Synthesis and Characterisation of Perovskite and Chalcogenide Semiconductor Materials for Opto-Electrical Applications

Fiona McGrath

Department of Chemical Sciences

University of Limerick

Academic Supervisor: Prof. Kevin M. Ryan

Submitted to the University of Limerick in fulfilment of the requirements of

Doctor of Philosophy (PhD)

June 2021

Table of Contents

Chapter 1 Introduction ...... 1-1

1.1 Research Questions ...... 1-2

1.2 Hypothesis ...... 1-2

1.3 Scope of the Thesis ...... 1-3

1.4 Bibliography ...... 1-4

Chapter 2 Introduction ...... 2-6

2.1 Electronic and Electro-optic Properties of NCs ...... 2-8

2.1.1 Quantum Confinement Effects ...... 2-10

2.1.2 Spin-Orbit Coupling...... 2-12

2.2 Solution Processing of Nanoparticles ...... 2-12

2.2.1 Hot Injection Method ...... 2-13

2.2.2 Ligand-Assisted Re-Precipitation (LARP) ...... 2-18

2.2.3 Other Perovskite Synthesis Techniquesructures ...... 2-21

2.2.4 Influencers to Nanoparticle Growth and Shape ...... 2-24

2.3 Crystal Structure ...... 2-26

2.3.1 Crystal Nucleation and Growth Mechanisms ...... 2-33

2.3.2 2D chalcogenide semiconductor – Chapter 6 ...... 2-34

2.4 Nucleation and Growth of NPs ...... 2-38

2.5 Surface Chemistry of Semiconductor NPs ...... 2-45

iii

2.5.1 Colloidal Stabilization and Purification ...... 2-46

2.5.2 Ligand Types ...... 2-48

2.5.3 Surface passivation of perovskite NCs – Chapter 3...... 2-61

2.5.4 Ligand Exchange ...... 2-62

2.6 Potential applications of nanocrystals ...... 2-66

2.7 Bibliography ...... 2-67

Chapter 3 Experimental Methods and Characterisation Techniques ...... 3-89

3.1 Synthesis Techniques ...... 3-89

3.2 Characterisation Techniques ...... 3-90

1.1.1 X-Ray Diffraction (XRD) ...... 3-90

1.1.2 Scanning Electron Microscopy (SEM) ...... 3-92

1.1.3 Energy-Dispersive X-Ray Spectroscopy (EDX) ...... 3-93

1.1.4 Transmission Electron Microscopy (TEM) ...... 3-93

1.1.5 UV-Vis Absorption Spectroscopy (UV-Vis) ...... 3-96

1.1.6 Solid State Diffuse Reflectance UV-Visible Spectroscopy (DRS)...... 3-98

1.1.7 Photoluminescence Spectroscopy ...... 3-99

1.1.8 XPS ...... 3-100

3.3 Bibliography ...... 3-101

Chapter 4 Synthesis and Dimensional Control of CsPbBr3 Perovskite Nanocrystals Using

Phosphorous Based Ligands ...... 4-103

4.1 Abstract ...... 4-103 iv

4.2 Introduction ...... 4-104

4.3 Experimental Section ...... 4-107

4.4 Results and Discussion ...... 4-108

4.5 Conclusion ...... 4-121

4.6 Bibliography ...... 4-122

Chapter 5 Colloidal Synthesis of 2D Monolayer Indium Chalcogenide Alloys Nanosheets

5-127

5.1 Abstract ...... 5-127

5.2 Introduction ...... 5-127

5.3 Experimental Section ...... 5-129

5.4 Results and discussion ...... 5-131

5.5 Conclusion ...... 5-141

5.6 Bibliography ...... 5-142

Chapter 6 Germanium-Tin Alloying at B site of Inorganic CsBX3 (X = Br, I) and Cs2BX6

Perovskite Systems...... 6-145

6.1 Abstract ...... 6-145

6.2 Introduction ...... 6-145

6.3 Experimental ...... 6-149

6.4 Results and discussion ...... 6-149

6.5 Conclusion ...... 6-161

6.6 Bibliography ...... 6-161 v

Chapter 7 Additive Effects on the Growth of Inorganic Caesium Germanium Bromide

Perovskite 7-166

7.1 Abstract ...... 7-166

7.2 Introduction ...... 7-166

7.3 Experimental Section ...... 7-169

7.4 Results ...... 7-171

7.4.1 Ammonium Additives ...... 7-175

7.4.2 Phosphorous Additives ...... 7-179

7.4.3 Polymer Additives ...... 7-181

7.4.4 Silicon Additives ...... 7-183

7.4.5 Sulphur Additives ...... 7-185

7.5 Discussion ...... 7-186

7.6 Conclusion ...... 7-188

7.7 Bibliography ...... 7-188

Chapter 8 Conclusions and Perspectives ...... 8-194

8.1 Conclusion and Future Directions from Chapter 4 ...... 8-194

8.2 Conclusion and Future Directions from Chapter 5 ...... 8-195

8.3 Conclusion and Future Directions from Chapter 6 ...... 8-196

8.4 Conclusion and Future Directions from Chapter 7 ...... 8-196

8.5 Answering the research questions ...... 8-197

8.6 Testing the research hypotheses ...... 8-199 vi

vii

Abstract

Colloidal synthesis strategies hold significant promise for a wide range of applications, specifically in optoelectronics as absorbers and phosphors. Colloidal nanocrystals (NCs) are now a fundamental building block in nanoscience due to the surfactant-assisted precision synthesis that provides an acutely narrow size distribution, highly regular morphologies, controllable surface chemistry and enhanced optical properties. Careful design is required to produce nanocrystals that are suitable for a wide range of applications. Building on the knowledge gained from research of metal chalcogenide quantum dots, metal-halide perovskite materials have become a dominant research area in recent years due to several factors, including exceptionally high optical absorption, long carrier lifetimes and diffusion lengths, and high defect tolerance. An example of a popular perovskite nanomaterial is CsPbBr3, which can be manipulated by varying synthesis parameters such as precursors and surfactants, to achieve different morphologies and properties. Herein, CsPbBr3 and a range of other materials are synthesised by solution-based protocols, both colloidally and otherwise. These materials are then systematically characterised structurally and optically to understand the effects of the synthesis parameters.

First, the manipulation of the surfactant system of colloidal CsPbBr3 perovskite nanocrystals is explored in Chapter 4. The typical oleic acid and oleylamine ligand combination employed in the synthesis leads to optically bright NCs with a narrow size distribution. However, they are easily displaced, leading to insufficient stability for real-world application. This study introduces phosphonic acids of various lengths into the system leading to small, near monodisperse NCs and a precise control parameter for particle size and hence material bandgap. Further, the substitution of the primary ligand oleylamine with trioctylphosphine

viii oxide leads to increased reaction yield and demonstrates the potential of phosphorous-based ligands in the perovskite NC synthesis.

Next, Chapter 5 focuses on the substitution of the B cation away from the toxicity of Pb towards

Sn and Ge. Moving up the group 14 column leads to more covalent bonds and a more stable perovskite structure which counteracts the fear that the Ge(II) will oxidise to Ge(IV). However,

Ge(II) does readily oxidise during processing making the material difficult to form using organic and colloidal synthesis strategies. Therefore, this system utilises an aqueous solvent and a reducing agent, which creates a stable perovskite. An array of Ge perovskites are formed successfully by incorporating Sn into the lattice up to 25%.

Chapter 6 continues the theme of alloying elements using organic colloidal synthesis. Indium chalcogenide materials are synthesised using a hot injection synthesis rather than the typical mechanical exfoliation route. The system requires high temperatures to dissolve elemental chalcogenide precursors in oleylamine however, it is a relatively unexplored phosphine free system that produces variations of In2(S1-xSex)3, In2(S1-xTex)3 and In2(Se1-xTex)3.

Finally, Chapter 7 explores attempts to form Ge based perovskites via a colloidal synthesis strategy, resulting instead in the preferential formation of alkali halides. The investigation then provides a detailed overview of the surfactant effect on CsGeBr3 microcrystals. Additive types based on amines, phosphorous, polymers, sulphur and silicon are all known to manipulate perovskite structures in various ways; thus, they are explored systematically and their effects are discussed. Amines and triblock copolymers show the most robust control over the perovskite morphology.

Declaration

I hereby declare that this work is the result of my investigations and that this report has not been submitted in this form or any other form to this or any other university in candidature for a higher degree.

______

Fiona McGrath

(Candidate)

______

Prof. Kevin Ryan

(Supervisor)

This thesis was defended on ______

Examination Committee:

Chairperson: Dr. Shalini Singh, University of Limerick

External Examiner: Dr. Christopher Barrett, Pacific Northwest Laboratories, US

Internal Examiner: Dr. Emmett O’Reilly, University of Limerick x

Acknowledgements

To my supervisor, Prof. Kevin Ryan, my first and biggest thank you must go to you. There are not enough words to express my gratitude, but here is a start; thank you for your excellent guidance, endless expertise, encouragement, support, patience, and wisdom throughout the years. Thank you for always believing in me, but more importantly, thank you for making me believe in me.

To my employment mentor, Dr. Aine Munroe, thank you for the opportunities, support and encouragement. Thank you to all in the Irish Research Council for your support of this research as part of your Employment-based programme and for always being a pleasure to deal with.

Dr. Uma Ghorpade, I could not have asked for a better postdoc! Thank you for all your support, for always making yourself available to answer my questions and for our chats. I am so grateful to have worked with you and learned from you. To all my colleagues at the University of Limerick, thank you all for sharing your expertise, for your support, advice and friendship.

Thank you to Dr. Vasily Lebdev, Dr. Sergey Beloshapkin and Dr. Robbie O’Connell for your help with TEM, Paula Olsthoorn for your SEM help. Thank you to Dr Wynette Redington for XRD training and advice and Dr. Karrina McNamarra for XPS and Raman. Thanks to Ning Liu for advice on optical expertise.

Thank you to Siobhan and Tom Moriarty for the use of their kitchen table in Kilkee. The majority of this thesis was written on it during the summer of 2020 during the Covid-19 pandemic.

To Orlaith and Ronan, undergraduate friends, postgraduate friends, St. Michaels friends - thank you all for being kind, funny, supportive, honest, encouraging, cheeky, loving, moral, caring, all of my favourite things. To my family, thank you all for your love and support. Thank you for always trying to understand my research and being the best parents I could ask for.

To my parents, Catherine and Michael, thank you, thank you, thank you for more than I will ever be able to put into words.

List of Publications

Peer-reviewed journal articles:

F McGrath, UV Ghorpade, KM Ryan; ‘Synthesis and Dimensional Control of CsPbBr3

Perovskite Nanocrystals Using Phosphorous Based Ligands’. J. Chem. Phys. 152, 239901

(2020)

Conference articles and presentations:

F. McGrath; ‘Investigation of Phosphine based Ligands for the Synthesis of Inorganic CsPbX3

Nanocrystals’. 71st Irish Universities Chemistry Research Colloquium (2019)

F. McGrath ‘Investigation of Phosphine based Ligands for the Synthesis of Inorganic CsPbX3

Nanocrystals’ Intel Inclusivity in Tech Conference 2019

F. McGrath; ‘Synthesis of Alkali halide on route towards germanium perovskite NCs’. Intel

Inclusivity in Tech Conference 2018

Poster:

F. McGrath; ‘Synthesis of Alkali halide on route towards germanium perovskite NCs’.

Proceedings of International Conference on Perovskite Thin Film Photovoltaics, Photonics and

Optoelectronics (ABXPV18PEROPTO) Rennes, France. 2018

xii

Nomenclature/Abbreviations

Abb. Meaning Abb. Meaning ACN Acetonitrile IPA Isopropylalcohol APTES 3-Aminopropyltriethoxysilane LARP Ligand-Assisted Reprecipitation BSE Backscattered Electrons LED Light emitting diode CB Conduction Band LMC Layered Metal Chalcogenides CCD Charge Coupled Device LUMO lowest unoccupied molecular orbital CdS Cadmium Sulphide MA Methylammonium CMC Critical Micelle Concentration MPA 3-mercaptopropionic acid CNT Classical Nucleation Theory MPA Methylphosphonic acid CsBr Caesium Bromide MPTMS (3mercaptopropyl)trimethoxysilane CsCl Caesium Chloride NC Nanocrystal CsI Caesium Iodide NIR near-infrared CsX Caesium Halide NMP N-methyl pyrrolidone CTAB Cetyltrimethylammonium Bromide NMR nuclear magnetic resonance DCB Dichlorobenzene NP Nanoplate DDAB Didecyldimethylammonium Bromide NR Nanorod DDPA Dodecylphosphoic Acid NW Nanowire DDT Dodecanethiol OAc Oleic acid DFT Density Functional Theory ODE Octadecane DFTEM Dark Field Tem ODS Octadecylsilane DMC Dimethylcarbonate ODT Octadecanethiol DMF Dimethylformamide OLA oleylamine DMSO Dimethyl Sulfoxide PAs Phosphonic acids DPSO Diphenylsilanediol PbE Lead Chalcogenide DRS Diffuse Reflectance Spectroscopy PCE Power Conversion Efficiency ECL Electrochemiluminescence PEG polyethylene glycol ED Electron Diffraction PLQY Photoluminescence Quantum Yield EDX energy dispersive Xray spectroscopy PVP Polyvinylpyrrolidone EELS Electron Energy Loss Spectroscopy QAS quaternary ammonium salts EQE External Quantum Efficiency QD quantum dots ETL Electron Transport Layer RA reducing agent FA Formamidinium SAED Selected Area Electron Diffraction FF Fill Factor SAM self-assembled monolayers FFT Fast Fourier Transform SE Secondary Electrons FWHM full width half maximum SEM Scanning Electron Microscopy GBL Γ-Butyrolactone STL Solution Temperature-Lowering

GeO2 Germanim Oxide TBP Triblock copolymers X Halide TBPO Tributylphosphine Oxide

H3PO2 Hypophosphorous acid TDPA Tetradecylphosphonic acid HBr Hydrobromic acid TEM Transmission Electron Microscope xiii

HCl Hydrochloric acid TOP Trioctylphosphine HI Hot Injection TOPO Trioctylphosphine Oxide HI Hydroiodic Acid UV Ultravisible HMDS Hexamethyldisilazane UV-Vis UV-Visible spectroscopy HOMO highest occupied molecular orbital VB valance band HRTEM High Resolution TEM VOC Voltage Open Circuit HTL hole transport layer XPS X-ray photoelectron spectroscopy HX Hydrohalic acid XRD X-Ray Diffraction InE Indium chalcogenide IPA Isopropyl alcohol

xiv

Chapter 1 Introduction

At the nanoscale, many materials exhibit interesting physical properties that differ profoundly from their bulk counterpart. Exceptionally high surface area to volume ratio and the possible appearance of quantum confinement effects are key features that affect their properties. [1] The exploitation of tuneable electronic structures, small exciton binding energy, high luminescence efficiency, and low thermal conductivity allow for a wide range of applications, including photovoltaics, LEDs, and transistors laser diodes, quantum computing and medical imaging.

[2-9] Many prototype devices have been produced, indicating how nanotechnology will support an efficient energy economy.

Colloidally synthesised nanomaterials hold significant promise for a wide range of applications, specifically in optoelectronics as absorbers and phosphors. The surfactant- assisted precision of colloidal synthesis provides nanocrystals (NCs) with an acutely narrow size distribution, highly regular morphologies, controllable surface chemistry and enhanced optical properties. Careful design is required to produce nanocrystals that are suitable for a wide range of applications. Building on the knowledge gained from research of metal chalcogenide quantum dots, metal halide perovskite materials have become a dominant research area in recent years due to several factors, including exceptionally high optical absorption, long carrier lifetimes and diffusion lengths, and high defect tolerance. An example of a popular perovskite nanomaterial is CsPbBr3, where manipulation types and concentrations of precursors and surfactants can vary the final crystal properties.

In this thesis, the central aim is to explore the physical effects of manipulating synthesis parameters on perovskite and chalcogenide materials and investigate the resulting properties.

1-1

1.1 Research Questions

This work aims to answer the following research questions:

Research Question 1: Can the ligand shell of perovskite nanomaterials be changed while the size and optical properties remain unaffected?

Research Question 2: Can indium chalcogenide nanomaterials undergo chalcogenide alloying in a system without trioctylphosphine?

Research Question 3: Can lead be substituted for tin and germanium while maintaining the bulk perovskite structure?

Research Question 4: Can germanium based perovskites be confined to the nanoscale using surfactants?

1.2 Hypothesis

Hypothesis 1: A different ligand shell on the perovskite nanocrystal surface may better bond with the NC surface than OA and OLA, resulting in preserved luminescence properties.

Hypothesis 2: Alloying of indium chalcogenide should be possible without using a phosphine- based solvent and still provide a route to optical property control.

Hypothesis 3: A less toxic B cation may hamper the perovskite's optical properties; however, alloying of two B cations may yield properties from both and maintain the efficacy in optoelectronic applications.

Hypothesis 4: An untested type of surfactant might be capable of controlling the formation of germanium based perovskites.

1-2

1.3 Scope of the Thesis

Chapter 2 looks at the theory surrounding nanocrystals and their properties, along with a discussion on the most common solution processing techniques used for perovskite and chalcogenide NCs. Next, there is an analysis of crystal structures and nucleation and growth characteristics for perovskite and chalcogenide materials, the two types of materials discussed in the experimental chapters. The final section of Chapter 2 discusses surface chemistry and ligand effects before briefly reviewing perovskite and chalcogenide materials' potential applications. Chapter 3 explains the synthesis routes used to produce perovskite and chalcogenide materials, including hot injection and the temperature lowering methods, followed by an overview of the characterisation equipment used.

Chapter 4 explores the manipulation of the surfactant system of colloidal CsPbBr3 perovskite nanocrystals by introducing phosphorous based ligands to expand the types of ligands used while maintaining the physical properties. The incorporation of phosphonic acids of various lengths into the system yields small, near monodisperse NCs and a precise control parameter for particle size and hence bandgap. Further substitution of the primary ligand oleylamine with trioctylphosphine oxide increases the reaction yield and demonstrates the potential value of using phosphorous based ligands for perovskite NC synthesis. Chapter 5 employs a phosphine free, hot-injection method for the alloying of indium chalcogenide materials. The phosphine free system produces variations of In2(S1-xSex)3, and so the synthesis is subsequently extended to In2(S1-xTex)3 and In2(Se1-xTex)3.

Chapter 6 substitutes the B cation by alloying Sn and Ge in a perovskite structure to combine

Ge's covalent and stable structure with Sn’s low bandgap and impressive absorption properties.

These alloys are presented, showing an upper limit of Sn incorporation into the CsGeI3 and

CsGeBr3 lattices up to 25% and 50%, respectively. 1-3

Chapter 7 attempts to form Ge based perovskites via a colloidal synthesis strategy. Introduction of TOP into the synthesis to maintain the 2+ oxidation state of Ge instead produces nanorods of CsI, which have not been seen before without a template growth regime. The investigation then moves to a temperature lowering synthesis and provides a detailed overview of the effects of different surfactants on the growth of CsGeBr3 microcrystals. Additive types based on amines, phosphorous, polymers, sulphur and silicon are all known to produce perovskite NC structures. Hence they are systematically explored and their effects are discussed. Amines and triblock copolymers show the most robust control over perovskite morphology. Finally,

Chapter 8 contains conclusive remarks and future directions, then revisits the research questions and hypotheses.

1.4 Bibliography

[1] A. B. Asha and R. Narain, "Nanomaterials properties," Polymer Science and Nanotechnology: Fundamentals and Applications, p. 343, 2020.

[2] M. A. Steiner et al., "High Efficiency Inverted GaAs and GaInP/GaAs Solar Cells With Strain- Balanced GaInAs/GaAsP Quantum Wells," Advanced Energy Materials, vol. 11, no. 4, p. 2002874, 2021/01/01 2021, doi: https://doi.org/10.1002/aenm.202002874.

[3] A. D. Wright et al., "Intrinsic quantum confinement in formamidinium lead triiodide perovskite," Nature Materials, vol. 19, no. 11, pp. 1201-1206, 2020.

[4] I. Devadoss, P. Sakthivel, and A. Krishnamoorthy, "Band gap tailoring and photoluminescence performance of CdS quantum dots for white LED applications: influence of Ba 2+ and Zn 2+ ions," Journal of Materials Science: Materials in Electronics, pp. 1-9, 2021.

[5] L. Mandal, B. Verma, J. Rani, and P. K. Patel, "Progressive advancement of ZnS-based quantum dot LED," Optical and Quantum Electronics, vol. 53, no. 1, pp. 1-20, 2021.

[6] R. Sridevi, J. C. Pravin, A. R. Babu, and D. Nirmal, "Investigation of Quantum Confinement Effects on Molybdenum Disulfide (MoS 2) Based Transistor Using Ritz Galerkin Finite Element Technique," Silicon, pp. 1-7, 2021.

[7] J. C. Norman, R. P. Mirin, and J. E. Bowers, "Quantum dot lasers—History and future prospects," Journal of Vacuum Science & Technology A: Vacuum, Surfaces, and Films, vol. 39, no. 2, p. 020802, 2021.

1-4

[8] C. R. Kagan, L. C. Bassett, C. B. Murray, and S. M. Thompson, "Colloidal Quantum Dots as Platforms for Quantum Information Science," Chemical Reviews, 2021.

[9] D. K. Patel, R. Kesharwani, and V. Kumar, "Nanoparticles: an emerging platform for medical imaging," in Nanoparticles in Analytical and Medical Devices: Elsevier, 2021, pp. 113-126.

1-5

Chapter 2 Introduction

This introductory chapter presents the prerequisite information of this thesis. Beginning with

NC properties, the chapter discusses typical syntheses strategies and the underlying methods and techniques involved in manipulating the formation and growth of crystals. Next, the various crystal structures are considered: a background into nucleation, growth mechanisms and individual characteristics. Then, crystallographic surface chemistry is explored, discussing the role of surfactants in stabilising the structures. Different types of ligands are examined with a view to the ones used elsewhere in this thesis. Lastly, various potential applications of the materials herein are briefly presented.

“One shouldn’t work on semiconductors, that is a filthy mess; who knows if they really exist?”

Wolfgang Pauli, a pioneer of quantum physics, 1931

Nanoscience and nanotechnology are the study and application of matter on the atomic, molecular and supramolecular scale.[1, 2] Its potential is shown in lab-scale research with some examples shown in Figure 2-1. Applications include affordable water filtration, monitoring of climate changes, energy generation (solar PV) and storage (batteries), in medicine (DNA- wrapped single-walled carbon nanotubes (DNA-SWCNTs) and biosensors, and in information storage [3-9]

An important method to form materials on the nanoscale is through colloidal synthesis, a bottom-up synthesis method that forms crystals from nuclei's growth in solution. Colloidal

2-6 semiconductor nanocrystals (NCs) grow from inorganic precursors with an outer layer of organic surfactant on their surface to stabilise them. The core displays unique properties dependent on and controllable by the composition, size, shape and surfactant coating. [10]

Metal chalcogenide NCs, such as CdE, ZnE, CuInE2, CuZnSnE, PbE (E = S, Se, Te), are the most significant semiconductor NCs of the last 30 years since Murray et al. used a hot injection synthesis method to produce CdE (E = S, Se, Te). [11-13]

Figure 2-1 Scale of State of the Art Nanotechnology © Encyclopaedia Britannica, Inc.

Meanwhile, the last decade has produced metal halide perovskite NCs, ABX3 (A =

+ + + methylammonium (MA, CH3NH3 ), formamidinium (FA, CH(NH2)2 ), and/or Caesium (Cs );

+ + + – – – B = Pb2 , Sn2 , and/or Ge2 ; X = Cl , Br , and/or I ). [14] These materials can show a Quantum

2-7

Confinement Effect as their size reduces to the nanoscale; a previously low photoluminescence quantum yield in bulk perovskite increases to almost unity in NCs. [14] Protesescu et al. reported the first hot injection synthesis of perovskite NCs in 2015, four years after the first thin-film solar cell based on the CH3NH3PbI3 absorber. [15, 16] Since then, perovskites have quickly caught up with chalcogenide research with synthesis improvements and various morphologies such as quantum dots (QDs), nanowires (NWs), nanosheets (NSs), nano-plates

(NPs), successfully being prepared. [17-26]

The synthesis of perovskite NCs is similar to that of traditional metal chalcogenides as they both use organic capping ligands to control the growth and passivate defects or voids on the

NC surface.[13] However, metal perovskites are ionic compounds, while metal chalcogenides have a strong covalent character. [13, 27] This difference leads to hugely significant differences in structure, synthesis methods, surface, and optical and electrical properties. While the chemistry of metal chalcogenides is well understood, the long term stability of perovskite NCs remains a challenge due to the highly dynamic nature of the ligand-surface bond. [28, 29]

2.1 Electronic and Electro-optic Properties of NCs

In any array of atoms, atomic orbitals form lower and higher energy bands of molecular orbitals. The most bonding and antibonding molecular orbitals are the band edges. Valence

(outermost) electrons fill the lower energy band resulting in the “valence band” (VB), which is the highest occupied molecular orbital (HOMO). The higher energy band is the lowest unoccupied molecular orbital (LUMO), is empty of any electrons and is called the conduction band (CB). Electrons promoted into the CB can freely move throughout the crystal leading to electric conduction. The space between the CB and VB is called the Bandgap energy (Eg). It is forbidden, meaning that electrons will not move out of the VB unless they have enough energy

2-8 to cross to the CB.[30] So, beginning from individual atoms' bonding with individual atomic orbitals, a semiconductor material is formed, as illustrated in Figure 2-2.

Figure 2-2 Formation of semiconductor bands through the combination of multiple atomic orbitals In a metal, electrons are not tightly bound to the atoms' nucleus, and the bandgap is small (<25 meV), i.e. the VB and CBs are close or overlap. Electrons can move freely from the full VB to the empty CB by absorbing energy from the surroundings. These materials are highly conductive, with conductivity is between 106 and 104 (Ωcm-1). A semiconductor material has a bandgap in which electrons can be excited into the CB by electromagnetic spectrum's energies

(~25 meV – 3 eV). Semiconductors are often black or brown as they only partially absorb visible light with conductivity between 104 and 10-10 (Ωcm-1). Suppose the electrons are tightly bound to the core of the crystal (i.e. strong potential barrier). In that case, the energy bands are further apart or “higher”, resulting in insulating materials that do not absorb any visible light.

Therefore, these are colourless materials. The conductivity of these materials is less than 10-10

(Ωcm-1).[31] Relaxation of the electron back to the ground state releases energy, generally as photon emission, i.e. radiative recombination or luminescence.

The conductivity of a semiconductor can be changed by several factors, including (a) chemical purity, meaning the concentration of impurities or doping, (b) temperature, increasing temperature increases conductivity, and (c) irradiation by light or high energy electron interactions.[31]

2-9

2.1.1 Quantum Confinement Effects

Figure 2-3 Illustration of Quantum Confinement of Semiconductor NCs [32] Quantum Dots (QDs) have dimensions small enough to be comparable with an incident electron wavelength. At this scale, Quantum Confinement occurs, resulting in a blue shift of the absorption wavelength because the nanomaterial bandgap is inversely proportional to the square of the crystal size, as explained in Figure 2-3. The energy levels become discrete, increasing or widening the bandgap between the HOMO and LUMO and increasing the bandgap energy. When excitons occur in nanomaterials, confinement effects trap their wave functions, and so the electronic structure is transformed from the CB to discrete or quantised

2-10 energy states. More energy is required to excite the electron, and more energy is released. As such, smaller quantum dots possess larger bandgaps, as shown in Figure 2-3. [33]

The Bohr radius is a statistically preferred distance of separation between the electron and hole in an exciton. A strong confinement regime occurs when the particle radius is smaller than the

Bohr radius resulting in confinement energy that is much higher than the Coulomb attraction energy (potential energy). Weak confinement occurs when the particle is larger than the Bohr radius and the interaction potential of the two charges is much lower than the Coulomb potential. An intermediate confinement regime occurs in a semiconductor where the electron and hole radii are very different lengths, meaning that particle size can be larger than one but not the other. Each of these quantum confinement regimes results in a separation of exciton energy levels and a blueshift of the material luminescence.[11]

In a semiconductor, electrons are immobilised in the VB unless they gain enough energy to reach the CB. A method for reaching the CB is “photon excitation”, which is the absorption of a photon (light) with higher energy than the semiconductor bandgap. The empty spot left by the excited electron in the VB is called a hole. If the photoinduced hole (+) and exited electron

(-) have Columbic interactions that attract the two together, the pair is known as an “exciton”.

The electron-hole pairs are free to recombine to irradiate photons (luminescence) or disassociate, resulting in conductance. The momentum of an electron-hole pair in an exciton must be close to zero because the original photon's momentum is small due to the conservation of momentum rule. The electron and hole need to have equal momentum in opposite directions.

The pair are known as free charge carriers if the attraction forces between the two are negligible.[30]

In Frenkel exciton, the Bohr radius is small (<5 Å), and the electron is tightly bound/confined within a single lattice constant. The electronic structure of a material containing an exciton is

2-11 described as a “quantum mechanical superposition of states, where it is equally probable that the exciton is associated with any atom in the crystal”. This type of exciton is most likely in alkali halide materials (e.g. CsI).[34]

In semiconductors with a large dielectric constant, the Columbic attraction between the electron and hole pair is reduced due to the electric field's force, resulting in a larger Bohr radius larger than the lattice (40-100 Å). Therefore, the exciton wave function crosses unit cells. These are known as “Wannier excitons”.

2.1.2 Spin-Orbit Coupling

Electrons orbiting an atom have their own axis of rotation and spin with angular momentum.

They exist in pairs with their momentum in opposite directions. These electrons' angular momentum may interact with the electron orbital momentum, causing its energy to split over different levels and give rise to different transition energies. This effect is called spin-orbit coupling. The interaction between electron spin and orbital momentums creates the total angular momentum.

The total angular momentum of individual electrons can merge. The whole atom's total angular momentum can have a range of values, depending on the orbitals' exact occupation. Lighter atoms have a weaker electron angular momentum while having stronger orbital angular momentum.

2.2 Solution Processing of Nanoparticles

Although the ancient Egyptians used metal chalcogenide nanoparticles (NPs) for dying their hair, the understanding of colloidal semiconductor quantum dots (QDs) began in the early

1980s. [35] Michael Grätzel and co-workers reported the first synthesis of colloidal cadmium

2-12 sulphide (CdS) NPs used to avoid the electrodes' photo-corrosion when studying the synthesis of solar fuels through photo-electrochemical devices using semiconductor electrodes.[36]

Solution-based approaches are attractive alternatives to conventional vacuum processing techniques due to lower cost combined with the high degree of control over the material's stoichiometry. Through producing NCs in a colloidal solution or “ink”, low-cost coating techniques (i.e. spin coating, printing and spraying) can be employed, reducing the overall device fabrication costs. [39-41] Solution syntheses eliminate secondary impurities that occur regularly in vacuum-deposited films, dramatically improving device performances.[37, 38]

2.2.1 Hot Injection Method

Murray et al. reported the first Hot Injection (HI) synthesis of cadmium chalcogenide colloidal quantum dots in 1993. La Mer’s theory, discussed below, describes the growth mechanism.

The high quality, monodisperse CdE (E=S, Se, Te) semiconductor NPs had size tuneable band- edge absorption and emissions, as shown by the CdS example in Figure 2-4 a-c.[11] This synthesis system has become the principal synthesis route to high-quality NCs and has led to nanotechnology's expansion into modern devices. [13, 42-44]

2-13

Figure 2-4 Comparison of CdSe and perovskite NCs synthesised through the hot injection route. a) TEM image of CdSe with photographs of samples and their PL luminescence shown in b+c) with the corresponding perovskite images shown from d) to f). [16, 45] The hot-injection method involves the injection of a combination of metal-organic precursors into a hot organic solvent solution containing coordinating phosphine chalcogenides ligands

(trioctylphosphine oxide: TOPO) at high temperatures (> 300 °C). In this method, a rapid burst of nucleation at a high temperature begins immediately following the injection that uses up the monomers, reducing the monomer concentration below the critical concentration and stopping nucleation. This stage is followed by slower, diffusion-controlled growth at a sufficient temperature that facilitates nanoparticles' formation with monodisperse sized particles. Another monomer injection at the same high temperature can focus on the growth of the NCs.[46]

Controlling the NC formation kinetics through adjusting the reaction time and temperature and

2-14 the precursors, ligands, and solvents ratio generates metal chalcogenide NCs with the desired morphology.[47-49] Since then, scientists have developed synthesis protocols that avoid phosphorous-containing chemicals.

These ligands' presence helps rectify the internal defects and yields uniform surface passivation, ensuring well-defined physical properties. Modifications of the utilised surfactants have allowed for decreased environmental impact by removing phosphines and resulted in this hot injection synthesis becoming the most widely employed chemical synthetic technique for

NC preparation.

In 2015, Protesescu et al. applied the hot injection approach to perovskite NCs. However, the reaction conditions and processes are different. In a typical inorganic perovskite colloidal synthesis, metal halide, cation-anion precursors are loaded into a three-neck flask with the non- coordinating organic solvent and heated to 120 °C, evacuated and kept under vacuum for 30-

60 minutes to remove any volatile impurities or water in the system. This step is essential in colloidal synthesis as some reactants may be pyrophoric or sensitive to air. After filling the system with argon, previously well-dried surfactants are then added to dissolve the precursor, and the system temperature is increased to the reaction temperature. Separately, the second reactive cationic precursor is prepared in a similar system, at 120°C under vacuum with surfactants included. This preheated precursor is then swiftly injected into the reaction system.

Figure 2-5 shows the schematic diagram.

The nucleation happens immediately, and quenching occurs by immersing the three-neck flask in a water bath to decrease temperature quickly. Depending on the precursors and surfactants, perovskite NP formation occurs between the temperature ranges of 170-220 °C. After the reaction, the as-synthesised NPs are isolated from the solution and suspended in a nonpolar

2-15 solvent. The results are compared with metal chalcogenide NCs in Figure 2-4 d-f. Section 2.5.3 below will examine the photoluminescence results shown in Figure 2-4.

Figure 2-5 Schematic of reaction set-up for a hot injection reaction. Following quenching, an anti-solvent such as IPA or ethanol that is miscible with the suspension solvent but has an unfavourable interaction with the capping groups is added to the solution. However, the typical OLA and oleic acid (OA) are not tightly bound to halide perovskites, so the antisolvent step is not performed.

Due to the material's ionic nature, the reaction kinetics are very fast in perovskite NCs meaning that investigating the nucleation and growth stages is difficult. Lignos et al. studied the formation kinetics by combining online photoluminescence and absorption measurements in a droplet-based microfluidic platform. The results show that the perovskite NPs form within 5s,

2-16 so examining the two individual stages is difficult.[50] Tuning the reaction temperature has been an effective way to control the NC size. Almedia et al. found that Ostwald ripening, which usually defocuses the size, is suppressed using a critical amount of ligands for metal halide precursor dissolution. This method even removed the need for the ice bath quenching.[18] A different size focusing method uses another metal halide precursor that adjusts the equilibrium between halide anions X- in the NCs and the solution allowing control via thermodynamic equilibrium. [51, 52]

Figure 2-6 Possible formation mechanism of perovskite nanoparticles.

Li et al. developed a microwave synthesis of CsPbBr3 NCs with slower reaction kinetics. [53]

The slower synthesis allowed them to examine the formation mechanism by studying the intermediate stages of growth. They show that a framework of corner-sharing PbBr6 octahedra develops first, followed by infilling Cs+ ions into the interstitial voids. All the voids do not need to be filled for the structure to adopt the perovskite lattice spacing’s, resulting in spherical particles that present broad prominent peaks of the perovskite pattern when examined by XRD.

As more Cs+ incorporates, the XRD peaks sharpen, and the weaker indices become defined, with TEM showing particles with a cube shape, shown in Figure 2-6. However, the PbX2 contains both a cation and an anion meaning that it should be insoluble in the reaction solution.

This may disprove the above formation mechanism.

2-17

While amines act as a ligand, they also form complexes with PbX2: (RNH3)2PbBr4, meaning that the amine competes with the Cs+ ion the A site space in the lattice. [14, 18, 54] As such,

+ Chen et al. propose that pre-formed (RNH3)2PbBr4 is the basic structure of the NC before Cs incorporation. [55] Conversely, Shamsi et al. reported cation exchange as a mechanism to form

CsPbX3 NCs where neither PbBr6 nor (RNH3)2PbBr4 were necessary for perovskite formation.

They also show that an alkali halide, CsX, can be used as the preliminary structure. [56] This is not an exhaustive list of proposed formation mechanisms for perovskite NCs as the exact mechanism remains uncertain.

Disadvantages of the hot injection method include the high temperatures and inert gas needed to grow the NPs and the injection step's complexity for scaling up the system. As such, for metal sulphide NCs, a heat-up method that uses thiols as both ligand and precursors has been developed. Unfortunately, this does not work for other chalcogenides nor perovskite NPs.

2.2.2 Ligand-Assisted Re-Precipitation (LARP)

Continually increasing the concentration of ions dispersed in a saturated solution leads to a supersaturated state, crossing the metastable limit and forming nuclei and crystallising the material. The solution concentration is manipulated by reducing the temperature, evaporating the solvent, or adding an anti-solvent/precipitant. The latter method reduces the solute's solubility in the original solvent, generating a supersaturation force. It is an advantageous technique when the solute to be crystalised is highly soluble, has solubility that is a weak function of temperature, is heat-sensitive, or is unstable in high temperatures.

Antisolvent crystallisation does not use as much energy as evaporative crystallisation, reducing the material's production costs. However, it strongly depends on mixing, the lack of which leads to high local supersaturation and antisolvent induction zones leading to primary

2-18 nucleation. A ligand assisted re-precipitation synthesis (LARP) can produce a much larger quantity of NPs than the hot injection method.[57]

Figure 2-7 Schematic of reaction set-up for a Ligand Assisted Re-Precipitation reaction. Antisolvent crystallisation of perovskite NPs used ligands to assist in the formation: Ligand-

Assisted Reprecipitation (LARP). Here, precursor salts dissolved in a polar solvent, dimethylformamide (DMF) or dimethyl sulfoxide (DMSO), are injected quickly into a non- polar antisolvent, toluene, containing the ligands. NPs form immediately with the assistance of the OA and OAm ligands. These NPs are monoclinic, differing from the hot injection method but show excellent PL (90%) and high photostability. [19]

This method's reaction kinetics might be even faster than that of the hot injection method, so it presents a similar difficulty for studying the reaction mechanism. The absence of amines causes

2-19 aggregation and therefore enlarging of the perovskite NCs. Thus, the amines control the kinetics of formation and the NC size, while carboxylic acid improves NP dispersity and stability. [58] Varying the precursor concentrations caused a photoluminescence redshift as the concentration increased due to larger final precipitate while higher ligand concentration reduced the reaction rate and improved surface passivation. [54] Variation of amine chain length can control the dimensions of perovskite growth. The ammonium cations bind preferentially to sheets of corner-sharing PbBr6 octahedra. Hence, the tuning of their concentration and chain length reduces monomeric units' supersaturation (Either A cations or more PbX6) on the surface, enabling precise control of the layer-by-layer growth of 2D nanoplatelets. [59] The LARP method of NC synthesis allows the synthesis of NCs in ambient conditions, which significantly simplifies the operation process and reduces the cost of preparing the perovskite and enabling large scale synthesis of NCs, opening an avenue to commercialisation.

This method does not work for metal chalcogenide NCs synthesis as the reaction kinetics for the covalently bonded materials are much slower. The technique leads to poor crystallisation of the particles. More energy is needed to form a high-quality NC lattice, so a higher temperature is necessary. [60, 61] Secondly, it is challenging to find a solvent that dissolves the metal cations and chalcogen anions without forming a metal chalcogenide directly.

The disadvantages of the LARP process for perovskite NCs is that the suitable solvents are coordination solvents: N,N-Dimethylformamide (DMF), N-methyl pyrrolidone (NMP), dimethylsulfoxide (DMSO) and N,N-dimethylacetamide (DMAc). [62, 63]

These coordination solvents form intermediates with the precursors, and so it is complicated to remove the solvent following the reaction entirely. The bonding mechanism's ionic nature then means that the NCs decompose gradually as these intermediates remain on their surface,

2-20 compromising the stability. These solvents are also toxic and unsafe in the environment, limiting the commercialisation of the synthesis. The use of non-coordinating solvents acetonitrile (ACN) and γ-butyrolactone (GBL) show success when replacing the coordinating solvents; however, they do not dissolve the precursors as well as the previous solvents and produce crystals with low crystallinity. Using a mixture of coordinating and non-coordinating solvents means that NCs grow with good crystallinity while reducing destructive intermediates' presence on the surface. [64]

As mentioned previously, high precursor concentration in the good solvent leads to precipitation, so the precursor concentration is limited. The low concentration of precursor in the good solvent still requires a high volume of bad solvent to synthesise the NCs leading to waste of expensive non-polar solvents with costly disposal. As a result, the majority of high- performance devices use NCs formed via the hot injection method.

2.2.3 Other Perovskite Synthesis Techniquesructures

Along with the first colloidal synthesis of CsPbBr3, the Kovalenko group first demonstrated the Ion Exchange process to tune the bandgap over the entire visible spectrum. This process was carried out at RT with excess halide and preserved the size and shape of the original

NC.[16] The B cation, Pb, can also be exchanged with other divalent cations (e.g. Zn, Cd, Mn) in the reaction system.[65-67] This method was beneficial as it avoided PL degradation due to additional purification processes.

An ultra-sonication approach is the simplest of the reported methods. Here, the precursors and ligands are dissolved together in octadecene and agitated with an ultrasonication tip. This approach produces cubic phase CsPbX3 with excellent air stability.[68, 69]

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The phase transformation from CsPbX3 to Cs4PbX6 or CsPb2X5 expands these perovskite materials as they have different structures and properties compared to the pure perovskite. As discussed in Section 2.2, briefly, the octahedra are entirely decoupled from one another in

+ Cs4PbX6, and Cs cations fill the holes. [70, 71] CsPb2X5 has a 2D structure formed by slicing

+ along, and inserting Cs into, the crystallographic planes. The structures Cs4PbX6, CsPb2X5, and CsPbX3, are termed 0D, 2D and 3D, respectively.

As described by Figure 2-8, CsPbX3 transforms into one of the two other CsxPbyXz structures under specific conditions. Cs4PbX6 is synthesised in Cs rich conditions in a hot injection reaction and can transform into the 3D analogue by introducing more metal halide postsynthesis. [72] Otherwise, removing Cs through addition of water to the system at an oil/water interface also transformed the 0D to 3D perovskite.[73] This process does result in size shrinkage of the NCs however they are more stable than CsPbX3 produced directly through the hot injection method due to passivation during the water treatment by an oxide shell. Light mediated dissociation converts between CsPbX3 and CsPb2X5, while thiols are reported to convert the perovskite between all three structures. [74, 75]

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Figure 2-8 Interconversion process between different compounds of CsPbBr3 perovskite [76] Morphological control is crucial in controlling the properties of the NCs. [77] The most- reported morphology of CsPbX3 is nanocubes, most likely due to their cubic crystal structure.

Varying the OLA/OA ratio and temperature and the ratio of lead to caesium precursors through the study of thermodynamic equilibrium and kinetics controls the NP size. Tuning the ratio of ligands to PbBr2 precursor and injecting Cs at a low temperature led to a nanosphere morphology.[18, 51] Substitution of the ligand system leads to morphologies tuneable across nanorods, spheres, cubes and plates. Time-dependent results show that orthorhombic nanowires grow from nanocubes to nanosheets, finally to uniform nanowires in 90 minutes.

Further increase of time to 180 minutes results in the formation of large crystals with no more

NPs. Using short acids with long amines tunes the diameter of these nanowires. Nanorods with dimensions less than 100nm are desirable for integrated optoelectric devices, yet more work is needed on their synthesis as reported nanorods are several 100 nm in length.

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2.2.4 Influencers to Nanoparticle Growth and Shape

The control of composition, size and shape of NCs are essential factors widely studied— tailoring the composition tunes the size and shape, bandgap, optical properties, structure, and stability. Parameters that affect this tailoring include time, temperature, surfactants/additives, solvents, seeds/templates, and pH.

Time and temperature are the simplest methods to control the nucleation and growth rate.

Increasing time and/or temperature results in an increased amount of monomers adsorbed onto the surface, leading to larger NCs. A high temperature indicates the energetic movement of molecules and ions, leading to instability of the reaction solution due to a high Gibbs free energy, resulting in the supersaturation rate increasing faster. Hence the nucleation and growth process will be accelerated and shortened. Also, temperature shifts the equilibrium between the

NPs and the participating species by varying the stabilisers or additives activity and chemical state. It is important to note that temperature is a relative concept. A high temperature for one system may be low for another. It is crucial to consider the boiling points of each of the species involved. Increasing reaction time allows for more diffusion of monomers to NC surfaces. It also promotes Ostwald ripening, where smaller NCs are dissolved and consumed into larger

NCs. (Section 2.4.1.3)

Most NPs tend to aggregate into undesirable assemblies due to their high surface energy.

Surfactants and additives can restrict this as they absorb onto a high energy facet and dramatically reduce surface energy, forming layers on the surface that protect against aggregation. With high concentration and strong binding energy to the NCs surface, larger surfactants can slow down the formation kinetics of NCs, leading to a lower adsorption rate of monomers onto the NC surface and reducing the average size of the resultant NCs. The surfactants are generally composed of functional or coordinating groups, key capping agents

2-24 for absorption. Common capping agents are small molecules and polymers. They also control the formation kinetics of the NCs. For example, lowering the surfactant concentration while increasing temperature and time forms larger NCs of Cu12Sb4S13. [78] For perovskite NCs, temperature and species concentration tune the NC size; however, time does not as the particles form so quickly.[50] Pan et al. showed that the lengthening of the carboxylic acid chain length reduces the size of the NCs due to steric hindrance.[79] These effects will be discussed in more detail in later chapters.

Seed-mediated or template-mediated growth provide a preformed surface for growth. The final particles can be homogeneous or heterogeneous structures. The seed size must be minimal so that the NP shape is affected by the seed shape. It is a beneficial method because it lowers the activation energy for precursor monomer addition vs nuclei formation from a homogenous solution.

As explored previously with coordinating and non-coordinating solvents in the LARP synthesis, solvents with different functional groups (e.g. ionic liquids) provide special coordination with the precursor monomer, which can be advantageous for the formation of inorganic NPs under thermodynamic or kinetic control due to the adjustment of the supersaturation increase or depletion rate. The mixture of solvents with different polarity or volatility (boiling point) enables the shape control of inorganic NPs.

+ - Finally, altering the pH through the addition of acid or alkali (H , OH or NH3) adjusts the state of the chemical species in solution and coordination bonding with ions in the precursor monomer solution to form a complex (e.g. HGeX3). Controlling the release rate of ions from the coordination bonding for supersaturation enables control of the initial nucleation rate for shape control.

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2.3 Crystal Structure

The crystal structure is the most significant contributor to the final physical and chemical properties of NCs.[80-85] Figure 2-9 (a-c) shows the structure of traditional metal chalcogenide

NCs as one of zinc-blende (ZnE, CdE, CuInE2), wurtzite (ZnE, CdE, CuInE2), and face-centred cubic (PbE) (the chalcopyrite structure, such as CuInE2, can be derived from the zinc-blende

+ + + structure by substituting the Zn2 with Cu and In3 ). [13, 86-93] These structures have highly covalent, strong chemical bonds, which stabilises them in common solvents like water, isopropanol and toluene etc., and at room temperature, enabling the as synthesised structure to be maintained for a long time in ambient temperature.

It is important to note that these different structures have similar optical properties. The two common structures of CdS, zinc-blende and wurtzite, both have an optical bandgap of about

2.4 eV, which, as a result of the stable structure, gives stable optoelectronic properties that enable durable devices. [94]

Secondly, the ionic radius of constituent atoms does not fundamentally affect the crystal structure of metal chalcogenide NCs. As a result, CdS and ZnS can exist in the same structure even though there is a significant difference in size between Cd (1.55 Å) and Zn (1.35 Å), while the same is true for a S/Se (1.00 Å / 1.15 Å) substitution. [95] Changes to the composition occur through various methods such as cation exchange between Zn2+ and Cd2+ ions, which can obtain nanocrystals shown in Figure 2-1(d-f) with homogeneous, gradient, or core-shell nanostructures without changing their crystal structure. Likewise, S, Se and Te can be substituted without changing the structure despite their large size differences. [96, 97] This means that metal chalcogenides can be produced using various synthesis methods with different compositions to obtain NCs with a homogeneous, gradient or core-shell architectures without changing the central structure. [98] 2-26

Figure 2-9 Metal chalcogenide crystal structures: a) zinc-blend, b) wurtzite, c) face-centred cubic. Metal chalcogenide with d) homogeneous, e) gradient, and f) core-shell architectures. g+h) Cu2ZnSnS4 with wurtzite-kesterite or wurtzite-stannite structures. i) Experimental XRD pattern of wurtzite Cu2ZnSnS4 compared with simulated kesterite and stannite patterns [55] A further property of metal chalcogenides is that alloyed cations' position is flexible in

+ + mutinary materials, for example, the Cu and In in CuInE2, while the anions remain fixed.

