UNIVERSITY of CALIFORNIA Los Angeles Interface Design of Halide

UNIVERSITY OF CALIFORNIA

Los Angeles

Interface Design of Halide Perovskite and Metal Oxide Semiconductor

for Optoelectronic Devices

A dissertation submitted in partial satisfaction of the

requirements for the degree Doctor of Philosophy

in Materials Science and Engineering

Pengyu Sun

2018

Pengyu Sun

2018

ABSTRACT OF THE DISSERTATION

Interface Design of Halide Perovskite and Metal Oxide Semiconductor

for Optoelectronic Devices

Pengyu Sun

Doctor of Philosophy in Materials Science and Engineering

University of California, Los Angeles, 2018

Professor Yang Yang, Co-Chair

Professor Eli Yablonovitch, Co-Chair

Thin film optoelectronics, including solar cells and thin film transistors, has been investigated extensively in recent years. Thin film solar cells have attracted substantial attention due to their low cost, ease of fabrication, flexible nature and potential to achieve high power conversion efficiency. We witnessed a rapid rise of halide perovskite solar cells with power conversion efficiency reaching 22.7% within a few years, exceeding Cu(InxGa1-x)Se2 and CdTe.

Perovskite materials have shown unique properties of defect tolerance and strong optical absorption, which makes the material a perfect candidate for photovoltaic applications. On the other hand, thin film transistors (TFT) have evolved at a fast pace with rapid industrialization.

InGaZnO (IGZO) material has shown a quick adoption in flat panel display (FPD) backplanes

ii since discovery in 2004. IGZO’s high mobility, good stability and uniformity enable the material to compete with traditional amorphous silicon (a-Si) or polycrystalline silicon (p-Si) technologies.

This thesis focuses on role of interface and fundamentals of optoelectronics for both perovskite solar cells and IGZO TFTs. In Chapter 2, I will discuss the theoretical background of optoelectronics where photoluminescence quantum efficiency (PLQE) plays an important role in perovskite thin film solar cells. I will also review the current progress in the field where people trying to optimize PLQE in perovskite material and future directions. In Chapter 3, I will demonstrate the impact of bottom contact on PLQE of CH3NH3PbI3 (MAPbI3) perovskite and its correlation with open circuit voltage of perovskite solar cells. It has been found that perovskite material shows improved crystal growth on NiOx bottom contact over PEDOT:PSS coated substrates. In Chapter 4, I will demonstrate the growth of formamidinium lead iodide (FAPbI3) single crystals with diameter up to 5 mm. The material shows phase transformation at elevated temperatures and a lower band gap compared with MAPbI3, making it more suitable for photovoltaics. FAPbI3 also shows high mobility compared with MAPbI3. In Chapter 5, I will demonstrate admittance spectroscopy for identification of defect states in MAPbI3 perovskite solar cells. Lastly, in Chapter 6, I will demonstrate stability enhancement for solution processed IGZO

TFTs with top surface passivation by parylene coating. Solution processed IGZO material shows comparable performance metrics with traditional sputtered IGZO with respect to mobility, on-off ratio, subthreshold swing (SS), and threshold voltage. However, due to the nature of sol-gel process, the thickness of solution processed IGZO is limited and the material is prone to surface environment. I will demonstrate an effective way of passivation on solution processed IGZO with

Parylene to minimize surface effect and the passivated TFT shows positive bias stress below 0.5

V drift in turn-on voltage.

iii

The dissertation of Pengyu Sun is approved.

Mark S. Goorsky

Eli Yablonovitch, Committee Co-Chair

Yang Yang, Committee Co-Chair

University of California, Los Angeles

2018

Table of Contents

Chapter 1: Introduction ...... 1 1.1 Basics of Solar Cells ...... 2 1.2 Lead halide perovskites ...... 6 1.3 Basics of thin film transistors ...... 9 1.4 Amorphous oxide semiconductor (AOS) ...... 11 References ...... 12

Chapter 2: Detailed balance and theoretical limit of perovskite solar cells ...... 15 2.1 Introduction ...... 15 2.2 Detailed balance and Shockley-Queisser limit...... 15 2.3 Theoretical limit of perovskite solar cells ...... 17 References ...... 18

Chapter 3: High photoluminescence quantum efficiency of CH3NH3PbI3 perovskite on NiOx hole contact...... 21 3.1 Introduction ...... 21 3.2 Experimental details ...... 24 3.2.1 Perovskite thin film and device fabrication ...... 24 3.2.2 Characterization and measurement ...... 24 3.3 Results and discussion ...... 25 3.4 Conclusion ...... 33 References ...... 33

Chapter 4: Optoelectronics of Single Crystal Formamidinium Lead Iodide (FAPbI3) .... 37 4.1 Introduction ...... 37 4.2 Experiment Details ...... 39

4.2.1 Synthesis of single crystal FAPbI3 ...... 39

4.2.2 Synthesis of FAPbI3 thin films and permittivity measurement ...... 39 4.2.3 Measurement and characterization ...... 40 v

4.3 Results and discussion ...... 41 4.4 Conclusion ...... 56 References ...... 57

Chapter 5: The identification and characterization of defect states in hybrid organic- inorganic perovskite photovoltaics ...... 62 5.1 Introduction ...... 62 5.2 Results and discussion ...... 63 5.3 Conclusions ...... 72 References ...... 73

Chapter 6: Enhanced Stability of Solution Processed InGaZnO Thin-Film Transistor with Parylene Passivation...... 76 6.1 Introduction ...... 76 6.2 Experimental details ...... 77 6.3 Results and discussion ...... 78 6.4 Conclusion ...... 83 References ...... 83

List of Figures

Figure 1.1 (a) Schematic diagram of p-n junction solar cells; (b) schematic diagram of p-in junction solar cells; (c) schematic diagram of charge transport in p-n junction solar cells; (d) schematic diagram of energy band diagram and charge separation in p-n junction solar cells...... 4

Figure 1.2 An example of current-voltage (I-V) curve. ISC = JSC × device area...... 5

Figure 1.3 Crystal structure of the halide perovskite material: A site is typically CH3NH3I

(MAI), NH2CHNH2I (FAI), Cs; B site is commonly Pb, Sn; X site could be Cl, Br or I...... 6

Figure 1.4 Three typical device structures of perovskite solar cells: (a) mesoporous, (b) regular planar structure, and (c) inverted planar structure ...... 8

Figure 1.5 Characteristic (a) output (b) transfer I-V curves of TFT devices...... 9

Figure 1.6 Schematic cross-sections of the four principle thin-film transistor (TFT) structures.

The carrier channel is schematically shown in red. (a) Bottom-gate (inverted) staggered TFT. (b)

Bottom- gate (inverted) coplanar TFT. (c) Top-gate staggered TFT. (d) Top-gate coplanar TFT.

...... 10

Figure 1.7 Conduction mechanisms of crystalline oxide semiconductor, crystalline silicon, amorphous oxide semiconductor and amorphous silicon...... 12

Figure 3.1 Steady state photoluminescence of perovskite films grown on different substrates (a)

PL spectra under one-sun illumination. (b) Log-log plot of measured PL intensity vs illumination power density. (c) Pump power dependence of internal PLQE ...... 27

Figure 3.2 Time Resolved PL decay and bi-exponential fitting of perovskite films on (a) Glass

(b) NiOx (c) PEDOT:PSS substrates, measured at spectrum peak wavelength and with three different pump fluences...... 29

vii

Figure 3.3 (a) Current density versus voltage (J-V) of solar cells based on NiOx and

PEDOT:PSS substrates. (b) EQE of corresponding devices ...... 31

Figure 3.4 SEM image of perovskite films on (a)glass (b)NiOx (c)PEDOT:PSS. (d) XRD patterns of the corresponding films ...... 32

Figure 4.1 Crystal structures of a) α-phase and b) δ-phase single crystal FAPbI3. The inset of a) and b) are images of α-phase and δ-phase single crystal, respectively. c) TGA-DSC curves of single crystal FAPbI3. d) and e) are experimental and calculated powder XRD patterns of α- phase single crystal FAPbI3 and δ-phase FAPbI3, respectively...... 44

Figure 4.2 δ-phase to α-phase FAPbI3 single crystal phase transition video capture images of annealing time interval 20 sec. a) 0 s, b) 20 s, c) 40 s, d) 60 s, e) 80 s, f) 100 s, g) 120 s and h)

140 s...... 44

Figure 4.3 Raman spectra of single crystal FAPbI3 with α-phase (black) and δ-phase (red)...... 45

Figure 4.4 Decay traces of time-dependent PL on a α-phase single crystal FAPbI3 at λ = 820 nm with the biexponential fitting showing a fast (τ1 = 32 ns, green) and a slow transient (τ2 = 484 ns, blue). The inset shows the emission PL peak of α-phase single crystal FAPbI3...... 46

Figure 4.5 PL time decay traces on the FAPbI3 films with the thickness of 300 nm at λ=805 nm with the biexponential fitting showing a fast (τ1=29 ns) and a slow transient (τ2=227 ns). The inset shows the PL spectrum of FAPbI3 thin film...... 47

Figure 4.6 UV-Vis absorption spectra of α-FAPbI3 single crystal and film...... 47

Figure 4.7 UV-Vis absorption spectrum of δ-FAPbI3. The inset shows the PL spectrum of δ-

FAPbI3 single crystal...... 48

Figure 4.8 Dark current-voltage curve of α-phase single crystal FAPbI3 for space charge limited current analysis...... 50

viii

Figure 4.9 Time-of-flight traces showing the transient current under various reverse biases in a bilogarithmic plot. Inset shows the charge transit time versus the reciprocal of bias; the solid line is a linear fit to the data...... 50

Figure 4.10 Dark current-voltage curve of δ-phase FAPbI3 single crystla for space charge limited current (SCLC) analysis...... 52

Figure 4.11 Dark current-voltage curve of α-phase FAPbI3 thin film for space charge limited current analysis...... 52

Figure 4.12 Schematic illustration of the photodetector fabricated on free-standing single crystal

FAPbI3...... 54

Figure 4.13 Device configuration and characteristics of the photodetector based on single crystal

FAPbI3. a) Photocurrent and photoresponsivity of devices for incident power densities. b)

Transient photocurrent response at a pulse frequency of 1 Hz...... 54

Figure 4.14 Photocurrent response during on–off illumination switching: rise and decay (inset) curves...... 56

Figure 5.1 A schematic band diagram of p-type semiconductor junction with a single trap level

Ed and two measurement energies Eω1 and Eω2 ...... 64

Figure 5.2 (a) Surface SEM image of an annealed perovskite layer, (b) cross-sectional SEM image of a perovskite cell and (c) current density-voltage characteristics ...... 65

Figure 5.3 Admittance spectroscopy: capacitance and conductance for different device architectures ...... 67

Figure 5.4 (a) Temperature dependence of capacitance, (b) the derivative of the capacitance spectra which shows the characteristic frequency at each temperature (c) Arrhenius plot of the

ix characteristic frequencies to extract the defect activation energy and (d) the defect energy distribution of CH3NH3PbI3 perovskite ...... 70

Figure 6.1 Transfer characteristics and saturation mobility of as fabricated IGZO TFT...... 79

Figure 6.2 Transfer characteristics of solution processed IGZO TFTs before and after 100 s,

10000 s stress test a) without, b) with parylene passivation. Insets show device structure...... 80

Figure 6.3 a) Cross-sectional schematic of AMOLED display driven by IGZO TFTs. b)

Schematic of one-transistor, one-diode circuit. c) Mini AMOLED display fabricated with solution processed IGZO TFTs. b) Line scan of the display ...... 82

List of Tables

Table 1 Summary of solar cell performances ...... 30

Acknowledgement

The dissertation and the work behind were completed with the selfless support and guidance of many. I would like to express my deep appreciation to them and my sincere wish that our friendships will continue to thrive even though the dissertation has come to a conclusion.

Firstly, I would like to express my sincerest gratitude to Prof. Yang Yang, my Ph.D. advisor, for giving me the opportunity to participate in the distinguished research of his group. His research guideline of “always take the challenge” has influenced me the greatest. During the years,

I have never been short of encouragement, trust, and inspiration from my advisor, for which I thank him the most.

I would also like to thank Prof. Eli Yablonovitch for being my co-advisor. His depth of knowledge and attitude towards scientific research inspired and motivated me along my Ph.D. journey. I also wish to thank my committee members, Prof. Dwight C. Streit and Prof. Mark S.

Goorsky for their mentorship.

I will not forget to thank my dearest senior partners in Prof. Yang’s Lab during my Ph.D. career: Dr. You Seung Rim, Dr. Qi Chen, Dr. Huanping Zhou, Dr. En-Ping Yao, Dr. Qifeng Han,

Dr. Lijian Zuo, Dr. Bowen Zhu, Dr. Wenchao Huang, Dr. Jin-Wook Lee, Dr. Min Cai, Dr. Shirong

Lu, Dr. Letian Dou, Dr. Samuel Duan, Dr. Jing Gao, Dr. Tze-Bin Song, Dr. Yang (Michael) Yang,

Dr. Chun-Chao Chen, Dr. Shenglin Ye and Dr. Wei-Hsuan Chang.

I am glad to have the chance to share the joys and hard times with my fellow colleges: Dr.

Meng Lei, Dr. Huajun Chen, Song Luo, Yao-Tsung Hsieh, Dr. Sanghoon Bae, Hongxiang Zhao,

Tony Chung, Eric Young, Hsin-Hua Wang, Shiqi Dong, Sheng-Yung Chang, Nicholas De Marco and Zhanlue Yang. I wish them the best in their life and career. xii

Lastly, I would like to thank my parents, my fiancée, and my other family members. I would not have reached this step without their belief in me.

xiii

VITA

2009-2013 B.Sci. Optical Science and Engineering, Fudan University, Shanghai, China

2013-2018 Graduate Student Researcher, Department of Materials Science and Engineering,

University of California, Los Angeles, CA, USA

Publications

1. En-Ping Yao, Zhanlue Yang, Lei Meng, Pengyu Sun, Shiqi Dong, Ye Yang, Yang Yang, High-

Brightness Blue and White LEDs based on Inorganic Perovskite Nanocrystals and their

Composites, Advanced Materials (2017).

