Halide Semiconductors

Research Collection

Doctoral Thesis

Novel Pb(II), Sn(II) and Bi(III) Halide Semiconductors

Author(s): Nazarenko, Olga

Publication Date: 2019-04

Permanent Link: https://doi.org/10.3929/ethz-b-000338638

Rights / License: In Copyright - Non-Commercial Use Permitted

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ETH Library Dissertation ETH No. 25772

Novel Pb(II), Sn(II) and Bi(III) Halide Semiconductors

A thesis submitted to attain the degree of DOCTOR OF SCIENCES of ETH ZÜRICH (Dr. sc. ETH Zürich)

Presented by Olga Nazarenko Master of Science in Chemistry Kiev National Taras Shevchenko University Born on 26.12.1989 Citizen of Ukraine

Accepted on the recommendation of Prof. Dr. Maksym V. Kovalenko, examiner Prof. Dr. Hansjörg Grützmacher, co-examiner

Department of Chemistry and Applied Biosciences Laboratory of Inorganic Chemistry 2019

Declaration | Olga Nazarenko

Declaration

I hereby confirm that I am the sole author of the written work herein enclosed and that I have compiled it in my own words. Parts excepted are corrections of form and content by the supervisor. Title of work: Novel Pb(II), Sn(II) and Bi(III) halide semiconductors

With my signature I confirm: − I have committed none of the forms of plagiarism. − I have documented all methods, data, and processes truthfully. − I have not manipulated any data. − I have mentioned all persons who were significant contributors to this work, as described below. − I obtained copyright permissions from the journal for reproducing the text and Figures in this thesis, where needed. − I am aware that the work may be screened electronically for plagiarism.

Place, date Signature(s)

Dr. S. Yakunin (ETH Zürich) performed the TR-PL, temperature-dependent PL, and photoconductivity measurements, as well as γ-rays detection measurements presented in Chapters 3-7 as well as contributed to the analysis of the data. Dr. E. Cuervo-Reyes (Empa) performed a computational study included in Chapter 4. Dr. M. D. Wörle (ETH Zürich) helped with the crystal structure determination of the compounds presented throughout this thesis. M. Aebli (ETH Zürich) performed SSNMR measurements and contributed to the analysis of the data presented in Chapter 5. L. Piveteau (ETH Zürich) performed SSNMR measurements and contributed to the analysis of the data presented in Chapters 4 and 7.

B. M. Benin (ETH Zürich) synthesized Cs8Sn6Br13I7 and characterized it with absorption and powder XRD techniques presented in Chapter 6. Dr. G. Raino (ETH Zürich) performed PL and TR-PL with spatial resolution measurements presented in Chapter 5 and temperature-dependent PL measurements included in Chapter 6. Dr. M. R. Kotyrba (ETH Zürich) carried out a part of the SC-XRD measurements presented in Chapters 4, 5 and contributed to the analysis of the results shown in Chapters 4-6. Dr. F. Krumeich (ETH Zürich) conducted SEM, EDXS measurements presented in Chapter 7. Dr. M. Kepenekian (ISCR, CNRS, Rennes), Prof. Dr. J. Even (INSA, CNRS, Rennes), Dr. Sc. C. Katan (ISCR, CNRS, Rennes) conducted a computational study included in Chapter 6. Dr. Y. Polyhach (ETH Zürich) performed EPR measurements for the project presented in Chapter 7.

Declaration | Olga Nazarenko

V. Morad, I. Cherniukh (ETH Zürich) grew some of the single crystals used for the γ-rays detection study, presented in Chapter 3.

Aknowledgements | Olga Nazarenko

Acknowledgments

Countless words of gratitude I would like to address to my supervisor Prof. Maksym V. Kovalenko for giving me an opportunity to work and study in his team, for his support and understanding, his supervision, guidance, and a good sense of humor. It was a pleasure to work and to grow in his group. I want to thank all the group members for creating a warm and pleasant atmosphere in the group. Special thanks to Dr. Loredana Protesescu, Dr. Maryna I. Bodnarchuk and Dr. Dmitry Dirin for teaching me the techniques of the colloidal synthesis when I came to the group. I think this is an excellent tradition in Prof. Kovalenko’s group that a newcomer can be taught all sorts of synthetic techniques used in the group. The transfer of knowledge gives a broader scope of possibilities for a new student to develop his pathway in the world of chemistry. I am grateful for the chance bestowed on me to make doctoral studies in the group of Prof. Maksym Kovalenko at ETH Zurich. The creative and supportive atmosphere at ETH Zurich makes it to a perfect place for effective learning and personal development. I want further to express my gratitude to Dr. Michel D. Wörle, who taught me so much about the X-ray diffraction. I am thankful for his patience, help, and support throughout my Ph.D.. As well as I am grateful to Michael for keeping the Smart in an excellent condition, as on this rather old device I could perform numerous single crystal XRD measurements. Special thanks to Sergii Yakunin for all his amusing life stories during coffee breaks, as well as for his support and collaboration. Through different periods of my Ph.D., I got to work with many people, with various devices, different synthetic techniques. Some of the most challenging things were welding of metal ampules, sealing quartz ampules, working with liquid ammonia. For these incredible skills and knowledge, I am gratified to Dr. Martin R. Kotyrba. I am grateful to my supporting and loving family for believing in me and to my boyfriend Martin for being there for me, for his encouragement, inspiration, funny chemistry related stories, and for bringing balance into my life. I want to thank Claudia Salameh-Ott and Chantal Hänni for support with administrative matters. I thank Prof. Dr. Hansjörg Grützmacher for his kind willingness to co-examine this thesis. Last but not least, I would like to thank the members of the examination committee for their efforts and time and also to you, the reader, for considering my work.

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Table of Contents | Olga Nazarenko

Table of Contents Declaration ...... i Acknowledgments ...... iii Abbreviations ...... vi Summary ...... viii Zusammenfassung ...... xi Chapter 1. Introduction ...... 1 1.1. A brief history of semiconductors ...... 1 1.2. Mixed organic-inorganic lead halide perovskites: basics and properties ...... 3 1.3. Bandgap theory of semiconductors ...... 8 1.4. The defect tolerance ...... 11 1.5. Layered lead halide perovskites ...... 13 1.6. The potential applications of LHPs ...... 16 1.7. On Sn(II) and Ag(I)/Bi(III) halide compounds with perovskite crystal structure ...... 19 1.8. Scope and outline of the dissertation ...... 22 Chapter 2. Methods and techniques ...... 26 2.1. The growth of single crystals ...... 26 2.2. Characterization methods ...... 28 Chapter 3. Single crystals of caesium-formamidinium lead halide perovskites: solution growth and gamma dosimetry ...... 34 3.1. Introduction ...... 34 3.2. Experimental section ...... 36 3.3. Results and discussion ...... 38 3.4. Conclusions ...... 47 Chapter 4. Luminescent and photoconductive layered lead halide perovskite compounds comprising mixtures of caesium and guanidinium cations ...... 49 4.1. Introduction ...... 49 4.2. Experimental section ...... 51 4.3. Results and discussion ...... 55 4.4. Conclusions ...... 69 Chapter 5. Guanidinium-formamidinium lead iodide: a layered perovskite-related compound with red luminescence at room temperature ...... 70 5.1. Introduction ...... 70 5.2. Experimental section ...... 72

Table of Contents | Olga Nazarenko

5.3. Results and discussion ...... 74 5.4. Conclusions ...... 81 Chapter 6. Guanidinium and mixed cesium-guanidinium tin(II) bromides: effects of quantum confinement and out-of-plane octahedral tilting ...... 82 6.1. Introduction ...... 82 6.2. Experimental section ...... 84 6.3. Results and discussion ...... 88 6.4. Conclusions ...... 98 Chapter 7. Rubidium copper and silver iodobismuthates ...... 99 7.1. Introduction ...... 99 7.2. Experimental section ...... 100 7.3. Results and discussion ...... 103 7.4. Conclusions ...... 108 Chapter 8. Conclusions and outlook ...... 110 8.1. Conclusions ...... 110 8.2. Outlook ...... 113 Chapter 9. Appendix...... 115 Appendix to Chapter 3 ...... 115 Appendix to Chapter 4 ...... 119 Appendix to Chapter 5 ...... 136 Appendix to Chapter 6 ...... 144 Appendix to Chapter 7 ...... 157 Literature ...... 161 A creative page ...... 179 Curriculum Vitae ...... 180

Abbreviations | Olga Nazarenko

Abbreviations

CN Coordination number

CCDC The Cambridge Crystallographic Data Centre

CBMin Conduction band minimum

DSC Differential scanning calorimetry

DFT The density functional theory

ELF The electron localization function

EDXS Energy dispersive X-ray spectroscopy

FWHM Full width at half maximum

GTF Goldschmidt tolerance factor

ICSD The Inorganic Crystal Structure Database

LHP Lead halide perovskites

PL Photoluminescence

PLQY Photoluminescence quantum yield

RT Room temperature

RoHS Restriction of Hazardous Substances

STE Self-trapped exciton

SSNMR Solid-state nuclear magnetic resonance

TR-PL Time-resolved photoluminescence

TG Thermogravimetry

UV Ultraviolet

VBMax Valence band maximum

WD-XRF Wavelength dispersive X-ray fluorescence

XRD X-ray diffraction

Abbreviations | Olga Nazarenko

0D, 1D, 2D, 3D Zero-, one-, two-, and three-dimensional

I Cs[C(NH2)3]PbI4

II Cs[C(NH2)3]PbBr4

III Cs2[C(NH2)3]Pb2Br7

1-3 IV (n-C4H9NH3)2PbI4

4 V [C(NH2)3]2PbI4

5-6 VI CsPbBr3

7-9 VII (n-C4H9NH3)2PbBr4

10 VIII CsPb2Br5

11 IX [C(NH2)3]2PbBr4

6 X CsPbI3

XI [C(NH2)3]2SnBr4

XII Cs[C(NH2)3]SnBr4

XIII Cs2[C(NH2)3]Sn2Br7

XIV A compound obtained in the Rb-Cu-Bi-I-O-H system

XV A compound obtained in the Rb-Ag-Bi-I-O-H system

+ MA, CH3NH3 Methylammonium

+ FA, CH(NH2)2 Formamidinium

+ G, C(NH2)3 Guanidinium

CZT CdZnTe

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Summary | Olga Nazarenko

Summary

Research on semiconductors goes with its roots into the 19th century. Nowadays, the field of semiconductors is very broad; it includes elements, as well as binary, ternary, and quaternary semiconductor compounds from different classes, such as halides, chalcogenides, oxides, etc.. In recent years, mixed organic-inorganic (denoted in the literature by hybrid) ternary lead-halide coordination compounds with perovskite and perovskite-like crystal structures (LHPs) and a general formula APbX3 (A – monovalent organic cation, X – Cl, Br, I) have attracted immense attention due to their tunable optoelectronic properties, high absorption coefficients, and excellent charge transport characteristics, such as long charge carrier recombination lifetimes and diffusion lengths. Because of these properties, such compounds are promising for applications in such fields as photovoltaics, hard radiation detection, lighting, etc.. Of these compounds, the ternary perovskites consisting of MA and FA lead iodides and bromides – CH3NH3PbX3 and CH(NH2)2PbX3 (X = Br, I) – have come under the spotlight. The crystal structure of CH3NH3PbI3 consists of lead iodide octahedral building units that assemble + into a 3D anionic framework through corner-sharing. The CH3NH3 cations then fill the voids + within this framework. CH3NH3 interacts with the anionic frame through ionic, as well as hydrogen bonds. The degree of crosslinking of the anionic network can be modified through + the choice of the organic cation, e.g., use of long alkyl-chain amines CnH2n+1NH3 with n ≥ 4 often results in layered (2D) A2PbX4 perovskite-related structures. LHPs have an electronic structure with CBMin and VBMax made of antibonding orbitals, with vacancies as major defects. These vacancies form shallow defect levels near the band edges or within the bands. Furthermore, the charge carriers interact with lattice distortions forming large polarons that reduce the probability of scattering and facilitate long charge-carrier diffusion lengths. As a result of their electronic structures and crystal lattice dynamics, LHP compounds are considered defect-tolerant semiconductors. However, on the way to high- performance and long-lived devices, certain obstacles exist, for instance, the thermodynamic and chemical instability of LHP compounds, as well as the toxicity of lead latest at the step of an industrial product, underlying additional restrictions. This project was a fundamental study on 2D and 3D LHPs with several aims: - to study intrinsic optical and electronic properties of LHPs and derived compounds; - to synthesize novel lead-free Sn(II) and Bi(III) based halide coordination compounds; - to determine and understand crystal structures and optoelectronic properties of the new compounds.

In the first project, single crystals of CsxFA1-xPbI3-yBry (x = 0-0.1, y = 0-0.6) were grown – based on previously published findings from research on substitutionally doped (on A and X positions) CH(NH2)2PbI3 – to study their intrinsic properties. Thus, it was shown that CsxFA1−xPbI3−yBry (x = 0.1, y = 0.2–0.6) exhibited improved kinetic stability compared to the parent compound. The mixed-cation and mixed-anion phases were stable for at least three months compared to the stability of one to seven days for CH(NH2)2PbI3. After one week, the black, 3D CH(NH2)2PbI3 transforms into a 1D, yellow polymorph. The elemental analysis

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Summary | Olga Nazarenko

(CHN, Br, I) of selected samples was performed to estimate the composition of the obtained crystals, and to compare it to the ratio of components in solution. The study unveiled a higher Br/I ratio in the crystals as compared to the solution. The incorporation of Cs+ and Br- increases the bandgap of the resulting compound from

1.43 eV for CH(NH2)2PbI3 to 1.63 eV for Cs0.1FA0.9PbI2.4Br0.6. The electronic properties were probed by the use of single crystals for γ-radiation detection, which is possible due to the high atomic number (Z) and density of CsxFA1−xPbI3−yBry (x = 0.1, y = 0.2–0.6). These single crystals showed a high charge carriers mobility per lifetime product (up to 1.2×10-1 cm2 V-1, which is comparable to the best quality CZT), which allows sensitive detection of hard radiation. As compared to 3D LHPs, layered LHP compounds offer a much larger library of A- type cations and exhibit some distinct properties such as: - higher exciton binding energies (leading to higher PLQY); - the possibility to tune the optoelectronic properties through modifications to the thickness and geometry of lead-halide layers through the choice of A-type cation (larger cations to fill the interlayer space, and small cations to fill the cavities within the lead-halide layers); - (de-) intercalation of the ions in the crystal.

In this project, new layered Cs[C(NH2)3]PbI4 and Csn[C(NH2)3]PbnBr3n+1 (n = 1, 2) were + synthesized from acidic solutions. The use of highly symmetric C(NH2)3 cation, unlike typical long alkyl-ammonium cations leads to denser packed Pb-X (X = Br, I) anionic frameworks, + + which are connected by the C(NH2)3 cation via halogen-hydrogen bonds. The flat C(NH2)3 additionally induces anisotropic out-of-plane tilting of the lead halide octahedra. These compounds were found to possess a good temperature stability with decomposition only occurring above 571 K. This research also showed that Cs[C(NH2)3]PbI4 and Cs2[C(NH2)3]Pb2Br7 are photoconductive; Cs[C(NH2)3]PbI4 and Cs[C(NH2)3]PbBr4 are green and blue luminescent respectively under UV light, at moderate cooling; and that

Cs2[C(NH2)3]Pb2Br7 is blue luminescent at RT. Sn(II) can form isostructural complexes to Pb(II), which often exhibit smaller band gaps. Therefore, such tin-based compounds can be used for the replacement of lead-containing compounds in optoelectronic devices. In the second part of the work, the analogous Cs-

C(NH2)3-Sn-Br system was investigated. Two phases were found to exist in this system: Cs[C(NH2)3]SnBr4 and Cs2[C(NH2)3]Sn2Br7, next to the ternary [C(NH2)3]2SnBr4 and CsSnBr3.

The Sn-Br anionic frameworks in Csn[C(NH2)3]SnnBr3n+1 (n = 1, 2) are layered as was confirmed by the flat band dispersion along the stacking direction. Cs2[C(NH2)3]Sn2Br7, which is isostructural to Cs2[C(NH2)3]Pb2Br7, has a smaller optical band gap. This might be explained by the difference in the E3 ionization energies [E3(Pb) = 31.94 eV and E3(Sn) = 30.5 eV] the effect of the additional polarizability in Pb (presence of filled diffuse d and f orbitals below the valence electrons), and/or distortions of the M-X (M = Sn, Pb) octahedra. The stereochemical activity of the Pb(II) and Sn(II) lone pairs could be observed in [C(NH2)3]2PbBr4 and 2- [C(NH2)3]2SnBr4, where the 1D chains of metal-halides consist of corner-sharing [MBr5] (M = Sn, Pb) square pyramids. [C(NH2)3]2SnBr4, unlike its lead analog, is luminescent with a PLQY of about 75±5 % at 77 K and 2 % at RT. [C(NH2)3]2SnBr4 exhibits broad-band emission resulting from STEs.

Summary | Olga Nazarenko

Alternatively, some Bi(III) based 3D semiconductors, such as Cs2AgBiX6 (X = Cl, Br) were previously reported in the literature. These are indirect bandgap semiconductors with bandgaps of 2.19 eV and 2.77 eV for Cs2AgBiBr6 and Cs2AgBiCl6, respectively. The iodide analog was not obtained as it is thermodynamically unstable. In this project, the compositional space of Rb-Cu-Bi-I and Rb-Ag-Bi-I was under investigation. The synthesis from acidic aqueous media results in compounds, which have water molecules and hydronium ions as intrinsic components. In each system, one common phase was found. These compounds are semiconductors with resistivity on the order of 108 Ω m and bandgaps of 1.71 and 2.00 eV for rubidium copper iodobobismuthate hydrate and rubidium silver iodobismuthate hydrate, correspondingly. Both compounds were found to possess interesting and complex crystal structures with Bi-I octahedra connected with Ag-I or Cu-I units via corner-sharing to form a + + 3D network with large, square tunnels. In these tunnels, Rb cations, H2O, hydronium (H3O ), 3- and isolated BiI6 octahedra are situated. Compounds obtained in the Rb-Cu-Bi-I-O-H and Rb- Ag-Bi-I-O-H systems were studied with various techniques, such as single crystal XRD, SSNMR, and for the determination of the composition WD-XRF, and EDXS were used. The ratio between elements, obtained with EDXS, WD-XRF, and single crystal XRD differs in the amount of Cu(Ag) and I, which could indicate different compositions throughout the crystal suggesting a complex crystal growth mechanism. Research in thin-film perovskite photovoltaics has recently focused on the simultaneous substitution of the A and X sites in APbX3 perovskites to increase both the chemical and thermodynamic stability of the final phase. Next, the C(NH2)3-CH(NH2)2-Pb-I system was investigated, and the common obtained phase was found to have the formula

[CH(NH2)2][C(NH2)3]PbI4. This compound consists of corrugated layers made of Pb-I corner- sharing octahedra separated by two compact organic cations. [CH(NH2)2][C(NH2)3]PbI4 is thermally stable up to 528 K and exhibits RT PL in the red region with a QY of 3.5%. Even though the broad emission appeared to be the convolution of different PL bands, various measurement techniques - temperature-dependent, spatially resolved PL, TR-PL, PL excitation, and SSNMR - could not fully reveal the origin of the emission in [CH(NH2)2][C(NH2)3]PbI4. The results nevertheless suggested that the emission bands might result from both free excitons and self-trapped excitonic states, defects, or color centers. This study further highlighted the structural diversity that exists within the compositional space typically used in perovskite photovoltaics.

Zusammenfassung | Olga Nazarenko

Zusammenfassung

Die Forschung an Halbleitern geht bis zurück ins 19. Jhdt. Heute umfasst dieses Forschungsfeld ein breites Spektrum an Elementen sowie binäre, ternäre und quaternäre Verbindungen aus unterschiedlichsten Verbindungsklassen; etwa Halide, Chalkogenide, Oxide, usw. In den letzten Jahren haben anorganisch-organisch gemischt aufgebaute (in der Literatur als hybrid bezeichnet) ternäre Bleihalogenidverbindungen der allgemeinen Formel APbX3 (A = organisches Kation, X = Halogenidion) mit Perovskit- und perovskitartiger Kristallstruktur – Bleihalogenidperovskite (LHPs) – ein reges Interesse auf sich gezogen. Grund hierfür sind die einstellbaren optoelektronischen Eigenschaften, grosse Absorptionskoeffizienten und ausgezeichnete Ladungsträgertransportcharakteristiken wie etwa lange Ladungsträgerrekombinationszeiten und Diffusionswege. Aufgrund dessen ist diese Verbindungsklasse vielversprechend für Anwendungen etwa in der Photovoltaik, die Messung ionisierender Strahlung, Beleuchtungszwecke, usw. Besonders aus der Klasse der ternären

Perovskite sind MA und FA Bleiiodid und bromid – CH3NH3PbX3 und CH(NH2)2PbX3 (X =

Br, I) zu erwähnen. Die Kristallstruktur von CH3NH3PbI3 besteht aus einem dreidimensionalem + anionischem Netzwerk von eckenverknüpften Bleiiodidoktaedern, in welchem die CH3NH3 - Kationen die Lücken innerhalb des Netzwerkes besetzen. Methylammonium wechselwirkt dabei sowohl über ionische als auch Wasserstoffbrückenbindungen mit den Halogenidionen. Die anionischen Netzwerke lassen sich dabei in ihrem Vernetzungsgrad durch die Wahl der A- Kationen beeinflussen, bspw. führt der Einbau von Kationen mit langen Alkylketten, etwa + gemäss CnH2n+1NH3 mit n ≥ 4, häufig zu zweidimensional vernetzten (schichtartigen), perovskit-abgeleiteten Strukturen mit der Formel A2PbX4. Die LHP’s weisen eine elektronische Struktur auf, bei welchen das Energieminimum des Leitfähigkeitsbandes als auch das Maximum des Valenzbandes aus antibindenden Orbitalen bestehen, mit Leerstellen als hauptsächlichem Defekttyp. Diese Leerstellen bilden schmale Energiebereiche nahe an der Bandkante oder innerhalb der Bänder. Weiter interagieren die Ladungsträger mit den Gitterschwingungen unter Ausbildung von starken Polaronen, wodurch die Wahrscheinlichkeit der Streuung reduziert wird und lange Ladungsträgerdiffusionswege resultieren. Die LHP’s werden aufgrund dieser Kombination ihrer elektronischen Strukturen mit der Dynamik der Kristallstrukturen als defekttolerante Halbleiter beschrieben. Die Anwendung von LHP’s in effizienten und langlebigen technischen Einrichtungen ist jedoch noch nicht möglich, da sowohl die thermodynamische als auch chemische Beständigkeit nicht ausreichend ist, als auch die Toxizität von Bleiverbindungen, spätestens bei der Verwendung für industrielle Produkte, die zusätzlichen Regulationen unterliegen. Die vorgelegte Arbeit versteht sich daher als eine Grundlagenstudie von 2D und 3D LHP’s mit folgenden Zielen: - Studie der intrinsischen optischen und elektronischen Eigenschaften von LHP’s und von diesen abgeleiteten Verbindungen; - Synthese neuartiger ternärer Sn(II)- und Bi(III)-Halide und - Untersuchung der Kristallstrukturen und optoelektronischen Eigenschaften der neu erhaltenen Verbindungen.