This flexibility is the reason that the crystal structure of the NCs is so stable. [93, 99-102] For example, in Cu2ZnSnS4 in Figure 2-9 (g-h), the S positions are fixed while the others can be different but still maintains the same structure. The XRD pattern of Figure 2-9i shows that the cations' movement does not affect the crystallographic structure.

Perovskite NCs differ on these fundamental points. Firstly, the crystal structure has a substantial relationship with the ionic radii of its constituent atoms. The ABX3 perovskite

2-27 structure consists of corner-sharing [BX3] octahedra with the A cation occupying the middle of the cube, formed by eight octahedra suggesting that the ionic radii are critical to the stability of the perovskite structure. [103-106]

The Goldschmidt Tolerance Factor t, shown in Equation 2-1 and the octahedral factor μ, are used extensively to predict perovskites' structural stability, and the ionic radii determine both.

(푟 + 푟 ) 휏 = 퐴 푋 ≈ 1 √2(푟퐵 + 푟푋) Equation 2-1

푟퐴, 푟퐵 and 푟푋 are the constituent ionic radii, and the ions are assumed to be hard spheres.

Therefore t is a measure of how well the spheres pack. The equation describes the ratio between

A-X and B-X's distance and estimates the perovskite's degree of distortion. A value close to unity suggests that the chosen elements will likely form a perovskite with semiconducting properties. Values slightly above or below 1 indicate a perovskite with distorted octahedra or ferroelectric distortions, respectively, while values further from 1 do not result in the perovskite structure forming. [107] Table 2-1 shows this relationship:

Table 2-1 Relationship between the Goldschmidt tolerance factor and the perovskite structure

Goldschmidt TF Structure Explanation

1-1.13 Hexagonal or tetragonal A ion too big, B ion too small

0.9-1 Cubic A and B ions are the ideal size

0.71-0.9 Orthorhombic/Rhombohedral A ion too small

<0.71 Different structure A and B: ionic radii, µ < 0.41

According to the Goldschmidt tolerance factor, only methylammonium (MA), formamidinium

(FA), and caesium (Cs) can stabilise the 3D metal halide framework to offer the desired

2-28 semiconducting properties. Larger or smaller cations favour the formation of lower- dimensional polymorphs (0D, 1D or 2D), with the octahedra becoming edge or face sharing, rather than corner-sharing, resulting in larger and often indirect bandgaps. [27, 108, 109]

These calculations mean that only Cs+, MA+, and FA+ form the 3D lead halide perovskite.

Moreover, Cs+ is slightly smaller, and FA+ is marginally larger than ideal. The 3D cubic phase transfer to the thermodynamically preferred yellow phase (1D orthorhombic or hexagonal structures, respectively) after a short time at room temperature. The solution to this is to mix them both into the A-site Cs1–xFAxPbI3 which prevents the phase transition and forms a stable

3D perovskite structure. [110]

Although the tolerance factor shows the perovskite formability, it is not enough to predict the structural stability. Equation 2-2 presents the octahedral factor μ, which describes the relationship between the ionic radii of the small cation B and the X anion.

푟퐴 휇 = ≈ 0.425 − 0.895 Equation 2-2 푟퐵

The octahedral factor is vital as the octahedron [BX6] is the repeated unit cell in the crystal

2- formation. A value between 0.425 and 0.895 predict the construction of the BX6 octahedra.

[111]

Figure 2-10 shows an overview of different types of perovskites that form when the structural factors or valencies are changed and when vacancies are introduced. Only the a) and c) perovskites are discussed further in this work.

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Figure 2-10 Overview of different perovskite structures. a) Cubic ABX3. b) Antiperovskite with A being a monovalent metal, B a halide and Y a chalcogenide. c) Orthorhombic and tetragonal perovskites. d) Vacant BX3 perovskite. e) Ordered perovskite where M(I) and M(III) metals replace two M(II) metals. f) Vacancy-ordered perovskites where M(III) or M(IV) and vacancies replace the metals. [119] In the past several years, the possible perovskite candidates with stable structures were predicted using the two structural factors. Many of these new combinations look to the B cation

3+ 2+ site. Cao et al. utilised bismuth (Bi ) to replace divalent lead (Pb ) to prepare stable Cs3Bi2I9

QDs with near Photoluminescence Quantum Yield (PLQY) of 37%. They also show successful

3+ electrochemiluminescence (ECL) response. [112] Antimony (Sb ) perovskite Cs3Sb2Br9 NCs were synthesised, showing a PLQY of 46%. [113] Subsequently, mixing of the B cation atoms has developed complex perovskites such as alloyed Cs2(Ag0.6Na0.4)InCl6 with 0.04% bismuth doping synthesised by Luo et al. that has shown white light emission with 86% PLQY and over

1000 h stability showing the benefit of mixing the B cation as well as the A cation. [114]

A difference between the chalcogenide and perovskite structures is the weak chemical bonding in the perovskite materials relative to the chalcogenides. These are ionic, resulting in poor stability in polar solvents such as water, ethanol, DMF. [115-117] This is because the AX salt 2-30 dissolves well in polar solvents, which decomposes the structure. Further, the polar solvent promotes metal ions' movement, distorting the perovskite octahedra. [118] This change in structure results in a change in optoelectronic properties in contrast to metal chalcogenides.

The phase stability of perovskite is also different from that of metal chalcogenides. For example, the orthorhombic phase CsPbI3 (γ-CsPbI3) is thermodynamically stable over the cubic phase CsPbI3 (α-CsPbI3) at a temperature lower than 320 °C. [79] As-synthesized α-CsPbI3

NCs transform into γ-CsPbI3 at room temperature after several weeks; however, thin films' phase transitions occur quickly. [120-123] Besides the ionic bonding of the internal structure, the capping ligands bind ionically. The shell nature is changeable even after the temperature is decreased and the NCs are washed. This c is in contrast to the strongly bound ligands on the surface of metal chalcogenide NCs.[29, 79, 124]

The position of both cations of perovskite NCs are relatively fixed (ignoring the surface dangling bond) compared with that of metal chalcogenide NCs. Unlike the metal chalcogenide

NCs, non-stoichiometry is relatively rare for perovskite NCs, though possible (Figure 2-10f).

[125] The structure is defect tolerant of vacancies. However, non-stoichiometry leads to degradation of the 3D perovskite structure into lower dimensions and therefore changes the optical properties. The balance of charge is crucial for the structural stability of perovskites.

In contrast, the position of the anions is not as fixed as for metal chalcogenide NCs. Anion exchange reactions result in the formation of CsPbBr3-xIx NCs via partial substitution of Br with I in CsPbBr3 NCs and are widely used to tune the optical properties of perovskite NCs.

However, the anions cannot form a stable framework as Sulphur and Selenium in metal chalcogenide NCs. [65, 126-129] Zhang et al. showed a reversible blue shift in photoluminescence when the perovskite is excited by either a laser or an applied bias. This shift suggests that a local electric field generates when excited carriers get trapped on the NC's

2-31 surface, which causes a strain on the lattice, breaking the iodide ion bonds. These then migrate away from the energy source or sublime from the surface. They suggest that this could be the origin of light-induced phase segregation and material degradation observed in many perovskite materials. [130] This anion flexibility is the source of the reduced structural stability of perovskite NCs compared with metal chalcogenide NCs. As a result, under standard operating conditions, phase transformations occur, along with surface instability towards hydration and ion migrations, leading to lower-dimensional counterparts.

-4 In 0D perovskites such as Cs4PbX6, the octahedra [PbX6] are entirely isolated from one another but surrounded by cations. They are, therefore, much more thermally and chemically stable versus the 3D CsPbBr3 analogue; the zero-dimensionality causes strong quantum confinement and strong exciton-phonon interactions. The quantum yield of these crystallites is

~2 orders of magnitude higher than the 3D material when in solid form, and interestingly, they behave like organic materials, contrary to their inorganic nature. These exceptional PL properties' origin has not yet been uncovered, nor even studied in the germanium perovskite counterpart.

One possible explanation is self-trapped excitons and structural distortion. [131] Another hypothesis is that there are intrinsic defect centres in the bandgap due to antisites or vacancies.

[132-134] There are suggestions that there are CsPbX3 impurities embedded inside the compound; however, the luminescence exists even when high-resolution characterisation cannot pick up any CsPbX3. [135, 136]

Meanwhile, 2D perovskite materials may have the composition CsPb2Br5. This material is an indirect bandgap material and so was not expected to show luminescence. Studies of CsPb2Br5 have shown photoluminescence properties hypothesised to originate from the different

2-32 amorphous lead halide ammonium complexes on the material surface due to the presence of oleylamine (OLA). [137]

2.3.1 Crystal Nucleation and Growth Mechanisms

High-quality NCs with low polydispersity are necessary to realise the full potential of semiconductor NCs in any application. As such, the growth mechanism and subsequent assembly must be fully understood. From this understanding, there have been a large number of synthesis techniques investigated.

In general, all nanoscale material synthesis are either “top-down” or “bottom-up”. Top-down methods involve the physical and lithographic principles of minimisation from a sizeable initial substrate. The bottom-up approach assembles ionic, atomic or molecular units into structures with improved structural purity, diverse shapes, sizes, compositions and surface properties.

[138]

In solution, the bottom-up synthesis of shape-controlled inorganic NPs occurs via NP precipitation, hot-injection, seed-mediated growth, polyol approach, template approach, electrochemical syntheses or photochemical synthesis routes. [138-142] Novel properties are achievable from bottom-up synthesis routes due to the ability to directionally control quantum confinement in 0D (isotropic structure), 1D, 2D, 3D (anisotropic structures) NPs.

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Figure 2-11 Typical NP morphologies according to dimensionality.[138] Figure 2-11 shows the most common morphologies in each dimensional set of NPs. In the case of 0D inorganic NPs, typical shapes include spherical, tetrahedral, cubic, or hollow structures.

1D morphologies of inorganic NPs include nanorods, nanoribbons, or nanowires, or hollow structures like nanotubes. Round disks, hexagonal plates or sheets, belts, mesoporous-hollow nanospheres, hollow rings, etc., belong to the 2D shape class of inorganic NPs. 3D morphologies of inorganic NPs are complex hierarchical nanostructures that gradually grow from one parent structure and include assemblies of lower-dimensional materials. [138, 143]

2.3.2 2D chalcogenide semiconductor – Chapter 6

Compared with 3D materials, 2D layered semiconductors have a much larger surface to volume ratio. Almost every atom is exposed in a single layer of 2D material, meaning that the surface area of 2D materials is significantly increased versus other forms of nanomaterials. This high surface area increases the physical and chemical reactivity, influencing the 2D wave function through quantum confinement effects.[144]

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Many materials can form monolayers, including silicene, germanene, black phosphorus, hexagonal boron nitride and Layered Metal Chalcogenides (LMCs).[145-149] LMCs are comprised of two close-packed chalcogenide planes sandwiching a metal layer. They exhibit numerous interesting properties ranging from insulators to semiconductors to metals and find applications in catalysis, photovoltaics, optoelectronics and spintronics. [150-153] Because of the sizeable photoreactive contact area and short migration distance of the generated carriers, the possibility of electron-hole recombination is depressed and photocatalytic performance is enhanced.

There are three main ways to prepare monolayer materials: exfoliation from the bulk, growth on a substrate, and colloidal synthesis. [154-157] Exfoliation can be mechanical (i.e. the tape method), chemically assisted (sonication in a solvent), or purely chemical (intercalation). These methods result in large quantities of flakes but allow for no control over the material's size, shape, or edge. Substrate growth allows more control of shape and size, but as only a single layer can grow, it is volume limited. The high vacuum deposition conditions make it cost- intensive. In this regard, the colloidal approach offers a low-temperature route to nanomaterials with unique morphological control. It is a known alternative for growing 2D materials for use in solution-processed electronic devices. [158-160]

Indium chalcogenide materials are ambient stable and flexible semiconductors with high charge carrier mobility. Indium selenide (InSe) has four covalently bonded Se-In-In-Se atoms with weak Van der Waals bonding between multiples of these quadruple layers.[161] There are several reports of outstanding efficiency of optoelectronic devices based on InSe. [162-

164] Field-effect transistors (FETs) with an active channel of just a few layers of InSe show room-temperature electron mobility values as high as 103 cm2/(V·s).[162] Furthermore, InSe

2-35 has potential for applications in photocatalysis, photovoltaics, strain engineering, and nonlinear optics. [165-168]

Figure 2-12 Common crystal structures of InSe. Unit cells are on the left with the side view on the right. a) ε-InSe, b) β-InSe, and c) γ-InSe. [169] Three common polytypes of InSe occur depending on the stacking sequence of the quadruple layers (ε, β, and γ) shown in Figure 2-12. The ε phase has a hexagonal lattice and has a non- centrosymmetric space group. It is formed by two quadruple layers stacked but with the second layer shifted along the horizontal plane, as shown in Figure 2-12a. The β phase is also hexagonal but belongs to a centrosymmetric space group but is rotated by 60° around the vertical axis, as shown in Figure 2-12b. The γ phase polytype has a rhombohedral lattice and is non-centrosymmetric. It has three layers, with each shifted along the horizontal plane, as shown in Figure 2-12c.

Considering that both the ε and β phases are hexagonal, XRD analysis does not complete the phase analysis. Raman spectroscopy can differentiate between centrosymmetric and non- 2-36 centrosymmetric vibrations when the sample is thin. Therefore, a combination of characterisation techniques determines the phase. [169] This phase determination is of high importance for InSe for optoelectronics because ε-InSe has an indirect bandgap of 1.4 eV. In contrast, both β- and γ-InSe have a direct bandgap with closely matching values (1.28 eV and

1.29 eV, respectively).[170-172] Accordingly, the direct bandgap β and γ phases of InSe are more beneficial for use in optoelectronics. [173]

Density functional theory (DTF) calculations concerning γ-InSe show a transition from direct to indirect bandgap as the number of material layers decreases. The photoluminescence spectra, which is acquired using micrometric spatial resolution, shows that reducing thickness shifts the spectra to higher photon energies.

Compared with other IIIE-VIE (MA, M = Ga, In, E = S, Se, Te) layered materials, InSe has a narrow bandgap in the visible spectrum, and so InSe offers efficient solar absorption.

Remarkably, bending of an InSe photodetector device does not significantly reduce performance, so InSe is an ideal candidate for advanced optoelectronic applications. [174]

Ambient stability is one of the most desirable properties of InSe however, surface vacancies enables decomposition through water intercalation at room temperature leading to p-type doping when in ambient air humidity. Therefore, the minimisation of defects by vacancy passivation using appropriate ligands is highly beneficial.[161]

Due to a mismatch of valence electron numbers, many complex forms exist InSe, In2Se3,

In4Se3, and In6Se7 as well as different phases of In2Se3, such as α-In2Se3, β-In2Se3, γ-In2Se3, κ-

In2Se3, and δ-In2Se3 phases. [175, 176] Usually, In2Se3 crystallises into double layers consisting of Se–In–Se–In–Se sheets stacked together through Se interactions along the c-axis. [177] The

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Meanwhile, In2S3 has three polymorphs: α-In2S3 defect cubic, β-In2S3 defect spinel, γ-In2S3 layered hexagonal. [178] However, only β-In2S3 is stable at ambient temperatures. The defect structures contain voids that can accommodate dopants, either electrical or magnetic, allowing further functionality.[179] It is, therefore, easier to identify In2S3 using XRD analysis.

2.4 Nucleation and Growth of NPs

There are countless reports on the effects of changing reaction parameters on NPs. However, there are no simple rules that determine the final shape of inorganic materials.[138] Reaction parameters alter the basic principle of nucleation and growth, including precursor concentration or super-saturation, reaction temperature, ageing time, and additives.

2.4.1.1 Classical Theory of Nucleation and Growth

The first step in NP growth is nucleation, a short, fast, temporary event followed by a slower growth phase.[180] Nucleation is a pure thermodynamic process and is usually described well by the Classical Nucleation Theory (CNT) (Figure 2-13), which acknowledges that a thermodynamic system needs to minimise its Gibbs free energy (entropy).[181] Solid nuclei generated from a homogeneous supersaturated bulk solution form via primary nucleation.

Conversely, nuclei generated in a supersaturated solution in the presence of other particles with the same or different components (impurities, grain boundaries, dislocations) undergo

“secondary nucleation” or “heterogeneous nucleation”.

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Figure 2-13 Dependence of cluster free energy ΔG, on the radius r according to the classical nucleation theory. A critical cluster size r* has maximum free energy, which defines the first stable nuclei [182] These have a lower energy barrier than primary nucleation due to the stable nucleating sites already present in the system. Heterogeneous nucleation occurs where nucleation sites on solid surfaces contact liquid or vapour, for example, in seed-mediated synthesis. At the preferential sites, such as phase boundaries or impurities, the adequate surface energy is lower, decreasing the activation energy for growth, making growth more likely. Conversely, homogeneous nucleation occurs spontaneously and randomly but requires a supercritical state such as supersaturation. [180, 182]

The Gibbs free energy of a nucleus is the sum of interfacial energy (bonding) between two bodies and the bulk's free surface energy in homogeneous reactions. The Gibbs free energy has a maximum at r*, with the energy barrier: Activation energy ΔG*. For clusters smaller than r*, growth is unfavourable, and dissolution is probable, while larger clusters form nuclei. This is called the Gibbs-Thompson Effect. [180, 182]

An NP's growth depends on two mechanisms: The surface reaction and the monomer diffusion to the nuclei. The precursor monomers precipitate immediately onto the nuclei surface through

2-39 the reaction medium within a high precursor monomer concentration solution. Here the rate- limiting step is the rate of diffusion of monomers to the nuclei surface. Where the monomer concentration is low, the growth becomes a reaction-limited process, and the kinetics of monomer formation is the rate-limiting step. As such, smaller particles should grow faster near larger particles, leading to a narrower size distribution.[180, 182]

2.4.1.2 LaMer’s Theory

Figure 2-14 The principle of NP nucleation and growth according to La Mer’s mechanism [182] LaMer’s nucleation theory (Figure 2-14) proposed that nucleation and growth of NPs occur in two distinct stages. Heterogeneous nucleation follows an initial burst of homogenous nucleation. Three steps that enable control of particle size distribution during growth describe the two-stage process. According to Figure 2-14, stage 1 involves increasing monomers' concentration, which eventually reaches a critical super-saturation Cs level where homogeneous nucleation is infinitely possible. Next, the saturation reaches Cmin, overcoming the activation energy for nucleation leading to a burst of self-nucleation. Due to this nucleation burst, the super-saturation lowers immediately, ending the self-nucleation period allowing heterogeneous nucleation or growth to occur, beginning stage 3.[180, 182]

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This model and its modifications are the only acceptable model of the general mechanism of

NP growth. However, it is not able to predict or characterise the evolution of size distributions.

2.4.1.3 Ostwald Ripening

Ostwald Ripening theorises that the change in solubility of NPs causes growth dependent on the size. Described by the Gibbs-Thompson relation, the solute concentration equilibrium state at a surface is lower for larger particles and higher for smaller ones. As such, smaller particles have increased solubility and surface energy, so, therefore, dissolve, supplying more monomers to the larger particles, resulting in a solute ion flow due to concentration gradient. It is a time- dependent process and is used to explain hollow nanostructures.[138]

Figure 2-15 Top views of the SEM images of (a) MAPbI3 and (b) MABr-treated MAPbI3 films inducing Ostwald ripening. Scale bars, 1 μm. Comparison of (c) UV-vis absorption spectra and (d) X-ray diffraction patterns of MAPbI3 films with and without MABr treatment. Figure 2-15 shows an example of Ostwald ripening in a perovskite film induced by the addition of MABr, which increases the grain size of the MAPbBr3 film. The figure compares the SEM images of the two films and the UV-Vis absorption and XRD data.

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Meanwhile, Digestive Ripening proposes that smaller particles grow at the expense of larger ones. Both processes eventually lead to almost monodisperse dispersions.[180, 182]

2.4.1.4 Finke-Watzky Two Step Mechanism

The Finke-Watzky mechanism proposes that nucleation is slow and happens simultaneously with an autocatalytic surface growth, which is not diffusion controlled. This mechanism is not explicitly proven; however, it is a good fit for many systems where the LaMer mechanism does not apply.[180, 182]

2.4.1.5 Intraparticle Growth

Intraparticle diffusion or ripening considers monomers' diffusion across a particle surface to change the particle's shape with time. (e.g. nanosphere to nanorod) For this to occur, the particle's surface energy is almost equal to the bulk solution's energy, so that the only instability is surface energy differential between the crystal facets. The high energy facets dissolve, and the low energy ones grow, leading to shape change.[180, 182]

2.4.1.6 Adsorption Growth

Surface energy and area determine surface free energy. Minimising this free energy results in the formation of anisotropic NPs. Particles with higher surface energy grow faster than those with lower surface energy to reduce the surface area until the surface equilibrizes. However, the facet of slowest growth determines the final morphology. Surfactants/ligands are used to adsorb selectively to these facets to prevent too much development on the high energy facets reducing the surface tension and altering the surface properties. The selective bonding of surfactants to specific facets occurs through an electronic charge, which adsorbs or coordinates with the surfactant's opposite charge. The concentration and nature of the surfactant determine this roll. [138] See 2.2.4 above.

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2.4.1.7 Coalescence and Oriented Attachment

Figure 2-16 Various types of nanoparticle assemblies. a) PbSe nanocubes assembled in a linear structure and b) hexagonal superlattice by changing the concentration of nanocrystals concentration and ligands, (c) A TEM image of Cu2–xSe nanoplate assembled into a columnar structure, (d) Honeycomb-like networks of ZnTe/CdTe tetrapods are formed by adding poly(methyl methacrylate) (PMMA) to the solution, (e) SEM image of a cluster assembled by CdSe/CdS. [183-186] Coalescence attachment forms crystal assemblies with randomly oriented crystals, while oriented attachment aligns along crystallographic planes. Oriented attachment involves spontaneous self-organisation of particles, sharing a common crystallographic orientation, and joining a planar interface where particles attach at their high-energy facets to eliminate surface energy. Possible intermediate steps may include collision or agglomeration and realignment processes. Once the higher energy facets fill, they fuse crystallographically to form a secondary complete anisotropic particle.[138] Other theories suggest that the particles undergo Columbic attraction to contact the particles, and Van der Wall’s interactions with anisotropic polarizability may occur.[180, 182] Oriented attachment is an uncontrollable process,

2-43 dependent on the reaction conditions. However, control of the reaction conditions allows controlled assemblies such as those in Figure 2-16.

2.4.1.8 Aggregation and Agglomeration

Aggregation is one of the most common difficulties in sample preparation for characterisation.

It is due to the surface energy of the tiny NPs. Two fundamental causes are Brownian and flow motion, which brings the particles together, and a net interparticle attractive force. It results in a direct mutual attraction between particles (Van der Walls, chemical bonding). Usually, aggregation occurs when a reaction solution has a high concentration of small NPs, a high temperature and a slow agitation speed. Increasing agitation speed and the addition of appropriate surfactants can avoid this issue. In cases where the aggregation is desirable, surfactants can control the process, leading to the formation of wires, polyhedral and other isotropic architectures. Agglomeration occurs when clusters aggregate to form the new structure. [138]

2.4.1.9 Self-Assembly

Self-assembly provides a development route for novel functional materials with molecular- and nanoscale-level feature control. [187] The balancing of repulsive interactions and attractive non-covalent forces affects particles' self-organisation and stabilizes discrete self-assembling superstructures. Growth stops when these two forces become equal, meaning that it is a thermodynamically controlled process, unlike in a template synthesis where templates exert kinetic control. Modifying environmental conditions such as temperature, ionic strength, pH and ligands tune properties such as composition and dimensions, as shown in Figure 2-17.

Intermolecular forces such as van der Waals, electrostatic, metal-ligand coordination and nucleic acid base pairing make up the interactions that limit the superstructures' growth. [187]

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Figure 2-17 Factors influencing the self-assembly of nanoparticles and its limits. [187] 2.5 Surface Chemistry of Semiconductor NPs

Surface atoms are a low proportion of bulk materials, and the ratio increases dramatically as the size of the material decreases. The number of surface Pb atoms per CsPbX3 perovskite particle rises from 36 to 223 atoms when the NC size ranges from 2 to 5 nm. In the 5 nm particle, 71% of the Pb atoms are in surface accessible sites. [54]

The main benefit of NPs is the high surface-to-volume ratio. However, surface atoms' un- coordination leads to surface dangling bonds and, therefore, facile loss of surface atoms resulting in surface traps. These traps then capture photo-generated carriers, causing nonradiative recombination, reducing electron-hole diffusion length and lowering the carrier lifetime.[188] These high energy surface sites also facilitate aggregation to other particles.

The primary strategy to deal with high energy sites is to over-saturate anionic dangling bonds with cations to suppress anion bond impact and then passivate the cations with surface acting agents (surfactants) ligands. These passivate the mid-gap (forbidden) electronic states and

2-45 enables luminescence. These can effectively bind via different types of interactions with NC surfaces, especially cations, and cause short-range repulsive forces driving the NCs to be colloidally stable. Any small soluble amphiphilic molecules, macromolecules or polymers with suitable structures are suitable as stabilizing agents. Simultaneously, only those that form a quasi-monolayer on the surface of NCs through covalent, dative or ionic bonds are suitable as surface capping agents or surface ligands. The ligands play multiple roles, from solubility regulation and component availability to post-synthetic minimisation of surface energy and encoding of NP functionality. They affect the growth, colloidal stability, and optical properties of the underlying NPs, making them crucial in nanoscience. [188-190]

2.5.1 Colloidal Stabilization and Purification

Fully formed NCs must be stable enough to be dispersed in a good solvent. There are two fundamental mechanisms of colloidal stabilization, steric and electrostatic. Sterically stabilized particles dispersed in a good solvent (e.g. OLA + Toluene) have negative free energy associated with mixing the hydrocarbon chain and the solvent, which causes the chains to repel one another (Figure 2-18a). As such, they don’t clump together and therefore stabilize the NC dispersion. Conversely, when dispersed in a bad solvent, the chain-solvent mixing energy becomes positive, driving the chains to minimise contact with the solvent. The carbon chains contract and stick to one another, causing the NCs to aggregate together. Typically suitable solvents for hydrocarbon capped NCs are nonpolar liquids (e.g. hexane, toluene, chloroform), whereas non-solvents are polar (e.g. ethanol, acetone, acetonitrile). X-ray scattering experiments and molecular dynamics simulations characterise the interaction between sterically stabilized NCs for various capping ligand lengths and solvent conditions.

Adsorption of charged species onto the surface electrostatically stabilizes NPs. The NC surface charge is balanced by oppositely charged ions located around the surface in an electron double

2-46 layer (Stern layer) around the particle (Figure 2-18b). Suitable solvents for electrostatic stabilisation of NCs are generally those with high dielectric constant, ε, (e.g. formamide (FA),

ε ≈ 110). In these solvents, the NC surfaces cannot approach one another due to the electrostatic repulsion caused by this condensation of counterions in the diffusion region. The addition of low dielectric constant solvents (e.g. toluene) causes the breakup of the mono-charged area around the particle, ultimately resulting in the flocculation of the NCs. The interaction potential between a pair of charge-stabilized NCs, including both the double-layer repulsion and Van der Waals core-to-core attraction, can be characterised using Derjaguin, Landau, Verwey and

Overbeek (DLVO) theory.

Figure 2-18 Stabilization of colloidal NC dispersions through (a) steric stabilization and (b) electrostatic stabilization. Once NC stabilization and the appropriate ligands are known, the purification process needs to be considered. This process usually uses anti-solvent to remove the NCs from the crude reaction solution and separate them from by-products. Unfortunately, this often removes ligands from the NC surface. Selection of the most applicable purification process can remove

2-47 the excess and optimise the surface ligand density, hence improving the charge transfer properties. If an unsuitable process is applied, aggregation and/or degradation occurs.

MCNCs are stable in the typical solvents (acetone, ethanol, methanol) due to the covalent bonding, which means that the amount of antisolvent used and the number of purification cycles does not cause concern of degradation. This stability is apparent through the high yields achieved after the purification of MCNCs. For example, PbS NCs achieve a yield of almost

100%.[191, 192]

In contrast, Perovskite Nanocrystals (PNCs) are ionic, which significantly restricts the purification process. The ionic bonding within the crystal results in a degradation in polar solvents. The ionic nature of ligand-surface bonding allows for ligand mobility and, therefore, rapid removal of the ligand layer during purification cycles.

Two successful antisolvents for PNCs are methyl- and ethyl acetate. [193, 194] Further, the volume ratio of crude solution to anti-solvent is limited to reaction solution:MeOAc as 1:3 or else too much of the ligand layer is removed, causing the cubic structure to revert to orthorhombic.[193] This restriction results in a poor isolation effect resulting in a significant amount of PNCs that remain in the solution leading to a poor NC yield. The reaction solution is centrifuged directly to avoid this effect; however, the yield remains lower than that of

MCNCs. These NCs are then dispersed in a non-polar solvent (chloroform, hexane, toluene), meaning that they only go through one purification cycle. [194]

2.5.2 Ligand Types

L-type ligands are neutral molecules that coordinate to the surface by donating a lone pair of electrons. Examples include amines, thiols, phosphine oxide, neutral carboxylic or phosphonic acids. An X-type ligand may be carboxylates (RCOO-), phosphonates (ROP(OH)O-), and

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- - - 3- thiolates (RS ), as well as inorganic ions (Cl , InCl4 , AsS3 ) interacts with the surface through an anionic moiety. They have an odd number of valence-shell electrons and require an electron from the NC surface to form a covalent bond. Metal to X-type ligand bonds often cleaves heterolytically, forming ionic, closed-shell fragments. Therefore X-type ligands may be neutral radicals binding to neutral surface sites or monovalent ions binding to oppositely charged sites at the NC surface. [195]

Both L- and X-type ligands are nucleophilic, electron-rich molecules that bind to electrophilic, electron-deficient surface sites with Lewis acidity. Other surfaces expose electron-rich Lewis base sites which bond with Z-type ligands such as Pb(OOCR)2 or CdCl2 via the metal atom as two-electron acceptors. Z-type ligands accept a pair of electrons. Therefore, according to the polarity of the NC surface, L- or X- ligands dominate.

2.5.2.1 Monofunctional Molecular Additives

In organo-metal perovskites, ammonium salts fill the A cation site. The most commonly used are methylammonium and formamidinium. However, the addition of secondary ammonium salts with a hydrophobic chain can serve as an efficient water-resistant layer, considerably improving moisture stability. The molecular structure of these additives is hugely important.

Figure 2-19 shows various materials examined. Some molecules have anti-conformation in their alkyl-chain and can induce phase transformation from 3D to 2D perovskites, which is unfavourable for charge transport and device performance due to its insulating nature. [196,

197] Meanwhile, other diammonium halide molecules stay at the lattice grain boundaries and bond to uncoordinated Pb2+ sites due to an anti-gauche isomerization, which increases the activation energy of the phase transformation from 3D to 2D. [197] These molecules interact with the perovskite surface without affecting its bulk properties, somewhat healing surface defects resulting in improved material efficiency and stability. Among these molecules that

2-49 passivate the surface, increased bulkiness increases the steric effects and forms the protective layer, hindering water molecules' absorption.[198, 199]

Figure 2-19 The chemical structures of monofunctional materials that have improved perovskite stability: aniline (A), benzylamine (BA), phenethylamine (PA) 1,3-diaminopropane (DAP), 1,6- diaminohexane (DAH), 1,8-diaminooctane (DAO), tetramethyl ammonium (TMA), tetraethyl ammonium (TEA), ethylammonium iodide (EAI), phenethylammonium iodide (PEAI),iso- butylammonium iodide (iBAI), and hexamethylenetetramine (HMTA)[199] The addition of hydrophobic tertiary and quaternary ammonium salts (QAS) have shown a similar effect; however, they tend not to enter the crystal structure like the diammonium halides. [198] The defect passivation of hybrid perovskite decreases charge trap density and extends the carrier recombination lifetime, increasing devices' performance and improving stability.[200] Benzyltrimethylammonium bromide (BTABr) reduced the turn-on voltage and caused enhanced maximum luminescence and current efficiency in an LED. [201] Meanwhile, a photovoltaic device showed better operational stability when Cetyltrimethylammonium bromide (CTAB) was incorporated. [199, 202] Indirect methods of introducing QAS dimethylammonium bromide (DMAB) to the surface through ligand exchange led to PLQY of

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CsPbX3 increase from 49% to 71%. The relatively short-chain ligand could facilitate carrier transport in LEDs, demonstrating much higher External Quantum Efficiency (EQE) and luminescence than OA/OLA capped NCs.[203]

The replacement of oleylamine with the Quaternary Ammonium Salt (QAS) CTAB leads to improved PLQYs in polar solvents. [204] So far, the most effective QAS has been didodecyldimethylammonium bromide (DDAB), which has been used as post-synthetic surface treatment and an additive during the synthesis.[203] DDAB is a relatively short-chain

(C12) ligand that facilitates carrier transport. Excess OA must protonate OLA further and facilitate its removal before the DDAB is injected; otherwise, quantum wells of (OLA)PbBr4 form.[203] Less bulky QASs ligands surround NCs the most efficiently, improve surface passivation and reduce aggregation compared to bulky ligands.[205] The passivation of undercoordinated Pb-atoms with Br- anions and strong affinity of the ammonium cation DDA+ to the surface A-site causes both electronic passivation and near-unity PLQY values, as well as improved stability and processability of devices. As such, the QAS are of interest as the sole ligand in the CsPbX3 synthesis. At temperatures around 100 °C, DDAX solubilizes the lead precursor by forming an intermediary species PbBr-3.[206] Song et al. showed that tetraoctylammonium bromide (TOAB) improves the PbBr2 solubility; however, it is too bulky to bind to the NC surface. [207] To avoid oleic acid in the system, the Cs-complex is achieved using diisooctylphosphonic acid and toluene. The resulting DDAB-NCs show improved stability towards polar solvents.[206]

Trioctylphosphine (TOP) coordinates with the metal halide precursor and leads to the eventual

NC's surface coordination.[208] The PLQY reached 100% for CsPbI3 samples prepared from this TOP-PbI2 precursor, signifying the complete elimination of surface trapping defects. A

TOP derivative, TOP-oxide (TOPO), also passivates and stabilizes CsPbX3 NCs, even in the

2-51 presence of polar solvents.[209] Tributylphosphine (TBP) improved PLQY when introduced as a post-synthesis treatment. On the surface, it acts as an L-type ligand that conferred a lone pair to a Pb ion, resulting in the removal of trap sites and reduction of charge extraction. [210]

Alkylphosphonic acid work in a similar mechanism and are studied in detail later in this thesis.

[211]

2.5.2.2 Bifunctional Additives

Ammonium functional groups are critical as they directly link with the crystal's A cations; however, other functional groups are needed to connect with other ions. For example, acid groups can interact with the perovskite's halide or with the metal oxide of a charge transport layer in a device. In this way, bifunctional groups become highly beneficial. Figure 2-20 shows a number of these. An amino-acid group can link between surface A cation vacancies on one crystal and the surface halide of another, linking two NPs, or linking the perovskite to the next layer of a device. This linking applies to grain boundaries of thin films as well as for the assembly of NPs.

5-Ammoniumvaleric acid (5-AVA) was the first such molecule to be included in a perovskite film, forming a mixed cation perovskite with lower defect concentration and better pore filling

+ - due to COOH forming hydrogen bonds with terminal -NH3 or I from PbI6 octahedra. It also improved contact with the metal oxide TiO2 electron transport scaffold, resulting in improved performance and stability for >1000 h. [212, 213] Since then, this architecture has shown stability of one year in standard conditions. [214]

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Figure 2-20 The chemical structures of bifunctional materials that have improved perovskite stability: 5-Ammoniumvaleric acid (5-AVA), 3-aminopropyltrimethoxysilane(organic silane), butylphosphonic acid 4-ammonium chloride (4-ABPACl),3-(decyldimethylammonio)-propane-sulfonate inner salt (DPSI/SB3-10),4-ethylamine phenylphosphate disodium salt (EAPP), 3-aminopropaniocacid (C3), 3- (1-pyridinio)-1-propanesulfonate (NDSB201), and 3-[2-hydrox-yethyl(dimethyl)azaniumyl] propane- 1-sulfonate (NDSB211). However, these amino acids' conductivity is low, resulting in a low fill factor (FF) in photovoltaic devices, resulting in poor Power Conversion Efficiency (PCE) of sunlight into electricity. Hu et al. designed a bifunctional conjugated organic molecule 4-(aminomethyl) benzoic acid hydroiodide (AB), which, compared with the previously mentioned amino-acids, show improved PCE of 15.6% and improved FF by facilitating charge transport. [215] Post synthetic introduction of ligand 2,2’-iminodibenzoic acid (IDA) showed passivation of CsPbI3

NPs surfaces due to the much stronger binding between CsPbI3 and IDA compared to OA. The

IDA-NCs displayed a PLQY of >95% and a slightly improved lifetime. This IDA ligand could separately bind to the two Pb atoms via the coordination of double carboxyl groups, leading to considerable binding energy (1.4 eV) compared to the single carboxyl group in OA (1.14 eV).

[79]

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Amino-silane materials work similarly, specifically showing optimized interface band alignments and enhanced charge lifetimes.[216-220] 3-aminopropyltriethoxysilane (APTES) has been added to the surface by both in-situ capping and post-synthetic methods and attaches to the surfaces of the NPs through hydrogen bonds between silanol and bromide dangling bonds. It, therefore, played an active role in the passivation of surface defects. This ligand is a precursor for a silicon oxide (SiOx) shell around the NP. Alkylcarboxylic acid ω-ammonium additives can also be used as templates. However, they are incorporated into the lattice forming

2D perovskites.

-6 Figure 2-21 Schematic illustration of two neighbouring grain structures; the PbI4 octahedra are shown in red, crosslinked by butylphosphonic acid 4-ammonium chloride (4-ABPACl) hydrogen- bonding interactions (O–H...I and N–H...I) of the iodide from the iodoplumbate complex with the + phosphonic acid (–PO(OH)2) and the ammonium (–NH3 ) end groups of the 4-ABPACl species.[221] Phosphonic acid ammonium additives (4-ABPACl) do not enter the perovskite lattice but bond with the perovskite through hydrogen bonds from both the phosphonic acid and the ammonium groups. Therefore this additive can passivate the surface thoroughly, rendering it immune to moisture.[221] A photovoltaic device PCE improved by 90% with the 4-ABPACl additive.[221] They also aid in crosslinking the separate crystals leading to improved film morphology, as shown in Figure 2-21. Here the ammonium species attach to the peripheral 2-54 octahedra using ionic bonds as standard, with the terminal phosphonic acid functional groups protruding from the surface.[29, 221] This phosphonic acid headgroup can form hydrogen bonds with the halide anions of another perovskite crystal grain, enhancing a film's morphology.[221]

Three different sulfobetaine-zwitterions have proven capable of stabilizing the α-phase (cubic) of CsPbI3 by slowing the crystallization of CsPbI3, forming nanoparticles of CsPbI3 instead of a complete film. The reduced size increased the particles' surface energy, stabilizing the phase, which resulted in a photovoltaic device with 11.4% PCE.[222] Sulfobetaine zwitterions contain a positively charged quaternary ammonium group and a negatively charged sulfonic group. 3-

(decyldimethylammonium) propane-sulfonate inner salt (DPSI) and 4-ethylamine phenyl phosphate disodium salt (EAPP) are two sulfobetaines that have shown to improve crystallization and passivate hybrid perovskite defects yielding improves PCE and moisture stability in devices. [223, 224] 3-(N,N-dimethyloctadecylammonio)-propane-sulfonate was used as the OA and OLA substitute for the synthesis of CsPbBr3 NCs and formed simultaneous coordination with both cations and anions on the surface. [225] This bonding is shown schematically in Figure 2-22.

Figure 2-22 Schematic illustration of (a) conventional OLA + OA capping ligands and (b) zwitterionic (C3-sulfobetaine) capping ligands.

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Preliminary work with thiol introduction to the inorganic perovskite was through ligand exchange. When thiols are added to CsPbBr3 for recrystallisation, 0D Cs4PbBr6 is formed rather than CsPbBr3. Short-chain thiols produced more stable 0D perovskites than those with long chains.[226] Baek et al. used ligand exchange to introduce thiol to the NC surface. The thiol resulted in improved photo-stability to UV radiation in both solution and thin film. [227]

A synthesis using an amino-thiol ligand system results in the formation of CsPb2Br5 nanowires, while a thiol-acid ligand system forms CsPb2Br5 nanosheets. The most promising thiol- containing ligand is cysteamine (1- or 2-aminoethanethiol). Cysteamine is a small bifunctional molecule containing both thiol and amino groups. It is widely used as a chelating ligand in coordination chemistry, biochemistry and forms functional self-assembled monolayers

(SAMs) on metal surfaces such as gold, silver and copper.[228-230] Ionization and solvation of the amine-group in aqueous solutions control the cysteamine linkage's binding properties to the solution phase species.[230] For the hybrid perovskite, cysteamine forms an interconnect between MAI and PbI2, forming a perovskite that is highly resistant to water immersion.[231]

This new ligand is potentially game-changing for the stability problems of hybrid perovskites.

Building on this hydrophobicity, Men et al. reported an aqueous synthesis of germanium perovskite using cysteammonium halide ligand through the reaction of cysteamine with the hydrohalic acid of choice. [232]

2.5.2.3 Multifunctional Molecular Additives

Targeting both the improved stability and the higher PCE of perovskite solar cells has led to a multifunctional molecule design that combines different functional groups' advantages. The

Pb2+ ion interacts with the Lewis acid while polar molecules that contain O-, S-, or N- donor groups interact with Lewis bases. It has been known for some time that the interaction between the weak Lewis acid PbI2 and dimethyl sulfoxide (DMSO) forms a beneficial intermediate

2-56 adduct in the perovskite synthesis, leading to the incorporation of DMSO, NMP, Urea or its derivatives into the DMF solvent system.[233-235] In general, Lewis bases' addition leads to improved perovskite film morphology with fewer traps and a longer lifetime. The introduction of Lewis bases improved the interface between the electron transport layer (ETL) and perovskite. Studies then looked to incorporate Lewis acids at the interface between the perovskite and the hole transport layer (HTL), where charge accumulates due to the undercoordinated halide anion, which acts as a hole trap. Abate et al. use iodopentafluorobenzene to passivate excess iodine on the perovskite surface and decrease nonradiative recombination.

[236] The addition of these acids and bases lead to higher charge collection efficiency, device performance as high as 20.62% and long-term ambient stability exceeding 2880 h without encapsulation for a passivated Ruddlesden-Popper/3D heterostructure perovskite film. [237]

Given the excellent properties of the interaction of electron-rich atoms (e.g., sulphur, thiophene

π-spacer, carbonyl, etc.) with perovskites, molecules containing such functional groups are being designed as additives to the perovskite synthesis. The systematic study of passivated functional groups' structures shows that carboxyl and amine groups can heal charged defects via electrostatic interactions while aromatic structures can reduce neutral iodine related defects.[238] D-4-tert-butylphenylalanine (D4TBP) combines 4-tert-butylphenyl, amine and carboxyl, functional groups, showing excellent passivation with prolonged PL lifetime and very low open-circuit voltage (VOC) deficit. [238]

Regarding the unreacted precursor, a diboron compound C12H10B2O4 (B2Cat2) can selectively react with FAI to reduce defect densities and mitigate nonradiative recombination between the perovskite absorber and the HTL and improve charge carrier extraction. Fluorinated molecules have excellent water vapour and heat resistance. Fluorinated perylenediimide (F-PDI) contains carbonyl groups that chelate with noncoordinating Pb2+ and the fluorine, forming hydrogen

2-57 bonds with MA. F-PDI is hydrophobic and conductive, resulting in better stability towards humidity while also improving defects. Imidazolium iodide dopants significantly increase operational stability due to the hydrophobic nature of the salt.

2.5.2.4 Polymer Matrix Encapsulation

Polymers protect perovskites from water because their dense polymer matrix can endow them with water-resistant features and prohibit an ion exchange.[77] Polystyrene (PS) is ultra hydrophobic and is used widely in industrial applications. PS micro-hemispheres/beads/fibres were introduced as the encapsulation shell for CsPbX3 to form hybridized composite particles with superior stability even when exposed to water. [239-241] This method has led to CsPbX3 perovskites with 80% fluorescence stability after 60 h. Long-chain insulating polymers, such as polyethylene glycol (PEG), are chemically inert towards the perovskite precursors and are also electrical insulators that guarantee the effective transport photogenerated carriers through channels in the perovskite. A PEG scaffold has stabilised the perovskite film, rendering a device resistant to moisture and displays a strong self-healing effect.[242] Ethyl Cellulose (EC) has been employed in a similar architecture, increasing the stability but decreasing the performance.[243] A polar polymer polyvinylpyrrolidone (PVP) has O atoms that form hydrogen bonds with H in the hybrid perovskite, wrapping crystal grains of perovskite in a protective air barrier. [244]

Triblock copolymers (TBPs) are templates for CsPbX3 nanowires in a composite ink where it is possible to define the nanowire alignment. This template approach expands device applications to optical nanocomposites that exhibit highly polarized absorption and emission properties leading to optical storage, encryption, sensing and full-colour display devices.[245]

Organic-inorganic perovskites have benefited from the same templated design where MAPbI3 and a block copolymer were spin-coated into a thin film which crystallised into various

2-58 nanoforms with dimensions between 40-72 nm leading to enhanced photoluminescence properties of the film.[246] Meanwhile, a copolymer dramatically improved the long term fluorescence intensity of CsPbBr3 quantum dots under thermal and chemical stressors as well as in both pure and salt-water immersion.[247, 248]

2.5.2.5 Inorganic ligands

In 2009, Kovalenko et al. completely changed the traditional view that capping ligands must

+4 be organic molecules by introducing metal chalcogenide complexes (MCCs) such as SnS4 ,

2- 3- 4 In2Se4 , SbSe4 , MoS4 as inorganic ligands.[249, 250] Solution phase inorganic ligand exchange reactions (ILERs) with a large variety of MCC-NC pairs is now a viable method to achieve high colloidal stability and improved conductivity and solubility in common polar solvents with high dielectric constant (e.g. water, methanol). These electrostatically stabilise suspensions, and NCs retain their size and shape, as well as their quantum-size dependent photo-physics. Solid-state exchanges have been demonstrated and patented by Talapin and co- workers, the group in which Kovalenko et al. first published the breakthrough. [251, 252] The ligand replacement can be characterised by infrared nuclear magnetic resonance (NMR) and/or x-ray photoelectron spectroscopies, while the photoluminescence or conductivity measurements indicate any property changes of the original material.[253]

Since this MCC discovery, there has been an extensive exploration of other inorganic moieties that can act as capping agents. [254, 255] Candidates including metal-free chalcogenides and

2- - 2+ - 2- - 2- hydro-chalcogenides (S , HS , Se , HSe , Te , and The ), mixed chalcogenides (TeS3 ), as

- - well as OH and NH2 . Furthermore, weaker nucleophiles’ such halides (Cl-, Br-, I-) and pseudo- halides (CN-, SCN-, N3 -) can exchange as the ligands directly.[250, 256, 257] Meanwhile, halometalates, such as methylammonium metal halide and simple metal halides (M = Pb, Cd,

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Zn, In, Fe, Sb, X = Cl, Br, I) may replace the organic ligands. Unlike the MCCs that act as hole traps that quench the PL, halometalates do not affect the PLQYs.[258]

MCC-functionalization of QDs has enabled breakthroughs in the electronic characteristics of

QD solids. This reproducible and straightforward approach is becoming increasingly popular, allowing solid-state devices with band-like transport, high electron mobility and high photoconductivity to be realised.[259] For example, photovoltaic power conversion efficiencies up to 12% and improved thermoelectric characteristics.[260]

Acids such as HBF4 and HPF6 are exceptional ligand removers that show that ligands' protonation is a critical step in their removal.[250] Strong acids such as these etch the inorganic material. However, Meerwein’s Salt (Et3OBF4) is effective while avoiding etching.[261] This removal results in “naked” NCs, all of which were dispersible in Dimethylformamide (DMF).