2. Pengyu Sun, Bowen Zhu, Le Cai, Guangwei Xu, Suhui Lee, Jin Jang and Yang Yang,

Enhanced Stability of Solution Processed InGaZnO Thin-Film Transistor with Parylene

Passivation for OLED Display, IMID (2018)

3. Xuxia Shai, Jinsong Wang, Pengyu Sun, Wenchao Huang, Peizhe Liao, Feng Cheng, Bowen

Zhu, Sheng-Yung Chang, En-Ping Yao, Yan Shen, Ling Miao, Yang Yang, Mingkui Wang,

Achieving ordered and stable binary metal perovskite via strain engineering, Nano Energy

(2018)

4. Xiaojuan Zhao, Leiming Tao, Hao Li, Wenchao Huang, Pengyu Sun, Jun Liu, Shuangshuang

Liu, Qiang Sun, Zhifang Cui, Lijie Sun, Yan Shen, Yang Yang, Mingkui Wang, Efficient

Planar Perovskite Solar Cells with Improved Fill Factor via Interface Engineering with

Graphene, Nano Letters (2018)

5. Sheng‐Yung Chang, Yu‐Che Lin, Pengyu Sun, Yao‐Tsung Hsieh, Lei Meng, Sang‐Hoon Bae,

Yu‐Wei Su, Wenchao Huang, Chenhui Zhu, Gang Li, Kung‐Hwa Wei, Yang Yang, High‐

xiv

Efficiency Organic Tandem Solar Cells With Effective Transition Metal Chelates

Interconnecting Layer, Solar RRL (2018)

6. Shiqi Dong, Yongsheng Liu, Ziruo Hong, Enping Yao, Pengyu Sun, Lei Meng, Yuze Lin,

Jinsong Huang, Gang Li, Yang Yang, Unraveling the High Open Circuit Voltage and High

Performance of Integrated Perovskite/Organic Bulk-Heterojunction Solar Cells, Nano Letters

(2017)

7. Xuxia Shai, Lijian Zuo, Pengyu Sun, Peizhe Liao, Wenchao Huang, En-Ping Yao, Hao Li,

Shuangshuang Liu, Yan Shen, Yang Yang, Mingkui Wang, Efficient planar perovskite solar

cells using halide Sr-substituted Pb perovskite, Nano Energy (2017)

8. Lei Meng, En‐Ping Yao, Ziruo Hong, Huajun Chen, Pengyu Sun, Zhanlue Yang, Gang Li,

Yang Yang, Pure Formamidinium‐Based Perovskite Light‐Emitting Diodes with High

Efficiency and Low Driving Voltage, Advanced Materials (2017)

9. Qifeng Han, Sang‐Hoon Bae, Pengyu Sun, Yao‐Tsung Hsieh, Yang Michael Yang, You Seung

Rim, Hongxiang Zhao, Qi Chen, Wangzhou Shi, Gang Li, Yang Yang, Single crystal

formamidinium lead iodide (FAPbI3): Insight into the structural, optical, and electrical

properties, Advanced Materials (2017)

10. Nicholas De Marco, Huanping Zhou, Qi Chen, Pengyu Sun, Zonghao Liu, Lei Meng, En-Ping

Yao, Yongsheng Liu, Andy Schiffer, Yang Yang, Guanidinium: a route to enhanced carrier

lifetime and open-circuit voltage in hybrid perovskite solar cells, Nano Letters (2016)

11. Hsin-Sheng Duan, Huanping Zhou, Qi Chen, Pengyu Sun, Song Luo, Tze-Bin Song, Brion

Bob, Yang Yang, The identification and characterization of defect states in hybrid organic–

inorganic perovskite photovoltaics, Physical chemistry chemical physics (2015)

Chapter 1: Introduction

Thin film electronics, including photovoltaics and transistors, plays an important role in modern industry. Currently, world energy consumption relies mainly on fossil fuels. Among various types of renewable energy sources, solar energy is one of the most promising candidates owing to its ubiquitous availability. Photovoltaic (PV) technologies that directly convert sunlight energy to electricity is one of the most promising long-term clean energy solutions.

The most mature and well-adopted PV cell is based on crystalline silicon, which occupies approximately 85% of the PV market. Thin film PV devices have been extensively researched on to reduce production costs to compete with traditional energy sources. Additionally, thin absorber layers enable fabrication of lightweight, flexible solar modules. Organic inorganic lead halide perovskite material is a rising star in thin film PV technology field, with optimal optoelectronic properties comparable with traditional materials.

On the other hand, thin film transistors (TFT) are well-adopted in flat panel display (FPD) industry. The switching characteristics of TFTs make it ideal for display pixel control.

Traditionally, with liquid crystal display (LCD), amorphous silicon (a-Si) technology is widely used. In recent progress of FPD industry, higher resolution, lower noise and fast switching have become the prerequisite for high quality display. Hence, a-Si’s low charge carrier mobility and on- off ratio renders it outdated for current technology. Low temperature poly-silicon (LTPS) and

InGaZnO (IGZO) emerged as new materials for TFT applications thanks to their high mobility and better performance. Here I will introduce IGZO material for TFT applications.

In section 1.1, I will introduce the basics of solar cells, including their working mechanisms and fundamentals. Then, in section 1.2, I will discuss fundamentals of perovskite material. In section 1.3, fundamentals of TFT will be introduced. Following the discussion, in section 1.4, a brief introduction to amorphous oxide semiconductor will be discussed.

1.1 Basics of Solar Cells

Two types of solar cell structures are currently well adopted in different material systems.

The first type of structure consists of back electrode/p-type/n-type/top electrode as shown in Figure

1.1a. This structure was developed in the early stages of the solar cells research on Si-based devices.

Currently, Cu(In,Ga)(S,Se)2 (CIGS), CdTe, hybrid Si/CNT solar cells, organic photovoltaic (OPV) and GaAs systems adopt this structure. A typical structure is based on a lightly doped p-type semiconductor material combined with a heavily doped n-type semiconductor material. Due to the poor carrier mobility of holes, the wider depletion region should occur in the p-type material and assist charge separation and transport. Most photogenerated carriers can be extracted efficiently with proper control of the film thickness and doping concentration.

Another type of solar cell structure consists of back electrode/p-type/i-type/n-type/top electrode as shown in Figure 1.1b. This structure was designed as an improvement on p-n junction structure of Si solar cells. To control and introduce a specific concentration of dopants over an entire Si wafer requires long-time processing and high-temperature post-annealing.

Moreover, the extrinsic doping induces more defects in the materials, resulting in more recombination. The p-i-n solar cell structure overcomes these technical and fundamental issues and further improves device performance. Currently, amorphous Si (a-Si), crystalline Si solar cells,

2 heterojunction with intrinsic thin layer (HIT) and halide perovskite solar cells adopt this structure.

Generally, the thin and highly doped p-type and n-type layers are used on both sides of the intrinsic layer for charge separation and collection and the intrinsic layer is the main photon absorber.

Though Si and HIT solar cells have achieved high efficiencies compared to other materials, their cost-effectiveness and preparation process are still under investigation. As a result, halide perovskite materials emerge as potential candidates for thin-film solar cell applications.

The schematic diagram for the working principle of photogeneration and carrier transport is shown in Figure 1.1c and d. When the light shines through the transparent electrode, a photon will excite an electron from the valence band to the conduction band and generate a bonded electron-hole pair (exciton). Then, the exciton will be separated by the built-in potential through the p-n junction into free electron and hole. The free electron and hole will be transported to the electrodes and then to the external circuit. The built-in potential from the p-n junction is established as shown in Figure 1.1d. The conduction band bending and valence band bending in the depletion region are the main driving force for charge separation and collection.

In the other case, the p-i-n structure device has two interfaces, p-i and i-n. These two interfaces will form built-in potentials and these built-in potentials can span across the entire i- type layer by tuning the thickness of i-type layer or carrier concentrations of the semiconductor layers. Controlling the built-in potential and the depletion width is critical for efficient separation of excitons and achieving good charge transport and collection.

The performance of solar cell is evaluated based on its ability to produce electrical power under illumination. The power conversion efficiency (PCE) of a solar cell is determined by its simultaneously produced voltage and current, which make up its power output. The typical characterization of a solar cell is extracted by applying an external voltage and measuring its current output at various bias levels under illumination. The equation below describes PCE:

푉 × 퐽 × 퐹퐹 푃퐶퐸 = 표푐 푠푐 (1.1) 푃𝑖푛

The open-circuit voltage (Voc) is the voltage at the open-circuit point; the short-circuit current density (Jsc) is the current density at the short-circuit point and the maximum power point is the maximum value of current times voltage. Pin is the input light power density. All these parameters are indicated in a sample current-voltage (I-V) curve shown in Figure 1.2. The fill factor (FF) is the ratio between V × J at maximum power point and Voc × Jsc that could also be represented as the area ratio between the dark dotted line and red dotted line in the 4th quadrant. The cell’s J-V characteristics depend on the various material properties and interactions that take place within the cell and thus is a far more complex matter, which has been and will continue to be the topic of intensive research.

Figure 1.2 An example of current-voltage (I-V) curve. ISC = JSC × device area.

1.2 Lead halide perovskites

Inorganic-organic hybrid perovskite solar cells show a great potential as a low-cost and high performance photovoltaic technology, owing to the effective combination of solution- processing and its unique material properties.[2, 3] Organic halide perovskites have ABX3 crystal structures, where A, B, X are organic cation, metal cation and halide anion, respectively (Figure

1.3), where the bandgap can be tuned from the ultraviolet to infrared region through varying these components.[3, 4] This family of materials exhibits a myriad of properties ideal for PV such as high dual electron and hole mobility, large absorption coefficients resulting from s-p anti-bonding coupling, a favorable band gap, a strong defect tolerance and shallow point defects, benign grain boundary recombination effects and reduced surface recombination.[5] After 5 years efforts, the power conversion efficiency (PCE) of perovskite solar cells has risen from about 3% to over

22%,[2, 5-15] which is close to that of traditional solar cell technologies such as Si, GaAs, and higher than CIGS and CdTe.

Figure 1.3 Crystal structure of the halide perovskite material: A site is typically CH3NH3I (MAI), NH2CHNH2I (FAI), Cs; B site is commonly Pb, Sn; X site could be Cl, Br or I.

The hybrid perovskite solar cell was initially discovered in a dye sensitized solar cell

(DSSC) using a mesoporous TiO2 structure.[2] Miyasaka and co-workers were the first to utilize the perovskite (CH3NH3PbI3 and CH3NH3PbBr3) nanocrystal as absorbers in DSSC structure, achieving an efficiency of 3.8% in 2009.[2] Park et al. improved the PCE of CH3NH3PbI3 to 6.54% by optimizing the TiO2 surface based on the same configuration.[5] However, these devices degraded very quickly due to liquid hole transporting material used.[5] In 2012, Park and Graetzel et al reported a solid state perovskite solar cell by using the solid hole transport layer 2,2’,7,7,’- tetrakis-(N,Ndimethoxyphenyl-amine)-9,9’-spirobifuorene (Spiro-OMeTAD), showing an efficiency of 9.7% with improved stability.[7] Meanwhile, Snaith et al. achieved a device efficiency up to 10.9% by using an insulating mesoporous aluminum oxide (Al2O3) instead of TiO2 scaffold.[8] After that, several milestones in device performance have achieved.[9, 10, 13-15]

However, these devices are all based on the mesoporous structure, which needs a high temperature sintering that could increase the processing time and cost of cell production.

As revealed independently by Sum’s group and Snaith’s group in 2013, solution processed methylammonium-based perovskites exhibit extremely long and balanced charge carrier diffusion lengths (~100 nm for CH3NH3PbI3 and ~1000 nm for CH3NH3PbI3-xClx).[16, 17] Recently, single crystals of CH3NH3PbI3 were found to reach diffusion lengths larger than 175 μm, indicating that charge can be easily transported within the bulk perovskite materials.[18] These properties makes it possible to reduce the thickness of mesoporous scaffold and ideally avoid the use of mesoporous structure altogether without sacrificing efficiency. Further studies demonstrated that perovskites exhibit ambipolar behavior, meaning that these materials can transport both electrons and holes between the cell terminals.[16] All of these results indicated that a planar structure was feasible.

The first successful demonstration of the planar structure can be traced back to the

7 perovskite/fullerene structure reported by Guo and Chen,[19] showing a 4% efficiency. The low efficiency reported at that time was due to the inferior film quality and inadequate absorption of the perovskite film.[19] The breakthrough of the planar perovskite structure was obtained by using a dual source vapor deposition, providing dense and high quality perovskite films that achieved

15.4% efficiency.[20] Later, the efficiency of the planar structure was pushed over 19% through interface engineering.[12] These results showed that the planar structure could achieve similar device performance to the mesoporous structure. The evolution of the structure of perovskite is shown in Figure 1.4. The planar structure can be divided into two categories depending on which selective contact is used on the bottom, i. e. regular (n-i-p) and inverted (p-i-n). The device structures are shown in Figure 1.4 (b) and (c), respectively. For the regular n-i-p structure, the bottom selective contact is an electron transport layer (ETL). This structure has been extensively studied and could be traced back to dye sensitized solar cells. In the p-i-n structure, the bottom charge transport layer is p-type, which is derived from the organic solar cell, and consequently involves the utilization of commonly used charge transport layers in organic photovoltaics.

Figure 1.4 Three typical device structures of perovskite solar cells: (a) mesoporous, (b) regular planar structure, and (c) inverted planar structure

1.3 Basics of thin film transistors

Thin film transistors are now a fundamental component in flat panel display products. The commercial application of TFT started with hydrogenated amorphous silicon (a-Si:H). A great variety of semiconductor material candidates then emerged in the field. The operation principle of

TFTs is very similar to traditional MOSFETs. Figure 1.5 shows representative output and transfer

I-V characteristics of a typical TFT device. Compared with MOSFET, TFT shows inferior performance characteristics. It is reasonable due to the fact that active layer in TFT processes much more defects. Also, gate insulator in TFT is usually deposited with plasma enhanced chemical vapor deposition (PECVD). As a result, TFT shows lower on-current, higher off-current and larger sub-threshold swing.

Figure 1.5 Characteristic (a) output (b) transfer I-V curves of TFT devices.[21]

Figure 1.6 shows common TFT device structures. Common TFT device structure comprises four parts: gate, gate insulator, active layer (semiconductor) and source/drain. If active layer lies in the middle of gate and source/drain electrodes, it is staggered structure, otherwise it is

9 coplanar structure. If gate is on the top of the structure, it is regular structure, otherwise it is inverted structure. In the actual fabrication process of TFTs, these structures have their own features and a lot of engineering issues will be taken into consideration.

Figure 1.6 Schematic cross-sections of the four principle thin-film transistor (TFT) structures. The carrier channel is schematically shown in red. (a) Bottom-gate (inverted) staggered TFT. (b) Bottom- gate (inverted) coplanar TFT. (c) Top-gate staggered TFT. (d) Top-gate coplanar TFT.[22]

In terms of operation principle, TFTs are very similar with MOSFETs. We can use square law model to describe TFT behaviors. When 푉퐺푆 < 푉푡ℎ,

퐼퐷푆 = 0 (1.2)

When 푉퐺푆 > 푉푡ℎ and 푉퐷푆 < 푉퐺푆 − 푉푡ℎ,

푊 1 퐼 = 휇 퐶 ( ) [(푉 − 푉 )푉 − 푉2 ] (1.3) 퐷푆 퐹퐸 𝑖 퐿 퐺푆 푡ℎ 퐷푆 2 퐷푆

When 푉퐺푆 > 푉푡ℎ and 푉퐷푆 > 푉퐺푆 − 푉푡ℎ,

푊 퐼 = 휇 퐶 ( ) (푉 − 푉 )2 (1.4) 퐷푆 퐹퐸 𝑖 2퐿 퐺푆 푡ℎ

Eq. 1.3 describes linear region and Eq. 1.4 describes saturation region. 휇퐹퐸 is the field

푊 effect mobility. 퐶 is gate insulator capacitance. is width to length ratio of channel. 푉 is 𝑖 퐿 푡ℎ threshold voltage.

1.4 Amorphous oxide semiconductor (AOS)

Amorphous oxide semiconductors are a promising class of TFT material that have made an impressive progress over the past decade. AOS represents a new type of material that combines high optical transparency, high electron mobility and amorphous microstructure. The nature that

AOSs have no grain boundaries make the materials bypass the limitation of carrier mobility in polycrystalline materials.

The research on oxide semiconductors started with binary oxides such as ZnO, In2O3 or

SnO2. In 2003, Nomura et al. reported complex semiconductor composition, InGaZnO, which boosted oxide TFT’s performance with effective mobility of 80 cm2/Vs, and on-off ratio of

106.[23] The next year Nomura reported amorphous IGZO that still yielded good performance and demonstrated application on flexible substrates.[24] It opened up a broad research interest in AOSs and TFT performance continued to improve in the following years and then commercialized by a group of companies. Figure 1.7 demonstrates conduction mechanisms of different semiconductor materials. In AOSs, electrons are transported through overlapped s-orbitals in the metal ions. This fast transport mechanism ensures high mobility of electrons. Compared with that, in a-Si dangling bonds are present to hinder electron transport, which is the reason for low electron mobility.

