C and annealed at 250°C. The onset of optical bandgap edge transition is 604 nm

(2.05eV). Compared with CsPbI3 (Eg~1.73eV)[8] and CsPbBr3 (Eg~2.25eV)[9], the optical bandgap of the mixed halide CsPbIBr2 shows a linear relationship of the content of Br in the mixed halide Cs perovskites which is in good agreement with the Vegard's law, see Figure 4.6.

Figure 4.5. Optical properties of CsPbIBr2 sample deposited at 75°C and annealed at 250°C for 10min:

(a) absorption coefficient and steady state PL spectrum, (b) PL image (contrast is indicator of PL

intensity), and (c) typical PL decay traces of a small-bright grain (blue line), a large-dim grain (red line),

and a dark grain (black line) in (b).

Figure 4.6. Optical bandgap comparison of CsPbI(3-x)Brx, demonstrating the bandgap of CsPbIBr2 has a

linear relationship of the content of Br in the mixed halide Cs perovskites.

The photoluminescence (PL) spectrum of the CsPbIBr2 shows a peak at 2.00eV (620nm), which is very close to the optical band edge and does not undergo shift under light soaking, see Figure 4.7. In the case of CH3NH3Pb(BrxI1-x)3, low energy PL features appear upon light soaking observed by Hoke et al. suggesting photo-induced halide segregation resulting in a reduction in the electronic bandgap and quasi-Fermi level splitting[10]. This causes the observed red-shift in PL and limits the achievable voltage as can be seen in MA lead mixed halide devices [11-13]. The lack of PL red-shift in Cs lead mixed halide perovskite which is also observed in other works is advantageous over the organic counterparts as the lack of bandgap change allows true bandgap tunability of Cs perovskites. [14] [15] [16]

Figure 4.7. PL Spectrum at different excitation intensity. Time interval for different intensity

measurement is 30-60 seconds. This figure shows that the film has no obvious peak shift at elevated

excitation intensity and no halide segregation under illumination.

Figure 4.5(b) shows the PL image of the perovskite film excited at 470 nm and detected at 620 nm. The grain size detected by the imaging is very similar to that of SEM images.

The small grains that are brighter indicate higher PL efficiency while the larger and typically dimmer and dark grains have lower PL efficiency. The typical PL decay traces of a small-bright grain; a large-dim grain; and a dark grains were measured using time correlated single photon counting (TCSPC) technique as shown in Figure 4.5(c). Using bi-exponential decay function = ( / ) + ( / ), the PL decay trace

1 1 2 2 of the large grains is fitted to determine푦 퐴 푒푥푝 the− decay푡 휏 times퐴 푒푥푝 of the−푡 fast휏 (τ1) and slow (τ2) components for each region, while the PL decay trace of the small gains is fitted with dingle exponential function with one decay time component, as tabulated in Table 1.

The presence of the fast component in the PL decay for the dim (and dark) regions of the CsPbIBr2 film indicates the presence of defect trapping. Otherwise, a typical lifetime of 9.35 and 17.7 ns found for the CsPbIBr2 film corresponds to the carrier 82 recombination time. This lifetime is comparable with that of CH3NH3PbBr3 film on c-

TiO2 (3.5ns)[17].

Table 1. Typical lifetimes extracted from TCSPC for different regions of a CsPbIBr2

film.

Lifetime Bright Dim Dark

Fast Component (ns) N/A 0.76 1.37

Slow Component (ns) 9.35 17.7 16.4

4.3.4 Thermal Stability of the CsPbIBr2 Thin Film

Preliminary thermal stability tests are also performed on CsPbIBr2 films which involves heat treatment at 200°C on a hot plate in a glove box. The 6-hour-interval results are shown in Figure 4. The XRD patterns in Figure 4.8(a) show no detectable phase change or impurity peaks in the films after the heat treatments. Figure 4.8(b) indicates no film degradation after the heat treatments. In addition, the top view SEM images in Figure

4.9 show no obvious morphology changes. All these provide evidences that CsPbIBr2 has good thermal stability. Change of colour of MAPbI3 film is reported when film has been heated for more than 30 mins at 150 °C in air[18]. We therefore test the thermal stability of CsPbIBr2 films in ambient conditions as shown in Figure 4.10. The results show that CsPbIBr2 films do not show obvious degradation after heated at 150°C on a hot plate for 120 min in air, indicating its better thermal stability than MAPbI3.

Figure 4.8. (a) XRD patterns of the CsPbIBr2 films with no thermal treatment, heated at 200°C for 6

hours and 12 hours in a glove box; (b) the corresponding photos of the CsPbIBr2 films.

Figure 4.9. Comparison of top view SEM images of the CsPbBrI2 film without heating (a) and heated

on a hot plate at 200 °C for 12 hours in a glovebox (b). There is no obvious morphology change

observed by SEM.

Figure 4.10. (a) XRD patterns of the CsPbBrI2 films without thermal treatment and heated at 150 °C on

a hot plate for 120 min in air; (b) Photos of the CsPbBrI2 films without thermal test, heated at 150 °C

for 20 min, 60min and 120min in air.

4.3.5 Electrical Characteristics of the CsPbIBr2 Solar Cells

The cross-sectional SEM image of a planar CsPbIBr2 cell and the corresponding energy band diagram are shown in Figure 4.11(a) and Figure 4.11(b). The electrical characteristics under illumination of the best performing HTM free planar CsPbIBr2 cell in this work are shown in Figure 4.11(c). It is fabricated at Ts= 75℃ and Tp= 250°C for

10min. The PCE of this cell measured under reverse scan at 1.2V/s is 4.7% with a VOC of

2 959 mV, JSC of 8.7 mA/cm , and fill factor of 56%. Under forward scan, PCE= 3.7%; JSC=

2 8.7 mA/cm ; VOC = 818 mV and FF=52%. This gives an average PCE of 4.2% which is 89% of the PCE measured under reverse scan.

Figure 4.11. (a) Cross-sectional SEM image of a planar CsPbIBr2 cell and (b) the energy band diagram,

250°C for 10 min.

Stabilized PCEs measured at maximum power points have also been performed.

Normalised results are given in Figure 4.12. They are within 87% to 95% of the highest

PCEs measured (usually under reverse scan). This is similar to what is observed in Ref.

[19] and by many prominent research groups that the stabilised PCEs are in good agreement with the averaged PCEs from reverse and forward scans.