Zusammenfassung | Olga Nazarenko

Für den ersten Teil wurden, basierend auf in der Literatur beschriebenen Ergebnissen über die substitutionelle Dotierung von CH(NH2)2PbI3 auf A- und X-Positionen, Einkristalle von CsxFA1-xPbI3-yBry (x = 0 – 0.1, y = 0 – 0.6) gezüchtet, mit dem Ziel die intrinsischen Eigenschaften zu studieren. Schwarzes FAPbI3 zeigt bei Raumtemperatur eine Phasenumwandlung aus der Perovskitstruktur mit dreidimensionalem Anionennetzwerk zu dem gelbem Polymorph mit eindimensionalem Anionennetzwerk. Es konnte gezeigt werden, dass CsxFA1-xPbI3-yBry (x = 0.1, y = 0.2 – 0.6) eine für die Anwendung wesentlich günstigere Kinetik hinsichtlich der Phasenumwandlung gegenüber der Stammverbindung aufweist. Die Phasenumwandlung bei den erhaltenen Phasen mit sowohl gemischten Kationen als auch gemischten Anionen benötigte mindestens drei Monate, während die Umwandlung von FAPbI3 sich innerhalb von sieben Tagen vollzieht. Die Zusammensetzung der gemischten Verbindungen wurde mithilfe der Elementaranalyse an Einkristallen bestimmt und erlaubte einen Vergleich mit dem molaren Ausgangsverhältnis der Komponenten in Lösung. Dabei zeigte sich ein höheres Br/I-Verhältnis im Kristall im Vergleich zur Lösung. Messungen zeigten, dass die Dotierung mit Cs+ und Br- die Bandlücke bei

Cs0.1FA0.9PbI2.4Br0.6 auf 1.63 eV gegenüber FAPbI3 mit 1.43 eV vergrössert. Die elektronischen Eigenschaften wurden über die Anwendung von Einkristallen als γ-Strahlendetektoren bestimmt, was aufgrund der grossen Ordnungszahlen (Z) und der Dichte von CsxFA1-xPbI3-yBry (x = 0.1, y = 0.2 – 0.6) möglich ist. Die Einkristalle zeigen ein gutes Produkt aus Ladungsträgermobilität × Rekombinationszeiten (bis zu 1.2×10-1 cm2V-1; vergleichbar mit CZT höchster Qualität), wodurch eine empfindliche Messung ionisierender Strahlung resultiert. Im Vergleich zu dreidimensional vernetzten LHP’s haben schichtartig aufgebaute LHP- Verbindungen eine grössere Auswahl an A-Typ-Kationen und zeigen weitere Eigenschaften wie: (i) Grössere Exzitonbindungsenergien und damit grössere Quantenausbeuten der Photolumineszenz (PLQY); (ii) Die Möglichkeit die optoelektronischen Eigenschaften über die Dicke und Geometrie der Metallhalidschicht zu modifizieren durch geeignete Wahl der A- Typ-Kationen (grössere Kationen besetzen Positionen zwischen den Schichten, kleinere Kationen innerhalb der Schichten); (iii) (De-)Interkalation von Ionen im bestehenden Kristall. In dieser Studie wurden die neuartigen, schichtartig aufgebauten Verbindungen

Cs[C(NH2)3]PbI4 und Csn[C(NH2)3]PbnBr3n+1 (n = 1, 2) aus sauren Lösungen erhalten. Die Verwendung des hochsymmetrischen und im Vergleich zu langen Alkylammoniumkationen + kompakten C(NH2)3 , führt zu dichter gepackten Pb-X-Anionennetzwerken (X = Br, I), welche durch das Guanidiniumkation über Halogen-Wasserstoffbrückenbindungen miteinander verbunden sind. Diese anisotrope Wechselwirkung des flachen Guanidiniumkations mit den Halidionen führt zu einer Auslenkung der Bleihalidoktaeder aus der Schicht. Diese Verbindungen zeigen eine ausgeprägte thermische Stabilität und zersetzen sich erst oberhalb

571 K. Weiter konnte gezeigt werden, dass Cs[C(NH2)3]PbI4 und Cs2[C(NH2)3]Pb2Br7 Photoleitfähigkeit aufweisen. Cs[C(NH2)3]PbI4 und Cs[C(NH2)3]PbBr4 zeigen grüne rsp. blaue

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Zusammenfassung | Olga Nazarenko

Lumineszenz bei Abkühlung unter Bestrahlung mit UV-Licht, Cs2[C(NH2)3]Pb2Br7 zeigt bereits bei Raumtemperatur blaue Lumineszenz unter Anregung. Sn(II)-Verbindungen sind häufig isostrukturell zu den entsprechenden Pb(II)- Verbindungen, wobei meistens Bandlücken geringerer Energie resultieren. Sie sind daher grundsätzlich für den Ersatz bleihaltiger Verbindungen in photoelektrischen und photolumineszenten Anwendungen geeignet. Für den zweiten Teil wurde das analoge quaternäre System Cs-C(NH2)3-Sn-Br untersucht. Es konnten zwei Phasen in diesen System nachgewiesen werden, Cs[C(NH2)3]SnBr4 und Cs2[C(NH2)3]Sn2Br7, nebst den ternären Phasen [C(NH2)3]2SnBr4 und CsSnBr3. Die Sn-Br-Anionennetzwerke in den Verbindungen

Csn[C(NH2)3]SnnBr3n+1 (n = 1, 2) sind 2D aufgebaut, was sich in den berechneten Bandstrukturen durch Bereiche geringer Banddispersion in Richtung der Stapelung bestätigt.

Cs2[C(NH2)3]Sn2Br7, welches isostrukturell zu Cs2[C(NH2)3]Pb2Br7 ist, weist eine geringere optische Bandlücke als sein Bleianalogon auf. Die Gründe hierfür dürften eine geringe 3.

Ionisierungsenergie (E3(Pb) = 31.94 eV gegenüber E3(Sn) = 30.5 eV), der Effekt der zusätzlichen Polarisierbarkeit bei Pb (gefüllte d- sowie f-Orbitalstände unterhalb der Valenzelektronen) und/oder Verzerrungen der M-X-Oktaeder (M = Sn, Pb) sein. Die Stereochemische Aktivität der nichtbindenden Elektronenpaaren von Pb(II) und Sn(II) konnte bei [C(NH2)3]2PbBr4 sowie [C(NH2)3]2SnBr4, welche eindimensionale Ketten aus 2- eckenverknüpften, quadratischen Pyramiden [MBr5] (M = Sn, Pb) aufweisen, beobachtet werden. [C(NH2)3]2SnBr4 ist im Vergleich zu seinem Bleianalog lumineszent mit einem PLQY von 75±5 % bei 77 K und 2 % bei Raumtemperatur. [C(NH2)3]2SnBr4 weist eine brandbandige Emission auf bedingt durch selbstgefangene Exzitonen. Als Alternative zur Verwendung von Sn(II)-Analoga wurden in der Literatur auch

Bi(III)-Verbindungen, wie etwa Cs2AgBiX6 (X = Cl, Br), mit dreidimensional vernetztem Anionennetzwerk als geeignete Halbleiter beschrieben. Sie weisen eine indirekte Bandlücke auf mit 2.19 für Cs2AgBiBr6 rsp. 2.77 eV für Cs2AgBiCl6. Das Iodidanaloga konnte aufgrund thermodynamischer Instabilität nicht erhalten werden. Als Ausgangspunkt wurden daher die Systeme Rb-Cu-Bi-I und Rb-Ag-Bi-I für die Suche nach neuen Phasen ausgewählt. Die Synthese aus sauren, wässrigen Lösungen ergab jedoch Verbindungen, welche sowohl Wasser als auch Hydroniumionen als strukturellen Bestandteil aufnehmen. In beiden konnten jeweils eine neue halbleitende Phase erhalten werden. Die elektrischen Widerstände der gefundenen Phasen liegen im Bereich um 108 Ωm und die Bandlücken bei 1.71 eV für Rubidiumkupferiodobismutathydrat rsp. 2.00 eV für Rubidiumsilberiodobismutathydrat. Beide Verbindungen besitzen sehr interessante und komplexe Kristallstrukturen, in welchen Bi-I- Oktaeder mit Ag-I- oder Cu-I-Einheiten via Eckenverknüpfung eine dreidimensionales Anionennetzwerk bilden. Dieses besitzt grosse, quadratische Tunnel, in welchen sich die Rb+- + 3- Kationen, H2O, Hydroniumionen (H3O ) und isolierte BiI6 -Oktaeder befinden. Die erhaltenen Verbindungen wurden nebst Röntgendiffraktometrie und Festkörper-NMR zur Bestimmung der Zusammensetzung auch mit WD-XRF und energiedispersiver EDXS vermessen. Das mit den einzelnen Methoden bestimmte Elementeverhältnis zeigt Abweichungen betreffend dem stöchiometrischen Anteil von Cu(Ag) und I, was möglicherweise ein Indiz für eine variierende Zusammensetzung innerhalb eines Kristalls und damit für einen komplexen Kristallwachstumsprozess ist.

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Zusammenfassung | Olga Nazarenko

Die Forschung an Perovskiten für Dünnfilmanwendungen hat seit einiger Zeit einen

Fokus auf die gleichzeitige Substitution von A- sowie X-Positionen in APbX3-Perovskiten gerichtet, mit dem Ziel die chemische und thermodynamische Stabilität der Produktphase zu erhöhen. In dieser Studie wurde daher das System C(NH2)3-CH(NH2)2-Pb-I als möglicher Kandidat für Dünnfilmschichten untersucht. Hier konnte die Verbindung

[C(NH2)3][CH(NH2)2]PbI4 gefunden werden. Die Verbindung besteht aus gewellten Pb-I- Oktaederschichten zwischen denen die beiden kompakten organischen Kationen eingebettet sind. [C(NH2)3][CH(NH2)2]PbI4 ist thermisch stabil bis 528 K und zeigt Lumineszenz im roten Spektralbereich bei Raumtemperatur mit einer Quantenausbeute von 3.5 %. Die breite Emission von [C(NH2)3][CH(NH2)2]PbI4 besteht aus mehreren Übergängen, deren genaue Ursprünge trotz Zuhilfenahme unterschiedlicher Techniken der Photolumineszenz-Spektroskopie und Festkörper-NMR nicht geklärt werden konnten.

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Introduction | Olga Nazarenko

Chapter 1. Introduction

1.1. A brief history of semiconductors The epoch of semiconductors began in 1833 with a work of Michael Faraday on the 12 conductivity of Ag2S. Michael Faraday observed a negative temperature coefficient of resistivity. For metals, for instance, this coefficient is a positive number, and resistivity of metals increases with temperature increase. Further, in 1839 Becquerel discovered a generation of photovoltage by shining light on the AgCl electrode that was immersed into an acidic solution and connected to Pt electrodes.12 The first works on semiconductor compounds were based on observations and descriptions of new phenomena, but no theory existed yet to explain the obtained results. Two other important characteristics of semiconductors: photoconductivity (observed on Selenium by Warren Smith in 1873) and rectification (Adams and Day, 1877) were discovered some years later.13 In 1874 Clerk Maxwell wrote: ’’I saw conductivity of Selenium as affected by light. It is most sudden. Effect of a copper heat insensible. That of the sun great.14’’ The observed phenomena of photoconductivity is a bulk effect that can be explained by the absorption of the photons of the incident light - whose energy exceeds a specific value, the bandgap energy - by the body of the semiconductor crystal generating an electron-hole pair. In such a way, a photon to energy conversion happens. The process of rectification appeared to be very practically important, allowing to convert alternating current into a direct current. Se and Cu-Cu2O were first rectifiers. In 1904 Hallwachs constructed a 15 solar cell base on Cu-Cu2O. The understanding of the process of rectification came only in the 1930s after Nevill F. Mott, Walter H. Schottky, and Boris Davydov developed barrier models for metal-semiconductor contacts (nowadays called Schottky junction), according to

Introduction | Olga Nazarenko which the rectification is a contact effect: a concentration of electrons on the surface of semiconductor forming a barrier to a current flow.12 During the first decades of research on the semiconductors, a challenge was to find proper measurements that would help to identify and characterize the compounds of this class. Therefore, one could hardly overestimate the importance of the first observation of the Hall effect in 1879. The Hall effect describes the difference of the potentials between two opposite faces of a crystal when a crystal is placed in a magnetic field. This effect gives means to differentiate the nature of the charge carriers, electrons (n-type semiconductor) or holes (p-type semiconductor), as well as to estimate their mobility (µ). Now it became possible to distinguish between metals and semiconductors using such characteristics as the density of charge carriers and their sign.16 The further advances were enabled by the development of the quantum theory. Thus, in 1900 Max Planck developed a theory of black body radiation and proposed that the electromagnetic waves are quantized and not continuous. In 1905 Albert Einstein described the photoelectric effect with the light-quantum; he concluded: “Monochromatic radiation of low density behaves, in the thermodynamic sense, as if it consisted of mutually independent radiation quanta of magnitude 푅훽휈/푁0 (Albert Einstein equaled (푅/푁0)훽 to the Plank constant h)”.17 Development of the quantum theory by Max Plank, Albert Einstein, Niels Bohr, and Arnold Sommerfeld, as well as the application of its concepts to the theory of metals in 1928 by Sommerfeld and Felix Bloch launched the development of the theory of the semiconductors. The first application of semiconductor properties was found in a cat’s whisker crystal radio receiver. It was constructed beginning of 20th century by Jagadish Bose, after an existence of nonlinear current-voltage (I-V) characteristics in metal sulfide crystals was reported by Karl F. Braun in 1874.18 The cat’s whisker crystal radio receiver had a simple construction: a metal wire held to the semiconductor crystal (typically galena, PbS) by a mechanical pressure. The rectification by the metal-semiconductor junction lies in the principle of work of such receiver. In 1909 Karl F. Braun together with Guglielmo Marconi (the father of long-distance wireless transmission) received a shared Nobel Prize for the contribution to the development of wireless telegraphy. After the first application attempts, the problem of the purity of the studied semiconductors was identified - the experiments were not reproducible. Yet, for the development of the quantum theory, an accurate comparison with practical observations was necessary. In 1931 Wolfgang Pauli expressed his opinion: “One shouldn’t work on semiconductors, that is a filthy mass; who knows whether any semiconductors exist.” A Polish scientist Jan Czochralski did a significant step in the direction of receiving pure highly crystalline semiconductors. Czochralski introduced in 1918 a method of growing single crystals by pooling them out from a melt. Further, a Zone refining was introduced in 1952 by William Pfann, a research metallurgist at Bell Laboratories, which enabled scientists to obtain a highly pure Ge.19 For this, a bar of Ge was molten from one side, and the molten zone was shifted to the other side of the bar. During this process, impurities concentrate in the melted zone and move up to the end of the Ge bar; this zone can next be removed and reprocessed. Though, zone-refining was proved ineffective for removal of impurities that contained B and Si, as these must be removed prior to zone-refining.20 Thus, the growth of single crystals of Ge by

Introduction | Olga Nazarenko

Czochralski method with the next purification by zone refining give about 1010 cm-3 impurity concentration.21 The electronics era began with the invention of the transistor in 1947 at Bell Laboratories by John Bardeen, Walter Brattain, and William Schockley.22-23 This invention was governed by the desire to substitute bulky, fragile vacuum tubes in telephone lines, and was based on the radar technologies of that time, in particular, on the experience of germanium crystals growth and a rectification principle of metal-semiconductor contact. The invention of zone-refining made the production of transistors feasible. Furthermore, another type of purification procedure for Si was developed at Siemens laboratories in 1960: chemical vapor deposition method (Siemens-process) for polycrystalline Si of high purity and floating-zone method for monocrystalline Si.24 Si with time substituted Ge in transistors, diodes, etc., as Si has a higher bandgap and needs smaller collector cutoff currents. Besides, Si withstands higher temperatures and voltages than Ge. Industrial demands stimulated progress in the field of electronics based on Si. Pure silicon itself is a bad conductor with a conductivity of 5×10-4 Ω-1 m-1 for single crystals,25 so controlled doping is necessary to increase the conductivity. The type of conductivity (n- or p-type) of the final doped Si can also be chosen. For example, doping with phosphorous results in an n-type semiconductor (102-103 Ω-1 m-1), whereas the introduction of boron leads to a p-type semiconductor (103 Ω-1 m-1).25 Furthermore, the development of planar technology (devices are made directly on the flat semiconductor plate) enabled the construction of microelectronic schemes. Si is also used in solar cells. In 1958, Gerald L. Peterson, Daryl M. Chapin, and Calvin S. Fuller patented a silicon solar cell with 8 % efficiency.15 These days the efficiency of Si (crystalline) solar cells is as high as 26.7±0.5 %.26 Among the other applications, semiconductor materials are used in hard radiation detection, lasers, and photodetectors.27-29

1.2. Mixed organic-inorganic lead halide perovskites: basics and properties

Perovskites are compounds adopting a crystal structure of a mineral calcium titanate, CaTiO3. This mineral was called Perovskite by the mineralogist Gustav Rose in honor of the Russian

Introduction | Olga Nazarenko

31 mineralogist Lev Perovski. A general formula of perovskite compounds is ABX3, where A and B are cations and X – anion, with qA + qB = -3qX, where q is a charge of the corresponding ion. In the case of halide perovskites, A is a monovalent and B – divalent cations. Ideal or aristotype perovskite crystal structure belongs to the cubic crystal system and consists of BX6 octahedra that condense by corner-sharing into a 3D network with voids filled by A-type cation

(Figure 1.2a). After the discovery of dielectric and ferroelectric properties of perovskite BaTiO3 in the 1940s, increased scientific interest was oriented on the search of compounds with the perovskite crystal structures, along with the study of their optoelectronic and magnetic properties.32-33 Subsequently, perovskite crystal structures were found among such classes of compounds as halides, cyanides, sulfides, oxynitrides, hydrides, etc..34 Complex organic- inorganic perovskites - CH3NH3PbX3 (X = Cl, Br, I and mixtures thereof) were reported in 1978 by Dieter Weber.35 The reported compounds are composed of a 3D lead halide framework - of corner-sharing [PbI6] octahedra and MA cations situated in the cavities within the lead halide anionic network (Figure 1.2b).35 Fascinating is an intense color of these LHPs - except for

MAPbCl3, which is colorless, but MAPbI3 is black, and MAPbBr3 is orange - and a broad mixing range of halides in mixed Br/Cl and Br/I systems.35

Fully inorganic CsPbX3 compounds exist in a perovskite phase only at higher temperatures, and a color change indicates a phase transition to the higher symmetry network. 6 This observation was reported by Christian K. Möller in 1959. CsPbI3 changed its color from yellow to black around 578-581 K and, based on the powder XRD, adopted a distorted perovskite crystal structure. The size of Cs+ cation is too small to maintain a 3D perovskite lattice in Pb-I system. CsPbI3 crystallizes at RT in the orthorhombic crystal system, space

Figure 1.2. (a) The aristotype perovskite crystal structure (the crystallographic data, cif file, was downloaded from the ICSD database, card 23076). (b) The crystal structure of LHP CH3NH3PbI3 at 253 K (ICSD card 238610), MA cations show a highly dynamic behavior in the lattice cavities. Two cations can build a 3D perovskite crystal structure in Pb-I system: + + CH3NH3 and CH(NH2)2 , shown on the right side of the picture. group Pnma. The crystal structure of CsPbI3 at RT consists of 1D chains of corner-sharing - + 36 [PbI6] octahedra, and Cs cations are situated in between these chains.

Introduction | Olga Nazarenko

In 1926 in ’’Die Gesetze der Krystallochemie’’ Victor M. Goldschmidt formulated laws of the dependence of a formed crystal structure on its chemical composition.37 These laws were derived after a careful analysis of the crystallographic data of numerous compounds by

Goldschmidt and his co-workers. Thus, analysis of crystallographic data of CaTiO3, BaTiO3,

SrTiO3, KMgF3, etc. lead to the discovery of the coefficient nowadays called Goldschmidt tolerance factor (GTF or α) that allows predicting the formation of the perovskite crystal structure:

푅퐴 + 푅푋 = 훼√2(푅퐵 + 푅푋) (1.1) where Ri is an ionic radius of the corresponding ion. When α lies between 0.8 and 1 a perovskite crystal structure forms, but when α is less than 0.8 typically corundum-type structures result and α > 1 leads to hexagonal structures of calcite-type. In the case of mixed organic-inorganic metal-halide perovskites, the calculation of the ionic radii of organic cations is challenging for the following reasons: (i) the organic cations, typically long chain alkyl-amines, have a non- spherical geometry and (ii) they often form hydrogen-halogen bonds with halide ions of the metal-halide framework. The reduced symmetry of the organic A-type cation is one of the important differences between fully inorganic and hybrid LHPs. G. Kieslich et al. used a set of crystallographic data to estimate the effective radii of organic cations by applying a rigid sphere model and assuming free rotational freedom around the center of mass of the organic molecule.38-39 G. Kieslich derived an equation for the evaluation of the effective radii of the organic cations:

푟퐴푒푓푓 = 푟푚푎푠푠 + 푟𝑖표푛 (1.2) where rmass is a distance between the center of mass of the molecule and a furthest away lying atom (hydrogen atoms excluded), rion is an ionic radius of this atom. Molecular anions such as - - CN , HCOO are typically considered as rigid cylinders of a certain height, hXeff, and radius, 39 rXeff. Based on GTF, cations listed in

Table 1.1 could form perovskite APbI3 compounds. Unfortunately, hydroxylamine is an explosive hazard when heated. In the case of azetidinium cation, C3H6NH2PbI3 crystallizes in the hexagonal crystal system, rhombohedral crystal lattice, a space group R-3. C3H6NH2PbI3 possesses a 0D crystal structure that consists of groups of three face-sharing Pb-I octahedra and 40 organic cations are situated in between these units. Imidazolium lead iodide, C3N2H5PbI3, crystallizes in the hexagonal crystal system, a space group P63/m. The crystal structure of

C3N2H5PbI3 is made of 1D chains of face-sharing lead iodide octahedra and organic cations 41 situated in between these chains. In the N2H5-Pb-I system, there were three phases reported so far. All three compounds have a yellow color, and neither of them crystallizes in a perovskite crystal structure.42 Within azetidinium-, imidazolium- and hydrazinium-Pb-I systems no perovskite crystal structures were found so far, which highlights that not only the sizes of the ions play the crystal structure directing role, but such parameters as the spatial geometry of the organic cations, their interactions with a metal-halide lattice and among each other are of importance.

Introduction | Olga Nazarenko

Table 1.1. The organic cations that should form a perovskite crystal structure in Pb-I system according to GTF.38

Organic cation Molecular formula Goldschmidt tolerance factor + hydroxylammonium NH3OH 0.909 + methylammine CH3NH3 or MA 0.912 + hydrazinium N2H5 0.912 + azetidinium C3H6NH2 0.980 + formamidinium CH(NH2)2 or FA 0.987 + imidazolium C3N2H5 0.997

MAPbI3 has a perovskite crystal structure and crystallizes in the tetragonal crystal 43 system, a space group I4/mcm (Figure 1.2). The tetragonal phase of MAPbI3 exists in about 160 – 327 K temperature range, and at ~327 K MAPbI3 undergoes a second-order phase 44 transition to cubic phase, whereas below 160 K MAPbI3 has an orthorhombic crystal structure. There is a dynamic disorder of MA cations within the Pb-I framework at RT. The reorientation (on the ps time scale) and wobbling of the MA cations within the cavities of the Pb-I lattice was studied with the quasielastic neutron scattering and infrared spectroscopy.45-46 After being discovered by Dieter Weber in 1978, MAPbI3 attracted the attention of scientists again in 2009, when it was first used as a light sensitizer in a liquid electrolyte solar cell with an efficiency of 3.8 %.47 The boost in the research on this and similar compounds started especially after in 48 2012 the efficiency of ~10 % was obtained for all solid-state solar cell based on MAPbI3. Since then, thousands of articles are published every year on LHPs, their optoelectronic properties, and application. Nowadays, MAPbI3 is the most studied compound in the field of research on hybrid metal-halide perovskites. The properties of MAPbI3 that attract close attention are: (1) intrinsic conductivity and a suitable band gap for instance for sunlight harvesting:

Eg(MAPbI3) was estimated to be in the range from 1.51 eV (estimated from absorption measurements performed on single crystals by Saidaminov et al. in Ref49) to 1.61 eV (on single crystals, reported in Ref50); (2) high absorption coefficients (104-105 cm-1 comparing to 103 cm-1 for Si, Figure 1.3);51- 52 (3) single crystals growth and films deposition is performed from a solution at low temperatures (< 423 K);

(4) long charge carrier diffusion lengths (LD >175 μm for electrons and holes, under one sun illumination);50 (5) long charge carriers recombination lifetime (μs level)50, 53 (6) the low density of trap states (3.3×1010 cm-3).53

54 The LD is defined by the average distance the relevant charge moves in the semiconductor, and can be calculated using the charge carriers recombination lifetime (휏푟) and they mobility (μ) according to the formula:

Introduction | Olga Nazarenko

퐿 = √푘퐵푇휇휏푟 (1.3) 퐷 푒 where 푘퐵 is a Boltzmann constant, T – temperature and e – the particle charge. Using the Hall 50 effect measurements, Q. Dong et al. showed a hole concentration in MAPbI3 single crystals (a measurement was performed using Au and Ga contacts) to be 9(±2)×109 cm-3 with a hole 2 -1 -1 mobility μp = 105±35 cm V s . Similarly, using Ga contacts, an electron density was 10 -3 2 -1 -1 measured to be 4.5×10 cm with a mobility μn = 24.8±4.1 cm V s , which is smaller than 2 2 -1 -1 50, 55 the mobility of electrons in GaAs - μn ≤ 85×10 cm V s . The 휏푟 was measured with impedance spectroscopy and transient photovoltaic to be τr = 82±5 and 95±8 μs correspondingly. From these measurements, according to the formula 1.3, Q. Dong, et al. 50 calculated LD for holes to be 175±25 μm under 1 sun. Charge carrier mobility, lifetime, a diffusion length, as well as a density of the trap states are crucial parameters for estimation of the electronic quality of the semiconductor. The amount of trap states was estimated by D. Shi 10 -3 et al. from the current-voltage (I-V) response of single crystals of MAPbI3 to be 3.3×10 cm , which is similar to a high quality crystalline n- and p-type Si (108 - 1015 cm-3).53, 56-57 The reason for the comparatively small amount of the trap states in LHPs single crystals will be discussed in the next section 1.3. 58 The best efficiency of solar cells based on MAPbI3 thin films is 19.71 %. The significant issues with MAPbI3 on the way to its practical application are its chemical stability and a phase transition at ~ 327 K leading to a change in the properties. The chemical stability problems originate from the degradation of MAPbI3 in the presence of moisture according to the scheme:59-60

→ + [(퐶퐻3푁퐻3)푃푏퐼3]푛 + 퐻2푂 ← [(퐶퐻3푁퐻3)푛−1(푃푏퐼3)푛[퐻3푂 ] + 퐶퐻3푁퐻2 (1.4) + → [(퐶퐻3푁퐻3)푛−1(푃푏퐼3)푛[퐻3푂 ] ← 퐻퐼 + 푃푏퐼2 + [(퐶퐻3푁퐻3)푃푏퐼3]푛−1 + 퐻2푂 (1.5)

Formamidinium lead iodide perovskite (α-FAPbI3) might be a better candidate for photovoltaic or other applications due to its higher chemical stability and suitable optoelectronic properties. Although, the challenge in work with α-FAPbI3 is a phase transition that it undergoes at RT from the black colored compound with a 3D perovskite crystal structure 61 - to yellow -FAPbI3 - a hexagonal 1D phase. A work on substitutional doping of FA and I ions in α-FAPbI3 and a study of the intrinsic properties of the obtained compounds were performed in this Ph.D. project, and this work is presented in Chapter 3. The LHPs single crystals are typically grown by (i) cooling the hot solution (often using hydrohalic acids water solution as a solvent), (ii) slow vapor diffusion, (ii) an inversed temperature solubility method [typically from N,N-dimethylformamide, γ-butyrolactone].49 Thin polycrystalline films are mostly deposited from N,N-dimethylformamide or dimethyl sulfoxide solutions by dissolving corresponding halide salts, for instance, MAI and PbI2 to obtain MAPbI3. Thus, LHPs are synthesized at rather low temperatures. LHP compounds are promising for application in various devices, such as solar cells, photodetectors, hard radiation detectors, which will be further discussed in section 1.6 of the introduction. The excellent semiconductor properties of LHPs (listed above) were elucidated through the examination of

Introduction | Olga Nazarenko

Figure 1.3. The effective absorption coefficient of a thin film of CH3NH3PbI3 comparing to other semiconductors, such as a-Si (amorphous), GaAs, Cu(In, Ga)Se2 (CIGS), CdTe, and c-Si (crystalline), measured at RT. Reprinted with permission from Ref.52.Copyright 2014, American Chemical Society. the electronic structure of these compounds, a study of the crystal lattice dynamics, as well as interactions of the charge carriers with phonons, which will be discussed in the following sections 1.3 and 1.4.