This process results in highly conductive films of NPs such as PbSe.[261]

However, complete ligand removal is usually the first step of a two-step process that allows the functionalisation of the NC surface with ligands that would not otherwise have taken the

3- 2 place of the native ligands, such as oxoanions, e.g. PO4 and MoO4 . The binding affinity of oxoanions is stronger with oxide NCs while chalcogenidometalate ligands have a stronger binding affinity towards metal chalcogenide NCs.

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2.5.3 Surface passivation of perovskite NCs – Chapter 3

Table 2-2 Degradation pathways in the perovskite the family

Degradation Mechanism

Light Charge trapping at the surface defects

Moisture Decomposition, phase transition, or phase segregation

Heat/Thermal Abnormal synergistic effects of A cation or X halide vacancy defects

The perovskite family of materials suffer from analogous degradation pathways: light, moisture, and thermal-induced degradation leading to phase segregation shown in Table 2-2.

Hydration of hybrid lead halide perovskite is thermodynamically preferred, where moisture

absorption breaks down the perovskite into HI and CH3NH3 gases within 3 hours under 70-

80% humidity.[262] As such, methods to improve hybrid perovskites' moisture stability are likely to enhance other perovskites' stability towards oxidation.

Since the first reports of colloidal CsPbX3, OLA and OA have become the most commonly used ligands for perovskite materials. However, their long chains and insulating features significantly weaken the conductivity and block charge transfer in devices. How to replace them with other ligands to enhance their properties without destabilizing remains the key challenge preventing the development and applications of CsPbX3 NCs. Generally, the acidic oleic acid ligand protonates the (OLA) ligand, leading to the ligand layer's degradation on the inorganic perovskite nanocrystals' surface, therefore poor stability against moisture, UV light and high temperatures.[204] While the material and the mechanism of degradation are different, the methods to combat both may be similar.

De Roo et al. showed that surface Cs and Pb ions bind to either oleate or Br- while surface halides bind to oleylammonium ions via hydrogen bonding or electrostatic interactions.

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Therefore, ligands that would remain attached during processing would form strong metal- ligand bonds or stable hydrogen bonds with halide ions. Oleylammonium ions act as capping agents by substituting Cs+ cations preferentially over hydrogen bonding to the halide ions.

[263] Phosphonic acids (PAs) bind strongly to divalent cadmium and lead cations, and they bind strongly to lead perovskites via hydrogen bonding. [221, 264, 265] The presence of

Tetradecylphosphonic acid (TDPA) produced colloidal perovskite NCs withstood washing in

H2O and is suitable for use in room temperature synthesis. [266, 267] Wang et al. [268] found that a phosphinic acid could stabilise the alpha phase of CsPbI3. They saw through NMR and

Diffusion Ordered Spectroscopy (DOSY) that the phosphinic acid exists in an ion pair with oleylammonium and does not play a surfactant role in stabilizing the α-CsPbI3, meaning that only the amines are the predominant surface ligands. Separately, washing with TOP stabilised

α-CsPbI3, leading to over a month's quantum yield efficiency stability.[269] In related studies, octylphosphonic acid (OPA) dramatically enhanced the perovskite stability towards polar solvents, producing an LED with an EQE of 7.74%. [270] Therefore phosphonic based ligands should be an attractive alternative to carboxylic acid ligands.

2.5.4 Ligand Exchange

Surface ligands with L- or X- type headgroup and long hydrocarbon tails (e.g. OLA, OA, TOP,

TOPO, dodecanethiol (DDT)) are excellent agents for control over NC nucleation and growth kinetics. However, L- or X- type ligands hinder charge transport between NCs and are often too large to heal all of the defects on the surface of MCNCs. Therefore, they must be replaced by ligands better suited to the end application.

For example, the X-type OLA and OA ligands are necessary for CsPbBr3 synthesis. They dissolve precursors and control the growing NCs however, the proton exchange between the two while coordinated induces their detachment from the surface. Therefore, the ligand density

2-62 at the surface typically decreases after isolation and purification. Re-adding a small amount of ligands is a prerequisite to achieving good colloidal stability.[29] However, excess amine destabilizes the PNCs, irreversibly degrading into Cs4PbX6.[271] Further, these ligands do not encourage dispersibility in polar solvents, thus limiting the application of PNCs.

Figure 2-23 An overview of all significant NC surface chemistries derived from the NCs capped with native organic capping ligands. All presented examples entail a solution-phase ligand exchange that preserves the colloidal state, either in the same or in a different solvent (through phase transfer reactions).[258]

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Complete removal of these organic ligands would be the ideal scenario for charge transport.

However, it reverses the surface passivation effect, causing surface dangling bonds to form in the forbidden band, acting as carrier traps in semiconductor NPs. Annealing the NPs can only partially remove organic ligands and usually creates unwanted carbonaceous products of pyrolysis or leads to sintering of NCs. Therefore, mild chemical removal of ligands and ligand- exchange reactions are the main surface chemistry adjustments: solution exchange and solid- state exchange. [258, 272-274]

Solid-state ligand exchange involves assembling a thin film of the NCs on a surface and immersion in a new ligand solution. However, the strategy is limited to thin films and is often incomplete and inhomogeneous throughout the film. Also, the replacement of very long ligands with short ones results in a drastic change in particle volume, leading to cracks or complete disintegration of the film. Therefore, exchange in the colloidal state, followed by film formation or other assemblies, is more desirable.

Figure 2-23 presents the surface states achievable from a colloidal solution. The incoming ligand replacement may exhibit a lower chemical affinity to the NP than the original ligand, or it may have the same functional group (e.g. thiol-for-thiol exchange). To overcome this, the

NPs must be exposed to a significant excess of replacement ligand to displace most original ligands.[275] These substitutions can take several seconds for strongly binding incoming ligands, such as thiols, or up to several days for weakly coordinating species, such as pyridine and may require moderate heating (≤100 °C). The solution-phase exchange often occurs across the phase boundary between mixtures of polar and non-polar solvents.[276]

There is a wide array of knowledge built up around ligand exchange for MCNCs. The most important of these have been short-chain organic ligands and inorganic ligands. The use of hydrophilic, deprotonated thiol ligands: 3-mercaptopropionic acid (MPA) and 3-mercapto-1-

2-64 propanol (MPOH), resulted in water-soluble CdS and CdSe NCs. The water solubility of these

NCs depended on the solution pH as well as the ligand shell. [277] These discoveries lead to far improved device efficiencies for a host of different metal chalcogenide materials. [5, 278-

280]

Deep trap states caused by interstitial and anti-site defects are rare in PNCs, so their defect tolerance is well known. However, shallow traps are common.[281] The poor stability of perovskite materials does lead to defects over time, meaning that maintaining strong surface coverage is vital to these materials' success. As discussed before, the ionic bond between the

NC surface and ligand leads to easy removal of the ligand and is the source of these defects.

[263] This ionic nature also limits the use of polar ligands as well as polar solvents.

Even so, there have been successes. Ligand exchanges during synthesis include Bis-(2,2,4- trimethylpentyl)phosphinic acid (TMPPA) as a suitable replacement of OA to synthesise

CsPbI3 NCs, resulting in improved stability.[268] Meanwhile, a trioctylphosphine-PbI2 (TOP-

PbI2) precursor lead to photoluminescence quantum yield (PLQY) reaching close to

100%.[282] Post synthesis exchanges have shown that the surface of PNCs can be substituted for thiocyanates NH4SCN, also improving quantum yield[283]. ZnBr2, sodium/ammonium tetrafluoroborate salt, and caesium cations all improve the surface of PNCs through ligand exchange, subsequently improving solar cell and PLQY performance.[207, 284, 285]

Unfortunately, other examples of ligand exchange for PNCs need developing. Ligand substitution during the reaction step has seen more success due to the ionic nature of the materials.

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2.6 Potential applications of nanocrystals

Colloidal NCs hold great promise for next-generation solution-processed devices through cost- effective methods such as spin coating, inkjet printing, slot-die coating, or spray coating, leading to large-scale processing as in roll-2-roll production. Perovskite NCs are regarded as good light emitters and have produced LEDs with tremendous electroluminescence efficiency.

[286-288] Figure 2-24a shows a schematic of a typical LED architecture fabricated with binary

2+ cation perovskite NCs, FAxMA1-xPbBr3, where FA = formamidinium, CH3(NH2) , and MA =

+ methylammonium, CH3NH3 , embedded into a PMMA layer with a hole transport layer (HTL), an emission layer (EML) and an electron transport layer (ETL). The contacts are generally a metal thin-film (Mo, Au) or highly conductive semiconductor (ITO, FTO). Figure 2-24b+c shows the device characteristics. This device has an efficiency of 7.37%.

Figure 2-24 a) Schematic of LED device architecture using perovskite quantum dots embedded in PMMA; b) Current density and luminescence as a function of driving voltage; c) Current efficiency and external quantum efficiency as a function of current density. [289] Other applications include solar cells and lasers. [194, 199, 290-296] Photo-detectors and photocatalysis applications are also growing fields in recent years. [297-300] Other applications include drug delivery, sensor, bio-imaging.[301-305] Perovskite photovoltaics have applications as the power source of spacecraft and satellites on space missions because of high specific power and high cosmic radiation resistance. [306]

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Layered III-VI metal chalcogenides have a similar variety of applications. Their anisotropic structure and bonding configuration result in high polarizability, optical uniformity and tuneable centrosymmetry. Their non-linear optical properties have potential in integrated optics, imaging techniques, and optical information communication. [307] Superior second harmonic generation performance is found in InSe and ε-InSe based alloys like InSe0.8S0.2 and

InSe0.8Te0.2, providing opportunities to develop 2D crystals over the abroad spectral range with high nonlinear optical efficiencies. [308] Single-photon emission of III-VI semiconductors is a focus for possible optical computing strategies in the field of quantum information science and quantum computing. Likewise, in condensed matter physics, group III-VI materials have been predicted to show novel physical properties such as superconductivity and the Bose-

Einstein condensation. Photocatalysis applications require materials with a bandgap above that of the free energy of the water-splitting 1.23 eV and a conduction band minimum that is higher

+ than the reduction potential of H /H2. The valence band maximum must be lower than the oxidation potential of O2/H2O. [309] Several theoretical studies show that the 2D layered III-

VI materials have the appropriate bandgaps and suitable band alignments that fulfil the above criteria.[310, 311]

Performance is enhanced by forming heterojunctions with a type II band alignment. The complete control from NC syntheses to device assembly paves the way for fabricating next- generation devices in the micro-and nano regime.

2.7 Bibliography

[1] G. Caruso, L. Merlo, and M. Caffo, Innovative Brain Tumor Therapy (no. 67). Woodhead Publishing, 2014. [2] V. Pokropivny, R. Lohmus, I. Hussainova, A. Pokropivny, and S. Vlassov, Introduction to nanomaterials and nanotechnology. Tartu University Press Ukraine, 2007.

2-67

[3] P. T.C, S. K. Sharma, and M. Kennedy, "Nanoparticles in household level water treatment: An overview," Separation and Purification Technology, vol. 199, pp. 260-270, 2018/06/30/ 2018, doi: https://doi.org/10.1016/j.seppur.2018.01.061. [4] F. Ejeian et al., "Biosensors for wastewater monitoring: A review," Biosensors and Bioelectronics, vol. 118, pp. 66-79, 2018/10/30/ 2018, doi: https://doi.org/10.1016/j.bios.2018.07.019. [5] S. Emin, S. P. Singh, L. Han, N. Satoh, and A. Islam, "Colloidal quantum dot solar cells," Solar Energy, vol. 85, no. 6, pp. 1264-1282, 2011/06/01/ 2011, doi: https://doi.org/10.1016/j.solener.2011.02.005. [6] Y. R. Kumar, K. Deshmukh, K. K. Sadasivuni, and S. K. Pasha, "Graphene quantum dot based materials for sensing, bio-imaging and energy storage applications: a review," RSC Advances, vol. 10, no. 40, pp. 23861-23898, 2020. [7] S. Daniel et al., "A review of DNA functionalized/grafted carbon nanotubes and their characterization," Sensors and Actuators B: Chemical, vol. 122, no. 2, pp. 672-682, 2007. [8] M. Holzinger, A. Le Goff, and S. Cosnier, "Nanomaterials for biosensing applications: a review," Frontiers in chemistry, vol. 2, p. 63, 2014. [9] I. Matsui, "Nanoparticles for electronic device applications: a brief review," Journal of chemical engineering of Japan, vol. 38, no. 8, pp. 535-546, 2005. [10] Y. Yin and A. P. Alivisatos, "Colloidal nanocrystal synthesis and the organic–inorganic interface," Nature, vol. 437, no. 7059, pp. 664-670, 2005/09/01 2005, doi: 10.1038/nature04165. [11] C. B. Murray, D. J. Norris, and M. G. Bawendi, "Synthesis and characterization of nearly monodisperse CdE (E = sulfur, selenium, tellurium) semiconductor nanocrystallites," Journal of the American Chemical Society, vol. 115, no. 19, pp. 8706-8715, 1993/09/01 1993, doi: 10.1021/ja00072a025. [12] M.-R. Gao, Y.-F. Xu, J. Jiang, and S.-H. Yu, "Nanostructured metal chalcogenides: synthesis, modification, and applications in energy conversion and storage devices," Chemical Society Reviews, vol. 42, no. 7, pp. 2986-3017, 2013. [13] C. Coughlan, M. Ibanez, O. Dobrozhan, A. Singh, A. Cabot, and K. M. Ryan, "Compound copper chalcogenide nanocrystals," Chemical reviews, vol. 117, no. 9, pp. 5865-6109, 2017. [14] J. Shamsi, A. S. Urban, M. Imran, L. De Trizio, and L. Manna, "Metal Halide Perovskite Nanocrystals: Synthesis, Post-Synthesis Modifications, and Their Optical Properties," Chemical Reviews, vol. 119, no. 5, pp. 3296-3348, 2019/03/13 2019, doi: 10.1021/acs.chemrev.8b00644. [15] H.-S. Kim et al., "Lead Iodide Perovskite Sensitized All-Solid-State Submicron Thin Film Mesoscopic Solar Cell with Efficiency Exceeding 9%," Scientific Reports, vol. 2, no. 1, p. 591, 2012/08/21 2012, doi: 10.1038/srep00591. [16] L. Protesescu et al., "Nanocrystals of Cesium Lead Halide Perovskites (CsPbX3, X = Cl, Br, and I): Novel Optoelectronic Materials Showing Bright Emission with Wide Color Gamut," Nano Letters, vol. 15, no. 6, pp. 3692-3696, 2015/06/10 2015, doi: 10.1021/nl5048779. [17] A. Dutta, R. K. Behera, P. Pal, S. Baitalik, and N. Pradhan, "Near-Unity Photoluminescence Quantum Efficiency for All CsPbX3 (X=Cl, Br, and I) Perovskite Nanocrystals: A Generic Synthesis Approach," Angewandte Chemie International Edition, vol. 58, no. 17, pp. 5552- 5556, 2019, doi: 10.1002/anie.201900374.

2-68

[18] G. Almeida et al., "Role of Acid-Base Equilibria in the Size, Shape, and Phase Control of Cesium Lead Bromide Nanocrystals," (in eng), ACS nano, vol. 12, no. 2, pp. 1704-1711, 2018, doi: 10.1021/acsnano.7b08357. [19] S. Sun, D. Yuan, Y. Xu, A. Wang, and Z. Deng, "Ligand-Mediated Synthesis of Shape-Controlled Cesium Lead Halide Perovskite Nanocrystals via Reprecipitation Process at Room Temperature," ACS Nano, vol. 10, no. 3, pp. 3648-3657, 2016/03/22 2016, doi: 10.1021/acsnano.5b08193. [20] X. Zhang et al., "Water‐assisted size and shape control of CsPbBr3 perovskite nanocrystals," Angewandte Chemie International Edition, vol. 57, no. 13, pp. 3337-3342, 2018. [21] J. Shamsi et al., "Colloidal Synthesis of Quantum Confined Single Crystal CsPbBr3 Nanosheets with Lateral Size Control up to the Micrometer Range," Journal of the American Chemical Society, vol. 138, no. 23, pp. 7240-7243, 2016/06/15 2016, doi: 10.1021/jacs.6b03166. [22] W. Zhai et al., "Solvothermal synthesis of ultrathin cesium lead halide perovskite nanoplatelets with tunable lateral sizes and their reversible transformation into Cs4PbBr6 nanocrystals," Chemistry of Materials, vol. 30, no. 11, pp. 3714-3721, 2018. [23] L. Dou et al., "Atomically thin two-dimensional organic-inorganic hybrid perovskites," Science, vol. 349, no. 6255, pp. 1518-1521, 2015. [24] D. Zhang, Y. Yu, Y. Bekenstein, A. B. Wong, A. P. Alivisatos, and P. Yang, "Ultrathin Colloidal Cesium Lead Halide Perovskite Nanowires," Journal of the American Chemical Society, vol. 138, no. 40, pp. 13155-13158, 2016/10/12 2016, doi: 10.1021/jacs.6b08373. [25] J. Zhao, S. Cao, Z. Li, and N. Ma, "Amino Acid-Mediated Synthesis of CsPbBr3 Perovskite Nanoplatelets with Tunable Thickness and Optical Properties," Chemistry of Materials, vol. 30, no. 19, pp. 6737-6743, 2018/10/09 2018, doi: 10.1021/acs.chemmater.8b02396. [26] M. C. Weidman, A. J. Goodman, and W. A. Tisdale, "Colloidal Halide Perovskite Nanoplatelets: An Exciting New Class of Semiconductor Nanomaterials," Chemistry of Materials, vol. 29, no. 12, pp. 5019-5030, 2017/06/27 2017, doi: 10.1021/acs.chemmater.7b01384. [27] Q. A. Akkerman, G. Rainò, M. V. Kovalenko, and L. Manna, "Genesis, challenges and opportunities for colloidal lead halide perovskite nanocrystals," Nature Materials, vol. 17, no. 5, pp. 394-405, 2018/05/01 2018, doi: 10.1038/s41563-018-0018-4. [28] M. G. Kanatzidis, "Discovery-synthesis, design, and prediction of chalcogenide phases," Inorganic chemistry, vol. 56, no. 6, pp. 3158-3173, 2017. [29] J. De Roo et al., "Highly Dynamic Ligand Binding and Light Absorption Coefficient of Cesium Lead Bromide Perovskite Nanocrystals," ACS Nano, vol. 10, no. 2, pp. 2071-2081, 2016/02/23 2016, doi: 10.1021/acsnano.5b06295. [30] C. E. Housecroft and E. C. Constable, Chemistry: an introduction to organic, inorganic and physical chemistry. Pearson education, 2010. [31] K. Seeger, Semiconductor physics. Springer Science & Business Media, 2013. [32] S. Jagtap, P. Chopade, S. Tadepalli, A. Bhalerao, and S. Gosavi, "A review on the progress of ZnSe as inorganic scintillator," Opto-Electronics Review, vol. 27, no. 1, pp. 90-103, 2019/03/01/ 2019, doi: https://doi.org/10.1016/j.opelre.2019.01.001. [33] A. C. Berends and C. de Mello Donega, "Ultrathin One- and Two-Dimensional Colloidal Semiconductor Nanocrystals: Pushing Quantum Confinement to the Limit," (in eng), The journal of physical chemistry letters, vol. 8, no. 17, pp. 4077-4090, 2017/09// 2017, doi: 10.1021/acs.jpclett.7b01640.

2-69

[34] N. Ashcroft and N. Mermin, "Solid State Physics Harcourt College Publishers," New Yoek, 1976. [35] P. Walter et al., "Early Use of PbS Nanotechnology for an Ancient Hair Dyeing Formula," Nano Letters, vol. 6, no. 10, pp. 2215-2219, 2006/10/01 2006, doi: 10.1021/nl061493u. [36] K. Kalyanasundaram, E. Borgarello, D. Duonghong, and M. Grätzel, "Cleavage of water by visible‐light irradiation of colloidal CdS solutions; inhibition of Photocorrosion by RuO2," Angewandte Chemie International Edition in English, vol. 20, no. 11, pp. 987-988, 1981. [37] Q. Cao et al., "Defects in Cu(In,Ga)Se2 Chalcopyrite Semiconductors: A Comparative Study of Material Properties, Defect States, and Photovoltaic Performance," Advanced Energy Materials, https://doi.org/10.1002/aenm.201100344 vol. 1, no. 5, pp. 845-853, 2011/10/01 2011, doi: https://doi.org/10.1002/aenm.201100344. [38] T. Unold and H.-W. Schock, "Nonconventional (non-silicon-based) photovoltaic materials," Annual Review of Materials Research, vol. 41, pp. 297-321, 2011. [39] P. Stathi, M. Solakidou, and Y. Deligiannakis, "Lattice Defects Engineering in W-, Zr-doped BiVO4 by Flame Spray Pyrolysis: Enhancing Photocatalytic O2 Evolution," Nanomaterials, vol. 11, no. 2, p. 501, 2021. [40] W.-H. Chou, A. Gamboa, and J. O. Morales, "Inkjet printing of small molecules, biologics, and nanoparticles," International Journal of Pharmaceutics, p. 120462, 2021. [41] R. Vignesh, V. B. Mathy, G. Geetha, R. Sivakumar, and C. Sanjeeviraja, "Temperature induced thermochromism of m-BiVO4 thin films prepared by sol-gel spin coating technique," Materials Letters, vol. 285, p. 129200, 2021. [42] C. de Mello Donegá, P. Liljeroth, and D. Vanmaekelbergh, "Physicochemical evaluation of the hot‐injection method, a synthesis route for monodisperse nanocrystals," Small, vol. 1, no. 12, pp. 1152-1162, 2005. [43] U. Ghorpade et al., "Towards environmentally benign approaches for the synthesis of CZTSSe nanocrystals by a hot injection method: a status review," Chemical communications, vol. 50, no. 77, pp. 11258-11273, 2014. [44] S. B. Brichkin and V. F. Razumov, "Colloidal quantum dots: synthesis, properties and applications," Russian Chemical Reviews, vol. 85, no. 12, p. 1297, 2016. [45] A. M. Smith and S. Nie, "Chemical analysis and cellular imaging with quantum dots," Analyst, 10.1039/B404498N vol. 129, no. 8, pp. 672-677, 2004, doi: 10.1039/B404498N. [46] X. Peng, M. C. Schlamp, A. V. Kadavanich, and A. P. Alivisatos, "Epitaxial growth of highly luminescent CdSe/CdS core/shell nanocrystals with photostability and electronic accessibility," Journal of the American Chemical Society, vol. 119, no. 30, pp. 7019-7029, 1997. [47] R. Xie, M. Rutherford, and X. Peng, "Formation of high-quality I− III− VI semiconductor nanocrystals by tuning relative reactivity of cationic precursors," Journal of the American Chemical Society, vol. 131, no. 15, pp. 5691-5697, 2009. [48] W. W. Yu and X. Peng, "Formation of high‐quality CdS and other II–VI semiconductor nanocrystals in noncoordinating solvents: tunable reactivity of monomers," Angewandte Chemie International Edition, vol. 41, no. 13, pp. 2368-2371, 2002. [49] J. J. Li et al., "Large-scale synthesis of nearly monodisperse CdSe/CdS core/shell nanocrystals using air-stable reagents via successive ion layer adsorption and reaction," Journal of the American Chemical Society, vol. 125, no. 41, pp. 12567-12575, 2003.

2-70

[50] I. Lignos et al., "Unveiling the shape evolution and halide-ion-segregation in blue-emitting formamidinium lead halide perovskite nanocrystals using an automated microfluidic platform," Nano letters, vol. 18, no. 2, pp. 1246-1252, 2018. [51] Y. Dong, T. Qiao, D. Kim, D. Parobek, D. Rossi, and D. H. Son, "Precise control of quantum confinement in cesium lead halide perovskite quantum dots via thermodynamic equilibrium," Nano letters, vol. 18, no. 6, pp. 3716-3722, 2018. [52] X. Chen, L. Peng, K. Huang, Z. Shi, R. Xie, and W. Yang, "Non-injection gram-scale synthesis of cesium lead halide perovskite quantum dots with controllable size and composition," Nano Research, journal article vol. 9, no. 7, pp. 1994-2006, July 01 2016, doi: 10.1007/s12274-016- 1090-1. [53] Y. Li, H. Huang, Y. Xiong, S. V. Kershaw, and A. L. Rogach, "Revealing the Formation Mechanism of CsPbBr3 Perovskite Nanocrystals Produced via a Slowed-Down Microwave-Assisted Synthesis," Angewandte Chemie International Edition, vol. 57, no. 20, pp. 5833-5837, 2018, doi: https://doi.org/10.1002/anie.201713332. [54] H. Huang et al., "Growth mechanism of strongly emitting CH3NH3PbBr3 perovskite nanocrystals with a tunable bandgap," Nature Communications, vol. 8, no. 1, p. 996, 2017/10/17 2017, doi: 10.1038/s41467-017-00929-2. [55] K. Chen et al., "The chemistry of colloidal semiconductor nanocrystals: From metal- chalcogenides to emerging perovskite," Coordination Chemistry Reviews, vol. 418, p. 213333, 2020/09/01/ 2020, doi: https://doi.org/10.1016/j.ccr.2020.213333. [56] J. Shamsi et al., "Colloidal CsX (X = Cl, Br, I) Nanocrystals and Their Transformation to CsPbX3 Nanocrystals by Cation Exchange," Chemistry of Materials, vol. 30, no. 1, pp. 79-83, 2018/01/09 2018, doi: 10.1021/acs.chemmater.7b04827. [57] H. Kasai et al., "A novel preparation method of organic microcrystals," Japanese journal of applied physics, vol. 31, no. 8A, p. L1132, 1992. [58] F. Zhang et al., "Brightly luminescent and color-tunable colloidal CH3NH3PbX3 (X= Br, I, Cl) quantum dots: potential alternatives for display technology," ACS nano, vol. 9, no. 4, pp. 4533- 4542, 2015. [59] J. Cho, Y.-H. Choi, T. E. O’Loughlin, L. De Jesus, and S. Banerjee, "Ligand-Mediated Modulation of Layer Thicknesses of Perovskite Methylammonium Lead Bromide Nanoplatelets," Chemistry of Materials, vol. 28, no. 19, pp. 6909-6916, 2016/10/11 2016, doi: 10.1021/acs.chemmater.6b02241. [60] L. Al Juhaiman, L. Scoles, D. Kingston, B. Patarachao, D. Wang, and F. Bensebaa, "Green synthesis of tunable Cu (In1− xGax) Se2 nanoparticles using non-organic solvents," Green Chemistry, vol. 12, no. 7, pp. 1248-1252, 2010. [61] X. Hu, Q. Zhang, X. Huang, D. Li, Y. Luo, and Q. Meng, "Aqueous colloidal CuInS2 for quantum dot sensitized solar cells," Journal of Materials Chemistry, vol. 21, no. 40, pp. 15903-15905, 2011. [62] Y. Jo et al., "High performance of planar perovskite solar cells produced from PbI2 (DMSO) and PbI2 (NMP) complexes by intramolecular exchange," Advanced Materials Interfaces, vol. 3, no. 10, p. 1500768, 2016. [63] D. Shen et al., "Understanding the solvent-assisted crystallization mechanism inherent in efficient organic–inorganic halide perovskite solar cells," Journal of Materials Chemistry A, vol. 2, no. 48, pp. 20454-20461, 2014.

2-71

[64] Y. Deng, C. H. Van Brackle, X. Dai, J. Zhao, B. Chen, and J. Huang, "Tailoring solvent coordination for high-speed, room-temperature blading of perovskite photovoltaic films," Science Advances, vol. 5, no. 12, p. eaax7537, 2019, doi: 10.1126/sciadv.aax7537. [65] F. Li et al., "High Br–Content CsPb (Cl y Br1–y) 3 Perovskite Nanocrystals with Strong Mn2+ Emission through Diverse Cation/Anion Exchange Engineering," ACS applied materials & interfaces, vol. 10, no. 14, pp. 11739-11746, 2018. [66] F. Li, Z. Xia, Y. Gong, L. Gu, and Q. Liu, "Optical properties of Mn 2+ doped cesium lead halide perovskite nanocrystals via a cation–anion co-substitution exchange reaction," Journal of Materials Chemistry C, vol. 5, no. 36, pp. 9281-9287, 2017. [67] W. van der Stam et al., "Highly Emissive Divalent-Ion-Doped Colloidal CsPb1–xMxBr3 Perovskite Nanocrystals through Cation Exchange," Journal of the American Chemical Society, vol. 139, no. 11, pp. 4087-4097, 2017/03/22 2017, doi: 10.1021/jacs.6b13079. [68] Y. Tong et al., "Highly Luminescent Cesium Lead Halide Perovskite Nanocrystals with Tunable Composition and Thickness by Ultrasonication," Angewandte Chemie International Edition, vol. 55, no. 44, pp. 13887-13892, 2016, doi: doi:10.1002/anie.201605909. [69] L. Rao et al., "Polar-Solvent-Free Synthesis of Highly Photoluminescent and Stable CsPbBr3 Nanocrystals with Controlled Shape and Size by Ultrasonication," Chemistry of Materials, vol. 31, no. 2, pp. 365-375, 2019/01/22 2019, doi: 10.1021/acs.chemmater.8b03298. [70] A. Bohun, J. Dolejší, and Č. Barta, "The absorption and luminescence of (PbCl 6) 4− and (PbBr 6) 4− complexes," Czechoslovak Journal of Physics B, vol. 20, no. 7, pp. 803-807, 1970. [71] S. Kondo, K. Amaya, and T. Saito, "Localized optical absorption in Cs4PbBr6," Journal of Physics: Condensed Matter, vol. 14, no. 8, p. 2093, 2002. [72] Q. A. Akkerman et al., "Nearly Monodisperse Insulator Cs4PbX6 (X = Cl, Br, I) Nanocrystals, Their Mixed Halide Compositions, and Their Transformation into CsPbX3 Nanocrystals," Nano Letters, vol. 17, no. 3, pp. 1924-1930, 2017/03/08 2017, doi: 10.1021/acs.nanolett.6b05262. [73] L. Wu et al., "From Nonluminescent Cs4PbX6 (X = Cl, Br, I) Nanocrystals to Highly Luminescent CsPbX3 Nanocrystals: Water-Triggered Transformation through a CsX-Stripping Mechanism," Nano Letters, vol. 17, no. 9, pp. 5799-5804, 2017/09/13 2017, doi: 10.1021/acs.nanolett.7b02896. [74] W. Shen, L. Ruan, Z. Shen, and Z. Deng, "Reversible light-mediated compositional and structural transitions between CsPbBr 3 and CsPb 2 Br 5 nanosheets," Chemical Communications, vol. 54, no. 22, pp. 2804-2807, 2018. [75] L. Ruan, W. Shen, A. Wang, A. Xiang, and Z. Deng, "Alkyl-Thiol Ligand-Induced Shape- and Crystalline Phase-Controlled Synthesis of Stable Perovskite-Related CsPb2Br5 Nanocrystals at Room Temperature," The Journal of Physical Chemistry Letters, vol. 8, no. 16, pp. 3853-3860, 2017/08/17 2017, doi: 10.1021/acs.jpclett.7b01657. [76] J. Li et al., "Inter-conversion between different compounds of ternary Cs-Pb-Br system," Materials, vol. 11, no. 5, p. 717, 2018. [77] D. Yang, M. Cao, Q. Zhong, P. Li, X. Zhang, and Q. Zhang, "All-inorganic cesium lead halide perovskite nanocrystals: synthesis, surface engineering and applications," Journal of Materials Chemistry C, 10.1039/C8TC04381G vol. 7, no. 4, pp. 757-789, 2019, doi: 10.1039/C8TC04381G. [78] K. Chen, J. Zhou, W. Chen, P. Zhou, F. He, and Y. Liu, "Size‐Dependent Synthesis of Cu12Sb4S13 Nanocrystals with Bandgap Tunability," Particle & Particle Systems Characterization, vol. 32, no. 11, pp. 999-1005, 2015.

2-72

[79] J. Pan et al., "Bidentate ligand-passivated CsPbI3 perovskite nanocrystals for stable near-unity photoluminescence quantum yield and efficient red light-emitting diodes," Journal of the American Chemical Society, vol. 140, no. 2, pp. 562-565, 2017. [80] J. Pei, J. Yang, T. Yildirim, H. Zhang, and Y. Lu, "Many-Body Complexes in 2D Semiconductors," Advanced Materials, vol. 31, no. 2, p. 1706945, 2019, doi: 10.1002/adma.201706945. [81] L. Wu et al., "2D Tellurium Based High-Performance All-Optical Nonlinear Photonic Devices," Advanced Functional Materials, vol. 29, no. 4, p. 1806346, 2019, doi: 10.1002/adfm.201806346. [82] J. Theerthagiri et al., "Recent advances in metal chalcogenides (MX; X= S, Se) nanostructures for electrochemical supercapacitor applications: a brief review," Nanomaterials, vol. 8, no. 4, p. 256, 2018. [83] P. Baláž, M. Baláž, M. Achimovičová, Z. Bujňáková, and E. Dutková, "Chalcogenide mechanochemistry in materials science: insight into synthesis and applications (a review)," Journal of Materials Science, vol. 52, no. 20, pp. 11851-11890, 2017. [84] S. Ananthakumar and S. M. Babu, "Progress on synthesis and applications of hybrid perovskite semiconductor nanomaterials—A review," Synthetic Metals, vol. 246, pp. 64-95, 2018. [85] L. Xu, S. Yuan, H. Zeng, and J. Song, "A comprehensive review of doping in perovskite nanocrystals/quantum dots: evolution of structure, electronics, optics, and light-emitting diodes," Materials Today Nano, vol. 6, p. 100036, 2019. [86] H. Ren et al., "Synthesis and Characterization of CuZnSe2 Nanocrystals in Wurtzite, Zinc Blende, and Core–Shell Polytypes," Chemistry of Materials, vol. 31, no. 24, pp. 10085-10093, 2019. [87] D. Pan et al., "Synthesis of Cu−In−S Ternary Nanocrystals with Tunable Structure and Composition," Journal of the American Chemical Society, vol. 130, no. 17, pp. 5620-5621, 2008/04/01 2008, doi: 10.1021/ja711027j. [88] D. Aldakov, A. Lefrançois, and P. Reiss, "Ternary and quaternary metal chalcogenide nanocrystals: synthesis, properties and applications," Journal of Materials Chemistry C, 10.1039/C3TC30273C vol. 1, no. 24, pp. 3756-3776, 2013, doi: 10.1039/C3TC30273C. [89] J.-J. Wang, Y.-Q. Wang, F.-F. Cao, Y.-G. Guo, and L.-J. Wan, "Synthesis of Monodispersed Wurtzite Structure CuInSe2 Nanocrystals and Their Application in High-Performance Organic−Inorganic Hybrid Photodetectors," Journal of the American Chemical Society, vol. 132, no. 35, pp. 12218-12221, 2010/09/08 2010, doi: 10.1021/ja1057955. [90] J. Zhou, M. Zhu, R. Meng, H. Qin, and X. Peng, "Ideal CdSe/CdS Core/Shell Nanocrystals Enabled by Entropic Ligands and Their Core Size-, Shell Thickness-, and Ligand-Dependent Photoluminescence Properties," Journal of the American Chemical Society, vol. 139, no. 46, pp. 16556-16567, 2017/11/22 2017, doi: 10.1021/jacs.7b07434. [91] Y. Chen, D. Chen, Z. Li, and X. Peng, "Symmetry-Breaking for Formation of Rectangular CdSe Two-Dimensional Nanocrystals in Zinc-Blende Structure," Journal of the American Chemical Society, vol. 139, no. 29, pp. 10009-10019, 2017/07/26 2017, doi: 10.1021/jacs.7b04855. [92] B. Koo, R. N. Patel, and B. A. Korgel, "Wurtzite−Chalcopyrite Polytypism in CuInS2 Nanodisks," Chemistry of Materials, vol. 21, no. 9, pp. 1962-1966, 2009/05/12 2009, doi: 10.1021/cm900363w. [93] S. Lei et al., "Spinel Indium Sulfide Precursor for the Phase-Selective Synthesis of Cu–In–S Nanocrystals with Zinc-Blende, Wurtzite, and Spinel Structures," Chemistry of Materials, vol. 25, no. 15, pp. 2991-2997, 2013/08/13 2013, doi: 10.1021/cm400848f.

2-73

[94] R. Malik et al., "Interesting structural transformation of CdS from zinc blende into wurtzite during minuscule loading of magnetic nanoparticles: emergence of heterojunctions with enhanced photocatalytic performance," New Journal of Chemistry, 10.1039/C7NJ02707A vol. 41, no. 23, pp. 14088-14102, 2017, doi: 10.1039/C7NJ02707A. [95] M. Winter. "WebElements." https://www.webelements.com (accessed. [96] W. W. Yu, L. Qu, W. Guo, and X. Peng, "Experimental Determination of the Extinction Coefficient of CdTe, CdSe, and CdS Nanocrystals," Chemistry of Materials, vol. 15, no. 14, pp. 2854-2860, 2003/07/01 2003, doi: 10.1021/cm034081k. [97] N. Razgoniaeva et al., "One-Dimensional Carrier Confinement in “Giant” CdS/CdSe Excitonic Nanoshells," Journal of the American Chemical Society, vol. 139, no. 23, pp. 7815-7822, 2017/06/14 2017, doi: 10.1021/jacs.7b02054. [98] M. D. Regulacio and M.-Y. Han, "Composition-Tunable Alloyed Semiconductor Nanocrystals," Accounts of Chemical Research, vol. 43, no. 5, pp. 621-630, 2010/05/18 2010, doi: 10.1021/ar900242r. [99] A. de Kergommeaux et al., "Synthesis, Internal Structure, and Formation Mechanism of Monodisperse Tin Sulfide Nanoplatelets," Journal of the American Chemical Society, vol. 137, no. 31, pp. 9943-9952, 2015/08/12 2015, doi: 10.1021/jacs.5b05576. [100] S. T. Connor, C.-M. Hsu, B. D. Weil, S. Aloni, and Y. Cui, "Phase transformation of biphasic Cu2S− CuInS2 to monophasic CuInS2 nanorods," Journal of the American Chemical Society, vol. 131, no. 13, pp. 4962-4966, 2009. [101] A. Singh, H. Geaney, F. Laffir, and K. M. Ryan, "Colloidal Synthesis of Wurtzite Cu2ZnSnS4 Nanorods and Their Perpendicular Assembly," Journal of the American Chemical Society, vol. 134, no. 6, pp. 2910-2913, 2012/02/15 2012, doi: 10.1021/ja2112146. [102] A. Singh, S. Singh, S. Levcenko, T. Unold, F. Laffir, and K. M. Ryan, "Compositionally Tunable Photoluminescence Emission in Cu2ZnSn(S1−xSex)4 Nanocrystals," Angewandte Chemie International Edition, vol. 52, no. 35, pp. 9120-9124, 2013, doi: 10.1002/anie.201302867. [103] Z. Wang, Z. Shi, T. Li, Y. Chen, and W. Huang, "Stability of perovskite solar cells: a prospective on the substitution of the A cation and X anion," Angewandte Chemie International Edition, vol. 56, no. 5, pp. 1190-1212, 2017. [104] X.-G. Zhao et al., "Design of Lead-Free Inorganic Halide Perovskites for Solar Cells via Cation- Transmutation," Journal of the American Chemical Society, vol. 139, no. 7, pp. 2630-2638, 2017/02/22 2017, doi: 10.1021/jacs.6b09645. [105] X.-G. Zhao et al., "Cu–In Halide Perovskite Solar Absorbers," Journal of the American Chemical Society, vol. 139, no. 19, pp. 6718-6725, 2017/05/17 2017, doi: 10.1021/jacs.7b02120. [106] D. W. Ferdani et al., "Partial cation substitution reduces iodide ion transport in lead iodide perovskite solar cells," Energy & Environmental Science, vol. 12, no. 7, pp. 2264-2272, 2019. [107] K. Y. Tsui, N. Onishi, and R. F. Berger, "Tolerance Factors Revisited: Geometrically Designing the Ideal Environment for Perovskite Dopants," The Journal of Physical Chemistry C, vol. 120, no. 40, pp. 23293-23298, 2016/10/13 2016, doi: 10.1021/acs.jpcc.6b09277. [108] B. Saparov and D. B. Mitzi, "Organic–Inorganic Perovskites: Structural Versatility for Functional Materials Design," Chemical Reviews, vol. 116, no. 7, pp. 4558-4596, 2016/04/13 2016, doi: 10.1021/acs.chemrev.5b00715. [109] M. V. Kovalenko, L. Protesescu, and M. I. Bodnarchuk, "Properties and potential optoelectronic applications of lead halide perovskite nanocrystals," Science, vol. 358, no. 6364, pp. 745-750, 2017, doi: 10.1126/science.aam7093. 2-74

[110] D. P. McMeekin et al., "A mixed-cation lead mixed-halide perovskite absorber for tandem solar cells," Science, vol. 351, no. 6269, pp. 151-155, 2016. [111] W. Travis, E. Glover, H. Bronstein, D. Scanlon, and R. Palgrave, "On the application of the tolerance factor to inorganic and hybrid halide perovskites: a revised system," Chemical Science, vol. 7, no. 7, pp. 4548-4556, 2016. [112] Y. Cao, Z. Zhang, L. Li, J.-R. Zhang, and J.-J. Zhu, "An Improved Strategy for High-Quality Cesium Bismuth Bromine Perovskite Quantum Dots with Remarkable Electrochemiluminescence Activities," Analytical Chemistry, vol. 91, no. 13, pp. 8607-8614, 2019/07/02 2019, doi: 10.1021/acs.analchem.9b01918. [113] J. Zhang et al., "High Quantum Yield Blue Emission from Lead-Free Inorganic Antimony Halide Perovskite Colloidal Quantum Dots," ACS Nano, vol. 11, no. 9, pp. 9294-9302, 2017/09/26 2017, doi: 10.1021/acsnano.7b04683. [114] J. Luo et al., "Efficient and stable emission of warm-white light from lead-free halide double perovskites," Nature, vol. 563, no. 7732, pp. 541-545, 2018. [115] D. Wang, M. Wright, N. K. Elumalai, and A. Uddin, "Stability of perovskite solar cells," Solar Energy Materials and Solar Cells, vol. 147, pp. 255-275, 2016. [116] J. W. Lee, D. H. Kim, H. S. Kim, S. W. Seo, S. M. Cho, and N. G. Park, "Formamidinium and cesium hybridization for photo‐and moisture‐stable perovskite solar cell," Advanced Energy Materials, vol. 5, no. 20, p. 1501310, 2015. [117] N.-G. Park, M. Grätzel, T. Miyasaka, K. Zhu, and K. Emery, "Towards stable and commercially available perovskite solar cells," Nature Energy, vol. 1, no. 11, pp. 1-8, 2016. [118] J.-K. Sun et al., "Polar solvent induced lattice distortion of cubic CsPbI3 nanocubes and hierarchical self-assembly into orthorhombic single-crystalline nanowires," Journal of the American Chemical Society, vol. 140, no. 37, pp. 11705-11715, 2018. [119] Q. A. Akkerman and L. Manna, "What Defines a Halide Perovskite?," ACS Energy Letters, vol. 5, no. 2, pp. 604-610, 2020/02/14 2020, doi: 10.1021/acsenergylett.0c00039. [120] K. Chen et al., "Short‐Chain Ligand‐Passivated Stable α‐CsPbI3 Quantum Dot for All‐Inorganic Perovskite Solar Cells," Advanced Functional Materials, vol. 29, no. 24, p. 1900991, 2019. [121] A. Dutta, S. K. Dutta, S. Das Adhikari, and N. Pradhan, "Phase‐Stable CsPbI3 Nanocrystals: The Reaction Temperature Matters," Angewandte chemie international edition, vol. 57, no. 29, pp. 9083-9087, 2018. [122] A. K. Jena, A. Kulkarni, Y. Sanehira, M. Ikegami, and T. Miyasaka, "Stabilization of α-CsPbI3 in Ambient Room Temperature Conditions by Incorporating Eu into CsPbI3," Chemistry of Materials, vol. 30, no. 19, pp. 6668-6674, 2018/10/09 2018, doi: 10.1021/acs.chemmater.8b01808. [123] Y. Wang, T. Zhang, M. Kan, and Y. Zhao, "Bifunctional Stabilization of All-Inorganic α-CsPbI3 Perovskite for 17% Efficiency Photovoltaics," Journal of the American Chemical Society, vol. 140, no. 39, pp. 12345-12348, 2018/10/03 2018, doi: 10.1021/jacs.8b07927. [124] W. Ke, I. Spanopoulos, C. C. Stoumpos, and M. G. Kanatzidis, "Myths and reality of HPbI 3 in halide perovskite solar cells," Nature communications, vol. 9, no. 1, pp. 1-9, 2018. [125] D. Yang, X. Li, and H. Zeng, "Surface Chemistry of All Inorganic Halide Perovskite Nanocrystals: Passivation Mechanism and Stability," Advanced Materials Interfaces, vol. 5, no. 8, p. 1701662, 2018, doi: 10.1002/admi.201701662.

2-75

[126] A. Pisanu, P. Quadrelli, and L. Malavasi, "Facile anion-exchange reaction in mixed-cation lead bromide perovskite nanocrystals," RSC Advances, 10.1039/C9RA01089K vol. 9, no. 23, pp. 13263-13268, 2019, doi: 10.1039/C9RA01089K. [127] G. Nedelcu, L. Protesescu, S. Yakunin, M. I. Bodnarchuk, M. J. Grotevent, and M. V. Kovalenko, "Fast anion-exchange in highly luminescent nanocrystals of cesium lead halide perovskites (CsPbX3, X= Cl, Br, I)," Nano letters, vol. 15, no. 8, pp. 5635-5640, 2015. [128] L. Dou et al., "Spatially resolved multicolor CsPbX₃ nanowire heterojunctions via anion exchange," Proceedings of the National Academy of Sciences, vol. 114, no. 28, pp. 7216- 7221, 2017, doi: 10.1073/pnas.1703860114. [129] Q. A. Akkerman et al., "Tuning the Optical Properties of Cesium Lead Halide Perovskite Nanocrystals by Anion Exchange Reactions," Journal of the American Chemical Society, vol. 137, no. 32, pp. 10276-10281, 2015/08/19 2015, doi: 10.1021/jacs.5b05602. [130] H. Zhang et al., "Phase segregation due to ion migration in all-inorganic mixed-halide perovskite nanocrystals," Nature Communications, vol. 10, no. 1, p. 1088, 2019/03/06 2019, doi: 10.1038/s41467-019-09047-7. [131] Z. Ma et al., "Pressure-induced emission of cesium lead halide perovskite nanocrystals," Nature communications, vol. 9, no. 1, pp. 1-8, 2018. [132] Q. A. Akkerman, A. L. Abdelhady, and L. Manna, "Zero-dimensional cesium lead halides: history, properties, and challenges," The Journal of Physical Chemistry Letters, vol. 9, no. 9, pp. 2326-2337, 2018. [133] Y. Zhang et al., "Zero-Dimensional Cs4PbBr6 Perovskite Nanocrystals," The Journal of Physical Chemistry Letters, vol. 8, no. 5, pp. 961-965, 2017/03/02 2017, doi: 10.1021/acs.jpclett.7b00105. [134] O. F. Mohammed, "Outstanding Challenges of Zero-Dimensional Perovskite Materials," The Journal of Physical Chemistry Letters, vol. 10, no. 19, pp. 5886-5888, 2019/10/03 2019, doi: 10.1021/acs.jpclett.9b00175. [135] J. Yin et al., "Point defects and green emission in zero-dimensional perovskites," The journal of physical chemistry letters, vol. 9, no. 18, pp. 5490-5495, 2018. [136] Y. Zhang et al., "Ligand-free nanocrystals of highly emissive Cs4PbBr6 perovskite," The Journal of Physical Chemistry C, vol. 122, no. 11, pp. 6493-6498, 2018. [137] P. Acharyya, P. Pal, P. K. Samanta, A. Sarkar, S. K. Pati, and K. Biswas, "Single pot synthesis of indirect band gap 2D CsPb2Br5 nanosheets from direct band gap 3D CsPbBr3 nanocrystals and the origin of their luminescence properties," Nanoscale, 10.1039/C8NR09349K vol. 11, no. 9, pp. 4001-4007, 2019, doi: 10.1039/C8NR09349K. [138] Z. Wu, S. Yang, and W. Wu, "Shape control of inorganic nanoparticles from solution," Nanoscale, 10.1039/C5NR07681A vol. 8, no. 3, pp. 1237-1259, 2016, doi: 10.1039/C5NR07681A. [139] T.-D. Nguyen, "From formation mechanisms to synthetic methods toward shape-controlled oxide nanoparticles," Nanoscale, vol. 5, no. 20, pp. 9455-9482, 2013. [140] M. Grzelczak, J. Pérez-Juste, P. Mulvaney, and L. M. Liz-Marzán, "Shape control in gold nanoparticle synthesis," Chemical Society Reviews, vol. 37, no. 9, pp. 1783-1791, 2008. [141] Y.-w. Jun, J.-s. Choi, and J. Cheon, "Shape Control of Semiconductor and Metal Oxide Nanocrystals through Nonhydrolytic Colloidal Routes," Angewandte Chemie International Edition, vol. 45, no. 21, pp. 3414-3439, 2006, doi: 10.1002/anie.200503821.