Figure 4.12. Stabilized PCE of five CsPbIBr2 cells measured at maximum power points. The cell

efficiency ranges from 2.1% to 4.7%. For all these cells, the PCE stabilizes to 87% ~94% of that

obtained by reverse scan in less than 10 sec.

The cell fabricated at Ts= 20°C and Tp= 250°C for 10min also achieves respectful

2 performance with a PCE of 3.6%, VOC of 594 mV, JSC of 10.0 mA/cm and fill factor of

61%, measured under the same condition. Its current density –voltage curve is shown in Figure 4.13. Compared with the two-step solution processed CsPbBr3 cells[7], our

CsPbIBr2 devices exhibit larger hysteresis, probably due to the use of planar structure and/or the lack of HTM resulting in worse carrier extracting interfaces[8]. The fill factor is also relatively low again due to the lack of electron blocking or hole transporting layer. To solve these problems, one can use inorganic HTM in future work. The open- circuit voltage is smaller than the expected from the absorber’s bandgap due to defects in the material and un-optimised interface. However, compared with the Voc of CsPbI3 cells (850 mV)[8], the higher VOC from our mixed halide devices demonstrates 87 the potential of tunable and larger open-circuit voltage by incorporating Br content. In addition, the short-circuit current of the cell can be further improved by increasing the thickness of CsPbIBr2 absorber to be beyond the current 190 nm.

Figure 4.13. Light current density-voltage curves of the device made at Ts 20°C and Tp 250°C for 10min.

The arrows show the scan directions.

Large area device is fabricated by this dual source thermal evaporation method at Ts=

75°C and Tp= 250°C for 10min as well, as shown Figure 4.14a. The active area of

1.5cm2 is defined by the area of the perovskite film covered by the Au contact. The photovoltaic performance is measured by an aperture with an opening area of around

1.5cm2, similar with the gold back contact shape. The results are presented in Figure

4.14b that by a reverse scan it gives a PCE of 3.71% with a Voc of 0.718V, a Jsc of

9.678mA/cm2 and a FF of 53% and by a forward scan it gives a PCE of 2.75% with a Voc of 0.597V, a Jsc of 9.677mA/cm2 and a FF of 48%. The performance of the large area device, which is not as good as the small area device, can be further improved by a front metal grid design to reduce the sheet resistance.

Figure 4.14. (a) Image of a large area device and (b) the photovoltaic performance.

4.3.6 Deposition of CsPbIBr2 on Silicon Substrate

With the potential as a top cell in a tandem solar cell, the deposition of CsPbIBr2 on silicon substrate is demonstrated in this work. Firstly, CsPbIBr2 is deposited onto a large

Si wafer with a standard industrialized size of 6 inch, as shown in Figure 4.15. It can be seen that the CsPbIBr2 coated Si wafer shows a uniform and specular surface, indicating the ability to fabricate uniform perovskite thin film on large substrate for scalable production by dual source thermal evaporation.

Figure 4.15. Deposition of CsPbIBr2 on 6 inch silicon wafer by dual source thermal co-evaporation

Also, the deposition of CsPbIBr2 on texture Si wafer is demonstrated by this dual source thermal co-evaporation method, as shown in Figure 4.16a. The top view SEM image in Figure 4.16b shows that the deposited CsPbIBr2 thin film has uniform and pin- hole free coverage on the textured Si wafer with pyramid size of around 8μm. Textured surface in a solar cell structure can reduce the reflection of the incident light by trapping the light in the solar cell and therefore improve the photovoltaic performance.

The texture design has been widely applied in the solar cell manufacture, particularly in the Si based solar cell. Therefore the demonstration of depositing CsPbIBr2 onto texture and large size Si substrate makes it promising in the application of tandem solar cells with Si.

Figure 4.16. (a) Image of CsPbIBr2 thin film on texture Si substrates with pyramid size of around 8μm

by dual source thermal co-evaporation and (b) the top view SEM image. 4.4 Conclusion

In summary, we demonstrate dual source thermal co-evaporated CsPbIBr2 HTM free planar perovskite solar cell devices for the first time. The films show a typical optical bandgap of 2.05 eV and stable PL emission at 2.00 eV. No light-induced PL shift is 90 observed in the Cs lead mixed halide perovskites that is otherwise observed in MA lead mixed halide perovskites indicating genuine bandgap tunability in mixed halide Cs perovskite. This shows the high voltage potential of the Cs lead mixed halide perovskites suitable for tandem silicon cells in the future as the fabrication techniques for these materials improve with time. The best performing device achieves a conversion efficiency of 4.7% under reverse scan and 3.7% under forward scan, the highest of its kind deposited by dual source co-evaporation without the use of hole transport material. The efficiency of this device can be further improved by inserting a

HTM layer such as P3HT between the perovskite layer and gold contact by which the open circuit voltage can be improved. The CsPbIBr2 films have also shown good thermal stability in N2 environment and ambient conditions. Large area device with an active area of 1.5 cm2 is fabricated as well, indicating the possibility in scalable production. The depositions of CsPbIBr2 thin film onto textured silicon substrate and large size (6 inch) silicon wafer have also been demonstrated in this work, showing the potential in the application of tandem solar cell on silicon bottom cell.

References:

1. Liu, M., M.B. Johnston, and H.J. Snaith, Efficient planar heterojunction perovskite solar cells by vapour deposition. Nature, 2013. 501(7467): p. 395-398. 2. Gao, C., et al., Formation of organic–inorganic mixed halide perovskite films by thermal evaporation of PbCl2 and CH3NH3I compounds. RSC Advances, 2015. 5(33): p. 26175- 26180. 3. Kim, B.-S., et al., Fully vacuum–processed perovskite solar cells with high open circuit voltage using MoO3/NPB as hole extraction layers. Organic Electronics, 2015. 17: p. 102-106. 4. Ono, L.K., et al., Fabrication of semi-transparent perovskite films with centimeter-scale superior uniformity by the hybrid deposition method. Energy & Environmental Science, 2014. 7(12): p. 3989-3993.