Crystalline oxide semiconductor Crystalline silicon Metal ns-orbital ((n-1)d10ns0 (n≥5))

ⓔ- ⓔ-

Oxygen 2p-orbital Silicon sp3-orbital Amorphous oxide semiconductor Amorphous silicon

ⓔ-

Figure 1.7 Conduction mechanisms of crystalline oxide semiconductor, crystalline silicon, amorphous oxide semiconductor and amorphous silicon.

References

[1] http://en.wikipedia.org/wiki/Theory_of_solar_cells.

[2] 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.

[3] Stranks, S.D. and H.J. Snaith, Metal-halide perovskites for photovoltaic and light-emitting devices. Nature Nanotechnology, 2015. 10(5): p. 391-402.

[4] 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.

[5] Yin, W.-J., et al., Halide perovskite materials for solar cells: a theoretical review. Journal of

Materials Chemistry A, 2015. 3(17): p. 8926-8942.

[6] Im, J.-H., et al., 6.5% efficient perovskite quantum-dot-sensitized solar cell. Nanoscale, 2011.

3(10): p. 4088-4093.

[7] Kim, H.-S., et al., Lead Iodide Perovskite Sensitized All-Solid-State Submicron Thin Film

Mesoscopic Solar Cell with Efficiency Exceeding 9%. Scientific Reports, 2012. 2.

[8] Lee, M.M., et al., Efficient Hybrid Solar Cells Based on Meso-Superstructured Organometal

Halide Perovskites. Science, 2012. 338(6107): p. 643-647.

[9] Burschka, J., et al., Sequential deposition as a route to high-performance perovskite-sensitized solar cells. Nature, 2013. 499(7458): p. 316-319.

[10] Im, J.-H., et al., Growth of CH3NH3PbI3 cuboids with controlled size for high-efficiency perovskite solar cells. Nature Nanotechnology, 2014. 9(11): p. 927-932.

[11] You, J., et al., Moisture assisted perovskite film growth for high performance solar cells.

Applied Physics Letters, 2014. 105(18).

[12] Zhou, H., et al., Interface engineering of highly efficient perovskite solar cells. Science, 2014.

345(6196): p. 542-546.

[13] Jeon, N.J., et al., Solvent engineering for high-performance inorganic-organic hybrid perovskite solar cells. Nature Materials, 2014. 13(9): p. 897-903.

[14] Jeon, N.J., et al., Compositional engineering of perovskite materials for high-performance solar cells. Nature, 2015. 517(7535): p. 474-480.

[15] Yang, W.S., et al., High-performance photovoltaic perovskite layers fabricated through intramolecular exchange. Science, 2015. 348(6240): p. 1234-1237.

[16] Xing, G., et al., Long-Range Balanced Electron- and Hole-Transport Lengths in Organic-

Inorganic CH3NH3PbI3. Science, 2013. 342(6156): p. 344-347.

[17] Stranks, S.D., et al., Electron-Hole Diffusion Lengths Exceeding 1 Micrometer in an

Organometal Trihalide Perovskite Absorber. Science, 2013. 342(6156): p. 341-344.

[18] Dong, Q., et al., Electron-hole diffusion lengths > 175 mu m in solution-grown CH3NH3PbI3 single crystals. Science, 2015. 347(6225): p. 967-970.

[19] Jeng, J.-Y., et al., CH3NH3PbI3 Perovskite/Fullerene Planar-Heterojunction Hybrid Solar

Cells. Advanced Materials, 2013. 25(27): p. 3727-3732.

[20] 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.

[21] Itoh, Takeki, et al. "Fabrication of InGaN thin-film transistors using pulsed sputtering deposition." Scientific reports 6 (2016): 29500.

[22] Klauk, Hagen. "Organic thin-film transistors." Chemical Society Reviews 39.7 (2010): 2643-

2666.

[23] K. Nomura, H. Ohta, K. Ueda, T. Kamiya, M. Hirano, H. Hosono, Science, 300, 5623, 1269-

1272 (2003).

[24] K. Nomura, H. Ohta, A. Takagi, T. Kamiya, M. Hirano, H. Hosono, Nature 2004, 432, 488.

Chapter 2: Detailed balance and theoretical limit of perovskite solar cells

2.1 Introduction

Solar cells have limited power conversion efficiency based on thermodynamic principles.

The driving force of solar energy conversion is the temperature difference of the sun (T=5800 K) and surface of the earth (T=300 K). In a single junction solar cells, light absorbing material can only absorb photons in the solar spectrum above its band gap, and those photons will also go through thermalization processes. These loss mechanisms are inevitable, so they set the upper limit of power conversion efficiency in solar cells. This limit was carefully investigated through detailed balance by William Shockley and Hans-Joachim Queisser at Shockley Semiconductor.[1] The work was later known as Shockley-Queisser limit and became well accepted. In chapter 2.2, I will discuss detailed balance and how theoretical limit of a single junction solar cell is calculated. In chapter 2.3, I will discuss theoretical limit of perovskite solar cells.

2.2 Detailed balance and Shockley-Queisser limit

Shockley-Queisser limit makes four primary assumptions: 1) the probability of absorption of photons with energy above the band gap is unity, and it is zero for photons with energy below the band gap. 2) All photo-generated charge carriers thermalize to the band edges. 3) The collection probability for all generated charge carriers is unity. 4) Non-absorbed photons and thermalization loss are the only mechanisms of loss.

In order to calculate the maximum short circuit current Jsc, we need incoming photon flux

휙𝑖푛푐 and absorptance A(E), which is the percentage of absorption at energy E. According to

Shockley-Queisser limit, A(E) should be a step function where A(E) = 1 for E > Eg and A(E) = 0 for E < Eg. Then we have

∞ ∞ 퐽푠푐 = 푞 ∫ 퐴(퐸)휙𝑖푛푐(퐸)푑퐸 = 푞 ∫ 휙𝑖푛푐(퐸)푑퐸 (2.1) 0 퐸푔 where q is fundamental charge.

Light absorption by generating free carriers and light emission by recombination of electron-hole pairs is closely related through detailed balance, where photon emission and photon absorption per unit time should be equal in ideal case. When Würfel’s generalization is applied,[2] we can describe radiative recombination current Jrec under applied voltage bias:

∞ 푞푉 퐽 = 푞 ∫ 휙 (퐸, 푇) exp ( ) 푑퐸 (2.2) 푟푒푐 푏푏 푘푇 퐸푔

where 휙푏푏 is blackbody radiation spectrum. Eq. (2.2) describes current density of a solar cell in the dark when only radiative recombination is considered at temperature T. The total current density will be the superposition of radiative recombination current density and short circuit current density. We should note that there is photon flux from ambient environment, so that we should replace the incoming photon flux with 휙푠푢푛 + 휙푏푏. Then we have:

∞ 푞푣 ∞ 퐽(푉) = 푞 ∫ 휙 (퐸)푑퐸 [exp ( ) − 1] − 푞 ∫ 휙 (퐸)푑퐸 (2.3) 푏푏 푘푇 푠푢푛 퐸푔 퐸푔

Eq. (2.3) describes a typical diode equation with additional photocurrent due to extra solar irradiation.

Shockley-Queisser limit states the situation where only radiative recombination is present.

In real cases, non-radiative recombination will degrade solar cell performance from ideal situation. 16

In 1967, Ross generalized quantitatively how non-radiative recombination reduces solar cell performance, or open circuit voltage to be more specific[3]:

푘푇 푉 = 푉 + ln(휂 ) (2.4) 표푐 표푐,푟푎푑 푞 푒푥푡

In his statement, 푉표푐,푟푎푑 is the radiative limit of open circuit voltage and 휂푒푥푡 is external photoluminescence quantum efficiency (PLQE). This is meaningful in a way that we should maximize external PLQE to achieve highest Voc possible. Thus, defect passivation to minimize

SRH recombination and photonic design for best luminescence extraction becomes important. This idea is validated in GaAs thin film solar cells that by maximizing external PLQE, very high solar conversion efficiency was achieved.[4]

2.3 Theoretical limit of perovskite solar cells

Perovskite materials demonstrate benign optoelectronic properties that non-radiative recombination is low compared with other polycrystalline materials. There has been extensive research effort on optimizing PLQE of perovskite thin films. Deschler et al reported over 70%

PLQE of CH3NH3PbI3 perovskite thin films in 2014, which was the first report of PLQE of perovskite material.[5] There has been a series of research that proved PLQE can be enhanced by defect passivation through Lewis base. Snaith et al reported pyridine can passivate perovskite surface and enhance PL.[6] Ginger et al reported that a selection of Lewis bases can enhance

PLQE, exhibiting PLQE as high as 35% and PL lifetimes over 8 µs.[7] Pazos-Outón reported photon recycling effect through an indirect experiment.[8]

There are some recent progresses with perovskite thin film PLQE. Braly et al reported over

90% internal PLQE of CH3NH3PbI3 perovskite film with TOPO passivation.[9] The work demonstrated over 20% external PLQE, which is among the highest reported. Abdi-Jalebi et al reported high PLQE based on alloyed perovskite composition,

((Cs0.06FA0.79MA0.15)Pb(I0.85Br0.15)3, through potassium passivation.[10] PLQE as high as 66% was reported. More significantly, they reported PLQE of perovskite p-i-n device stack to be over

20%. This result matches previous GaAs solar cell record,[11] indicating optimal optoelectronic properties perovskite material has competing with the best solar material known, GaAs.

With these progresses on minimizing non-radiative recombination of perovskite thin films, we can estimate efficiency limit of perovskite solar cells.[12,13] Considering simplest MAPbI3 perovskite with a band gap of 1.6 eV, by application of Shockley-Queisser’s model, we can calculate maximum efficiency to be 31%. By considering another intrinsic aspect of perovskite material, Auger recombination will bring the efficiency limit down to 30.5%. Both numbers are approaching best single junction solar cell efficiency limit of 33%. There is a big room for improvement in perovskite solar cell efficiency.

References

[1] Shockley, William, and Hans J. Queisser. "Detailed balance limit of efficiency of p‐n junction solar cells." Journal of applied physics 32.3 (1961): 510-519.

[2] Würfel, Peter. Physics of solar cells. Vol. 1. Weinheim: Wiley-vch, 2005.

[3] Ross, Robert T. "Some thermodynamics of photochemical systems." The Journal of Chemical

Physics 46.12 (1967): 4590-4593.

[4] Miller, Owen D., Eli Yablonovitch, and Sarah R. Kurtz. "Strong internal and external luminescence as solar cells approach the Shockley–Queisser limit." IEEE Journal of Photovoltaics

2.3 (2012): 303-311.

[5] Deschler, Felix, et al. "High photoluminescence efficiency and optically pumped lasing in solution-processed mixed halide perovskite semiconductors." The journal of physical chemistry letters 5.8 (2014): 1421-1426.

[6] Noel, Nakita K., et al. "Enhanced photoluminescence and solar cell performance via Lewis base passivation of organic–inorganic lead halide perovskites." ACS nano 8.10 (2014): 9815-9821.

[7] deQuilettes, Dane W., et al. "Photoluminescence lifetimes exceeding 8 μs and quantum yields exceeding 30% in hybrid perovskite thin films by ligand passivation." ACS Energy Letters 1.2

(2016): 438-444.

[8] Pazos-Outón, Luis M., et al. "Photon recycling in lead iodide perovskite solar cells." Science

351.6280 (2016): 1430-1433.

[9] Ginger, David S., and Hugh W. Hillhouse. "Hybrid perovskite films approaching the radiative limit with over 90% photoluminescence quantum efficiency." Nature Photonics (2018).

[10] Abdi-Jalebi, Mojtaba, et al. "Maximizing and stabilizing luminescence from halide perovskites with potassium passivation." Nature 555.7697 (2018): 497.

[11] Green, Martin A. "Radiative efficiency of state‐of‐the‐art photovoltaic cells." Progress in Photovoltaics: Research and Applications 20.4 (2012): 472-476.

[12] Sha, Wei EI, et al. "The efficiency limit of CH3NH3PbI3 perovskite solar cells." Applied

Physics Letters 106.22 (2015): 221104.

[13] Pazos-Outón, Luis M., T. Patrick Xiao, and Eli Yablonovitch. "Fundamental efficiency limit of lead iodide perovskite solar cells." The journal of physical chemistry letters 9.7 (2018): 1703-

1711.

Chapter 3: High photoluminescence quantum efficiency of CH3NH3PbI3

perovskite on NiOx hole contact

3.1 Introduction

Organic inorganic halide hybrid perovskite solar cells have been intensively investigated in the recent years due to the surging power conversion efficiency which has passed 22% in just several years of research.[1] Contradictory to traditional high efficiency solar cells that utilize high quality semiconductor materials, fabricated through high temperature or high vacuum processes, hybrid perovskites can be synthesized in tremendously easy solution process while maintaining a comparably high material quality with traditional semiconductors. Within a wide variety of compositions of perovskites, CH3NH3PbI3 with a bad gap of 1.6 eV is the mostly investigated one.

The material demonstrates superior semiconductor properties including high absorption coefficient,[2] long charge carrier lifetime and diffusion lengths,[3] low Urbach energy and high photoluminescence quantum efficiency (PLQE).[4,5] The good optoelectronic properties perovskites have make it a promising material in a wide range of optoelectronic applications.[6]

Perovskite solar cells are made with two common structures: regular n-i-p structure, commonly with TiO2 as bottom contact, and inverted p-i-n structure. There is a selection of hole transport materials that we can fabricate perovskite onto. The most common one is PEDOT:PSS, which is the common material adopted from OPV. The material is easy to process and features rather low annealing temperature. However, perovskite solar cells fabricated with PEDOT:PSS suffer from a huge Voc loss compared with traditional TiO2 structure. The Voc is below 1 V compared with over 1.1 V in TiO2 structure. Recently, solution processed NiOx material emerged as a good hole transport layer in perovskite solar cells. A wide range of methods have been used

21 to synthesize NiOx compact layer, ranging from low temperature nanoparticle method to high temperature sol-gel method. Both methods can yield perovskite solar cell with a good power conversion efficiency above 17%. Hou et al. have demonstrated a scalable, facile-processed device that achieves 17% using a low-temperature NiOx.[7] Laser pulsed deposition, as Park et al utilize, is also a viable processing method for achieving PCEs above 17%.[8] Including Voc, NiOx has the potential to benefit other device parameters. As shown in the work of Shihe Yang, NiO was modified by a diethanolamine interlayer resulting in a FF of 80, with a PCE of 16%.[9] Yang

Yang’s group was also able to achieve a device performance of 16% using NiO with a 90% retention of PCE over 60 days in open air.[10] A solution processed NiO has also proven useful in achieving beyond 16% PCE for flexible solar cells.[11] More importantly, compared with

PEDOT:PSS structure, NiOx based perovskite solar cells generally suffer less loss in Voc, with

Voc well above 1.05 V. Photoluminescence is a powerful tool for analyzing material properties for device applications. PLQE serves as the primary indicator for low non-radiative recombination and semiconductor quality. Especially for photovoltaic applications, PLQE is in a direct relationship with solar cell open circuit voltage (Voc).[12] A high PLQE in room temperature is necessary for low loss in photo-voltage output, which is essential for maximizing solar cells power conversion efficiency (PCE). It has been reported that organic inorganic hybrid perovskite films show a 70% PLQE in room temperature and near 100% in low temperature, measured with integrating sphere.[4,13] Defects states caused by electronic structure disorder have been shown to negatively impact film photoluminescence,[14-16] which has led to an increased interest in characterizing microstrucutres and carrier recombination dynamics.[17-19] Using photoconductive atomic force microscopy (pcAFM), Kutes et al were able to map key device parameters, such as Voc and Isc, across nanoscopic scales, showing how defect states in disordered

22 areas correlate with poor device performance.[20] Along a similar vein, Zhao et al combined

Kelvin probe force microscopy with pcAFM to characterize three films with varying amounts of microstructure disorder, suggesting that poor device performance was related to small grain sizes and high densities of grain boundaries.[21] Asbury et al propose that perovskite surface morphology is more significant than bulk morphology when considering nonradiative recombination. Unwanted recombination centers are more prolific amongst film surfaces, hence passivation efforts should be focused towards forming an optimal surface chemistry.[22] High electroluminescence (EL) can also indicate high Voc in perovskite solar cells, as recently reported by Bi et al.[23] Sutter-Fella et al shows perovskite films with high internal PLQE with band gaps tunable between 1.6 eV and 2.3 eV.[24]

In this work, we demonstrate how the bottom contact in inverted perovskite solar cell structure impact PLQE and Voc. We found a clear difference in PLQE of CH3NH3PbI3 perovskite films on NiOx and PEDOT:PSS coated ITO substrates. Also, both films show a far lower PLQE compared with films grown on glass substrates, indicating a loss of Voc in the films on contacts.