2-76

[142] K. Manthiram, B. J. Beberwyck, D. V. Talapin, and A. P. Alivisatos, "Seeded synthesis of CdSe/CdS rod and tetrapod nanocrystals," (in eng), J Vis Exp, no. 82, pp. e50731-e50731, 2013, doi: 10.3791/50731. [143] "Synthesis, Characterization, and Applications of Three-Dimensional (3D) Nanostructures," in Synthesis and Applications of Inorganic Nanostructures, pp. 363-520. [144] K. Khan et al., "Recent developments in emerging two-dimensional materials and their applications," Journal of Materials Chemistry C, 10.1039/C9TC04187G vol. 8, no. 2, pp. 387- 440, 2020, doi: 10.1039/C9TC04187G. [145] P. Vogt et al., "Silicene: compelling experimental evidence for graphenelike two-dimensional silicon," Physical review letters, vol. 108, no. 15, p. 155501, 2012. [146] M. Dávila, L. Xian, S. Cahangirov, A. Rubio, and G. Le Lay, "Germanene: a novel two- dimensional germanium allotrope akin to graphene and silicene," New Journal of Physics, vol. 16, no. 9, p. 095002, 2014. [147] J. Kang et al., "Solvent exfoliation of electronic-grade, two-dimensional black phosphorus," ACS nano, vol. 9, no. 4, pp. 3596-3604, 2015. [148] Q. H. Wang, K. Kalantar-Zadeh, A. Kis, J. N. Coleman, and M. S. Strano, "Electronics and optoelectronics of two-dimensional transition metal dichalcogenides," Nature nanotechnology, vol. 7, no. 11, p. 699, 2012. [149] L. Liu et al., "Heteroepitaxial growth of two-dimensional hexagonal boron nitride templated by graphene edges," Science, vol. 343, no. 6167, pp. 163-167, 2014. [150] D. Deng, K. Novoselov, Q. Fu, N. Zheng, Z. Tian, and X. Bao, "Catalysis with two-dimensional materials and their heterostructures," Nature nanotechnology, vol. 11, no. 3, p. 218, 2016. [151] G. Eda and S. A. Maier, "Two-dimensional crystals: managing light for optoelectronics," ACS nano, vol. 7, no. 7, pp. 5660-5665, 2013. [152] M. Bernardi, M. Palummo, and J. C. Grossman, "Extraordinary sunlight absorption and one nanometer thick photovoltaics using two-dimensional monolayer materials," Nano letters, vol. 13, no. 8, pp. 3664-3670, 2013. [153] L. Li et al., "Functionalized graphene for high-performance two-dimensional spintronics devices," ACS nano, vol. 5, no. 4, pp. 2601-2610, 2011. [154] X. Li et al., "Synthesis of thin-film black phosphorus on a flexible substrate," 2D Materials, vol. 2, no. 3, p. 031002, 2015. [155] V. Nicolosi, M. Chhowalla, M. G. Kanatzidis, M. S. Strano, and J. N. Coleman, "Liquid exfoliation of layered materials," Science, vol. 340, no. 6139, p. 1226419, 2013. [156] H. Li et al., "Mechanical exfoliation and characterization of single‐and few‐layer nanosheets of WSe2, TaS2, and TaSe2," Small, vol. 9, no. 11, pp. 1974-1981, 2013. [157] S. Jeong, D. Yoo, J.-t. Jang, M. Kim, and J. Cheon, "Well-defined colloidal 2-D layered transition- metal chalcogenide nanocrystals via generalized synthetic protocols," Journal of the American Chemical Society, vol. 134, no. 44, pp. 18233-18236, 2012. [158] M. Nasilowski, B. Mahler, E. Lhuillier, S. Ithurria, and B. Dubertret, "Two-Dimensional Colloidal Nanocrystals," Chemical Reviews, vol. 116, no. 18, pp. 10934-10982, 2016/09/28 2016, doi: 10.1021/acs.chemrev.6b00164. [159] D. Yoo, M. Kim, S. Jeong, J. Han, and J. Cheon, "Chemical Synthetic Strategy for Single-Layer Transition-Metal Chalcogenides," Journal of the American Chemical Society, vol. 136, no. 42, pp. 14670-14673, 2014/10/22 2014, doi: 10.1021/ja5079943. 2-77

[160] D. D. Vaughn, R. J. Patel, M. A. Hickner, and R. E. Schaak, "Single-Crystal Colloidal Nanosheets of GeS and GeSe," Journal of the American Chemical Society, vol. 132, no. 43, pp. 15170-15172, 2010/11/03 2010, doi: 10.1021/ja107520b. [161] D. W. Boukhvalov et al., "The advent of indium selenide: synthesis, electronic properties, ambient stability and applications," Nanomaterials, vol. 7, no. 11, p. 372, 2017. [162] S. Sucharitakul et al., "Intrinsic electron mobility exceeding 103 cm2/(V s) in multilayer InSe FETs," Nano letters, vol. 15, no. 6, pp. 3815-3819, 2015. [163] S. Lei et al., "An atomically layered InSe avalanche photodetector," Nano letters, vol. 15, no. 5, pp. 3048-3055, 2015. [164] Z. Chen, J. Biscaras, and A. Shukla, "A high performance graphene/few-layer InSe photodetector," Nanoscale, vol. 7, no. 14, pp. 5981-5986, 2015. [165] Q. Peng et al., "Computational mining of photocatalysts for water splitting hydrogen production: two-dimensional InSe-family monolayers," Catalysis Science & Technology, vol. 7, no. 13, pp. 2744-2752, 2017. [166] C. H. Ho and Y. J. Chu, "Bending photoluminescence and surface photovoltaic effect on multilayer InSe 2D microplate crystals," Advanced Optical Materials, vol. 3, no. 12, pp. 1750- 1758, 2015. [167] Y. Ma, Y. Dai, L. Yu, C. Niu, and B. Huang, "Engineering a topological phase transition in β-InSe via strain," New Journal of Physics, vol. 15, no. 7, p. 073008, 2013. [168] M. Yüksek, H. G. Yaglioglu, A. Elmali, E. M. Aydın, U. Kürüm, and A. Ateş, "Nonlinear and saturable absorption characteristics of Ho doped InSe crystals," Optics Communications, vol. 310, pp. 100-103, 2014. [169] I. Grimaldi et al., "Structural investigation of InSe layered semiconductors," Solid State Communications, p. 113855, 2020. [170] S. Lei et al., "Evolution of the electronic band structure and efficient photo-detection in atomic layers of InSe," ACS nano, vol. 8, no. 2, pp. 1263-1272, 2014. [171] B. Gürbulak, M. Şata, S. Dogan, S. Duman, A. Ashkhasi, and E. F. Keskenler, "Structural characterizations and optical properties of InSe and InSe: Ag semiconductors grown by Bridgman/Stockbarger technique," Physica E: Low-dimensional Systems and Nanostructures, vol. 64, pp. 106-111, 2014. [172] G. Mudd et al., "Quantum confined acceptors and donors in InSe nanosheets," Applied Physics Letters, vol. 105, no. 22, p. 221909, 2014. [173] L. Viti et al., "Black phosphorus terahertz photodetectors," Advanced Materials, vol. 27, no. 37, pp. 5567-5572, 2015. [174] P.-H. Ho et al., "High-mobility InSe transistors: the role of surface oxides," Acs Nano, vol. 11, no. 7, pp. 7362-7370, 2017. [175] C.-Y. Lu et al., "Laser and electrical current induced phase transformation of In 2 Se 3 semiconductor thin film on Si (111)," Applied Physics A, vol. 93, no. 1, pp. 93-98, 2008. [176] Y. Li et al., "Thermal phase transformation of In 2 Se 3 nanowires studied by in situ synchrotron radiation X-ray diffraction," Journal of Materials Chemistry, vol. 21, no. 19, pp. 6944-6947, 2011. [177] T. Zhai et al., "Morphology-tunable In2Se3 nanostructures with enhanced electrical and photoelectrical performances via sulfur doping," Journal of Materials Chemistry, vol. 20, no. 32, pp. 6630-6637, 2010. 2-78

[178] R. Diehl and R. Nitsche, "Vapour growth of three In2S3 modifications by iodine transport," Journal of Crystal Growth, vol. 28, no. 3, pp. 306-310, 1975. [179] F. Horani and E. Lifshitz, "Unraveling the Growth Mechanism Forming Stable γ-In2S3 and β- In2S3 Colloidal Nanoplatelets," Chemistry of Materials, vol. 31, no. 5, pp. 1784-1793, 2019/03/12 2019, doi: 10.1021/acs.chemmater.9b00013. [180] N. T. K. Thanh, N. Maclean, and S. Mahiddine, "Mechanisms of Nucleation and Growth of Nanoparticles in Solution," Chemical Reviews, vol. 114, no. 15, pp. 7610-7630, 2014/08/13 2014, doi: 10.1021/cr400544s. [181] V. I. Kalikmanov, "Classical nucleation theory," in Nucleation theory: Springer, 2013, pp. 17- 41. [182] J. Polte, "Fundamental growth principles of colloidal metal nanoparticles – a new perspective," CrystEngComm, 10.1039/C5CE01014D vol. 17, no. 36, pp. 6809-6830, 2015, doi: 10.1039/C5CE01014D. [183] M. Zanella, G. Bertoni, I. R. Franchini, R. Brescia, D. Baranov, and L. Manna, "Assembly of shape-controlled nanocrystals by depletion attraction," Chemical Communications, vol. 47, no. 1, pp. 203-205, 2011. [184] H. Shen et al., "Fabrication of “strong” columnar Cu 2− x Se superstructures assisted by inorganic ligands," Nanoscale, vol. 4, no. 8, pp. 2741-2747, 2012. [185] K. Miszta et al., "Hierarchical self-assembly of suspended branched colloidal nanocrystals into superlattice structures," Nature materials, vol. 10, no. 11, pp. 872-876, 2011. [186] W. H. Evers et al., "Low-dimensional semiconductor superlattices formed by geometric control over nanocrystal attachment," Nano letters, vol. 13, no. 6, pp. 2317-2323, 2013. [187] E. Piccinini, D. Pallarola, F. Battaglini, and O. Azzaroni, "Self-limited self-assembly of nanoparticles into supraparticles: towards supramolecular colloidal materials by design," Molecular Systems Design & Engineering, 10.1039/C6ME00016A vol. 1, no. 2, pp. 155-162, 2016, doi: 10.1039/C6ME00016A. [188] J. Owen, "The coordination chemistry of nanocrystal surfaces," Science, vol. 347, no. 6222, pp. 615-616, 2015, doi: 10.1126/science.1259924. [189] L. Jing et al., "Aqueous Based Semiconductor Nanocrystals," Chemical Reviews, vol. 116, no. 18, pp. 10623-10730, 2016/09/28 2016, doi: 10.1021/acs.chemrev.6b00041. [190] A. Heuer-Jungemann et al., "The Role of Ligands in the Chemical Synthesis and Applications of Inorganic Nanoparticles," Chemical Reviews, vol. 119, no. 8, pp. 4819-4880, 2019/04/24 2019, doi: 10.1021/acs.chemrev.8b00733. [191] M. P. Hendricks, M. P. Campos, G. T. Cleveland, I. Jen-La Plante, and J. S. Owen, "A tunable library of substituted thiourea precursors to metal sulfide nanocrystals," Science, vol. 348, no. 6240, p. 1226, 2015, doi: 10.1126/science.aaa2951. [192] M. P. Campos et al., "A Library of Selenourea Precursors to PbSe Nanocrystals with Size Distributions near the Homogeneous Limit," Journal of the American Chemical Society, vol. 139, no. 6, pp. 2296-2305, 2017/02/15 2017, doi: 10.1021/jacs.6b11021. [193] A. Swarnkar et al., "Quantum dot–induced phase stabilization of α-CsPbI3 perovskite for high- efficiency photovoltaics," Science, vol. 354, no. 6308, pp. 92-95, 2016. [194] J. Li et al., "50-Fold EQE Improvement up to 6.27% of Solution-Processed All-Inorganic Perovskite CsPbBr3 QLEDs via Surface Ligand Density Control," Advanced Materials,

2-79

https://doi.org/10.1002/adma.201603885 vol. 29, no. 5, p. 1603885, 2017/02/01 2017, doi: https://doi.org/10.1002/adma.201603885. [195] M. A. Boles, D. Ling, T. Hyeon, and D. V. Talapin, "The surface science of nanocrystals," Nature Materials, vol. 15, no. 2, pp. 141-153, 2016/02/01 2016, doi: 10.1038/nmat4526. [196] A. Lemmerer and D. G. Billing, "Lead halide inorganic–organic hybrids incorporating diammonium cations," CrystEngComm, vol. 14, no. 6, pp. 1954-1966, 2012. [197] T. Zhao, C.-C. Chueh, Q. Chen, A. Rajagopal, and A. K. Y. Jen, "Defect Passivation of Organic– Inorganic Hybrid Perovskites by Diammonium Iodide toward High-Performance Photovoltaic Devices," ACS Energy Letters, vol. 1, no. 4, pp. 757-763, 2016/10/14 2016, doi: 10.1021/acsenergylett.6b00327. [198] S. Yang, Y. Wang, P. Liu, Y.-B. Cheng, H. J. Zhao, and H. G. Yang, "Functionalization of perovskite thin films with moisture-tolerant molecules," Nature Energy, vol. 1, no. 2, p. 15016, 2016/01/18 2016, doi: 10.1038/nenergy.2015.16. [199] Z. Zhang et al., "Progress in Multifunctional Molecules for Perovskite Solar Cells," Solar RRL, vol. 4, no. 2, p. 1900248, 2020. [200] X. Zheng et al., "Defect passivation in hybrid perovskite solar cells using quaternary ammonium halide anions and cations," Nature Energy, vol. 2, no. 7, pp. 1-9, 2017. [201] H. Zhang et al., "Bright perovskite light-emitting diodes with improved film morphology and reduced trap density via surface passivation using quaternary ammonium salts," Organic Electronics, vol. 67, pp. 187-193, 2019/04/01/ 2019, doi: https://doi.org/10.1016/j.orgel.2019.01.030. [202] S. Wang, Y. Zhu, C. Wang, and R. Ma, "Interface modification by a multifunctional ammonium salt for high performance and stable planar perovskite solar cells," Journal of Materials Chemistry A, 10.1039/C9TA02631B vol. 7, no. 19, pp. 11867-11876, 2019, doi: 10.1039/C9TA02631B. [203] J. Pan et al., "Highly efficient perovskite‐quantum‐dot light‐emitting diodes by surface engineering," Advanced Materials, vol. 28, no. 39, pp. 8718-8725, 2016. [204] Y. Cai, L. Wang, T. Zhou, P. Zheng, Y. Li, and R.-J. Xie, "Improved stability of CsPbBr3 perovskite quantum dots achieved by suppressing interligand proton transfer and applying a polystyrene coating," Nanoscale, 10.1039/C8NR06607H vol. 10, no. 45, pp. 21441-21450, 2018, doi: 10.1039/C8NR06607H. [205] J. H. Park et al., "Surface Ligand Engineering for Efficient Perovskite Nanocrystal-Based Light- Emitting Diodes," ACS Applied Materials & Interfaces, vol. 11, no. 8, pp. 8428-8435, 2019/02/27 2019, doi: 10.1021/acsami.8b20808. [206] Y. Shynkarenko et al., "Direct Synthesis of Quaternary Alkylammonium-Capped Perovskite Nanocrystals for Efficient Blue and Green Light-Emitting Diodes," ACS Energy Letters, vol. 4, no. 11, pp. 2703-2711, 2019/11/08 2019, doi: 10.1021/acsenergylett.9b01915. [207] J. Song et al., "Room‐temperature triple‐ligand surface engineering synergistically boosts ink stability, recombination dynamics, and charge injection toward EQE‐11.6% perovskite QLEDs," Advanced Materials, vol. 30, no. 30, p. 1800764, 2018. [208] F. Wang, S. Bai, W. Tress, A. Hagfeldt, and F. Gao, "Defects engineering for high-performance perovskite solar cells," npj Flexible Electronics, vol. 2, no. 1, p. 22, 2018/08/14 2018, doi: 10.1038/s41528-018-0035-z.

2-80

[209] L. Wu et al., "Improving the Stability and Size Tunability of Cesium Lead Halide Perovskite Nanocrystals Using Trioctylphosphine Oxide as the Capping Ligand," Langmuir, vol. 33, no. 44, pp. 12689-12696, 2017/11/07 2017, doi: 10.1021/acs.langmuir.7b02963. [210] M. R. Alpert et al., "Colloidal nanocrystals as a platform for rapid screening of charge trap passivating molecules for metal halide perovskite thin films," Chemistry of Materials, vol. 30, no. 14, pp. 4515-4526, 2018. [211] Y. Tan et al., "Highly Luminescent and Stable Perovskite Nanocrystals with Octylphosphonic Acid as a Ligand for Efficient Light-Emitting Diodes," ACS Applied Materials & Interfaces, vol. 10, no. 4, pp. 3784-3792, 2018/01/31 2018, doi: 10.1021/acsami.7b17166. [212] A. Mei et al., "A hole-conductor–free, fully printable mesoscopic perovskite solar cell with high stability," science, vol. 345, no. 6194, pp. 295-298, 2014. [213] X. Li et al., "Outdoor Performance and Stability under Elevated Temperatures and Long-Term Light Soaking of Triple-Layer Mesoporous Perovskite Photovoltaics," Energy Technology, vol. 3, no. 6, pp. 551-555, 2015, doi: 10.1002/ente.201500045. [214] G. Grancini et al., "One-Year stable perovskite solar cells by 2D/3D interface engineering," Nature Communications, vol. 8, no. 1, p. 15684, 2017/06/01 2017, doi: 10.1038/ncomms15684. [215] Y. Hu et al., "Improved Performance of Printable Perovskite Solar Cells with Bifunctional Conjugated Organic Molecule," Advanced Materials, vol. 30, no. 11, p. 1705786, 2018, doi: 10.1002/adma.201705786. [216] F. Zhang and M. Srinivasan, "Self-assembled molecular films of aminosilanes and their immobilization capacities," Langmuir, vol. 20, no. 6, pp. 2309-2314, 2004. [217] Y.-Q. Wang, S.-B. Xu, J.-G. Deng, and L.-Z. Gao, "Enhancing the efficiency of planar heterojunction perovskite solar cells via interfacial engineering with 3-aminopropyl trimethoxy silane hydrolysate," Royal Society open science, vol. 4, no. 12, p. 170980, 2017. [218] P. Fu, X. Guo, S. Wang, Y. Ye, and C. Li, "Aminosilane as a molecular linker between the electron-transport layer and active layer for efficient inverted polymer solar cells," ACS applied materials & interfaces, vol. 9, no. 15, pp. 13390-13395, 2017. [219] K. Xu et al., "Synergistic Surface Passivation of CH3NH3PbBr3 Perovskite Quantum Dots with Phosphonic Acid and (3-Aminopropyl)triethoxysilane," Chemistry – A European Journal, vol. 25, no. 19, pp. 5014-5021, 2019, doi: 10.1002/chem.201805656. [220] J. Zhu, Y. Zhu, J. Huang, Y. Gong, J. Shen, and C. Li, "Synthesis of CsPbBr3 perovskite nanocrystals with the sole ligand of protonated (3-aminopropyl)triethoxysilane," Journal of Materials Chemistry C, 10.1039/C9TC02089F vol. 7, no. 24, pp. 7201-7206, 2019, doi: 10.1039/C9TC02089F. [221] X. Li et al., "Improved performance and stability of perovskite solar cells by crystal crosslinking with alkylphosphonic acid ω-ammonium chlorides," Nature Chemistry, Article vol. 7, p. 703, 08/17/online 2015, doi: 10.1038/nchem.2324 https://www.nature.com/articles/nchem.2324#supplementary-information. [222] Q. Wang, X. Zheng, Y. Deng, J. Zhao, Z. Chen, and J. Huang, "Stabilizing the α-Phase of CsPbI₃ Perovskite by Sulfobetaine Zwitterions in One-Step Spin-Coating Films," Joule, vol. 1, no. 2, pp. 371-382, 2017, doi: 10.1016/j.joule.2017.07.017.

2-81

[223] X. Zheng et al., "Dual Functions of Crystallization Control and Defect Passivation Enabled by Sulfonic Zwitterions for Stable and Efficient Perovskite Solar Cells," Advanced Materials, vol. 30, no. 52, p. 1803428, 2018, doi: 10.1002/adma.201803428. [224] T. Zhang et al., "Multi-functional organic molecules for surface passivation of perovskite," Journal of Photochemistry and Photobiology A: Chemistry, vol. 355, pp. 42-47, 2018/03/15/ 2018, doi: https://doi.org/10.1016/j.jphotochem.2017.11.031. [225] F. Krieg et al., "Colloidal CsPbX3 (X = Cl, Br, I) Nanocrystals 2.0: Zwitterionic Capping Ligands for Improved Durability and Stability," ACS Energy Letters, vol. 3, no. 3, pp. 641-646, 2018/03/09 2018, doi: 10.1021/acsenergylett.8b00035. [226] Z. Liu et al., "Ligand Mediated Transformation of Cesium Lead Bromide Perovskite Nanocrystals to Lead Depleted Cs4PbBr6 Nanocrystals," Journal of the American Chemical Society, vol. 139, no. 15, pp. 5309-5312, 2017/04/19 2017, doi: 10.1021/jacs.7b01409. [227] S. Baek, Y. Kim, and S.-W. Kim, "Highly photo-stable CsPbI3 perovskite quantum dots via thiol ligand exchange and their polymer film application," Journal of Industrial and Engineering Chemistry, vol. 83, pp. 279-284, 2020/03/25/ 2020, doi: https://doi.org/10.1016/j.jiec.2019.11.038. [228] W. Dai, Y. Lei, M. Xu, P. Zhao, Z. Zhang, and J. Zhou, "Rare-Earth Free Self-Activated Graphene Quantum Dots and Copper-Cysteamine Phosphors for Enhanced White Light-Emitting-Diodes under Single Excitation," Scientific Reports, vol. 7, no. 1, p. 12872, 2017/10/09 2017, doi: 10.1038/s41598-017-13404-1. [229] Z. Salehi, R. V. Parish, and R. G. Pritchard, "A polymeric cationic copper(I) complex involving a quadruply bridging, zwitterionic thiolate ligand: {[Cu8Cl6(SCH2CH2NH3)6]Cl2}," Journal of the Chemical Society, Dalton Transactions, 10.1039/A705217K no. 22, pp. 4241-4246, 1997, doi: 10.1039/A705217K. [230] L. Riauba, G. Niaura, O. Eicher-Lorka, and E. Butkus, "A Study of Cysteamine Ionization in Solution by Raman Spectroscopy and Theoretical Modeling," The Journal of Physical Chemistry A, vol. 110, no. 50, pp. 13394-13404, 2006/12/01 2006, doi: 10.1021/jp063816g. [231] B. Li et al., "Constructing water-resistant CH3NH3PbI3 perovskite films via coordination interaction," Journal of Materials Chemistry A, 10.1039/C6TA06892H vol. 4, no. 43, pp. 17018- 17024, 2016, doi: 10.1039/C6TA06892H. [232] L. Men, B. A. Rosales, N. E. Gentry, S. D. Cady, and J. Vela, "Lead-Free Semiconductors: Soft Chemistry, Dimensionality Control, and Manganese-Doping of Germanium Halide Perovskites," ChemNanoMat, vol. 5, no. 3, pp. 334-339, 2019, doi: 10.1002/cnma.201800497. [233] N. Ahn, D.-Y. Son, I.-H. Jang, S. M. Kang, M. Choi, and N.-G. Park, "Highly Reproducible Perovskite Solar Cells with Average Efficiency of 18.3% and Best Efficiency of 19.7% Fabricated via Lewis Base Adduct of Lead(II) Iodide," Journal of the American Chemical Society, vol. 137, no. 27, pp. 8696-8699, 2015/07/15 2015, doi: 10.1021/jacs.5b04930. [234] J.-W. Lee et al., "Tuning Molecular Interactions for Highly Reproducible and Efficient Formamidinium Perovskite Solar Cells via Adduct Approach," Journal of the American Chemical Society, vol. 140, no. 20, pp. 6317-6324, 2018/05/23 2018, doi: 10.1021/jacs.8b01037. [235] L. Zhu et al., "Investigation on the role of Lewis bases in the ripening process of perovskite films for highly efficient perovskite solar cells," Journal of Materials Chemistry A, 10.1039/C7TA05378A vol. 5, no. 39, pp. 20874-20881, 2017, doi: 10.1039/C7TA05378A.

2-82

[236] A. Abate et al., "Supramolecular Halogen Bond Passivation of Organic–Inorganic Halide Perovskite Solar Cells," Nano Letters, vol. 14, no. 6, pp. 3247-3254, 2014/06/11 2014, doi: 10.1021/nl500627x. [237] T. Niu et al., "High performance ambient-air-stable FAPbI3 perovskite solar cells with molecule-passivated Ruddlesden–Popper/3D heterostructured film," Energy & Environmental Science, 10.1039/C8EE02542H vol. 11, no. 12, pp. 3358-3366, 2018, doi: 10.1039/C8EE02542H. [238] S. Yang et al., "Tailoring passivation molecular structures for extremely small open-circuit voltage loss in perovskite solar cells," Journal of the American Chemical Society, vol. 141, no. 14, pp. 5781-5787, 2019. [239] H. Zhang, Q. Liao, Y. Wu, J. Chen, Q. Gao, and H. Fu, "Pure zero-dimensional Cs 4 PbBr 6 single crystal rhombohedral microdisks with high luminescence and stability," Physical Chemistry Chemical Physics, vol. 19, no. 43, pp. 29092-29098, 2017. [240] Y. Wei et al., "Perovskite Quantum Dots: Enhancing the Stability of Perovskite Quantum Dots by Encapsulation in Crosslinked Polystyrene Beads via a Swelling–Shrinking Strategy toward Superior Water Resistance (Adv. Funct. Mater. 39/2017)," Advanced Functional Materials, vol. 27, no. 39, 2017, doi: 10.1002/adfm.201770230. [241] X. Yang et al., "Preparation of CsPbBr3@PS composite microspheres with high stability by electrospraying," Journal of Materials Chemistry C, 10.1039/C8TC01408F vol. 6, no. 30, pp. 7971-7975, 2018, doi: 10.1039/C8TC01408F. [242] Y. Zhao et al., "A polymer scaffold for self-healing perovskite solar cells," Nature Communications, vol. 7, no. 1, p. 10228, 2016/01/06 2016, doi: 10.1038/ncomms10228. [243] J. He, C. F. Ng, K. Young Wong, W. Liu, and T. Chen, "Photostability and Moisture Stability of CH3NH3PbI3‐based Solar Cells by Ethyl Cellulose," ChemPlusChem, vol. 81, no. 12, pp. 1292- 1298, 2016. [244] Y. Guo, K. Shoyama, W. Sato, and E. Nakamura, "Polymer stabilization of lead (II) perovskite cubic nanocrystals for semitransparent solar cells," Advanced Energy Materials, vol. 6, no. 6, p. 1502317, 2016. [245] N. Zhou et al., "Perovskite nanowire–block copolymer composites with digitally programmable polarization anisotropy," Science Advances, vol. 5, no. 5, p. eaav8141, 2019, doi: 10.1126/sciadv.aav8141. [246] H. Han et al., "Highly Photoluminescent and Environmentally Stable Perovskite Nanocrystals Templated in Thin Self-Assembled Block Copolymer Films," Advanced Functional Materials, vol. 29, no. 26, p. 1808193, 2019, doi: 10.1002/adfm.201808193. [247] J. Park et al., "Mussel-Inspired Polymer Grafting on CsPbBr3 Perovskite Quantum Dots Enhancing the Environmental Stability," Particle & Particle Systems Characterization, vol. 36, no. 12, p. 1900332, 2019, doi: 10.1002/ppsc.201900332. [248] S. M. Lee, H. Jung, W. I. Park, Y. Lee, E. Koo, and J. Bang, "Preparation of Water-Soluble CsPbBr3 Perovskite Quantum Dot Nanocomposites via Encapsulation into Amphiphilic Copolymers," ChemistrySelect, vol. 3, no. 40, pp. 11320-11325, 2018, doi: 10.1002/slct.201802237. [249] M. V. Kovalenko, M. Scheele, and D. V. Talapin, "Colloidal Nanocrystals with Molecular Metal Chalcogenide Surface Ligands," Science, vol. 324, no. 5933, pp. 1417-1420, 2009, doi: 10.1126/science.1170524.

2-83

[250] A. Nag, M. V. Kovalenko, J.-S. Lee, W. Liu, B. Spokoyny, and D. V. Talapin, "Metal-free Inorganic Ligands for Colloidal Nanocrystals: S2–, HS–, Se2–, HSe–, Te2–, HTe–, TeS32–, OH–, and NH2– as Surface Ligands," Journal of the American Chemical Society, vol. 133, no. 27, pp. 10612- 10620, 2011/07/13 2011, doi: 10.1021/ja2029415. [251] D. V. Talapin, W. Yuanyuan, and H. Zhang, "Photoactive, inorganic ligand-capped inorganic nanocrystals," ed: Google Patents, 2020. [252] R. Tangirala, J. L. Baker, A. P. Alivisatos, and D. J. Milliron, "Modular inorganic nanocomposites by conversion of nanocrystal superlattices," Angewandte Chemie International Edition, vol. 49, no. 16, pp. 2878-2882, 2010. [253] Z. Huang, Q. Zeng, Z. Bai, and S. Qin, "Regulating the Fluorescence Emission of CdSe Quantum Dots Based on the Surface Ligand Exchange with MAA," Polymers for Advanced Technologies, vol. 31, no. 11, pp. 2667-2675, 2020, doi: https://doi.org/10.1002/pat.4993. [254] F. Qiao and Y. Xie, "Strategies for enhancing conductivity of colloidal nanocrystals and their photoelectronic applications," Journal of Energy Chemistry, vol. 48, pp. 29-42, 2020/09/01/ 2020, doi: https://doi.org/10.1016/j.jechem.2019.12.022. [255] F. Qiao, Y. Xie, Z. Weng, and H. Chu, "Ligand engineering of colloid quantum dots and their application in all-inorganic tandem solar cells," Journal of Energy Chemistry, vol. 50, pp. 230- 239, 2020/11/01/ 2020, doi: https://doi.org/10.1016/j.jechem.2020.03.019. [256] H. Zhang, J. Jang, W. Liu, and D. V. Talapin, "Colloidal Nanocrystals with Inorganic Halide, Pseudohalide, and Halometallate Ligands," ACS Nano, vol. 8, no. 7, pp. 7359-7369, 2014/07/22 2014, doi: 10.1021/nn502470v. [257] Z. M. Norman, N. C. Anderson, and J. S. Owen, "Electrical Transport and Grain Growth in Solution-Cast, Chloride-Terminated Cadmium Selenide Nanocrystal Thin Films," ACS Nano, vol. 8, no. 7, pp. 7513-7521, 2014/07/22 2014, doi: 10.1021/nn502829s. [258] L. Protesescu, "Novel luminescent colloidal nanocrystals and studies on nanocrystal surface chemistry," ETH Zurich, 2016. [259] J.-S. Lee, M. V. Kovalenko, J. Huang, D. S. Chung, and D. V. Talapin, "Band-like transport, high electron mobility and high photoconductivity in all-inorganic nanocrystal arrays," Nature Nanotechnology, vol. 6, no. 6, pp. 348-352, 2011/06/01 2011, doi: 10.1038/nnano.2011.46. [260] M. G. Panthani et al., "High Efficiency Solution Processed Sintered CdTe Nanocrystal Solar Cells: The Role of Interfaces," Nano Letters, vol. 14, no. 2, pp. 670-675, 2014/02/12 2014, doi: 10.1021/nl403912w. [261] J. Huang et al., "Surface functionalization of semiconductor and oxide nanocrystals with small inorganic oxoanions (PO43–, MoO42–) and polyoxometalate ligands," Acs Nano, vol. 8, no. 9, pp. 9388-9402, 2014. [262] Y. Hou et al., "Enhanced moisture stability of metal halide perovskite solar cells based on sulfur–oleylamine surface modification," Nanoscale Horizons, 10.1039/C8NH00163D vol. 4, no. 1, pp. 208-213, 2019, doi: 10.1039/C8NH00163D. [263] V. K. Ravi et al., "Origin of the Substitution Mechanism for the Binding of Organic Ligands on the Surface of CsPbBr3 Perovskite Nanocubes," The Journal of Physical Chemistry Letters, vol. 8, no. 20, pp. 4988-4994, 2017/10/19 2017, doi: 10.1021/acs.jpclett.7b02192. [264] J. Y. Woo et al., "Air-Stable PbSe Nanocrystals Passivated by Phosphonic Acids," Journal of the American Chemical Society, vol. 138, no. 3, pp. 876-883, 2016/01/27 2016, doi: 10.1021/jacs.5b10273.

2-84

[265] R. Gomes et al., "Binding of Phosphonic Acids to CdSe Quantum Dots: A Solution NMR Study," The Journal of Physical Chemistry Letters, vol. 2, no. 3, pp. 145-152, 2011/02/03 2011, doi: 10.1021/jz1016729. [266] W.-k. Koh, S. Park, and Y. Ham, "Phosphonic Acid Stabilized Colloidal CsPbX3 (X=Br, I) Perovskite Nanocrystals and Their Surface Chemistry," ChemistrySelect, vol. 1, no. 13, pp. 3479-3482, 2016, doi: doi:10.1002/slct.201600809. [267] T. Xuan et al., "Highly stable CsPbBr3 quantum dots coated with alkyl phosphate for white light-emitting diodes," Nanoscale, 10.1039/C7NR04179A vol. 9, no. 40, pp. 15286-15290, 2017, doi: 10.1039/C7NR04179A. [268] C. Wang, A. S. R. Chesman, and J. J. Jasieniak, "Stabilizing the cubic perovskite phase of CsPbI3 nanocrystals by using an alkyl phosphinic acid," Chemical Communications, 10.1039/C6CC08282C vol. 53, no. 1, pp. 232-235, 2017, doi: 10.1039/C6CC08282C. [269] C. Lu et al., "Enhanced stabilization of inorganic cesium lead triiodide (CsPbI3) perovskite quantum dots with tri-octylphosphine," Nano Research, vol. 11, no. 2, pp. 762-768, 2018/02/01 2018, doi: 10.1007/s12274-017-1685-1. [270] A. A. M. Brown et al., "Self-assembly of a robust hydrogen-bonded octylphosphonate network on cesium lead bromide perovskite nanocrystals for light-emitting diodes," Nanoscale, 10.1039/C9NR02566A vol. 11, no. 25, pp. 12370-12380, 2019, doi: 10.1039/C9NR02566A. [271] M. Liu, A. Matuhina, H. Zhang, and P. Vivo, "Advances in the Stability of Halide Perovskite Nanocrystals," (in eng), Materials (Basel), vol. 12, no. 22, p. 3733, 2019, doi: 10.3390/ma12223733. [272] I. Lokteva, N. Radychev, F. Witt, H. Borchert, J. Parisi, and J. Kolny-Olesiak, "Surface Treatment of CdSe Nanoparticles for Application in Hybrid Solar Cells: The Effect of Multiple Ligand Exchange with Pyridine," The Journal of Physical Chemistry C, vol. 114, no. 29, pp. 12784- 12791, 2010/07/29 2010, doi: 10.1021/jp103300v. [273] M. Law, J. M. Luther, Q. Song, B. K. Hughes, C. L. Perkins, and A. J. Nozik, "Structural, Optical, and Electrical Properties of PbSe Nanocrystal Solids Treated Thermally or with Simple Amines," Journal of the American Chemical Society, vol. 130, no. 18, pp. 5974-5985, 2008/05/01 2008, doi: 10.1021/ja800040c. [274] G. Bree, C. Coughlan, H. Geaney, and K. M. Ryan, "Investigation into the Selenization Mechanisms of Wurtzite CZTS Nanorods," ACS applied materials & interfaces, vol. 10, no. 8, pp. 7117-7125, 2018. [275] C. B. Murray, C. R. Kagan, and M. G. Bawendi, "Synthesis and Characterization of Monodisperse Nanocrystals and Close-Packed Nanocrystal Assemblies," Annual Review of Materials Science, vol. 30, no. 1, pp. 545-610, 2000/08/01 2000, doi: 10.1146/annurev.matsci.30.1.545. [276] N. Gaponik, D. V. Talapin, A. L. Rogach, A. Eychmüller, and H. Weller, "Efficient Phase Transfer of Luminescent Thiol-Capped Nanocrystals: From Water to Nonpolar Organic Solvents," Nano Letters, vol. 2, no. 8, pp. 803-806, 2002/08/01 2002, doi: 10.1021/nl025662w. [277] J. Aldana, N. Lavelle, Y. Wang, and X. Peng, "Size-Dependent Dissociation pH of Thiolate Ligands from Cadmium Chalcogenide Nanocrystals," Journal of the American Chemical Society, vol. 127, no. 8, pp. 2496-2504, 2005/03/01 2005, doi: 10.1021/ja047000+. [278] E. H. Sargent, "Colloidal quantum dot solar cells," Nature photonics, vol. 6, no. 3, pp. 133-135, 2012.

2-85

[279] G. H. Carey, A. L. Abdelhady, Z. Ning, S. M. Thon, O. M. Bakr, and E. H. Sargent, "Colloidal quantum dot solar cells," Chemical reviews, vol. 115, no. 23, pp. 12732-12763, 2015. [280] Z. Pan, K. Zhao, J. Wang, H. Zhang, Y. Feng, and X. Zhong, "Near Infrared Absorption of CdSexTe1–x Alloyed Quantum Dot Sensitized Solar Cells with More than 6% Efficiency and High Stability," ACS Nano, vol. 7, no. 6, pp. 5215-5222, 2013/06/25 2013, doi: 10.1021/nn400947e. [281] J. Kang and L.-W. Wang, "High defect tolerance in lead halide perovskite CsPbBr3," The journal of physical chemistry letters, vol. 8, no. 2, pp. 489-493, 2017. [282] F. Liu et al., "Highly Luminescent Phase-Stable CsPbI3 Perovskite Quantum Dots Achieving Near 100% Absolute Photoluminescence Quantum Yield," ACS Nano, vol. 11, no. 10, pp. 10373-10383, 2017/10/24 2017, doi: 10.1021/acsnano.7b05442. [283] B. A. Koscher, J. K. Swabeck, N. D. Bronstein, and A. P. Alivisatos, "Essentially trap-free CsPbBr3 colloidal nanocrystals by postsynthetic thiocyanate surface treatment," Journal of the American Chemical Society, vol. 139, no. 19, pp. 6566-6569, 2017. [284] T. Ahmed, S. Seth, and A. Samanta, "Boosting the Photoluminescence of CsPbX3 (X = Cl, Br, I) Perovskite Nanocrystals Covering a Wide Wavelength Range by Postsynthetic Treatment with Tetrafluoroborate Salts," Chemistry of Materials, vol. 30, no. 11, pp. 3633-3637, 2018/06/12 2018, doi: 10.1021/acs.chemmater.8b01235. [285] Q. Zhou et al., "High-Performance Perovskite Solar Cells with Enhanced Environmental Stability Based on a (p-FC6H4C2H4NH3)2[PbI4] Capping Layer," Advanced Energy Materials, vol. 9, no. 12, p. 1802595, 2019, doi: https://doi.org/10.1002/aenm.201802595. [286] X. Zhao and Z.-K. Tan, "Large-area near-infrared perovskite light-emitting diodes," Nature Photonics, vol. 14, no. 4, pp. 215-218, 2020. [287] W. Xu et al., "Rational molecular passivation for high-performance perovskite light-emitting diodes," Nature Photonics, vol. 13, no. 6, pp. 418-424, 2019. [288] M.-H. Park, J. S. Kim, J.-M. Heo, S. Ahn, S.-H. Jeong, and T.-W. Lee, "Boosting efficiency in polycrystalline metal halide perovskite light-emitting diodes," ACS Energy Letters, vol. 4, no. 5, pp. 1134-1149, 2019. [289] S. Kumar, J. Jagielski, T. Tian, N. Kallikounis, W.-C. Lee, and C.-J. Shih, "Mixing Entropy-Induced Layering Polydispersity Enabling Efficient and Stable Perovskite Nanocrystal Light-Emitting Diodes," ACS Energy Letters, vol. 4, no. 1, pp. 118-125, 2019/01/11 2019, doi: 10.1021/acsenergylett.8b02013. [290] E. Yassitepe et al., "Amine-Free Synthesis of Cesium Lead Halide Perovskite Quantum Dots for Efficient Light-Emitting Diodes," Advanced Functional Materials, vol. 26, no. 47, pp. 8757- 8763, 2016, doi: 10.1002/adfm.201604580. [291] Z.-K. Tan et al., "Bright light-emitting diodes based on organometal halide perovskite," Nature Nanotechnology, vol. 9, no. 9, pp. 687-692, 2014/09/01 2014, doi: 10.1038/nnano.2014.149. [292] J. C. Norman, R. P. Mirin, and J. E. Bowers, "Quantum dot lasers—History and future prospects," Journal of Vacuum Science & Technology A: Vacuum, Surfaces, and Films, vol. 39, no. 2, p. 020802, 2021. [293] J. Xu et al., "Halide Perovskites for Nonlinear Optics," Advanced Materials, vol. 32, no. 3, p. 1806736, 2020, doi: 10.1002/adma.201806736. [294] S. Li et al., "Water-resistant perovskite nanodots enable robust two-photon lasing in aqueous environment," Nature Communications, vol. 11, no. 1, p. 1192, 2020/03/04 2020, doi: 10.1038/s41467-020-15016-2. 2-86

[295] A. A. Tabrizi, H. Saghaei, M. A. Mehranpour, and M. Jahangiri, "Enhancement of absorption and effectiveness of a perovskite thin-film solar cell embedded with Gold nanospheres," Plasmonics, pp. 1-14, 2021. [296] L. Xie, T. Zhang, and Y. Zhao, "Stabilizing the MAPbI3 perovksite via the in-situ formed lead sulfide layer for efficient and robust solar cells," Journal of Energy Chemistry, vol. 47, pp. 62- 65, 2020/08/01/ 2020, doi: https://doi.org/10.1016/j.jechem.2019.11.023. [297] T. Wang, T. Fang, X. Li, L. Xu, and J. Song, "Controllable Transient Photocurrent in Photodetectors Based on Perovskite Nanocrystals via Doping and Interfacial Engineering," The Journal of Physical Chemistry C, 2021. [298] Z. Dai et al., "Capillary-bridge mediated assembly of aligned perovskite quantum dots for high- performance photodetectors," Journal of Materials Chemistry C, 10.1039/C9TC01104H vol. 7, no. 20, pp. 5954-5961, 2019, doi: 10.1039/C9TC01104H. [299] K. S. Schanze, P. V. Kamat, P. Yang, and J. Bisquert, "Progress in perovskite photocatalysis," ed: ACS Publications, 2020. [300] Y. Dai, C. Poidevin, C. Ochoa‐Hernández, A. A. Auer, and H. Tüysüz, "A supported bismuth halide perovskite photocatalyst for selective aliphatic and aromatic C–H bond activation," Angewandte Chemie International Edition, vol. 59, no. 14, pp. 5788-5796, 2020. [301] M. A. Jahangir et al., "Nanocrystals: Characterization Overview, Applications in Drug Delivery, and Their Toxicity Concerns," Journal of Pharmaceutical Innovation, pp. 1-12, 2020. [302] D. Lombardo, M. A. Kiselev, and M. T. Caccamo, "Smart Nanoparticles for Drug Delivery Application: Development of Versatile Nanocarrier Platforms in Biotechnology and Nanomedicine," Journal of Nanomaterials, vol. 2019, p. 26, 2019, Art no. 3702518, doi: 10.1155/2019/3702518. [303] Z. Wu et al., "An excellent impedance-type humidity sensor based on halide perovskite CsPbBr3 nanoparticles for human respiration monitoring," Sensors and Actuators B: Chemical, p. 129772, 2021. [304] G. Li et al., "A “Turn-on” fluorescence perovskite sensor based on MAPbBr3/mesoporous TiO2 for NH3 and amine vapor detections," Sensors and Actuators B: Chemical, vol. 327, p. 128918, 2021. [305] A. Chizhov et al., "Photoresistive gas sensor based on nanocrystalline ZnO sensitized with colloidal perovskite CsPbBr3 nanocrystals," Sensors and Actuators B: Chemical, vol. 329, p. 129035, 2021. [306] J. Yang, Q. Bao, L. Shen, and L. Ding, "Potential applications for perovskite solar cells in space," Nano Energy, vol. 76, p. 105019, 2020/10/01/ 2020, doi: https://doi.org/10.1016/j.nanoen.2020.105019. [307] H. Cai, Y. Gu, Y.-C. Lin, Y. Yu, D. B. Geohegan, and K. Xiao, "Synthesis and emerging properties of 2D layered III–VI metal chalcogenides," Applied Physics Reviews, vol. 6, no. 4, p. 041312, 2019/12/01 2019, doi: 10.1063/1.5123487. [308] Q. Hao et al., "Phase identification and strong second harmonic generation in pure ε-InSe and its alloys," Nano letters, vol. 19, no. 4, pp. 2634-2640, 2019. [309] S. Chen and L.-W. Wang, "Thermodynamic Oxidation and Reduction Potentials of Photocatalytic Semiconductors in Aqueous Solution," Chemistry of Materials, vol. 24, no. 18, pp. 3659-3666, 2012/09/25 2012, doi: 10.1021/cm302533s.

2-87

[310] H. L. Zhuang and R. G. Hennig, "Single-Layer Group-III Monochalcogenide Photocatalysts for Water Splitting," Chemistry of Materials, vol. 25, no. 15, pp. 3232-3238, 2013/08/13 2013, doi: 10.1021/cm401661x. [311] Y. Sun, R. Wang, and K. Liu, "Substrate induced changes in atomically thin 2-dimensional semiconductors: Fundamentals, engineering, and applications," Applied Physics Reviews, vol. 4, no. 1, p. 011301, 2017.

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Chapter 3 Experimental Methods and Characterisation

Techniques

This experimental chapter revises the main synthesis techniques following their brief discussion in Chapter 2. The characterisation techniques used in the following experimental chapters are presented, discussing how they work and the information they relay about the analysed materials.

“Before anything else, preparation is the key to success.”

Alexander Graham Bell, 1847-1922

3.1 Synthesis Techniques

There are two main synthesis techniques used throughout this work. The organic, colloidal synthesis is a variation of the “Hot-Injection” method first developed by Murray et al. in

1993.[1] This approach involves the thermal decomposition of organometallic precursors in a high boiling point organic solvent, with surfactants employed to instigate size and shape control over the resultant NCs. (Refer to Chapter 2, Section 1.2.1)

In this work, perovskite NCs are produced in Chapter 4 using phosphorous based ligands to exert control over the NCs and open a different route to the production of these particles than the traditional, limited OLA/OA combination. Chapter 5 utilises the hot injection method to produce 2D indium chalcogenide nanosheets and nanoribbons, alloying the chalcogenide site.

Following the hot injection and LARP method's failures to produce germanium perovskites, a

Solution Temperature-Lowering (STL) Method was employed. The STL method is a classic crystallisation method that creates supersaturation by lowing the temperature. It is suitable to

3-89 grow perovskite in the HX-based precursor solution because its solubility dramatically decreases as the temperature goes down. Chapter 6 explores the synthesis of a tin-germanium alloyed perovskite with an iodide and bromide halide. This method does not form nanocrystals; however, to control the crystallisation of the crystals, a range of surfactants are added in

Chapter 7.

3.2 Characterisation Techniques

A range of characterisation techniques is necessary to understand material properties

(structural, chemical, optical, etc.). In any wet chemical synthesis, the parameters (temperature, growth time, solvent use, etc.) discussed previously can significantly reduce chemical precursors into the desired compound, resulting in impurities and defects in the product. The concentration and type of surfactants can also affect the coating efficiency on the nanocrystal surface that influences surface oxidation or luminescence quenching.

Powder X-ray Diffraction (XRD) is one of the most common analytical techniques used to determine the crystal phase and purity. Using the Scherrer equation, Bragg reflections in the

XRD pattern can also indicate NC products' crystallite sizes. The preliminary technique to confirm the material compound's chemical makeup is energy dispersive spectroscopy (EDS).