5. Kim, B.-S., et al., Composition-controlled organometal halide perovskite via CH3NH3I pressure in a vacuum co-deposition process. Journal of Materials Chemistry A, 2016. 4(15): p. 5663-5668. 6. Zhao, D., et al., Annealing-free efficient vacuum-deposited planar perovskite solar cells with evaporated fullerenes as electron-selective layers. Nano Energy, 2016. 19: p. 88- 97. 7. Kulbak, M., D. Cahen, and G. Hodes, How important is the organic part of lead halide perovskite photovoltaic cells? Efficient CsPbBr3 cells. The journal of physical chemistry letters, 2015. 6(13): p. 2452-2456. 8. Eperon, G.E., et al., Inorganic caesium lead iodide perovskite solar cells. Journal of Materials Chemistry A, 2015. 3(39): p. 19688-19695. 9. Stoumpos, C.C., et al., Crystal growth of the perovskite semiconductor CsPbBr3: a new material for high-energy radiation detection. Crystal Growth & Design, 2013. 13(7): p. 2722-2727. 10. Hoke, E.T., et al., Reversible photo-induced trap formation in mixed-halide hybrid perovskites for photovoltaics. Chemical Science, 2015. 6(1): p. 613-617. 11. Noh, J.H., et al., Chemical management for colorful, efficient, and stable inorganic– organic hybrid nanostructured solar cells. Nano letters, 2013. 13(4): p. 1764-1769. 12. Suarez, B., et al., Recombination study of combined halides (Cl, Br, I) perovskite solar cells. The journal of physical chemistry letters, 2014. 5(10): p. 1628-1635. 13. Kulkarni, S.A., et al., Band-gap tuning of lead halide perovskites using a sequential deposition process. Journal of Materials Chemistry A, 2014. 2(24): p. 9221-9225. 14. Akkerman, Q.A., et al., Tuning the optical properties of cesium lead halide perovskite nanocrystals by anion exchange reactions. Journal of the American Chemical Society, 2015. 137(32): p. 10276-10281. 15. Song, J., et al., Quantum Dot Light‐Emitting Diodes Based on Inorganic Perovskite Cesium Lead Halides (CsPbX3). Advanced materials, 2015. 27(44): p. 7162-7167. 16. Sharma, S., N. Weiden, and A. Weiss, Phase diagrams of quasibinary systems of the

type: ABX3—a′ BX3; ABX3—AB′ X3, and ABX3—ABX′ 3; X= halogen. Zeitschrift für Physikalische Chemie, 1992. 175(1): p. 63-80. 17. Sheng, R., et al., Methylammonium lead bromide perovskite-based solar cells by vapor- assisted deposition. The Journal of Physical Chemistry C, 2015. 119(7): p. 3545-3549. 18. Eperon, G.E., et al., Formamidinium lead trihalide: a broadly tunable perovskite for efficient planar heterojunction solar cells. Energy & Environmental Science, 2014. 7(3): p. 982-988. 19. Jeon, N.J., et al., Compositional engineering of perovskite materials for high- performance solar cells. Nature, 2015. 517(7535): p. 476-480.

Chapter 5 Inorganic CsPbI2Br Perovskite

Solar Cells by Dual Source Thermal Co-

Evaporation

5.1 Introduction

In Chapter 4, we have demonstrated planar HTM-free CsPbIBr2 solar cells with an efficiency of 4.7%[1]. The CsPbIBr2 perovskite film fabricated by dual source thermal evaporation have shown to be stable for 12 hours at 200 °C in a nitrogen glovebox and for 2 hours at 150 °C in atmosphere condition.

According to the efficiency limits calculated by detailed balance approach for multi- junction solar cells [2], the bandgap of CsPbIBr2 is higher than the optimum necessitating the development of lower bandgap Cs-based perovskite solar cells. This can be achieved by increasing the iodine content. However, increase in iodine content in Cs based perovskite affects its structural stability as the films convert to non- perovskites phase rapidly when exposed to air. Previous work has shown that as x increases from 1 to 2 in CsPbIxBr(3-x) film, air-stability deteriorates requiring the films and devices to be encapsulated for characterizations [1, 3, 4]. When x=3, CsPbI3 has even poorer stability [5, 6] requiring quantum dot–induced phase stabilization [7]. This shows that more work needs to be done towards phase stabilization of inorganic Cs based perovskite films.

In this chapter, we firstly demonstrate the fabrication of CsPbI2Br perovskite cells by dual source thermal co-evaporation method. CsPbI2Br has a more suitable (lower) bandgap (1.9 eV) than CsPbIBr2 (2.05 eV) as a photovoltaic energy harvesting material and a top cell in tandem solar cells. However, increase in iodine content reduces structural stability due to the preference towards the non-perovskite orthorhombic phase when the film is exposed to air. Then, we investigate the effect of the stoichiometry on the stability of CsPbI2Br thin films fabricated by dual source thermal co-evaporation, as well as the grain size, presence of impurities and nature of impurity grains, photoluminescence, morphology, and elemental distribution. The films are deposited by co-evaporation of CsBr and PbI2 followed by a post annealing at 300°C on a hot plate for 10 min in a N2 glovebox. Details on how to vary the stoichiometry by varying the deposition rates of CsBr and PbI2 during the dual source thermal co- evaporation process are presented. The film stoichiometry and morphology are analyzed by the energy-dispersive X-ray spectroscopy (EDS) and scanning electron microscopy (SEM). Interestingly, it is found that slight changes in CsPbI2Br stoichiometry can affect air stability significantly. The CsBr-rich film shows the best stability while the PbI2-rich sample is the worst.

The electrical performance of planar Glass/FTO/c-TiO2/perovskite/P3HT/Au perovskites solar cells using CsPbI2Br with different stoichiometry is compared. The stoichiometrically balanced CsPbI2Br cell achieves the best photovoltaic performance with a PCE of 7.7% for an active area of 0.159cm2 and 6.8% for a 1.2cm2 cell under reverse scan condition. The 1.2cm2 active area is the largest in inorganic caesium lead halide perovskites solar cells reported to date. This device remains 90% of the initial performance after 240 hours’ damp and heat test at 85 °C and 85% relative humidity. 94

5.2 Experimental Section

5.2.1 Device Substrate Preparation

FTO-coated glass (TEC15, 15Ω/☐sheet resistivity) is patterned by laser scribing, followed by cleaning by sonication in solutions of 2% Hellmanex in deionized water, acetone and isopropanol for 15 minutes respectively. After drying, the substrate is treated by UV ozone cleaner for 10min. To form the compact TiO2 layer, a solution of titanium diisopropoxide bis(acetylacetonate) in ethanol is deposited on the cleaned substrate by spray pyrolysis at 450°C and the substrate is subsequently annealed on a hot plate at 450°C for 10 min.