1.3. Bandgap theory of semiconductors The difference between metals, semiconductors, and insulators is found in their conductivity: metals are excellent conductors with conductivity on order of 108 Ω-1 m-1 at 1 K, and insulators do not conduct or conduct very poorly with as low conductivity as 10-20 Ω-1 m-1 for an extreme insulator. Semiconductors lie in the middle between the two extremes indicated above with conductivity in the range of 10-7 – 1 Ω-1 m-1.62 The conductivity (σ) can be estimated using the next formula:

휎 = (푛휇푒 + 푝휇ℎ)푒 (1.6) where n and p are numbers of electrons and holes. The change of conductivity with temperature is defined mainly by the number of charge carriers, as charge (e) is temperature independent, and mobility (μ) decreases slightly with increasing the temperature due to the increase in lattice vibrations. The dependence of the concentration of charge carriers on the temperature in the semiconductor can be described via a Fermi-Dirac distribution:

푛푝 = 푛0 푝0 푒푥푝(−퐸푔/푘퐵푇) (1.7) 63 At T = 300 K kBT = 26 meV. For an intrinsic semiconductor with n = p the equation (1.7) becomes:

1/2 푛 = 푝 = (푛0 푝0) 푒푥푝(−퐸푔/2푘퐵푇) (1.8) The difference in n for metals and semiconductors can be understood from the bandgap theory. There are two theoretical approaches to the bandgap theory: from the point of chemistry (based on molecular orbitals, MOs, theory) and from the point of physics when metal is considered to be a potential well or a lattice with a periodic potential. Chemists describe a band structure

Introduction | Olga Nazarenko using the MOs approach. Taking the example from Roald Hoffmann’s “Solids and surfaced: a chemist’s view of bonding in extended structures, ”64 a chain of equally spaced hydrogen atoms can be considered as a piece of the ring with an imperceptible bending. The MOs of rings of various sizes are presented in Figure 1.4 (the figure was reproduced from Ref64). According to the MO theory, the number of combined orbitals is equal to the number of the newly formed orbitals, and each orbital is occupied with a maximum of two electrons. Combination of two 1s orbitals in H2 molecule results in two sigma (σ) orbitals: σg, lying lower in energy – bonding * orbital, and σg lying higher in energy – antibonding (Figure 1.4). In cyclic H3 or cyclopropane, one bonding orbital and two degenerate orbitals with an antibonding character form. When orbitals are in anti-phase and can’t constructively overlap a node is formed. Node is a zero electron density between the two nuclei; nodes number is growing with an energy increase. The lowest bonding orbitals have no nodes. A MO diagram for a cycle with an infinite amount of atoms shows the lowest bonding nodes-free orbital followed by pairs of degenerate orbitals and the highest energy orbital with the largest amount of nodes. As the amount of orbitals increases the difference in energy between them decreases generating bands of energy levels that can overlap or create energy gaps. This approach can be extended to 3D structures by writing orbitals using translational symmetry.64 The MOs can be described as a linear combination of atomic orbitals (LCAO), which are derived using irreducible representations:

푘푛푎 휓푘 = ∑푛푒 휒푛 (1.9) where 휓푘 is a MO wavefunction, k is an index that labels an irreducible representation, a – the lattice constant. In physics, instead of LCAO, the Bloch functions are used:

𝑖푘·푟 휓푘(푟) = 푢푘(푟)푒 (1.10) 𝑖푘·r where 휓k(푟) are the eigenfunctions of the Schrödinger equation for a periodic potential, 푒 - a plane wave, 푢푘(푟) - a periodic function. k is a wave vector in the Brillouin zone. The range of unique k values defines the first Brillouin zone: -π/a ≤ k < π/a.

Introduction | Olga Nazarenko

In summary, k is a wave vector, symmetry label, and a node counter. Plotting E(k) versus k gives a graph called the band structure. The shape and the run of the curve is influenced by the topology of the overlap of the MOs.64 The other important characteristic of the band structure, described by Roald Hoffmann in his work,64 is a dispersion of the band that is called a bandwidth and depends on the overlap between the orbitals of different unit cells; the larger is the overlap, the more dispersed the band is. The occupied orbitals of semiconductors form a valence band (VB), whereas higher energy empty orbitals form a conduction band (CB). In semiconductor compounds, the VB and the CB do not overlap and create a bandgap. The bandgap is estimated as minimum energy (typically expressed in eV) necessary to promote an electron from the highest occupied valence orbitals to the lowest unoccupied orbitals of the conduction band. In metals, on the other hand, the VB and the CB overlap, allowing high conductivities. To induce a conductivity in the semiconductor, energy has to be given (for instance, from light photons) to promote electrons to the CB. On the place of electrons, holes are formed. Depending on the k value of the VBMax and CBMin from E-k graph direct and indirect bandgap semiconductors can be distinguished (Figure 1.5a). Elements that belong to the class of semiconductors are Ge, Si, Se, as was already mentioned in a first subchapter, Te, as well as grey α-Sn (diamond structure), with the smallest bandgap among semiconductors of ~0 eV. Examples of binary semiconductors include compounds formed by elements from groups III-V (cubic and hexagonal phases of BN, AlP,

Figure 1.5. (a) A scheme of direct and indirect bandgap semiconductors: when k values of CBMin and VBMax are equal, a direct bandgap forms (for instance, direct bandgap semiconductors are GaAs, PbI2), but if the k values of CBMin and VBMax are different, a semiconductor has an indirect band gap (for example, Si, Ge). (b) A MOs diagram for MAPbI3. AlAs, AlSb, GaAs, InP, etc.), II-VI (ZnO, MgO, ZnSe, HgTe, etc.), I-VII (CuX, X = F, Cl, Br, I, AgI), IV-IV (SiC), as well as IV-VI (PbS, PbSe, SnTe, etc.). To ternary semiconductors belong I-III-VI2 compounds, such as CuAlSe2 and CuInS2. Cu2ZnGeS4 is an example of quarternary semiconductors. In semiconductor compounds, the bond nature lies between covalent and ionic. It is typically accepted that semiconductors are compounds with a band gap 65 66 <3 eV. Although, AlN (Eg ~ 6 eV) and diamond (5.5 eV) are considered to be broadband semiconductors.67 The VB in LHPs is composed of the Pb 6s2 and I 5p orbitals (Figure 1.5b). The mixing of these orbitals leads to an antibonding character of the VBMax. The CBMin has an

Introduction | Olga Nazarenko

Society. (b,c) The band structure and the density of states of tetragonal CH3NH3PbI3. The density of states graph shows a contribution of atoms to the valence and conduction bands. Reprinted with permission from Ref.69. Copyright 2014, American Chemical Society. antibonding character as well. The CB is made of Pb 6p and I 5p orbitals (Figure 1.5b).

According to the electronic structure calculation, MAPbI3 is a direct bandgap semiconductor with a bandgap at a Γ point - the center of the Brillouin Zone (Figure 1.6a, b). The Γ point coincides with a position of the Pb atoms in the lattice. The overlap of the Pb 6s and 6p orbitals with the I 5p orbitals can be observed from the density of states calculations (Figure 1.6c).70 Such overlaps contribute to the high dispersion of the conduction and valence bands. Moreover, 2 due to relativistic effects, the Pb 6s pair is stabilized, and 6p orbitals are split into 6p1/2 and 6p3/2 as a result of the spin-orbit coupling, generating a more dispersed conduction band. Therefore, during the calculations of the electronic properties of the LHPs, the spin-orbit coupling has to be taken into account. Organic cations do not contribute to the frontier orbitals. Although, due to spatial geometry and hydrogen-halogen bonds formation, organic cations influence the distortions of Pb-Hal octahedra and, consequently, a topology of the overlap of lead and halide orbitals.

1.4. The defect tolerance As was mentioned in section 1.1 of the introduction, the study on the semiconductors requires that they are highly pure to obtain reproducible results of the measurements. Besides, a highly crystalline semiconductor compound gives better performance. At first, this was concluded based on practical experiments, observations, and a common sense. Thanks to further development of the theory of semiconductors, nowadays, the requirements of high purity and crystallinity can be explained from the bandgap structure of semiconductor compounds. Thus, impurities and lattice imperfections (defects) create additional energy levels within the bandgap

Introduction | Olga Nazarenko and influence optoelectronic properties of semiconductors. Interestingly, some compounds inherently have a low probability of deep defects formation. Instead, shallow defects appear near the CBMin or the VBMax. This effect is called a defect tolerance. In LHPs, for instance, from such point defects as vacancies, interstitials, and antisites, mainly vacancies form.71 Vacancies have low formation energy comparing to other point defects (Figure 1.7a). Vacancies in LHPs form in the valence and conduction bands or close to them (Figure 1.7b), owing to a specific electronic structure of LHPs. On the other hand, interstitials and antisites would generate energy levels deep in the bandgap of LHPs, but they have high formation energies (it is energetically demanding to misplace ions in the lattice), and therefore, a low probability of formation. In comparison, in PbS (Eg = 0.42 eV) lead vacancies have a shallow character, while sulfur vacancies form as deep color centers.72 The VBMax and CBMin in LHPs are made of antibonding orbitals (Figure 1.5b, Figure 1.7b). The defect tolerance was also observed in a binary compound Cu3N, where VBMax is made of antibonding orbitals, and CBMin has a bonding character.73 In defect-intolerant semiconductors, such as CdSe, the band gap consists of bonding VBMax, and defect states typically lie deep in the band gap (Figure 1.7b). The dominance of the vacancies as point defects near the band edges or in the electronic bands in LHPs do not yet fully explain the long diffusion lengths and recombination lifetimes of the charge carriers that have moderate mobility. The dynamics of the LHPs lattices were investigated to understand the charge carriers’ behavior, in particular phonons (lattice vibrations) and a formation of polarons. The mobility of charge carriers can be limited by the effective masses of the free carriers and by scattering. The effective masses of LHPs are rather small, me and mh are 0.23mo, and 0.29mo correspondingly (spin-orbit coupling effects taken into account), as calculated by G. Giorgi et al. for MAPbI3, and they are as small as in silicon - 74 me = 0.26mo and mh = 0.29mo. The charge carriers can potentially scatter on phonons, impurities, and defects. With infrared spectroscopy and the temperature dependent PL measurements, it was concluded that scattering on the phonons is the main scattering mechanism.75-76 At the same time, the long diffusion lengths and recombination lifetimes of the charge carriers suggest that the scattering is not a frequent event. It was suggested that large polarons (larger than a unit cell) are the reason (Figure 1.7c).77 A strong electron-phonon coupling results in a large polaron particle, with a bigger effective mass as the effective mass of the charge carrier initially, protecting the charge carriers from scattering. The electron- phonon interaction is characterized by certain energy, which reduces the probability of electron- hole recombination.77 X.-Y. Zhu and V. Podzorov estimated the effective carrier masses based on the experimental results, in particular, on the carrier mobilities in single crystals of MAPbI3:

(푒휆)2 휇 = 푒휆⁄ → 푚∗ = (1.11) √2푘푇푚∗ 2푘푇휇2 where 푚∗ is an effective carrier mass, 푘푇 – a thermal energy, λ – a mean-free path. From Hall 2 -1 -1 measurements, 휇 of the single crystals of MAPbI3 fall into the range of 10-60 cm V s , ∗ leading to rather heavy charge carriers, 푚 ≈ 10푚푒 − 300푚푒, which supports the idea of large polarons formation in LHPs.77 As was mentioned in section 1.2, MA cations are dynamically disordered in the LHP lattice: they wobble and rotate by changing orientation every few ps.45- 46 MA has a dipole moment (2.3D) and its reorientation contributes to the dielectric response.46

Introduction | Olga Nazarenko

Due to the direct bandgap and the p-p absorption transition - from I p orbitals that give a significant contribution to VBMax to Pb p orbitals that mostly constitute the CBMin - high absorption coefficients are observed for LHPs (Figure 1.3). From the absorption spectrum, it is also possible to judge the presence of the deep trap states, specifically, from the decay of the absorption coefficient. For example, the absorption coefficient of MAPbI3 undergoes a purely exponential decay.52 A long wavelength edge of the absorption slope with exponential decay is called an Urbach tail, which is a broadening of the absorption edge due to the electron (or exciton)-phonon interactions. A structural disorder typically results in the broadening of the

Urbach energy. The sharp absorption edge of MAPbI3 (Figure 1.3) points towards a good electronic quality.

Figure 1.7. (a) Most common defects in LHPs: vacancies, interstitial, and antisite atoms, pictured in order of increasing formation energy and their depths in the bandgap. (b) A scheme of the electronic band structure of typical defect-tolerant, like LHPs, and defect-intolerant (such as CdSe, GaAs) semiconductors. (c) A scheme of polaron formation: the charge carrier (electron or hole) interacts with a local structural deformation of the Pb-Br framework. Reprinted with permission from Ref.78.Copyright 2018, Springer Nature.

1.5. Layered lead halide perovskites Use of long alkyl-chain ammonium cations as an A-type cation often results in a formation of layered LHPs. Moving to lower dimensionality compounds lifts a limitation on the size of the organic amines and a broader scope of possibilities appears regarding a choice and functionalization of the organic molecular cations. A comprehensive review of 0-2D perovskite-related structures was written by B. Saparov and D. B. Mitzi in 2016, where an existing variety of metal-halide perovskite-related structures was presented.79 Thus, the lower dimensionality perovskite structures can be depicted as derived from the aristotype 3D perovskite lattice: using molecular scissors (large organic cations) the 3D perovskite lattice is cut along certain crystallographic planes, which typically are (100), (110) and (111) (Figure 79 1.8). The general formula of layered LHP compounds is A2PbX4. The organic cations, which are usually long alkyl-chain amines, are situated in between the Pb-X layers and, due to the van der Waals interactions between the organic molecules, hold the inorganic layers together.

In 1991 J. Calabrese et al. found that 3D perovskites, for instance, MAPbI3, and 2D + compounds, (RNH3)2PbI4 (RNH3 – typically a long alkyl-chain ammonium cation), are two

Introduction | Olga Nazarenko extreme cases, whereas series of mixed (RNH3)2MAn-1PbnI3n+1 (n corresponds to the thickness of Pb-I layers) compounds can be obtained with a different thickness of Pb-I layers (Figure 1.8 on the right side).80 J. Calabrese et al. synthesized a bright red compound of

(C9H19NH3)2MAPb2I7 with an absorption feature at 2.21 eV, which is larger than the optical 80-81 bandgap of MAPbI3, 1.52 eV. The highest 2D member so far was found by C. C. Stoumpos 82 et al. in 2017 (n-BuNH3)2MA4Pb5I16 with n = 5. Fascinating about such compounds is a possibility to tune their optoelectronic properties by changing the thickness of the metal-halide slabs and the atomic number of the halide ion. For example, the excitonic absorption of (n-

BuNH3)2(MA)Pb2I7 with n = 2 is 2.17 eV and of (n-BuNH3)3(MA)2Pb3I10 with n = 3 is 2.03 eV, whereas changing to bromide - (n-BuNH3)3(MA)2Pb3Br10 results in the bandgap of about 2.43 eV.2, 83 Fine tunings of the bandgap can be achieved through the variation in the length of the alkyl chain in the alkyl-chain ammonium cations, for instance, the bandgap of C4H9NH3PbBr4 84 is 3.3 eV and of C7H15NH3PbBr4 – 3.38 eV. Therefore, such systems are very interesting for fundamental studies. Within this Ph.D. project (presented in Chapter 5) research was conducted + + on layered Pb-X (X = Br, I) systems with mixed inorganic (Cs ) and small organic [C(NH2)3 ] cations. Similar to oxide perovskites, among layered LHPs Ruddlesden-Popper (RP) and Dion- Jacobson (DJ) phases can be distinguished.2, 85 RP phases consist of perovskite layers parallel to (100) plane, with a general formula of An+1BnX3n+1 and the simplest member adopting K2NiF4 crystal structure, with a displacement of the layer comparing to the neighboring layers by (a+b)/2 vector (where a and b are ideal cubic cell vectors). It should be mentioned that often in 2 LHPs, for instance, in (n-BuNH3)2MA2Pb3I10, layers are displaced by an a/2 vector, making such compounds rather RP-related phases. When 2A cations are substituted by one A’ cation,

Figure 1.8. Layered perovskites depicted as derived from the 3D perovskite crystal structure, on the example of MAPbI3 (in the middle, crystallographic data for this compound was downloaded from the ICSD database with a code 238610). (n-BuNH3)2MA2Pb3I10 (ICSD code 3 252316) and (n-BuNH3)2PbI4 (Ref ) are obtained by cutting the 3D perovskite lattice along (100) crystallographic plane (on the right side). (C6H13N3)2PbBr4 (CCDC code 283056) results from the cut of the 3D perovskite lattice along (110) plane (on the left side).

Introduction | Olga Nazarenko

DJ phases result, with a general formula A′An−1BnX3n+1. L. Mao et al. obtained DJ series using 3-(and 4-) (aminomethyl)piperidinium (3AMP and 4AMP) as an A’ type cation: 85 AMP(MA)n−1PbnI3n+1, where n = 1−4. Separated by isolating large organic ammonium cations (the barrier), M-X slabs can be viewed as 2D quantum wells. The energy of the electron-hole pair (exciton) within the M-X layers is known to be four times larger than in extended 3D networks. The exciton binding energy (Eb) is also influenced by the dielectric constants of the well and the barrier. Thus, εb < 30 εw results in Eb(2D) > 4Eb(3D). This is a result of dielectric confinement - penetration of the electromagnetic field into the barrier medium with smaller dielectric constant. The Eb can be estimated from the absorption spectrum, more precisely, from the excitonic enhancement of the absorption edge, using Elliot’s theory of Wannier excitons, as was described in the work of M. 86 Saba et al. Another way is to measure an absorption spectrum at low T, where both Eb and Eg can be estimated.30, 87 A low T absorption spectrum of 2D LHPs typically shows a peak and a step-like absorption feature. The difference in Eg (estimated from the step-like feature) and a maximum of the peak gives Eb. Eb for layered LHP is typically on the scale of hundreds of meV and, therefore, these stable excitons are observed even at RT (higher than thermal energy at RT, which is ~25 meV). Excitons in 2D LHPs were proven to be of Wannier type (the exciton radius exceeds the lattice constant).88 A formation of stable excitons leads to higher probabilities of radiative recombination. Consequently, these compounds might be promising for light-emitting applications. The bandgap of layered LHP compounds is determined by the nature of the halide ions, the thickness of B-X layers, and distortions of BX6 octahedra (Figure 1.9a, b). The hydrogen- halogen interactions between metal-halide framework and organic cations influence a tilting and rotation of the B-X octahedra. B-X-B angles, as was studied on the example of tin(II) iodide perovskites in the work of J. L. Knutson et al., have the most profound influence on the electronic structure of hybrid metal-halide perovskite compounds.89 Ease of deformability of the lattice of layered LHPs, especially the ones with corrugated Pb-X layers, results in an increased probability of trapping of excitons on deformations created in the excited state. Such events lead to STEs and typically resolve in a broadband emission.90 Thus, among 2D LHPs few white-light emitting compounds were discovered, for instance, 2,2’-

(ethylenedioxy)bis(ethylammonium) lead bromide – (EDBE)PbBr4 with PLQY of 9 % at 300 K and 85 % at 105 K.91 The emission from STEs is characterized with long lifetimes, on the scale of hundreds of nanoseconds. The mechanism of formation of STE is not yet understood. 2+ It was suggested that analogously to PbBr2 crystals, ST electrons could be localized on Pb 3+ 2+ - forming a Pb2 dimer with another Pb , whereas ST holes could localize on Br and form Br2

Introduction | Olga Nazarenko

Figure 1.9. (a, b) Distortions in the metal-halide perovskite layers (on the example of Cs[C(NH2)3]PbI4 presented in Chapter 4 of this Ph.D. project). Deviation of B-X-B angles from 2+ ideal 180 ° can be described as in- (in) and out-of-plane (out) angles, while B cations define the plane. This deviations, especially those characterized by out, correlate with a change in the bandgap values, as they describe the efficiency of the overlap of metal-halide orbitals.90

1.6. The potential applications of LHPs These days, a lot of scientific efforts are oriented on the work on renewable energy technologies. According to the BP statistical review of the world energy, the part of power obtained from renewable resources grew by 17 % in 2017, which is higher than the 10-year average.92 Harvesting of sun-energy is, in particular, enabled by semiconductor materials that absorb the electromagnetic radiation from the sun and convert it into the electrical current. A boost in research on the solar energy harvesting with metal-halide coordination complexes was given by success of the LHPs as a photon to energy converters. Thus, according to the efficiency chart presented by the National Center for Photovoltaics, solar cells based on mixed organic- inorganic lead halides reached an efficiency of 23.3 %.93 The mystery of the success of LHPs in solar cells was uncovered (although, not yet all the aspects are fully understood) after about 6 years of intense research. Thousands of groups all around the world were actively working on this topic. Multi-cation and multi-anion systems were studied in solar cells with various device architectures that use light into energy (and vice versa) conversion. As was discussed in the above-presented sections, LHPs have suitable band gaps for a solar cell application, for instance, about 1.6 eV for MAPbI3 and 1.4 eV for FAPbI3. These are direct bandgap semiconductors with high absorption coefficients, and a clear exponential decay of absorption coefficients, implying a high absorption of the solar radiation with thin compounds layers and high electronic quality of the compounds correspondingly. The Eb < 25 eV (thermal energy at RT), therefore excitons dissociate into free carriers at RT. As the CBMin and VBMax are made of antibonding orbitals, it was calculated that the vacancies that have low formation energy would be the major defect type; vacancies lie in the bands or near the band edges creating shallow defects. LHPs is a very dynamic system: the anion exchange happens very fast, the halide ions migrate under bias (although, this results in I-V hysteresis), the phonons couple with charge carriers forming large polarons that prevent charge carriers from scattering resulting in a long lifetime and diffusion lengths of the charge carriers - an efficient charge transport. The mobility in LHPs is moderate due to polarons formation. The reorientation of the organic cation on the ps scale gives an impact to the dielectric response. Therefore, due to a combination of all these properties, LHPs thin films are showing good performance in solar energy harvesting.