It is usually coupled with an electron microscope. The Transmission Electron Microscope

(TEM)

1.1.1 X-Ray Diffraction (XRD)

X-Ray Diffraction (XRD) is an analytical tool used to detect internal information about the composition, crystallography and grain size of nanocrystals, both qualitatively and quantitatively. German physicist Max von Laue received a Nobel Prize in Physics in 1914 for his discovery of the diffraction of X-rays by a crystal lattice. This discovery, combined with

3-90 the introduction of Bragg’s Law, which relates X-ray scattering with reflections from atomic planes, lead to considerable developments in the field of crystallography.

The wavelength of the incident beam (λ) and the spacing (푑ℎ푘푙) between two adjacent planes where scattering occurs determines the direction of the diffracted beam. The diffracted beam's intensity is a function of atomic lattice distribution and the crystal planes' orientation relative to the incident beam (θ). Diffraction occurs only if the wave is scattered by two atomic planes resulting in a total path difference (2∆푃), an integer (푛) multiple of the wavelength of incident

λ. Equation 3-1a and Figure 3-1 (a) Beam path difference in Bragg's Law; (b) Beam path of an

XRD show Bragg’s Law:

푛휆 = 2Δ푃 = 2푑ℎ푘푙 sin θ Equation 3-1

The diffracted beam's intensity is measured by converting the electron charge into bias voltage, which is proportional to the X-Ray beam intensity (Figure 3-1b).

Figure 3-1 (a) Beam path difference in Bragg's Law; (b) Beam path of an XRD The X-Ray target used in this study is copper with emits 8 keV with a wavelength of 1.54 Å.

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1.1.2 Scanning Electron Microscopy (SEM)

Scanning Electron Microscopy is one of several techniques used to characterise nanomaterials that use a high energy electron beam to gain insights into the material's surface. It is similar to optical microscopy, except for photons' substitution with electrons due to the fine tunability of wavelengths that a magnetic field can control. An electron beam interacts with solid material in numerous ways described in Figure 3-2. The SEM is a non-destructive imaging tool where the Secondary Electrons (SE) and Backscattered Electrons (BSE) are specifically collected from the microscope.

EDX is discussed later in the chapter; however, when a detector is present in an SEM, the characteristic X-rays originate deeper within the sample than SE and BSE, so a higher penetration depth is needed. This depth is facilitated by using higher beam energy. Typically

20 keV for materials containing caesium.

Figure 3-2 Interactive phenomena of an electron beam with a solid material. 3-92

1.1.3 Energy-Dispersive X-Ray Spectroscopy (EDX)

The atomic structure is essentially the fingerprint of an element. Its number of protons, orbits, number of electrons and excited electrons are unique. When an accelerated beam reaches the sample surface, a ground-state electron is excited to a free electron state or a higher unoccupied energy level, creating a hole in the ground state. Hence, a higher state electron falls to this hole by releasing energy in the form of an X-ray or Auger electron (Figure 3-2). These signals provide quantitative chemical, electrical, and structural information. The amount of emitted X- rays is counted as a function of photon energy using an energy-dispersive spectrometer. EDX is better suited to heavy elements due to the fluorescence associated with Auger emission for lighter elements. The EDX detector can be installed on an SEM or TEM.

1.1.4 Transmission Electron Microscopy (TEM)

The Transmission Electron Microscope (TEM) is a powerful tool in nanotechnology research, capable of investigating the morphology, crystal structure (through electron diffraction) and chemical composition (through EDX or EELS) of a wide range of nanomaterials. [2] A high voltage electron gun emits a monochromatic electron beam that travels down the column through the condenser lenses and aperture, which control the spot size and intensity. It transmits through the sample as a localised probe (Figure 3-2). The condenser aperture removes electrons diffracted to a high angle. Subsequently, the beam passes through the objective lenses and apertures that ultimately focus the electrons onto the fluorescent viewing screen or Charge

Coupled Device (CCD) camera at the bottom of the column. Figure 3-3 shows the simplified apparatus. The entire system operates under vacuum to prevent air from interfering with the electrons. [2]

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Figure 3-3 Simplified cross-section of a TEM Typically, a copper grid coated with a thin coating of carbon holds a solution deposited sample after the solvent evaporated completely. High contrast images form when the sample is highly crystalline and orientated to the beam, resulting in many diffracted electrons. Higher contrast is associated with atoms of a large atomic number as they scatter more of the incident beam, so the image appears darker. If both of these phenomena happen simultaneously, Dark Field

TEM (DFTEM) is used to provide more information. In DFTEM, the direct beam is blocked, leaving only the diffracted electrons to pass through the objective aperture, similar to the SEM but with a much higher resolution.

1.1.4.1 Electron Diffraction (ED)

Electrons are a type of ionising radiation, meaning they can remove tightly bound inner shell electrons from the atom, resulting in the wide range of secondary signals shown in Figure 3-2.

These enable structural analysis, Energy-Dispersive X-ray spectrometry (EDX) and Electron

Energy Loss Spectroscopy (EELS).

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Electron Diffraction (ED), first used in a TEM in the 1940s, is a powerful technique for characterising thin materials' structure. It is advantageous over other methods (e.g. X-Ray or

Neutron diffraction) due to the extremely short wavelength (~2 pm), strong atomic scattering, and the ability to examine tiny volumes of matter (~10 nm3) [3]. Electrons experience more scattering than X-rays or neutrons because they interact with the nucleus and the scattering atoms' electrons through Coulomb forces.

The scattering of electrons in a sample creates an interference pattern unique to the composition of the material. The diffraction pattern appears in reciprocal space above the image that forms in real space, shown in Figure 3-4. Therefore, a large lattice has a corresponding small distance between the ED pattern spots.

Complete structural determination is possible as the angles between the spots correspond to angles between lattice planes. And the interplanar spacing is revealed by a Fast Fourier

Transform (FFT), which transforms between the real and reciprocal space.

Figure 3-4 Optical ray diagram with an optical objective lens showing the principle of the imaging process in a TEM Inserting a Selected Area Electron Diffraction (SAED) aperture and using a parallel incident beam gives diffraction information for an area as small as 100 nm in diameter. Convergent beam electron diffraction uses a conical beam, providing information that allows the unique

3-95 determination of all the point groups and most space groups of the sample. SAED works at low resolution on almost every grain in a polycrystalline sample, meaning that it is possible to obtain the crystal structure and orientation from crystals that would otherwise be too small for

X-ray or neutron diffraction, such as nanocrystals.

1.1.5 UV-Vis Absorption Spectroscopy (UV-Vis)

UV-Vis is one of the oldest and widely used techniques used to investigate photonic energy's effect on compounds' optical properties. When monochromatic incident light interacts with electrons in the sample, part of the energy is absorbed (Figure 3-5b), and while the remained reaches the detector (Figure 3-5 (a) Spectroscopic Principles, (b) Figure 3-5a). The amount of energy detected is mathematically transformed to determine the absorbance value of the solution.

Figure 3-5 (a) Spectroscopic Principles, (b) UV-Vis Instrumentation

The absorption intensity (퐴) is directly proportional to the molar absorptivity (휀), the path length of the sample (푙) and analyte concentration (푐) via the Beer-Lambert law:

퐴 = 휀푙푐 Equation 3-2

The absorptivity is highly related to the structure of the analysed molecule. In a semiconductor, absorption of a photon lifts an electron from the valance band (VB) to the conduction band

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(CB), and the wavelength maxima (λmax) of the absorption peak correlates to the bandgap energy of the material. The bandgap energy increases as crystallite size decreases due to quantum confinement, resulting in a blue shift of the λmax relative to the bulk. The spectra are generally broad due to several vibrational and rotational levels within the electronic levels.

Extrapolation of the absorption onset to the baseline intersection point shows the bandgap energy. Otherwise, the Tauc analysis accounts for the difference in electronic transitions of direct and indirect bandgap semiconductors.

The Tauc analysis demonstrates the proportionality between the absorption coefficient and the available density of states (ℎ푣 − 퐸𝑔) by Equation 3-3:

푛 (푎ℎ푣) = 퐴 (ℎ푣 − 퐸𝑔) Equation 3-3

Where A is the proportionality coefficient, hv is the photon energy, and Eg is the bandgap. The value of n (1/3, 1/2, 2 or 2/3) is determined by the interband transitions of indirect forbidden, indirect allowed, direct allowed, or directly forbidden, respectively. The bandgap can be obtained from the baseline intersection, as explained previously.

In this work, UV-Vis spectroscopy examined the size-dependent optical bandgap of the resulting semiconductor NCs. Typically, purified samples were dispersed in a non-polar medium (hexane, toluene) and contained in quartz cuvettes for analysis. A Cary 5000 UV-Vis-

NIR Spectrometer (Agilent Technologies) equipped with a deuterium source collected all absorption spectra. Instrumental parameters such as slit width and cell path length were kept constant through each chapter's experiments.

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1.1.6 Solid State Diffuse Reflectance UV-Visible Spectroscopy (DRS)

The optical phenomenon known as diffuse reflectance is used in the UV-visible, near-infrared

(NIR), and mid-infrared (sometimes called DRIFT or DRIFTS) regions to obtain molecular spectroscopic information from powders. When the individual particles' size is comparable to the wavelength (0.2-3 µm), it is impossible to distinguish between reflection, refraction, and diffraction as the source of light scattering. A reflectance spectrum is collected and analysed according to surface-reflected electromagnetic radiation as a function of frequency ν, wavenumbers, cm-1 or wavelength λ, (nanometers, nm). Two different reflection types can occur: specular reflection, usually associated with reflection from smooth, polished surfaces like mirrors, and diffuse reflection, associated with absorption, reflection, and diffraction from rough solid surfaces, like powders (Figure 3-6). These reflection beams are collected as they retain sample information.

Figure 3-6 Types of reflection from a sample surface.[4] The Kubelka-Munk function describes the reflectance, R, from the powder in terms of a molar absorption coefficient, K, and a scattering coefficient, S. It is applied to obtain absorption information for the material.

(1 − 푅)2 퐾 푓(푅) = = 2푅 푆

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For this thesis experiments, a Cary 5000 UV–vis–NIR spectrophotometer (Agilent

Technologies) equipped with a praying mantis DFS attachment performed the reflectance measurements. This technique serves as a complementary one to the solution UV-Vis absorption technique to explore the optical bandgaps of materials.

1.1.7 Photoluminescence Spectroscopy

Photoluminescence (PL) refers to the emission of light by exciting molecules that absorb energy from an incident excitation source (Figure 3-5a). It is also considered the result of radiative recombination of electron-hole pairs in a semiconductor. In general, both the ground state and excited state consist of many vibrational and rotational energy levels. Absorption of a photon of energy excites electrons from the VB into the CB, creating an electron-hole pair followed by relaxation from the CB back to the VB with photons' release. This radiative process results in fluorescence energy that is characteristic of the material. If the energy absorbed is more than this transition, the relaxation from higher levels manifests as heat.

PL Quantum Yields (PLQYs) are measured relative to a standard dye. The NC samples' optical densities and the dye need to be equal (expected value of 0.08) and the emission spectra recorded from the same excitation wavelength. The emission spectra are integrated, and the values compared to give a percentage yield.

Alongside absorption spectroscopy, the effect of size on the optical properties of the NCs can be investigated using PL spectra. Also, relative quantum efficiency can be determined using this technique. All PL measurements were performed using a Cary Eclipse fluorescence spectrophotometer (Agilent Technologies) equipped with a xenon arc lamp.

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Figure 3-7 Set up of a photoluminescence spectrometer. 1.1.8 XPS

X-ray photoelectron spectroscopy (XPS) is a surface analytical technique used to determine the species and chemical bonding of atoms in the surface region of a solid. The basis of photoelectron spectroscopy is the photon (from X-ray beam, Al Kα or Mg Kα) interacts with the atom on the surface of the specimen. Kinetic energy (Ekin) is imparted onto an electron by the incoming photons (whose energy is hν) to overcome the binding energy of the electron within that orbit (EB). Φ is the sample work function shown in

Figure 3-8 Schematic of the photoemission process involved in XPS surface analysis. Φ is the work function, determined by the properties of the sample and spectrometer.

An electron detector records the XPS spectrum by measuring the kinetic energy Ekin and the number of escaped electrons under ultra-high vacuum. The energies of the escaped electrons

3-100 are determined by the chemical environment, indicating the Coulomb interaction with other electrons and with the nuclei. Because each element has a unique structure and the electron binding energy (EB) increases as a function of atomic number (except Hydrogen and Helium), the adjacent elements throughout the periodic table can be easily distinguished.

When the X-ray beam is insufficient to excite the electron in the K shell, it creates vacancies on the L or M shells instead. The photoemission from p, d and f electronic states with non-zero orbital angular momentum produces spin-orbit splitting lines, quickly readable from the spectrum. A notable chemical shift is observed from the XPS spectrum when the elements' chemical bonding states are changed. The spatial redistribution of the valence electron charges and the atom's changed potential causes this shift. When the atom exhibits greater electronegativity (accepts more electrons), the electronic charge's displacement becomes significant. Thus the binding energies are higher in the spectrum. [5]

XPS measurements in this work used a Kratos Axis 165 spectrometer. High-resolution spectra used monochromated Al Kα radiation energy of 1486.6 eV at fixed pass energy of 20 eV. Peak simulation used a mixed Gaussian-Lorentzian function with Shirley-type background subtraction. Samples were flooded with low energy electrons for efficient charge neutralization. The binding energies were determined using C 1s at 284.8 eV as the charge reference.

3.3 Bibliography

[1] C. B. Murray, D. J. Norris, and M. G. Bawendi, "Synthesis and characterization of nearly monodisperse CdE (E = sulfur, selenium, tellurium) semiconductor nanocrystallites," Journal of the American Chemical Society, vol. 115, no. 19, pp. 8706-8715, 1993/09/01 1993, doi: 10.1021/ja00072a025. [2] D. B. Williams and C. B. Carter, "The Transmission Electron Microscope," in Transmission Electron Microscopy: A Textbook for Materials Science. Boston, MA: Springer US, 1996, pp. 3- 17.

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[3] L. A. Bendersky and F. W. Gayle, "Electron Diffraction Using Transmission Electron Microscopy," Journal of research of the National Institute of Standards and Technology, vol. 106, no. 6, pp. 997-1012, 2001, doi: 10.6028/jres.106.051. [4] J. K. Devineni and H. S. Dhillon, "Ambient Backscatter Systems: Exact Average Bit Error Rate Under Fading Channels," IEEE Transactions on Green Communications and Networking, vol. 3, no. 1, pp. 11-25, 2019, doi: 10.1109/TGCN.2018.2880985. [5] J. Moulder, W. Stickle, P. Sobol, and K. Bomben, "A reference book of standard spectra for identification and interpretation of XPS data," Perkin-Elmer Corporation, 1995.

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Chapter 4 Synthesis and Dimensional Control of CsPbBr3

Perovskite Nanocrystals Using Phosphorous Based Ligands

This chapter is presented as published in the Journal of Chemical Physics, with the inclusion

of Supporting Information figures in the main text for ease of reading.

F McGrath, UV Ghorpade, KM Ryan; J. Chem. Phys. 152, 239901 (2020)

4.1 Abstract

Nanocrystals of the inorganic perovskite, CsPbBr3, display outstanding photophysical properties, making them ideal for next-generation optical devices. However, the typical combination of oleic acid and oleylamine ligands employed in their synthesis is easily displaced, leading to poor stability that can hinder their applicability. This work looks to replace oleic acid and amine with phosphorus-based ligands. CsPbBr3 nanocrystals with near- perfect monodispersity were synthesised with an oleylamine/alkyl phosphonic acid combination with the ability to tune the bandgap by varying the alkyl chain length. We further investigate the replacement of the oleylamine giving a ligand combination of alkyl phosphonic acid/trioctylphosphine oxide (TOPO) for perovskite nanocrystal nucleation and growth. This combination is typical for the widely studied metal chalcogenide synthesis and, when used in our study with CsPbBr3, yields a pure phase perovskite.

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4.2 Introduction

Lead halide perovskites (CsPbX3, where X = I, Br, Cl) have been the focus of intense research over the past decade due to their potential in photovoltaics and light-emitting devices. [1, 2] In particular, all inorganic CsPbX3 NCs display outstanding photophysical properties, such as high photoluminescence quantum yield (PLQY) and narrow full width half maximum

(FWHM), which sets them above traditional semiconducting nanomaterials. [3, 4] Their encouraging properties originate from the defect-free electronic structure of the bulk perovskite, combined with the ability to tune their bandgap energy via compositional and nanoscale size variations resulting in quantum confinement effects. [5]

There are many reported syntheses of different shapes of CsPbX3 nanoparticles with focused size distributions.[6-12] A challenge is the chemical and colloidal instability of CsPbX3 NCs when exposed to polar solvents limits their widespread utilisation. The perovskite crystal structure exhibits a predominantly ionic behaviour compared to metal chalcogenide nanocrystals with a partial covalent character. This ionicity weakens their interactions with organic ligands increasing their vulnerability to ligand loss during washing, ultimately leading to chemical instability, ion migration, and worsening carrier transport. [13, 14]

The nanoparticle's ligand layer is the key parameter controlling the solvent dispersibility, effective volume, adhesion to substrates, functionalization, and surface charge. [15, 16]

CsPbX3 NCs are conventionally synthesised using non-coordinating solvents, employing a mixture of primary amines and carboxylic acids to solvate PbX2, inhibit particle aggregation, and determine the extent of hydrophilicity on the surface of the NPs. The oleylamine (OAm), oleylammonium halides, and oleylammonium oleates are the primary surface-capping agents of CsPbX3 NCs. Their unsaturated nature provides a high dispersibility in organic solvents.

[16, 17]. Alkylammonium ligands are much more mobile than carboxylate ligands. They are 4-104 susceptible to detachment from the NC surface during centrifugation or contact with polar solvents, making it difficult to retain the crystals' PL properties. [18] Excess amines drive the dissolution of perovskite NPs and their transformation into an insulating phase, Cs4PbBr6. [19]

Further, the carboxylates are less effective than the alkylammonium ions in modulating the size of the NCs. [18] Therefore, stable perovskites require alternative surfactants.

De Roo et al. showed that surface Cs and Pb ions bind to either oleate or Br- while surface halides bind to oleylammonium ions via hydrogen bonding or electrostatic interactions.

Oleylammonium ions act as capping agents by substituting Cs+ cations preferentially over hydrogen bonding to the halide ions.[20] Therefore, ligands that would remain attached during processing would form strong metal-ligand bonds or stable hydrogen bonds with halide ions.

Phosphonic acids (PAs) bind strongly to divalent cadmium and lead anions and undergo strong hydrogen bonding to perovskites. [21-23] The presence of Tetradecylphosphonic acid (TDPA) produced colloidal perovskite NCs which withstand washing in H2O and has since been shown as suitable for use in room temperature synthesis. [24, 25] Wang et al. [16] found that a phosphinic acid stabilizes the alpha phase of CsPbI3. They saw through NMR and Diffusion

Ordered Spectroscopy (DOSY) that the phosphinic acid exists in an ion pair with oleylammonium and does not play a surfactant role in stabilizing the α-CsPbI3, meaning that only the oleylammonium iodide, from OAm is the dominant surface ligand. Separately, α-

CsPbI3 was stabilized through washing with trioctylphosphine (TOP), leading to over a month's quantum yield efficiency stability. [26] In related studies, Octylphosphonic acid (OPA) dramatically enhanced the perovskite stability towards polar solvents, producing an LED with an EQE of 7.74%. [27]

The use of branched capping ligands is a route for improving NC stability as they can provide a strong steric effect. [28] By incorporating TOPO into the reaction system, monodisperse

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CsPbX3 NCs form at higher temperatures than achievable using the traditional ligands and with the advantage of stability towards washing with ethanol. Almeida et al. conducted an amine- free synthesis using TOPO and OA to investigate how the system's acidity affects the reactivity of the PbBr2 and the size of the NCs. [29] Notably, through NMR and FTIR study, they found that TOPO was virtually absent from the particle ligand shell. The addition of TOPO to the toluene antisolvent in a ligand assisted reprecipitation synthesis (LARP) significantly reduced the surface defects of the NPs and increased the photoluminescence lifetimes as well as stability of LED devices. [30] [31]. An alternative amine-free synthesis involving quaternary alkylammonium halides in place of alkylamines worked well to produce NCs. However, the

LED-based on this work only yielded a peak EQE of 0.325% due to moderate electroluminescence intensity.[32]

The variety within these reports suggests that phosphorus-based ligand syntheses may be realistic routes towards a stable, reproducible, high yield production of perovskite NCs. Here, this work synthesized CsPbBr3 NCs using a phosphine ligand-based protocol. For the first time, alkyl phosphonic acids' effect as a replacement for OA is systematically studied, displaying an evident influence over NC size. We use PAs of increasing chain length, in conjunction with

OAm, to examine the feasibility of replacing OA in the precursor solution to find a more stable surfactant. Examination of the morphology, chemical structure and optical properties presents small near monodisperse NCs able to tune the size and material bandgap. Following encouraging findings, we examined the use of TOPO as the second ligand in place of OAm.

Similar characterisations demonstrated the feasibility of moving away from the traditional

OA/OAm ligand combination. These results show that PAs offer numerous possibilities to broaden the ligand chemistry of perovskites, extending to morphology control, stability, and expanded reaction regimes.

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4.3 Experimental Section

Materials. Octylphosphonic acids (>99%, OPA), Decylphosphonic acid (99%, DPA),

Dodecylphosphonic acid (>99%, DDPA), Hexadecylphosphonic acid (>99%, HDPA),

Octadecylphosphonic acid (>99%, ODPA) and tetradecylphosphonic acid (>99%, TDPA) were purchased from PCI Synthesis. Acetone (99.5%) and toluene (99%) were purchased from

Lennox. Caesium carbonate (Cs2CO3, 99%), oleylamine (70%, OAm), oleic acid (90%, OA),

1-octadecene (90%), PbBr2 (≥99%) and Trioctylphosphine oxide (99%, TOPO) were purchased from Sigma-Aldrich. All chemicals were used without any further purification unless otherwise stated.

Preparation of Cs Precursors. Cs-Oleate (0.15 M) in 1-Octadecene. Cs2CO3 (0.407 g, 1.25 mmol) and OA (1.7 mL, 5.4 mmol) were degassed in 20.0 mL of ODE in a three-neck round- bottomed flask under vacuum at 100 °C for 1 h followed by reacting under argon at 150 °C until all Cs2CO3 dissolves. The Cs-oleate was stored in an argon atmosphere.

Synthesis of CsPbBr3 Nanoparticles (NPs). In a typical synthesis, PbBr2 (69 mg, 0.16 mmol) and PA (0.5 mmol): OPA (100 mg), DPA (110 mg), DDPA (125 mg), TDPA (140 mg), HDPA

(150 mg), or ODPA (170 mg), were dried at to 120°C in ODE (6.0 mL) for 1 h. Under argon,

OAm (0.25 ml) was injected, and the temperature increased to 210 °C to form a solution. The system was cooled to the desired reaction temperature (180°C), and Cs-oleate was injected (0.6 mL, preheated to 120°C). The TOPO reaction was identical, except for the replacement of

OAm with TOPO (1.0 g, 0.88 g/mL, 2.59 mmol), which was incorporated into the flask at the beginning.

After 5 s of growth, the vial was plunged into a water bath to quench the reaction. 5ml of anhydrous toluene was added to the quenched solution to wash the NCs. The suspension was

4-107 then centrifuged at 13000 rpm and washed once more before re-dispersion in anhydrous hexane. All washing occurred in an inert atmosphere.

Transmission Electron Microscopy (TEM). NC dispersions were drop-cast on carbon-coated

200 mesh copper grids. We acquired bright-field TEM images on a JEOL TEM-2100 microscope (W filament) operating at an accelerating voltage of 200 kV.

X-ray Diffraction (XRD). Dried and milled NC powder was analysed on a zero diffraction silicon substrate. We conduct XRD measurements on a PANalytical Empyrean X-ray diffractometer equipped with a 1.8 kW Cu Kαceramic X-ray tube and PIXcel3D 2×2 area detector, operating at 40 kV and 40 mA.

Steady-State UV−Vis Extinction Spectroscopy and Steady-State Photoluminescence

Spectroscopy. We recorded optical extinction and photoluminescence spectra of anhydrous toluene dispersions in quartz cuvettes with a 1 cm path length, employing an Agilent Cary 5000

UV−vis spectrophotometer and a Varian Cary Eclipse fluorescence spectrophotometer, respectively. Samples were stored in Argon before measuring.

4.4 Results and Discussion

Octylphosphonic acid (OPA) was the first and shortest PA investigated as an OA replacement.

Figure 4-1a shows a TEM image of NCs produced by the hot injection of Cs-Oleate into a solution of PbBr2 dissolved in OAm, OPA and ODE. The particles were nearly monodisperse with an average NC diameter of 6.2 ± 0.3 nm. Selected Area Electron Diffraction (SEAD) of multiple NCs verifies the sample crystallinity, with the three most intense diffraction rings of monoclinic CsPbBr3 highlighted in Figure 4-1b. This data agrees with the High-Resolution

TEM (HRTEM) in Figure 4-1c, which shows a d-spacing of 0.58 nm corresponding to the

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(100) plane. Figure 4-1d displays the Fast Fourier Transform (FFT) of one of the particles showing the (110) and (100) planes of monoclinic CsPbBr3 (Ref: 00-054-0751).

Figure 4-1a) ) TEM analysis of OPA/OAm capped CsPbBr3. b) SAED pattern of the image in a. c+d) HRTEM of the NCs and the corresponding FFT of a single particle (bottom right). e) XRD pattern of the OPA/OAm capped perovskite showing two phases along with the reference patterns of both. The corresponding X-Ray Diffraction (XRD) pattern of the OPA capped NCs (Figure 4-1e) show the expected peaks for monoclinic CsPbBr3 in agreement with the SAED. However, additional impurities are present corresponding to lead depleted perovskite, Cs4PbBr6. PA ligands bind with metal Pb ions and halide ions while the oleylammonium ion, formed from

OAm, binds to the halide. [17] The smaller amount of lead in the lead depleted sample may

4-109 correlate with a reduction of the OAm ligand coating on the particle surface and hence lower dispersibility than the CsPbBr3 NCs that were collected for TEM.

We can learn from the development of chalcogenide research for the synthetic control of perovskite materials with phosphorus-based ligands. [33-38] Several reports have shown excellent control of size and morphology across the entire particle distribution.[39-42] There is an understanding of the importance of different ligand classifications on the growth of NCs.

[43-47] As an example, Wang et al. [48] showed that varying PA length resulted in incremental size effects. In the interest of developing other shape and size control of CsPbBr3 NCs, PAs with chain lengths between 8 and 18 carbon atoms were employed in this synthesis alongside

OAm (See the Experimental Section for details.) Figure 4-2a-f shows the resulting TEM images. Figure 4-2g shows the samples' size distribution as measured (N = 60 particles) determined from the TEM images' profile analysis in Gatan Digital Micrograph. Short-chain ligands OPA and DPA allowed the formation of nearly monodisperse NCs of about 6.3 nm and

7.5 nm, respectively. As the chain length increased to C12, the polydispersity increased, and in the DDPA/OAm sample, a bimodal distribution of crystal size was identified. An outlier to this study was the TDPA/OAm sample, which produced relatively small nanoparticles (7 nm).

However, the samples' general trend is an increase in size as the PA chain length increases.

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Figure 4-2 a-f) NCs capped with each PA and OAm. g) Particle size distribution corresponding to the TEM images. As a comparison, Supporting Figure 4-3 shows the syntheses' SEM image involving traditional ligands OA and OAm form NCs with a size of around 20 nm.

Figure 4-3 Supporting image of NCs synthesised using OAm and OA ligands.

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Compared with these, the PA capped NCs were considerably smaller. According to a study of

PbSe, phosphorous based ligands have a strong binding affinity towards lead; P-O- moieties passivates the Pb on the surface of the final PbSe NC. [21, 49] The study by Brown et al. [27] reported the improved LED performance with OPA incorporation, finding that OPA ligands

- attach strongly to the CsPbBr3. Octylphosphonates bind preferentially through P-O , resulting in hydrogen bonds between neighbouring octylphosphonate chains. This P-O- moiety bonds to lead from the PbSe study, and so a hydrogen network likely forms around the PbBr2 precursor before the caesium precursor injection. Therefore, the formation of the perovskite lattice undergoes slower reaction kinetics. This decelerated growth mechanism would be responsible for the smaller nanocrystals in comparison to the carboxylic case.

Figure 4-4 shows the XRD analysis of the samples showing the CsPbBr3 perovskite but also the presence of Cs4PbBr6. This lead depleted, vacancy ordered Cs4PbBr6 perovskite has

4- outstanding thermal, chemical and photo-stability. The individual octahedra of [PbX6] are perfectly isolated by caesium cations in the crystal lattice. [50, 51] However, DFT calculations show that it has a theoretical bandgap of 3.84 eV and might be a beneficial impurity that increases CsPbBr3 stability, but not as a stand-alone luminescent material.[51]

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Figure 4-4 XRD patterns of the six samples capped with each PA and OAm. The CsPbBr3 peaks are identified with the *, with the remaining peaks corresponding to the Cs4PbBr6. Researchers have shown that excess OAm concentration results in the formation of lead depleted Cs4PbBr6.[52] Here, in our studies, OAm is kept at a constant minimal concentration sufficient to dissolve the PbBr2. This progression to a multiphase sample may result from the lack of Pb-oleate in the precursor solution. The carboxylic acid has one bonding site, while the phosphonic acid has two sites on which it can bond with the lead precursor. This extra bonding site may result in a more difficult reaction between the lead and other perovskite precursors, leading to a slower reaction mechanism.

UV-Vis absorption and Photoluminescence (PL) spectroscopy characterise the optical properties of the samples. Figure 4-5 shows their spectra with Tauc plots and bandgap data in

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Supporting Figure 4-6. The OPA/OAm sample's bandgap is 2.44 eV, after which the lengthening of the PA generally shifts the absorbance onset towards lower energies to 2.37 eV for the ODPA/OAm sample. This shift correlates with the increase in particle size seen in

Figure 4-1.

Figure 4-5a) UV-Vis (dash line) and PL (full line) of sample 1-6.

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Figure 4-6 Supporting figure of Tauc plots and Trendline data showing Bandgap for PA+OAm

The slight redshift of the absorption onset as the PA's chain length increases directly contrasts with Pan et al., where a redshift occurs with the shortening of the carboxylic acid chain length.

[18] However, in their case, as the acid length increased, the NC size decreased, whereas here, the NC size increases with ligand length. Therefore, the redshift for both ligand types occurs as the NCs increase in size. Reducing particle size down to the exciton Bohr radius should lead to quantum size effects and a consequent blue shift of the absorption onset and fluorescence.[53] This effect is seen here with the shortening of the PA chain length and explains the impact of the ligand length on the optical properties of CsPbBr3. The

Photoluminescence (PL) properties of the samples show a less obvious progression. The peak shift is very slight but is still in agreement with the absorption data.

As shown by these results, PAs are a viable alternative to carboxylic acid ligands, allowing other size control methods. The amine ligand, which deprotonates to take the form of oleylammonium halides, is proposed as the primary surface-capping agents of CsPbX3 NCs

[16, 17]. Yang et al. [54] showed that NCs with branched ligands are more dispersible than

NCs with long straight chains such as OAm. The amine interaction with the precursors and surfactants in the synthesis is a significant component of the PbBr2 precursor system. Changing

4-115 from a carboxylic acid to a PA is likely to have consequences for all the precursors involved.

Although the amine concentration was consistently low in these reactions, OA removal may have caused effects that replicated the increase of amine that causes the formation of the lead depleted Cs4PbBr6 phase. In the work of Liu et al., they found that a pure OAm system, with a considerable increase (x1000) of amine relative to carboxylic acid, lead to a transformation of

CsPbBr3 to Cs4PbBr6. Yet, the addition of a thiol ligand markedly decreased the amount of amine required to drive the change. Interestingly, they found that the thiol ligand alone did not cause the transformation at all. [52]

Meanwhile, TOPO is a branched capping ligand that can provide a strong steric effect to the

NC. It has been widely used in the synthesis for group II-VI semiconductors and has many benefits. [55, 56] Its high boiling point allows reactions to proceed at high temperatures and is compatible with organic solvents, allowing an utterly inert reaction environment and air- sensitive precursors. TOPO can dissolve PbBr2 and so can be used here to replace OAm.[29]

Figure 5a-f shows the TEM images of the six samples made with TOPO and PAs. There are several differences immediately apparent. The particles are easier to wash than the OAm coated samples, leading to a more facile characterisation procedure. The 3:1 mix of anhydrous toluene and acetone was sufficient for both sets of samples; however, in the OAm case, the precipitate remained as a paste and was more difficult to characterise by XRD. While the samples were cleaner when using TOPO; and they aggregated in solution much more readily. The NCs are a similar size and shape to the OAm samples, yet the polydispersity increased, especially in

TOPO/DPA and TOPO/ODPA combinations. It is interesting to note that the DDPA and TDPA samples are again out of agreement, similar to the OAm combinations.

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Figure 4-7a-f) TEM images of samples made with PA 1-6 and TOPO. g) XRD patterns of the six samples with the reference pattern of orthorhombic CsPbBr3. h) Optical data of the six samples and i) Particle size distributions of the six samples.

Figure 4-8 Supporting figure of Tauc plots and Trendline data showing Bandgap for PA+TOPO The TOPO/PA system's true benefit is evident in the XRD patterns of Figure 5g. In all cases, the patterns are pure CsPbBr3 in the cubic phase without any evidence of Cs4PbBr6 impurities.

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It appears that the window for the complete dissolution of the precursors without the formation of lead depleted perovskite is more significant when using TOPO instead of OAm.

The UV-Vis absorption and PL data shown in Fig 5h show a slight redshift with increasing PA chain length. The bandgap decreases from 2.22 eV to 2.13 eV, considerably lower energies than the OAm case. Figure 4-8 displays the Tauc plots of the absorption data along with bandgap calculations. This movement corresponds well with the size distribution shown in

Figure 4-7i and can be taken together to show the size and quantum confinement effect of lengthening PA ligands. However, the size distribution shows that the particles lose their monodispersity to some extent, especially in the DPA/TOPO, DDPA/TOPO and ODPA/TOPO samples.

Interestingly, this system was discovered when the reaction time increased to 30 s, with double

Cs concentration. While there was little change in most samples, the reactions using

TOPO/DDPA and TOPO/HDPA yielded rods as shown in Figure 4-9a+c, respectively.

Figure 4-9 One dimensional nanomaterials synthesised using DDPA-TOPO and then HDPA-TOPO with the HRTEM and FFT of b) shown in c+d).

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Nanoparticles formed in the reaction using the surfactant combination DDPA/TOPO are shown in Figure 4-9a, displaying nanorods with aspect ratios between 3 and 5. The corresponding FFT shown in Figure 4-9b proves that the nanorods are rhombohedral Cs4PbBr6 (JCPDS: 01-073-

2478). A higher concentration of Cs relative to Pb in the reaction flask results in the synthesis of lead depleted Cs4PbBr6; however, this does not usually form 1D structures [29]. In Figure

4-9c+d, the HDPA-TOPO sample produced 1D nanoribbons with an aspect ratio of 5 and corresponded to the orthorhombic perovskite phase presented in the previous XRD patterns of

Figure 4-7g (JCPDS: 96-451-0746).

The vital difference here is that the batch of TOPO was different to the batch used throughout the rest of the study. In 2009, Wang et al. reported an in-depth study of TOPO impurities and their effects on 1D growth of CdSe nanowires. [57] Unfortunately, the batch of TOPO is no longer available to test using NMR or other techniques to understand what exact impurities it contained; however, we can test new batches.

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Figure 4-10 CsPbBr3 synthesised in different batches of TOPO

TOPO Lot # Purity Brand A MKCF0650 90% Sigma Aldrich B MKBP5115V 90% Sigma Aldrich C MKBV6061V 90% Sigma Aldrich D MKCF0650 99% Sigma Aldrich E 10217124 98% Alfa Aesar Original MKBP5119V 90% Sigma Aldrich

Table 4-1 TOPO batch details Figure 4-10 display TEM images of materials synthesised under identical conditions to that of the DDPA-TOPO sample displayed in Figure 4-9a but using different TOPO batches (Table

4-1). Unfortunately, there was no repetition of the 1D growth. However, the wide range of polydispersity and growth between the samples is significant. Specifically, the high purity

(98%) TOPO sample (Figure 4-10E) sourced from Alfa Aesar resulted in large misshapen material showing that the impurities are beneficial. An addition of purchasable phosphonic acid

4-120 impurities (1%) yielded similar anisotropic growth in the CdSe system. Namely,

Methylphosphonic acid [58], Octylphosphonic acid [57], Decylphosphonic acid [57],

Ethlyenediphosphonic acid [59], Diisooctylphosphinic acid [60] and Propylphosphonic acid.[61] 1% (10 mg) of impurity was included in the reaction flask with TOPO batch #1.

Figure 4-11 shows the resulting TEM images. Again, no 1D structures form. However, in most cases, polydispersity decreased. These images show that careful tuning of impurity species in the reaction system significantly improves the resultant dispersion.

Figure 4-11 TEM images of CsPbBr3 synthesised with 1% TOPO impurity 4.5 Conclusion

In this work, we present the viability of alkyl phosphonic acids as a replacement for oleic acid in the synthesis of CsPbBr3 NCs. When used as a ligand combined with oleylamine, phosphonic acids provide excellent control over the morphology of CsPbBr3 NCs, leading to

4-121 monodisperse nanoparticles with the option to increase their size by increasing the length of the phosphonic acid chain in the system. However, the system shows an increased propensity to form the lead depleted Cs4PbBr6 in conjunction with CsPbBr3, even with limited use of oleylamine. The subsequent replacement of oleylamine with trioctylphosphine oxide allowed for a pure monoclinic phase without any impurities. This purely phosphine based ligand system resulted in an overall lower bandgap of the perovskite material and varied the bandgap through phosphonic acid. As such, this work presents a systematic way in which to alter the optical properties of CsPbBr3.

A study across a wide range of TOPO batches of varying purity, from different and inclusion of numerous additives, was unable to pinpoint the exact impurity responsible for the 1D growth of Cs4PbBr6 nanorods seen in a specific Sigma Aldrich batch. Like CdSe, it is vital to discover which impurities are adventitious and use them in controlled amounts to allow rational and reproducible nanocrystal synthesis.

4.6 Bibliography

[1] M. A. Green, A. Ho-Baillie, and H. J. Snaith, "The emergence of perovskite solar cells," Nature Photonics, Review Article vol. 8, p. 506, 06/27/online 2014, doi: 10.1038/nphoton.2014.134. [2] Z.-K. Tan et al., "Bright light-emitting diodes based on organometal halide perovskite," Nature Nanotechnology, vol. 9, no. 9, pp. 687-692, 2014/09/01 2014, doi: 10.1038/nnano.2014.149. [3] G. Li, J. Huang, H. Zhu, Y. Li, J.-X. Tang, and Y. Jiang, "Surface Ligand Engineering for Near-Unity Quantum Yield Inorganic Halide Perovskite QDs and High-Performance QLEDs," Chemistry of Materials, vol. 30, no. 17, pp. 6099-6107, 2018/09/11 2018, doi: 10.1021/acs.chemmater.8b02544. [4] Y. Tan et al., "Highly Luminescent and Stable Perovskite Nanocrystals with Octylphosphonic Acid as a Ligand for Efficient Light-Emitting Diodes," ACS Applied Materials & Interfaces, vol. 10, no. 4, pp. 3784-3792, 2018/01/31 2018, doi: 10.1021/acsami.7b17166. [5] F. Wang, S. Bai, W. Tress, A. Hagfeldt, and F. Gao, "Defects engineering for high-performance perovskite solar cells," npj Flexible Electronics, vol. 2, no. 1, p. 22, 2018/08/14 2018, doi: 10.1038/s41528-018-0035-z. [6] S. Aharon and L. Etgar, "Two Dimensional Organometal Halide Perovskite Nanorods with Tunable Optical Properties," Nano Letters, vol. 16, no. 5, pp. 3230-3235, 2016/05/11 2016, doi: 10.1021/acs.nanolett.6b00665.

4-122

[7] Q. A. Akkerman et al., "Tuning the Optical Properties of Cesium Lead Halide Perovskite Nanocrystals by Anion Exchange Reactions," Journal of the American Chemical Society, vol. 137, no. 32, pp. 10276-10281, 2015/08/19 2015, doi: 10.1021/jacs.5b05602. [8] Q. A. Akkerman et al., "Solution Synthesis Approach to Colloidal Cesium Lead Halide Perovskite Nanoplatelets with Monolayer-Level Thickness Control," Journal of the American Chemical Society, vol. 138, no. 3, pp. 1010-1016, 2016/01/27 2016, doi: 10.1021/jacs.5b12124. [9] X. Chen, L. Peng, K. Huang, Z. Shi, R. Xie, and W. Yang, "Non-injection gram-scale synthesis of cesium lead halide perovskite quantum dots with controllable size and composition," Nano Research, journal article vol. 9, no. 7, pp. 1994-2006, July 01 2016, doi: 10.1007/s12274-016- 1090-1. [10] H. Deng et al., "Growth, patterning and alignment of organolead iodide perovskite nanowires for optoelectronic devices," Nanoscale, 10.1039/C4NR06982J vol. 7, no. 9, pp. 4163-4170, 2015, doi: 10.1039/C4NR06982J. [11] T. Paul, B. K. Chatterjee, N. Besra, S. Thakur, S. Sarkar, and K. K. Chattopadhyay, "Fabrication of all-inorganic CsPbBr3 perovskite nanocubes for enhanced green photoluminescence," Materials Today: Proceedings, vol. 5, no. 1, Part 2, pp. 2234-2240, 2018/01/01/ 2018, doi: https://doi.org/10.1016/j.matpr.2017.09.224. [12] L. Protesescu et al., "Nanocrystals of Cesium Lead Halide Perovskites (CsPbX3, X = Cl, Br, and I): Novel Optoelectronic Materials Showing Bright Emission with Wide Color Gamut," Nano Letters, vol. 15, no. 6, pp. 3692-3696, 2015/06/10 2015, doi: 10.1021/nl5048779. [13] D. Yang, X. Li, and H. Zeng, "Surface Chemistry of All Inorganic Halide Perovskite Nanocrystals: Passivation Mechanism and Stability," Advanced Materials Interfaces, vol. 5, no. 8, p. 1701662, 2018, doi: 10.1002/admi.201701662. [14] R. Grisorio et al., "Exploring the surface chemistry of cesium lead halide perovskite nanocrystals," Nanoscale, 10.1039/C8NR08011A vol. 11, no. 3, pp. 986-999, 2019, doi: 10.1039/C8NR08011A. [15] J. T. Kopping and T. E. Patten, "Identification of Acidic Phosphorus-Containing Ligands Involved in the Surface Chemistry of CdSe Nanoparticles Prepared in Tri-N-octylphosphine Oxide Solvents," Journal of the American Chemical Society, vol. 130, no. 17, pp. 5689-5698, 2008/04/01 2008, doi: 10.1021/ja077414d. [16] C. Wang, A. S. R. Chesman, and J. J. Jasieniak, "Stabilizing the cubic perovskite phase of CsPbI3 nanocrystals by using an alkyl phosphinic acid," Chemical Communications, 10.1039/C6CC08282C vol. 53, no. 1, pp. 232-235, 2017, doi: 10.1039/C6CC08282C. [17] J. De Roo et al., "Highly Dynamic Ligand Binding and Light Absorption Coefficient of Cesium Lead Bromide Perovskite Nanocrystals," ACS Nano, vol. 10, no. 2, pp. 2071-2081, 2016/02/23 2016, doi: 10.1021/acsnano.5b06295. [18] A. Pan et al., "Insight into the Ligand-Mediated Synthesis of Colloidal CsPbBr3 Perovskite Nanocrystals: The Role of Organic Acid, Base, and Cesium Precursors," ACS Nano, vol. 10, no. 8, pp. 7943-7954, 2016/08/23 2016, doi: 10.1021/acsnano.6b03863. [19] F. Palazon et al., "Postsynthesis Transformation of Insulating Cs4PbBr6 Nanocrystals into Bright Perovskite CsPbBr3 through Physical and Chemical Extraction of CsBr," ACS Energy Letters, vol. 2, no. 10, pp. 2445-2448, 2017/10/13 2017, doi: 10.1021/acsenergylett.7b00842.