5.2.2 Dual Source Thermal Co-Evaporation of the Perovskite

Absorber

The dual source thermal evaporation of the two precursors caesium bromide (CsBr) and lead iodide (PbI2), are carried out in a thermal evaporation system (Kurt J. Lesker

Mini Spectros) integrated in a glove box. CsBr and PbI2 (Alfa Aesar) are loaded in separate crucible heaters and the sample substrates are fixed on a rotatable substrate holder with the compact TiO2 side facing towards the precursor sources. After the

-6 pressure of the evaporator chamber is pumped down to 10 mbar, CsBr and PbI2 are then heated to the set temperature of 300°C and 160°C, respectively. Once the temperatures are reached, the shutter for each source is opened to commence deposition. To produce a stoichiometrically balanced CsPbI2Br thin film, the deposition

−1 −1 rates of CsBr and PbI2 are set at 0.14 Å s and 0.22 Å s respectively. To achieve

−1 different stoichiometry, we fix the CsBr deposition rate at 0.14 Å s and vary the PbI2

−1 −1 −1 deposition rate from 0.22 Å s to 0.17 Å s or 0.27 Å s to produce CsBr-rich and PbI2- rich CsPbI2Br film, respectively. After the evaporation, the samples are annealed on a hot plate at 300°C for 10min in the glove box. The hole transporting layer is deposited on CsPbI2Br by spin coating a 15mg/ml P3HT (Sigma Aldrich) solution in chlorobenzene

(Sigma Aldrich) at 2000rpm for 60s. To complete device fabrication, 100nm gold is thermally evaporated on the samples to form rear contacts. Some of the samples are encapsulated to minimise degradation for some of the characterisations, e.g. photoluminescence (PL) imaging, time-resolved PL spectroscopy (TRPL) and current-voltage

(I-V) measurement of completed photovoltaic devices.

5.2.3 Characterization

Energy dispersive X-ray spectroscopy (EDS) measurements are carried out using a FEI

Verios equipped with an Oxford EDX detector and the data are collected with an accelerating voltage of 10 kV.

X-ray diffraction (XRD) patterns are measured using a PANalytical Xpert Materials

Research diffractometer system with a Cu Kα radiation source (λ=0.1541 nm) at 45 kV and 40mA.

Top view scanning electron microscopic (SEM) images are obtained at 5KV using a field emission SEM (NanoSEM 230). The SEM images in the EDS analysis parts are obtained by the EDS settings as described above.

The optical absorption spectra are measured using Perkin Elmer Lambda1050

UV/Vis/NIR spectrophotometer.

PL spectra are measured using a back scattering geometry with 405 nm laser excitation and a thermo-electrically cooled Si-CCD detector. The PL images are obtained by Leica

TCS SP5 microscopy. The time-resolve PL (TRPL) decay traces are measured by the

Microtime-200 (PicoQuant). Both PL imaging and TRPL are measured with 470 nm excitation and detection through a 620/40 nm band pass filter.

The current density–voltage (J–V) measurements are performed using a solar cell I–V testing system from Abet Technologies, Inc. (using class AAA solar simulator) under an illumination power of 100 mW cm-2 and a scan rate of 40mV s-1. All measurements are undertaken at room temperature in ambient condition.

The damp heat tests are carried out in an environmental chamber with conditions of

85 and 85% relative humidity. The sample is taken out for the I-V test every day and returned℃ into the chamber afterwards. The tests were performed by a collaborator Dr

Jizhong Yao. 5.3 Results and Discussions

5.3.1 Elemental and Optical Analysis

Details of material preparation, test structure and cell fabrication are listed in the

Experimental Section. The stoichiometry of the CsPbI2Br perovskite film is adjusted by varying the evaporation rate of PbI2 with the evaporation of CsBr fixed at 0.14 Å/s. The elemental composition of the evaporated films was analysed by EDS. The EDS results of three films deposited with different PbI2 evaporation rates are listed in Table 5.1.

When PbI2 is evaporated at a rate of 0.17 Å/s, a CsBr-rich CsPbI2Br film is produced. At a rate of 0.22 Å/s for PbI2, the CsPbI2Br film is stoichiometrically balanced. The slightly higher than expected iodine and bromine values is partly attributed to error in the EDS measurement. When PbI2 is evaporated at a rate of 0.27 Å/s, a PbI2-rich CsPbI2Br film is produced. Figure 5.1 shows the absorption and PL spectra of these three films. All of 97 the films have the same absorption onset at around 645 nm (1.9 eV) which is consistent with the reported bandgap of CsPbI2Br [3, 8]. The three films also have similar PL emissions at around 643nm. These results confirm that the change of PbI2 evaporation rate in the range of 0.17 Å/s to 0.27 Å/s does not change the optical absorption and PL emission position of the CsPbI2Br perovskite. However the absorption intensities of the three films are slightly different possibly due to the different surface roughness.

Table 5.1. Average atomic ratio of the samples deposited at different rates of PbI2 with

the evaporation rate of CsBr fixed at 0.14 Å/s

Sample PbI2 rate ( Å/s) Cs Pb I Br

CsBr-rich 0.17 1.3 1.0 2.0 1.6

Stoichiometrically

0.22 1.0 1.0 2.6 1.2 balanced

PbI2-rich 0.27 0.7 1.0 2.4 1.3

Figure 5.1. (a) and (b) The PL and absorption spectra of CsBr-rich, stoichiometrically balanced and

PbI2-rich CsPbI2Br perovskite films with a thickness around 230nm. 5.3.2 The Effect of Stoichiometry on the Air Stability

However, we have observed that these three films manifest very different air stability.

Photos of the CsBr-rich, stoichiometrically balanced and PbI2-rich CsPbI2Br films showing different rates of film degradation in air are shown in Figure 5.2a. PbI2-rich sample degrades completely in 5 min in ambient. The stoichiometrically balanced film starts to discolour after 10 min when left in air and becomes colourless after 40 min.

The CsBr-rich film on the other hand does not change colour after 40 min in air, indicating its better air stability.