Introduction | Olga Nazarenko

Owing to a high Z (in particular those based on iodide and bromide anions), density, high resistivity, high μτ product, as well as little charge trapping, LHPs are promising compounds for X-rays and γ-rays radiation detection.94-100 The hard radiation detection is a necessity in the research, as well as in industry (for instance, for medical diagnostics, security, and in a field of nuclear energy).101 Typically, the energy of the radiation source, its amount, and location should be determined. To solid radiation detectors belong semiconductors (for instance, CZT, and CdTe) and scintillators (such as NaI:Tl).101-102 Interaction of X-rays and γ- rays with high-Z semiconductors results in an electron-positron pair, photoionization, and Compton scattering.103 The optimal bandgap values for semiconductor detectors fall into 1.5 – 5 eV range.103 Nowadays, commercial devices are based on CZT crystals that are grown from the melt by Czochralski method. The challenges met during the crystal growth of CZT are non- stoichiometry, and inclusions,104 as well as high production costs. Therefore, a search for new low-cost hard radiation detector materials is highly relevant. LHP single crystals can be grown 95, 97 99 from solution, or likewise from a melt (by Bridgman method for fully inorganic CsPbBr3). One of the important working parameters of the hard radiation detector is an applied voltage.

CZT, for instance, can work at 500 – 1000 V, whereas solution grown MAPbX3 (X= Br, I) and 94 241 FAPbI3 (cubic phase) withstand 20-30 V. Nevertheless, energy-resolved signal from Am 94 (Eγ = 59.6 keV) source was obtained using 3 mm FAPbI3 single crystals at a bias of 23 V. FAPbI3 appeared to have smaller dark currents than MAPbI3 and a higher μτ product than best -2 2 -1 -2 2 -1 -2 2 -1 quality CZT and MAPbI3 - 1.8×10 cm V versus 0.91×10 cm V and 1×10 cm V 94, 105 correspondingly. Unfortunately, cubic FAPbI3 is not stable at RT, but a partial substitution of FA for Cs+ and I- for Br- ions leads to higher stability of the cubic phase at RT. An efficient 241 137 γ-rays counting ( Am and Cs sources) was obtained with Cs0.1FA0.9PbI2.8Br0.2 single crystals studied within this Ph.D. project presented in Chapter 3.95 A good spectral resolution, which is comparable to the commercial CZT (3.8-3.9 % versus 4.1 % accordingly) was reported 99 by Y. He et al. in 2018 for CsPbBr3 grown by Bridgman method. CsPbBr3 is a p-type -3 2 -1 semiconductor with μhτ = 1.34×10 cm V . The scientists used an asymmetric electrode design - Au and Ga electrodes - that largely (about three orders of magnitude) decreased the dark currents.99 The chemical nature of contacts plays an important role in the final device performance; for instance, the use of Ag contacts leads within some time to a formation of insulating AgX salts. The absorption of hard radiation by scintillators produces light emission. Therefore, such parameters as high quantum efficiency and short decay times are the most critical requirements for the scintillating material. The mechanism of scintillation is complex; in a few words, it can be described as a down-conversion of hard radiation to UV-visible photons. The most studied systems are based on a transparent insulating compound (matrix) and inclusion, a luminescence center, e.g., NaI:Tl, ZnO:Ga, YAlO3:Ce, etc.. From LHPs layered compounds were probed as scintillators, as the exciton binding energy is > 4 times higher than the Eb in 3D LHPs. The decay time of 0.75 ns (fast component) was detected for free-excitonic transitions in (C6H13NH3)2PbI4 and a 16 ns decay time for a slow component (assigned to the emission from trapped excitons on defects), which is still ten times faster than in NaI:Tl. Although, if the 106 relative output for NaI:Tl is 100 % then for (C6H13NH3)2PbI4 it is 11 %.

Introduction | Olga Nazarenko

Lower energy electromagnetic radiation (visible region) is absorbed by LHPs when the energy of the photons is ≥ Eg generating a photocurrent. At RT in LHPs, the electron-hole exists as free-charges, as mentioned above because the exciton binding energies are low (few meV in 107 MAPbI3 ). Depending on the way of charge separation in the semiconductor, two types of photodetectors can be distinguished: a photoconductor (the semiconductor is placed between two ohmic contacts, and applied bias ensures a charge collection) and a photovoltaic photodetector (work is based on the p-n junction, a p-i-n construction, or Schottky junction).108 There are numerous figures of merit that describe a photodetector’s performance, such as the external quantum efficiency (EQE - photon-electron conversion efficiency), the photoresponsivity, the detectivity, the response time, etc..109 LHPs are easy to obtain as thin films and single crystals from solution, they have high absorption coefficients, tailorable optoelectronic properties, excellent transport characteristics (discussed in section 1.2-1.4), and are highly compatible with flexible substrates, which make them attractive compounds for photodetectors. The photoresponse of MAPbX3 (X = Cl, Br, I and mixtures thereof), CsPbBr3 as thin films, single crystals, and nanoparticles was studied for potential use in photodetectors using various device architectures.110 For instance, the study of S. Yakunin et al.111 showed that a detector made of single crystals of MAPbCl3, MAPbBr3, and MAPb(Br/I)3 stacked on top of each other (for an effective photocurrent generation) is promising for color sensing. The crystals were isolated from each other with a polymer film to prevent an anion exchange. Due to high absorption coefficients, the pixel size and layer thickness of the MAPbX3 detector would be much smaller than Si n- or p-type doped layers commercially used as color sensors (Figure 1.10a,b). LHPs are probed for application in various devices, not only as a photon to current converters but also as emitters. According to O. D. Miller, E. Yablonovitch, and S. R. Kurtz, a good solar cell should also be a good light emitting diode (LED), meaning that light photons

Figure 1.10. (a) Absorption spectra of MAPbX3 films comparing to Si films. (b) The minimum pixel sizes (defined by the penetration depth of light) of stacked Si layers comparing to MAPbX3 LHPs (scale bar – 10 μm). 111 Reprinted with permission granted by the Creative Commons Attribution 4.0 International License: NPG Asia Mater. 2017, 9, e431.

Introduction | Olga Nazarenko absorption at the open circuit should result in the efficient radiative recombination.112 The commercial LED - the most energy-efficient lightning source nowadays - are based on direct bandgap III-V group semiconductors such as GaN, InGaN, GaAs. Regrettably, they are expensive to produce. Low-cost, efficient lightening sources could reduce the overall energy consumption. Organic LEDs and colloidal quantum dot LEDs are following III-V compounds regarding efficiency, although there are still some unresolved issues, such as non-radiative recombination processes and high production costs.113 The light emission in LEDs comes from the recombination of the electrically injected carriers. Therefore, the LED materials should have a high EQE, an excellent current efficiency (CE) and brightness. LHPs are highly color pure (tunable wavelength of the emission and narrow FMHM) and high PLQY.114-115 LHPs have tunable bandgaps, good charge transport characteristics and charge-injection characteristics meet the necessary requirements. 2D and 3D, as well as nanoparticles of LHPs were probed for the LED application.113, 116 H. Cho et al. reached in their study 8.53 % EQE, CE of 42.9 cd A-1 -2 117 and luminance 15 000 cd m for a green LED based on 3D MAPbBr3 thin films in 2015. The films were prepared using excess of MABr. Whereas, F. Yan et al. used MAPbBr3 nanoparticles and obtained LEDs with 10.67 % EQE and 25’410 cd m-2 luminance.118 Mixed halide compositions are often employed to achieve different emission colors, although, unfortunately, phase segregation occurs under illumination and bias.119-120 A green-emitting perovskite LED

(based on thin layers of CsPbBr3, MAPbBr3, and MABr) with a record EQE of 17 % and 14’000 cd m-2 luminance and low turn-on voltages (2.7 V) was reported by K. Lin et al. in Nature in 2018.115 On the other hand, blue perovskite LEDs have so far little efficiencies, for instance, a blue LED based on a 2D hybrid lead bromide compound has a luminance of 33’000 cd m-2 and 121 2.6 % EQE.

1.7. On Sn(II) and Ag(I)/Bi(III) halide compounds with perovskite crystal structure The materials based on lead are integrated into the modern life, these are lead-acid batteries (invented in 1859 by Gaston Planté), pigments (for instance, yellow PbCrO4), solder, as well as shielding from X-rays. Although, due to high toxicity of lead and its compounds when penetrating the human body as a solution, dust or vapor, the amount of lead in the products is regulated. According to the RoHS, the limit of Pb is maximum 1000 ppm (or 0.1 %) by weight in the homogeneous material. To the list of Pb exemptions - according to the Annex III accepted by the European Commission in March 2018 - belong alloys of lead with steel (0.35 % Pb by weight), Al (up to 0.4 % Pb by weight), Cu (up to 4 % Pb by weight), lead in glass and ceramic, and few others. Among the other restricted by RoHS substances are Hg (<100 ppm), Cd (<100 ppm), etc.. Thus, there are sterner demands regarding Cd metal content comparing to Pb. According to the calculations, for application of LHP nanoparticles in the TV displays, the condition of <1000 ppm of Pb by weight is easily met, as minimal quantities are sufficient for the efficient device performance.78 On the other hand, in the solar cells based on LHPs absorbing layer, the amount of lead is more than 10 % by weight of the device.122 Use of LHPs in tandem solar cells might be an answer to the reduction of the mass percent of the lead in the

Introduction | Olga Nazarenko final product. The danger of LHPs is in their solubility in water. Therefore, good encapsulation should be developed before LHPs can be used in commercial devices. One way to deal with LHPs toxicity issue is to develop a good encapsulating material, as stated above, on the other hand, alternative less toxic metal-halide perovskites are searched that would possess similar optoelectronic properties, such as based on Sn(II), Sb(III), Bi(III) or a mixture of heterovalent Ag(I)/Bi(III) cations. Such requirements as the possibility of film deposition from a solution, a direct bandgap, high absorption coefficient, low carriers effective masses, and a defect tolerance are issued to the metal-halide semiconductors to substitute LHPs.123 Naturally, the first choice to substitute Pb(II) lies on the cations with analogous electron configuration to Pb(II) and within the same group - Sn(II) and Ge(II). The study on 124 hydride tin(II) halide perovskites goes back to 1988 when MASnBr3 was first obtained. Unfortunately, ease of oxidation of Sn(II) and especially Ge(II) introduce a necessity to perform synthesis, further manipulations with a product, and a film deposition under inert conditions.

The other problem is hydrolysis, for instance, SnCl2 hydrolyzes in aqueous solutions at pH 6-7 according to the next equation:125

2+ − → + 2푆푛 + 3퐻2푂 + 퐶푙 ← 푆푛(푂퐻)퐶푙 + 푆푛(푂퐻)2 + 3퐻 (1.12) Despite problems with oxidation and hydrolysis, hybrid tin(II) halide perovskites are interesting compounds for research, e.g., they can have semiconductor or even metallic conductivities.126- 127 A study of the conductivity of (n-BuNH3)2MAn-1SnnI3n+1 (n = 1–5) series (p-type semiconductors) was performed by D. Mitzi’s scientific group in the 90th.127 The scientists concluded that when the thickness of Sn-I layers is growing to n > 3, the resistivity drops and a + conductivity gets a metallic character. Interestingly, if n-BuNH3 cation facilitates a formation of flat (100) Sn-I layers in (n-BuNH3)2MAn-1SnnI3n+1 (n = 1–5), iodoformamidinium forms in the Sn-I system 2D [CI(NH2)2]2MAmSnmI3m+2 (m = 2-4) with corrugated (a zig-zag shape) (100) Sn-I layers.126 The difference in these crystal structures underlines a structure directing role of the organic cation. The largest zig-zag step was achieved in α-[NH3(CH2)5NH3]SnI4 <330> that forms when CH3NH3I is added to the synthetic solution of SnI2 and [NH3(CH2)5NH3]I2 in hydriodic acid (Figure 1.11a).128 Another interesting example of the templating role of organic cations was observed on 2-trimethylammonioethylammonium tin iodide – + [(CH3)3N(CH2)2NH3]SnI4. The organic cation with two amino groups has one bulky (CH3)3N functional group that causes tilting of Sn-I octahedra (Figure 1.11b).129 Sn-I layers accommodate the organic cations in the manner that short I···I interlayer distances – 4.19 Å are enabled (van der Waals distance 4.00 Å).

A study on the partial substitution of Sn(II) in the MASnI3 semiconductor by Pb(II) revealed that the bandgaps change nonlinearly and even smaller bandgaps were observed in mixed compounds (1.17 eV for MAPb0.5Sn0.5I3) comparing to parental MAPbI3 (1.55 eV) and 130 MASnI3 (1.30 eV). According to the recent study, the reason lies in the lattice distortions and 131 spin-orbit interaction effects. Interestingly, the VBMax in MAPbxSn1-xI3 (x < 0.875) was calculated to consist mostly of Sn 5s and I 5p orbitals.131

Introduction | Olga Nazarenko

Figure 1.11. (a) A crystal structure of α-[NH3(CH2)5NH3]SnI4 (the crystallographic data, cif file, downloaded from CCDC database, deposition number 105852) and

[(CH3)3N(CH2)2NH3]SnI4 (b) (the crystallographic data, cif file, downloaded from Z. Xu, D. B. Mitzi, D. R. Medeiros, Inorg. Chem. 2003, 42, 1400-1402).

+ Recently, a raw of compounds A2(Ag/Bi)X6 (A = Cs , X = Br, Cl) and (MA)2AgBiBr6 were synthesized, where a 3D network is formed of corner-sharing Bi-X and Ag-X (X = Br, Cl) octahedra alternating in a periodic manner, with Cs+ and MA cations filling the voids within the lattice (Figure 1.12). These compounds are called double perovskites. The inspiration for their synthesis was a Cs2NaBiCl6 compound with an elpasolite (mineral K2NaAlF6) crystal structure,132 where Na+ (crystal ionic radius 116 pm) was substituted for Ag+ (129 pm).

According to the calculations, the CBMin of A2(Ag/Bi)X6 are made of antibonding Bi 6p/X p orbitals and VBMax - of antibonding Ag 4d/X p orbitals. The Bi 6s2 pair has a small contribution to the VBMax, and Ag 5s orbitals make a small contribution to CBMin.133 Unfortunately, these double perovskites have indirect bandgaps that lead to smaller absorption coefficients than in the direct bandgap LHPs. According to calculations, a reason for an indirect bandgap is an electronic mismatch between Bi and Ag. Cs2AgBiBr6 has a bandgap energy of 2.19 eV.134 This compound can be crystallized from hydrobromic acid water solution as orange truncated octahedra. It is stable under ambient conditions in the dark, but under ambient conditions and exposure to light for some weeks the compound decomposes. According to calculations, Cs2AgBiI6 should possess a smaller bandgap and smaller me comparing to its Br 135 136 and Cl analogs, but Cs2AgBiI6 is thermodynamically unstable and decomposes:

2퐶푠2퐴푔퐵푖퐼6 → 퐶푠3퐵푖2퐼9 + 2퐴푔퐼 + 퐶푠퐼 (1.13) It was calculated that if Ag+ is substituted by s elements, such as Tl+ or In+ (elements with valence s states), a direct bandgap will result.136 Compounds with indirect bandgaps might still be good absorbers if there is a direct transition possible a bit higher in energy. In Cs2AgBiX6 (X = Br, Cl) according to the calculations, such a direct transition is > 0.5 eV higher in energy.136 + 3+ + 3+ The study on substitutional doping of Ag and Bi sites in Cs2AgBiBr6 by Tl and Tl 137-138 correspondingly was performed by A. H. Slavney et al.. Cs2AgTl(III)Br6 is a black colored compound with a very small direct optical bandgap of 0.95 eV. Cs2AgTl(III)Br6 is an n-type semiconductor conditioned by cationic bromide vacancies. The toxicity of Tl is an issue

Introduction | Olga Nazarenko for practical applications of these compounds, but the study showed a possibility to synthesize direct and narrow bandgap perovskite semiconductors. The most common way to obtain smaller bandgaps was to change to the higher atomic number halide, but the later work shows another way: alignment of energy bands. The A2(Ag/Bi)X6 system is rather complicated, and preparation of films of a reasonably pure phase is challenging, therefore, until this moment, a limited work was done on solar cells based on double perovskites. Although, in a recent publication by C. Wu et al. and E. Greul et al. films of pure Cs2AgBiBr6 phase were obtained after the deposited films were additionally annealed.139-140 The best efficiency obtained so far 140 for solar cells based on films of Cs2AgBiBr6 is 2.5 %. A part of this Ph.D. project was dedicated to searching for an alternative to LHPs metal- halide semiconductors. This study is presented in Chapters 6 and 7.

Figure 1.12. The crystal structure of Cs2AgBiBr6 (the crystallographic data, cif file, downloaded from ICSD database, the card number 239875).

1.8. Scope and outline of the dissertation Essentially, this Ph.D. project was dedicated to fundamental research on mixed organic-inorganic lead(II) halide compounds with perovskite and perovskite-like crystal structures, as well as on ternary and quaternary tin(II) and bismuth(III) halide compounds. In particular, the research was directed on the determination of the crystal structures of novel semiconductors and studying their optical and electronic properties along with testing the selected synthesized compounds for hard radiation detection and photoconductivity. The purpose of the Introduction chapter was to provide a concise tour of the field of semiconductors. In particular, a short historical outline, an overview of the dynamic field of hybrid LHPs, an introduction into the bandgap theory of semiconductors, as well as a short overview of organic-inorganic tin halide perovskites and heterometallic silver/bismuth halide double perovskites, possible alternatives to LHP semiconductor compounds. In the following Chapter 2 methods and techniques are described that were used to synthesize and characterize the novel semiconductor compounds in this Ph.D. project. The results are presented

Introduction | Olga Nazarenko in Chapters 3-7, four of which are based on first author publications by the Ph.D. candidate. Lastly, in Chapter 8 conclusions to the obtained results, as well as possible further research directions are outlined. A list of first author and joined publications is given in the Curriculum Vitae. Single crystals of cesium-formamidinium lead halide perovskites: solution growth and gamma dosimetry (Chapter 3). A significant boost in the research on LHPs happened in

2012 when solar cells based on MAPbI3 absorber were fabricated with the resulting efficiency 141 of 10 %. However, due to chemical stability issues of MAPbI3, more stable MAFAPb(XX’)3 142-143 and FAPbI3 were found to be more promising candidates for photovoltaics application.

FAPbI3 has excellent optical and electronic properties, but its perovskite phase (α-FAPbI3) is not thermodynamically stable at RT and undergoes a phase transition within one to seven days to a broad bandgap 1D polymorph, -FAPbI3. The reason lies in the spatial geometry and dynamics of FA cation within the cavities of Pb-I lattice causing destabilization of the 3D perovskite structure. The researches endeavored to stabilize thin films α-FAPbI3 at RT by modifying the GTF. Thus, FA and I- were partially substituted by Cs+ and Br- that were added to the solution.144-145 These thin films were characterized with kinetic stability (defined as the time before the phase transition), homogeneity, and probed in solar cells. This research was an inspiration for the study of intrinsic properties of single crystals of FA1-xCsxPbI3-yBry (x = 0- 0.1, y = 0-0.6). The composition of selected single crystals was accurately determined with the elemental analysis (CHN, Br, I). The first part of this Ph.D. project was, therefore, dedicated to the growth of the single crystals of various FA1-xCsxPbI3-yBry (x = 0-0.1, y = 0-0.6) compositions from γ-butyrolactone solution, and studying their photophysical properties. The powder XRD method was chosen to monitor the beginning of the phase transition. The absorption and PL studies were used to estimate the change in the band gap of the resulting compounds dependent on the Cs and Br concentrations. The photoresponse measurements were used to estimate the charge-carriers mobility × lifetime product and further the compounds were characterized with a specific resistivity and tested for use as γ-rays detectors. Luminescent and photoconductive layered lead halide perovskite compounds comprising mixtures of cesium and guanidinium cations (Chapter 4). After the success of 3D LHPs as excellent light to energy converters, attention was as well drawn to lower dimensionality compounds. Layered ternary (and higher) LHPs have higher exciton binding energies, and the charge transport is limited to two directions. Moving to 2D systems gives a larger choice of A-type cations and the possibility to tune the bandgap not only through the halide change but also through the change in a thickness of lead-halide layers. Fine tuning of the bandgap of 2D LHPs is also possible through the change in the length of the alkyl chain in long alkyl-chain ammonium cations, and additional interactions of organic cations, through incorporating different functional groups, with inorganic layers. At the end of 2015, research + in the field of 2D LHPs was largely concentrated on (RNH3)2PbX4 compounds (RNH3 – typically a long alkyl-chain ammonium cation). The idea of this part of the Ph.D. project was to obtain layered compounds with minimum interlayer distance to facilitate a higher density of electronic states and therefore, possibly smaller bandgaps, as well as to modify the thickness of the anionic slabs of obtained compounds. For this study, an inorganic Cs+ and a high symmetry

Introduction | Olga Nazarenko organic cation guanidinium were chosen. Three layered CsGPbI4, Csn-1GPbnBr3n+1 n = 1, 2 compounds were obtained that crystallize as Ruddlesden-Popper related phases. Both inorganic Cs+ and organic G cations are incorporated in the crystal structure, situated in the interlayer space, and are periodically alternating each other. In the Cs-G-Pb-Br system, it was possible to + obtain a Cs2GPb2Br7 compound with two Pb-Br octahedra thick layers, and Cs is filling the voids within the Pb-Br layers. The synthesized compounds were characterized with single crystal XRD, powder XRD, as well as their thermal stability was studied. The band gaps were estimated from the absorption spectra using Eliot’s theory. The crystal structure of CsGPbI4 was compared to (n-BuNH3)2PbI4 to understand the influence of the density of the structure and the distortions of the Pb-I octahedra on the final bandgap of the compound. Thereupon, this study included a detailed crystallographic analysis of the obtained compounds, exploration of their optical properties, and photoconductivity, as well as calculations of their electronic structures. Guanidinium-formamidinium lead iodide: a layered perovskite-related compound with red luminescence at room temperature (Chapter 5). A unique combination of excellent electronic properties in LHPs allowed for exceptional performance in solar cells, LEDs, photodetectors. To optimize the performance, researches used multi-cation and multi-anion systems: different mixtures of MA, FA, G, Cs+, Rb+ with Pb(II)XX’, where X = I- and X’ = Br- or small quantities of Cl-. The part of this Ph.D. project was oriented on probing such compositional ranges for the formation of mixed phases with other than 3D perovskite crystal structures. Thus, Chapter 5 of this Ph.D. project is dedicated to FA-G-Pb-I system, where a common FAGPbI4 phase was found. FAGPbI4 has a layered crystal structure, where Pb-I layers made of Pb-I corner-sharing octahedra and have a stair-like arrangement. FAGPbI4 was characterized with single crystals XRD, PL, photoconductivity, and thermal analysis. The emission properties of FAGPbI4 appeared to be interesting and complex. A PL spectrum is a convolution of several emission bands with PLQY of maximum 3.5 % at RT. To explain the origin of these bands, a 207Pb SSNMR was measured to detect if any possible amorphous or nanocrystalline contaminants are present in the FAGPbI4 sample and give the origin to some of the PL bands. The temperature dependent PL was measured to understand the behavior of the emission processes with temperature. Additional measurements, such as TR-PL, special resolution PL and TR-PL, and photoconductivity were employed to learn the time of the decay of the PL and if the PL spectrum differs from spot to spot on the single crystals. Guanidinium and mixed caesium-guanidinium tin(II) bromides: effects of quantum confinement and anisotropic lattice distortions (Chapter 6). Due to lead toxicity issues, research encircled other metal-halides coordination compounds based on Sn2+, Sb3+, Bi3+, Bi3+/Ag+. As discussed in the Introduction, section 1.2, according to the RoHS, the limit of Pb in the homogeneous material by weight is maximum 1000 ppm, with some exceptions, like lead alloys, ceramics, glass, and few others. This limit poses restrictions on the amount of LHPs that can be used in devices. The concern with LHPs lies in their solubility, for instance, the solubility of PbI2 (a product of degradation of MAPbI3) is about 0.076 g in 100 ml water at 298 K.146 Therefore, research on less toxic metal-halide semiconductors is desired. In Chapter 6 of this Ph.D. project, investigation of phases within the G-Sn(II)-Br system, the influence of the additional Cs+ cation on the crystal structure, and a comparison with analogous lead