4-123

[20] V. K. Ravi et al., "Origin of the Substitution Mechanism for the Binding of Organic Ligands on the Surface of CsPbBr3 Perovskite Nanocubes," The Journal of Physical Chemistry Letters, vol. 8, no. 20, pp. 4988-4994, 2017/10/19 2017, doi: 10.1021/acs.jpclett.7b02192. [21] J. Y. Woo et al., "Air-Stable PbSe Nanocrystals Passivated by Phosphonic Acids," Journal of the American Chemical Society, vol. 138, no. 3, pp. 876-883, 2016/01/27 2016, doi: 10.1021/jacs.5b10273. [22] X. Li et al., "Improved performance and stability of perovskite solar cells by crystal crosslinking with alkylphosphonic acid ω-ammonium chlorides," Nature Chemistry, Article vol. 7, p. 703, 08/17/online 2015, doi: 10.1038/nchem.2324 https://www.nature.com/articles/nchem.2324#supplementary-information. [23] R. Gomes et al., "Binding of Phosphonic Acids to CdSe Quantum Dots: A Solution NMR Study," The Journal of Physical Chemistry Letters, vol. 2, no. 3, pp. 145-152, 2011/02/03 2011, doi: 10.1021/jz1016729. [24] W.-k. Koh, S. Park, and Y. Ham, "Phosphonic Acid Stabilized Colloidal CsPbX3 (X=Br, I) Perovskite Nanocrystals and Their Surface Chemistry," ChemistrySelect, vol. 1, no. 13, pp. 3479-3482, 2016, doi: doi:10.1002/slct.201600809. [25] T. Xuan et al., "Highly stable CsPbBr3 quantum dots coated with alkyl phosphate for white light-emitting diodes," Nanoscale, 10.1039/C7NR04179A vol. 9, no. 40, pp. 15286-15290, 2017, doi: 10.1039/C7NR04179A. [26] C. Lu et al., "Enhanced stabilization of inorganic cesium lead triiodide (CsPbI3) perovskite quantum dots with tri-octylphosphine," Nano Research, vol. 11, no. 2, pp. 762-768, 2018/02/01 2018, doi: 10.1007/s12274-017-1685-1. [27] A. A. M. Brown et al., "Self-assembly of a robust hydrogen-bonded octylphosphonate network on cesium lead bromide perovskite nanocrystals for light-emitting diodes," Nanoscale, 10.1039/C9NR02566A vol. 11, no. 25, pp. 12370-12380, 2019, doi: 10.1039/C9NR02566A. [28] L. Wu et al., "Improving the Stability and Size Tunability of Cesium Lead Halide Perovskite Nanocrystals Using Trioctylphosphine Oxide as the Capping Ligand," Langmuir, vol. 33, no. 44, pp. 12689-12696, 2017/11/07 2017, doi: 10.1021/acs.langmuir.7b02963. [29] G. Almeida et al., "The Phosphine Oxide Route toward Lead Halide Perovskite Nanocrystals," Journal of the American Chemical Society, vol. 140, no. 44, pp. 14878-14886, 2018/11/07 2018, doi: 10.1021/jacs.8b08978. [30] Y. Yao et al., "Efficient Quantum Dot Light-Emitting Diodes Based on Trioctylphosphine Oxide- Passivated Organometallic Halide Perovskites," ACS Omega, vol. 4, no. 5, pp. 9150-9159, 2019/05/31 2019, doi: 10.1021/acsomega.9b00464. [31] X. Yang et al., "Efficient green light-emitting diodes based on quasi-two-dimensional composition and phase engineered perovskite with surface passivation," Nature Communications, vol. 9, no. 1, p. 570, 2018/02/08 2018, doi: 10.1038/s41467-018-02978-7. [32] E. Yassitepe et al., "Amine-Free Synthesis of Cesium Lead Halide Perovskite Quantum Dots for Efficient Light-Emitting Diodes," Advanced Functional Materials, vol. 26, no. 47, pp. 8757- 8763, 2016, doi: 10.1002/adfm.201604580. [33] P. Liu et al., "Assembling Ordered Nanorod Superstructures and Their Application as Microcavity Lasers," Scientific Reports, Article vol. 7, p. 43884, 03/08/online 2017, doi: 10.1038/srep43884

4-124 https://www.nature.com/articles/srep43884#supplementary-information. [34] J.-J. Wang and K. M. Ryan, "Colloidal synthesis of Cu 2 SnSe 3 nanocrystals with structure induced shape evolution," CrystEngComm, vol. 18, no. 18, pp. 3161-3169, 2016. [35] P. Liu, S. Singh, G. Bree, and K. M. Ryan, "Complete assembly of Cu₂ZnSnS₄ (CZTS) nanorods at substrate interfaces using a combination of self and directed organisation," 2016. [36] C. Coughlan, M. Ibanez, O. Dobrozhan, A. Singh, A. Cabot, and K. M. Ryan, "Compound copper chalcogenide nanocrystals," Chemical reviews, vol. 117, no. 9, pp. 5865-6109, 2017. [37] S. Singh, M. Brandon, P. Liu, F. Laffir, W. Redington, and K. M. Ryan, "Selective Phase Transformation of Wurtzite Cu2ZnSn (SSe) 4 (CZTSSe) Nanocrystals into Zinc-Blende and Kesterite Phases by Solution and Solid State Transformations," Chemistry of Materials, vol. 28, no. 14, pp. 5055-5062, 2016. [38] U. V. Ghorpade et al., "Colloidal wurtzite Cu2SnS3 (CTS) nanocrystals and their applications in solar cells," Chemistry of Materials, vol. 28, no. 10, pp. 3308-3317, 2016. [39] G. Bree, H. Geaney, K. Stokes, and K. M. Ryan, "Aligned copper zinc tin sulfide nanorods as lithium-ion battery anodes with high specific capacities," The Journal of Physical Chemistry C, vol. 122, no. 35, pp. 20090-20098, 2018. [40] G. Bree, H. Geaney, and K. M. Ryan, "Electrophoretic Deposition of Tin Sulfide Nanocubes as High‐Performance Lithium‐Ion Battery Anodes," ChemElectroChem, vol. 6, no. 12, pp. 3049-3056, 2019. [41] G. Bree, C. Coughlan, H. Geaney, and K. M. Ryan, "Investigation into the Selenization Mechanisms of Wurtzite CZTS Nanorods," ACS applied materials & interfaces, vol. 10, no. 8, pp. 7117-7125, 2018. [42] C. Coughlan, Y. Guo, S. Singh, S. Nakahara, and K. M. Ryan, "Synthesis of Curved CuIn1–x Ga x (S1–y Se y) 2 Nanocrystals and Complete Characterization of Their Diffraction Contrast Effects," Chemistry of Materials, vol. 30, no. 23, pp. 8679-8689, 2018. [43] S. Singh, A. Singh, K. Palaniappan, and K. M. Ryan, "Colloidal synthesis of homogeneously alloyed CdSexS1−x nanorods with compositionally tunable photoluminescence," Chemical Communications, 10.1039/C3CC45497E vol. 49, no. 87, pp. 10293-10295, 2013, doi: 10.1039/C3CC45497E. [44] A. Singh, H. Geaney, F. Laffir, and K. M. Ryan, "Colloidal Synthesis of Wurtzite Cu2ZnSnS4 Nanorods and Their Perpendicular Assembly," Journal of the American Chemical Society, vol. 134, no. 6, pp. 2910-2913, 2012/02/15 2012, doi: 10.1021/ja2112146. [45] A. Singh et al., "Controlled semiconductor nanorod assembly from solution: influence of concentration, charge and solvent nature," Journal of Materials Chemistry, 10.1039/C1JM14382D vol. 22, no. 4, pp. 1562-1569, 2012, doi: 10.1039/C1JM14382D. [46] A. Singh, R. D. Gunning, A. Sanyal, and K. M. Ryan, "Directing semiconductor nanorod assembly into 1D or 2D supercrystals by altering the surface charge," Chemical Communications, 10.1039/C0CC01455A vol. 46, no. 38, pp. 7193-7195, 2010, doi: 10.1039/C0CC01455A. [47] U. V. Ghorpade et al., "Wurtzite CZTS nanocrystals and phase evolution to kesterite thin film for solar energy harvesting," Physical Chemistry Chemical Physics, vol. 17, no. 30, pp. 19777- 19788, 2015. [48] W. Wang, S. Banerjee, S. Jia, M. L. Steigerwald, and I. P. Herman, "Ligand Control of Growth, Morphology, and Capping Structure of Colloidal CdSe Nanorods," Chemistry of Materials, vol. 19, no. 10, pp. 2573-2580, 2007/05/01 2007, doi: 10.1021/cm0705791.

4-125

[49] B. Zhang et al., "Alkyl Phosphonic Acids Deliver CsPbBr3 Nanocrystals with High Photoluminescence Quantum Yield and Truncated Octahedron Shape," Chemistry of Materials, 2019/10/14 2019, doi: 10.1021/acs.chemmater.9b03529. [50] O. F. Mohammed, "Outstanding Challenges of Zero-Dimensional Perovskite Materials," The Journal of Physical Chemistry Letters, vol. 10, no. 19, pp. 5886-5888, 2019/10/03 2019, doi: 10.1021/acs.jpclett.9b00175. [51] Y.-M. Chen et al., "Cs4PbBr6/CsPbBr3 Perovskite Composites with Near-Unity Luminescence Quantum Yield: Large-Scale Synthesis, Luminescence and Formation Mechanism, and White Light-Emitting Diode Application," ACS Applied Materials & Interfaces, vol. 10, no. 18, pp. 15905-15912, 2018/05/09 2018, doi: 10.1021/acsami.8b04556. [52] Z. Liu et al., "Ligand Mediated Transformation of Cesium Lead Bromide Perovskite Nanocrystals to Lead Depleted Cs4PbBr6 Nanocrystals," Journal of the American Chemical Society, vol. 139, no. 15, pp. 5309-5312, 2017/04/19 2017, doi: 10.1021/jacs.7b01409. [53] J. A. Sichert et al., "Quantum Size Effect in Organometal Halide Perovskite Nanoplatelets," Nano Letters, vol. 15, no. 10, pp. 6521-6527, 2015/10/14 2015, doi: 10.1021/acs.nanolett.5b02985. [54] Y. Yang et al., "Entropic Ligands for Nanocrystals: From Unexpected Solution Properties to Outstanding Processability," Nano Letters, vol. 16, no. 4, pp. 2133-2138, 2016/04/13 2016, doi: 10.1021/acs.nanolett.6b00730. [55] D. V. Talapin, A. L. Rogach, A. Kornowski, M. Haase, and H. Weller, "Highly luminescent monodisperse CdSe and CdSe/ZnS nanocrystals synthesized in a hexadecylamine− trioctylphosphine oxide− trioctylphospine mixture," Nano letters, vol. 1, no. 4, pp. 207-211, 2001. [56] M. Green, "The nature of quantum dot capping ligands," Journal of Materials Chemistry, vol. 20, no. 28, pp. 5797-5809, 2010. [57] F. Wang, R. Tang, J. L. F. Kao, S. D. Dingman, and W. E. Buhro, "Spectroscopic Identification of Tri-n-octylphosphine Oxide (TOPO) Impurities and Elucidation of Their Roles in Cadmium Selenide Quantum-Wire Growth," Journal of the American Chemical Society, vol. 131, no. 13, pp. 4983-4994, 2009/04/08 2009, doi: 10.1021/ja900191n. [58] Y.-w. Jun, S.-M. Lee, N.-J. Kang, and J. Cheon, "Controlled Synthesis of Multi-armed CdS Nanorod Architectures Using Monosurfactant System," Journal of the American Chemical Society, vol. 123, no. 21, pp. 5150-5151, 2001/05/01 2001, doi: 10.1021/ja0157595. [59] A. G. Kanaras, C. Sönnichsen, H. Liu, and A. P. Alivisatos, "Controlled Synthesis of Hyperbranched Inorganic Nanocrystals with Rich Three-Dimensional Structures," Nano Letters, vol. 5, no. 11, pp. 2164-2167, 2005/11/01 2005, doi: 10.1021/nl0518728. [60] A. D. Dukes, J. R. McBride, and S. J. Rosenthal, "Synthesis of Magic-Sized CdSe and CdTe Nanocrystals with Diisooctylphosphinic Acid," Chemistry of Materials, vol. 22, no. 23, pp. 6402-6408, 2010/12/14 2010, doi: 10.1021/cm102370a. [61] J. Huang, M. V. Kovalenko, and D. V. Talapin, "Alkyl Chains of Surface Ligands Affect Polytypism of CdSe Nanocrystals and Play an Important Role in the Synthesis of Anisotropic Nanoheterostructures," Journal of the American Chemical Society, vol. 132, no. 45, pp. 15866- 15868, 2010/11/17 2010, doi: 10.1021/ja105132u.

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Chapter 5 Colloidal Synthesis of 2D Monolayer Indium

Chalcogenide Alloys Nanosheets

5.1 Abstract

Nanosheets of indium chalcogenides are ideal for next-generation optoelectronic devices due to their outstanding photophysical properties. While mechanical exfoliation is the most widespread method of formation, colloidal wet-chemistry methods are particularly promising.

The as-synthesized colloidal dispersions are directly applicable in solution-based processing, opening a gateway for a straightforward and cost-efficient introduction into various technology platforms. This chapter presents a hot injection synthesis of alloyed indium chalcogenide nanoribbons. They are characterised both structurally and optically. A substitution limit of 25%

S substitution was found for In2Se3 and 50% Se into In2S3, after which a phase change occurs.

The In2(SeS)3 wires are highly anisotropic. Substitution of tellurium into the reactions lead to the formation of an InTe-Se alloy.

5.2 Introduction

A tremendous development of layered 2D nanomaterials has occurred since the discovery of graphene in 2004. For example, indium chalcogenide nanosheets have application in electronics, biomedical engineering and electrocatalysis.[1-3] In a single layer of 2D material, almost every atom is exposed, meaning that the surface to volume ratio of 2D materials is significantly increased versus other forms of nanomaterials.[3] The high surface-to-volume ratio of these 2D nanomaterials exposes abundant unsaturated surface atoms, providing many active sites for electrochemical and photochemical reactions. [1-3] Moreover, this significantly increases the lateral conduction and the physical and chemical reactivity, influencing the 2D wave function through quantum confinement effects.[4-6] 2D materials are processed into thin- 5-127 film devices through printing techniques such as ink-jet printing and screen printing.[7]

Mechanical exfoliation is a widespread method for the layer separation of 2D materials; however, the emergence of solution processing has added value for their economic viability, offering scalable manufacturing for future industries. [8, 9]

Many colloidal 2D semiconductor nanomaterial syntheses show morphological control in many compositions, including CdSe and PbS. Their properties are desirable in applications such as wearable electronics, and so they cannot contain these toxic elements. [10-14] Indium chalcogenides are non-toxic, Van der Walls layered III-VI structures. They have atomically smooth lateral surfaces and are thermally stable up to 660 °C.[13] InSe has promising optoelectronic properties, high mobility in FETs (1x103 cm2/(V s) ), and outperform materials such as MoS2, GaS GaSe in photodetectors. [15-20] Indium sulphide, In2S3, has high photosensitivity and photoconductivity and long term stability. [14, 21, 22] In2S3 has applications ranging from LED displays, photocatalysis, and photodetectors to lithium storage cells. [23-25] Meanwhile, In2Se3 has a narrow bandgap with attractive properties for photovoltaics, thermoelectrics, photocatalysis, ionic batteries, optoelectronics, and phase- change memory devices. [26-31] There are few reports of telluride based indium chalcogenides in colloidal synthesis. [32, 33] They have potential in nonlinear optics, photocatalysis, and optoelectric devices. [34-38]

Varying the distribution of different chalcogenides enables the manipulation of the material bandgap.[39] For example, selenium incorporation into CZTS solar cells modulates the bandgap from 1.54 to 1.47 eV. In comparison, selenium inclusion into CIGS lowers the bandgap from ~1.5 to 1.05 eV and increases a solar cell's efficiency from ~1% to 7.2%.[40,

41] This change is attributed to Se incorporation improving grain growth and passivating the grain surfaces. [40]

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In this work, indium chalcogenide (InE, E = Se, S, Te) alloys are synthesised by hot injection.

Chalcogenide anion alloying manipulates the bandgap of indium chalcogenide materials. Se incorporates into In2S3 up to 25%, and S incorporates into In2Se3 but not into InSe. The incorporation of Te into the InSe and InS materials was then further studied. XRD shows that sulphur and tellurium do not alloy in this system.

5.3 Experimental Section

Materials Indium (III) Chloride (InCl3, 99%), Selenium powder (99.99%), Sulphur powder

(99.98%) Tellurium power (99.99%), Sulphur (>99%), Oleylamine (OLA, 70%),

Trioctylphosphine (TOP, 90%), Dimethyl carbonate (DMC, 99%), and Dichlorobenzene

(DCB, 99%) were purchased from Sigma Aldrich. Toluene, Methanol, IPA and Acetone were purchased from Lennox Chemicals. All except OLA were used without further purification.

Synthesis of indium precursor solution: 1 mmol (221 mg) InCl3 was dissolved in 5 ml oleylamine and heated to 200 °C to form a solution.

Synthesis of indium chalcogenide: Mixtures of chalcogenide powder of 1 mmol was added to a three-neck round bottom flask in an argon-filled glove box and then connected to a Schlenk line and equipped with a temperature finger and a septum cap. In this reaction system, oleylamine (OLA) is both solvent and surfactant with elemental selenium, tellurium and sulphur powders as the chalcogenide sources. Work by Park et al. showed that selenium needs to be reacted at a temperature of 250 °C in OLA to dissolve entirely before reacting with the indium chloride precursor.[42] Here, a temperature of 300 °C was needed for 1h to dissolve the tellurium precursor. The appropriate ratio of chalcogenide powders to maintain 1 mmol is heated in 5 ml of dried OLA at the proper temperature and time to dissolve the precursors before cooling to 110 °C. A solution of 1 mmol InCl3:OLA is injected, and the system is

5-129 returned to 215 °C (Se + S) or 300 °C (Te) for the desired amount of time. The colour change to black indicated the formation of InE particles.

Table 5-1 Table of chalcogenide concentrations. Selenium/Sulphur is highlighted in green, tellurium/sulphur is highlighted in red and tellurium/selenium is highlighted in blue.

Ratio Selenium Thiourea Ratio Tellurium Sulphur

In+Se1-xSx x=0 36 mg 0 mg In+Te1-xSx x=0 130 mg 0 mg

In+Se1-xSx x=0.125 31.5 mg 9.5 mg In+Te1-xSx x=0.25 97.5 mg 8 mg

In+Se1-xSx x=0.25 27 mg 19 mg In+Te1-xSx x=0.5 65 mg 16 mg

In+Se1-xSx x=0.375 22.5 mg 28.5 mg In+Te1-xSx x=0.75 32.5 mg 24 mg

In+Se1-xSx x=0.5 18 mg 38 mg In+Te1-xSx x=1 0 mg 32 mg

In+Se1-xSx x=0.625 13.5 mg 47.5 mg Ratio Tellurium Selenium

In+Se1-xSx x=0.75 9 mg 57 mg In+Te1-xSex x=0 130 mg 0 mg

In+Se1-xSx x=0.875 4.5 mg 66.5 mg In+Te1-xSex x=0.25 97.5 mg 9 mg

In+Se1-xSx x=1 0 mg 76 mg In+Te1-xSex x=0.5 65 mg 18 mg

In+Te1-xSex x=0.75 32.5 mg 27 mg

In+Te1-xSex x=1 0 mg 36 mg

Purification of InE Nanosheets: Following the allotted reaction time, the heat was removed, and the system cooled to 50 °C. 5 ml of toluene was added directly to the nascent, followed by

5ml isopropanol (IPA). The product was centrifuged at 5,000 rpm for 5 minutes. The product was washed with dimethyl carbonate, IPA and toluene in the ratio three times until the particles disperse well in toluene. Further purification steps were required to achieve clear XRD patterns.

Transmission Electron Microscopy (TEM): Bright- and dark-field TEM images, high- resolution TEM (HRTEM) and selected area electron diffraction (SAED) patterns were acquired on samples prepared by drop-casting colloidal solutions in toluene onto carbon-coated

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200 mesh copper grids. The machine used was a JEOL TEM-2100 microscope (W filament) operating at an accelerating voltage of 200 kV.

X-ray Diffraction (XRD): The XRD patterns of dried powders were acquired on a zero- diffraction silicon substrate in a PANalytical Empyrean X-ray diffractometer equipped with a

1.8 kW Cu Kαceramic X-ray tube, operating at 40 kV and 40 mA.

Steady-State UV−Vis Extinction Spectroscopy: Optical extinction spectra of anhydrous toluene dispersions were recorded in quartz cuvettes with a 1 cm path length, employing an

Agilent Cary 5000 UV−vis spectrophotometer. Samples were stored in Argon.

5.4 Results and discussion

Figure 5-1 shows the TEM images of the InE (E = Se, S). The first material shown in Figure

5-1a is a pure InSe. Oleylamine is used as a ligand here which also acts as a reducing agent and reduces the In3+ to In2+ and the elemental Se to Se2-, forming indium monoselenide.[43]

Measuring the InSe nanosheets' side lengths is complicated by the screw dislocation of some particles, which causes growth along the c axis, as shown most transparently by the nanosheet at the top of Figure 5-1a. However, long to short sides' ratios were generally between 2.3 and

2.8, with a few almost triangular particles, consisting of three sides >5x longer than the short sides. No impurities, such as In2O3, InS, or In(OH)3, are detected.

Table 5-1 shows experimental details of sulphur concentration increases of 12.5% increments.

The TEM images in Figure 5-1b shows that alloying in In+Se1-xSx x=0.125 formed highly anisotropic nanoribbons (length: >10 µm, width: 6.5 ± 0.7 nm). There are also a small number of hexagonal plates throughout the sample in Figure 5-1b. These plates disappear as the amount of sulphur precursor is increased to 25%, In+Se1-xSx x=0.25, and the nanoribbons make up all of the sample (length: >10 µm, width: 7.0 ± 0.8 nm) seen clearer in the inset of Figure 5-1c. Of 5-131 the eight rings of the SAED pattern for In+Se1-xSx x=0.125: (0.4756, 0.378, 0.2419, 0.226,

0.2016, 0.1307, 0.1194 and 0.1038 nm), five correspond to ICDD pattern 00-040-0911:

In2S0.6Se2.4, while four correspond to the hexagonal InSe ICDD: 00-034-1431 with one peak overlapping. As such, one can extrapolate from the first image that the hexagonal plates are

InSe while the long nanoribbons are an alloyed In2S0.6Se2.4 material. This pattern (ICDD: 00-

035-1056) contains the peaks of rhombohedral In2Se3 but shifted to higher angles. Therefore the material is identified as rhombohedral In2Se3 with sulphur alloying. It is therefore improbable that any sulphur atoms infiltrated into the hexagonal indium monoselenide nanoplates. In Figure 5-1c, the In+Se1-xSx x=0.25 sample does not contain any InSe peaks in the SAED nor XRD pattern of Figure 5-2, while the existing peaks correspond to the

In2S0.6Se2.4 alloy.

In In+Se1-xSx x=0.375, the widths of the nanoribbons increase in size Figure 5-1e (w: 20 ± 5.6 nm) and the cubic phase of In2S3 becomes prominent. The SAED pattern corroborates the increase by identifying the cubic In2S3 phase. Therefore the narrow and long nanowires in

Figure 5-1b+c can be identified as alloyed In2S0.6Se2.4 with a phase change occurring. The material in Figure 5-1d-g a cubic phase of the alloy. The materials formed when the sulphur concentration is 87.5%, and 100% are hexagonal nanosheets with nanoribbons occurring as very thin wires in the InSe1-xSx x = 0.875 sample. Interestingly, the SAED still matches with the ICDD of cubic beta In2S3.

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Figure 5-1 TEM images with respective SAED insets of the array of Indium chalcogenide materials synthesised with increasing amounts of sulphur anion relative to selenium. All scale bars 500 nm except inset of c), which is 50 nm.

The TEM image of the In+Se1-xSx x=0.5 sample, Figure 5-1e, shows a material with a very low polydispersity (length ~ 2 µm, w: 44 ± 2.7 nm). Further increase of sulphur to In+Se1-xSx x=0.875 results in very long and narrow nanoribbons that measure multiple micrometres in length but with a width of 4 nm ± 0.9 nm. Nanoplates also occur in In+Se1-xSx x=0.875 and

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In+Se1-xSx x=1. Nanoplates are present in the pure indium sulphide sample, suggesting that the nanosheets are pure β-In2S3, with some nanosheets elongating into nanoribbons. This suggests that elongation occurs following nanoplatelets formation rather as a result of alloying in the

In2S3 system.

Figure 5-1 indicates that the alloys form in the 1D plane rather than as 2D hexagons. The width of the 1D nanoribbon is controllable with the anion content, with the In2S3 alloys being much wider than the In2Se2.4S0.6 material.

Figure 5-2 shows the XRD patterns of the resultant InE (E = S, Se) nanomaterials which correspond to the TEM images shown in Figure 5-1. The sample InSe is a pure indium monoselenide material with a hexagonal structure and space group P63/mmc, #194. InSe is identified by both the XRD pattern (ICDD: 00-034-1431) and agrees with the SAED pattern inset of Figure 5-1a.

Figure 5-2 XRD pattern of the alloyed materials ranging from pure indium selenide to pure indium sulphide. The coloured * indicates the phases, with red indicating In2S3 and blue indicating In2S0.6Se2.4.

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The sample named In+Se1-xSx X=0.25 corresponds to the In2Se2.4S0.6 material with the reference pattern code ICDD: 00-040-0911. This pattern contains all of the peaks of rhombohedral In2Se3 ICDD:00-035-1056 but with peak shifts to higher angles. Therefore, the unidentified phase-type of In2Se2.4S0.6 is a rhombohedral In2Se3 with S alloying. Once the selenium ratio to sulphur is 50:50, i.e. In+Se1-xSx x=0.5, the XRD pattern shows peaks of

In2S0.6Se2.4 and In2S3. The In2S0.6Se2.4 peaks do not shift their position indicating that further substitution of sulphur into the In2S0.6Se2.4 material does not occur. The SAED pattern inset into Figure 5-1e shows rings accounted for by the In2S3 pattern. The XRD peaks of In2S3 are shifted to a higher angle consecutively from In+Se1-xSx x=0.5 to In+Se1-xSx x=1, showing that selenium does alloy into the In2S3 material. The remaining sample substitutions correspond to

In2S3 with selenium substitution decreasing, indicated by the shifting of the XRD peak positions until they correspond to the cubic β-In2S3, space group Fd-3m, #227. The XRD shows that sulphur and selenium alloy in the indium chalcogenide materials but undergo phase changes. Even with selenium alloying, the cubic In2S3 phase occurs when sulphur content is

50% and above. However, when there are equal parts sulphur and selenium, two phases occur,

In2S3 and In2S0.6Se2.4, with alloying in both phases.

Figure 5-3 shows the absorption properties of the InE (E = Se – S) samples. The inset shows the Tauc plot used to find the bandgap by comparing the amount of light absorption at each wavelength to the incident light's energy. Figure 5-3 shows this data. From this, the direct bandgap of InSe is 1.6 eV. The bulk bandgap of this material is Eg ~ 1.3 eV which demonstrates quantum confinement effects that are causing blue shifting of the band edges.[43]

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Figure 5-3 Normalised absorption data for the Indium chalcogenide alloys to In+Se1-xSx x=0 to 1, InSe to In2S3; with Tauc plot inset. Bandgaps are in the attached legend. The break at 800 nm is due to detector changeover.

The immediate conversion from InSe into the In2Se3 phase when sulphur is introduced is apparent in the swift increase of the absorption onset from 1.48 eV to 2.17 eV. Here the bandgap increases slightly from 2.17 eV to 2.28 eV as sulphur concentration increases. This shift is more significant than the expected bulk bandgap of In2Se3, consistent with an alloying effect from sulphur incorporation and agrees with the TEM and XRD data findings. The second phase change into cubic In2S3 when sulphur content reaches 62.5% is shown by the significant increase from 2.28 eV to 2.61 eV.

The patterns from In+Se1-xSx x=0.5 to x=1 display a bandgap ranging from 2.61 to 3.09 eV massively increased from the expected value of 2.3 eV for bulk In2S3.[44] These samples all

5-136 display the peaks of β-In2S3 in the XRD patterns of Figure 5-2a, although shifted to lower angles with increasing amounts of selenium, as shown by the red star in Figure 5-2b. The bandgaps consistency agrees with the findings of XRD where these five samples, In+Se1-xSx x=0.625 to x=1, show peaks of In2S3 with Se alloying. This shift shows that Se incorporation into In2S3 decreases the optical bandgap from 2.61 eV to 3.09 eV.

As 25% of each chalcogenide alloys into the materials, 25% tellurium was introduced into the

InSe and InS systems. Figure 5-4 shows the resulting TEM images. For the dissolution of tellurium in OLA, the temperature required is increased to 300°C and maintained for 1h. This temperature requires more energy than selenium dissolution in OLA, which dissolves at 250

°C in 30 minutes.

Figure 5-4a+b presents the materials with tellurium and selenium combined. Figure 5-4a shows the In+Te1-xSex x=0.25 sample, which contains largely anisotropic structures and also smaller, spherical particles agglomerated together. The SAED pattern shows orthorhombic In4Te3

(ICDD: 00-029-0679). In In+Te1-xSex x=0.75, in Figure 5-4b, the polydisperse nanocrystals are shown as hexagonal In2Se3 (ICDD: 00-012-0117) through the inset SAED pattern. The inset of a different area on the TEM grid shows nanorods structures and the In2Se3 nanocrystals.

Figure 5-4c+d are the TEM images of the tellurium and sulphur samples. In In+Te1-xSx x=0.25,

Figure 5-4c, the sample is made up of nanoplatelets with undefined edges due to agglomeration on the flat plane. The SAED identifies only the tetragonal InTe phase (ICDD: 00-007-0112).

Interestingly, when the reaction time increases to 24 h, nanoribbons occur, as seen by Figure

5-4e. The FFT of a closeup nanoribbon matches the crystallographic pattern of cubic β-In2S3,

ICDD: 00-032-0456. These 1D nanoforms do not appear in any other sample after 3 or 24h.

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Figure 5-4 TEM images of the a-b) Indium telluride-selenide materials c-d) indium telluride-sulfide materials, e) 24h reaction time of In+Te0.75S0.25 material, and finally, f) InTe material. Insets are the SAED patterns for the materials except for e), an FFT of a single nanowire.

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When the concentration of tellurium is reduced further, in Figure 5-4d, In+Te1-xSx x=0.75, nanoplatelets continue to be the morphology of the formed material. The SAED pattern interestingly only identifies peaks of InTe, with no InS nor elemental sulphur peaks.

Figure 5-4f, InTe is synthesised using only the tellurium chalcogenide. The image shows a nanosheet with edge lengths of between 2-3.5 µm. The SAED pattern inset identifies a single crystal of tetragonal InTe, ICDD: 00-007-0112. This tetragonal InTe phase is identical to the

InTe of Figure 5-4d; however, the single crystal morphology is very different. This pattern suggests that incorporating sulphur into the reaction system changes the growth properties and preferred orientation.

Figure 5-5 XRD pattern of the alloyed indium telluride and selenide materials ranging from pure indium telluride to pure indium selenide. The purple * identifies elemental Te, blue * identifies In4Te3, black identifies In2Se3 and red * identifies InSe. Figure 5-5 presents the XRD patterns of selenium-tellurium material alloys. The patterns show the InTe material with elemental Te precursor. InTe is a pure mono-chalcogenide, indium telluride (ICDD: 00-007-0112) with a tetragonal phase, crystallised in space group I4/mcm.

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The SAED patterns identify this pattern in the TEM Figure 5-4. When selenium is introduced,

In+Te1-xSex x=0.25 shows an XRD pattern with shifted peak positions of the InTe phase indicating selenium alloying into InTe. A second orthorhombic In4Te3 phase, ICDD: 01-083-

0040, also occurs, identified in the SAED pattern of Figure 5-4a. The blue markers identify the

In4Se3 pattern, and elemental Te (ICDD: 01-085-0555), identified by a purple *. When the

In+Te1-xSex x=0.75 pattern is analysed, hexagonal In2Se3 (ICDD; 00-040-1407), seen in the

SAED pattern, and rhombohedral InSe (ICDD: 01-071-0447) both occur with no tellurium phases present. In the pure selenium sample, a pure phase of hexagonal InSe forms, as discussed earlier in Figure 5-2. These XRD patterns show that there is some alloying of selenium into tetragonal InTe but that the sample becomes a phase mixture and no other alloying is presented.

Figure 5-6 XRD pattern of the alloyed indium telluride and sulphide materials ranging from pure indium telluride to pure indium sulphide. The blue * identifies InTe in the middle pattern, while red * identifies In2S3.

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Figure 5-6 shows the XRD pattern of the indium and sulphur/tellurium samples. When sulphur is introduced as a 25 molar mass percentage, i.e. In+Te1-xSx x=0.25, the patterns show the InTe phase without any peak shift, suggesting that there is no sulphur incorporation, similar to the lack of sulphur infiltration into the hexagonal InSe earlier. A second phase also occurs that is identified as cubic β-In2S3 (ICDD:00-032-0456). This β-In2S3 phase becomes the dominant phase when the concentration of sulphur increases to 75% and above. These four XRD patterns show that there is no alloying between sulphur and tellurium when the chalcogenides are mixed in the reaction system.

5.5 Conclusion

Alloyed In+Se1-xSx nanomaterials composed of nanoribbons and nanosheets are synthesized in a phosphine free colloidal system. The concentration of chalcogenide precursor has a significant influence on the morphology of the resultant nanomaterials. Se incorporates into

In2S3 up to 25%, and S incorporates into In2Se3 but not into InSe. Mixing of the chalcogenides immediately changes the nano-hexagonal sheets into highly anisotropic nanoribbons that broaden as the concentration of Sulphur increases and the main phase becomes In2S3 with Se alloying. Selenium incorporation into the In2S3 materials reduces the bandgap from 3.19 to

3.04 eV, which are all more significant than the 2.1 eV expected for pure In2S3. Then, this synthesis's power to mix the chalcogenide material in In+S/Te and In+Se/Te is investigated.

The materials formed are nanoplatelets with much more agglomeration than the selenium- sulphur alloys, indicating the lesser ability for OLA to fulfil both the solvent and surfactant roles. XRD data shows that some selenium atoms alloy into InTe. However, tellurium does not alloy into InSe. Likewise, there is no alloying of sulphur and tellurium. Instead, separate phases of InTe and In2S3 occur.

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5.6 Bibliography

[1] S. M. Oh, S. B. Patil, X. Jin, and S.-J. Hwang, "Recent Applications of 2D Inorganic Nanosheets for Emerging Energy Storage System," Chemistry – A European Journal, vol. 24, no. 19, pp. 4757-4773, 2018/04/03 2018, doi: https://doi.org/10.1002/chem.201704284. [2] S. Zhang et al., "The latest development of CoOOH two-dimensional materials used as OER catalysts," Chemical Communications, 10.1039/D0CC05876A vol. 56, no. 98, pp. 15387-15405, 2020, doi: 10.1039/D0CC05876A. [3] X. Kong, X. Shen, C. Zhang, S. N. Oliaee, and Z. Peng, "Engineering active sites of two- dimensional MoS2 nanosheets for improving hydrogen evolution," Inorganic Chemistry Frontiers, 10.1039/C6QI00334F vol. 3, no. 11, pp. 1376-1380, 2016, doi: 10.1039/C6QI00334F. [4] K. Khan et al., "Recent developments in emerging two-dimensional materials and their applications," Journal of Materials Chemistry C, 10.1039/C9TC04187G vol. 8, no. 2, pp. 387- 440, 2020, doi: 10.1039/C9TC04187G. [5] I. Ben Amara, R. Bennaceur, S. Jaziri, and H. Ben Abdallah, "Selenium alloying of indium sulfide: Ab-initio study of structural, electronic and optical features," Materials Science in Semiconductor Processing, vol. 31, pp. 56-67, 2015/03/01/ 2015, doi: https://doi.org/10.1016/j.mssp.2014.10.036. [6] S. Chen et al., "Metal selenide photocatalysts for visible-light-driven Z-scheme pure water splitting," Journal of Materials Chemistry A, 10.1039/C9TA00768G vol. 7, no. 13, pp. 7415- 7422, 2019, doi: 10.1039/C9TA00768G. [7] K. Hassan et al., "Functional inks and extrusion-based 3D printing of 2D materials: a review of current research and applications," Nanoscale, 10.1039/D0NR04933F vol. 12, no. 37, pp. 19007-19042, 2020, doi: 10.1039/D0NR04933F. [8] Y. Huang et al., "Universal mechanical exfoliation of large-area 2D crystals," Nature communications, vol. 11, no. 1, pp. 1-9, 2020. [9] G. Hu et al., "Functional inks and printing of two-dimensional materials," Chemical Society Reviews, 10.1039/C8CS00084K vol. 47, no. 9, pp. 3265-3300, 2018, doi: 10.1039/C8CS00084K. [10] S. Delikanli et al., "Ultrathin Highly Luminescent Two-Monolayer Colloidal CdSe Nanoplatelets," Advanced Functional Materials, vol. 29, no. 35, p. 1901028, 2019, doi: 10.1002/adfm.201901028. [11] D. Son et al., "Colloidal Synthesis of Uniform-Sized Molybdenum Disulfide Nanosheets for Wafer-Scale Flexible Nonvolatile Memory," Advanced Materials, vol. 28, no. 42, pp. 9326- 9332, 2016, doi: 10.1002/adma.201602391. [12] X. Li, X. Zuo, X. Jiang, D. Li, B. Cui, and D. Liu, "Enhanced photocatalysis for water splitting in layered tin chalcogenides with high carrier mobility," Physical Chemistry Chemical Physics, 10.1039/C9CP00088G vol. 21, no. 14, pp. 7559-7566, 2019, doi: 10.1039/C9CP00088G. [13] J. Lauth et al., "Solution-Processed Two-Dimensional Ultrathin InSe Nanosheets," Chemistry of Materials, vol. 28, no. 6, pp. 1728-1736, 2016/03/22 2016, doi: 10.1021/acs.chemmater.5b04646. [14] F. Horani and E. Lifshitz, "Unraveling the Growth Mechanism Forming Stable γ-In2S3 and β- In2S3 Colloidal Nanoplatelets," Chemistry of Materials, vol. 31, no. 5, pp. 1784-1793, 2019/03/12 2019, doi: 10.1021/acs.chemmater.9b00013.

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[15] S. Lei et al., "Optoelectronic memory using two-dimensional materials," Nano letters, vol. 15, no. 1, pp. 259-265, 2015. [16] S. Sucharitakul et al., "Intrinsic electron mobility exceeding 103 cm2/(V s) in multilayer InSe FETs," Nano letters, vol. 15, no. 6, pp. 3815-3819, 2015. [17] W. Feng, W. Zheng, W. Cao, and P. Hu, "Back gated multilayer InSe transistors with enhanced carrier mobilities via the suppression of carrier scattering from a dielectric interface," Advanced materials, vol. 26, no. 38, pp. 6587-6593, 2014. [18] S. R. Tamalampudi et al., "High performance and bendable few-layered InSe photodetectors with broad spectral response," Nano letters, vol. 14, no. 5, pp. 2800-2806, 2014. [19] S. Lei et al., "Evolution of the electronic band structure and efficient photo-detection in atomic layers of InSe," ACS nano, vol. 8, no. 2, pp. 1263-1272, 2014. [20] S. Lei et al., "An atomically layered InSe avalanche photodetector," Nano letters, vol. 15, no. 5, pp. 3048-3055, 2015. [21] R. Mane and C. Lokhande, "Studies on structural, optical and electrical properties of indium sulfide thin films," Materials chemistry and physics, vol. 78, no. 1, pp. 15-17, 2003. [22] K. Hara, K. Sayama, and H. Arakawa, "Semiconductor-sensitized solar cells based on nanocrystalline In2S3/In2O3 thin film electrodes," Solar Energy Materials and Solar Cells, vol. 62, no. 4, pp. 441-447, 2000. [23] W. Chen, J.-O. Bovin, A. G. Joly, S. Wang, F. Su, and G. Li, "Full-color emission from In2S3 and In2S3: Eu3+ nanoparticles," The Journal of Physical Chemistry B, vol. 108, no. 32, pp. 11927- 11934, 2004. [24] F. Ye et al., "Facile and rapid synthesis of RGO–In 2 S 3 composites with enhanced cyclability and high capacity for lithium storage," Nanoscale, vol. 4, no. 23, pp. 7354-7357, 2012. [25] R. Jayakrishnan, "Photoluminescence in Spray Pyrolysis Deposited β-In2S3 Thin Films," Journal of Electronic Materials, vol. 47, no. 4, pp. 2249-2256, 2018/04/01 2018, doi: 10.1007/s11664- 017-6047-y. [26] R. B. Jacobs-Gedrim et al., "Extraordinary photoresponse in two-dimensional In2Se3 nanosheets," ACS nano, vol. 8, no. 1, pp. 514-521, 2014. [27] J. Cui, L. Wang, Z. Du, P. Ying, and Y. Deng, "Correction: High thermoelectric performance of a defect in α-In2Se3-based solid solution upon substitution of Zn for In," Journal of Materials Chemistry C, 10.1039/C5TC90163D vol. 3, no. 37, pp. 9750-9750, 2015, doi: 10.1039/C5TC90163D. [28] Y. Jiang et al., "Construction of In2Se3/MoS2 heterojunction as photoanode toward efficient photoelectrochemical water splitting," Chemical Engineering Journal, vol. 358, pp. 752-758, 2019. [29] B. Xue, F. Xu, B. Wang, and A. Dong, "Shape-controlled synthesis of β-In 2 S 3 nanocrystals and their lithium storage properties," CrystEngComm, vol. 18, no. 2, pp. 250-256, 2016. [30] Y. Hu et al., "Temperature-dependent growth of few layer β-InSe and α-In2Se3 single crystals for optoelectronic device," Semiconductor Science and Technology, vol. 33, no. 12, p. 125002, 2018. [31] S. Wan et al., "Nonvolatile Ferroelectric Memory Effect in Ultrathin α‐In2Se3," Advanced Functional Materials, vol. 29, no. 20, p. 1808606, 2019.

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[32] S. Douman et al., "New Generation Nanoelectrochemical Biosensors for Disease Biomarkers: 1. Indium Telluride Quantum Dots Signaling of Telomerase Cancer Biomarker," Journal of Nanoscience and Nanotechnology, vol. 16, no. 12, pp. 12844-12850, 2016. [33] J. Tabernor, P. Christian, and P. O'Brien, "A general route to nanodimensional powders of indium chalcogenides," Journal of Materials Chemistry, 10.1039/B600921B vol. 16, no. 21, pp. 2082-2087, 2006, doi: 10.1039/B600921B. [34] Y. Fu, J. Zhou, H.-H. Zou, F. A. Almeida Paz, X. Liu, and L. Fu, "Unique two-dimensional indium telluride templated by a rare wheel-shaped heterobimetallic Mn/In cluster," Inorganic chemistry, vol. 59, no. 9, pp. 5818-5822, 2020. [35] Q. Zhang, I. Chung, J. I. Jang, J. B. Ketterson, and M. G. Kanatzidis, "A polar and chiral indium telluride featuring supertetrahedral T2 clusters and nonlinear optical second harmonic generation," Chemistry of Materials, vol. 21, no. 1, pp. 12-14, 2009. [36] R. Chen et al., "A new polymorph telluridoindate [In (en) 3][In5Te9 (en) 2] with photocatalytic properties," Inorganic Chemistry Communications, vol. 28, pp. 55-59, 2013. [37] R. Biswas, P. Deb, and S. Das, "Novel optoelectronic properties in barbed wire nanophotonic structures of indium telluride," Optical Materials, vol. 47, pp. 586-588, 2015. [38] V. Sowjanya, "Preparation of Indium Telluride Thin Films for Device Applications," National Institute of Technology Karnataka, Surathkal, 2019. [39] S. C. Riha, B. A. Parkinson, and A. L. Prieto, "Compositionally Tunable Cu2ZnSn(S1–xSex)4 Nanocrystals: Probing the Effect of Se-Inclusion in Mixed Chalcogenide Thin Films," Journal of the American Chemical Society, vol. 133, no. 39, pp. 15272-15275, 2011/10/05 2011, doi: 10.1021/ja2058692. [40] Y. Qu, G. Zoppi, and N. S. Beattie, "Selenization kinetics in Cu2ZnSn(S,Se)4 solar cells prepared from nanoparticle inks," Solar Energy Materials and Solar Cells, vol. 158, pp. 130-137, 2016/12/01/ 2016, doi: https://doi.org/10.1016/j.solmat.2015.12.016. [41] Q. Guo et al., "Fabrication of 7.2% Efficient CZTSSe Solar Cells Using CZTS Nanocrystals," Journal of the American Chemical Society, vol. 132, no. 49, pp. 17384-17386, 2010/12/15 2010, doi: 10.1021/ja108427b. [42] Y. Zou, X. Su, and J. Jiang, "Phase-Controlled Synthesis of Cu2ZnSnS4 Nanocrystals: The Role of Reactivity between Zn and S," Journal of the American Chemical Society, vol. 135, no. 49, pp. 18377-18384, 2013/12/11 2013, doi: 10.1021/ja405962k. [43] M. A. Airo, S. Gqoba, F. Otieno, M. J. Moloto, and N. Moloto, "Structural modification and band-gap crossover in indium selenide nanosheets," RSC Advances, 10.1039/C6RA00262E vol. 6, no. 47, pp. 40777-40784, 2016, doi: 10.1039/C6RA00262E. [44] M. Li et al., "Controlled growth of vertically aligned ultrathin In 2 S 3 nanosheet arrays for photoelectrochemical water splitting," Nanoscale, vol. 10, no. 3, pp. 1153-1161, 2018.

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Chapter 6 Germanium-Tin Alloying at B site of Inorganic CsBX3

(X = Br, I) and Cs2BX6 Perovskite Systems.

6.1 Abstract

Perovskite optoelectric devices have changed the research landscape in recent years. However, they suffer from the presence of lead halide, a weakly bound toxic compound that is highly water-soluble, allowing it to enter groundwater quickly. Given this, substitution of Pb with more suitable elements is desirable. Ge substitution over Pb in perovskites increases stability in aqueous environments by forming a more covalent halide bond, while Sn substitution lowers the bandgap and maintains desired optical properties of Pb perovskites. This work presents the first solution-processed Sn and Ge alloyed perovskite. Systematically varying Sn's ratio to Ge in the synthesis of bromide and iodide perovskites identifies an upper limit of Sn substitution into CsGeX3. An Sn-based vacancy ordered double perovskite with no Ge alloying is produced above this concentration. Instead, Ge incorporation leads to secondary growth on the Sn perovskite that includes the formation of nanorods. Both iodide and bromide perovskites are examined, and the optical properties of the alloyed Sn-Ge perovskites show suitability for photovoltaics and light-emitting diodes (LEDs), respectively. They are stable for many weeks in the air and are a more environmentally friendly alternative to the Pb perovskite material.

6.2 Introduction

Alloys of semiconductor materials are excellent materials in solar cells, LEDs, and battery anodes.[1-4] The resultant bandgap engineering enables materials to reach the desired optical and electrical properties while reducing toxic elements' reliance. For example, the alloy of sulphur and selenium in CdSexS1-x enables the tuning of absorption and photoluminescence emission, significantly improving solar cell Power Conversion Efficiency (PCE) and light-

6-145 emitting diodes.[5-8] Combining two active materials like germanium and silicon (Ge-Si), can result in improved conduction pathways and higher capacity retention in Li-ion batteries.[1, 9-

12] Therefore, investigation of new semiconductor alloys with hybrid properties will pave the way to future devices that surpass today's materials.

Halide perovskites have focused on intense research over the past decade due to their potential in solar cells, with photovoltaic properties reaching 25% in 2020 and 29.52% in a silicon tandem cell.[13-17] They also demonstrate promising properties suitable for a range of other optoelectric applications, including light-emitting diodes (LEDs), lasers and photodetectors.[18-20] The origin of these characteristics is the intrinsic and strong spin-orbit coupling (SOC) effects that stem from the group 14 elements (Pb, Sn or Ge) which significantly lower the optical bandgap down towards the NIR region and preserve the optical absorption despite local defects in the lattice.[21] In the absence of inversion symmetry, these SOC effects give rise to Rashba effects, where electric fields can manipulate spin-dependent properties in spintronics. [22] As a result, perovskites display a bandgap between direct and indirect in nature as it displays characteristics from both.[23] Direct semiconductors have high absorption coefficients, while indirect semiconductors have much slower recombination of electron-hole pairs, both of which are defining properties of halide perovskites.

Single perovskites of ABX3 composition are not the optimal option for applications as demonstrated by the limited stability: decomposition into binary components (e.g. MAPbI3 degrades to MAI + PbI2), oxidation (e.g. CsSnI3 converting into Cs2SnI6) loss of halide, and possible phonon (dynamic) instability. [24-31] The properties of alloys can change significantly vs the constituents' respective properties opening the possibility of alloyed perovskites solving these problems. [32] The record-breaking solar cells based on perovskites contain mixtures. Studies have shown that mixing caesium (Cs), Rubidium (Rb), and

6-146 formamidinium (FA) with methylammonium (MA) enhances the thermal stability of hybrid perovskites. [33, 34] Meanwhile, halide anions have been mixed in perovskite nanocrystals to tune the bandgap.[35-37] Likewise, substitution and alloying of the B cation site of halide perovskites can remove toxic Pb and increase structural stability. [38-44] Suitable alternative metal ions have an electronic configuration of ns2 such as Sn2+, Ge2+, Bi3+, Sb3+, In+, and Tl+, similar to Pb2+.[45, 46] Higher in group 14 atoms make the metal-halogen bond more covalent due to their reduced electronegativity. In Sn-iodide perovskites, the bandgap reduces to close to the ideal, while Ge perovskites are more stable due to covalency. [2] An ideal bandgap between 1.15 to 1.4 eV absorbs light throughout the visible and into the NIR region of the electromagnetic spectrum, paving the way to high-efficiency tandem photovoltaic devices.[2]

The majority of B-cation alloying reports focus on reducing rather than eliminating Pb or using a B cation pair of mono- and tri-valent atoms that cause indirect bandgaps. These do not achieve the record-breaking optoelectric properties of divalent atoms, and so this work investigates alloying of divalent Ge and Sn.

Figure 6-1 shows the bandgaps of various Pb-free, Sn, and Ge based perovskites. Variations of organo-tin iodide perovskites reach the highest PCE of 8.12%.[47] Caesium cation reduces the instability in humid conditions brought about by volatile organic cations.[48] This substitution also lowers the 2.0 eV bandgap of CH3NH3GeI3 to 1.6 eV of CsGeI3. Studies of this inorganic

CsSnI3 material found that incorporation of divalent Sn halide additives (SnF2, [49] SnCl2,[50] and SnI2[51]) is an effective strategy for the prevention of oxidation of Sn(II) to Sn(IV). [52]

However, Sn perovskite's metallic conductivity hampers the development of efficient CsSnI3 based solar cells. [50, 53, 54] This has garnered interest in the compound Cs2SnI6, a double perovskite crystal structure that contains ordered vacancies due to the +4 oxidation state of

Sn.[25, 53] The corresponding Cs2GeX6 materials are not theoretically stable. [55]

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Figure 6-1 Bandgaps of inorganic perovskites with B-cation substitution by group 14 elements Sn and Ge.[43, 48, 56]

Chen et al. discovered that alloying Ge(II) into CH3NH3SnI3 in a solid solution forms a highly stable and air-tolerant thin film. This material displays a Goldschmidt tolerance factor of 0.94, and an octahedral factor of 0.4 explains the alloy's structural stability. Added to this, the higher high oxidation activity of Ge(II), which makes it challenging to work with alone, forms an ultrathin (<5 nm) native oxide layer. Through melt-crystallization and vapour deposition, Chen et al. produced a solar cell device with a PCE of 7.11%, stable under continuous illumination of 1-sun for 500 h.

There are no reports of these alloyed materials synthesised by solution-based processes. Here, this work investigates a series of Pb-free, inorganic mixed Ge and Sn perovskites based on bromide and iodide halides. Both perovskites reach an alloying limit, found through characterisation of the morphology and structural properties. Sn incorporates into the rhombohedral germanium perovskite up to 25% in iodide perovskites and up to 50% in bromide perovskites. XRD analyses the phase changes, and SEM provides information on the structure.

XPS examines the materials' oxidation states, and finally, diffuse reflectance spectroscopy investigates the optical properties.