The X-ray diffraction (XRD) patterns of un-encapsulated fresh films and air-exposed- films are shown in Figure 5.2b. The main diffraction peaks at 14.70°，20.95° and 29.65° correspond to the (100), (110) and (200) planes for the CsPbI2Br perovskite cubic phase[4, 8, 9]. The perovskite crystallinity of the fresh PbI2-rich CsPbI2Br sample is not as good as the fresh CsBr-rich and stoichiometrically balanced samples. Impurity peaks at 11.18° and 28.18° are present in the PbI2-rich CsPbI2Br film, and impurity peaks at

12.11° and 27.63° are present in the CsBr-rich CsPbI2Br film. These impurity peaks don’t correspond to any single compound of CsBr, PbI2 or CsPbI2Br and therefore they

99 could be from the mixtures of these compounds. Any grains containing these impurities may not be radiative and not contribute to the photovoltaic performance, as will be discussed later. It can be seen that the CsBr-rich and stoichiometrically balanced samples both have better perovskite phase initially. After 40 min of air- exposure, the stoichiometrically balanced and PbI2-rich samples experienced transition to non-perovskite orthorhombic phase which is found to be reversible up on heating at

300°C, while the CsBr-rich sample remains essentially the same, indicating better air stability.

Figure 5.2. (a) Photos showing different rates of degradation for un-encapsulated CsBr-rich,

stoichiometrically balanced and PbI2-rich CsPbI2Br films when exposed to air. (b) XRD patterns of the fresh and air-exposed CsBr-rich, stoichiometrically balanced and PbI2-rich CsPbI2Br films. The films are

100

deposited on FTO glass substrates. “*”corresponds to peaks from CsPbI2Br perovskite cubic phase,

“#”corresponds to peaks from CsPbI2Br non-perovskite orthorhombic phase, and “ ”corresponds to

peaks from the FTO substrates.

The morphology of the CsPbI2Br films are analysed by top view SEM, as shown in

Figure 5.3. The grain size of the CsBr-rich is notably smaller. The CsBr-rich films also have more brighter impurity grains [10]. The impurity grains in the PbI2-rich films have a different appearance. They are darker in the centre surrounded by brighter perimeter. The size of impurity grains is also different in CsBr-rich and PbI2-rich samples. Those in the CsBr-rich sample have a size of few hundred nanometres while those in the PbI2-rich sample are around 1μm in size. Upon air-exposure, both the stoichiometrically balanced and the PbI2-rich film experiences morphology change while the morphology of the CsBr-rich remains the same. The impurity grains in PbI2- rich sample become darker and bright perimeters disappear upon air-exposure. Upon closer examination of the PbI2-rich film, see high resolution SEM image in Figure 5.4, pin holes can be seen to emerge upon air-exposure. These results show that the

CsPbI2Br films with different stoichiometry have different surface morphology and the

CsBr-rich film have better air stability than the stoichiometrically balanced and PbI2- rich films.

101

Figure 5.3. Top view SEM images of un-encapsualted fresh (top row) and air-exposed (bottom row)

CsBr-rich, stoichiometrically balanced and PbI2-rich CsPbI2Br films. The films are deposited on FTO

glass substrates.

Figure 5.4. Top view SEM images of (a) fresh and (b) air-exposed PbI2-rich CsPbI2Br films. Damages

(pinholes) can be seen on the grains after air exposure.

5.3.3 The Effect of Stoichiometry on the Film Morphology and

Carrier Lifetime

102

To further understand the relationship between the impurity grains and the normal grains in the CsPbI2Br films, PL imaging is performed on the samples by using confocal

PL microscopy (Leica TCS SP5). Figure 5.5a shows the PL images detected through a

620/40 nm band pass filter, corresponding to the emission peak of the CsPbI2Br perovskite films. As shown in the figures, the stoichiometrically balanced film exhibits uniform PL emission across the film and the size of the grains observed under PL imaging is very similar to that observed under SEM. On the other hand, both CsBr-rich and PbI2-rich films exhibit non-uniform PL emission across the film. The brightest regions correspond to highest PL efficiency and they are typically grains with the smallest size. The less bright regions correspond to lower PL efficiency which are typically larger grains. There are regions that appear dark that are PL inactive and they correspond to impurity grains as observed under the SEM. The dark regions in the PbI2- rich films may correspond to the impurity peaks at 11.18° and 28.18° in the XRD pattern and those in the CsBr-rich films may correspond to the impurity peaks at 12.11° and 27.63°. As these impurity peaks cannot be assigned to the typical cubic phase of

CsPbI2Br perovskite and are non-radiative at 620/40 nm, they do not contribute to the photovoltaic performance.

The time-resolved PL (TRPL) measurement is also performed on the samples. Results are shown in Figure 5.5b. For the PL decay traces, a triple exponential function

(Equation 1) allows the best fit. The effective lifetime (Equation 2) is then extracted for the following quantitative analysis [11, 12]. y = A exp + A exp + A exp (1) t t t 1 �− τ1� 2 �− τ2� 3 �− τ3� = (2) A1τ1+A2τ2+A3τ3 τeff A1+A2+A3 103

It is calculated that PL effective lifetimes of the CsBr-rich, stoichiometrically balanced and PbI2-rich films are 1.15 ns, 7.15 ns and 4.05ns, respectively. The long lifetime observed in the stoichiometrically balanced film is likely due to the absence of the non- radiative impurity grains. It is later found out that stoichiometrically balanced film also produces higher device performance, as shall be discussed later.

Figure 5.5. (a) PL images and (b) PL decay traces of the CsBr-rich, stoichiometrically balanced and PbI2-

rich CsPbI2Br film detected at 640 nm. The intensity bar on the left of (a) is in the range of 0 to 4095

a.u. and the indicator in the bar is at 350 a.u.

To investigate the non-perovskite phase/impurities in the CsPbI2Br films, EDS mapping is carried out to analyse the chemical composition of the different grains. As shown in

Figure 5.6, the brighter impurity grains in the SEM image of CsBr-rich film contain more

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Cs element and less Pb element. EDS point analysis on the grains as shown in Figure

5.7a confirms that the impurities grains (position 1 and 2) are Cs-rich. The stoichiometrically balanced film has uniform distribution of Cs and Pb. This is confirmed by the point analysis as shown in Figure 5.7b. The impurity grains of PbI2- rich film are rich in Pb element and poor in Cs element (see Figure 5.6c and grains 3 and 4 in Figure 5.7c).

Figure 5.6. SEM images and the corresponding EDS mappings of Pb and Cs elements in (a) CsBr-rich, (b)

stoichiometrically balanced, and (c) PbI2-rich samples.

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Figure 5.7. EDS point analysis on the (a) CsBr-rich, (b) stoichiometrically balanced and (c) PbI2-rich

CsPbI2Br films (atomic ratio normalized to Pb element).