Introduction | Olga Nazarenko compounds (presented in Chapter 4) were conducted. The study showed that in the G-Cs-Sn-

Br system there is a complex equilibrium of two common phases Csn-1GSnnBr3n+1 n = 1, 2 with CsSnBr3, and G2SnBr4. A change in the overall concentration or the concentration of one of the components results in the different ratio of the compounds in the precipitate. However, it was possible to obtain a Cs2GSn2Br7 as a major phase. G2SnBr4 and Csn-1GSnnBr3n+1 n = 1, 2 were characterized with a single crystal XRD, powder XRD, absorption, PL, and their electronic structures were calculated. Obtained new ternary and quaternary Sn(II) halide complexes: Csn- 1GSnnBr3n+1 n = 1, 2 and G2SnBr4 were compared to their lead analogs regarding crystal structures, optical properties, and electronic band structures. In this project, the addition of I- ions to Cs-Sn-Br system was also explored. Thus, the addition of a small quantity of I- to the synthetic solution with [I]/[Sn]=19/1, lead to the formation of the new layered mixed halide compound Cs8Sn6Br13I7, that was studied with single crystal XRD, as well as EDXS (to prove the compositional homogeneity). Rubidium copper and silver iodobismuthates (Chapter 7). Another alternative to LHPs is ternary and higher Bi(III) halide coordination complexes. Bi(III) has the same electronic structure as Pb(II), but a different charge. E. T. McClure et al. showed a possibility to synthesize a 3D double-perovskite semiconductor Cs2BiAgX6 (X = Br, Cl) by substituting + + 134 + Na in the elpasolite crystal structure Cs2NaBiCl6 for Ag . Ag forms octahedral units with Br-, Cl- ions that are further connected through corner-sharing with Bi-X octahedra. The compound is typically described as having a double perovskite crystal structure: B(II) cation in the ABX3 network is substituted for a mixture of B’(III) and B’’(I) cations. The synthesis of iodide representative of Cs2BiAgX6 would be desired due to a smaller expected bandgap. However, no common phase was obtained so far, as it is thermodynamically unstable and typically Cs3Bi2I9 (0D) precipitates first. The aim of this part of the Ph.D. project was to study mixed phases in the Rb-Bi-Ag(Cu)-I system. Rb+ was chosen, as it facilitates a formation of 2D 147 Rb3Bi2I9, where Bi-I layers are made of corner-sharing octahedra. The objective was to link Bi-I anionic species with Ag(Cu)-I building blocks (tetrahedra, which are regular coordination + + polyhedra of Ag and Cu , or also octahedra of AgX6). The synthesis was performed in the water solution of hydriodic acid; and consequently, some water and hydronium molecules are present in the structure. Thereupon, one compound was obtained in each Rb-Bi-Ag-I-O-H and Rb-Bi-Cu-I-O-H system. The crystal structures were characterized by single crystal XRD. Solving the crystal structures was challenging, as a significant static disorder is present in these phases. The Cu, Ag and some I atoms were assigned based on the type of the coordination surrounding and the length of the bonds. The resulting compositions are not charge neutral. Therefore, further techniques, such as WD-XRD, EDXS were used to study the composition of the obtained compounds. The 63Cu SSNMR and EPR for Ag and Cu were used to learn more about the oxidation states of Cu and Ag. As the water molecules are present in the structure, the charge neutrality might be reached if water molecules are in fact hydronium ions. Further, 1H, 87Rb SSNMR were measured to understand the complex Rb-Bi-Ag-I and Rb-Bi-Cu-I systems better. It is possible to grow a few mm size crystals of the obtained compounds in Rb-Bi-Ag-I- O-H and Rb-Bi-Cu-I-O-H systems. Additionally, the compounds were characterized by absorption and specific resistivity.

Methods and techniques | Olga Nazarenko

Chapter 2. Methods and techniques

2.1. The growth of single crystals The novel compounds reported in this Ph.D. project were grown from a solution. In particular, an inversed temperature crystallization method (the crystals grow upon increasing the T, unlike often used crystallization upon cooling) was used to grow single crystals of

CsxFA1-xPbI3-yBry (x = 0-0.1, y = 0-0.6), presented in Chapter 3. The crystals were grown from

γ-butyrolactone (Tb = 477 K) by dissolving corresponding halides and heating the solution from about 353 K to 403 K with a speed of ~ 5 K per hour. The O. M. Bakr’s scientific group was the first one to apply the inverse temperature crystallization method to the growth of LHPs.49, 148 It was suggested that upon addition of γ-butyrolactone to the halide salts solvated ions as well as colloids result (that are not removed after the filtration through a 0.2 μm syringe filter). The colloids then dissociate upon heating creating an oversaturated solution.149 According to the study by P. K. Nayak et al. an important role in crystal growth plays a product of the hydrolysis of γ-butyrolactone - γ-butyric acid, which promotes the dissociation of the colloids.149 The hybrid tin(II) halide compounds, presented in Chapter 6 of this Ph.D. project were crystallized under inert conditions using the Schlenk technique (Figure 2.1). Crystals grow when the amount of the solute is larger than the equilibrium amount – supersaturated solution. Typically, a supersaturated state can be reached by 1) changing the temperature of the solution depending on the temperature coefficient of solubility; 2) slow solvent evaporation; 3) using a temperature gradient. In this Ph.D. project mostly method 1 was used. The nucleation is a first-order phase transition from metastable (supersaturated) phase of solute to a stable new phase. It can be homogeneous (without preferential nucleation sites and

Methods and techniques | Olga Nazarenko

Figure 2.1. Basic Schlenk setups used for the work with Sn(II) salts. The Sn(II) halide compounds were synthesized under inert conditions using the Schlenk vessel (a). Next, crystalline powders were separated from the solution by vacuum filtration under an Ar flow using a glass filter (porosity 3) with two joints (b). The obtained crystalline powder was then dried on the sand bath under vacuum using the setup shown on the right (c). random) or heterogeneous (occurring at specific nucleation sites, such as introduced seeds, surfaces, as a scratch on the wall of the reaction vessel).150 A solid body of the crystal is held together via bonds formed between atoms and molecules. The nature of the bonds can be described by four forces: covalent, ionic, metallic, and Van der Waals. The crystallization is driven by the decrease in the total energy of the system. The change in energy can be described by the equation: ∆퐺 = ∆퐻 − 푇∆푆 (2.1) where ∆퐻 is an enthalpy and ∆푆 entropy change of the system. For solutions it is convenient to use:

푛 ∆퐺 = −푆푑푇 + 푉푑푝 + ∑𝑖 휇𝑖푑푁𝑖 (2.2) where n is a number of constituents with i species and Ni particles, μ is a chemical potential. The free energy of the formed crystal seed is defined by the energy of the bulk and the surface:

∆퐺 = ∆퐺푏 + ∆퐺푠 (2.3) The energy of the surface is typically large; therefore, when the seed is smaller than a certain critical size, it dissolves. There are three mechanisms of crystal growth: 1) via 2D nucleation, 2) spiral growth with screw dislocations, and 3) a layer by layer growth.151 At a very high degree of oversaturation, nucleation might also go through the condensation of 3D particles in a random orientation into a polycrystalline block. A moving force in the crystal growth is a gradient of chemical potentials between the solid and the environment. This gradient is reduced by means of diffusion as the system tends to reach an equilibrium state.151-152 Although, only

Methods and techniques | Olga Nazarenko the ions at the interface take part in the crystal growth, therefore, intensification of the diffusion from the solution phase to the surface of the crystal will facilitate faster crystal growth. A motion of the ions in solution to the center of crystallization might occur due to molecular diffusion, natural or forced convection. When optimal conditions for crystal growth are found high-quality crystals with a fast growth rate can be obtained.

2.2. Characterization methods Powder X-ray diffraction patterns were collected in transmission mode (Debye- Scherrer-Geometrie) with a STADI P diffractometer (STOE&Cie GmbH), equipped with a curved Ge (111)-Monochromator (CuKα1 = 1.54056Å) and a silicon strip MYTHEN 1K Detector (Fa. DECTRIS). For the measurement, depending on the air-sensitivity of the prepared compounds, the ground powder was placed between the adhesive tape, Mylar foil or sealed in the glass capillary (Hilgenberg GmbH) with 0.1-0.5 mm diameter and 0.01 mm wall thickness. Measurements and data evaluation were conducted with WinXPOW (Version 3.0.1.13, STOE & Cie GmbH) and Match! (Version 1.11j, Crystal Impact GbR).

Single crystal X-ray diffraction measurements were conducted on Bruker Smart Platform diffractometer equipped with an Apex I CCD detector and molybdenum (MoKα = 0.71073 Å) sealed tube as an X-ray source. Crystals were a tip–mounted on a micromount with paraffin oil. Crystals of air-sensitive compounds were covered with paraffin oil during selection and measured under a liquid nitrogen stream at temperatures 100 - 250 K. The data was processed with the APEX3 (v2015.5-2, Bruker software)153 or CrysAlisPro, the structure solution and refinement were performed with SHELXS154 and SHELXL,155 respectively, which are embedded in Olex2 (v1.2, OlexSys Ltd.).156 For twin indexing, the cell_now algorithm (a part of the Bruker APEX3 software package) was used.

UV-Vis absorbance spectra of the microcrystalline powders were collected using a Jasco V670 spectrophotometer equipped with deuterium (D2) lamp (190–350 nm) for use in UV, a halogen lamp (330–2700 nm) for use in UV/NIR, and an integrating sphere (ILN-725) with a working wavelength range of 220–2200 nm. A white standard was a barium sulfate

(BaSO4). The measurements and analysis were performed using Spectra Manager software (Version 2.13.00, JASCO Corporation) including Spectra Management (Version 2.5.0.1) and Spectra Analysis (Version 2.15.2.1). The diffuse reflectance data were transformed into the 2 Kubelka-Munk function: F(R∞) = α/S = (1−R∞) /2R∞, where F(R∞) - Kubelka-Munk units, α - absorption coefficient, S - the scattering coefficient, and R∞ - the reflectance of an infinitely thick layer. The absorbance spectra of samples, presented in Chapter 4 were also estimated from reflectance and transmittance spectra collected from a thin layer of a powder dispersed in an optically transparent Teflon grease.

Photoluminescence measurements (measured by Dr. S. Yakunin) for samples described in Chapter 3 were performed with a Fluorolog iHR 320 Horiba Jobin Yvon spectrofluorometer equipped with a Xe lamp and a photomultiplier tube.

Methods and techniques | Olga Nazarenko

PL spectra for compounds presented in Chapters 4 and 5 were measured in a Joule– Thomson cryostat (MMR Technologies) operated in the temperature range of 78 - 300 K. PL emission was recorded with a heating rate of approximately 5 K/min. A 355 nm excitation source (a frequency-tripled, picosecond Nd:YAG laser, the model Duetto from Time- Bandwidth) and a CW diode laser with the excitation wavelength of 405 nm were used. Scattered laser emission was filtered out using dielectric long-pass filters with edges at 400 and 450 nm, respectively. The emission from the samples presented in Chapter 4 was collimated to an optical fiber and recorded at 1 K intervals with a spectrograph SP-2300 from Princeton Instruments coupled with a CCD array (LC100/M from Thorlabs). The emission from the samples presented in Chapter 5 was collimated to an optical fiber and recorded at 1 K intervals with a CCD spectrometer LR1 (Aseq Instruments). The PL spectra were corrected to the spectral sensitivity of the setup using Planck irradiation from a calibrated halogen lamp. PL spectra for compounds presented in Chapter 6 spectra were measured with a CCD fiber spectrometer (LR1, Aseq Instruments) with a 355 nm excitation source (frequency-tripled, picosecond Nd:YAG laser, Duetto from Time-Bandwidth). PL emission from the samples passed through a long-pass optical filter with an edge at 400 nm in order to reject the excitation laser line. PL spectra were corrected to the spectral sensitivity of the detection system.

For temperature-dependent PL measurements, a sample of CsSnBr3 presented in Chapter 6 was placed atop of a 4-stage Peltier cooling/heating element in an evacuated chamber with a quartz window. The sample temperature was adjusted and stabilized with an accuracy of 0.25 °C by a home-made electronic scheme based on an Arduino microcontroller and thermocouple sensor. The current through the Peltier was reversible. Thus the setup provided a wide working temperature range of -40 – 120 °C. This is an open-source project by the authors, deposited and described in details at https://www.researchgate.net/project/High-power- thermoelectric-cooler-TEC-controller-with-4-stage-Peltier-refrigerator-heater. In Chapter 6, for measurements of PL and PLE spectra at low temperature (77 K), the sample was encapsulated in a quartz tube filled with Ar gas and placed in a homemade cryostat. Low temperature (77 K) and RT absolute PLQY were measured with excitation at 340 nm using Quantaurus-QY spectrometer from Hamamatsu. The sample was encapsulated in a quartz tube filled with Ar gas.

PL measurements at lower temperatures (in the case of Cs2GSn2Br7, presented in Chapter 6), down to 5 K, were conducted in a helium cryostat and PL spectroscopy was performed by exciting the sample with a frequency-doubled Ti:Sapphire mode-locked laser delivering pulses of about 150 fs duration at 400 nm and a repetition rate of 80 MHz. The time- integrated PL was analyzed using a CCD-coupled grating spectrometer, whereas TR-PL traces were recorded with a streak camera.

Time-resolved photoluminescence measurements (measured by Dr. S. Yakunin) were performed using a time-correlated single photon counting (TCSPC) setup, equipped with an SPC-130-EM counting module (Becker & Hickl GmbH) and an IDQ-ID-100-20-ULN avalanche photodiode (Quantique) for recording the decay traces. The emission for compounds presented in Chapter 4 was excited by a frequency-tripled (λ = 355 nm), picosecond Nd:YAG laser, the model Duetto from Time-Bandwidth), externally

Methods and techniques | Olga Nazarenko triggered at a 78 kHz repetition rate. PL emission from the samples passed through sets of long- and short-pass optical filters selecting wavelength ranges 500-600 nm and 650-800 nm. The emission for a compound presented in Chapter 5 was excited by a BDL-488-SMN laser (Becker & Hickl) with a pulse duration of 50 ps and wavelength of 486 nm, the CW power equivalent of ~0.5 mW, externally triggered at a 1 MHz repetition rate. PL emission from the samples passed through a long-pass optical filter with an edge at 500 nm to reject the excitation laser line. 2 2 Σ푖=1휏푖 퐴푖 Average radiative lifetimes were determined as: 휏푎푣푔 = 2 , where Ai and τi are Σ푖=1휏푖퐴푖 corresponding amplitudes and exponential decay parameters in biexponential analysis. Absolute values of PLQY was measured at Quantaurus-QY spectrometer from Hamamatsu in powder mode.

Micro-photoluminescence spectra (measured by Dr. G. Raino) were recorded in a home-built setup by Gabriele Raino. The sample presented in Chapter 5, i.e., a single crystal of

FAGPbI4, was mounted on xyz positioning stages (SmarAct, SLS 32:32; 1 nm resolution). A 405 nm laser diode (PicoQuant, LDH-D-C-405; repetition rate of 2.5 MHz; excitation power density of 90 μJcm-2) was focused on the sample by an oil immersion objective (Olympus, 100X; NA = 1.4). PL was recorded by the same objective and analyzed by a monochromator (500 mm focal length; 300 lines/mm grating) coupled to an EMCCD camera (Princeton Instruments ProEM: 1600x400B_eXcelon3). For time-resolved PL, a single photon avalanche diode (PicoQuant PDM, 50 ps time resolution) was coupled at the exit port of the monochromator. TR-PL traces were recorded using a time-correlated single photon counting (TCSPC) module (PicoQuant, PicoHarp 300).

Photoconductivity measurements (performed by Dr. S. Yakunin) were performed with illumination from a tungsten lamp dispersed by an Acton SP2150 (Roper Scientific) spectrograph/monochromator. A mechanical chopper modulated the light at a frequency of 9 Hz. The sample was contacted by a soft conductive rubber from opposite sides along the shorter axis of the crystal and biased at 100 (or 50) V through these contacts using a Keithley 236 SMU. The signal, amplitude, and phase were measured across series of resistance by a Stanford Research 830 lock-in amplifier. The light intensity was controlled by a calibrated power detector (UM9B-BL, Gentec-EO). The current-voltage (IV) characteristics were measured using a Keithley 236 SMU.

Thermal analysis (TG and DSC) was performed using a Netzsch Simultaneous Thermal Analyzer (STA 449 F5 Jupiter). Measurements and data evaluation were conducted with a Netzsch software package (Version 6.1.0, NETZSCH-Geraetebau GmbH). A powdered sample (5-25 mg) was placed in an alumina crucible (without a lid) and heated under Ar gas flow (40 or 50 ml/min) to 800-850 °C with a rate of 10 °C/min.

207Pb solid-state nuclear magnetic resonance (measured by M. Aebli). All experiments were performed on a Bruker spectrometer (16.4 T) equipped with a 2.5 mm triple-channel solid-

Methods and techniques | Olga Nazarenko state probe head and an Avance III console. A 2.5 mm zirconia rotor was used. Chemical shifts were referenced to PbMe4. A Hahn echo pulse sequence was used. The 90-degree pulse had a length of 5 μs. The echo delay was set to 1 rotor cycle and the recycle delay to 0.5-2 seconds. For every spectrum, between 5’000 and 786’432 scans were recorded. The spinning frequency was set to 20 kHz except for α-FAPbI3, which was measured in static mode.

1H solid-state nuclear magnetic resonance (measured by L. Piveteau). All experiments were performed on a Bruker spectrometer (16.4 T) equipped with a 2.5 mm double-channel solid- state probe head and an Avance III console. A 2.5 mm zirconia rotor was used. Chemical shifts were referenced to tetramethylsilane (TMS). A one-pulse experiment was conducted. The 90-degree pulse had a length of 2.6 μs. The recycle delay was set to 4 seconds. For every spectrum, four scans were recorded. The spinning frequency was set to 11.8 kHz.

79Br solid-state nuclear magnetic resonance (measured by L. Piveteau). Experiments were performed on a 16.4 T Bruker spectrometer equipped with a 4 mm double-channel solid- state probe head and an Avance III console. A 4 mm zirconia rotor was used. Experiments were performed under static conditions without spinning the sample. Chemical shifts were referenced to 0.01 M NaBr in D2O. Wideband, uniform rate and smooth truncation (WURST) pulses were used in a Carr-Purcell-Meiboom-Gill (CPMG) echo-train sequence to excite wide frequency ranges at once while maximizing detected signal intensities. WURST-80 pulses of 50 µs, 1 MHz excitation width, and 63 W power were used. 64 echos were acquired with delays of 0.045 ms, resulting in spikelet separations of 5 kHz. 4096 transients were acquired for every sub- spectrum and recycle delays of 0.1 s were applied.

87Rb solid-state nuclear magnetic resonance (measured by L. Piveteau). Experiments were performed on a 16.4 T Bruker spectrometer equipped with a 4 mm double-channel solid- state probe head and an Avance III console. A 4 mm zirconia rotor was used. Spectra were measured under static conditions and spinning at 5 kHz at the magic angle. Chemical shifts were referenced to 0.01 M RbCl in D2O. One-pulse and Hahn echo pulse-sequences were used. 90° pulses lasted 8.25 µs and 180° pulses accordingly 16.5 µs. Echo delays varied from 0.087 – 0.788 ms. 256-5120 transients were acquired and recycle delays of 0.5 s were applied.

Electronic structures calculations (performed by Dr. E. Cuervo-Reyes, included in Chapter 4). Two implementations of density functional theory (DFT) were employed. For the calculations of the band structure and density of states, the Dmol3 package was used within the Materials Studio suite.157-158 For the real space representation of the bonding, the electron localization function (ELF)159-160 was computed using Savin’s implementation within the TB- LMTO-ASA code developed at MPI Stuttgart.161 The ELF plots contain information of both ELF and electron density over a given plane and were obtained using an in-house developed code. ELF values and electron density values are shown as the color of the pixels and the number of colored pixels, respectively, over a black background. Scalar-relativistic corrections were included in all calculations because they are expected to play an important role in the presence of heavy elements. The LMTO is an all-electron method and in Dmol3 including all

Methods and techniques | Olga Nazarenko electrons equally in the scalar-relativistic calculation, meaning that no pseudopotential was used. Within Dmol3, the tolerance for self-consistency was set to about 10-6 Ha for the total energy, using the number of k-points that results in a grid with a separation smaller than 0.03 A-1 in reciprocal space. The BOP (Becke-One Parameter) exchange-correlation functional was employed with which in general better outcomes were obtained for the bandgaps than with the commonly used PBE. For the LMTO calculations, the Langreth–Mehl–Hu exchange- correlation functional was employed. The self-consistency tolerance was set to 10-5 Ry for the total energy, and 10-5 e for the atomic charges. Taking advantage of the speed of the LMTO code, the k-space was sampled over 64-times denser grids than those used within Dmol3. The investigation of the effects on the bandgaps due to spin-orbit coupling and non-local exchange interactions, which require the extensive use of semi-empirical methods, should be the subject of future (and more computationally oriented) works.

First-principles calculations (performed by Dr. M. Kepenekian, Prof. Dr. J. Even, Dr. Sc. C. Katan, included in Chapter 6) were performed with experimental crystal structures using density functional theory (DFT) as implemented in the SIESTA package.162-163 Calculations have been carried out on experimental structures with the GGA functional in the PBE form.164 Core electrons were described with Troullier-Martins pseudopotentials,165 while valence wavefunctions were developed over double-ζ polarized basis set of finite-range numerical pseudoatomic orbitals .166 Spin-orbit coupling was taken into account through the on-site approximation as proposed by Fernández-Seivane, et al.167 In all cases, an energy cut-off of 150 Ry for real-space mesh size was used.

Elemental analysis. Elemental analysis (CHN) was performed by the standard combustion method, where carbon (as CO2) and hydrogen (as H2O) are analyzed quantitatively by infrared spectroscopy and nitrogen (N2) is determined by a thermal conductivity detector using LECO TRUSpec Micro instrument. The halogens (Br and I) were quantified by the Schöniger method, where they are collected in an absorbing liquid medium and then analyzed by titration. The work was done by the Laboratory for Microelemental Analysis, ETH Zürich.

Photoresponse measurement (performed by Dr. S. Yakunin). For the evaluation of the mobility-lifetime product for compounds presented in Chapter 3, μτ, a current-voltage measurement was obtained by a Keithley 236 SMU in the dark and then under infrared light at λ = 850 nm.

Gamma-ray absorption measurements (measurements performed by Dr. S. Yakunin). The absorption of gamma radiation was measured for a set of perovskite single crystals with a thickness of 0.2 to 15 mm. Gamma radiation from 241Am and 137Cs sources was collimated with a lead aperture having a diameter of 6 mm and a thickness of 5 cm. The intensity of gamma radiation transmitted through the perovskite single crystal was measured with a CZT detector (eV-Products, model B1758).

Methods and techniques | Olga Nazarenko

Gamma-detection measurements using perovskite single crystals (performed by Dr. S. Yakunin). A custom-made test fixture was used for energy-resolved measurements, connected to an A250CF CoolFET charge sensitive preamplifier (Ametek), coupled with an amplifier- shaper (Model 572, EG&G Ortec) and a digital multichannel analyzer MCA-8000D (Ametek). The high bias voltage was applied through a Keithley 236 SMU that was also used for monitoring the current through the single crystal detector. 241Am and 137Cs gamma sources with activities of 0.4 MBq and 2.2 MBq, respectively, were used for recording the energy spectra. For demonstration purposes, a prototype of a portable dosimeter based on a perovskite single crystal was connected to a preamplifier (eV Products) and a custom-made amplifier based on one stage of an LM358N chip. The gamma pulses were recorded with a digital counting board using an Arduino microprocessor set, based on an open-source project (https://sites.google.com/site/diygeigercounter/). Detector single crystals were biased at 9 V by a NiMH battery.

Scanning electron microscopy (SEM, performed by Dr. Frank Krumeich) was done on a Quanta 200F microscope (Thermo Fisher Scientific) operated at an acceleration voltage

Vacc = 20 kV. EDXS was performed with an Octane SDD detector (EDAX (Ametec)) attached to the microscope column. For spectra recording and quantification (ZAF correction), the software Gemini (EDAX) was used.