The trend of decreasing PLQE is in agreement with carrier lifetimes from time-resolved PL data.

We compare two different inverted device structures, ITO/NiOx/Perovskite/PC61BM/Al and

ITO/PEDOT:PSS/Perovskite/PC61BM/Al, in terms of their Voc and external PLQE (ext-PLQE).

With 0.1 V lower Voc, PEDOT:PSS device structure also shows much lower ext-PLQE. We performed SEM imaging and XRD study to explain this difference. It has been found that there is a dramatic morphological change in the perovskite film from SEM imaging when comparing film quality on different substrates. However, the films show a same XRD pattern, which indicates the same chemical and crystal structure of the films. These results indicate the lower quality in

23 perovskite films and perovskite/contact interfaces is decreasing PL, thus penalizing Voc output from perovskite solar cells.

3.2 Experimental details

3.2.1 Perovskite thin film and device fabrication Perovskite films are fabricated on three different types of substrates: soda-lime glass,

PEDOT:PSS and NiOx coated indium doped tin oxide (ITO) glass. PEDOT:PSS (Clevious) is spin- coated on ITO glass at 2000 rpm for 30s, then annealed at 150 ºC for 15min. NiOx film is fabricated through reported method.[10] Perovskite film is fabricated through two-step method. PbI2 (Sigma-

Aldrich) is dissolved in dimethylformamide (DMF) with 1 mol/L concentration and spin-coated on substrate at 3000 rpm. CH3NH3I (MAI) is dissolved in isopropyl alcohol (IPA), 1 mol/L with

10:1 CH3NH3Cl (MACl) doping, and spin-coated on PbI2 at 4000 rpm. The double layer precursor is annealed in atmosphere for 2 min at 130 ºC to form perovskite layer of 300 nm. To finish the device fabrication process, Phenyl-C61-butyric acid methyl ester (PC61BM) solution at 20 mg/ml is spin coated at 1500 rpm for 30 s on top of perovskite layer to make electron contact. Then, a layer of Aluminum Zinc Oxide (AZO) is deposited by spin coating of AZO nanoparticle solution

(Sigma-Aldrich) at 3000 rpm for 30 s and annealed at 85 oC for 45 s. Finally, a layer of 100 nm Al is thermally evaporated on the top to complete the device.

3.2.2 Characterization and measurement UV-Vis absorption spectra of the perovskite films were obtained using a U-4100 spectrophotometer (Hitachi) equipped with integrating sphere, in which monochromatic light was incident to the substrate side. X-ray diffraction (XRD) patterns of the films were recorded by X- ray diffractometer (PANalytical) with Cu kα radiation at a scan rate of 4 o/min. Surface and cross-

24 sectional microscopic images of the films and devices were acquired by scanning electron microscopy (SEM, Nova Nano 230). Steady-state photoluminescence (PL) measurement was carried out using Horiba Jobin Yvon system, in which a 640 nm monochromatic laser was used as an excitation fluorescence source. Time-resolved PL decay profiles of the perovskite films were investigated by a Picoharp 300 with time-correlated single-photon counting capabilities. A 640 nm monochromatic pulsed laser with a repetition frequency of 100 kHz was generated from a picosecond laser diode head (PLD 800B, PicoQuant). PLQE is measured through a homemade

PLQE system by comparing luminescence intensity and diffuse reflection from a near perfect diffuse reflector.

3.3 Results and discussion

We fabricated MAPbI3 perovskite films on NiOx and PEDOT:PSS coated substrates, through two-step sequential deposition of PbI2 and MAI solutions. PEDOT:PSS hole transport layer is made with spin-coating of commercially available solution. NiOx film is made by our previously reported method. To address the impact of bottom hole contact on perovskite films, we performed steady state PLQE measurement on as fabricated films. For comparison, we also fabricated perovskite film on soda-lime glass. Figure 3.1 demonstrates the steady state photoluminescence. Perovskite films grown on different substrates show similar spectrum shape, with a slight variation of peak wavelength. However, the PL intensity under same excitation intensity varied dramatically for these films, indicating a substantial difference in photoluminescence quantum efficiency. From pump-power dependent PL intensities of perovskite films (Figure 3.1b), we can calculate internal and external PLQE under various excitation light power densities (Figure 3.1c). The methodology we use to measure PL and light extraction model

25 we use to extract int-PLQE is elaborated in Supporting Information. The highest int-PLQE in three film structures is found to be on the glass substrate, with peak int-PLQE to be 78% with illumination intensity a bit over one-sun. Other substrates show much lower PLQE, indicating a substantial loss of charge carriers through non-radiative recombination. PLQE under one-sun intensity illumination is of special interest for photovoltaic application. To equal the absorbed photon flux of AM 1.5G with 532 nm CW laser we use to perform PL experiments, only 72 mW/cm2 green laser is used to simulate one-sun condition. Under one-sun illumination, int-PLQE of perovskite on glass, NiOx and PEDOT:PSS is found to be 61%, 6.8% and 0.75%. With int-

PLQE one and two orders of magnitude lower in NiOx and PEDOT:PSS substrates compared with on glass, we see clearly that non-radiative recombination is severe in perovskite films with contact, especially on PEDOT:PSS. The low PLQE would compromise solar cell Voc dramatically.

(a) 200 1.0 180 Glass NiOx 160 PEDOT:PSS 0.5 140

120 (Normalized) PL 0.0

100 700 800 900

Wavelength (nm) 80 60

40 Photoluminescence (a.u.) Photoluminescence 20 0 700 750 800 850 900 Wavelength (nm)

(b) 104 White 103 Glass NiOx PEDOT:PSS 102

101

100

-1 Illumination (mW/cm2) Illumination 10

10-2 10-3 10-2 10-1 100 101 102 103 104 105 106 luminescence (nW)

Number of suns (c) 0.01 0.1 1 10 100

Glass int-PLQE (%) int-PLQE 0.1 NiOx PEDOT:PSS 0.01 0.1 1 10 100 1000 Illumination intensity (mW/cm2)

Figure 3.1 Steady state photoluminescence of perovskite films grown on different substrates (a) PL spectra under one-sun illumination. (b) Log-log plot of measured PL intensity vs illumination power density. (c) Pump power dependence of internal PLQE To further determine why the difference in A further study in time resolved photoluminescence (TRPL) shows the minority charge carriers lifetime agrees with steady state

PL trend. In Figure 3.2, TRPL decay is carried out with 3 pump fluences, resulting in a charge

27 carrier density of 8×1016 cm-3, 1.6×1016 cm-3 and 3×10-15 cm-3. A clear trend of decreasing minority carrier lifetime with increasing charge carrier density is observed with perovskite on every substrates, which has also been reported in other research results. The decrease of charge carrier lifetimes is attributed to multi charge carrier processes. When comparing TRPL of perovskite films on different substrates, it shows a distinguishable difference. It shows longest lifetime on glass substrate and shortest lifetime on PEDOT:PSS substrate, with NiOx films in the middle. Longer minority carrier lifetimes indicate less defect states in the band gap that act as non-radiative recombination centers. The trend is in good agreement with steady state PLQE results.

(a) 17

) 10 -3

1016

1015

14 Charge carrier density (cm density carrier Charge 10 0 200 400 600 800 1000 Time (ns)

(b) 17

) 10 -3

1016

1015

14 Charge carrier density (cm density carrier Charge 10 0 100 200 300 400 Time (ns)

(c)

1017

) -3

1016

1015

1014 Charge carrier density (cm density carrier Charge 0 100 200 Time (ns)

Figure 3.2 Time Resolved PL decay and bi-exponential fitting of perovskite films on (a) Glass (b) NiOx (c) PEDOT:PSS substrates, measured at spectrum peak wavelength and with three different pump fluences.

Solar cell devices are made based on NiOx and PEDOT:PSS substrates as hole transport electrodes. On top of perovskite layer, we apply the commonly used electron transport material

29 phenyl-C61-butyric acid methyl ester (PC61BM) and evaporated Al as electron transport material.

Device performance can be seen in Table 1.

Table 1 Summary of solar cell performances Voc (V) Jsc (mA/cm-2) Fill Factor (%) PCE (%) Ext-PLQE

−4 NiOx 1.08 21.4 74.5 17.3 6.6 × 10

PEDOT:PSS 0.95 21.3 73 14.7 1.3 × 10−4

With similar Jsc and fill factor, NiOx based solar cell shows dramatically higher Voc. The dramatic drop of Voc is in correspondence with a big decrease of int-PLQE in perovskite films.

We also measured external PLQE on perovskite solar cell devices under one-sun illumination and we found that even with top electron selective contact, NiOx based devices still show a much higher ext-PLQE. Ext-PLQE has a direct connection with solar cell Voc defined by Eq. (1):

푉표푐 = 푉표푐,푚푎푥 − 푘푇|ln 휂푒푥푡| (1)

In Eq. (1), Voc,max is the maximum Voc we can obtain from any solar material. It is 0.3 V lower than solar material’s band gap, where the reduce comes from entropy consumption.

Perovskite as a solar material with a band gap of 1.6 eV, the maximum Voc that can be extracted from this material would be 1.3 V. This value is further penalized with external PLQE lower than unity. As a result, with a much lower PLQE value, perovskite on PEDOT:PSS shows a much lower Voc, which originated from consumption of charge carriers through non-radiative recombination.

(a) 0

-5 NiOx

PEDOT:PSS )

2 -10

-15

J (mA/cm J -20

-25 0.0 0.2 0.4 0.6 0.8 1.0 1.2 Voltage (V)

(b) 100

40 EQE (%) NiOx PEDOT:PSS 20

300 400 500 600 700 800 Wavelength (nm)

Figure 3.3 (a) Current density versus voltage (J-V) of solar cells based on NiOx and PEDOT:PSS substrates. (b) EQE of corresponding devices

To understand the origin of reduction in PLQE of perovskite on contacts, SEM images of the three perovskite film stacks are shown in Fig. 4. MAPbI3 fabricated on glass Fig. 4(a), NiO 31 figure 3.4(b), and PEDOT figure 3.4(c) resulted in grain sizes ranging up to 1um, 600nm, and

200nm, respectively. The superior PL of MAPbI3 on glass and NiOx suggests that improved crystallinity results in a more luminescent film. We envisage that this correlation between film morphology and PL is due to the reduction of grain boundaries as grain sizes increase. While the connection between perovskite grain boundaries and non-radiative recombination centers is still in contention, Ginger et al. demonstrated tenable evidence that grain boundaries reduce the PL of

MAPbI3.[15] As Figure 4(d) shows, XRD patterns for each of these stacks remained the same, ruling out the possibility that an increase in PL is due to an altered MAPbI3 composition.

(a) (b)

(110) PEDOT

(220)

(213)

Intensity (a.u.)

10 20 30 40 50 2 (deg.)

Figure 3.4 SEM image of perovskite films on (a)glass (b)NiOx (c)PEDOT:PSS. (d) XRD patterns of the corresponding films 32

3.4 Conclusion

In conclusion, we demonstrated the direct relationship of perovskite solar cell Voc with

PLQE. Perovskite films fabricated on NiOx and PEDOT:PSS substrates show int-PLQE of 6.8% and 0.75% respectively under one-sun illumination, both much lower than 61%, on glass substrate.

The dramatic difference is also found in TRPL experiment, in which glass substrates show longer minority charge carrier lifetimes. Solar cells fabricated from NiOx and PEDOT:PSS substrates with PC61BM/Al electron selective contact show a difference in Voc while Jsc and fill factor remain similar. We performed PLQE measurement on as fabricated devices and find the same trend. A further study of film morphology and crystallization is done by SEM imaging and XRD patterns. We find that perovskite film on different substrates show same XRD pattern but very different crystal morphology, with crystal size of perovskites much bigger on glass than on NiOx and PEDOT:PSS. The surface energy and defect states on the interface contribute to the distinction.

References

[1] W. S. Yang, B.-W. Park, E. H. Jung, N. J. Jeon, Y. C. Kim, D. U. Lee, S. S. Shin, J. Seo, E.

K. Kim, J. H. Noh, S. I. Seok, Iodide management in formamidinium lead halide-based perovskite layers for efficient solar cells, Science, 356, 167–171 (2017).

[2] Xiao, Z. et al. Solvent annealing of perovskite-induced crystal growth for photovoltaic-device efficiency enhancement. Adv. Mater. 26, 6503–6509 (2014).

[3] Stranks, Samuel D., et al. "Electron-hole diffusion lengths exceeding 1 micrometer in an organometal trihalide perovskite absorber." Science 342.6156 (2013): 341-344.

[4] Deschler, Felix, et al. "High photoluminescence efficiency and optically pumped lasing in solution-processed mixed halide perovskite semiconductors." The journal of physical chemistry letters 5.8 (2014): 1421-1426.

[5] De Wolf, Stefaan, et al. "Organometallic halide perovskites: sharp optical absorption edge and its relation to photovoltaic performance." The journal of physical chemistry letters 5.6

(2014): 1035-1039.

[6] Stranks, Samuel D., and Henry J. Snaith. "Metal-halide perovskites for photovoltaic and light-emitting devices." Nature nanotechnology 10.5 (2015): 391.

[7] Hou, Yi, et al. "Overcoming the Interface Losses in Planar Heterojunction Perovskite‐Based

Solar Cells." Advanced Materials 28.25 (2016): 5112-5120.

[8] Park, Jong Hoon, et al. "Efficient CH3NH3PbI3 Perovskite Solar Cells Employing

Nanostructured p‐Type NiO Electrode Formed by a Pulsed Laser Deposition." Advanced

Materials 27.27 (2015): 4013-4019.

[9] Bai, Yang, et al. "Effects of a molecular monolayer modification of NiO nanocrystal layer surfaces on perovskite crystallization and interface contact toward faster hole extraction and higher photovoltaic performance." Advanced Functional Materials 26.17 (2016): 2950-2958.

[10] You, Jingbi, et al. "Improved air stability of perovskite solar cells via solution-processed metal oxide transport layers." Nature nanotechnology 11.1 (2016): 75.

[11] Yin, Xingtian, et al. "Highly efficient flexible perovskite solar cells using solution-derived

NiO x hole contacts." ACS nano 10.3 (2016): 3630-3636.