5.3.4 Discussion on the Improved Air Stability

Combining the SEM, PL imaging and the degradation images, it can be concluded that as the CsBr-to-PbI2 evaporation rate ratio increases for the CsPbI2Br film deposition, the grains become smaller while the air stability improves. This is consistent with the findings in previous work [6, 9], that reported higher number of smaller grains in inorganic Cs based perovskite result in more stable film with higher instances of black perovskite phase. A closer look at the XRD pattern of CsBr-rich film (Figure 5.2b and

Figure 5.8) shows that the (110) peak at 20.95° has a small split peak at 21.58° and the

(200) peak at 29.65° has a shoulder peak at 29.22°. These split peak and shoulder

107 peaks are not present in the XRD pattern of PbI2-rich and stoichiometrically balanced samples. This is consistent with the peak splitting for the (110) and (200) peaks observed by Eperon et al. in [6]. This suggests that the smaller grain size observed in the CsBr-rich film is caused by the crystal lattice strain which contributes to better air stability in the CsBr-rich film.

Figure 5.8. XRD pattern of the fresh CsBr-rich and PbI2 rich CsPbI2Br films showing that the (110) peak at 20.95° has a small split peak at 21.58° and the (200) peak at 29.65° has a shoulder peak at 29.22° in

the CsBr-rich film, as indicated by “*”, no such peaks in the PbI2 rich films

5.3.5 Electrical Characteristics of the CsPbI2Br Solar Cells with

Different Stoichiometry

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To compare the effect of varying the stoichiometry of CsPbI2Br absorber on cell performance, planar devices of Glass/FTO/c-TiO2/CsPbI2Br/P3HT/Au are fabricated and encapsulated. The energy band diagram of this solar cell structure is shown in

Figure 5.9. Device performance for the 0.159cm2 and 1.2cm2 cells are shown in Table

5.2 and Figure 5.10. Although the CsBr-rich cells have respectable performances, they have poorer stabilised efficiencies. Cells made of stoichiometrically balanced CsPbI2Br film give the best performance with highest currents. On the other hand, PbI2-rich

CsPbI2Br cells produce the lowest currents. The poorer performance in the CsBr-rich and PbI2-rich devices are due to the presence of non-perovskite grains that do not contribute to the photovoltaic performance as reflected in the lower carrier lifetime measured by TRPL previously. Hysteresis exists in these devices as can be seen from the stabilized efficiencies which are generally lower than the efficiencies by reverse scan. The hysteresis may come from the architecture of planar device in which the interface between perovskite and TiO2 is not sufficient enough for the carrier extraction. Also, the film quality plays a role on the hysteresis that less defects in the crystalline structure result in less carrier traps and thus smaller hysteresis, which can be verified by the better performance of stoichiometrically balanced CsPbI2Br (better thin film quality from PL results) devices. All of the 1.2cm2 devices have lower fill factor due to higher resistance as the cell area increases. It can be concluded that the instability of the stoichiometrically balanced CsPbI2Br cells can be overcome by encapsulations as shown in Figure 5.11.

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Figure 5.9. Energy band diagram of the planar CsPbI2Br perovskite solar cell structure

Table 5.2. Electrical performance of CsBr-rich, stoichiometrically balanced and PbI2- rich CsPbI2Br devices.

Active area Jsc Voc FF Eff Stabilized Eff Sample [cm2] [mA/cm2] [mV] [%] [%] [%]

0.159 11.2 970 63 6.8 4.4 CsBr-rich 1.2 10.2 1043 57 6.1 4.2

Stoichiometrically 0.159 11.5 1005 67 7.7 6.7

balanced 1.2 11.5 1019 58 6.8 5.5

0.159 9.9 1002 69 6.8 5.4 PbI2-rich 1.2 9.4 1004 60 5.7 4.1

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Figure 5.10. (a) Light current density-voltage (JV )curves under reverse scan (VOC to JSC) at a scan speed

of 40mV/s and (b) stabilized efficiencies of 0.159cm2 (solid lines) and 1.2cm2 (dash lines) CsBr-rich

(red), stoichiometrically balanced (black) and PbI2-rich (blue) CsPbI2Br planar cells. 111

Figure 5.11. Photos showing the lack of degradation for encapsulated stoichiometrically balanced

CsPbI2Br films upon exposure to air. It can be seen that with encapsulation the colour of the

perovskite film remains the same even after 6 months.

Damp heat test has been performed on the stoichiometrically balanced encapsulated

CsPbI2Br cell. Results are shown in Figure 5.12. It is found that CsPbI2Br cells can retain

90% of their initial performance after 240 hours damp heat testing at 85 and 85% relative humidity. The drop in fill factor is the main reason for performance℃ drop, although the VOC has a slightly increasing trend over time. Further work will be carried

112 out to fully understand the cause for VOC increase and FF decrease.

Figure 5.12. Results of damp heat test on the encapsulated stoichiometrically balanced CsPbI2Br cell.

Test condition: 85 and 85% relative humidity.

℃ 5.4 Conclusion

We have investigated the effect of varying the stoichiometry of CsPbI2Br on material property and photovoltaic performance. By varying the deposition rate ratio of CsBr to

PbI2 during dual source thermal co-evaporation, CsPbI2Br films that are CsBr-rich, stoichiometrically balanced or PbI2-rich films can be produced. While the absorption onset remains the same for these films, the grain property, level of impurities, elemental distribution, morphology and stability of these films are very different as the stoichiometry is varied. The presence of impurities in the CsBr-rich and PbI2-rich films result in lower photoluminescence and PL lifetimes and consequently lower cell

113 performance. In terms of air stability, CsBr-rich sample is the best likely due to smaller grain size that result in crystal lattice strain hampering degradation. However, the less stable stoichiometrically balanced CsPbI2Br devices once encapsulated produce the best cell performance compared to CsBr and PbI2 rich CsPbI2Br devices, with a PCE of

7.7% under reverse scan and a stabilized PCE of 6.7% for a 0.159cm2 device. A PCE at

6.8% under reverse scan and stabilized PCE at 5.5% are also obtained for the 1.2cm2 device. This is the largest inorganic caesium lead halide perovskite solar cell reported to date. This work enriches the understanding of air stability of CsPbI2Br perovskite film and the effectiveness of stoichiometry engineering on film property. CsPbI2Br perovskite thin film with a bandgap of 1.9 eV is more suitable for sunlight harvesting in terms of photovoltaic application and tandem solar cells than CsPbIBr2 with a bandgap of 2.05 eV. An efficiency of as high as 24% could possibly be achieved by a CsPbI2Br/Si tandem solar cell when the top CsPbI2Br cell and the tandem structure are optimized with a Jsc of 17mA/cm2, Voc of 1.8 V and a fill factor of 80%. This tandem cell will have the advantage of low cost and light weight because the thickness of the bottom Si cell can be reduced due to the reduced current required.