Wavelength dispersive X-ray Fluorescence. The pellets were examined with Rigaku Primus IV WD-XRF instrument with a standardless calculation of the element mass content. The measuring time was 10 min, with an energy of a 4 kW X-ray tube and an adjusted aperture of 1 mm diameter was used. A relative error of 7 - 10 % must be expected for standardless semiquantitative WD-XRF analyzes. The work was done by the Laboratory for Advanced Analytical Technologies, EMPA Dübendorf.

Continuous wave Electron Paramagnetic Resonance measurements (performed by Dr. Yevhen Polyhach) were performed at X-band frequency (~9.7 GHz) using a commercial X-band Elexsys E500 spectrometer (Bruker) equipped with a super high Q resonator (Bruker). Spectra were acquired at 10 K using 2 mW incident microwave power and 0.3 mT magnetic field modulation amplitude. Single crystal samples were placed into 3 mm quartz EPR tubes. Per sample, several random orientations with respect to external magnetic field were tested. Baseline measurements were performed using exactly the same settings on the same empty 3 mm tube.

Single crystals of caesium-formamidinium lead halide perovskites: solution growth and gamma dosimetry | Olga Nazarenko

Chapter 3. Single crystals of caesium- formamidinium lead halide perovskites: solution growth and gamma dosimetry

3.1. Introduction Lead halide semiconductors with the perovskite crystal structure,168 having the overall + + + - - - composition of APbX3 [A = Cs , CH3NH3 or CH(NH2)2 ; X = Cl , Br , I or mixture thereof], comprise an emerging class of optoelectronic materials for applications spanning inexpensive solar cells48, 51, 169-170 [with certified power conversion efficiencies exceeding 22% (http://www.nrel.gov/ncpv/images/efficiency_chart.jpg)], light-emitting diodes,171 lasers,172-173 broadband and narrowband photodetectors operating in the ultraviolet-visible-near infrared regions,174-175 soft X-ray detectors100, 176-177 or even gamma detectors.50, 94 Such a plethora of applications is enabled, on the one hand, by the material’s unique defect-tolerant photophysics,73, 178-180 i.e., the existence of a low density of carriers (109 - 1011 cm-3),50, 53, 130, 181 low densities of traps (109 - 1010 cm-3),53, 182 high carrier mobilities (25 - 100 cm2 V-1 s-1),50, 179 and high exciton diffusion lengths,50, 183 all despite a significant degree of structural imperfection on an atomistic scale (vacancies) and at the microscale (grain boundaries, polycrystallinity, etc.). On the other hand, LHPs can also be conveniently deposited from their solutions in the form of large single crystals,50, 53, 184-185 thin-films186-187 or various nanostructured forms.188-189 Of these forms, our focus herein is on single crystals for two distinctly different reasons. First, perovskite single crystals are naturally the most structurally perfect samples for fundamental studies of their intrinsic electronic and optical properties. Thus far, solution-grown single crystals of MAPbX3, FAPbX3, and CsPbX3 have been reported, and their photophysical properties were characterized.49, 53, 61, 96, 148, 181, 185, 190-192 The second, rather practical motivation lies in the unique applicability of perovskite single crystals as inexpensive, solid-state semiconductor detectors of high-energy photons (X-rays and gamma-rays).50, 94, 100, 176-177 Because most gamma radiation originates from the decay of radioactive nuclei, low-cost,

Single crystals of caesium-formamidinium lead halide perovskites: solution growth and gamma dosimetry | Olga Nazarenko sensitive and RT operational radiation detectors are clearly needed for a variety of applications related to the monitoring or neutralization of major threats from the operation of nuclear reactors, terrorism or limited nuclear war, or simply for occupational safety in working with radioactive isotopes. Perovskites hold great promise as hard radiation detectors due to the compelling combination of the aforementioned electronic characteristics, the foremost of which is a high mobility-lifetime product (µ), and the high absorptivity of hard radiation by the high Z lead and iodine atoms. Such high atomic number semiconductors are rarely available in the single-crystalline form, explaining why CdTe and CZT, grown using the complex and expensive Czochralski method, remain essentially the only materials used in commercial solid- state gamma-detectors.103, 193-194 In previous investigations of perovskite single crystals for gamma photon detection 94 focus lied on MAPbI3, MAPb(Br/I)3 and FAPbI3. The lattermost material, among all of those tested, was characterized by the highest µ product and lowest noise level and dark current, leading to the highest gamma-counting rate. This is consistent with other reports on FAPbI3 10 −3 single crystals, presenting a low trap density (nt = 1.13 × 10 cm ), high carrier mobility (µ = 35-40 cm2 V−1 s−1), 181, 191 carrier diffusion length of 6.6 µm and, importantly, low dark carrier density of 3.9 × 109 cm-3 (obtained by the space-charge-limited current calculation 181 technique). FAPbI3 is broadly regarded as a more chemically and thermally robust 61, 81, 145, 191 compound than MAPbI3, which decomposes to gaseous methylamine and hydrogen iodide, as described in Chapter 1. These beneficial attributes, along with narrower bandgap energy, have been recognized by the photovoltaic community, leading to a gradual shift of 143, 145, 195 research focus from MAPbX3 thin-film absorption layers to FA-based counterparts.

However, all devices employing pure -FAPbI3 (Figure 3.1a) eventually suffer from its thermodynamic instability towards conversion into a wide-bandgap hexagonal () phase (Figure 3.1b). In single crystals, this cubic-to-hexagonal transformation occurs very fast, typically within 24 hours after the growth of the single crystal.148, 181 In thin–films, this phase transformation requires from several hours to several weeks, being faster under humid atmosphere storage.196-197 This instability has recently been fully resolved for thin-films via the + formation of mixed-ionic compositions with Cs and MA, such as in Cs0.17FA0.83(PbI1-xBrx) (x 142, 144, 198-199 = 0-1) or (FAPbI3)1-x(MAPbBr3)x (x = 0-0.3). Such compositional tuning can stabilize the cubic structure via adjustment of the GTF,123 and determined by the radii of the constituting ions, and via the entropy of mixing.144 The aim of the study was to obtain and characterize FA-based perovskite single crystals and study their photophysical properties. Taking lessons from the photovoltaic research 142, 144, 198-199 discussed above, quaternary CsxFA1-xPbI3 and quinary CsxFA1-xPbI3-yBry compositions (x = 0-0.1, y = 0-0.6) were targeted. After growing cm-scale single crystals (Figure 3.1c), and examining their structure (Figure 3.1d) and thermal stability (Figure 3.1e), their gamma-counting capabilities were explored. All of the single crystals obtained were crystallized in the cubic perovskite lattice; they were found to be phase pure and possess composition-dependent bandgap energies and luminescence maxima in the range of 755-832 nm, blue-shifted compared to FAPbI3 (840 nm). These compositions were also found to exhibit

Single crystals of caesium-formamidinium lead halide perovskites: solution growth and gamma dosimetry | Olga Nazarenko

Figure 3.1. Growth and characterization of caesium-formamidinium lead halide single crystals. (a) Crystal structure of 3D cubic -FAPbI3 (showing orientational disorder of the FA 200 molecules). (b) Crystal structure of 1D hexagonal -FAPbI3 (showing disorder of the 81 nitrogen atoms down the 3-fold symmetry axis). (c) Photographs of typical 0.5-1 cm CsxFA1- xPbI3-yBry single crystals; the smallest cell size of the paper is 1mm1mm. (d) Powder XRD measurements of the (100) reflection as a function of composition for CsxFA1-xPbI3-yBry (x = 0-

0.1, y = 0-0.6). (e) Thermal analysis of Cs0.1FA0.9PbI2.7Br0.3 single crystals by TG (top) and DSC (bottom) in flowing Ar atmosphere. The derivative of the TG and DSC curves (solid lines) are shown as dashed lines. significantly improved phase-stability, with shelf lives (the time before hexagonal phase impurities could be detected) of up to 20 days for quaternary CsxFA1-xPbI3 single crystals and of >4 months for quinary CsxFA1-xPbI3-yBry single crystals. These single crystals possess outstanding electronic quality, manifested by an unusually high carrier mobilitylifetime product (µ) of up to 10-1 cm2 V-1, electronic stability under an applied bias of up to 30 V and sensitive gamma detection. Using single crystals of CsxFA1-xPbI3-yBry, an inexpensive prototype of a gamma-counting was built.

3.2. Experimental section

Chemicals and reagents. Lead (II) iodide (PbI2, 99%), formamidine acetate [FA(OAc), 99%] and -butyrolactone (≥99%) were purchased from Sigma-Aldrich. Lead (II) bromide

(PbBr2, 98+%) was purchased from Acros. Caesium iodide (CsI, 99.9%) and hydrogen iodide (HI, 57%, stabilized with 1.5% hypophosphorous acid) were purchased from ABCR. All chemicals were used as received without further purification. Synthesis of formamidinium iodide (FAI). FAI was synthesized via the reaction of FA(OAc) with an excess of HI (57%) upon stirring in an ice bath for two hours. Subsequently,

Single crystals of caesium-formamidinium lead halide perovskites: solution growth and gamma dosimetry | Olga Nazarenko the solution was dried in a rotary evaporator under vacuum and mild heating (50 C). The crude product was washed with diethyl ether and recrystallized from ethanol. After drying in a vacuum oven (overnight at 55 C) a white powder was obtained. Single crystals of FAI, grown via the diffusion of diethyl ether into an ethanol solution of FAI, were used for structure determination (Figure 3.2a, for crystallography tables, see Tables S3.1-3.4). The overall purity of the obtained FAI material was confirmed by X-ray powder diffraction (Figure 3.2b).

The growth of perovskite single crystals. CsxFA1-xPbI3-yBry (x = 0-0.1, y = 0-0.6) were grown using the inverse temperature crystallization technique, initially reported for FAPbI3 by 148 Bakr et al. Typically, for CsxFA1-xPbI3-yBry (x = 0-0.1, y = 0-0.4) a 0.9 M (referring to Pb) solution of CsI/FAI/PbX2 in γ-butyrolactone, prepared simply by dissolving the powders of three compounds, was used. All solutions were filtered through a 0.2 µm PTFE

(polytetrafluoroethylene) syringe filter. In order to grow Cs0.1FA0.9PbI2.4Br0.6 single crystals, the concentration was lowered to 0.8 M and the solution was not filtered. Typically, 4 ml of the precursor solution were filled into a 20 ml vial with a cap and the vial was placed into a preheated (80 C) glycerol bath. The temperature was elevated in small steps of 5 C/hour to 125-130 C, and maintained at this point for one hour. Single crystals nucleated and grew during this heating ramp. To avoid overheated areas on the bottom of the vial (to obtain fewer seeds), the vial was lifted off the bottom of the bath (fixed in a clamp), and the glycerol level was adjusted to be a few mm lower than the level of the crystallization solution. The temperature was monitored using a thermometer or thermocouple connected to the hotplate. Powder XRD were collected from ground crystals in transmission geometry (Debye- Scherrer Geometry) with a STADI P diffractometer (STOE & Cie GmbH), equipped with a silicon strip MYTHEN 1K Detector (Fa. DECTRIS) with a curved Ge (111)- Monochromator

(CuKα1, λ = 1.54056Å). Single crystal XRD measurements were conducted on Bruker Smart Platform diffractometer equipped with an Apex I CCD detector and molybdenum (MoKα, λ = 0.71073 Å) sealed tube as an X-Ray source. Crystals were a tip – mounted on a micromount with paraffin oil. Data was processed and refined with APEX2 (Bruker software) and Olex2 software (Durham University). Thermal analysis (TG and DSC) were performed using a Netzsch Simultaneous Thermal Analyzer (STA 449 F5 Jupiter). A powdered sample (10 mg) was placed in an alumina crucible under Ar/N2 gas flow (50 ml/min) and heated to 850 C (at 10 C min-1). UV-Vis absorption spectra of powdered samples held between two glass slides were collected using a Jasco V670 spectrometer equipped with an integrating sphere. PL measurements were performed with a Fluorolog iHR 320 Horiba Jobin Yvon spectrofluorometer equipped with a Xe lamp and a photomultiplier tube. Elemental analysis

(CHN) was performed by the standard combustion method, where carbon (as CO2) and hydrogen (as H2O) are analyzed quantitatively by infrared spectroscopy and nitrogen (N2) is determined by a thermal conductivity detector. The halogens (Br and I) were quantified by the Schöniger method, where they are collected in an absorbing liquid medium and then analyzed by titration. Photoresponse measurement: for the evaluation of the mobility-lifetime product, μτ, a current-voltage measurement was obtained by a Keithley 236 SMU in the dark and then under infrared light at λ = 850 nm. Being at the low-energy tail of the absorption spectrum, this

Single crystals of caesium-formamidinium lead halide perovskites: solution growth and gamma dosimetry | Olga Nazarenko

Figure 3.2. (a) The crystal structure of formamidinium iodide. (b) A powder XRD pattern of formamidinium iodide. wavelength corresponds to a large penetration depth of ~1 cm, similar to that of gamma photons with an energy of 300 keV. The absorption of gamma radiation was measured for a set of perovskite single crystals with a thickness of 0.2 to 15 mm. Gamma radiation from 241Am and 137Cs sources was collimated with a lead aperture having a diameter of 6 mm and a thickness of 5 cm. The intensity of gamma radiation transmitted through the perovskite single crystal was measured with a CZT detector (eV-Products, model B1758). Gamma-detection measurements using perovskite single crystals. A custom-made test fixture was used for energy-resolved measurements, connected to an A250CF CoolFET charge sensitive preamplifier (Ametek), coupled with an amplifier-shaper (Model 572, EG&G Ortec) and a digital multichannel analyzer MCA-8000D (Ametek). The high bias voltage was applied through a Keithley 236 SMU that was also used for monitoring the current through the single crystal detector. 241Am and 137Cs gamma sources with activities of 0.4 MBq and 2.2 MBq, respectively, were used for recording the energy spectra. For demonstration purposes, a prototype of a portable dosimeter based on a perovskite single crystal was connected to a preamplifier (eV Products) and a custom-made amplifier based on one stage of an LM358N chip. The gamma pulses were recorded with a digital counting board using an Arduino microprocessor set, based on an open-source project (https://sites.google.com/site/diygeigercounter/). Detector single crystals were biased at 9 V by a NiMH battery.

3.3. Results and discussion

Solution-growth, structure and thermal analysis of CsxFA1-xPbI3-yBry (x = 0-0.1, y = 0-0.6, Figure 3.1). The single crystals were grown at 95-130 ⁰C in γ-butyrolactone, using a modified inverse temperature crystallization (ITC) method (see details in the Experimental

Section), similarly to the original ITC method reported for FAPbI3 and FAPbBr3 by Bakr et al.148 In this method, the nucleation and fast growth of crystals occur upon heating, as opposed to commonly used cooling-induced supersaturation techniques. In ITC, the rare occurrence of

Single crystals of caesium-formamidinium lead halide perovskites: solution growth and gamma dosimetry | Olga Nazarenko a decrease in solubility upon heating (retrograde solubility) is exploited to induce 201 crystallization. For FAPbI3 in γ-butyrolactone, such unusual solubility behavior was suggested to be caused by the dissociation of solvent-precursor complexes at elevated temperatures, causing oversaturation.49, 148 Recent success in thinness- and shape-controlled growth of MAPbX3 single crystals for inch-sized photoelectronic devices on the inch-scale speaks volumes for the versatility of an ITC method.202-203 The modification of an ITC method, used for growth of FAPbX3, implemented in this work, lies in the use of caesium iodide as the Cs-source and lead bromide as the Br-source. The overall concentration of A-type cations (FA+Cs), Pb cations and halide anions (Br+I) was maintained at 0.9 M, 0.9 M and 2.7 M, respectively. Experimentally, a notable challenge was to prevent the formation of a yellow, non-perovskite CsPbI3 phase, often occurring for x  0.1. Neat γ-butyrolactone remained the best solvent, following trials using dimethyl sulfoxide or N,N-dimethylformamide as additives or concentrated HI as a solvent. Chemical compositions were estimated by standard CHN combustion analysis for the organic part and by the Schöniger method for Br and I (by titrimetric analysis, Table S3.5). The resulting compositions of single crystals are proportional to the mixing ratios in the mother solution. Thermal analysis of Cs0.1FA0.9PbI2.7Br0.3 indicates satisfactory stability upon heating in an inert atmosphere (Figure 3.1e). For instance, 61, 81 decomposition does not occur until above 300 C, a similar case to pure FAPbI3. Two steps of mass loss were observed. A total gravimetric loss of 20 % occurs between 316-370 C, corresponding to the removal of FA-halide (e.g., conversion into Cs0.1PbX2+x). Lead halides then evaporate above 500 C, leaving CsX residue. In comparison, the removal of the organic 204-206 part from the MA counterpart, CH3NH3PbI3, already occurs at 150-200 C. The powder

XRD pattern of the parent FAPbI3 single crystal (Figure 3.3) indicates phase-pure α-FAPbI3 with cubic symmetry, consistent with previously reported crystallographic data.181, 200

All compositions of CsxFA1-xPbI3-yBry (x = 0-0.1, y = 0-0.6) exhibited the same general XRD pattern, without any detectable crystalline impurities (Figure 3.4). The formation of solid solutions is clearly visible by the shifts of the XRD reflections, as exemplified in Figure 3.1d for the (100) reflection. With the simultaneous replacement of FA and I- ions with smaller Cs+ and Br- ions, an expected systematic shift to higher angles without any apparent broadening was observed, confirming the decrease in the lattice constant.

Single crystals of caesium-formamidinium lead halide perovskites: solution growth and gamma dosimetry | Olga Nazarenko

Figure 3.3. Powder XRD pattern of as-prepared α-FAPbI3 indexed according to the cubic FAPbI3 phase. The simulated pattern is based on the experimental single-crystal data reported by M. T. Weller et al. [J. Phys. Chem. Lett. 6, 3209-3212 (2015)].

Figure 3.4. (a) Powder XRD patterns of CsxFA1-xPbI3-yBry (x = 0-0.1, y = 0-0.6). (b) Selected 2θ-range from the powder XRD patterns of a chosen synthesized compound

(Cs0.1FA0.9PbI2.6Br0.4) and Si powder. (c) A comparison of FWHM values for perovskite single crystals and reference, highly-crystalline Si. Down to the instrument resolution, the crystallinity of perovskite single crystals is at least as good as of reference Si.

Single crystals of caesium-formamidinium lead halide perovskites: solution growth and gamma dosimetry | Olga Nazarenko

To within the instrumental resolution of our powder X-ray diffractometer, grounded perovskite single crystals, obtained in this work, exhibit the same degree of crystallinity as highly-crystalline Si (Figure 3.4b, c). Such a shift is also detectable, although to a lesser extent, with the exchange of FA by Cs+ only (i.e., without alteration of the halide composition, Figure S3.1). Powder XRD was then also used to monitor the phase stability of the single crystals grown in this work. In these experiments, all samples were handled under ambient conditions but stored in a desiccator over calcium chloride in order to ensure standardized humidity conditions during storage. Single crystals of α-FAPbI3 transformed into a yellow hexagonal phase within 24 hours, in agreement with earlier reports.81, 181 The replacement of 10 % of the FA ions by Cs+ led to much higher stability, with the first signs of degradation (minor XRD - peaks corresponding to hexagonal FAPbI3) seen only after 20 days of storage. When also Br ions are added, i.e. to form quinary compositions, the stability could be further improved to at least 2 months of storage (e.g., at y = 0.2-0.4, Figure 3.5).

Optical properties. The facile synthesis of solid solutions of CsxFA1-xPbI3-yBry (x = 0- 0.1, y = 0-0.6) lends itself to measurements of their optical properties, wherein the continuous compositional variation manifests in the continuous evolution of optical absorption and PL spectra (varying x in Figure 3.6, and varying y shown in Figure 3.7a-c). A bandgap of 1.43 eV (the narrowest measured herein) and a PL maximum at 840 nm is found for the reference sample 61, 143, 145 of -FaPbI3 (Figure 3.7a, b), in agreement with the literature. The incorporation of either Cs, Br or both increases the bandgap energy (Figure 3.7b). In general, the bandgap of the

3D APbX3-perovskites scales either with the nature of the halide ion (the electronegativity effect) or with geometric factors (e.g., the tilting angle between PbX6 octahedra, compressive/tensile strain, etc.).207 The effect of substitutional Cs-doping (Figure 3.6, Figure

3.7c) in FAPbI3 can be ascribed to fine geometric effects. The replacement of FA with Cs may also influence the spin-orbit effects in this lattice.208 The dependence of bandgap on composition reported in this work is also consistent with numerous recent thin-film studies in photovoltaic research.123, 144, 198-199 It should be emphasized that the convenient tuning of the absorption onset to shorter wavelengths of 700-750 nm is potentially highly advantageous for the construction of tandem solar cells.209-212

Single crystals of caesium-formamidinium lead halide perovskites: solution growth and gamma dosimetry | Olga Nazarenko

Figure 3.5. (a, b) Phase stability/purity of formamidinium lead iodide single crystals substitutionally doped with Cs or Cs/Br as shown by powder XRD. Asterisks indicate the formation of the yellow, hexagonal FAPbI3 phase.

Figure 3.6. Optical properties of CsxFA1-xPbI3 (x=0-0.1): (a) PL spectra and (b) absorption spectra.

Single crystals of caesium-formamidinium lead halide perovskites: solution growth and gamma dosimetry | Olga Nazarenko

Figure 3.7. Optical properties of CsxFA1- xPbI3-yBry (x = 0-0.1, y = 0-0.6) single crystals. (a) RT PL spectra of grounded single crystals. (b) Absorption spectra of single crystals held between two glass plates. (c) Bandgap energy as a function of Br content for compositions of fixed Cs content (x = 0.1).

Absorption and detection of gamma photons in hybrid perovskite single crystals. Upon the substitutional

doping of FAPbI3 single crystals by Cs and Br, not only does the crystalline phase stability increase and the optical absorption shift to lower wavelength, but also the electronic parameters such as dark current and shot noise are affected. In particular, the higher the bandgap is, the greater are the dark currents (see specific resistivity of single crystals of various compositions in Figure S3.2) and noise levels in the studied systems (not generally expected). Taking into account this trade-off between the doping level and the corresponding electronic characteristics, a focus lied on single crystals with a composition of x = 0.1 and y = 0.2 (i.e., 10% of Cs and 10% of

PbBr2) in subsequent gamma-detection studies. The absorption of X-rays and gamma radiation in high-Z semiconductor materials can be described by several processes: photoionization, Compton scattering and electron-positron pair creation.103 A calculation of the absorption (α) of high energy photons by single crystals with the composition

Cs0.1FA0.9PbI2.8Br0.2 in the photon energy range of 20-1000 keV (Figure 3.8a) was performed. The resulting absorptivity spectrum contains two resonant absorption peaks due to transitions involving the core electron levels of constituting atoms, so-called Kα lines, a long, monotonically decreasing photoelectric tail and a component responsible for electron-positron pair creation at higher energies (>1 MeV). The accuracy of these calculations was verified by measurements of the absorption of gamma radiation from 2 sources (241Am and 137Cs) by several 0.2-15 mm thick single crystals, and the results are presented as two points in Figure

Single crystals of caesium-formamidinium lead halide perovskites: solution growth and gamma dosimetry | Olga Nazarenko

Figure 3.8. Response of Cs0.1FA0.9PbI2.8Br0.2 single crystal to gamma rays. (a) Absorption and penetration depth as a function of photon energy, showing the calculated spectrum (green line) and measured values (red and blue symbols). The shaded green area indicates the effective range of charge collection, within which the penetration depth corresponds to the charge collection distance at the maximum applied bias of 30 V. (b) Evaluation of μτ from the dependence of photocurrent on the bias, and the corresponding effective distance of charge collection. (c, d) Energy resolved spectra of gamma detection from an 241Am and an 137Cs source, respectively.