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6.3 Experimental

Materials: Caesium Iodide (CsI, 99.999%), Caesium bromide (CsBr, 99.999%), Germanium

(IV) Oxide (GeO2, 99.99%), Tin Chloride (SnCl2, 99.99%), Tin Bromide (SnBr2, 99%), Tin

Iodide (SnI2, 99.99%), Hydrochloric acid (HCl, 37%), Hydrobromic acid (HBr, 48%),

Hydriodic acid (HI, 57%, stabilized, 99.95%), Hypophosphorous Acid (H3PO2%) and IPA were sourced from Sigma Aldrich and used as received.

Synthesis: Caesium metal halide perovskites were prepared by a temperature lowering method.

Briefly, germanium (IV) oxide and/or tin (II) halide (Up to 1 mmol), hydrohalic acid (5 mL), and hypophosphorous acid (2 mL) were stirred at 120 °C in an argon atmosphere in a 25 ml 3- neck-flask until the mixture became homogeneous. Caesium halide (1 mmol in 2 ml hydrohalic acid) was injected at 120 °C, and then the solution was immediately cooled to room temperature, with precipitate forming between 50 – 70 °C. The precipitate was collected by centrifugation for 5 min at 4500 rpm, washed twice with IPA then stored in IPA to prevent metal oxidation.

6.4 Results and discussion

The materials are grown using the temperature lowering method as the perovskite solubility in an acid-halide solvent is moderate at room temperature but increases considerably with temperature. This method often uses a seed crystal dipped into a precursor solution; however, monomer supersaturation induces nucleation and crystal growth here. The slower reduction of temperature allows increased crystallisation. Figure 6-2a shows a schematic of the synthesis method used in this reaction. The Cs precursor solution, comprised of CsBr dissolved in HBr, is injected into a solution of SnBr2 and/or GeBr2 in HBr with H3PO2 reducing agent at 120 °C.

Powering off or complete removal of the heating mantle leads to precipitate formation as the

6-149 system is cooled, becoming visible to the naked eye close to 70 °C. Once cooled, isopropanol is used as a miscible antisolvent with water, and the precipitate is collected by centrifuging at

5000 rpm for 5 minutes. IPA is used to rinse away the acid and isolate the material from the air.

Figure 6-2 a) Reaction setup schematic. b) Crystal structure of rhombohedral CsBBr3 distorted perovskite. c) Diagram of cubic double perovskite Cs2BBr6. d-h) SEM images with 20µm scale of (d) Sn bromide perovskite the following cooling to RT slowly. (e) Sn bromide perovskite following fast quenching with an ice water bath. Figure 6-2b+c show the schematic diagrams of the two possible perovskite structures for these materials. Figure 6-2b shows a rhombohedral distorted perovskite, while Figure 6-2c shows a cubic, vacancy ordered double perovskite associated with the oxidation of the B site cation from 2+ to 4+. Figure 6-2d shows an SEM image of a sample made using CsBr, SnBr2 and

HBr, referred to as CsGexSn1-xBr (Gex = 0). Here, a room temperature water bath quenches the system. The sample in Figure 6-2e is cooled slowly by switching off the heating mantle but leaving it in place. The material grows into crystals with lower polydispersity in the slowly cooled sample than with fast cooling. These images illustrate the effect of quenching speed on the reaction without surfactants.

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A series of iodide perovskites were first examined for the synthesis of alloyed B site materials comprised of CsGexSn1-xI (Gex = 0, 0.25, 0.5, 0.625, 0.75, 0.875, 1). Figure 6-3 outlines the structural and morphological characterisation of these iodide based perovskites. Figure 6-3a displays the XRD patterns for the array of samples, with Figure 6-3b showing higher resolution between 42-44°. In CsGexSn1-xI (Gex = 0), the XRD pattern contains peaks from two patterns:

Cubic Cs2SnI6 peaks at 26.75° (222), 30.9° (400), 44.2° (440), 52.3° (622) and 54.8° (444) are highlighted by the blue arrows, ICDD: 96-433-5643. The orange arrows identify peaks of orthorhombic CsSnI3 (yellow, room temperature phase) ICDD: 96-411-7957. These phases are prominent in CsGexSn1-xI (Gex = 0, 0.25, 0.5, 0.625), the top four patterns, (purple/dark blue) and there is little structural change as Ge concentration is increased however, when the molar ratio of Ge is increased above that of Sn, CsGexSn1-xI (Gex = 0.5, 0.625) to 50 and 62.5% Ge, a third phase appears: Cubic CsI (ICDD: 01-089-4257) with peaks at 27.8 (011), 39.68 (020),

49.05° (121), highlighted in green. There is an unidentified cubic phase in these XRD patterns preventing Reitveld refinement due to a hypophosphite salt leftover from the oxidation of the

H3PO2 reducing agent.

Figure 6-3b show the three germanium rich samples, CsGexSn1-xI (Gex = 0.75, 0.875, 1), contain pure hexagonal CsGeI3 (ICDD: 01-085-1274). Peak shifting shows the alloying as Sn concentration increases. The blue pattern of CsGexSn1-xI (Gex = 0.625) shows a mixture of phases, containing CsGeI3, CsI and defect perovskite Cs2MI6 as well as GeO2 and GeI4. This pattern establishes the limit of 25% Sn substitution into CsGeI3, after which the phase changes to a vacancy ordered perovskite and does not incorporate all of the CsI precursors. This Cs2SnI6 phase pattern does not shift peak position in response to Ge incorporation.

2+ 4+ In the germanium free sample, CsSnI3 does form. However, some Sn oxidises to Sn creating a vacancy ordered perovskite Cs2SnI6. The octahedra do not connect in this structure, and every

6-151 second B site is left vacant. Neither CsI precursor reduction nor the addition of excess H3PO2

2+ 4+ reducing agent prevents the oxidation of some Sn to Sn and does not stop the formation of

A2BX6. These results show that the initial CsX to BX ratio of precursors is different for Ge and

Sn perovskites, resulting in CsI precursor remaining and different reaction kinetics.

Figure 6-3 Characterisation of iodide based perovskite alloys. a) XRD of the Sn-based materials with increasing Ge percentages and phase identification through colour-coded arrows. b) Inspection of Sn incorporation into CsGeI3 between 42 and 45°. C) XPS pattern for Iodide perovskite materials. d) Identification of Sn oxidations states e) Identification of germanium oxidation states.

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XPS was used to study the electronic structure of the iodide materials. Figure 6-3c-e shows the core line data collected for the pure Sn and 50% Ge samples. In agreement with the XRD,

4+ 2+ pattern Ge is present without any Ge . While Cs2GeX6 is theoretically unstable, this XPS result might support Ge's substitution into some of the B cation sites of Cs2SnI6. However, the

4+ 4+ XRD pattern contains the GeO , i.e. GeO2 pattern, which accounts for the Ge .

Figure 6-4 shows the SEM images of the iodide perovskites. The tin rich materials CsGexSn1- xI (Gex = 0, 0.25, 0.5) Figure 6-4a-c contain 1D rods with lengths longer than 100 µm in all three cases. These rods seem cylindrical at low magnification; however, as seen in Figure 6-4g- i, they are cuboids without complete growth to the edges. These edges are high energy active sites. The rods stop occurring in the SEM images when Cs2SnI6 stops appearing in the XRD patterns. Therefore the rods are identified as Cs2SnI6. The images also contain smaller 3D particles with sizes between 1-5 µm and morphology resembling truncated octahedra, or materials with a high number of facets. These strongly resemble the single crystals of other iodide ABX3 perovskites, including the dominant photovoltaic materials, organo-metallic

MAPbI3 and MASnI3.[57-59] The truncated octahedra may be identified as the ABX3 perovskite, CsSnI3, as there is no other phase identified in the XRD patterns.

Figure 6-4g-i show a considerable increase in secondary growth in all samples that contain the

Cs2SnI6 phase as Ge concentration increases. In CsGexSn1-xI (Gex = 0.5), the 50% Ge sample, there are small 1D rods in batches throughout the material, most apparent in the 1st highlighted area of Figure 6-4i.

This group of rods have a diameter of 0.13 ± 0.2 µm. There is an angular crystal seen clearer in the 2nd highlighted area of Figure 6-4i, on the top of each rod that is larger than the rod diameter, suggesting that particles might nucleate as seed particles on the surface of a large facet and a rod grows between the large original particle and the smaller secondary particle,

6-153 similar to the growth of nanowires form a seed on a substrate.[60] Figure 6-4i shows rods of various lengths, sometimes with no rods, such as the first highlighted box in the image; however, no rods form without the capping crystal.

Figure 6-4 a-f) SEM images corresponding to the XRD patterns shown above. Scale bars of 50µm with an inset scale bar of 500µm. g-l) High-resolution SEM images of iodide perovskites showing secondary growth. Scale bars are 5 µm.

The CsGexSn1-xI (Gex = 0.625) in Figure 6-4j contains many small particles on the crystal vertices. However, they do not form into rods. As vertices fill, the surface energy lowers and becomes equal to that of the edges, meaning both positions are equally likely to nucleate

6-154 growth.[61] As Ge concentration increases, the growth continues as normal, increasing the size of the crystals. However, as Ge concentration decreases below 50%, 1D structures grow. The growth may occur due to the isopropyl alcohol washing etching and precipitating material or due to uneven temperature profiles. Cooling the system could lead to temperature differentials in the solvent as stirring stopped. It is possible that as the temperature drops and crystals form, the temperature increases again, dissolving the material. As it cools again, monomers increase, and further nucleation can occur at high-energy sites. Defects on the existing crystals act as nucleation centres from which these new growths occur. This mechanism would explain secondary growth from both the large 1D rods and the smaller octahedra.

In the Ge rich samples, the formation of CsGeI3 is preferred over the 1D Cs2MI6 as Cs2GeX6 is theoretically unstable. These CsGeI3 crystals have a truncated octahedral morphology and crystallite sizes between 5-10 µm. The remaining Ge samples consist of increasingly larger 3D particles.

In the bromide perovskite investigation, the GeBr2 precursor was not available throughout the

Covid-19 pandemic. GeO2 was used in its place with the H3PO2 reducing agent forming GeBr2 within the system. Figure 6-5 shows the XRD and XPS characterisation.

Similar to the iodide perovskite, in these bromide perovskites, the CsGexSn1-xBr (Gex = 0), sample forms both cubic CsSnBr3 perovskite (orange arrows) ICDD: 01-070-1645, and a cubic

Cs2SnBr6 pattern, a vacancy ordered perovskite, (blue arrows) ICDD: 00-101-0559. This pattern is the only instance of pure Sn perovskite, and further examination shows a reduction and eventual disappearance of this phase as time in air increases due to the expected oxidation of Sn.

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Figure 6-5 Characterisation of bromide based perovskite alloys. a) XRD of the Sn-based materials with increasing Ge percentages and phase identification through colour-coded arrows. b) Inspection of Sn incorporation into CsGeBr3 between 44 and 47°. c-h) SEM images corresponding to the XRD patterns shown above. Scale bars of 50µm.

Substitution of 25% Ge into the system in place of Sn, CsGexSn1-xBr (Gex = 0.25), the pattern shows only a pure vacancy ordered perovskite with no other phases present. Further substitution to 50%, CsGexSn1-xBr (Gex = 0.5) again shows a mixed-phase pattern. Here, the main (110) peak of CsBr occurs at 29.5°, along with (100) at 20.8°, (111) at 36.35°, (200) at

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42.19° and (211 at 52.3°. These are identified by the green arrow, ICDD: 01-089-3628. When

Ge concentration is above 50%, a rhombohedral phase of CsGeBr3 (ICDD: 01-085-1273) is the main phase with the cubic's continued presence CsBr phase. Figure 6-5b shows the peak shifting due to the Sn-Ge substitution from 44 to 47°.

Figure 6-5c-e shows the XPS data of the bromide perovskites. Changing from bromide to iodide in the 100% Sn perovskites decreases the BEs of Cs 3d5/2 and Sn 3d5/2 by 3.75 eV and

0.65 eV, respectively. This decrease agrees with Karim et al. and Han et al., who both report the Sn 3d5/2 peak is at higher energy in Cs2SnBr6 than in Cs2SnI6. [27, 35] To determine the separation of Sn2+ and Sn4+ binding energies, the valence band spectra are needed. Iodide use decreases the BE of Ge4+ by 0.15 eV. However, there is a blue peak in the bromide material, indicating the presence of Ge2+ at 30.67 eV. These results agree with XRD results, suggesting

2+ 4+ that AB X3 is the main phase of the 50:50 bromide perovskite while A2B X6 is the main phase of the 50:50 iodide material.

Figure 6-6 shows the SEM images of these bromide perovskite materials. The caesium tin bromide sample shown in Figure 6-6a shows octahedra with large polydispersity. Sizes range from less than 1 to 3 µm. The material is polydisperse but highly regular. The addition of 25%

Ge forms larger octahedra that are similarly polydisperse. There are many instances of crystals less than 1 µm. However, the majority are between 2-4 µm with some larger particles >5 µm.

50% germanium changes the morphology from entirely octahedral crystals to almost entirely large >5 µm cubic crystals with few octahedra. This morphology is throughout the remaining samples as germanium concentration increases, which agrees with the XRD patterns shown in

Figure 6-5a. CsGeBr3 is the dominant phase, with Sn alloying and no tin perovskite phases.

CsBr appears in all of the patterns except the pure germanium perovskite pattern, which

6-157 corresponds to the sample with no secondary growth indicating that the CsBr may be the smaller material between the cubes in Figure 6-6c-f.

Figure 6-6 SEM images of bromide based tin-germanium alloyed perovskites. Scale bars are 20 µm Figure 6-7a+b shows the normalised diffuse reflectance and subsequent KMF data of the iodide perovskites. UV-Vis diffuse reflectance spectra (DRS) identified the bandgap energy of the powdered samples. Figure 6-7a presents the DRS of iodide perovskites as a function of wavelength in the range of 400-1200 nm. The normalised reflectance shows a gradual shift of the absorption onset towards lower energy as Sn concentration is increased, from ~740 nm to

~965 nm. The Kubelka-Munk theory identified the direct bandgap of the materials. At any wavelength, the Kubelka-Munk Function (KMF), F(R) equation is:

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퐾 (1 − 푅)2 Equation 6-1 퐹(푅) = = 푆 2푅

Where K and S are scattering and absorbing coefficients, R is the reflectance. As the scattering factor is independent of wavelength, F(R) is proportional to S. [62]. Plotting F(R) from

Equation 6-1 against photon energy (hv) finds bandgap energies through plotting the trend line.

Figure 6-7b shows these bandgap energies.

The XRD of these single cation perovskites CsGexSn1-xI (Gex = 1, 0) shows mixed phases; however, the single cation perovskites show bandgaps very close to the theoretical for Sn (1.3 eV) and Ge (1.6 eV) iodide perovskites. [63, 64] This indicates that a large proportion of pure

ABX3 perovskite in the samples. In contrast to the XRD data, Ge addition to the Sn perovskite causes a blue shift of 45 meV, proving metal site substitution, which showed minimal change.

When Ge concentration increases over Sn in the iodide perovskites, the bandgap increases above the Sn perovskite's theoretical bandgap. From here, each 12.5% substitution affords large redshifts in the bandgaps observed. Until the CsGexSn1-xBr (Gex = 0.875, 1) samples, both of which show similar bandgaps attributed to CsGeI3 material without Sn incorporation.

The Shockley-Queisser Limit is the maximum theoretical efficiency that a solar cell can achieve, and it describes the ideal bandgap to reach this limit as 1.33 eV. The 62.5% Ge sample

CsGexSn1-xBr (Gex = 0.625) has a bandgap of 1.34 eV, extraordinarily close and implying massive potential for this sample. This sample is a mix of 3 phases, ABX3, A2BX6 and AX.

Meanwhile, the 75% Ge sample presents a 1.4 eV bandgap, and according to the XRD and

SEM data from Figure 6-3, it is a pure rhombohedral ABX3 phase and is significantly more stable than the Sn counterpart, evidenced by the lack of vacancy ordered material in the XRD pattern.

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Figure 6-7 UV-Vis Diffuse Reflectance Data for a) the iodide perovskites series with the K-M function applied in b). c) shows the bromide perovskite reflectance spectra with d) the K-M function. The insets of the KM plots display the bandgap values. Figure 6-7c+d shows the bromide perovskites' optical data with the bandgaps listed in the inset of Figure 6-7b. Here, the CsGexSn1-xBr (Gex = 1) perovskite has a bandgap of 2.407 eV, larger than the expected 2.2 eV for known CsGeBr3 materials. The bandgap increases slightly but remains close to 2.5 eV as Sn is introduced up to 37.5% Sn, CsGexSn1-xBr (Gex = 0.875, 0.75,

0.625).

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In CsGexSn1-xBr (Gex = 1), a clear absorption edge occurs at 3.1 eV which is indicative of

Cs2SnBr6 which has a less known bandgap reported between 2.9-3.0 eV. There is also an absorption edge located at 1.82 eV, which indicates the presence of CsSnBr3. In CsGexSn1-xBr

(Gex = 0.5), the bandgap is recorded at 2.6 eV, attributed to the CsGeBr3 shown in the XRD pattern. The remaining Ge rich materials remain close to this value. Further investigation using photoluminescence spectroscopy would examine the applicability of these materials for use in

LED devices.

6.5 Conclusion

This work presents Sn substitution into two Ge perovskites through a solution based synthesis for the first time. SnI2 can be incorporated up to 25% before a breakdown of the ABX3 structure into different phases in the caesium germanium iodide-based materials. SnBr2 incorporates up to 50% into caesium germanium bromide. Applications for these materials include photovoltaics for the iodide alloy, specifically the 25% Sn incorporation into CsGeI3, i.e.

CsGexSn1-xI (Gex = 0.75), as this material is in a pure phase and presents a bandgap of 1.406 eV. There are potential applications for luminescent devices using bromide alloys.

Future work for these materials should include an investigation varying the cooling time on the crystal morphology. The effect of the acid concentration, temperature and alcohol polarity on the secondary growth effect needs research. This route might form nanosized A2BX6 and ABX3 materials when appropriate reaction conditions and surfactants are employed.

6.6 Bibliography [1] K. Stokes, H. Geaney, G. Flynn, M. Sheehan, T. Kennedy, and K. M. Ryan, "Direct Synthesis of Alloyed Si1–xGex Nanowires for Performance-Tunable Lithium Ion Battery Anodes," ACS Nano, vol. 11, no. 10, pp. 10088-10096, 2017/10/24 2017, doi: 10.1021/acsnano.7b04523. [2] T. Lei et al., "Electrical and Optical Tunability in All-Inorganic Halide Perovskite Alloy Nanowires," Nano Letters, vol. 18, no. 6, pp. 3538-3542, 2018/06/13 2018, doi: 10.1021/acs.nanolett.8b00603. 6-161

[3] G. Bree, H. Geaney, and K. M. Ryan, "Common Battery Anode Testing Protocols Are Not Suitable for New Combined Alloying and Conversion Materials," ChemElectroChem, vol. 5, no. 23, pp. 3757-3763, 2018, doi: https://doi.org/10.1002/celc.201800990. [4] E. Kasper, M. Kittler, M. Oehme, and T. Arguirov, "Germanium tin: silicon photonics toward the mid-infrared," Photonics Research, vol. 1, no. 2, pp. 69-76, 2013. [5] A. Singh, S. Singh, S. Levcenko, T. Unold, F. Laffir, and K. M. Ryan, "Compositionally Tunable Photoluminescence Emission in Cu2ZnSn(S1−xSex)4 Nanocrystals," Angewandte Chemie International Edition, vol. 52, no. 35, pp. 9120-9124, 2013, doi: 10.1002/anie.201302867. [6] J. Doherty et al., "Germanium tin alloy nanowires as anode materials for high performance Li- ion batteries," Nanotechnology, vol. 31, no. 16, p. 165402, 2020/01/28 2020, doi: 10.1088/1361-6528/ab6678. [7] G. Bree, H. Geaney, K. Stokes, and K. M. Ryan, "Aligned copper zinc tin sulfide nanorods as lithium-ion battery anodes with high specific capacities," The Journal of Physical Chemistry C, vol. 122, no. 35, pp. 20090-20098, 2018. [8] A. Singh, H. Geaney, F. Laffir, and K. M. Ryan, "Colloidal Synthesis of Wurtzite Cu2ZnSnS4 Nanorods and Their Perpendicular Assembly," Journal of the American Chemical Society, vol. 134, no. 6, pp. 2910-2913, 2012/02/15 2012, doi: 10.1021/ja2112146. [9] G. Flynn, Q. M. Ramasse, and K. M. Ryan, "Solvent vapor growth of axial heterostructure nanowires with multiple alternating segments of silicon and germanium," Nano letters, vol. 16, no. 1, pp. 374-380, 2016. [10] K. Stokes, G. Flynn, H. Geaney, G. Bree, and K. M. Ryan, "Axial Si–Ge Heterostructure Nanowires as Lithium-Ion Battery Anodes," Nano letters, vol. 18, no. 9, pp. 5569-5575, 2018. [11] S. Foley et al., "Tin-Based Oxide, Alloy, and Selenide Li-Ion Battery Anodes Derived from a Bimetallic Metal–Organic Material," The Journal of Physical Chemistry C, vol. 125, no. 2, pp. 1180-1189, 2021/01/21 2021, doi: 10.1021/acs.jpcc.0c06395. [12] M.-G. Ju, J. Dai, L. Ma, and X. C. Zeng, "Lead-Free Mixed Tin and Germanium Perovskites for Photovoltaic Application," Journal of the American Chemical Society, vol. 139, no. 23, pp. 8038-8043, 2017/06/14 2017, doi: 10.1021/jacs.7b04219. [13] S. D. Stranks et al., "Electron-Hole Diffusion Lengths Exceeding 1 Micrometer in an Organometal Trihalide Perovskite Absorber," Science, vol. 342, no. 6156, pp. 341-344, 2013, doi: 10.1126/science.1243982. [14] M. A. Green, A. Ho-Baillie, and H. J. Snaith, "The emergence of perovskite solar cells," Nature Photonics, Review Article vol. 8, p. 506, 06/27/online 2014, doi: 10.1038/nphoton.2014.134. [15] A. Mei et al., "A hole-conductor–free, fully printable mesoscopic perovskite solar cell with high stability," science, vol. 345, no. 6194, pp. 295-298, 2014. [16] F. Zhang and K. Zhu, "Additive engineering for efficient and stable perovskite solar cells," Advanced Energy Materials, p. 1902579, 2019. [17] D. Bi et al., "Multifunctional molecular modulators for perovskite solar cells with over 20% efficiency and high operational stability," Nature Communications, vol. 9, no. 1, p. 4482, 2018/10/26 2018, doi: 10.1038/s41467-018-06709-w. [18] Y. Yao et al., "Efficient Quantum Dot Light-Emitting Diodes Based on Trioctylphosphine Oxide- Passivated Organometallic Halide Perovskites," ACS Omega, vol. 4, no. 5, pp. 9150-9159, 2019/05/31 2019, doi: 10.1021/acsomega.9b00464.

6-162

[19] J. R. Harwell, G. L. Whitworth, G. A. Turnbull, and I. D. W. Samuel, "Green Perovskite Distributed Feedback Lasers," Scientific Reports, vol. 7, no. 1, p. 11727, 2017/09/15 2017, doi: 10.1038/s41598-017-11569-3. [20] Z. Dai et al., "Capillary-bridge mediated assembly of aligned perovskite quantum dots for high- performance photodetectors," Journal of Materials Chemistry C, 10.1039/C9TC01104H vol. 7, no. 20, pp. 5954-5961, 2019, doi: 10.1039/C9TC01104H. [21] M. T. Pham et al., "Origin of Rashba Spin-Orbit Coupling in 2D and 3D Lead Iodide Perovskites," Scientific Reports, vol. 10, no. 1, p. 4964, 2020/03/18 2020, doi: 10.1038/s41598-020-61768- 8. [22] A. Manchon, H. C. Koo, J. Nitta, S. M. Frolov, and R. A. Duine, "New perspectives for Rashba spin–orbit coupling," Nature Materials, vol. 14, no. 9, pp. 871-882, 2015/09/01 2015, doi: 10.1038/nmat4360. [23] Y. Zhang et al., "Direct-indirect nature of the bandgap in lead-free perovskite nanocrystals," The Journal of Physical Chemistry Letters, vol. 8, no. 14, pp. 3173-3177, 2017. [24] G. Nagabhushana, R. Shivaramaiah, and A. Navrotsky, "Direct calorimetric verification of thermodynamic instability of lead halide hybrid perovskites," Proceedings of the National Academy of Sciences, vol. 113, no. 28, pp. 7717-7721, 2016. [25] B. Lee et al., "Air-stable molecular semiconducting iodosalts for solar cell applications: Cs2SnI6 as a hole conductor," Journal of the American Chemical Society, vol. 136, no. 43, pp. 15379- 15385, 2014. [26] W. Zhu et al., "Deciphering the degradation mechanism of the lead-free all inorganic perovskite Cs2SnI6," npj Materials Degradation, vol. 3, no. 1, p. 7, 2019/02/15 2019, doi: 10.1038/s41529-019-0068-3. [27] X. Han et al., "Lead‐Free Double Perovskite Cs2SnX6: Facile Solution Synthesis and Excellent Stability," Small, vol. 15, no. 39, p. 1901650, 2019. [28] A. Kaltzoglou et al., "Optical-Vibrational Properties of the Cs2SnX6 (X = Cl, Br, I) Defect Perovskites and Hole-Transport Efficiency in Dye-Sensitized Solar Cells," The Journal of Physical Chemistry C, vol. 120, no. 22, pp. 11777-11785, 2016/06/09 2016, doi: 10.1021/acs.jpcc.6b02175. [29] B.-w. Park et al., "Understanding how excess lead iodide precursor improves halide perovskite solar cell performance," Nature communications, vol. 9, no. 1, pp. 1-8, 2018. [30] R. X. Yang, J. M. Skelton, E. L. Da Silva, J. M. Frost, and A. Walsh, "Spontaneous octahedral tilting in the cubic inorganic cesium halide perovskites CsSnX3 and CsPbX3 (X= F, Cl, Br, I)," The journal of physical chemistry letters, vol. 8, no. 19, pp. 4720-4726, 2017. [31] O. Yaffe et al., "Local polar fluctuations in lead halide perovskite crystals," Physical review letters, vol. 118, no. 13, p. 136001, 2017. [32] G. M. Dalpian, X.-G. Zhao, L. Kazmerski, and A. Zunger, "Formation and composition- dependent properties of alloys of cubic halide perovskites," Chemistry of Materials, vol. 31, no. 7, pp. 2497-2506, 2019. [33] N. Pellet et al., "Mixed‐organic‐cation Perovskite photovoltaics for enhanced solar‐light harvesting," Angewandte chemie, vol. 126, no. 12, pp. 3215-3221, 2014. [34] A. Binek, F. C. Hanusch, P. Docampo, and T. Bein, "Stabilization of the trigonal high- temperature phase of formamidinium lead iodide," The journal of physical chemistry letters, vol. 6, no. 7, pp. 1249-1253, 2015.

6-163

[35] M. M. S. Karim et al., "Anion Distribution, Structural Distortion, and Symmetry-Driven Optical Band Gap Bowing in Mixed Halide Cs2SnX6 Vacancy Ordered Double Perovskites," Chemistry of Materials, vol. 31, no. 22, pp. 9430-9444, 2019/11/26 2019, doi: 10.1021/acs.chemmater.9b03267. [36] D. P. McMeekin et al., "A mixed-cation lead mixed-halide perovskite absorber for tandem solar cells," Science, vol. 351, no. 6269, pp. 151-155, 2016. [37] E. A. Tsiwah et al., "One-pot scalable synthesis of all-inorganic perovskite nanocrystals with tunable morphology, composition and photoluminescence," CrystEngComm, 10.1039/C7CE01749A vol. 19, no. 46, pp. 7041-7049, 2017, doi: 10.1039/C7CE01749A. [38] M. Wei et al., "Combining efficiency and stability in mixed tin–lead perovskite solar cells by capping grains with an ultrathin 2D layer," Advanced Materials, vol. 32, no. 12, p. 1907058, 2020. [39] G. Kapil et al., "Strain relaxation and light management in tin–lead perovskite solar cells to achieve high efficiencies," ACS Energy Letters, vol. 4, no. 8, pp. 1991-1998, 2019. [40] T. Leijtens et al., "Tin–lead halide perovskites with improved thermal and air stability for efficient all-perovskite tandem solar cells," Sustainable Energy & Fuels, vol. 2, no. 11, pp. 2450- 2459, 2018. [41] D. Zhao et al., "Low-bandgap mixed tin–lead iodide perovskite absorbers with long carrier lifetimes for all-perovskite tandem solar cells," Nature Energy, vol. 2, no. 4, pp. 1-7, 2017. [42] W. Li, J. Li, J. Li, J. Fan, Y. Mai, and L. Wang, "Addictive-assisted construction of all-inorganic CsSnIBr2 mesoscopic perovskite solar cells with superior thermal stability up to 473 K," Journal of Materials Chemistry A, 10.1039/C6TA08332C vol. 4, no. 43, pp. 17104-17110, 2016, doi: 10.1039/C6TA08332C. [43] S. Nagane et al., "Lead-Free Perovskite Semiconductors Based on Germanium–Tin Solid Solutions: Structural and Optoelectronic Properties," The Journal of Physical Chemistry C, vol. 122, no. 11, pp. 5940-5947, 2018/03/22 2018, doi: 10.1021/acs.jpcc.8b00480. [44] P. Cheng, T. Wu, J. Liu, W.-Q. Deng, and K. Han, "Lead-Free, Two-Dimensional Mixed Germanium and Tin Perovskites," The Journal of Physical Chemistry Letters, vol. 9, no. 10, pp. 2518-2522, 2018/05/17 2018, doi: 10.1021/acs.jpclett.8b00871. [45] J.-L. Xie et al., "New lead-free perovskite Rb7Bi3Cl16 nanocrystals with blue luminescence and excellent moisture-stability," Nanoscale, 10.1039/C9NR00600A vol. 11, no. 14, pp. 6719-6726, 2019, doi: 10.1039/C9NR00600A. [46] A. Swarnkar, V. K. Ravi, and A. Nag, "Beyond Colloidal Cesium Lead Halide Perovskite Nanocrystals: Analogous Metal Halides and Doping," ACS Energy Letters, vol. 2, no. 5, pp. 1089-1098, 2017/05/12 2017, doi: 10.1021/acsenergylett.7b00191. [47] X. Liu et al., "Efficient and stable tin perovskite solar cells enabled by amorphous- polycrystalline structure," Nature Communications, vol. 11, no. 1, p. 2678, 2020/05/29 2020, doi: 10.1038/s41467-020-16561-6. [48] M. Chen et al., "Highly stable and efficient all-inorganic lead-free perovskite solar cells with native-oxide passivation," Nature Communications, vol. 10, no. 1, p. 16, 2019/01/03 2019, doi: 10.1038/s41467-018-07951-y. [49] S. J. Lee et al., "Fabrication of efficient formamidinium tin iodide perovskite solar cells through SnF2–pyrazine complex," Journal of the American Chemical Society, vol. 138, no. 12, pp. 3974- 3977, 2016.

6-164

[50] K. Marshall, M. Walker, R. Walton, and R. Hatton, "Enhanced stability and efficiency in hole- transport-layer-free CsSnI 3 perovskite photovoltaics," Nature Energy, vol. 1, no. 12, pp. 1-9, 2016. [51] T.-B. Song, T. Yokoyama, S. Aramaki, and M. G. Kanatzidis, "Performance enhancement of lead-free tin-based perovskite solar cells with reducing atmosphere-assisted dispersible additive," ACS Energy Letters, vol. 2, no. 4, pp. 897-903, 2017. [52] I. Chung et al., "CsSnI3: semiconductor or metal? High electrical conductivity and strong near- infrared photoluminescence from a single material. High hole mobility and phase-transitions," Journal of the American Chemical Society, vol. 134, no. 20, pp. 8579-8587, 2012. [53] M.-G. Ju et al., "Earth-Abundant Nontoxic Titanium(IV)-based Vacancy-Ordered Double Perovskite Halides with Tunable 1.0 to 1.8 eV Bandgaps for Photovoltaic Applications," ACS Energy Letters, vol. 3, no. 2, pp. 297-304, 2018/02/09 2018, doi: 10.1021/acsenergylett.7b01167. [54] N. Wang et al., "Heterojunction‐Depleted Lead‐Free Perovskite Solar Cells with Coarse‐ Grained B‐γ‐CsSnI3 Thin Films," Advanced Energy Materials, vol. 6, no. 24, p. 1601130, 2016. [55] D. Liu, L. Liang, and R. Sa, "First-principles calculations of structural, electronic, and optical properties of double perovskites Cs2Sn1-xBxI6 (B = Si, Ge; x = 0, 0.25, 0.50, 0.75, 1)," Chemical Physics, vol. 542, p. 111075, 2021/02/01/ 2021, doi: https://doi.org/10.1016/j.chemphys.2020.111075. [56] W. Ke and M. G. Kanatzidis, "Prospects for low-toxicity lead-free perovskite solar cells," Nature Communications, vol. 10, no. 1, p. 965, 2019/02/27 2019, doi: 10.1038/s41467-019-08918-3. [57] S. Li, C. Zhang, J.-J. Song, X. Xie, J.-Q. Meng, and S. Xu, "Metal halide perovskite single crystals: from growth process to application," Crystals, vol. 8, no. 5, p. 220, 2018. [58] Y. Dang, D. Ju, L. Wang, and X. Tao, "Recent progress in the synthesis of hybrid halide perovskite single crystals," CrystEngComm, 10.1039/C6CE00655H vol. 18, no. 24, pp. 4476- 4484, 2016, doi: 10.1039/C6CE00655H. [59] Y. Dang et al., "Formation of Hybrid Perovskite Tin Iodide Single Crystals by Top-Seeded Solution Growth," Angewandte Chemie International Edition, vol. 55, no. 10, pp. 3447-3450, 2016, doi: https://doi.org/10.1002/anie.201511792. [60] E. Mullane, T. Kennedy, H. Geaney, C. Dickinson, and K. M. Ryan, "Synthesis of Tin Catalyzed Silicon and Germanium Nanowires in a Solvent–Vapor System and Optimization of the Seed/Nanowire Interface for Dual Lithium Cycling," Chemistry of Materials, vol. 25, no. 9, pp. 1816-1822, 2013/05/14 2013, doi: 10.1021/cm400367v. [61] C. N. Nanev, "7 - Theory of Nucleation," in Handbook of Crystal Growth (Second Edition), T. Nishinaga Ed. Boston: Elsevier, 2015, pp. 315-358. [62] R. Köferstein, L. Jäger, and S. G. Ebbinghaus, "Magnetic and optical investigations on LaFeO3 powders with different particle sizes and corresponding ceramics," Solid State Ionics, vol. 249, pp. 1-5, 2013. [63] C. Yu et al., "Temperature dependence of the band gap of perovskite semiconductor compound CsSnI3," Journal of Applied Physics, vol. 110, no. 6, p. 063526, 2011, doi: 10.1063/1.3638699. [64] C. C. Stoumpos et al., "Hybrid Germanium Iodide Perovskite Semiconductors: Active Lone Pairs, Structural Distortions, Direct and Indirect Energy Gaps, and Strong Nonlinear Optical Properties," Journal of the American Chemical Society, vol. 137, no. 21, pp. 6804-6819, 2015/06/03 2015, doi: 10.1021/jacs.5b01025.

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Chapter 7 Additive Effects on the Growth of Inorganic Caesium

Germanium Bromide Perovskite

7.1 Abstract

Inorganic germanium perovskite nanomaterials are a compelling class of materials due to the benign nature of the germanium atom compared to lead, paving the way for non-toxic yet high- performance materials with outstanding photophysical properties, ideal for next-generation optical devices. This work looks to replace lead with germanium in CsPbBr3 perovskite by examining different reaction techniques and surfactants. The study shows that utilising germanium precursors in the conventional hot injection synthesis method produces alkali halide NCs while the LARP method produces nanorods that have previously only formed through a template approach. Therefore, an aqueous, low pH system employing hypophosphorous acid as a reducing agent enables the formation of germanium based perovskite that proves stable in the air due to the covalent nature of the internal bonding. Here, by examining a range of surfactant ligand additives and triblock co-polymer templates, we manipulate inorganic germanium halide perovskites' size and morphology. However, restriction to the nanoscale is unsuccessful.

7.2 Introduction

Lead halide perovskites have attracted intense research over the past decade due to their potential in optoelectronic devices with photovoltaic efficiencies topping 25% in 2019. [1-3]

In particular, nanocrystals (NCs) of CsPbX3 (X = Cl, Br, I) display outstanding photo-physical properties such as high photoluminescence quantum yield (PLQY), which sets them above traditional semiconducting nanomaterials. [4, 5] There are many procedures reported that examine the synthesis of different shapes of CsPbX3 NCs with focused size distributions. [6-

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12] However, despite the superior optoelectronic properties and long carrier diffusion lengths, stability and toxicity are problematic as the ionic bond between the lead and halogen atoms is easily broken, leading to entry into the soil and groundwater. [13-16]

The metal-halogen bond can be stabilised by substitution of lead with atoms higher in group

14 due to the reduced electronegativity of these elements increasing the bond covalency. While tin perovskites have shown >10% power conversion efficiencies (PCE) in solar cells, they suffer from a phase transition to the “yellow phase” (non-perovskite structure) at room temperature. However, germanium perovskites maintain their structural phase. [17, 18] As such, this project examines the synthesis of germanium based perovskite NCs.

The ligand layer on the NC is a critical parameter for controlling the dispersibility of NCs in a solvent, their effective volume, adhesion to substrates, functionalization, and surface charge.

[19, 20] Conventionally, CsPbX3 NCs are synthesised in organic systems using non- coordinating solvents, employing a mixture of primary amines and carboxylic acids to solvate the PbX2 precursor, inhibit aggregation, and limit the growth of the NCs. [20, 21] These synthesis conditions transfer to tin perovskites with relatively few adjustments; however, germanium (II) halide precursors oxidise Ge(II) to Ge(IV) and prevents the formation of the perovskite structure. Instead, it forms monodisperse caesium halide (CsX) NCs and germanium oxide. One such product, caesium iodide, is a scintillation material with applications in high- energy X-ray detectors.

An alternative synthesis strategy is therefore needed to form a germanium based perovskite

NCs. Aqueous phase synthesis can incorporate the acidic reducing agent (RA) known to induce and preserve germanium's bivalency: hypophosphorous acid (HPA). Within this polar system, the precursors are soluble, and so the sole role of surfactant is the manipulation of crystal growth. This study examines the ligand types that have exerted control over perovskite NC

7-167 formation and stability in literature, going beyond the CsPbX3 system to hybrid organometallic

NC synthesis and the passivation of perovskite thin films. The study also expands to ligands used in the synthesis of non-perovskite, germanium based NCs of both oxidation states: Ge(IV) and Ge(II). This work aims to use surfactant additives within this reaction system to restrict germanium perovskite morphology to the nanoscale.

Successful surfactants that control perovskite morphology include the conventional, long-chain organic amines used in the hot injection synthesis become hydrophilic on reaction with concentrated hydrohalic acid by forming an ammonium salt: oleylammonium halide.

Interestingly, the main application of oleic acid in the hot injection organic synthesis is to protonate oleylamine (OLA), and so the acidic solvent fulfils this role. [6-8, 22, 23] Therefore, usually hydrophobic amines are applicable in an aqueous environment by reacting with a strong acid and used in an attempt to form NPs.

Like their interaction with a Pb atom, lone pairs of some atoms (e.g. S and N atoms) interact strongly with germanium ions, making them promising surfactant candidates. [24] While short- chain primary amines such as methylammonium and formamidinium take the place of the A cation, longer chain amines can cause the formation of 2D, Ruddlesden-Popper perovskite. [25,

26] Oleylamine is a primary fatty amine as it has C18 hydrocarbons and does not change the formation of the 3D perovskite crystal but still halts the growth to the nanoscale. This work thoroughly analyses the effect of amine chains.

A benefit of ligand bulkiness is the steric effects and the formation of a protective layer, which hinders water molecules' absorption. [27, 28] Work by Park et al. showed that less bulky ammonium salts surround the surface more efficiently than bulky counterparts, while chain length affects the electrical properties due to the organic tail's insulating nature. Quaternary ammonium salts (QAS) passivate the NC surface through concomitant passivation of

7-168 uncoordinated Pb atoms with bromide and the strong affinity of the ammonium-cation to the surface A cation site leading to passivation and near-unity PLQY. [29] This passivation shows the benefit of bifunctional additives, which interact with two sites on the NCs. Ammonium functional groups are critical as they directly link with the surface A cation sites of the crystal; however, other functional groups are needed to connect with other sites. Amino-thiol and amino-silane are two examples of bifunctional additives that have improved surface passivation of perovskite crystals. [30-38]

Many surfactants families are known to interact with perovskite surfaces and limit the crystal growth to the nanoscale. Specifically, Men et al. successfully employed cysteamine to produce polydisperse NCs of germanium perovskite, reasoning that it is the short-chain and the disulphide bond that made it successfully limit the growth. [38] Hence, this chapter examines the effects of short-chain amines before increasing to longer chains and expanding to a range of sulphur, phosphorous and silicon-based additives to limit CsGeBr3 growth to the nanoscale.

This study systematically analyses the morphological effect on the perovskite.

7.3 Experimental Section

Materials: (3-Mercaptopropyl)Trimethoxysilane (MPTMS), Caesium Bromide (CsBr,

99.99%), Caesium Iodide (CsI, 99.999%), Cesium Carbonate (Cs2CO3),

Cetyltrimethylammonium Bromide (CTAB), Cyanamide, Cysteamine,

Didecyldimethylammonium Bromide (DDAB), Dioctylamine, Diphenylsilanediol (DPSO),

Dodecanethiol (DDT), Dodecylphosphoic Acid (DDPA), Ethanolamine, Germanium (II) iodide (GeI2), Germanium (IV) Oxide (GeO2, 99.99%), Gold (III) Chloride (AuCl3),

Hexamethyldisilazane (HMDS), Hydriodic Acid (HI, 57%, Stabilized, 99.95%), Hydrobromic

Acid (HBr, 48%), Hypophosphorous Acid (HPA, H3PO2, 50%), Isobutylamine,

Methylphosphonic Acid (MPA), Octadecane, Octadecanethiol (ODT) Octadecylsilane 7-169

(ODS), Octylamine, Oleic acid, Oleylamine (OLA), Polyvinylpyrrolidone (PVP),

Tetraoctylammonium Bromide (TOA), Thiourea, Tributylphosphine Oxide (TBPO),

Trioctylamine, Trioctylphosphine Oxide (TOPO) and Tripropylamine were purchased from

Sigma Aldrich/Merck. Pluronic® 25R4, P85, F127, F88 were purchased from BASF. Isopropyl

Alcohol (IPA, 99.7%) and Toluene were purchased from Lennox.

Synthesis: Hot injection Synthesis of CsX Preparation of Cs-oleate: 0.407 g of Cs2CO3 (2.5 mmol, Aldrich, 99.9%) was loaded into a 50 mL 3-neck flask along with 20 mL of octadecene

(ODE) (ODE, Sigma-Aldrich, 90%) and 1.55 mL of oleic acid (5 mmol, oleic acid, Sigma-

Aldrich, 90%), dried under vacuum for 1 h at 120 °C, and then heated under Ar to 150 °C until all Cs2CO3 dissolved. Since Cs-oleate is insoluble in octadecene at room temperature, it must be preheated to 100 °C before injection. Final concentration: 0.116 M.

2+ 3+ CsX synthesis, 0.5 mmol MX2 (M= Ge or Au ; X = I, Cl) is dried in 6ml octadecene (ODE) under vacuum for 1 h. 0.2 ml of each ligand (OA and OAm) are added under argon, and the system temperature is raised to 180 °C. 0.4 ml of Cs-oleate solution is then swiftly injected and the reaction allowed to proceed for 5 minutes before cooling with a water bath. The produce is collected through centrifuging for 5 minutes at 7,000 rpm and washed with toluene twice before suspending in toluene.

Aqueous Synthesis of CsGeBr3: 0.1 mmol of additive was dissolved in 5 ml of hydrohalic acid with stirring. This solution was then added to a stirring solution of germanium (IV) oxide

(104 mg, 1 mmol), hydrohalic acid (3 ml) and HPA (2 ml) in a 3-neck flask under argon. These were then stirred at 120 °C until solubilised. CsX (0.75 mmol) was dissolved in 1.5 ml hydrohalic acid in a separate three-necked flask under vacuum until the CO2 and H2O evolution ceased. This room-temperature solution was then injected into the germanium solution at 120

°C, and then the flask was immediately immersed in a water bath. Black (CsGeI3) or orange

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(CsGeBr3) precipitate forms as the system cools. The precipitate was collected by centrifugation for 5 min at 4500 rpm and washed with isopropanol to remove the acid and stored in IPA.

7.4 Results

Two main solution-phase routes for the synthesis of lead halide NCs are the hot-injection method, and the ligand assisted reprecipitation method (LARP). Here, germanium based NC perovskite is attempted using a hot injection synthesis. This methodology is based on

Protesescu and Jellicoe's work to synthesise lead and tin perovskite NCs respectively.[12, 39]

Subsequently, the LARP method is applied, based on the work by Huang et al.[40] Figure 7-1 characterises both methods.

The synthesis's key factor is the metal precursor's stability as CsPbI3 perovskite forms when using a divalent, stable precursor such as PbX2. However, unstable B cations result in the formation of the alkali halide as shown by using Ge2+ and further tested by introducing Au3+.

The XRD patterns in Figure 7-1a represent the hot injection, organic syntheses using GeI2 and

AuCl3 as unstable B cation precursors. Injection of preheated Cs-oleate (120 °C) into a metal halide solution with ligands at 180 °C nucleates and grows the material before quenching the reaction with an ice bath. The red pattern shows cubic CsCl formed when using a gold chloride precursor, AuCl3, while the black pattern shows cubic CsI developed using germanium iodide,

GeI2. The pure phase CsCl shows an absence of impurity phases as the sample is centrifuged at 5000 rpm, the minimum speed at which precipitate separates from the nascent solution.

Meanwhile, the pattern of CsI contains an impurity peak (Figure 7-1a red) as the nascent solution is centrifuged at the highest speed (13,500 rpm). The XRD patterns in Figure 7-1a black detects the main (101) peak of hexagonal GeO2.

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The optical properties of CsI produced through hot injection of Cs into GeI2 solution are shown in Figure 7-1b. Caesium halides are insulator materials and should not show a bandgap above

200 nm. The bandgap of bulk CsI is 6.4 eV, so the onset of absorption for CsI would be about

195 nm.[41] The hexagonal crystalline quartz phase, c-GeO2, BG=2.77 eV is the broad, shallow peak located at 352 nm, while the double peak at 267 nm and 280 nm are characteristic of CsI in thin films, attributed to F-centres in the bandgap resultant from an anionic vacancy occupied by one or more unpaired electron. [42, 43] The Tauc plot inset shows the energy of light hv vs the quantity of light energy absorbed by (αhν)1/2, where α is the material absorption coefficient. From this plot, the bandgap of the CsI is 3.43 eV which is close to the theoretical value of bulk CsI (3.84 eV). A reason for the decrease is that as CsI ages, the absorption peak broadens, and so the slope will broaden to lower energies. [44] CsI is highly ionic with a low dielectric constant, leading to large binding energies of the excitons.[41] The absorption band is independent of NC size due to the strong exciton binding energy and small exciton Bohr radius. Therefore quantum confinement has a negligible effect on CsI.

Germanium iodide (II), GeI2, and oxidises to Ge(IV) in the hot injection synthesis. Particles of

GeO2, identified in the UV-Vis spectra at 360 nm, are not found on the TEM grid. The particles shown in the TEM image of Figure 7-1c are all identified as CsI with no germanium oxide material, possibly due to precipitation from suspension before the sample is drop-cast onto the

TEM grid. Oleic acid forms an oleate attachment with the A cation, caesium, while the oleylamine forms an oleylammonium-to-halide attachment, leaving the germanium oxide without ligands. FFT of a single crystal identifies cubic CsI NCs in Figure 7-1c, as does the

XRD pattern.

In the hot injection of perovskite NCs that utilise octadecene (ODE) as a non-coordinating solvent, oleic acid (OA) and oleylamine (OAm) act as both solubilizers of the metal halide and

7-172 ligands for the NCs. Therefore a germanium stabilizer or reducing agent is needed.

+ Trioctylphosphine (TOP) maintains the 2 divalent state of tin for colloidal CsSnX3 NCs and dissolves PbCl2 for the synthesis of CsPbCl3 NCs. Figure 7-1d shows the NCs produced when

TOP is incorporated with the metal halide precursor, showing hexagonally shaped, structurally cubic CsI NCs. The CsI hexagons are up to 10x larger with TOP inclusion, suggesting that

TOP may increase the reaction's kinetics but does not maintain the Ge2+ divalent state.