[12] Miller, Owen D., Eli Yablonovitch, and Sarah R. Kurtz. "Strong internal and external luminescence as solar cells approach the Shockley–Queisser limit." IEEE Journal of Photovoltaics

2.3 (2012): 303-311.

[13] Stranks, Samuel D., et al. "Recombination kinetics in organic-inorganic perovskites: excitons, free charge, and subgap states." Physical Review Applied 2.3 (2014): 034007.

[14] Draguta, Sergiu, et al. "Spatially non-uniform trap state densities in solution-processed hybrid perovskite thin films." The journal of physical chemistry letters 7.4 (2016): 715-721.

[15] Vorpahl, Sarah M., et al. "Impact of microstructure on local carrier lifetime in perovskite solar cells." Science (2015): aaa5333.

[16] Zhang, Wei, et al. "Photo-induced halide redistribution in organic–inorganic perovskite films." Nature communications 7 (2016): 11683.

[17] Zhu, X-Y., and V. Podzorov. "Charge carriers in hybrid organic–inorganic lead halide perovskites might be protected as large polarons." (2015): 4758-4761.

[18] Nie, Wanyi, et al. "High-efficiency solution-processed perovskite solar cells with millimeter- scale grains." Science 347.6221 (2015): 522-525.

[19] Johnston, Michael B., and Laura M. Herz. "Hybrid perovskites for photovoltaics: charge- carrier recombination, diffusion, and radiative efficiencies." Accounts of chemical research 49.1

(2015): 146-154.

[20] Kutes, Yasemin, et al. "Mapping the photoresponse of CH3NH3PbI3 hybrid perovskite thin films at the nanoscale." Nano letters 16.6 (2016): 3434-3441.

[21] Jiang, Chun-Sheng, et al. "Carrier separation and transport in perovskite solar cells studied by nanometre-scale profiling of electrical potential." Nature communications 6 (2015): 8397.

[22] Stewart, Robert J., et al. "Approaching bulk carrier dynamics in organo-halide perovskite nanocrystalline films by surface passivation." The journal of physical chemistry letters 7.7 (2016):

1148-1153.

[23] Bi, Dongqin, et al. "Efficient luminescent solar cells based on tailored mixed-cation perovskites." Science advances 2.1 (2016): e1501170.

[24] Sutter-Fella, Carolin M., et al. "High photoluminescence quantum yield in band gap tunable bromide containing mixed halide perovskites." Nano letters 16.1 (2015): 800-806.

Chapter 4: Optoelectronics of Single Crystal Formamidinium Lead Iodide

(FAPbI3)

4.1 Introduction

Recently, hybrid organolead trihalide perovskite (OTP) solar cells have developed as a promising candidate in photovoltaics due to their excellent properties including a direct bandgap,[1] strong absorption coefficient,[2] long carrier lifetime,[3] and high mobility.[4, 5] Basically, the

OTP structure is ABX3 where A and B represent an organic ammonium cation and Pb, respectively, and X is a halide such as I or a combination of Cl, Br and I.[6-12] Among these candidates for

+ solar cell application, the most popular is methylammonium (CH3NH3 or MA) lead iodide

(MAPbI3), reaching nearly 20% power conversion efficiency.[13, 14] Most recently,

+ formamidinium (NH2CH=NH2 or FA) lead iodide (FAPbI3) has attracted significant attention due to several advantages: (1) the larger organic FA cation can replace the MA cation and form a more symmetric crystal structure, (2) the smaller bandgap of FAPbI3 allows for near infrared (NIR) absorption, and (3) FAPbI3 has an elevated decomposition temperature and thus potential to improve stability.[15-18] Perovskite solar cells incorporating FAPbI3 in absorber layer

((FAPbI3)0.85(MAPbBr3)0.15) have achieved record certified power conversion efficiency of

20.1%.[19] Despite of many studies on FA based perovskite structures in thin film solar cells, the intrinsic electrical, optical, and structural properties of FAPbI3 still require further study.[20, 21]

Typically, it is difficult to understand the fundamental properties of FAPbI3 in thin films due to a favorable phase transforming at room temperature and the large amount of recombination sites due to grain boundaries, voids, and surface defects. Single crystals, on the other hand, provide an excellent model system to study the intrinsic electrical and optical properties of these materials

37 due to their high purity. This is particularly important to understand the limits of these materials.

So far, several approaches of the single crystal growth of MAPbI3 and MAPbBr3 have been studied.

Huang et al. reported a 10 mm-sized growth of single crystal MAPbI3 using a top-seeded solution method with a temperature gradient and obtained a carrier diffusion length over 175 μm.[22] Bakr et al. reported an anti-solvent vapor-assisted crystallization approach that enabled sizable crack- free MAPbX3 (X = I or Br) single crystals with volumes exceeding 100 cubic millimeters.[23] In addition, Tao et al. and Lin et al. used a cooling solution method to grow MAPbI3 single crystals in HI solution.[24, 25] Later Bakr et al. and Liu et al. separately reported using an inverse temperature crystallization method to grow MAPbX3 perovskite single crystals.[26, 27] Single crystal FAPbI3 has hardly been reported. Kanatzidis et al. reported the fundamental study of 50

μm sized single crystal FAPbI3, which is significantly smaller than the 10 mm size MAPbI3.[28]

Bakr et al. also grew FAPbX3 single crystals using the similar methods as they grew MAPbX3 single crystals.[29]

Here, we demonstrate for the first time the growth of 5 millimeter-sized single crystal

FAPbI3. The crystal was grown using a modified inverse temperature crystallization method. We used a cooling solution method to first grow the FAPbI3 seed crystal, followed by placing the seed crystals in an inverse temperature crystallization precursor to obtain larger crystals. The thermal, optical and electrical properties of single crystal FAPbI3 were explored. In addition, we fabricated single crystal FAPbI3 photodetectors and investigated photoconductivity and photoresponse. Low noise current and high switching speeds were confirmed.

4.2 Experiment Details

4.2.1 Synthesis of single crystal FAPbI3

Lead (II) acetate trihydrate (Pb(ac)2·3H2O, 99%), Fomamidine acetate salt (FAac, 99%),

Lead iodide (PbI2, 99.999%), Hydriodic acid (HI) (57% w/w aq. soln., stab with 1.5% hypophosphorous acid) and gamma-butyrolactone (GBL, 99%) were purchased from Sigma

Aldrich. Formamidinium iodide (FAI) was purchased from Dyesol Limited (Australia). All reagents and solvents were used as received without any further purification.

2.5g Pb(ac)2·3H2O was fully dissolved in 15 ml of HI in a 100 ml flask which was put into a 105°C oil bath to heat up the mixed acid solution. A blend solution of 1.5 ml HI solution and 0.7 g of FAac was added to the mixed acid solution. Then the temperature of mixed solution was decreased to 70 °C and kept for 6 h for the precipitation of seed crystal FAPbI3 (~1 mm in size).

The seed crystals were washed by diethyl ether and dried in vacuum.

At the beginning, 1.0 M solution containing PbI2 and FAI (1:1) was dissolved in GBL at

60°C overnight. Then, the solutions were filtered using polytetrafluoroethylene (PTFE) filter with

0.2-μm pore size. By placing the seed into this GBL solution in oil bath around 100 to 105°C for

3 h, the seed crystals grow into a large one. A larger crystal was formed by using this large crystal as the new seed put into a fresh precursor again. By repeating the above process three times, a large single crystal FAPbI3 was synthesized.

4.2.2 Synthesis of FAPbI3 thin films and permittivity measurement

450 mg of PbI2 was dissolved in 1 ml dimethylformamide (DMF) and then was spin-coated at indium-tin-oxide (ITO) substrates at 2,500 rpm for 30 s. Then, FAI (dissolved in 2-propanol) was spin-coated on the top of dried PbI2 layer at room temperature at 3,000 rpm. for 30 s in the

39 dry air. All films were annealed in the air at 150 °C for desired time. Impedance spectroscopy of

4 6 parallel plate capacitor from 10 to 10 Hz was used to measure permittivity of the FAPbI3.

The permittivity was calculated from the following equation,[30]

푑 −1 휀푟(푓) = (4.1) 퐴휀0 2휋푓퐼푚(푍)

where ε0 is the permittivity of vacuum, A is the capacitor area, and d is the distance between the two electrodes.

4.2.3 Measurement and characterization TGA-DSC was performed on a TA SDT-Q600. The crystal structures of the nanocrystals were characterized by D8 Focus X-ray diffraction and Raman spectra (Renishaw inVia Raman spectrometer) with an incident laser of 514.5 nm. Steady-state photoluminescence (PL) of α-phase single crystal was measured using Horiba Jobin Yvon system with an excitation of 640 nm, and the δ-phase single crystal PL spectrum is measured using a NKT SuperK Extreme lase with an excitation of 475 nm and New Focus Si fW detector. Time-dependent photoluminescence was acquired using the time-correlated single-photon counting technique (Picoharp 300), and the excitation was provided by using a picosecond diode laser at the wavelength of 640 nm with a repetition frequency of 1 MHz (PDL 800B). The absorption spectrum was measured using U-4100

Hitachi UV-Vis spectrophotometer. All current-voltage measurements were carried out by Agilent

4155C semiconductor parameter analyzer. ToF measurement was conducted by illuminating the devices with Laser Innovations DUO-220 dye laser. SR-570 current pre-amplifier with a bandwidth of 1 MHz was used to amplify the weak photocurrent, which recorded using a Tektronix

DPO 4104 Digital Phosphor Oscilloscope. The Tektronix AFG 3252 function generator was used

40 for frequency response measurements. Newport Model 1830-C was employed to gauge actual light intensity.

4.3 Results and discussion

MAPbI3 has a stable perovskite phase during growth at the temperature range from 60 ºC to 110 ºC for various growth methods. However, a small growth temperature variation could affect the growth process in the case of FAPbI3. In order to obtain a FAPbI3 seed crystal, we first attempted the inverse temperature crystallization method. Homogenous nucleations of α-phase

FAPbI3 started after 15 min at 100 ºC. The crystal growth speed was subsequently enhanced, leading to thermal fluctuation particularly at the solution/air interface as a result of the endothermic reaction. Light yellow δ-phase FAPbI3 and some needle like NH4PbI3 appeared at this time impeded the α-phase crystal growth.[29, 31] Therefore, it is difficult to obtain high quality single crystal FAPbI3 seeds directly using the inverse temperature crystallization method. In order to minimize the side effects, we use a new process combining the cooling solution and inverse temperature crystallization methods to grow single crystal FAPbI3. Firstly, we used the cooling solution method to grow single crystals seeds. Lead (II) acetate trihydrate (Pb (ac)2·3H2O) and

Fomamidine acetate salt (FAac) firstly dissolved in HI at 105°C and then gradually cooled to 70

°C. The seed crystals were obtained after 6 h. The details could be found in Supporting

Information. By placing a few seed crystals into the inverse temperature crystallization precursor solution and keeping it at 100°C, the seed crystal grew larger. This heterogeneous nucleation process effectively reduced NH4PbI3 and δ-phase FAPbI3 growth comparing to homogeneous nucleation, which is attributed to a small number of nucleation centers that reduce the precipitation rate and could decrease temperature fluctuation within the solution. In addition, we gradually

41 raised the temperature from 100 to 105 ºC after crystals began to grow. This compensated for the heat loss caused by crystal growth. After 1 h, the growth speed gradually decreased, and we turned back the temperature to 100 ºC to grow the crystal. By repeating the inverse temperature crystallization process 3 times, a large FAPbI3 single crystal was synthesized.

Interestingly, the color of FAPbI3 gradually changed from black (Figure 4.1a) to yellow

(Figure 4.1b) after ten days, regardless of storage in vacuum or inert gas. The inset of Figure 4.1 showed the photos of the same 5mm size single crystal FAPbI3 before and after the phase change at room temperature. The yellow colored crystal could also rapidly recover back to the black color without any visible shape change after the annealing at 185°C for only 2 min as shown in Figure

4.2. This change of single crystal FAPbI3 was monitored by heating a small yellow crystal from room temperature to 550°C at the rate of 5 ºC/min in Ar atmosphere using thermogravimetric analysis and differential scanning calorimetry (TGA-DSC) (Figure 4.1c). An endothermic peak was observed at 185ºC without mass loss, corresponding to the color change related to the phase transformation of the single crystal FAPbI3. A mass loss (over 20%), associated with the strong exothermic peak, occurred at 320–360 ºC and is related to the decomposition of hydrogen iodide

(HI).[27] From the temperatures of 375°C to 420°C, a mass loss (about 6%) occurred. This could be attributed to the decomposition of the formamidinium. This means that the formamidinium is more stable than iodide ion in FAPbI3 single crystal. The decomposition temperature of the FAPbI3

[27] single crystal is higher than MAPbI3 single crystal, which indicates that the thermal stability of

FAPbI3 is better than MAPbI3.

Figure 4.1 Crystal structures of a) α-phase and b) δ-phase single crystal FAPbI3. The inset of a) and b) are images of α-phase and δ-phase single crystal, respectively. c) TGA-DSC curves of single crystal FAPbI3. d) and e) are experimental and calculated powder XRD patterns of α-phase single crystal FAPbI3 and δ- phase FAPbI3, respectively.

Figure 4.2 δ-phase to α-phase FAPbI3 single crystal phase transition video capture images of annealing time interval 20 sec. a) 0 s, b) 20 s, c) 40 s, d) 60 s, e) 80 s, f) 100 s, g) 120 s and h) 140 s.

In order to understand the differences between the black and yellow crystals, X-ray diffraction (XRD) was performed on ground black and yellow single crystals to identify phase change. The results are shown in Figure 4.1d and e. For the black single crystal, all peaks were indexed to a trigonal perovskite phase (P3m1 space group, α-phase) with a = b = 0.898 nm and c

= 1.101 nm. On the other hand, the yellow crystal showed the hexagonal system (P63mc space group, δ-phase) with a = b = 0.866 nm and c = 0.790 nm.[28] The two phases were also studied by

Raman spectroscopy (Figure 4.3) using a 514.5 nm laser, where the main Raman peak red-shifted

-1 -1 from 135 cm to 111 cm while the FAPbI3 single crystal changed from α-phase to δ-phase, which might be related to the vibration of FA molecule.[32, 33]

Figure 4.3 Raman spectra of single crystal FAPbI3 with α-phase (black) and δ-phase (red).

The photoluminescence (PL) properties of FAPbI3 were measured at room temperature using a red light source at 633 nm. An emission peak of single crystal α-phase FAPbI3 at 820 nm was observed (the inset of Figure 4.4). In a 300-nm thick FAPbI3 thin film, the emission peak occurred at 805 nm (inset of Figure 4.5), resulting in a blue shift of the emission peak. The shift of absorption edge is also observed in the UV-Vis spectra shown in Figure 4.6. This means that

45 the single crystal structure has a lower bandgap than that of the thin film. Single crystals are predominantly free of grain boundaries or structural defects, which could reduce the localized defect states. Thus the red-shift of single crystal FAPbI3 PL can be related to the higher material quality than thin film FAPbI3. This indicates that the crystal quality in the current FAPbI3 thin film, and the resulting performance of FAPbI3-based solar cells, still have room to improve. The larger bandgap of single crystal δ-FAPbI3 was confirmed by using a UV-Vis absorption and PL spectrum as shown in Figure 4.7.

Figure 4.6 UV-Vis absorption spectra of α-FAPbI3 single crystal and film.

Figure 4.7 UV-Vis absorption spectrum of δ-FAPbI3. The inset shows the PL spectrum of δ-FAPbI3 single crystal.