3.8a. From the penetration depth (calculated as the inverse of absorption), it is apparent that 241 emission from the Am source (Eγ = 59.6 keV) at rather a low gamma photon energies can be absorbed by a 1 mm sized single crystal. Comparatively, a 10 times larger (1 cm) single crystal 137 absorbed only ~30 % of the intensity of gamma radiation from a Cs source (Eγ = 662 keV). With such large single crystals, the efficiency of carrier separation and collection becomes a major concern. It was recently shown that a drift length in the mm range can be accomplished 100 in MAPbBr3 single crystals. Charge collection efficiency can be evaluated from the dependence of the photocurrent on the bias (Figure 3.8b), with the aid of the Hecht model.213- 214 This directly yields a figure-of-merit known as the carrier mobility-lifetime product (μτ).103 This parameter increases drastically for single crystals15 as compared to polycrystalline perovskites.176 In this study, a median value of μτ 0.04 cm2 V-1 was found for single crystals of 2 - Cs0.1FA0.9PbI2.8Br0.2, whereas the champion single crystal showed a value of μτ = 0.12 cm V 1. These values correspond well to the best reported μτ values for CdTe-based single crystals. This and a previous study94 show that with LHPs the charge collection efficiency is primarily limited by the bias-stability of the single crystals. Due to ionic migration in perovskites, also

Single crystals of caesium-formamidinium lead halide perovskites: solution growth and gamma dosimetry | Olga Nazarenko known to cause hysteresis and other instabilities in photovoltaic solar cells,215 the highest bias at which the dark current and noise are not significantly affected is in the range of 5-30 V, being highest for FA-based single crystals and lowest for MAPbI3 single crystals. Based on these values, a drift length can be estimated (Figure 3.8b), also known as a schubweg,103 calculated as μτE, where E is the applied electric field. For optimal charge collection, the penetration depth of a gamma photon of a certain energy must not exceed the drift length; that is, αμτE ≥ 1. This assumes that the electric field (i.e., the trajectories of the collected carriers) is parallel to the gamma flux. At the higher allowed biases in FA perovskites, the calculated drift length reaches 1.7 cm, meaning that at photon energies lower than 400 keV the penetration depth will not exceed the drift length. The condition of αμτE ≥ 1 for drift lengths of 1.7 cm is graphically represented as a shaded green area in Figure 3.8a. Only within this area should a given single crystal be able to efficiently collect the charge carrier. Efficient charge carrier collection is also a precondition for obtaining energy-resolved spectra, based on the pulse-height analysis. The latter aspect can be illustrated by performing a pulse-height analysis of the signals from two gamma sources: 241Am (Figure 3.8c) and 137Cs (Figure 3.8d). Clearly, the better energy- resolving capability of perovskite single crystals is obtained at lower energies, i.e. within the range of efficient charge collection. Discussion ABX3 perovskites are composed of BX6 octahedra, connected by corner-sharing in three dimensions. One such compound, -FAPbI3 exhibits a highly symmetric cubic Pm-3m lattice with a unit cell parameter of a = 6.3620(8) Å (Figure 3.1a).181, 200 However, at RT this lattice undergoes a phase transformation into a hexagonal P63mc crystal structure, consisting of single

1D chains of face-sharing PbI6 octahedra and FA cations situated in between the chains (Figure 3.1b).81 The 3D electronic delocalization is therefore broken, deteriorating its semiconductive properties and drastically increasing the bandgap. Sometimes this phase is simply referred to as a “yellow phase” due to its color. Only at elevated temperatures of 185 C does the hexagonal 61 FAPbI3 transform into the black -FAPbI3. This black phase is obtained as a single crystal upon synthesis from a hot solution at 130 C, but then quickly undergoes a phase transition to the yellow phase at RT.61, 148, 181 The primary goal of this work was to prevent this phase transition through compositional variations within the FAPbI3 lattice: by replacing FA cations with up to 10% of Cs, and I anions with up to 30% of Br. The compositionally-dependent stability of perovskites can be semi-quantitatively rationalized using the GTF concept, initially proposed for metal-oxide perovskites37 and recently extended to metal-halides.38-39, 123, 216 The

GTF of a 3D perovskite with ABX3 composition and idealized cubic lattice can be calculated as: (푟 + 푟 ) 퐺푇퐹 = 퐴 푥 √2(푟퐵 + 푟푥) + where rA, rB and rX represent the ionic radii of each lattice site constituent (in this case, rFA = 2+ - 253 pm, rPb = 119 pm and rI = 220 pm). Stable cubic perovskites of formula ABX3 are 38, 79, 123 expected to exhibit a GTF between 0.8 and 1, which explains the instability of -FAPbI3

Single crystals of caesium-formamidinium lead halide perovskites: solution growth and gamma dosimetry | Olga Nazarenko at RT (GTF  1). It should be noted that the non-spherical geometry of the FA cation complicates the analysis.217 Clearly, the GTF can be decreased by replacing FA cations by smaller Cs+ ions, and/or by replacing I- anions with smaller Br- ions, likely leading to higher stabilities. These possibilities have been successfully tested in this work, leading to greatly increased phase stabilities of the perovskite single crystals, as evidenced by an extended shelf- life under ambient conditions of up to at least several months without any detectable traces of other crystalline phases (Figure 3.5). The most preferred scenario in the context of gamma- sensing is FA-to-Cs exchange, as it has minimal effect on the bandgap energy and incorporates a heavy atom, which is highly desirable for the absorption of gamma photons. Additionally, up to 30% of the I sites could be simultaneously substituted with Br- ions. Both cationic and anionic exchanges manifested themselves as changes in the bandgap energy. Cs0.1FA0.9PbI3 exhibits a

0.05 eV higher bandgap than -FAPbI3, whereas a further increase by 0.2 eV can be seen upon the incorporation of 30% Br, i.e. for Cs0.1FA0.9PbI2.4Br0.6. Besides Cs- and concomitant Cs/Br substitutional doping, doping of the halide site only, e.g. in the formation of FAPbI3-xBrx, has been successful as well and will be reported elsewhere. There is a good proportionality between

Figure 3.9. A prototype of a gamma dosimeter based on Cs0.1FA0.9PbI2.8Br0.2 single crystal. (a) Operational stability (crosses) and storage stability (circles) of Cs0.1FA0.9PbI2.8Br0.2 single crystals compared to MAPbI3 and FAPbI3 single crystals. (b) Detail photograph of the detector fixture with an installed FAPbI3 or Cs0.1FA0.9PbI2.8Br0.2 single crystal. Photographs of the complete device during measurement (c) in the absence of a gamma radiation source and (d) in the presence of a 137Cs source (activity = 2.2 MBq). The labels correspond to 1 – detector crystal fixture, 2 – preamplifier (from eV Products), 3 – amplifier, 4 – digital counting board using an Arduino microprocessor set, based on an open-source project (https://sites.google.com/site/diygeigercounter/) and 5 – 137Cs source.

Single crystals of caesium-formamidinium lead halide perovskites: solution growth and gamma dosimetry | Olga Nazarenko the I:Br and Cs:FA ratios in the mother solution and in the resulting single crystals (Table S3.5), with the general trend being that bromide contents are higher in the single crystals.

Next gamma-photon-sensing properties of CsxFA1-xPbI3-yBry single crystals were tested 94 and the results were compared to an earlier work on MAPbI3. Besides the thermodynamic phase stability and chemical stability that both determine the shelf-life of the single crystals, the operational stability due to the effects of polarization and ionic motion218-219 also becomes an important issue for practical gamma detection. CsxFA1-xPbI3-yBry single crystals exhibit a much higher operational stability than MAPbI3, as can be seen from the continuous gamma counting measurement under applied voltage bias (Figure 3.9a). The counting rates of all three compositions compared in Figure 3.9a fall within the same order of magnitude, with CsxFA1- xPbI3-yBry single crystals exhibiting systematically 2-3 times higher values. The subtle decrease in sensitivity can be quickly recovered by applying a bias of opposite sign, suggesting that ionic migration is the major cause for the sensitivity decrease. Such improvements inspired us to design a simple and low-cost prototype of a gamma dosimeter, comprising the advantages of solution-grown single crystals and the swiftly-developing field of open-source microprocessor controllers. An open-source “do-it-yourself” (DIY) project (https://sites.google.com/site/diygeigercounter/) was applied that uses standard Geiger gas tube as a detector and Arduino microprocessor platform for digital counting board. The Geiger counted with perovskite single crystal solid-state gamma detector (Figure 3.9b) and commercial preamplifier (Figure 3.9c). The resulting assembly exhibits a strong response to a 137Cs gamma radiation source of rather low activity (2.2 MBq), as illustrated by Figures 3.9c,d. These devices were assembled and handled under ambient conditions. The long-term operation of such dosimeters is directly determined by the chemical integrity of the perovskite single crystals, ranging from, at best, several days for FAPbI3 to several months for CsxFA1-xPbI3-yBry single crystals.

3.4. Conclusions In this chapter a simple solution-growth method was presented to synthesize quaternary

CsxFA1-xPbI3 and quinary CsxFA1-xPbI3-yBry (x  0.1, y  0.6) single crystals with excellent long-term thermodynamic stability. Such materials present an attractive alternative to MAPbX3 59 compounds, which are prone to decomposition into PbX2 and MAX. With very high mobility- lifetime products of up to 1.210-1 cm2 V-1, such high-Z semiconductor single crystals exhibit a high sensitivity to gamma photons, allowing single photon counting and inexpensive dosimetry. Besides hard radiation detection as presented herein, owing to their compositionally tunable bandgap energies, these single crystals might attract further research activities toward investigating their applications in photodetectors operating in the visible and infrared regions, as well as in fundamental studies of the charge transport and photophysics of these novel semiconductors.

Single crystals of caesium-formamidinium lead halide perovskites: solution growth and gamma dosimetry | Olga Nazarenko

Reproduced with modifications from: O. Nazarenko, S. Yakunin, V. Morad, I. Cherniukh, M. V. Kovalenko. Single crystals of caesium formamidinium lead halide perovskites: solution growth and gamma dosimetry. NPG Asia Mater., 2017, 9 (4), e373. Copyright 2017, Springer Nature. https://www.nature.com/articles/am201745

Luminescent and photoconductive layered lead halide perovskite compounds comprising mixtures of caesium and guanidinium cations | Olga Nazarenko

Chapter 4. Luminescent and photoconductive layered lead halide perovskite compounds comprising mixtures of caesium and guanidinium cations

4.1. Introduction Organic-inorganic layered lead halide compounds with perovskite-like structures are widely studied due to their fascinating electronic and optical properties and their rich structural chemistry.2-3, 9, 34, 79, 91, 207, 220-233 Fundamental to the construction of perovskites is the formation - - - of PbX6 octahedra (X = Cl , Br , I and mixtures thereof) and their condensation into crystals that possess 2D and 3D corner sharing. Such extended 2D and 3D delocalization of electronic bands impart semiconductive properties. Compounds with 3D arrangements and the generic + + + formula of APbX3 are formed when very small A-site cations are utilized [A = Cs , CH3NH3 , + 49, 98, 148, 181, 234-237 CH(NH2)2 ]. By contrast, larger A-site cations do not fit within the highly confined space between the octahedra. When a larger A-site organic cation is employed, the most likely outcome is the formation of 2D compounds, such as the compounds with the generic formula A2PbX4 consisting of single layers of PbX6 octahedra sharing four of their corners, with the A cations in the interlayer space. Most of the known layered (2D) perovskite derivatives have mono- or di-ammonium cations, i.e., (RNH3)2PbX4 and (NH3RNH3)PbX4, + respectively, where “R” is commonly an alkyl chain. The non-spherical geometry of RNH3 cations plays a significant role in the crystal symmetry of (RNH3)2PbX4 compounds, in addition 1, to distortions and tilting of the PbX6 octahedra, which are often caused by lone-pair effects. 238 The orientation of the alkylammonium cations between the inorganic layers allows the ammonium groups (-NH3) to form hydrogen bonds with neighboring halogen ions.

Luminescent and photoconductive layered lead halide perovskite compounds comprising mixtures of caesium and guanidinium cations | Olga Nazarenko

These A2PbX4 compounds can, therefore, be considered strongly quantum-confined systems, i.e., quantum wells, wherein electron motion is limited to two dimensions. An interlayer organic entity acts as the potential wall. The variation in the carbon chain length in

(RNH3)2PbI4 (from 4 to 14 carbon atoms) has a rather modest effect on the observed bandgap energies. For example, the energies of the PL maxima are 2.35 eV for C4 and 2.4 eV for C10, confirming a highly localized quantum-well-like behavior.1-2, 222 Engineering the chemistry of these compounds at the interlayer space is of increasing interest,230 with examples including intercalations of molecules such as hexane, dichlorobenzene, iodine, etc..239-240 However, engineering the inorganic slab remains the principal approach to adjusting the absorption edges and PL maxima. For example, bandgaps of 3.75, 3.17 and 2.55 eV are found for X- = Cl-, Br-, - 241 and I , respectively, in (C10H21NH3)2PbX4. Another strategy to tune the electronic structure is to increase the thickness of the lead halide slabs from a single layer to multilayers, thereby tuning the bandgap energies between the two limiting cases of 2D A2PbX4 and 3D APbX3. Creating thicker slabs must involve at least two A-site cations, wherein the larger cation remains as an interlayer cation, and the second cation fills in the smaller inter-slab cavities (which are of similar size as in the limiting case of APbX3 compounds). Structurally, these compounds are analogues to the Ruddlesden–Popper phase, as in the case of (RNH3)2(CH3NH3)n-1PbnX3n+1 (R - - - 80 + = C9H19-, Ph–CH2CH2–; X = Br , I ). Herein, a small CH3NH3 fills voids within the lead halide slabs, and the long alkyl-chain ammonium ions are situated between the slabs, thus determining the inter-slab distance and defining the crystallographic orientation of the perovskite sheets. These phases can be seen as 2D slices of a 3D APbX3 with various crystallographic orientations. For further crystallographic details of such compounds, recent review and research articles can be consulted.79, 242-243

In (C4H9NH3)2(CH3NH3)n-1PbnI3n+1, the bandgap energies change from 2.43 to 1.91 eV as the thickness is adjusted from n = 1 to n = 4.2 Due to such tunability, these 2D systems are presently being intensively investigated for their optical properties and applications in light emission and in solar cells.1, 9, 82, 127, 222, 244 In particular, the archetypical 3D compound,

CH3NH3PbI3 (n = ∞), had been reported recently as an outstanding photovoltaic material with 58, 195, 236, 245 power conversion efficiencies repeatedly exceeding 20 %. However, CH3NH3PbI3 is known to undergo irreversible degradation due to the volatile nature of the decomposition products (methylamine, HI, etc.), which is accelerated upon exposure to moisture, as in ambient 246 air. Higher chemical stability can be found within the family of (PEA)2(CH3NH3)n-1PbnI3n+1 + (PEA = C6H5(CH2)2NH3 ), wherein compounds with n = 1 - 3 have been synthesized and 87, 247 characterized in detail. (PEA)2(CH3NH3)2Pb3I10 (n = 3, bandgap of 2.1 eV) exhibits higher 247 chemical stability than CH3NH3PbI3. The relevance of such structural motifs extends beyond lead-based halides; for example, in the tin iodide compounds (C4H9NH3)2(CH3NH3)n-1SnnI3n+1, for n ˃ 3 metallic behavior is observed, while compounds with n < 3 are semiconductors.127 However, tin-based compounds are highly prone to oxidation to Sn4+, limiting their practical utility. In this work, the full elimination of protonated ammonium salts was made, such as those in primary, secondary or ternary alkyl amines, from the construction of layered LHPs. To create + + monolayer A2PbX4-like compounds, the unstable CH3NH3 was replaced with a mixture of Cs

Luminescent and photoconductive layered lead halide perovskite compounds comprising mixtures of caesium and guanidinium cations | Olga Nazarenko

+ and guanidinium [C(NH2)3] ions. Guanidinium is an ideal alternative to the primary ammonium ions due to its thermodynamic stability, high basicity (pKa = 13.6), and strong hydrogen-bonding capabilities.248-249 The high basicity of the guanidine molecule can be understood considering a resonance stabilization of the guanidinium ion by 6-8 kcal/mol.250 The guanidinium cation has been previously incorporated into a variety of perovskite-like salts - II with the formate (HCOO ) anion: [C(NH2)3][M (HCOO)3], where M is Mn, Co, Fe, Ni, Cu, Zn, or Cd.251-252 This suggests that guanidinium facilitates stable, highly symmetric structures.

Although stable [C(NH2)3]2Pb(Sn)I4 compounds have been already reported (with corrugated structure),4, 253 our interest was to test three further possibilities: (i) whether it is possible to mix such a stable organic cation with an inorganic cation at the interlayer and what structural effects accompany this combination, (ii) whether this mixture can also enable the formation of slabs thicker than one monolayer and (iii) whether non-corrugated structure can form. The small Cs ion is universally suited for both interlayer locations and for filling the voids within the slab. By contrast, the guanidinium ion can be accommodated only in the interlayer space, despite + + 39 being only slightly larger than CH3NH3 and CH(NH2)2 . Three-layered perovskite compounds with n = 1 - 2: Cs[C(NH2)3]PbI4 (I),

Cs[C(NH2)3]PbBr4 (II), and Cs2[C(NH2)3]Pb2Br7 (III) were obtained, all of which contain (2D) slabs of corner-sharing octahedra and possess satisfactory thermal stability (up to 300 C). Compounds I-II are luminescent upon moderate cooling, and compound III is emissive at RT. Photoresponsivity in the range of 1-10 mA·W-1 and with narrow-band characteristics were measured for compounds I and III, confirming their extended semiconductor nature.

4.2. Experimental section

Figure 4.1. The photographs of Cs[C(NH2)3]PbI4 and Cs2[C(NH2)3]Pb2Br7 compounds under the daylight at RT and under UV illumination in liquid nitrogen.

Chemicals, reagents and synthesis procedures. Lead (II) iodide (PbI2, 99%) and cesium bromide (CsBr, 99.9%) were purchased from Sigma-Aldrich. Lead (II) bromide (PbBr2, 98+%), hydrobromic acid (HBr, 48% water solution) and guanidinium carbonate (99+%) were

Luminescent and photoconductive layered lead halide perovskite compounds comprising mixtures of caesium and guanidinium cations | Olga Nazarenko purchased from Acros. Cesium iodide (CsI, 99.9%) and hydriodic acid (HI, 57%, stabilized with 1.5% hypophosphorous acid, H3PO2) were received from ABCR. Diethyl ether (˃99%) was dried over molecular sieves. All chemicals were used as received without further purification.

Cs[C(NH2)3]PbI4 (I), Cs[C(NH2)3]PbBr4 (II), and Cs2[C(NH2)3]Pb2Br7 (III) were synthesized from concentrated hydriodic and hydrobromic acids. The compounds were yellow (II, III) and red (I) microcrystalline powders (Figure 4.1). A large excess (≥6-fold) of guanidinium ions was found to be necessary to obtain the desired compounds.

Cs[C(NH2)3]PbI4 (I). CsI (1.55 mmol), [C(NH2)3]2CO3 (6 mmol) and PbI2 (1.55 mmol) were loaded into a 25-ml round- bottom flask, and 10 ml of HI acid (57% in water) was added carefully. Strong gas evolution was observed, and a red precipitate immediately formed. The precipitate was re-dissolved upon heating in a glycerol bath, yielding a clear yellow solution. Next, the flask was cooled down under a cold water stream, causing the formation of a red crystalline precipitate of I. The solution remained undisturbed for another few hours to allow further precipitation. The precipitate was separated by vacuum filtration, washed with diethyl ether, and dried in a vacuum oven at 60 C. A yield of 72% was estimated relative to Pb. To obtain larger crystals of I for photoconductivity measurements (0.5-2 mm3), the precursor solution was slowly evaporated at 60 C instead of cooling.

Cs[C(NH2)3]PbBr4 (II). CsBr (2 mmol), [C(NH2)3]2CO3 (6 mmol), PbBr2 (2 mmol) were loaded into a 25-ml round- bottom flask, and 10 ml of HBr (48%) acid was carefully added. Strong gas evolution was observed, and a yellow precipitate formed. Subsequent manipulations of the solution were identical to the procedure for compound I. A yield of 63% was estimated relative to Pb.

Cs2[C(NH2)3]Pb2Br7 (III). CsBr (1 mmol), [C(NH2)3]2CO3 (3 mmol), PbBr2 (2 mmol) were used for this synthesis. Upon addition of 12 ml of HBr (48%) acid, strong gas evolution was observed, and a yellow precipitate formed. Subsequent manipulations for the solution were identical to the procedure for compound I. A yield of 27% was estimated relative to Pb. To obtain larger crystals of III for photoconductivity measurements (0.5-2 mm3), the precursor solution was slowly evaporated at 60 C. 2 (n-C4H9NH3)2PbI4 (IV) was prepared according to Stoumpous et al.

[C(NH2)3]2PbI4 (V) was obtained as yellow platelets from a hot acidic solution using [C(NH2)3]2CO3 (3 mmol), PbI2 (2 mmol) and 7.5 ml of HI (57%). The crystals were separated by vacuum filtration, washed with diethyl ether and dried in a vacuum oven at 60 C.

CsPbBr3 (VI) was crystallized from aqueous HBr using corresponding bromides in stoichiometric quantities (0.15 M PbBr2, 0.15 M CsBr). Upon cooling to RT, orange crystals precipitated. Subsequent manipulations of the solution were identical to the procedure for compound I.

The synthesis of (n-C4H9NH3)2PbBr4 (VII) from hydrobromic acid was analogous to that of IV, using n-butylamine and PbO (in stoichiometric quantities, 0.5 M PbO, 1 M

C4H9NH2). Upon cooling to RT, colorless transparent plates crystallized. Subsequent manipulations of the solution were identical to the procedure for compound I.

Luminescent and photoconductive layered lead halide perovskite compounds comprising mixtures of caesium and guanidinium cations | Olga Nazarenko

CsPb2Br5 (VIII) was crystallized from stoichiometric quantities of corresponding bromides in hydrobromic acid. Typically, CsBr (0.7 mmol) and PbBr2 (1.4 mmol) were dissolved in 6 ml of hot HBr acid. Upon cooling to RT, colorless crystalline plates precipitated, followed by vacuum filtration, washing with diethyl ether and drying in the vacuum oven at 60 C.

[C(NH2)3]2PbBr4 (IX) was grown from a hot acid solution. [C(NH2)3]2CO3 (4 mmol), PbBr2 (4 mmol) were loaded into a 25-ml round -bottom flask with 10 ml of HBr (48%) acid. The precipitate was re-dissolved upon heating in a glycerol bath, yielding a yellowish solution. Upon cooling to RT, colorless transparent crystals appeared within a few hours.