References:

1. Ma, Q., et al., Hole transport layer free inorganic CsPbIBr2 perovskite solar cell by dual source thermal evaporation. Advanced Energy Materials, 2016. 6(7). 2. Bremner, S., et al., Optimum band gap combinations to make best use of new photovoltaic materials. Solar Energy, 2016. 135: p. 750-757. 3. Sutton, R.J., et al., Bandgap‐Tunable Cesium Lead Halide Perovskites with High Thermal Stability for Efficient Solar Cells. Advanced Energy Materials, 2016. 4. Beal, R.E., et al., Cesium lead halide perovskites with improved stability for tandem solar cells. The journal of physical chemistry letters, 2016. 7(5): p. 746-751.

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5. Chen, C.Y., et al., All‐Vacuum‐Deposited Stoichiometrically Balanced Inorganic Cesium Lead Halide Perovskite Solar Cells with Stabilized Efficiency Exceeding 11%. Advanced Materials, 2017. 29(12). 6. Eperon, G.E., et al., Inorganic caesium lead iodide perovskite solar cells. Journal of Materials Chemistry A, 2015. 3(39): p. 19688-19695. 7. Swarnkar, A., et al., Quantum dot–induced phase stabilization of α-CsPbI3 perovskite for high-efficiency photovoltaics. Science, 2016. 354(6308): p. 92-95. 8. Niezgoda, J.S., et al., Improved Charge Collection in Highly Efficient CsPbBrI2 Solar Cells with Light-Induced Dealloying. ACS Energy Letters, 2017. 2(5): p. 1043-1049. 9. Protesescu, L., et al., Nanocrystals of Cesium Lead Halide Perovskites (CsPbX3, X = Cl, Br, and I): Novel Optoelectronic Materials Showing Bright Emission with Wide Color Gamut. Nano Letters, 2015. 15(6): p. 3692-3696. 10. Yonezawa, K., et al., Annealing effects on CsPbI3-based planar heterojunction perovskite solar cells formed by vacuum deposition method. Japanese Journal of Applied Physics, 2017. 56(4S): p. 04CS11. 11. Wen, X., et al., Ultrafast electron transfer in the nanocomposite of the graphene oxide– Au nanocluster with graphene oxide as a donor. Journal of Materials Chemistry C, 2014. 2(19): p. 3826-3834. 12. Chen, S., et al., Spatial Distribution of Lead Iodide and Local Passivation on Organo- Lead Halide Perovskite. ACS Applied Materials & Interfaces, 2017. 9(7): p. 6072-6078.

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Chapter 6 Conclusion and Future Work

Emerging as a new class of photovoltaic semiconductor material, perovskites have attracted attention in the photovoltaic research community in recent years due to their rapid efficiency improvement, low embedded energy and their suitability as a top cell in tandem solar cells. The commonly used solution fabrication processes for perovskites, however, can make the large scale production of perovskite solar cells challenging especially when high volumes of solvents are required. This thesis develops scalable vapour based deposition methods to fabricate perovskite solar cells and thermally more stable perovskite materials with suitable bandgaps for tandem applications. These aims are well accomplished by the results in Chapters 3, 4 and 5, as will be summarised in this chapter. Based on these results, future work is planned towards the scale-up production of perovskite solar cells and perovskite/silicon tandem solar cells. 6.1 Conclusion

Chapter 3 presented the fabrication of MAPbIXCl3-x perovskite by vapour based deposition methods. Dual source thermal co-evaporation was introduced for the deposition of MAPbIXCl3-x. However, the accurate control of the evaporation rate of

PbCl2 and MAI at the same time was not successful in our evaporator, because the evaporation of MAI is not directional and inevitably affects the observed rate of PbCl2 deposition. To solve this problem, a sequential evaporation method in which the PbCl2 and MAI were evaporated onto the substrates separately was proposed. Solar cells fabricated by this method had PCE up to 6.3% according to reverse scanned J-V curve.

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Another problem associated with MAI is the condensation of the vapour on pumping system corroding the turbomolecular pump blades. To address this issue, a novel vapour-assisted (atmospheric pressure) evaporation method was proposed to fabricate MAPbIXCl3-x perovskite thin films. In this method, the PbCl2 was firstly coated onto the substrate by vacuum-thermal evaporation. A solar cell with a PCE of 11.0% by reverse scan was fabricated by this method. MAPbIXCl3-x perovskite devices with active area ~3 cm2 were also demonstrated by this vapour-assisted (atmospheric pressure) evaporation method.

Chapter 4 focused on the fabrication of CsPbIBr2 perovskite thin films and solar cells to explore stable perovskite materials with bandgaps suitable for application in tandem solar cells. Among the three vapour-based deposition methods described in Chapter 3, simultaneous dual source thermal evaporation of CsI and PbBr2 was employed for the deposition of CsPbIBr2 because inorganic CsI is difficult to vaporise at low temperature by the atmospheric pressure evaporation method. Co-evaporation of CsI and PbBr2 is simpler and more efficient than sequential evaporation of these precursors. The

CsPbIBr2 films fabricated by this method have an optical bandgap of 2.05 eV and a stable PL emission at 2.00 eV making them suitable for a top cell in 3-junction silicon based tandem cell. Planar HTM-free perovskite solar cells based on CsPbIBr2 were fabricated and the best performing device achieved a conversion efficiency of 4.7% under reverse scan and 3.7% under forward scan, which is the highest PCE reported for a cell deposited by dual source co-evaporation without the use of a hole transport material. The CsPbIBr2 films are more thermally stable under N2 and room ambient conditions when compared with MA based perovskites. A device with an active area of

1.5 cm2 was also fabricated, showing its potential in scalable production. The 117 deposition of CsPbIBr2 thin films on textured silicon substrates and large size (6 inch) silicon wafers were also demonstrated in this chapter, showing the adaptability of dual source co-evaporation towards the fabrication of silicon based tandem cells at industrial scale.