Carrier lifetime (τ) of single crystal FAPbI3 was investigated by using time-dependent PL

(TDPL) measurement. Transient fluorescent decay for the peak emission occurred from the photo- excitation at 820 nm as shown in Figure 4.4. The two-exponential decay showed fast (τ1 = 32 ns) and slow (τ2 = 484 ns) carrier lifetimes of the single crystal. The shorter lifetime originates from the high trap density related to the crystal surface conditions, and the longer lifetime represents the carrier transportation in bulk crystal having fewer defects.[23, 34] The single crystal lifetime is much longer than that of the thin film (τ1 = 29 ns and τ2 = 227 ns, Figure 4.5), which indicates that the single crystal FAPbI3 has much lower quantity of defects. It is expected that enhanced crystallinity of the FAPbI3 would contribute to improvement of the power conversion efficiency of the thin film solar cell. Compared to single crystal MAPbI3, FAPbI3 has relatively shorter carrier lifetime,[22, 23] which might be related to the fact that pure FAPbI3 based solar cells have relatively low efficiency below 15%.[35]

Current-voltage (I-V) measurement was performed using the single crystal FAPbI3 deposited two indium electrodes on opposite sides in the dark (Figure 4.8). Ohmic and quadratic regions were clearly separated at a bias voltage of 33 V. The result fits well with the Mott’s Space charge-limited current (SCLC) theory, i.e., beyond the bias voltage of 33 V, the current follows

Mott-Gurney’s square law:

9 V 2 J = 0b (4.2) D 8L3

where JD, ε, ε0, μ, Vb and L are the dark current, the relative dielectric constant, the vacuum permittivity, the mobility of single crystal FAPbI3, applied voltage, and the thickness of the single crystal, respectively. We can generally estimate the carrier mobility to be 4.4 cm2·V-1·s-1 from quadratic region, using relative dielectric constant ε = 46.9 measured in our lab. Time-of-flight shown in Figure 4.9 was also used to measure the carrier mobility of 1.07± 0.25 cm2 V-1 S-1 according to the equation μ=L2/Vτ, where μ, L, V and τ are the mobility, the thickness, applied voltage and the transit time of α-phase FAPbI3 single crystal, respectively. This mobility value is smaller but of the same order of magnitude comparing to the result from SCLC measurement.

The applied voltage at the kink point of current increasing linearly and non-linearly versus voltages is the trap-filled limit voltage (VTFL), which is determined by the defect density:[22, 23]

2 entL VTFL= (4.3) 2εε0

where e is elementary charge, nt is defect density, L is the thickness of FAPbI3 material, ε0 is vacuum permittivity, ε is relative dielectric constant. The defect density of α-phase FAPbI3 single crystal are 6.2 × 1011 cm-3, which is an order of magnitude higher than that of single crystal

[23] MAPbI3. 49

Figure 4.8 Dark current-voltage curve of α-phase single crystal FAPbI3 for space charge limited current analysis.

Using the linear region data, the conductivity and carrier concentration of single crystal

-7 -1 11 -3 FAPbI3 can be calculated to be σ = 1.1 × 10 (ohm·cm) and n = 1.5 × 10 cm , respectively.

Interestingly, the conductivity of the single crystal FAPbI3 is an order of magnitude higher than

-8 -1 that of single crystal MAPbI3 (σ= 1 × 10 (ohm·cm) ),[23] which is mainly attributed to the higher carrier concentration. The higher carrier concentration could be related to lower bandgap for

FAPbI3, possibly the difference in crystal structure, the electronic structure and defect density, which clearly deserves further study.

The mobilities, conductivities, carrier concentrations and defect density of δ-phase FAPbI3

2 -1 -1 -9 -1 11 -3 single crystal are μ = 0.179 cm V S , σ = 8.9 × 10 (Ω cm) , n = 3.1 × 10 cm and nt = 2.6×

1012 cm-3, respectively, which are also calculated by SCLC theory shown in Figure 4.10. The conductivity decreases compare to α-phase FAPbI3 single crystal could be attributed to the higher bandgap and defect density than α-phase FAPbI3 single crystal. Moreover, the conductivity of α- phase FAPbI3 thin film and single crystal are similar, which is consistent to MAPbI3 result.[36]

However, the mobility (1.1 × 10-3 cm2 V-1 S-1) decreases and defect density (1.3× 1016 cm-3) increases comparing to its single crystal as shown in Figure 4.11.

Figure 4.10 Dark current-voltage curve of δ-phase FAPbI3 single crystla for space charge limited current (SCLC) analysis.

Figure 4.11 Dark current-voltage curve of α-phase FAPbI3 thin film for space charge limited current analysis.

Combining the Einstein relation, D = μKBT/q and the definition of diffusion length

LD=  , where D is the diffusion constant, KB is Boltzmann's constant, T is the absolute temperature, and q is the elementary charge, we can calculated the carrier diffusion length of α- phase FAPbI3. The short and long diffusion lengths calculated from the short and long lifetimes, were 0.5 μm and 2.2 μm, respectively. Using a similar measurement method, Bakr et al. observed similar short (2 μm) and long (8 μm) components in diffusion length for the MAPbI3 single crystal.

Adopting the methodology of this group, here the long and short diffusion lengths were assigned to the bulk and surface components of FAPbI3 single crystals, respectively. Since both MAPbI3 and FAPbI3 have relatively long diffusion length, and FAPbI3 has a broader light absorption ability, it is reasonable to expect that a mixed perovskite absorption layer such as (FA/MA)PbX3

(X can be a pure or combination of halides - Cl, Br and I) can improve solar cell performance as reported.[37]

The excellent material properties such as the large light absorption coefficient, high carrier mobility, long carrier lifetime and long diffusion length of the FAPbI3 single crystal strongly suggest it’s a suitable material for photodetector. Therefore, a photoconductive type photodetector based on single crystal FAPbI3 was fabricated to prove the concept. The schematic illustration of the photodetector is shown in Figure 4.12. The devices were evaluated under the power intensity and the wavelength variations. Figure 4.13a shows a plotted linear response of single crystal

-4 FAPbI3-based photodetectors within the incident light power density range from 5 × 10 to 5 ×

10-1 mW·cm-2 at 380 nm (incident photon flux density range from 9.56 × 1011 to 9.56 × 1014 number·s-1·cm-2) under 0.1 V bias. The photoresponsivity decreased linearly from 27.57 to 0.68

A·W-1 when the light intensity increased, which is attributed to the enhanced charge recombination under higher light intensity. 53

Figure 4.12 Schematic illustration of the photodetector fabricated on free-standing single crystal FAPbI3.

Figure 4.13 Device configuration and characteristics of the photodetector based on single crystal FAPbI3. a) Photocurrent and photoresponsivity of devices for incident power densities. b) Transient photocurrent response at a pulse frequency of 1 Hz. Figure 4.13b represents the photo-switching of the devices at a frequency of 1 Hz, a wavelength of 380 nm, and a light intensity of 1 mW·cm-2. The photocurrent rapidly increased to a high level and recovered to the dark state. The rise-time (tr) of the device (from 10% to 90% of the saturated value) is 17ms, whereas the fall-time (tf) (from 90% to 10% peak value) is 21ms.[38]

The rise and decay dynamic response of this photodetector could also be described by equation (2) and (3), respectively.[39, 40]

Ilight = Idark + A [exp(t/T1)] + B [exp(t/T2)] (4.4)

Ilight = Idark + A [exp(-t/T3)] + B [exp(-t/T4)] (4.5) 54 where Idark is the dark current, A and B the scaling constants, t is the time when the light was turned on or off, and T1, T2, T3 and T4 are the time constants. T1 and T3 should be related to the carrier generation and recombination processes in the crystal, and T2 and T4 should be mainly related to carrier trapping and release processes. By fitting the experimental data with the two equations, the rise and fall times were calculated to be 12.4 and 17.2 ms, respectively (shown in Figure 4.14),

Thus the two approaches agree well, while the second one gives a more physical insight into the device. This response of the device is faster than perovskite–graphene hybrid photodetectors and

ZnO/Au nanoparticle transparent photodetectors,[41, 42] which reported values of larger than 87 ms and 10 s. The result is also comparable to the state-of-the-art photoconductor results reported in an organic planar-type photodetectors (≥10 ms) and nanocrystals-type photoconductivity photodetectors (about 20 ms).[43, 44]

One of the figure-of-merit factors in photodetectors, noise equivalent power (NEP), is a key parameter defining the device sensitivity. It can be expressed by the equation below:

NEP = (Af)1/2/D* = in/R (4.6)

where A is the effective area of the detector, f is the electrical bandwidth, D* is the detectivity, in is the noise current, and R is responsivity.[45] A smaller NEP value corresponds to higher photosensitivity. To calculate NEP, we measured the noise current using a lock-in amplifier.

Interestingly, the noise current of the devices is only 0.13 pA/Hz-1/2 at 1 Hz. The calculated NEP of the devices is about 2.6×10-14 W at 380 nm, which is better than the perovskite photodiode type

55 photodetector.[45] The very small noise current and low NEP therefore lays a solid foundation for constructing highly sensitive photodetectors.

Figure 4.14 Photocurrent response during on–off illumination switching: rise and decay (inset) curves.

4.4 Conclusion

In summary, a large 5-millimeter sized single crystal FAPbI3 has been successfully grown using a novel liquid based crystallization method. The single crystal FAPbI3 demonstrated a δ- phase to α-phase transition with a color change from yellow to black when heated to 185°C within approximately two minutes. The crystal structures of the two phases were identified and the PL emission peak of the α-phase FAPbI3 (820 nm) shows clear red-shift compared to the FAPbI3 thin film (805 nm). The FAPbI3 single crystal shows a long carrier lifetime of 484 ns, a high carrier mobility of 4.4 cm2·V-1·s-1, and even more interestingly a conductivity of 1.1 × 10-7(ohm·cm)-1, which is approximately one order of magnitude higher than that of the MAPbI3 single crystal.

Finally, high performance photoconductivity type photodetectors were successfully demonstrated

56 using the single crystal FAPbI3, which paves the way for future optoelectronic device applications.

The study of underlying properties of single crystal FAPbI3 is important for improving the perovskite solar cell technology since FAPbI3 has a broad absorption spectrum and good electronic properties.

References

[1] A. Kojima, I. Masashi, T. Kenjiro, M. Tsutomu, Chem. Lett. 2012, 41, 397.

[2] Q. Chen, N. D. Marco, Y. Yang, T. B. Song, C. C. Chen, H. Zhao, Z. Hong, H. Zhou, Y.

Yang, Nano Today 2015, 10, 355.

[3] C. Wehrenfennig, G. E. Eperon, M. B. Johnston, H. J. Snaith, L. M. Herz, Adv. Mater. 2014,

26, 1584.

[4] H. Oga, A. Saeki, Y. Ogomi, S. Hayase, S. Seki, J. Am. Chem. Soc. 2014, 136, 13818.

[5] W. Zhang, M. Saliba, D. Moore, S. Pathak, M. T. Hörantner, T. Stergiopoulos, S. D. Stranks,

G. E. Eperon, J. A. Alexander-Webber, A. Abate, A. Sadhanala, S. Yao, Y. Chen, R. H. Friend,

L. A. Estroff, U. Wiesner, H. J. Snaith, Nat. Commun. 2015, 6, 6142.

[6] S. Kazim, M. K. Nazeeruddin, M. Grätzel, S. Ahmad, Angew Chem Int Ed Engl. 2014, 53,

2812.

[7] M. Liu, M. B. Johnston, H. J. Snaith, Nature 2013, 501, 395.

[8] J. Burschka, N. Pellet, S. J. Moon, R. Humphry-Baker, P. Gao, M. K. Nazeeruddin, M.

Grätzel, Nature 2013, 499, 316.

[9] H. J. Snaith, J. Phys. Chem. Lett. 2013, 21, 3623.

[10] N. G. Park, Mater. Today 2015, 18, 65.

[11] M. Gratzel, Nat. Mater. 2014, 13, 838.

[12] M. A. Green, A. Ho-Baillie, H. J. Snaith, Nat. Photonics 2014, 8, 506.

[13] H. Zhou, Q. Chen, G. Li, S. Luo, T. B. Song, H. S. Duan, Z. Hong, J. You, Y. Liu, Y. Yang,

Science 2015, 345, 542.

[14] N. Ahn, D. Y. Son, I. H. Jang, S. M. Kang, M. Choi, N. G. Park, J. Am. Chem. Soc. 2015,

137, 8696.

[15] S. Pang, H. Hu, J. Zhang, S. Lv, Y. Yu, F. Wei, T. Qin, H. Xu, Z. Liu, G. Cui, Chem. Mater.

2014, 26, 1485.

[16] G. E. Eperon, S. D. Stranks, C. Menelaou, M. B. Johnston, L. M. Herz, H. J. Snaith, Energy

Environ. Sci. 2014, 7, 982.

[17] T. M. Koh, K. Fu, Y. Fang, S. Chen, T. C. Sum, N. Mathews, S. G. Mhaisalkar, P. P. Boix,

T. Baikie, J. Phys. Chem. C 2013, 118, 16458.

[18] N. Pellet, P. Gao, G. Gregori, T. Y. Yang, M. K. Nazeeruddin, J. Maier, M. Grätzel, Angew.

Chem. Int. Ed. 2014, 53, 3151.

[19] W. S. Yang, J. H. Noh, J. J. Nam, Y. C. Kim, S. Ryu, J. Seo, S. I. Seok, Science 2015, 348,

1234.

[20] S. Pang, H. Hu, J. Zhang, S. Lv, Y. Yu, F. Wei, T. Qin, H. Xu, Z. Liu, G. Cui, Chem. Mater.

2014, 26, 1485.

[21] A. Binek, F. C. Hanusch, P. Docampo, T. Bein, J. Phys. Chem. Lett. 2015, 6, 1249.

[22] Q. Dong, Y. Fang, Y. Shao, P. Mulligan, J. Qiu, L. Cao, J. Huang, Science 2015, 347, 967.

[23] D. Shi, V. Adinolfi, R. Comin, M. Yuan, E. Alarousu, A. Buin, Y. Chen, S. Hoogland, A.

Rothenberger, K. Katsiev, Y. Losovyj, X. Zhang, P. Dowben, O. F. Mohammed, E. H. Sargent,

O. M. Bakr, Science 2015, 347, 519.

[24] J. Su, D. P. Chen, C. T. Lin, J. Cryst. Growth 2015, 422, 75.

[25] Y. Dang, Y. Liu, Y. Sun, D. Yuan, X. Liu, W. Lu, G. Liu, H. Xia, X. Tao, CrystEngComm

2015, 17, 665.

[26] M. I. Saidaminov, A. L. Abdelhady, B. Murali, E. Alarousu, V. M. Burlakov,

W. Peng, I. Dursun, L. Wang, Y. He, G. Maculan, A. Goriely, T. Wu, O. F. Mohammed, O. M.

Bakr, Nature Commun. 2015, 6, 7586.

[27] Y. Liu, Z. Yang, D. Cui, X. Ren, J. Sun, X. Liu, J. Zhang, Q. Wei, H. Fan, F. Yu, X. Zhang,

C. Zhao, S. Liu, Adv. Mater. DOI: 10.1002/adma.201502597

[28] C. C. Stoumpos, C. D. Malliakas, M. G. Kanatzidis, Inorg. Chem. 2013, 52, 9019.

[29] M. I. Saidaminov, A. L. Abdelhady, G. Maculana, O. M. Bakr, Chem. Commun. 2015, DOI:

10.1039/C5CC06916E.

[30] S. Y. Leblebici, T. L. Chen, P. Olalde-Velasco, W. Yang, B Ma, ACS Appl. Mater.

Interfaces 2013, 5,10105.