Figure 7-1 Caesium iodide formed from hot injection organic synthesis. a) XRD of cubic caesium iodide (black) and caesium chloride (red). b) UV-Vis absorption plot of the sample showing a peak at 352 nm characteristic of GeO2 and a double peak at 267 nm and 280 nm characteristic of CsI c) TEM image and FFT inset of CsI made from hot injection synthesis d) TEM image of hot injection synthesis incorporating TOP RA, e) antisolvent precipitation in DMF f) antisolvent precipitation with RA. All TEM scales are 200 nm.

As the hot injection method does not produce NCs of CsGeBr3, the Ligand Assisted

Reprecipitation (LARP) synthesis method is employed as no thermal energy is used. Here, the precursors are solubilised together and then introduced dropwise into an antisolvent (toluene) and ligand solution. Figure 7-1e+f shows the TEM characterisation. The CsI is again in the cubic phase but has formed 1D short nanorods. The aspect ratio of these particles is consistently in the realm of 2:1 or 3:1. TOP's addition increases the aspect ratio, significantly narrowing the

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NRs with likely oriented attachment at the short edge. Again the FFT pattern matches well with the cubic phase of the CsI material. The only previous reports of the growth of 1D CsI nanocolumns involve an AlO2 template and high temperatures, while this antisolvent method requires no heating.

HPA is a phosphorous oxyacid and strong RA, known to reduce Ge4+ to Ge2+ effectively. It has been used in an aqueous, low pH synthesis, resulting in germanium-based perovskites per reports by Men et al. and Stoumpos et al.[26, 38, 45] In this work, the solvent employed is the desired perovskite halide's hydrohalic acid. For example, in CsGeBr3, HBr is used as both the solvent and the Br source. The RA, GeO2 and HBr form a solution, and the system is heated to

120 °C on a Schlenk line using a condenser. After injection, the heat is removed, and the flask is slowly cooled to room temperature, forming the perovskite material.

Figure 7-2 a-c) XRD patterns of cubic CsGeCl3, rhombohedral CsGeBr3 and hexagonal CsGeI3 synthesised using HPA RA and d-f) the corresponding SEM images with 20 μm scale bars

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After the caesium precursor injection, no precipitate of CsGeBr3 forms until the system cools below 70 °C. For the CsGeI3, precipitate forms immediately on injection, while for CsGeCl3,

70 °C is used as the injection temperature as no precipitate forms at all when the injection temperature is 120 °C. The precipitate is white in the case of cubic CsGeCl3, bright orange in the case of rhombohedral CsGeBr3 and black in the case of hexagonal CsGeI3. Figure 7-2 shows the XRD patterns and SEM images. The SEM image of CsGeBr3 shows the 1D aggregation and attachment of the particles.

Investigation of the solvent medium for CsX showed that CsGeI3 and CsGeBr3 form when using H2O, but CsGeCl3 does not form. Replacing the H2O with just hydrochloric acid does create the perovskite. This formation in acidic conditions may be due to a higher solubility of

CsGeCl3 in H2O than the other perovskites. As such, hydrohalic acid is the solvent in all cases.

7.4.1 Ammonium Additives

Ammonium ligands constitute the main shell around perovskite NCs in conventional syntheses.[21, 40, 46-48] While acidic ligands (e.g. oleate) bind to the A cation, their primary function is the protonation of oleylamine to octylammonium. This octylammonium subsequently attaches to the halide anion and forms the bulk of the ligand shell around the growing perovskite crystals as there are three halide atoms to every A cation. Usually, these long-chain amines cannot dissolve in aqueous environments; however, they are protonated and slowly dissolve exothermically in concentrated hydrohalic acids. As such, an array of amines, primary and quaternary amine solutions in HBr are introduced into the germanium precursor system to limit the nanoscale's growth. Figure 7-3 shows the results below.

Starting with cyanamide, CH2N2, and increasing the chain length to tetraoctylammonium bromide C32H68NBr allows for a complete examination of the effect of amine length. Other

7-175 factors considered include the primary, secondary and tertiary nature of the amines, using octylamine, C8H19N, dioctylamine, C16H35N, and trioctylamine, C24H51N. Secondary functional groups are examined here and throughout the work, for example, ethanolamine,

C2H7NO containing an alcohol group and cysteamine, C2H7NS containing sulphur and HMDS, containing silicon.

Cyanamide is the shortest amine investigates. Its addition to the synthesis results in a significant increase in particle dimensions (Figure 7-3a). The crystals grow to large cuboid particles of over 100 µm in length with widths between 20-30 µm. Ethanolamine, with an alcohol functional group, produces crystals that are larger on average than those without an additive

(Figure 7-3b). However, the crystals are less defined, even following multiple IPA washes.

The material formed using tripropylamine shows particles very similar in size to the pure sample, without any additive shown in Figure 7-3l, but with reduced facet clarity. There are a lot of smaller growths in this SEM image. XRD shows that this is one of the only samples where a secondary phase of CsBr is present, and so this small scale growth is unlikely to be perovskite material. Isobutylamine additive into the reaction system produces CsGeBr3 crystals with no change compared to the additive-less sample.

Octylamine, dioctylamine and trioctylamine are the clearest examinations of the effect of primary, secondary and tertiary amine types (Figure 7-3e-g). These have an extra carbon chain in each additive. Interestingly, these three perovskite samples are quite different from the perovskite made with other ligands. Each of the three additives results in flat sheets of material.

The octylamine additive results in the smallest sheets, measuring less than 20 µm at the smallest sections. When converted to octylammonium halide, octylamine can fill the A cation space in place of some or all Cs+ atoms, which causes it to act as a spacer between the octahedra layers.

[49]

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The perovskite sheets made when incorporating dioctylamine are much larger, over 50 µm wide in places. Trioctylamine results in samples of increased diameter also, but here one can see the sheets' edges, which measure between 0.5 – 1 µm. The thinness of the material suggests that the octylamine materials causes a 2D growth. It is common to see 2D, ruddlesden popper

CsPbX3 perovskites synthesised using amines such as n-butylammonium (n-BA) and 2- phenylethylammonium (PEA). [50] However, there is no report of using octylamine as an additive or for germanium-based perovskites.

In each of the XRD patterns shown in Figure 7-3m, CsGeBr3 peaks are present, with CsBr present in the patterns of the materials made with trioctylamine and tripropylamine additives.

All materials show the rhombohedral phase of CsGeBr3 with some instances of cubic CsBr

(ICDD: 98-006-3021) with peaks at 29.4, 42.0 and 52.1 degrees. This CsBr impurity is likely due to competition from the additive in the perovskite lattice, leaving considerably more A cation atoms unconsumed. There is an excess of halide in all reactions due to the hydrohalic acid solvent.

The perovskites synthesised with the octylamine and tripropylamine additives were more challenging to synthesize than those with other amine additives. In many cases, AGeBr3 did not form, indicated by the immediate precipitation of colourless precipitate upon injection, which XRD identifies as cubic CsBr. Dissolving the amines in HBr overnight overcame this challenge, allowing the formation of an octylammonium bromide solution precursor. Synthesis of alkylammonium halides for the colloidal synthesis of 2D lead perovskites dissolves them in hydrohalic acid and alcohol. Following recrystallization, they are incorporated into the organic precursor solution. As the method employed here uses hydrohalic acid as the reaction solvent, the alkylammonium halide precursors are not recrystallized but are used directly from the amine-acid solution.

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Figure 7-3 a-k) CsGeBr3 synthesised using amines of increasing chain length and compared to the pure synthesis without any ligand. The * identifies peak positions of CsBr impurity. All scale bars are 20 μm. l) XRD patterns of the perovskites.

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The final primary amine to be examined was the long-chain oleylamine with 18 carbons in a single chain. This addition resulted in a highly crystalline material with the sharpest edges, seen in Figure 7-3h, but with no nanomaterial identified.

Figure 7-3j+k show the SEM images of material formed with QAS additives. The addition of

DDAB and TOAB results in substantial crystalline particles measuring over 100 µm. TOAB forms long, flat particles that reduce in thickness at their edges to a point. These are not sheets as seen when using trioctylamine as broken edges reveal the thickness in the sheets' centre as over 10 µm.

Figure 7-3l shows the XRD patterns of each of these samples. The addition of these quaternary ammonium halide salt additives does not limit the nanoscale's growth nor affect the crystallographic phase as rhombohedral CsGeBr3 forms in all cases.

7.4.2 Phosphorous Additives

Chapter 4 of this thesis investigated phosphorous additives in the synthesis of CsPbBr3.

Phosphonic acids (PAs) bind strongly to divalent cadmium and lead anions and have strong hydrogen bonding to perovskites. [51-53] PA additives have produced colloidal perovskite

NCs that withstood washing in H2O and are suitable for room temperature synthesis. [54, 55]

Therefore, they are employed here to limit the germanium perovskite growth to the nanoscale range.

Wang et al. [20] found that PAs were capable of stabilizing the alpha phase of CsPbI3. α-CsPbI3 was stabilised through washing with trioctylphosphine (TOP), leading to over a month's quantum yield efficiency stability. [56] A long chain PA dramatically enhanced the perovskite stability towards polar solvents, producing an LED with an EQE of 7.74%. [57]

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Figure 7-4 (a-d) CsGeBr3 synthesised using silicon-based additives of increasing chain length and compared to the pure synthesis without any ligand (e). The * identifies peak positions of CsBr impurity. All scale bars are 20 μm. b) XRD patterns of the perovskites. Figure 7-4a+b shows the SEM images of materials incorporating a long and short-chain PA into the reaction system. They both reduce the polydispersity shown in the pure sample. Here all the crystals are ~5 µm, the smaller, more numerous particles shown in Figure 7-4e.

Branched capping additives are a route to improving CsPbBr3 NC stability as they can provide a strong steric effect. [58] By incorporating TOPO into the reaction system, monodisperse

CsPbX3 NCs form at higher temperatures than achievable using the conventional ligands with the further advantage of stability towards washing with ethanol.

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Figure 7-4c shows that the TBPO additive here does not result in defined crystals. TOPO's impact in this system is shown in Figure 7-4d, where the addition of TOPO decreases the polydispersity. The image shows cubic particles measuring 6-11 µm with some larger particles measuring up to 20 µm. These results indicate that phosphorous additives don’t reduce the

CsGeBr3 perovskite to nanoscale dimensions, but TOPO does reduce the polydispersity.

7.4.3 Polymer Additives

Template approaches have established success in perovskite synthesis and the reduction of the perovskite crystal size to the nanoscale. This section incorporates non-ionic triblock copolymers (Pluronic®) in the synthesis of CsGeBr3. These are composed of a hydrophobic chain

(polyoxypropylene) chain surrounded by two hydrophilic chains (polyoxyethylene) of variable length. The names of the Pluronic materials identify the state at room temperature (p=paste,

F=flake/solid), with the first one or two digits multiplied by 300, gives the approximate molecular weight, and the last digit multiplied by 10 gives the percentage polyoxyethylene content (e.g. F88 indicates a flaky polyoxypropylene molecular mass of 2400 g/mol and an

80% polyethene oxide content). At low concentration and temperature, individual block copolymers are present. As the concentration increases above the critical micelle concentration

(CMC) or the temperature increases above the critical micelle temperature (CMT), the individual monomers aggregate in a “micellization” process. This aggregation minimises the less soluble PPO blocks' interaction with the solvent, leading to a core of insoluble polyoxypropylene with the soluble polyoxyethylene forming the shell. Micelles can take the shape of spheres or become elongated, depending on polymer composition. [59, 60] These formed nanowires of perovskites for various applications. [61-63]

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Figure 7-5 (a-e) CsGeBr3 synthesised using Pluronic® polymers to increase chain length and the pure synthesis without any ligand. All scale bars are 20 μm unless otherwise stated. f) XRD patterns of the perovskites.

Figure 7-5 shows that Pluronic® polymers affect the formation of CsGeBr3. Pluronic 25R4 additive causes CsGeBr3 growth to form large particles with no defined shape. F127 addition creates 1D cuboid structures over the entire sample, as shown by the inset in Figure 7-5b. These particles vary in length over 10-100s of micrometres, often reaching 200 µm, with a consistent width between 10-12 µm, meaning that they have an aspect ratio well over 1:10. F88 addition

7-182 reduced the polydispersity compared with no additive. Regarding the study's aim, F88 results in small, less polydisperse CsGeBr3.

PVP is not a triblock copolymer but is a template for the growth of various types of NCs. Using

PVP as the additive reduces the particle size, as seen in Figure 7-5d; however, it does not achieve significant morphological control and reduces the crystal quality.

7.4.4 Silicon Additives

The use of silicon-based materials as surfactant is beneficial over some organic polymers as, under ultraviolet (UV) light, these polymers have relatively poor stability and lose their ability against photo-oxidation after a short time. Inorganic silica shells have a low diffusion rate of atoms or ions, as do titania and alumina, making them highly stable under these conditions.

These inorganic shells protect organic LEDs and phosphors from moisture/oxygen degradation. [64] Encapsulated perovskite NCs using silica shells result in improved stability and reducing the anion exchange in mixed perovskites. There are no reports of silicon-based surfactants controlling the morphology of the perovskite NCs.

(3-mercaptopropyl)trimethoxysilane (MPTMS) (Figure 7-6a) produces water-resistant perovskite@silca NCs that maintain luminescence in water for over six weeks. [65] When applied here, the silica appears not to have significantly impacted the large perovskite structure's growth. However, there are many more small particles coating the surface. It is interesting to note the significant proportion of cubic CsBr present in this material's XRD pattern. From this analysis, the smaller growth seen in the SEM image is likely CsBr.

Hexamethyldisilazane (HMDS) has shown promise as a stability agent in perovskite synthesis.

It has hydrophobic -CH3 functional groups that modify the surface and tune the lead perovskite's morphology. Its addition forms large particles with no defined shape.

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Figure 7-6 (a-d) CsGeBr3 synthesised using silicon-based additives of increasing chain length and compared to the pure synthesis without any ligand (e). All scale bars are 50 μm. b) XRD patterns of the perovskites. Diphenylsilanediol (DPSD) is another encapsulation material that increases thermal conductivity, stability and transmittance of white LEDs based on lead halide perovskites. [66]

The SEM image in Figure 7-6c shows a material with crystals of over 20 µm diameter, larger than the size of the crystals that form without any additive; however, the shapes are not clear.

The addition of Octadecylsilane (ODS) forms a similarly sized material (Figure 7-6d); however, the shape is inconsistent and unclear again. Figure 7-6f show XRD patterns of

CsGeBr3 in all cases, yet there is no evidence of size restriction to the nanoscale.

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7.4.5 Sulphur Additives

Sulphur additives are the last additive group investigated. Thiourea is a promising addition as it has two amino groups and one S-donor, of which the amino group can form hydrogen bonds with the halide, and the S-donors can form Lewis acid-base adducts with Ge atoms in the perovskite.[67] Figure 7-7a shows that thiourea addition affects the morphology, reducing the polydispersity of CsGeBr3. However, the particles are still 10s of µm in size, so the thiourea is not enough to halt the nanoscale's growth.

Cysteamine previously limited the growth of germanium perovskite to the nanoscale.[38]

Unfortunately, it had no significant morphological restriction here, as shown by the large,

26µm particles shown in Figure 7-7b.

Alkylthiol ligands induce phase transformation from CsPbBr3 to CsPb2Br5 and morphologically control perovskite growth, influencing change from nanosheets to nanowires when used with alkyl-amines and alkyl-acids. [68] Also, post-synthetic treatment of CsPbBr3

NCs with dodecanethiol (DDT) yields highly stable NCs with near-unity PL due to thioether reaction with the organic solvents and ligands in the system, which binds to Pb atoms on the surface. Further, DDT stabilises Ge's surface through the formation of self-assembled monomers on its surface, following removal of oxide using concentrated hydrohalic acid. [69]

DDT and octadecanethiol (ODT) are therefore implemented as additives in the CsGeBr3 synthesis yielding significant effects on perovskite growth. DDT limited the majority of the growth to 5 μm sized particles, and ODT reduced the particles further to between 1-3 μm, as shown in Figure 7-7c. XRD shows no CsBr impurity in the sample, so these are the smallest

CsGeBr3 crystals achieved in this additive study.

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Figure 7-7 a-e)CsGeBr3 synthesised using sulphur additives of increasing chain length and pure synthesis without any ligand (e). All scale bars are 20 μm. f) XRD patterns of the perovskites. 7.5 Discussion

Table 7-1 consolidates the results below. The colour separates the results according to the additive group incorporated. Each synthesis injects 2 mmol of CsBr dissolved in HBr into 2 mmol of GeO2 dissolved in HBr and H3PO2 with 1 mmol of the appropriate additive. This additive concentration is above the critical micelle concentration (CMC) for all materials in this 5 ml water volume. The carbon chain length of the additives increases downwards in each section.

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Table 7-1 Synthesis parameters and crystallographic properties, colour coded according to additive type. All sizes are in µm.

In all cases, the material formed is identified as rhombohedral CsGeBr3 by XRD. The impurity cubic CsBr occurs with no additive and many amines, polymer templates and (3- mercaptopropyl)trimethoxysilane. The last four columns of the table describe the dimensions of the CsGeBr3. 3D particles are sized according to the smallest and largest crystals in the sample in Size W1 and Size W2, respectively. Size L describes the long length of the 1D particles or the thickness of the 2D sheets.

From this table, the amines and polymer additives result in the most CsBr impurity. The majority of materials formed with the octylamine group of amines as additives result in 2D growth, while cyanimide, TOAB, Pluronic F127, and silicon additives result in the formation of large 1D CsGeBr3 cuboids.

Amine additives produce nanoscale perovskites; however, in this synthesis, the amino group does not significantly impact the size of CsGeBr3. The most successful additives, which reduced both size and polydispersity, is the sulphur group's long-chain thiols and Pluronic F88.

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While the aim of this work was to produce nanoscale CsGeBr3, the microstructures produced here show that additives exert morphological manipulation on the material produced. A future study could examine the electrical characteristics of this material through contact with ITO coated glass.

7.6 Conclusion

The ligand assisted re-precipitation synthesis using unstable metal precursors is a highly controllable and malleable way to form alkali halides with control over size and dimensionality, producing nanorods without a template for the first time. Focusing on the formation of nanoscale germanium perovskite, an HPA reducing agent is necessary to keep the Ge2+ for perovskite formation with its application resulting in the formation of CsGeX3 (X = Cl, Br, I).

While CsGeBr3 and CsGeI3 form on cooling from 120 °C, CsGeCl3 forms on heating from room temperature in a one-pot synthesis and cannot withstand temperatures as high as 120 °C.

Following these discoveries, additives are introduced, and their effect characterised according to how they affect the growth of CsGeBr3.

The two long-chain alkanethiols employed reduced the polydispersity of the CsGeBr3 more than the other 25 additives. Pluronic F88 formed crystals between 1.5-4 µm, and octadecanethiol reduced the size to between 1-3 µm. While the addition of surfactants did not achieve restriction to these germanium perovskite materials' nanoscale, study of the electrical characteristics of the CsGeBr3 microstructures will be an interesting future study.

7.7 Bibliography

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[3] N. R. E. Laboratory, "Best research-cell efficiencies," P. E. Chart, Ed., ed, 2019. [4] G. Li, J. Huang, H. Zhu, Y. Li, J.-X. Tang, and Y. Jiang, "Surface Ligand Engineering for Near-Unity Quantum Yield Inorganic Halide Perovskite QDs and High-Performance QLEDs," Chemistry of Materials, vol. 30, no. 17, pp. 6099-6107, 2018/09/11 2018, doi: 10.1021/acs.chemmater.8b02544. [5] Y. Tan et al., "Highly Luminescent and Stable Perovskite Nanocrystals with Octylphosphonic Acid as a Ligand for Efficient Light-Emitting Diodes," ACS Applied Materials & Interfaces, vol. 10, no. 4, pp. 3784-3792, 2018/01/31 2018, doi: 10.1021/acsami.7b17166. [6] S. Aharon and L. Etgar, "Two Dimensional Organometal Halide Perovskite Nanorods with Tunable Optical Properties," Nano Letters, vol. 16, no. 5, pp. 3230-3235, 2016/05/11 2016, doi: 10.1021/acs.nanolett.6b00665. [7] Q. A. Akkerman et al., "Tuning the Optical Properties of Cesium Lead Halide Perovskite Nanocrystals by Anion Exchange Reactions," Journal of the American Chemical Society, vol. 137, no. 32, pp. 10276-10281, 2015/08/19 2015, doi: 10.1021/jacs.5b05602. [8] Q. A. Akkerman et al., "Solution Synthesis Approach to Colloidal Cesium Lead Halide Perovskite Nanoplatelets with Monolayer-Level Thickness Control," Journal of the American Chemical Society, vol. 138, no. 3, pp. 1010-1016, 2016/01/27 2016, doi: 10.1021/jacs.5b12124. [9] X. Chen, L. Peng, K. Huang, Z. Shi, R. Xie, and W. Yang, "Non-injection gram-scale synthesis of cesium lead halide perovskite quantum dots with controllable size and composition," Nano Research, journal article vol. 9, no. 7, pp. 1994-2006, July 01 2016, doi: 10.1007/s12274-016- 1090-1. [10] H. Deng et al., "Growth, patterning and alignment of organolead iodide perovskite nanowires for optoelectronic devices," Nanoscale, 10.1039/C4NR06982J vol. 7, no. 9, pp. 4163-4170, 2015, doi: 10.1039/C4NR06982J. [11] T. Paul, B. K. Chatterjee, N. Besra, S. Thakur, S. Sarkar, and K. K. Chattopadhyay, "Fabrication of all-inorganic CsPbBr3 perovskite nanocubes for enhanced green photoluminescence," Materials Today: Proceedings, vol. 5, no. 1, Part 2, pp. 2234-2240, 2018/01/01/ 2018, doi: https://doi.org/10.1016/j.matpr.2017.09.224. [12] L. Protesescu et al., "Nanocrystals of Cesium Lead Halide Perovskites (CsPbX3, X = Cl, Br, and I): Novel Optoelectronic Materials Showing Bright Emission with Wide Color Gamut," Nano Letters, vol. 15, no. 6, pp. 3692-3696, 2015/06/10 2015, doi: 10.1021/nl5048779. [13] F. Giustino and H. J. Snaith, "Toward lead-free perovskite solar cells," ACS Energy Letters, vol. 1, no. 6, pp. 1233-1240, 2016. [14] S. Yang, W. Fu, Z. Zhang, H. Chen, and C.-Z. Li, "Recent advances in perovskite solar cells: efficiency, stability and lead-free perovskite," Journal of Materials Chemistry A, vol. 5, no. 23, pp. 11462-11482, 2017. [15] T. Zheng, J. Wu, D. Xiao, and J. Zhu, "Recent development in lead-free perovskite piezoelectric bulk materials," Progress in materials science, vol. 98, pp. 552-624, 2018. [16] A. Abate, "Perovskite solar cells go lead free," Joule, vol. 1, no. 4, pp. 659-664, 2017. [17] X. Liu et al., "Efficient and stable tin perovskite solar cells enabled by amorphous- polycrystalline structure," Nature Communications, vol. 11, no. 1, p. 2678, 2020/05/29 2020, doi: 10.1038/s41467-020-16561-6.

7-189

[18] C. C. Stoumpos, C. D. Malliakas, and M. G. Kanatzidis, "Semiconducting tin and lead iodide perovskites with organic cations: phase transitions, high mobilities, and near-infrared photoluminescent properties," Inorganic chemistry, vol. 52, no. 15, pp. 9019-9038, 2013. [19] J. T. Kopping and T. E. Patten, "Identification of Acidic Phosphorus-Containing Ligands Involved in the Surface Chemistry of CdSe Nanoparticles Prepared in Tri-N-octylphosphine Oxide Solvents," Journal of the American Chemical Society, vol. 130, no. 17, pp. 5689-5698, 2008/04/01 2008, doi: 10.1021/ja077414d. [20] C. Wang, A. S. R. Chesman, and J. J. Jasieniak, "Stabilizing the cubic perovskite phase of CsPbI3 nanocrystals by using an alkyl phosphinic acid," Chemical Communications, 10.1039/C6CC08282C vol. 53, no. 1, pp. 232-235, 2017, doi: 10.1039/C6CC08282C. [21] J. De Roo et al., "Highly Dynamic Ligand Binding and Light Absorption Coefficient of Cesium Lead Bromide Perovskite Nanocrystals," ACS Nano, vol. 10, no. 2, pp. 2071-2081, 2016/02/23 2016, doi: 10.1021/acsnano.5b06295. [22] G. Almeida et al., "The Phosphine Oxide Route toward Lead Halide Perovskite Nanocrystals," Journal of the American Chemical Society, vol. 140, no. 44, pp. 14878-14886, 2018/11/07 2018, doi: 10.1021/jacs.8b08978. [23] F. McGrath, U. V. Ghorpade, and K. M. Ryan, "Synthesis and dimensional control of CsPbBr3 perovskite nanocrystals using phosphorous based ligands," The Journal of Chemical Physics, vol. 152, no. 17, p. 174702, 2020, doi: 10.1063/1.5128233. [24] T. Li, Y. Pan, Z. Wang, Y. Xia, Y. Chen, and W. Huang, "Additive engineering for highly efficient organic–inorganic halide perovskite solar cells: recent advances and perspectives," Journal of Materials Chemistry A, 10.1039/C7TA01798G vol. 5, no. 25, pp. 12602-12652, 2017, doi: 10.1039/C7TA01798G. [25] M. Kazes, T. Udayabhaskararao, S. Dey, and D. Oron, "Effect of Surface Ligands in Perovskite Nanocrystals: Extending in and Reaching out," Accounts of Chemical Research, vol. 54, no. 6, pp. 1409-1418, 2021/03/16 2021, doi: 10.1021/acs.accounts.0c00712. [26] C. C. Stoumpos et al., "Hybrid Germanium Iodide Perovskite Semiconductors: Active Lone Pairs, Structural Distortions, Direct and Indirect Energy Gaps, and Strong Nonlinear Optical Properties," Journal of the American Chemical Society, vol. 137, no. 21, pp. 6804-6819, 2015/06/03 2015, doi: 10.1021/jacs.5b01025. [27] S. Yang, Y. Wang, P. Liu, Y.-B. Cheng, H. J. Zhao, and H. G. Yang, "Functionalization of perovskite thin films with moisture-tolerant molecules," Nature Energy, vol. 1, no. 2, p. 15016, 2016/01/18 2016, doi: 10.1038/nenergy.2015.16. [28] Z. Zhang et al., "Progress in Multifunctional Molecules for Perovskite Solar Cells," Solar RRL, vol. 4, no. 2, p. 1900248, 2020. [29] Y. Shynkarenko et al., "Direct Synthesis of Quaternary Alkylammonium-Capped Perovskite Nanocrystals for Efficient Blue and Green Light-Emitting Diodes," ACS Energy Letters, vol. 4, no. 11, pp. 2703-2711, 2019/11/08 2019, doi: 10.1021/acsenergylett.9b01915. [30] F. Zhang and M. Srinivasan, "Self-assembled molecular films of aminosilanes and their immobilization capacities," Langmuir, vol. 20, no. 6, pp. 2309-2314, 2004. [31] Y.-Q. Wang, S.-B. Xu, J.-G. Deng, and L.-Z. Gao, "Enhancing the efficiency of planar heterojunction perovskite solar cells via interfacial engineering with 3-aminopropyl trimethoxy silane hydrolysate," Royal Society open science, vol. 4, no. 12, p. 170980, 2017.

7-190

[32] P. Fu, X. Guo, S. Wang, Y. Ye, and C. Li, "Aminosilane as a molecular linker between the electron-transport layer and active layer for efficient inverted polymer solar cells," ACS applied materials & interfaces, vol. 9, no. 15, pp. 13390-13395, 2017. [33] K. Xu et al., "Synergistic Surface Passivation of CH3NH3PbBr3 Perovskite Quantum Dots with Phosphonic Acid and (3-Aminopropyl)triethoxysilane," Chemistry – A European Journal, vol. 25, no. 19, pp. 5014-5021, 2019, doi: 10.1002/chem.201805656. [34] J. Zhu, Y. Zhu, J. Huang, Y. Gong, J. Shen, and C. Li, "Synthesis of CsPbBr3 perovskite nanocrystals with the sole ligand of protonated (3-aminopropyl)triethoxysilane," Journal of Materials Chemistry C, 10.1039/C9TC02089F vol. 7, no. 24, pp. 7201-7206, 2019, doi: 10.1039/C9TC02089F. [35] W. Dai, Y. Lei, M. Xu, P. Zhao, Z. Zhang, and J. Zhou, "Rare-Earth Free Self-Activated Graphene Quantum Dots and Copper-Cysteamine Phosphors for Enhanced White Light-Emitting-Diodes under Single Excitation," Scientific Reports, vol. 7, no. 1, p. 12872, 2017/10/09 2017, doi: 10.1038/s41598-017-13404-1. [36] Z. Salehi, R. V. Parish, and R. G. Pritchard, "A polymeric cationic copper(I) complex involving a quadruply bridging, zwitterionic thiolate ligand: {[Cu8Cl6(SCH2CH2NH3)6]Cl2}," Journal of the Chemical Society, Dalton Transactions, 10.1039/A705217K no. 22, pp. 4241-4246, 1997, doi: 10.1039/A705217K. [37] L. Riauba, G. Niaura, O. Eicher-Lorka, and E. Butkus, "A Study of Cysteamine Ionization in Solution by Raman Spectroscopy and Theoretical Modeling," The Journal of Physical Chemistry A, vol. 110, no. 50, pp. 13394-13404, 2006/12/01 2006, doi: 10.1021/jp063816g. [38] L. Men, B. A. Rosales, N. E. Gentry, S. D. Cady, and J. Vela, "Lead-Free Semiconductors: Soft Chemistry, Dimensionality Control, and Manganese-Doping of Germanium Halide Perovskites," ChemNanoMat, vol. 5, no. 3, pp. 334-339, 2019, doi: 10.1002/cnma.201800497. [39] T. C. Jellicoe et al., "Synthesis and Optical Properties of Lead-Free Cesium Tin Halide Perovskite Nanocrystals," Journal of the American Chemical Society, vol. 138, no. 9, pp. 2941-2944, 2016/03/09 2016, doi: 10.1021/jacs.5b13470. [40] S. Sun, D. Yuan, Y. Xu, A. Wang, and Z. Deng, "Ligand-Mediated Synthesis of Shape-Controlled Cesium Lead Halide Perovskite Nanocrystals via Reprecipitation Process at Room Temperature," ACS Nano, vol. 10, no. 3, pp. 3648-3657, 2016/03/22 2016, doi: 10.1021/acsnano.5b08193. [41] J. Shamsi et al., "Colloidal CsX (X = Cl, Br, I) Nanocrystals and Their Transformation to CsPbX3 Nanocrystals by Cation Exchange," Chemistry of Materials, vol. 30, no. 1, pp. 79-83, 2018/01/09 2018, doi: 10.1021/acs.chemmater.7b04827. [42] B. Walker, C. C. Dharmawardhana, N. Dari, P. Rulis, and W.-Y. Ching, "Electronic structure and optical properties of amorphous GeO2 in comparison to amorphous SiO2," Journal of Non- Crystalline Solids, vol. 428, pp. 176-183, 2015. [43] W. Song, X. Wu, Q. Di, T. Xue, J. Zhu, and Z. Quan, "Morphologically controlled synthesis of ionic cesium iodide colloidal nanocrystals and electron beam-induced transformations," RSC Advances, 10.1039/C8RA02582G vol. 8, no. 33, pp. 18519-18524, 2018, doi: 10.1039/C8RA02582G. [44] A. Farzaneh, M. R. Abdi, K. R. E. Saraee, M. Mostajaboddavati, and A. Quaranta, "Cesium- iodide-based nanocrystal for the detection of ionizing radiation," Optical Materials, vol. 55, pp. 22-26, 2016/05/01/ 2016, doi: https://doi.org/10.1016/j.optmat.2016.03.014.

7-191

[45] M. Chen et al., "Highly stable and efficient all-inorganic lead-free perovskite solar cells with native-oxide passivation," Nature Communications, vol. 10, no. 1, p. 16, 2019/01/03 2019, doi: 10.1038/s41467-018-07951-y. [46] A. Pan et al., "Insight into the Ligand-Mediated Synthesis of Colloidal CsPbBr3 Perovskite Nanocrystals: The Role of Organic Acid, Base, and Cesium Precursors," ACS Nano, vol. 10, no. 8, pp. 7943-7954, 2016/08/23 2016, doi: 10.1021/acsnano.6b03863. [47] A. Heuer-Jungemann et al., "The Role of Ligands in the Chemical Synthesis and Applications of Inorganic Nanoparticles," Chemical Reviews, vol. 119, no. 8, pp. 4819-4880, 2019/04/24 2019, doi: 10.1021/acs.chemrev.8b00733. [48] M. Liu, A. Matuhina, H. Zhang, and P. Vivo, "Advances in the Stability of Halide Perovskite Nanocrystals," (in eng), Materials (Basel), vol. 12, no. 22, p. 3733, 2019, doi: 10.3390/ma12223733. [49] X. Gao et al., "Ruddlesden–Popper Perovskites: Synthesis and Optical Properties for Optoelectronic Applications," Advanced Science, https://doi.org/10.1002/advs.201900941 vol. 6, no. 22, p. 1900941, 2019/11/01 2019, doi: https://doi.org/10.1002/advs.201900941. [50] C. Lan, Z. Zhou, R. Wei, and J. C. Ho, "Two-dimensional perovskite materials: From synthesis to energy-related applications," Materials Today Energy, vol. 11, pp. 61-82, 2019/03/01/ 2019, doi: https://doi.org/10.1016/j.mtener.2018.10.008. [51] J. Y. Woo et al., "Air-Stable PbSe Nanocrystals Passivated by Phosphonic Acids," Journal of the American Chemical Society, vol. 138, no. 3, pp. 876-883, 2016/01/27 2016, doi: 10.1021/jacs.5b10273. [52] X. Li et al., "Improved performance and stability of perovskite solar cells by crystal crosslinking with alkylphosphonic acid ω-ammonium chlorides," Nature Chemistry, Article vol. 7, p. 703, 08/17/online 2015, doi: 10.1038/nchem.2324 https://www.nature.com/articles/nchem.2324#supplementary-information. [53] R. Gomes et al., "Binding of Phosphonic Acids to CdSe Quantum Dots: A Solution NMR Study," The Journal of Physical Chemistry Letters, vol. 2, no. 3, pp. 145-152, 2011/02/03 2011, doi: 10.1021/jz1016729. [54] W.-k. Koh, S. Park, and Y. Ham, "Phosphonic Acid Stabilized Colloidal CsPbX3 (X=Br, I) Perovskite Nanocrystals and Their Surface Chemistry," ChemistrySelect, vol. 1, no. 13, pp. 3479-3482, 2016, doi: doi:10.1002/slct.201600809. [55] T. Xuan et al., "Highly stable CsPbBr3 quantum dots coated with alkyl phosphate for white light-emitting diodes," Nanoscale, 10.1039/C7NR04179A vol. 9, no. 40, pp. 15286-15290, 2017, doi: 10.1039/C7NR04179A. [56] C. Lu et al., "Enhanced stabilization of inorganic cesium lead triiodide (CsPbI3) perovskite quantum dots with tri-octylphosphine," Nano Research, vol. 11, no. 2, pp. 762-768, 2018/02/01 2018, doi: 10.1007/s12274-017-1685-1. [57] A. A. M. Brown et al., "Self-assembly of a robust hydrogen-bonded octylphosphonate network on cesium lead bromide perovskite nanocrystals for light-emitting diodes," Nanoscale, 10.1039/C9NR02566A vol. 11, no. 25, pp. 12370-12380, 2019, doi: 10.1039/C9NR02566A. [58] L. Wu et al., "Improving the Stability and Size Tunability of Cesium Lead Halide Perovskite Nanocrystals Using Trioctylphosphine Oxide as the Capping Ligand," Langmuir, vol. 33, no. 44, pp. 12689-12696, 2017/11/07 2017, doi: 10.1021/acs.langmuir.7b02963.

7-192

[59] P. Alexandridis and T. Alan Hatton, "Poly(ethylene oxide) poly(propylene oxide) poly(ethylene oxide) block copolymer surfactants in aqueous solutions and at interfaces: thermodynamics, structure, dynamics, and modeling," Colloids and Surfaces A: Physicochemical and Engineering Aspects, vol. 96, no. 1, pp. 1-46, 1995/03/10/ 1995, doi: https://doi.org/10.1016/0927-7757(94)03028-X. [60] F. Aydin et al., "Self-Assembly and Critical Aggregation Concentration Measurements of ABA Triblock Copolymers with Varying B Block Types: Model Development, Prediction, and Validation," The Journal of Physical Chemistry B, vol. 120, no. 15, pp. 3666-3676, 2016/04/21 2016, doi: 10.1021/acs.jpcb.5b12594. [61] S. M. Lee, H. Jung, W. I. Park, Y. Lee, E. Koo, and J. Bang, "Preparation of Water-Soluble CsPbBr3 Perovskite Quantum Dot Nanocomposites via Encapsulation into Amphiphilic Copolymers," ChemistrySelect, vol. 3, no. 40, pp. 11320-11325, 2018, doi: 10.1002/slct.201802237. [62] J. Park et al., "Mussel-Inspired Polymer Grafting on CsPbBr3 Perovskite Quantum Dots Enhancing the Environmental Stability," Particle & Particle Systems Characterization, vol. 36, no. 12, p. 1900332, 2019, doi: 10.1002/ppsc.201900332. [63] N. Zhou et al., "Perovskite nanowire–block copolymer composites with digitally programmable polarization anisotropy," Science Advances, vol. 5, no. 5, p. eaav8141, 2019, doi: 10.1126/sciadv.aav8141. [64] F. Zhang et al., "Silica coating enhances the stability of inorganic perovskite nanocrystals for efficient and stable down-conversion in white light-emitting devices," Nanoscale, 10.1039/C8NR07022A vol. 10, no. 43, pp. 20131-20139, 2018, doi: 10.1039/C8NR07022A. [65] S. Li et al., "Water-resistant perovskite nanodots enable robust two-photon lasing in aqueous environment," Nature Communications, vol. 11, no. 1, p. 1192, 2020/03/04 2020, doi: 10.1038/s41467-020-15016-2. [66] P. M. Reddy, C.-J. Chang, C.-F. Lai, M.-J. Su, and M.-H. Tsai, "Improved organic- inorganic/graphene hybrid composite as encapsulant for white LEDs: Role of graphene, titanium (IV) isopropoxide and diphenylsilanediol," Composites Science and Technology, vol. 165, pp. 95-105, 2018/09/08/ 2018, doi: https://doi.org/10.1016/j.compscitech.2018.06.017. [67] S. Wang et al., "High-Performance Perovskite Solar Cells with Large Grain-Size obtained by using the Lewis Acid-Base Adduct of Thiourea," Solar RRL, vol. 2, no. 6, p. 1800034, 2018, doi: https://doi.org/10.1002/solr.201800034. [68] L. Ruan, W. Shen, A. Wang, A. Xiang, and Z. Deng, "Alkyl-Thiol Ligand-Induced Shape- and Crystalline Phase-Controlled Synthesis of Stable Perovskite-Related CsPb2Br5 Nanocrystals at Room Temperature," The Journal of Physical Chemistry Letters, vol. 8, no. 16, pp. 3853-3860, 2017/08/17 2017, doi: 10.1021/acs.jpclett.7b01657. [69] Q. Cai et al., "1-Dodecanethiol based highly stable self-assembled monolayers for germanium passivation," Applied Surface Science, vol. 353, pp. 890-901, 2015/10/30/ 2015, doi: https://doi.org/10.1016/j.apsusc.2015.06.174.

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Chapter 8 Conclusions and Perspectives

8.1 Conclusion and Future Directions from Chapter 4

Chapter 4 presents a synthesis protocol of CsPbBr3 perovskite NCs, facilitating a systematic study of the effects of a phosphorous-based ligand shell on the perovskite's structural and optical properties. Synthesis of CsPbBr3 nanocrystals is achieved using an oleylamine/alkylphosphonic acid combination with near-perfect monodispersity with the ability to tune the bandgap by varying the alkyl chain length. Further, oleylamine is replaced by TOPO, resulting in a ligand combination of alkylphosphonic acid/ Trioctylphosphine oxide

(TOPO) for perovskite nanocrystal nucleation and growth. This combination is typical for the widely studied metal chalcogenide synthesis, and our study with CsPbBr3 yields a pure phase perovskite, unlike previous ligands. However, the material displays a higher degree of polydispersity, as proven in TEM and XRD characterisation. An aspect of using TOPO in the synthesis is the possibility of impurities. Investigating across various batch numbers presents differing size distributions while one specific batch resulted in the formation of nanorods of

Cs4PbBr6. The introduction of a range of impurities into other batches did not replicate these formations.

Further investigations should examine the impurities in TOPO batches and look to their synthesis and subsequent addition to this system. Reproduction of the monodisperse nanorods of CsPbBr3 would enable their directed assembly onto an electrode through electrophoretic deposition, onto patterned films at high deposition rates using an easily scalable apparatus for production. An LED device based on this deposition technique would ideally be assembled on an ITO electrode coated with a buffer hole injection layer. The perovskite would then be coated with a TPBi electron transport layer and a LiF/Al metal top contact. An overcoat of TBTB

8-194 would be beneficial to heal halide vacancy defects. Complete characterisation of luminescence and electroluminescence efficiency of the device should occur.

8.2 Conclusion and Future Directions from Chapter 5

Chapter 5 examined the synthesis of a series of alloyed indium chalcogenide materials composed of nanoribbons and nanosheets without phosphine-based precursors. The concentration of chalcogenide precursor has a significant influence on the morphology of the resultant nanomaterials. Se incorporates into In2S3 up to 25%, and S incorporates into In2Se3 but not into InSe. Mixing of the chalcogenides immediately changes the nano-hexagonal sheets into highly anisotropic nanoribbons that broaden as the concentration of Sulphur increases and the main phase becomes In2S3 with Se alloying. Selenium incorporation into the In2S3 materials reduces the bandgap from 3.19 to 3.04 eV, which is considerably more significant than the 2.1 eV expected for pure In2S3. Further investigation of this synthesis's power to mix the chalcogenide material in In+S/Te and In+Se/Te is then investigated. The materials formed are nanoplatelets with much more agglomeration than the selenium-sulphur alloys, indicating the lesser ability for OLA to fulfil both the solvent and surfactant roles. XRD data shows that some selenium atoms alloy into InTe. However, tellurium does not alloy into InSe. Likewise, there is no alloying of sulphur and tellurium. Rather separate phases of InTe and In2S3 occur.

Future studies should examine the use of the alloyed indium chalcogenide materials as absorber materials for photocatalysis applications. In particular, alloys of indium with selenium and tellurium alloy has been identified as a promising photocatalyst material. Deposition of the alloys onto an electrode such as ITO on glass would enable electrocatalysis investigations. An ideal photocatalysis setup would use triethanolamine as the hole scavenger for photocatalytic

H2 production with 1 wt% Pt co-catalyst. Triethanolamine promotes the separation of

8-195 photoexcited electron/holes at the reduction/oxidation sites when 1 sun, simulated sunlight irradiation is applied. The photocatalytic activity should then be investigated.

8.3 Conclusion and Future Directions from Chapter 6

Chapter 6 investigated the substitution of Sn into two Ge perovskite materials through a solution-based synthesis. SnI2 can be incorporated up to 25% before a breakdown of the ABX3 structure into different phases in the caesium germanium iodide-based materials. SnBr2 incorporates up to 50% into the caesium germanium bromide perovskite. Applications for these materials include photovoltaics for the iodide alloy, specifically the 25% Sn incorporation into CsGeI3, i.e. CsGexSn1-xI (Gex = 0.75), as this material is a pure phase and presents a bandgap of 1.406 eV. There are potential applications for luminescent devices using bromide alloys.

Future work for these materials may include an investigation into the cooling time variation to vary the crystal size and improve polydispersity with an analysis of the effect of the acid concentration, temperature and impact of alcohol etching on the secondary growth effect.

When these synthesis parameters are understood further, solar cell and LED devices should be examined as lead-free perovskite devices would be less harmful to the environment.

8.4 Conclusion and Future Directions from Chapter 7

The ligand assisted re-precipitation synthesis using unstable metal precursors is a highly controllable and malleable way to form alkali halides with control over size and dimensionality, producing nanorods of CsI without a template for the first time. Focusing on the formation of nanoscale germanium perovskite, HPA reducing agent is necessary to keep the Ge2+ for perovskite formation with its application resulting in the formation of CsGeX3 (X = Cl, Br, I).

While CsGeBr3 and CsGeI3 form on cooling from 120 °C, CsGeCl3 forms on heating from 8-196 room temperature in a one-pot synthesis and cannot withstand temperatures as high as 120 °C.

Following these discoveries, additives' effect is characterised with regards to how they affect the growth of CsGeBr3.

The two alkanethiols employed reduced the polydispersity of the CsGeBr3 more than any of the other 25 additives. The size was reduced to between 1-3 µm by 1 mmol of octadecanethiol and 1.5-4 µm by Pluronic F88. However, no nanoscale restriction was achieved for these germanium perovskite materials.

Future investigations should focus on the use of alkyl-halide nanomaterials as scintillation materials. Complete washing away of the ligand layer and the nanomaterial attachment to the end of a fibre optic cable using resin will allow the measuring of the scintillation properties of these materials. The nano-dimensions of these caesium iodide nanorods speeds up the X-ray excited luminescence decay time, leading to fast x- or γ- ray detection.

Future attempts to limit the growth of germanium perovskite to the nanoscale should focus on the temperature reduction process by introducing ace baths or liquid nitrogen to quench the reaction and the surfactant incorporation. Separately, study of the electrical characteristics, such as impedance or conductivity measurements, of micro-CsGeBr3 would be an interesting future direction.

8.5 Answering the research questions

Here the research questions posed in Chapter 1 will be revisited and answered.

Research Question 1: Can the ligand shell of perovskite nanomaterials be changed while the size and optical properties remain unaffected?

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Specifically, no. Phosphorous based ligands can form the ligand shell; however, the luminescence properties are not maintained. The yield dramatically increases when TOPO replaces oleylamine.

Research Question 2: Can indium chalcogenide nanomaterials undergo chalcogenide alloying in a system without trioctylphosphine?

Yes, it is possible to produce alloys of InE (E=S, Se, Te) by dissolving the elemental precursors in oleylamine at an elevated temperature, similar to Top-E preparation but without the use of trioctylphosphine.

Research Question 3: Can lead be substituted for tin and germanium while maintaining the bulk perovskite structure?

Yes. Germanium perovskite alloys with up to 25% Sn in CsGeI3, while in CsGeBr3, up to 50%

Sn can be substituted. As Sn concentrations increases, the structure preferentially forms the

4+ vacancy ordered, double perovskite Cs4SnX6, with oxidised Sn as the only B cation incorporated. This type of perovskite is energetically impossible in the case of Ge.

Research Question 4: Can germanium based perovskites be confined to the nanoscale using surfactants?

No, through systematically studying three NC formation methods and examining five surfactant family systems, morphological variation of CsGeBr3 is possible across 1D, 2D and 3D particles. Polydispersity can be varied and specifically decreased. Octadecanethiol resulted in the smallest perovskite particles with the lowest polydispersity resulting in crystals measuring

1-3 µm. Unfortunately, nanoscale restriction was not achieved.

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8.6 Testing the research hypotheses

This section revisits the research hypotheses.

Hypothesis 1: A different ligand shell on the perovskite nanocrystal surface may better bond with the NC surface than OA and OLA, resulting in preserved luminescence properties.

Yes and no. Chapter 4 produced perovskite NCs with a different shell of phosphonic acids and oleylamine in place of oleic acid. However, the acids take the place of oleic acid and bond with the A cation, not the B cation. The optical properties vary according to the acid chain length; however, the luminescence achieved with oleic acid is not preserved.

Hypothesis 2: Alloying of indium chalcogenide should be possible without using a phosphine- based solvent and still provide a route to optical property control.

Yes, this hypothesis was correct, and oleylamine was used in place of TOP to dissolve the elemental chalcogenide precursors.

Yes, the use of germanium as a B cation did dramatically increase the perovskite's stability in air and over a long period. Alloying of Sn into Ge perovskite allowed tuning of the bandgap of

CsGeI3 to 1.406 eV, lower than that of Pb perovskites and closer to the theoretical ideal 1.33 eV for a solar cell.

Hypothesis 4: An untested type of surfactant might be capable of controlling the formation of germanium based perovskites.

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This hypothesis was incorrect. Surfactants were able to manipulate the degree of polydispersity and the morphology of the material but in no case were any nanoscale materials identified.

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