Chapter 5 demonstrated the fabrication of CsPbI2Br by the dual source thermal co- evaporation method in which the bandgap was lowered from 2.05 eV to 1.9 eV by increasing the iodine-to-bromine ratio from a composition of CsPbIBr2 to CsPbI2Br. This demonstrates the bandgap tuneability of perovskite materials provided by engineering the halide composition. The effects of varying the stoichiometry of CsPbI2Br on material property and photovoltaic performance were investigated and particularly the air stability was addressed. By varying the deposition rate ratio of CsBr to PbI2 during dual source thermal co-evaporation, CsPbI2Br films that were CsBr-rich, stoichiometrically balanced or PbI2-rich were produced. The properties of these three

CsPbI2Br films were compared. The CsBr-rich and PbI2-rich films had lower photoluminescence and PL lifetimes and thus lower cell performance, while CsBr-rich samples were the most stable in air possibly due to smaller grain size. When encapsulated, however, the less stable stoichiometrically balanced CsPbI2Br devices produced the best cell performance compared to either the CsBr- or PbI2-rich CsPbI2Br devices, with a PCE of 7.7% under reverse scan and a stabilized PCE of 6.7% for a 0.159 cm2 device. Larger (1.2 cm2) devices also performed well with 6.8% PCE demonstrated under reverse scan and a stabilized PCE of 5.5%. The work in this chapter helps the understanding of air stability of CsPbI2Br perovskite films and the stoichiometry engineering on film properties. CsPbI2Br perovskite thin films (bandgap 1.9 eV) are

118 more preferred in solar photovoltaic applications than CsPbIBr2 (bandgap 2.05 eV), as shown in the tandem cell designs in Figure 1.3. 6.2 Future Work

Based on the results obtained, several aspects of perovskite cell technology could be explored in the future, such as; improving on the performance of the current devices, addressing the moisture and light stability, fabricating larger devices and modules with industrial production equipment and fabricating tandem solar cells.

• For organic metal halide perovskite MAPbIXCl3-x, the efficiency of 11% and

relatively low Voc of 0.86 V can be improved by using spray pyrolysis deposited

compact TiO2 instead of spin-coated compact TiO2 as the spray pyrolysis

method can produce a more uniform compact TiO2 layer with better film

quality. The hysteresis can be reduced by interface modification, particularly

the interface between the perovskite and the n-type TiO2 layers. This can be

achieved by soaking the deposited compact TiO2 into a TiCl4 solution to

passivate the surface and reduce the defects. Also, a thin layer of mesoporous

TiO2 (100~200nm) can be inserted between the compact TiO2 and perovskite

layer to possibly reduce the hysteresis. In order to fabricate perovskite devices

on 6-inch silicon wafers, uniform heating source is required for the vapour-

assisted (atmospheric pressure) evaporation method for the deposition of MAI.

In addition, large area HTM deposition techniques such as doctor blading or

printing can be developed as the current spin-coating method is not suitable

for large area deposition. Alternatively, inorganic HTMs that are compatible

with the vacuum evaporation technique can be used. To fabricate large area

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devices with comparable fill factors to the small devices, the large series

resistance contribution from the FTO or ITO needs to be reduced, for example,

by using a front metal grid design[1].

• The performance of caesium lead halide perovskite solar cells can be improved

further. For both CsPbIBr2 and CsPbI2Br, Jsc can be improved by increasing the

film thickness from ~200 nm to 300-400 nm to maximally absorb the incident

light without voltage loss. In this work, the open circuit voltages of 0.959 V and

1.005 V for CsPbIBr2 and CsPbI2Br solar cells, respectively, are relatively low

compared to the corresponding bandgaps of these materials; 2.05 eV and 1.9

eV. The loss in potential or voltage deficit (the difference in energy between

the bandgap and the open circuit voltage) of the best reported organic-

inorganic perovskite solar cell was less than 0.4 eV as opposed to 1.1 eV and 0.9

for the CsPbIBr2 and CsPbI2Br solar cells. The low open circuit voltage of the

planar HTM-free CsPbIBr2 solar cell might be improved by simply inserting an

HTM layer between the perovskite layer and the Au contact. For CsPbI2Br solar

cells, the low open circuit voltage could be explained by the improper band

alignments at the interfaces between the perovskite layer, ETL and HTM, as

shown in Figure 5.9. The VOC can be possibly improved by replacing the ETL and

HTM materials with ones having more suitable conduction band (Ec) and

valence band (Ev) to reduce the cliff of Ec between perovskite and ETL and Ev

between perovskite and HTM, as illustrated in Figure 6.1. The possible ETL

substitute material could be PCBM and HTM substitute material could be NiOx,

CuSCN or CuI. Furthermore, the performance could be improved by optimizing

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the deposition conditions and thus producing defect free Cs perovskite thin

films.

Figure 6.1. Modification of the energy band positions of current ETL and HTM layer. The black solid lines show the current band positions and the dashed lines indicate the band positions of an ETL and

HTM suitable for the achieving higher Voc.

• Towards the commercialization of perovskite solar cells, the thermal stability,

moisture stability and light stability have to be addressed. The Cs lead halide

perovskites have shown improved thermal stability compared to the organic

perovskite counterparts. Stability in a moist environment can be improved

markedly by encapsulating the devices from the atmosphere, or exploring new

perovskite materials which are stable in moist environments. Not much work

has been done to investigate the light stability of inorganic perovskite materials

warranting detail study.

• In this thesis, CsPbIBr2 and CsPbI2Br perovskite with a bandgap of 2.05 eV and

1.9 eV respectively which are suitable to work as a top cell in a 3-junction 121

tandem cell on a 1.1 eV silicon bottom cell have been systematically

investigated. The next step would be the fabrication of 2 or 3-junction tandem

solar cells based on CsPbIBr2 and CsPbI2Br. For the 3-junction tandem solar cell

the middle cell could be CIGS solar cells or another kind of perovskite solar cell

with a bandgap around 1.5 eV. 2-junction tandem cells with a structure of 1.7

eV top cell on 1.1 eV silicon bottom cell are also promising to improve the

efficiencies of solar cells and drive down the production cost. CsPbI3 with a

bandgap of 1.73 eV could be a good option except the structural instability in

cubic or tetragonal phase (only cubic and tetragonal phase of perovskites are of

photovoltaic use). Based on the understanding of Chapter 5, this instability

could possibly be solved by the stoichiometry control or the grain size control

as smaller grains/crystallites result in better thin film stability in air.

Alternatively, one can replace Cs by or incorporate other inorganic cations such

as Tl and K, or explore other stable perovskites with organic cations.

References:

1. Kim, J., et al., Overcoming the Challenges of Large-Area High-Efficiency Perovskite Solar Cells. ACS Energy Letters, 2017: p. 1978-1984.

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