[31] L. Q. Fan, J. H. Wu, Acta Crystallogr. Sect. E 2007, 63, i189.

[32] R. Gottesman, L. Gouda, B. S. Kalanoor, E. Haltzi, S. Tirosh, E. Rosh-Hodesh,Y. Tischler,

A. Zaban, J. Phys. Chem. Lett. 2015, 6, 2332.

[33] C. Quarti, G. Grancini, E. Mosconi, P. Bruno, J. M. Ball, M. M. Lee, H. J.Snaith, A.

Petrozza, F. De. Angelis, J. Phys. Chem. Lett. 2014, 5, 279.

[34] T. S. Horfinyi, T. Pavelka, P. Tutto, Appl. Surf. Sci. 1993, 63, 306.

[35] G. E. Eperon, S. D. Stranks, C. Menelaou, M. B. Johnston, L. M. Herz, H. J. Snaith, Energy

Environ. Sci. 2014, 7, 982.

[36] K. Sveinbjӧrnsson, K. Aitola, X. Zhang, M. Pazoki, A. Hagfeldt, G. Boschloo, E. M. J.

Johansson, J. Phys. Chem. Lett. 2015, 6, 4259.

[37] N. J. Jeon, J. H. Noh, W. S. Yang, Y. C. Kim, S. Ryu, J. Seo, S. I. Seok, Nature 2014, 517,

476.

[38] J. M. Liu, Photonic Devices, Cambridge University Press, The Edinburgh Building,

Cambridge CB2 2RU, UK 2005.

[39] B. Chitara, L. Panchakarla, S. Krupanidhi, C. Rao, Adv. Mater. 2011, 23, 5419.

[40] Z. Sun, Z. Liu, J. Li, G. A. Tai, S. P. Lau, F. Yan, Adv. Mater. 2012, 24, 5878.

[41] Y. Lee , J. Kwon , E. Hwang , C. H. Ra , W. J. Yoo , J. H. Ahn , J. H. Park , J. H. Cho, Adv.

Mater. 2015, 27, 41.

[42] Z. Jin, L. Gao, Q. Zhou, J. Wang, Sci. Rep. 2014, 4, 4268.

[43] B. Mukherjee, M. Mukherjee, Y. Choi, S. Pyo, J. Phys. Chem. C 2009, 113, 18870.

[44] S. Hinds, L. Levina, E. J. D. Klem, G. Konstantatos, V. Sukhovatkin, E. H. Sargent, Adv.

Mater. 2008, 20, 1.

[45] L. Dou, Y. Yang, J. You, Z. Hong, W. H. Chang, G. Li, Y. Yang, Nature Commun. 2014, 5,

5404.

Chapter 5: The identification and characterization of defect states in hybrid

organic-inorganic perovskite photovoltaics

5.1 Introduction

During the last few decades, numerous promising solar cell concepts, ranging from single- crystal silicon to thin-film technologies, have been developed and are being researched intensely by a growing number of scientific groups and companies. Thin film photovoltaic cells based on hybrid organic/inorganic perovskite absorbers, such as methylammonium lead halide

CH3NH3PbX3 (X is Cl, Br and I), have received great attentions recently because of their extraordinary power conversion efficiencies (PCEs).[1–3] Initial PCEs of around 10% have been rapidly replaced by higher values ranging from 12% to over 15% as materials processing and device architectures are improved, making perovskite solar cells one of the most exciting new photovoltaics technologies available today.[4–13]

In contrast to the fast progress in device efficiency, the fundamental understanding concerning the defect properties, which play major roles in controlling the overall device performance of crystalline semiconductors, remains unknown. This is a particularly sensitive issue when a semiconductor such as a perovskite is doped by its own crystalline defects,[14] since their energy levels and spatial distributions are extremely important factors that determine perovskite’s electrical properties and success as a photovoltaic absorber. Specifically, the defects with transition deep in the bandgap will attract the carriers and act as Shockley-Read-Hall non-radiative recombination centers. Thus, the investigation of the defect properties in halide perovskite photovoltaics is urgently needed to provide the guidance for further improvement of the efficiency

62 to approach the Shockley-Queisser limit in which is only constrained by the radiative recombinations in the absorber.[15]

Both the concentration of defects and their effectiveness as dopants can be linked to junction capacitance, making it an invaluable quantity to observe and to correlate with other device parameters. In solar cell science and technology, admittance spectroscopy is commonly used to extract the distribution of defect density and the energetic position of the defects within the bandgap by tracing the junction capacitance.[16–18] In this paper, we conducted the admittance spectroscopy measurements on different device structures to identify the junction capacitance.

Furthermore, the results of temperature dependent admittance spectra reveal the defect energy distribution within the bandgap in the perovskite absorber. We found that the obtained defect distribution of halide perovskite shows similar features to those high efficiency Cu(In,Ga)Se2,[19] which provides the explanation of the remarkable progress in device efficiency of perovskite cells.

5.2 Results and discussion

Admittance spectroscopy involves the measurement of the complex admittance Y(ω, T) =

G(ω, T) + iωC(ω, T) and both the conductance G and the capacitance C are the functions of frequency ω and temperature T. In a typical p-n junction solar cell, the measured capacitance is mainly originated from the charging and discharging of the defect at a location where the defect energy level crosses the Fermi level, as illustrated in Figure 5.1. Trap occupancy is determined by the position of Fermi level. Electron traps above the Fermi level are assumed filled and those below empty. When the frequency of AC voltage signal varies, only states which can release their charge within the AC period will be able to contribute to the admittance. The cut-off frequency is governed by the energy depth of the trap Ed. In Figure 5.1, Eω1 and Eω2 represent measurement

63 conditions that by modulating the AC frequency and the temperature, the traps with different activation energies can respond to the junction capacitance accordingly.

Figure 5.1 A schematic band diagram of p-type semiconductor junction with a single trap level Ed and two measurement energies Eω1 and Eω2 The admittance measurements were conducted on the perovskite devices in which the absorber layer was made by vapor-assisted solution processing (VASP).[6] We have recently developed this approach to prepare the perovskite absorber which is highlighted with its superior film quality. By growing the perovskite film via in situ reaction of pre-coated PbI2 film on the substrate with CH3NH3I vapor, a CH3NH3PbI3 film with full surface coverage as shown in Figure

5.2(a). Thus, the cells can be fabricated without using the mesoporous TiO2 layer, and still show decent performance. Figure 5.2(c) shows the current density-voltages (J-V) characteristics of the device under investigation and the cell has yielded 10.2% power conversion efficiency under 1 sun

AM 1.5G illumination. The corresponding cross-sectional SEM image is shown in Figure 5.2(b).

As indicated by the SEM image, the perovskite absorber has a thickness of roughly 350 nm, constituted with uniform grains across the film thickness, and this confirms the effective

64 intercalation of CH3NH3I vapor into the inorganic PbI2 framework during the vapor-assisted annealing. Such grain structure provides complete surface coverage and significantly reduces the potential shunting paths between the electrodes due to pinholes in the absorber.

Figure 5.2 (a) Surface SEM image of an annealed perovskite layer, (b) cross-sectional SEM image of a perovskite cell and (c) current density-voltage characteristics To assure a precise determination of the junction capacitance and to avoid it from being overwhelmed by other parasitic capacitances within the cell, three different device structures were constructed and compared: (1) Spiro-OMeTAD deposited directly on a fluorine-doped tin oxide

(FTO) glass substrate (Figure 5.3(a)), (2) Spiro-OMeTAD deposited on compact TiO2-coated FTO substrate (Figure 5.3(b)) and (3) a complete device that uses Spiro-

OMeOTAD/perovskite/compact TiO2 (Figure 5.3(c)), where the perovskite absorber is prepared by vapor-assisted solution process. The first and second scenarios model the direct interfaces between (1) electrode and hole transporting layer (HTL) and (2) HTL and compact TiO2, respectively. By comparing the admittance spectrum of each of the different configurations, the true junction capacitance can be resolved and is used to characterize the defect properties of the perovskite layer.

Note that building the perovskite cell in the absence of mesoporous TiO2 is critical to our measurements. Evolving from the dye-sensitized solar cells (DSSC), most perovskite solar cells often have utilized mesoporous TiO2 scaffolds as a porous substrate to grow the perovskite layer.

In addition to serving as structural support for the absorber deposition, the TiO2 scaffold is able to extract and transport photogenerated electrons to the electrode. However, for mesoporous titania, it has been found that the tails of its density of states extends into the bandgap. This results in its increased carrier storing capability; a phenomenon is usually termed as chemical capacitance.[20]

Although the thickness of mesoporous TiO2 is relatively thin (350 nm) in the present perovskite cells compared to that used in traditional DSSC (several μm), the resulting chemical capacitance remains significant. Thus, a transmission line model has to be added into the conventional solar circuit model to reflect the effects attributed to the chemical capacitance and the diffusion- recombination limited transport of mesoporous TiO2 due to its sub-band gap state. More importantly, because of the considerable value of the chemical capacitance of mesoporous TiO2, it is challenging to distinguish whether the capacitance is originated from the p-n junction or the scaffold titania, as seen in the recent report by Bisquert et al.[21] In our case, the transmission line behavior of TiO2 can be likely excluded because of the very thin TiO2 layer employed in the devices.

0.022 8 (a) 0.021 6

G(S) F)

Au -8 0.020 4 Spiro C(10 0.019 2 FTO 0 0.018 103 104 105 106 F(Hz) 6.0 0.030 (b) 0.025 Au 4.5

G(S)

F) 0.020

Spiro -7 3.0

TiO2 0.015 FTO C(10 1.5 0.010 0.0 103 104 105 106 (c) F(Hz) 0.08 Au 4 0.06 Spiro 3

G(S) F) 0.04 PVSK -8 2

0.02 TiO2 C(10 FTO 1 0.00 0 103 104 105 106 F(Hz)

Figure 5.3 Admittance spectroscopy: capacitance and conductance for different device architectures In Figure 5.3, it is visible that the capacitance of the complete device (Figure 5.3(c)) starts to decay at relatively higher frequency (500 kHz), while the capacitances of other systems (without a perovskite layer) all begin to decrease at lower frequency. The frequency responses of conductance from each system (blue curves) are also consistent with their corresponding capacitance. The conductance spectra increase up to its peak value and then decrease for all cases, indicating the dependence of the carrier losses on different perturbation AC frequencies. It is likely that when the AC signal matches to the time constant of the defect states relaxation, losses are

67 reduced as a consequence of decreasing carrier trapping ability of the defects. Thus, the loss peak is a function of frequency and depends on the transition energy of the electrically active defects in the material.[16]

In addition, the conductance at low frequency shows a good agreement with the shunt resistance deduced from the J-V measurement. Based on the comparison of the different systems shown in Figure 5.2, the frequency response of the capacitance of Figure 5.2(c) is distinguished from the others, and it thus likely represents the junction capacitance due to the carrier occupation of defect states in the bandgap of the perovskite, consistent with the observation of the charge accumulation in the perovskite layer previously. This indicates the presence of density of states in the perovskite material, which is in sharp contrast to the absence of this behavior in the organic dyes of dye-sensitized solar cells (DSSC).[21] Note that the junction capacitance is able to be resolved here mainly because of the clear separation between HTL and TiO2 compact layer by the

VASP perovskite and the absence of the mesoporous TiO2. Otherwise, the capacitance from the subgap states in mesoporous TiO2 and from the interface between HTL and TiO2 compact layer can dominate the profile of the frequency response and prevent the information of the junction interface from being completely traceable. For other two scenarios, the capacitances of the HTL

Spiro-OMeTAD is found to dominate the admittance spectra. The capacitance of Spiro-OMeTAD has been observed by other groups,[22–25] and it is believed that the trap states in Spiro-OMeTAD gives rise to the chemical capacitance shown in Figure 5.2(a) and (b), and this chemical capacitance can described as

휕푝 퐶 = 푒2 (5.1) 휕휇푝

where p is the hole density and 휇푝 is the electrochemical potential of the holes in the OMeTAD. 68

To further obtain the defect distribution which results in the observed junction capacitance, temperature dependent admittance spectra were recorded. Figure 5.4(a) displays the capacitance spectra measured at various temperatures (T = 120K - 300K) in the dark from 102 to 106 Hz. An

AC voltage of 30 mV was used as an excitation signal and DC bias was kept at zero during the measurement. Several observations can be made from the spectra. First, the capacitance spectra converge at high frequency to a low capacitance value, which is likely the geometrical capacitance

Cg of the device. In other words, in this high-frequency regime the conductivity of perovskite is too low for the system to respond fast enough to the high frequency AC excitation, making the system behave like an insulator. The capacitance approaches its geometrical value

휀퐴 퐶 = (5.2) 𝑔 푡 where ε is the dielectric constant, A is the area and t is the thickness of the absorber layer.

(a) (b) 5 2.5

4 2.0

T F) T -8

3 1.5

-8

2 1.0 C(10

1 -F*dC/dF(10 0.5

0 0.0 102 103 104 105 3 4 5 Cg 10 10 10 F(Hz) F(Hz) (c) (d) 4

1017 )

) 2

-1

0 -3 16 )(s

2 10

(cm

 T

-2 N ln( 1015 -4 0.004 0.005 0.006 0.007 0.05 0.10 0.15 0.20 0.25 0.30 0.35 -1 1/T(K ) E-E (eV) v

Figure 5.4 (a) Temperature dependence of capacitance, (b) the derivative of the capacitance spectra which shows the characteristic frequency at each temperature (c) Arrhenius plot of the characteristic frequencies to extract the defect activation energy and (d) the defect energy distribution of CH3NH3PbI3 perovskite By determining the absorber layer thickness through scanning electron microscopy, together with the measured values of A and Cg, the dielectric constant of the perovskite absorber has estimated to be 30, which is consistent with other reports.[26,27] Secondly, steps are observed in each capacitance spectrum taken at different temperatures. The transition frequency at each step is related to the rate of carrier emission and carrier capture from the defect states in the bandgap.

In addition, as expected, these features in the spectra shift to higher frequency with increasing

70 temperature according to the temperature dependence of the detrapping time constant. Assuming the occupancy of the states is in thermal equilibrium, which is determined by the Fermi-Dirac distribution, and thus the thermal emission depth of the defect Ea and the characteristic transition frequencies ω0 can be expressed in the equation (5.3):16,17,28

−퐸 −퐸 휔 = 2휋푓 = 2휋푁 휐 휎exp ( 푎) = 2 푇2exp ( 푎) (5.3) 0 0 퐶,푉 푡ℎ 푘푇 0 푘푇

where 푁퐶,푉 is the effective density of states in the conduction band, 휐푡ℎ is the thermal velocity and

휎 is the hole capture cross section.

3 3 2 2 Here we approximate 푁퐶,푉 ∝ 푇 , 휐푡ℎ ∝ 푇 , and 0 is used to comprise all temperature independent parameters. Therefore, in the case of a p-type perovskite semiconductor, the occupation of the hole trap within the space charge region, having an energetic distance Ea = ET -

EV above the valence, can follow the applied AC signal to be charged and discharged at frequencies

ω < ω0. To determine the position of the capacitance step, i.e. the transition frequency ω0, we take the derivative of the capacitance spectra as shown in Figure 5.4(b). This transition reflects the carrier freeze-out in the perovskite absorber layer, which occurs because the frequencies of AC excitation ω > ω0 that the emission rate of the defect is not fast enough to respond.

By converting the x-axis from frequency to energy, which uses equation (5.3) and the fitting of the Arrhenius plot of the transition frequencies in Figure 5.4(c), differentiated capacitance spectra at each temperature are superimposed to yield the energetic defect distribution (Figure

5.4(d)). The calculation of the rescaling factor of the y-axis, including the band-bending and built- in voltage, uses typical values reported in literature.[27] The obtained defect energy distribution exhibits a broad spectrum with the maximum at 0.167 eV, and the moderate density of deeper level

71 defect (above 0.3 eV) is also observed. The span of defect states centered at 0.167 eV can be fitted as a Gaussian distribution (the dashed line) and the integrated defect density is ~1016 cm-3. This defect energy distribution resembles the N1 defect in CIGS absorber, which has be found commonly in Cu(In,Ga)Se2 photovoltaics.[29] The high performance of perovskite photovoltaics can thus be partially attributed to the similar defect properties with CIGS absorbers. To interpret the defect density spectra through admittance spectroscopy of perovskite solar cells, we proposed that they could be possibly ascribed to several point defects with low formation energies, such as lead vacancies, iodide interstitial sites or methylammonium interstitial sites, according to the first- principle calculations.[14] We note that the presence of defect states in the vicinity of heterojunction (e.g. perovskite-transporting layers junctions) can also lead to the charging of interface traps at a location where the trap energy equals the Fermi level EF.[30] Further investigation to bridge the structure and the electronic properties of the perovskite cells will be critical for obtaining complete knowledge of the device, involving the understanding how defect distribution reﬂects the atomic and nanoscale structure of the device, and how it inﬂuences electronic transport, and carrier recombination.

5.3 Conclusions

In this article, we showed for the first time the identification of the defect energy distribution in perovskite CH3NH3PbI3 solar cells employing the temperature-dependent admittance spectroscopy. A comparison of the perovskite containing devices to that of the transporting layer only devices confirms the presence of defect response in the observed admittance spectra. This study is an important step forward in characterizing the defect properties of perovskite absorber and deepens the knowledge of the electronic structure in the perovskite

72 devices. We believe this understanding of the electronic defects will therefore facilitate the improved design of the perovskite materials and the continued development of this emerging photovoltaic technology.

References

[1] N.-G. Park, J. Phys. Chem. Lett., 2013, 4, 2423–2429.

[2] H. Snaith, J. Phys. Chem. Lett., 2013, 4, 3623–3630.

[3] H. Kim, S. Im, and N. Park, J. Phys. Chem. C, 2014, 118, 5615–5625.

[4] J. H. Noh, S. H. Im, J. H. Heo, T. N. Mandal, and S. Il Seok, Nano Lett., 2013, 13, 1764–9.

[5] J. Burschka, N. Pellet, S.-J. Moon, R. Humphry-Baker, P. Gao, M. K. Nazeeruddin, and M.

Grätzel, Nature, 2013, 499, 316–9.

[6] Q. Chen, H. Zhou, Z. Hong, S. Luo, H.-S. Duan, H.-H. Wang, Y. Liu, G. Li, and Y. Yang, J.

Am. Chem. Soc., 2014, 136, 622–5.

[7] H.-S. Kim, C.-R. Lee, J.-H. Im, K.-B. Lee, T. Moehl, A. Marchioro, S.-J. Moon, R. Humphry-

Baker, J.-H. Yum, J. E. Moser, M. Grätzel, and N.-G. Park, Sci. Rep., 2012, 2, 591.

[8] M. Liu, M. B. Johnston, and H. J. Snaith, Nature, 2013, 501, 395–8.

[9] J. H. Heo, S. H. Im, J. H. Noh, T. N. Mandal, C.-S. Lim, J. A. Chang, Y. H. Lee, H. Kim, A.

Sarkar, M. K. Nazeeruddin, M. Gratzel, and S. Il Seok, Nat. Photonics, 2013, 7, 486.

[10] D. Liu and T. L. Kelly, Nat. Photonics, 2014, 8, 133–138.

[11] W. Zhang, M. Saliba, S. D. Stranks, Y. Sun, X. Shi, U. Wiesner, and H. J. Snaith, Nano Lett.,

2013, 13, 4505–10.

[12] J. M. Ball, M. M. Lee, A. Hey, and H. J. Snaith, Energy Environ. Sci., 2013, 6, 1739.

[13] M. M. Lee, J. Teuscher, T. Miyasaka, T. N. Murakami, and H. J. Snaith, Science, 2012, 338,

643–7.

[14] W.-J. Yin, T. Shi, and Y. Yan, Appl. Phys. Lett., 2014, 104, 063903.

[15] W. Shockley and H. J. Queisser, J. Appl. Phys., 1961, 32, 510–519.

[16] T. Walter, R. Herberholz, C. Muller, and H. W. Schock, J. Appl. Phys., 1996, 80, 4411–4420.

[17] R. Herberholz, M. Igalson, and H. W. Schock, J. Appl. Phys., 1998, 83, 318–325.

[18] H.-S. Duan, W. Yang, B. Bob, C.-J. Hsu, B. Lei, and Y. Yang, Adv. Funct. Mater., 2013, 23,

1466–1471.

[19] I. L. Repins, B. J. Stanbery, D. L. Young, S. S. Li, W. K. Metzger, C. L. Perkins, W. N.

Shafarman, M. E. Beck, L. Chen, V. K. Kapur, D. Tarrant, M. D. Gonzalez, D. G. Jensen, T. J.

Anderson, X. Wang, L. L. Kerr, B. Keyes, S. Asher, A. Delahoy, and B. Von Roedern, Prog.

Photovoltaics Res. Appl., 2006, 14, 25–43.

[20] J. Bisquert, Phys. Chem. Chem. Phys., 2003, 5, 5360.

[21] H.-S. Kim, I. Mora-Sero, V. Gonzalez-Pedro, F. Fabregat-Santiago, E. J. Juarez-Perez, N.-G.

Park, and J. Bisquert, Nat. Commun., 2013, 4, 2242.

[22] F. Fabregat-Santiago, J. Bisquert, E. Palomares, S. A. Haque, and J. R. Durrant, J. Appl.

Phys., 2006, 100, 034510.

[23] A. Dualeh, T. Moehl, N. Tétreault, J. Teuscher, P. Gao, M. Nazeeruddin, and M. Grätzel,

ACS Nano, 2014, 8, 4053.

[24] A. Dualeh, T. Moehl, M. K. Nazeeruddin, and M. Gratzel, ACS Nano, 2013, 7, 2292–2301.

[25] V. Gonzalez-Pedro, E. J. Juarez-Perez, W.-S. Arsyad, E. M. Barea, F. Fabregat-Santiago, I.

Mora-Sero, and J. Bisquert, Nano Lett., 2014, 14, 888–93.

[26] A. Poglitsch and D. Weber, J. Chem. Phys., 1987, 87, 6373.

[27] W. A. Laban and L. Etgar, Energy Environ. Sci., 2013, 6, 3249.

[28] A. Jasenek, U. Rau, V. Nadenau, and H. W. Schock, J. Appl. Phys., 2000, 87, 594–602.

[29] F. Pianezzi, A. Chirilă, P. Blösch, S. Seyrling, S. Buecheler, L. Kranz, C. Fella, and A. N.

Tiwari, Prog. Photovoltaics Res. Appl., 2012, 20, 253–259.

[30] H. J. Snaith, A. Abate, J. M. Ball, G. E. Eperon, T. Leijtens, N. K. Noel, S. D. Stranks, J. T.-

W. Wang, K. Wojciechowski, and W. Zhang, J. Phys. Chem. Lett., 2014, 5, 1511–1515.

Chapter 6: Enhanced Stability of Solution Processed InGaZnO Thin-Film

Transistor with Parylene Passivation

6.1 Introduction

Thin-film transistor (TFT) has been widely used in flat-panel displays (FPD) as a backbone of large area electronics. Amorphous silicon (a-Si) TFT showing the field-effect mobility of 0.5 cm2/V∙s is widely used for active-matrix liquid-crystal display (AMLCD) TV. Poly-Si TFTs are being used for high- resolution AMLCD and active-matrix organic light-emitting diode

(AMOLED) displays due to its high mobility over 100 cm2/V∙s. Recently, amorphous oxide semiconductor (AOS) emerged as new type of material for TFT applications, represented by indium gallium zinc oxide (IGZO). Since discovered,[1] IGZO material has shown great potential to replace amorphous-Si (a-Si) or polycrystalline Si (Poly-Si) as an active layer in TFTs, with outstanding mobility, stability, and low cost. IGZO material is traditionally fabricated with magnetron sputter, which is associated with high maintenance cost and low fabrication throughput.

Solution process, including slot-die coating and spray coating, shows the possibility of roll-to-roll processing of IGZO film, which allows for potential large throughput manufacturing and better composition control.[2]

For practical applications, stability of oxide TFT is an important issue. A lot of stress conditions can introduce instability of oxide TFTs, including bias, temperature, illumination etc.

There are a lot of reports illustrating degradation of device performance after stressing oxide TFTs.

Thus, improving stress stability is a key issue for IGZO TFTs.

In this work, we passivated solution processed IGZO TFT with parylene coating to achieve good stability under positive bias stress test. We show that solution processed IGZO TFTs suffer 76 from stress instability initially. Introduction of parylene will passivation IGZO material without degrading device performance. Positive bias stress stability was improved to 0.5 V with parylene passivation. We also fabricated a 5 cm x 5 cm, 144 pixels display to demonstrate effectiveness of

IGZO TFTs fabricated.

6.2 Experimental details

Synthesis of metal oxide solutions. The 0.1 M IGZO solution was synthesized by dissolving 225.6 mg of indium nitrate hydrate (In(NO3)3 ∙ xH2O, Aldrich, 99.999%), 21.3 mg of gallium nitrate hydrate (Ga(NO3)3 ∙ xH2O, Aldrich, 99.999%), and 31.5 mg of zinc nitrate hydrate

(Zn(NO3)2 ∙ xH2O, Aldrich, 99.999%) in 10 mL of 2-methoxyethanol (2ME, Aldrich, 99%). Then,

20-200 μL of acetylacetone (aq) (AcAc, Aldrich, 99%) and 7-70 μL of ammonium hydroxide (aq)

(NH4OH, 28.0% NH3 in water, Aldrich, 99.99%) were added in IGZO solutions. The total mole ratio was 9:1:2 of In, Ga, and Zn. After being stirred vigorously for 24 h at room temperature, the solutions appeared as light yellow transparent and homogeneous. All solutions were filtered through a 0.2 μm PTFE syringe filter (GE, Trevose, PA, USA).

Fabrication of TFTs. The substrates were cleaned in acetone and isopropyl alcohol

++ sequentially. IGZO solutions were spin-coated on SiO2 (1000 Å)/boron-doped (p )- doped Si wafers at 3000 rpm for 30 s and patterned via the direct light patterning process.[3] Spin-coated substrates were prebaked at 100 ℃ for 1 min, and then samples were exposed to DUV (184.9 nm

(10%) and 253.7 nm (90%)) irradiation through a quartz mask for 10 min in air. After the samples were exposed to DUV, they were placed in a 20 mL methanol (CH4O) + 1 mL acetic acid

(CH3COOH) solution for 5 s at room temperature and cleaned by DI water. The samples were then annealed at the desired temperatures (350 ℃) for 3 h. The aluminum source and drain (S/D)

77 electrodes (thickness = 100 nm) were deposited by thermal evaporation. The channel region was defined with a width (W) of 1000 μm and a length (L) of 200 μm.

Fabrication of the display. Glass substrates with 50 mm x 50 mm dimensions were cleaned in acetone and isopropyl sequentially. 30 nm Molybdenum is deposited on glass through magnetron sputtering and patterned with wet etching (etchant is a solution of HCl, HNO3 and

H2O2). A 150 nm of SiO2 gate insulator is deposited on top through plasma enhanced chemical vapor deposition (PECVD). IGZO channel is deposited by sol gel solution process and patterned with HCl etching. A top source/drain electrode of 30 nm Mo is deposited through magnetron sputtering and patterned with lift-off process to complete TFT structure. A 200 nm of ITO is deposited and patterned with lift-off connecting with source as anode for OLED pixels. On top of

ITO, sequential deposition of NPB, Alq3, LiF and Al was completed with thermal evaporation for

OLED. The completed structure was then encapsulated with a glass cavity covering active area with an oxygen/moisture getter inside.

6.3 Results and discussion

IGZO material is synthesized with reported sol-gel chemistry method.[4] TFT is fabricated

Ig -4 10 Drain Current 6 Saturation Mobility

10-6 4 10-8

Current(A) 2

10-10 SaturationMobility (Cm2/V*s) 10-12 0 -20 -10 0 10 20 30 Gate Voltage (V)

Figure 6.1 Transfer characteristics and saturation mobility of as fabricated IGZO TFT.

a) 10-2

10-5

10-8

Initial 100 s 10000 s

10-11 Source-Drain Current (A)

10-14 -10 0 10 20 Gate Voltage (V)

10-4

10-7

Initial 10-10 100 s

10000 s Source-Drain Current (A)

10-13 -10 0 10 20 Gate Voltage (V)

Figure 6.2 Transfer characteristics of solution processed IGZO TFTs before and after 100 s, 10000 s stress test a) without, b) with parylene passivation. Insets show device structure. with a bottom gate, top source-drain contact structure. It shows good performance with carrier mobility of 6 cm2/V∙s, on-off ratio over 108, and subthreshold swing of 100 mV/dec (Figure 6.1).

The device shows low leakage current level. Without passivation, the transfer characteristics of

TFT shows huge drift of around 10 V under 20 V positive bias stress (PBS) for 10000 s (Figure

6.2). The phenomenon is commonly observed in solution processed metal oxide TFTs due to oxygen vacancy deficiency. We then passivated the IGZO TFT with 2 µm of Parylene-C deposited by STS Parylene Coater, which enabled passivation at ambient temperature and low vacuum level.

As shown in Figure 6.2 b), passivated TFT shows substantially enhanced stability performance where Von drift is reduced to 0.5 V. It indicates low permeability of air in parylene film and low damage to the material during deposition and proves compatibility for encapsulation of solution processed IGZO.

Figure 6.3 a) Cross-sectional schematic of AMOLED display driven by IGZO TFTs. b) Schematic of one- transistor, one-diode circuit. c) Mini AMOLED display fabricated with solution processed IGZO TFTs. b) Line scan of the display To demonstrate the application of passivated IGZO TFT, we fabricated a 5 cm x 5 cm, 144- pixel AMOLED display based on the TFT device (Figure 6.3). Solution processed IGZO shows ample driving current for OLED pixels and uniformity potential for large area coating.

6.4 Conclusion

In conclusion, we have been able to improve the stability of solution processed IGZO TFT by parylene passivation, which shows large potential in application of OLED displays. More specifically, the device has demonstrated possibility of large area processing with solution process to enable high quality IGZO roll-to-roll processing.

References

[1] Nomura, H. Ohta, K. Ueda, T. Kamiya, M. Hirano, H. Hosono, Science, 300, 5623, 1269-1272

(2003).

[2] Kim, Si Joon, Seokhyun Yoon, and Hyun Jae Kim. "Review of solution-processed oxide thin- film transistors." Japanese Journal of Applied Physics 53.2S (2014): 02BA02.

[3] Rim, You Seung, et al. "Direct light pattern integration of low-temperature solution-processed all-oxide flexible electronics." ACS nano 8.9 (2014): 9680-9686.

[4] Y. S. Rim, H. Chen, X. Kou, H. Duan, H. Zhou, M. Cai, H. J. Kim, and Y. Yang, Adv. Mater.,

26, 4273-4278 (2014).

[5] M. Fakhri, N. Babin, A. Behrendt, T. Jakob, P. Görrn, T. Riedl, Adv. Mater. 25,2821-2825

(2013).