CsPbI3 (X) was crystallized in the form of orange needles from HI acid using corresponding iodides in stoichiometric quantities as precursors (0.1 M PbI2, 0.1 M CsI). Powder XRD patterns were collected in the transmission mode (Debye–Scherrer geometry) using a STADI P diffractometer (STOE& Cie GmbH) equipped with a silicon strip

MYTHEN 1K Detector (Fa. DECTRIS) with a curved Ge (111)-Monochromator (CuKα1 = 1.54056 Å). Single crystal XRD measurements were conducted using a Bruker Smart Platform diffractometer equipped with an Apex I CCD detector and a molybdenum (MoKα = 0.71073 Å) sealed tube as an X-ray source. The crystals were tip-mounted on a micromount with paraffin oil. The data were processed using the APEX3 (Bruker software), and the structure solution and refinement was performed with SHELXT and SHELXL programs embedded in the Olex2 software package. For twin refinement, the cell_now algorithm (a part of the Bruker APEX2 software package) was used. The obtained crystal structures were deposited as CIF files into the CCDC database with the numbers 1552605 (I), 1552602 (III), 1552603 (II) and 1552604 (I’). UV–Vis absorbance spectra of the microcrystalline powders were collected using a Jasco V670 spectrophotometer equipped with deuterium (190-350 nm) and halogen (330-2700 nm) lamps and an integrating sphere (ILN-725, working wavelength range of 220-2200 nm). The absorbance spectra were estimated from reflectance and transmittance spectra collected from a thin layer of powder dispersed in an optically transparent Teflon grease and by diffuse reflectance transformed into absorbance using the Kubelka–Munk relation. The bandgaps were estimated from the Kubelka-Munk function [F(R∞)], by subtracting excitonic peaks from the absorption edge. First, the excitonic transition was fitted with the Gaussian function. The 2 residual spectrum was used to calculate [F(R∞) h] (where h is the incident photon energy). 2 Plotting [F(R∞) h] versus energy (h) gave a spectrum for estimation of bandgaps. Although the fitting with Gaussian function does not give an impeccable fit for the sub-bandgap region of the Urbach tail, this fit gives a possibility to estimate the bandgap values. Better understanding of the spectral region at the Urbach tail can be attained by modeling according to Eliot approach.86, 254 PL spectra were measured in a Joule–Thomson cryostat (MMR Technologies) operated in the temperature range of 78-300 K. PL emission was recorded with a heating rate of approximately 5 K/min. A 355 nm excitation source (a frequency-tripled, picosecond Nd:YAG laser, the model Duetto from Time-Bandwidth) and a CW diode laser with the excitation wavelength of 405 nm were used. Scattered laser emission was filtered out using dielectric long-pass filters with edges at 400 and 450 nm, respectively. The emission from the

Luminescent and photoconductive layered lead halide perovskite compounds comprising mixtures of caesium and guanidinium cations | Olga Nazarenko samples was collimated to an optical fiber and recorded at 1 K intervals with a spectrograph SP-2300 from Princeton Instruments coupled with a CCD array (LC100/M from Thorlabs). PL spectra were corrected to the spectral sensitivity of the setup using Planck irradiation from a calibrated halogen lamp. TR-PL measurements were performed using a time-correlated single photon counting (TCSPC) setup, equipped with a SPC-130-EM counting module (Becker & Hickl GmbH) and an IDQ-ID-100-20-ULN avalanche photodiode (Quantique) for recording the decay traces. The emission was excited by a frequency-tripled (λ = 355 nm), picosecond Nd:YAG laser, the model Duetto from Time-Bandwidth), externally triggered at a 78 kHz repetition rate. PL emission from the samples passed through sets of long- and short-pass optical filters selecting wavelength ranges 500-600 nm and 650-800 nm. Thermal analysis (TG and DTA) was performed using a NETZSCH STA 409 C/CD in an alumina crucible under argon flow (40 ml/min) with a heating rate of 10 K/min. Photoconductivity measurements were performed with illumination from a tungsten lamp dispersed by an Acton SP2150 (Roper Scientific) spectrograph/monochromator. The light was modulated by a mechanical chopper at a frequency of 13 Hz. The sample was gently pressed from two opposite sides by conductive rubber contacts in a custom-made holder and biased at 50 V through these contacts using a Keithley 236 SMU. The bias voltage was adjusted to obtain stable dark currents in the range of 1-10 nA. The signal, amplitude, and phase were measured across series of resistance by a Stanford Research 830 lock-in amplifier. The light intensity was controlled by a calibrated power detector (UM9B-BL, Gentec- EO). The current-voltage (IV) characteristics were measured using a Keithley 236 SMU. 79Br SSNMR. Experiments were performed on a 16.4 T Bruker spectrometer equipped with a 4 mm double-channel solid-state probe head and an Avance III console. A 4 mm zirconia rotor was used. Experiments were performed under static conditions without spinning the sample. Chemical shifts were referenced to 0.01 M NaBr in

D2O. Wideband, uniform rate and smooth truncation (WURST) pulses were used in a Carr- Purcell-Meiboom-Gill (CPMG) echo-train sequence in order to excite wide frequency ranges at once while maximizing detected signal intensities. WURST-80 pulses of 50 µs, 1 MHz excitation width, and 63 W power were used. 64 echos were acquired with delays of 0.045 ms, resulting in spikelet separations of 5 kHz. 4096 transients were acquired for every sub-spectrum and recycle delays of 0.1 s were applied. Electronic structures calculations. Two implementations of density functional theory (DFT) were employed. For the calculations of the band structure and density of states, the Dmol3 package was used within the Materials Studio suite.157-158 For the real space representation of the bonding, the electron localization function (ELF)159-160 was computed using Savin’s implementation within the TB-LMTO-ASA code developed at MPI Stuttgart.161 The ELF plots contain information of both ELF and electron density over a given plane and were obtained using an in-house developed code. ELF values and electron density values are shown as the color of the pixels and the number of colored pixels, respectively, over a black background. Scalar-relativistic corrections were included in all calculations because they are expected to play an important role in the presence of heavy elements. The LMTO is an all- electron method and in Dmol3 all electrons were included equally in the scalar-relativistic calculation, meaning that no pseudopotential was used. Within Dmol3 the tolerance for self-

Luminescent and photoconductive layered lead halide perovskite compounds comprising mixtures of caesium and guanidinium cations | Olga Nazarenko consistency was set to ~10-6 Ha for the total energy, using the number of k-points that results in a grid with a separation smaller than 0.03 A-1 in reciprocal space. The BOP (Becke-One Parameter) exchange-correlation functional was employed with which better outcomes were obtained for the gaps than with the commonly used PBE. For the LMTO calculations, the Langreth–Mehl–Hu exchange-correlation functional was employed. The self-consistency tolerance was set to 10-5 Ry for the total energy, and 10-5 e for the atomic charges. Taking advantage of the speed of the LMTO code, the k-space was sampled over 64-time denser grids than those used within Dmol3. The investigation of the effects on the bandgaps due to spin- orbit coupling and non-local exchange interactions, which require the extensive use of semi- empirical methods, should be the subject of future (and more computationally oriented) works.

4.3. Results and discussion Crystal structure of I, II, and III. The synthesized compounds were characterized by single-crystal XRD, powder XRD, thermal analysis, optical absorption, and PL spectroscopies. In addition, electronic structures were calculated by DFT methods. The crystal structures were solved with direct methods; light elements (C, N) were located in the difference Fourier map, and hydrogen atoms were placed at the calculated positions. Some of the crystals under investigation were twinned. Twin indexing and integration was performed with the Bruker Apex3 Program software package. The compounds crystallize in the orthorhombic crystal system (centrosymmetric space groups Pnnm, Imma, and Cmmm for I, II, and III, respectively;

Table S4.1). The crystal structures of I and II feature single layers (n = 1) of PbX6 octahedra, whereas compound III consists of a bilayer arrangement of the octahedra (Figure 4.2, Figure

4.3). To obtain the compound with n = 2, i.e., Cs2[C(NH2)3]Pb2Br7, the molar ratio of Cs:C(NH2)3:Pb was adjusted to 1:6:2 during the synthesis. Further varying the molar ratios or thermal growth profiles during the crystallization did not yield compounds other than I-III. 2D layers were separated with a mixture of organic (guanidinium) and inorganic (cesium) cations in 1:1 ratio, alternating in a periodic matter. As is typical for layered A2BX4 systems (n = 1), neighboring layers are shifted by half of the octahedra with respect to each other (Figure 4.2a), maximizing the X-to-X distance for minimal electrostatic repulsion.255 In all three compounds, the PbX6 octahedra tilt, e.g., they deviate from the idealized Pb–X–Pb angle of 180, to better + + accommodate the small Cs and large C(NH2)3 cations. Guanidinium ions closely pack, with intermolecular d(C···N) in the range of 3-3.2 Å for all three compounds (Table S4.2), and form hydrogen–halogen bonds with neighboring halogen ions (table of hydrogen-halogen bonds as

Luminescent and photoconductive layered lead halide perovskite compounds comprising mixtures of caesium and guanidinium cations | Olga Nazarenko

Figure 4.2. (a) Crystal structure of 2D Cs[C(NH2)3]PbI4 (I); (b) Hydrogen bonding between protons of guanidinium cations and iodine atoms of two 2D Pb–I layers; (c) Distorted Pb–I octahedral coordination. well as additional crystallographic data for I-III are gathered in Tables S4.3 - S4.11 and a Figure

S4.1). Cs[C(NH2)3]PbI4 (I, Figure 4.2) exhibits a somewhat higher structural disorder and stronger distortions than its bromide analog. Both bridging iodides (along the c axis) of PbI6 octahedra can be represented as split into two positions (Figure 4.2c). All Pb–I bond lengths range from 3.101(20) to 3.226(26) Å. The average Pb–I distance is 3.156(3) Å, which is slightly 36 longer than that in the cubic structure of CsPbI3, where d(Pb–I) = 3.144(7) Å, but smaller than 1 in the closely related (C4H9NH3)2PbI4, where d(Pb–I) = 3.184 Å was estimated. Below 250 K, the crystals of I undergo extensive cracking into multiple new domains, prohibiting structural description of the newly formed phase(s). However, at exactly 250 K, crystal structure refinement is still possible, and certain conclusions can be drawn from the electron density map (this material is denoted by I’). In particular, the disorder of the bridging iodides (c direction) is reduced, and they can be described by a single atomic position (Figure 4.2 c). The Pb–I distances in I’ are in the range of 3.144(95)-3.227(17) Å. Additional crystallographic information is given in the Tables S4.12-S4.14.

Compound II is a 2D compound with n = 1 (Figure 4.3a). Bridging bromides (Br3) (along the b direction) are disordered at two positions in the Pb–Br octahedra (Figure 4.4a). The distortion of the lead iodide octahedra could be described by Pb–Br3–Pb angles, which are equal to 157.488(3). Two apical and two bridging (along the a axis) bromides can be depicted as lying in the same plane with ideal 90 Br–Pb–Br’ angles, d[Pb–Br1] = 3.030(63) Å and d[Pb– Br2] = 3.002(35) Å. Bridging bromide ions (along the b direction) are asymmetrically tilted at 92.7 with respect to that plane (Figure 4.4) and have equal Pb–Br distances of 2.953(54) Å.

Luminescent and photoconductive layered lead halide perovskite compounds comprising mixtures of caesium and guanidinium cations | Olga Nazarenko

Guanidinium ions form hydrogen bonds with halogen atoms of two different layers, d(H–Br) = 2.744(72)-2.869(42) Å (Figure 4.4b). The cesium ions occupy one type of position with a CN of 10 in the smaller voids between the Pb–Br octahedra. Compound III represents a succession of Cs[C(NH2)3]PbBr4 with n = 2 (Figure 4.3b). The crystal structure of III consists of double 2D layers of Pb–Br octahedra with an average Pb–Br distance of 2.977 Å, which is slightly 6 longer than that in cubic CsPbBr3 [d(Pb–Br) = 2.937 Å]. Cesium occupies two types of positions in the structure: in the cavities within the double layers and in the space between the layers. The corresponding CNs for the two positions are 8 and 10, respectively (d(Cs–Br) = 3.775(95)-4.272(01) Å) (d(C1–N1’) = 3.014(66) Å). Guanidinium is situated in the void created by the PbBr6 octahedra of the two 2D layers and forms hydrogen bonds with the neighboring Br ions of these two layers with two sets of H–Br distances equal to 2.803(8) and 2.846(7) Å, generating a 3D network (Figure 4.3c). Bridging bromide ions (along the c direction) are now disordered at a single position (Figure 4.3d). According to the crystal structure of II, there are 3 types of bromide ions: bridging Br1, Br3 (Br3 is disordered on two positions), and terminating Br2 ions (Figure 4.4a). Solid state 79Br NMR is a powerful tool for studying the structure of the metal-halide coordination networks. Although, the signal-to-noise ratio in the obtained 79Br SSNMR spectrum is rather low, the data quality was sufficient to clearly distinguish a minimal number of two Br species

(Figure 4.5). A third Br species with a quadrupole constant (CQ) slightly larger than CQ of species 2 is suggested from the spectrum, but it could not be simulated due to an insufficient signal-to-noise ratio. CQ describes the strength of the quadrupolar interaction and is directly correlated to the electric field gradient. Therefore, the gradient in electron density is larger for species 2 than for species 1. Moreover, the CQ of species 2 is comparable to the one observed for lead-bromide perovskites (133 MHz for CsPbBr3, 141 MHz for MAPbBr3 and 149 MHz for FAPbBr3), which would permit to assume a similar coordination. Therefore, species 2 can be assigned to the bridging Br ions: Br1 and/or Br3. Meanwhile, species 1 exhibits a span comparable to PbBr2. Unlike Br1 and Br3, Br2 is a terminating ion and has a different coordination surrounding (Figure 4.4a) and should exhibit a different CQ than Br1 and Br3. Therefore, species 1, most probably, could be assigned to Br2 with a smaller CQ. Powder XRD patterns of the synthesized compounds were compared to a theoretical pattern (Figure 4.6a-c) for testing the overall purity of the substances. For compound I, no impurities were detected in the powder XRD pattern (Figure 4.6a). Compounds II and III often contain few percents of each other as an impurity.

Luminescent and photoconductive layered lead halide perovskite compounds comprising mixtures of caesium and guanidinium cations | Olga Nazarenko

Figure 4.3. (a, b) Crystal structures of 2D Cs[C(NH2)3]PbBr4 (II) and Cs2[C(NH2)3]Pb2Br7 (III) perovskites; (c) Hydrogen–halogen bonds between protons of guanidinium cations and bromides of 2D layers of Cs2[C(NH2)3]Pb2Br7; (d) Distortion of Pb–Br octahedral coordination in III; bromide split position in the crystal structure of Cs2[C(NH2)3]Pb2Br7 .

Figure 4.4. (a) Octahedral coordination of lead ions in Cs[C(NH2)3]PbBr4. Split-positions are indicated for the bridging Br ions (Br3). (b) Hydrogen bonding between the protons of guanidinium and halide ions of the perovskite layer.

Luminescent and photoconductive layered lead halide perovskite compounds comprising mixtures of caesium and guanidinium cations | Olga Nazarenko

79 Figure 4.5. (a) Br SSNMR spectra of Cs[C(NH2)3]PbBr4, spectrum acquired under static conditions (without spinning the sample) by step-wise acquisition in order to cover the wide frequency range of signals.

Figure 4.6. Powder XRD patterns of (a) Cs[C(NH2)3]PbI4, (b) Cs[C(NH2)3]PbBr4, (c) Cs2[C(NH2)3]Pb2Br7: comparison with simulated data from single-crystal analysis. Thermal stability analysis of I (d), II (e) and III (f).

Luminescent and photoconductive layered lead halide perovskite compounds comprising mixtures of caesium and guanidinium cations | Olga Nazarenko

Thermogravimetric analysis was performed under inert conditions in an argon atmosphere (40 ml/min, ambient pressure), in alumina (Al2O3) crucibles, at a heating rate of 10 K/min. Compounds I, II and III exhibit high thermal stabilities (Figure 4.6d-f). Decomposition starts at approximately 300 C. Before decomposition, an endothermic process without a mass loss occurs for all three compounds: at 213 C for II and III and at 283 C for I, which is closer to the decomposition point. These endothermic processes indicate phase transitions, which could be melting or another structural transformation, which in the case of compound I is followed by the decomposition of the compound. Absorption. Having established the crystal structures of compounds I-III, their electronic structures were examined using optical absorption spectroscopy (Figure 4.7a-b). In particular, the bandgap energies were compared with those for the known compositionally and/or structurally related compounds IV-X (Figure 4.7a-c). These reference compounds were previously reported by other groups: IV,1-3 V,4 VI,5-6 VII,7-9 VIII,10 IX,11 and X.6 In this work, compounds IV-X were grown from acidic solutions (see the Experimental section). The purity of compounds I-X was assessed by comparing their powder XRD patterns with the simulated powder patterns generated using single-crystal data (see Figure 4.6a-c, S4.2-S4.8). Single- crystal data for some reference compounds were taken from ICSD cards 92045 (for V), 97851 (for VI), and 27979 (for X) as well as from Ref.1 (for IV) and from Ref.7 (for VII). Compounds

VIII and IX, namely, CsPb2Br5 and [C(NH2)3]2PbBr4, have not been structurally characterized prior to this study. Therefore, their crystal structures were determined by single-crystal XRD (for crystallographic details see Figure 4.8 and Tables S4.15-S4.20 and ICSD entries 432533 and 432534) and found good agreement with the powder patterns (Figures S4.6, S4.7).

CsPb2Br5 (VIII, Figure 4.8 b,c) is a 2D compound with layers constructed of Pb–Br face- sharing distorted square antiprisms and cesium ions located in the interlayer space. This compound is often reported for Cs–Pb–Br systems, along with 0D perovskite Cs4PbBr6 and 3D 5, 50, 256 perovskite CsPbBr3. The single-crystal diffraction data for CsPb2Br5 are consistent with 10 a powder diffraction pattern reported in Ref. and in PDF2 (00-025-0211). [C(NH2)3]2PbBr4 (IX, Figure 4.8 a) crystallizes in triclinic P-1 space group with a structural motif similar to 257 [C(NH2)3]2SnCl4, as reported by Szafranski et. al. [C(NH2)3]2PbBr4 consists of 1D chains of 2 corner-sharing [PbBr5] square pyramids, wherein the CN of Pb(II) is 6, including the lone 6s pair of electrons. The guanidinium cations are located between the inorganic chains and form hydrogen–halogen bonds, resulting in a 3D supramolecular structure. This structural analysis underlines the importance of Cs+ ions for building a stable 2D perovskite of compound II.

Interestingly, the iodide compositional analog, [C(NH2)3]2PbI4 (V), does form a 2D lattice, but unlike I, it is made of corrugated (zig-zag) layers (see the structural comparison between I, V and IV in Figure 4.9 a-c).1, 4 Figure 4.7 a-b and Figure 4.10 present the normalized Kubelka–

Munk (KM) functions, F(R∞), for all ten compounds. This function is derived from the diffuse reflectance spectra of the powdered samples. This function is estimated as F(R∞) = /S = (1- 2 R∞) /2R∞, where  is the absorption coefficient, S is the scattering coefficient, and R∞ is the reflectance of an infinitely thick layer. Bandgap energies were estimated by plotting [F(R∞) 2 h] (where F(R∞) is in Kubelka–Munk units, and h is the incident photon energy) versus

Luminescent and photoconductive layered lead halide perovskite compounds comprising mixtures of caesium and guanidinium cations | Olga Nazarenko energy (h). The excitonic peaks were subtracted from the Kubelka-Munk function for the bandgap determination due to the excitonic enhancement of the absorption edge (Figure 4.10 b).86 Conventional optical absorption spectra for all compounds I-X are presented in Figure S4.9.

The most important comparison is between (n-C4H9NH3)2PbX4 (IV and VII) and

Cs[C(NH2)3]PbX4 (I and II) because these compounds consist of flat sheets of PbX6 (unlike the corrugated sheets in V). The bandgaps for the Cs–guanidinium-based compounds are smaller

2 Figure 4.7. (a, b) Kubelka–Munk function F(R∞) = (1-R∞) /2R∞ (R∞ - diffusive reflectance) of 2 I-IX. (c) Table of compounds and their bandgaps determined from [F(R∞)hν] , where F(R∞) is in Kubelka–Munk units, and hν is the incident photon energy.

Figure 4.8. (a) The crystal structure of [C(NH2)3]2PbBr4: 1D chains of corner-sharing square pyramids. (b, c) The crystal structure of CsPb2Br5.

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Figure 4.9. (a) The crystal structure of Cs[C(NH2)3]PbI4 (I). (b) The crystal structure of (n- C4H9NH3)2PbI4 [IV, (crystallographic information (cif) downloaded from D.G. Billig, A. Lemmer, Acta Crystallogr. Sect. B 2007, 63 (5), 735-747]. (c) The crystal structure of [C(NH2)3]2PbI4 (V, cif.-file was downloaded from M. Szafranski and A. Katrusiak, Phys. Rev. 2000, B62 (2), 8787, and related ICSD card 92045). by 0.1-0.25 eV. This observation correlates well with the smaller interlayer ions reducing the average thickness of one layer (center-to-center distance) and, hence, also reducing the electronic isolation of the PbX6 layers. For the case of iodides, the average thickness of one layer is 9.78 Å in I (13.8 Å in IV); correspondingly, the specific density of I is 1.54 greater than that of IV (4.162 g cm-3 vs. 2.702 g cm-3), whereas their molecular weights are similar (907.79 g mol-1 and 863.109 g mol-1, respectively); pointing to much smaller molar volumes obtained with smaller interlayer cations. Interestingly, despite its similar density, compound V shows a wider bandgap of 2.44 eV (Figure 4.7 c), arguably due to the corrugated 2D structure (Figure 4.9 c). Another important factor is a distortion of the Pb-X octahedra. It can be described by a deviation of Pb-X-Pb angles from 180 ⁰. Moreover, in-plane (in the plane of propagation of inorganic layers) distortions were shown to have a more significant influence on the bandgap 89, 258 than the out-of-plane deviations. When comparing (n-C4H9NH3)2PbI4 and Cs[C(NH2)3]PbI4 (Figure 4.11), the difference in the bandgaps is consistent with a degree of distortion of Pb-I octahedra. The in-plane deviation of the Pb-I-Pb angle from 180 ⁰ for (n-

C4H9NH3)2PbI4 is as significant as 24.36 ⁰, while for Cs[C(NH2)3]PbI4 the maximum angle of deviation is 18.43 ⁰. The Kubelka–Munk absorption functions of 2D perovskites of variable- slab-thickness II and III point to smaller bandgap, as expected, for thicker Pb–Br slabs of III. In addition, the effect of the interlayer spacing is clearly seen in the approximately 0.25 eV wider gap of VII (flat 1D sheets) than that of II. Both II and III exhibit intermediate bandgap values compared to those of compounds with higher and lower dimensionalities, i.e., CsPbBr3 (VI, 3D) and [C(NH2)3]2PbBr4 (IX, 1D), respectively. Photoluminescence. Optical emission spectra (Figure 4.12) were recorded in the range of 77-300 K for I, II and III. Only compound III exhibits a measurable PL intensity at RT (Figure 4.12 c). All three compounds are brightly emissive to the naked eye under UV light at

Luminescent and photoconductive layered lead halide perovskite compounds comprising mixtures of caesium and guanidinium cations | Olga Nazarenko

77 K. For this temperature, Figure 4.12 d-f present the luminescence spectra, while the entire temperature dependence is plotted as color maps in Figure 4.12 a-c. The presence of minor emission peaks in Figure 4.12 b, c suggests that compounds II and III contain small impurities of each other, as also concluded from the powder XRD patterns. For the PL spectrum of I, in addition to the main emission peak at 535 nm, a broader, low-intensity feature at 650-800 nm was detected. In order to understand the origin of this feature, we performed TR-PL at various temperatures (Figure 4.13). There can be seen that while the emission at 535 nm decays very fast, within 10 ns, the long wavelength band relaxes much slower (sub-µs range) and becomes even slower upon cooling to 200 K. These features of the spectrally broad and long-living emission are consistent with the phenomenon of STEs that was recently observed in 2D layered perovskites.240, 258-259 Often, the emission of STEs accompanies narrow-band shorter– wavelength excitonic emission, such as in the present case.240, 258 Thus, we believe that we observed an excitonic emission at 535 nm and STEs at 650-800 nm.

2 Figure 4.10. (a) Kubelka–Munk function F(R∞)=(1-R∞) /2R∞ (R∞ - diffusive reflectance) of Cs[C(NH2)3]PbI4 (I) and CsPbI3 (X). (b) Kubelka–Munk function for Cs[C(NH2)3]PbI4 (I), a Gauss fit of the excitonic peak (green line), the Kubelka-Munk function without the excitonic 2 enhancement (red line) and a Tauc plot (insert) represented as [F(R∞) hν)] (where hν is the incident photon energy) versus energy (hν) for I.

Luminescent and photoconductive layered lead halide perovskite compounds comprising mixtures of caesium and guanidinium cations | Olga Nazarenko

Figure 4.11. (a) Crystal structure of Cs[C(NH2)3]PbI4 (I) and (n-C4H9NH3)2PbI4 (IV) (b). (c) A table of the distortions of Pb-X-Pb angles for I and IV, where Din = 180 ⁰- Ɵin was calculated as reported by Smith et. al. [Chem. Sci. 2017, 8, 4497-4504].

Figure 4.12. (a-c) Normalized temperature-dependent PL spectra. Color bar (on the right) represents intensity for all PL matrices. (d-f) Spectra at 77 K for Cs[C(NH2)3]PbI4 (I), Cs[C(NH2)3]PbBr4(II), and Cs2[C(NH2)3]Pb2Br7 (III).

Luminescent and photoconductive layered lead halide perovskite compounds comprising mixtures of caesium and guanidinium cations | Olga Nazarenko

Figure 4.13. Time-resolved PL (TR-PL) traces recorded in spectral ranges: 500-600 nm (a) and 650-800 nm (b) at various temperatures. Insert shows early decay for 500-600 nm emission. Photoconductivity. Semiconductive properties of the obtained compounds are manifested in the production of free charge carriers under light excitation and in the photoconductivity effects. The dark specific resistivity values are approximately 3108 Ω∙cm and 5109 Ω∙cm for compounds I and III, respectively. The photoconductivity spectra (Figure 4.14 a) of two single crystals of compounds I and III exhibit similar features. First, a shoulder on the shorter wavelength side corresponds to bandgap absorption. Second, sharp peaks appear -1 -1 at approximately 1 mA∙W for Cs[C(NH2)3]PbI4 (I) and 9 mA∙W for Cs2[C(NH2)3]Pb2Br7 46, 50, 96 (III). Such photopeaks have been recently reported for APbX3 perovskite single crystals and can be explained by the thickness (depth)-dependent absorption and charge transport phenomena. Two competing effects determine the photocurrent in this case. First, the charges produced in the crystal volume are more effectively collected on the electrodes, while the charges produced on or close to the surface demonstrate decreased charge transport characteristics, likely due to the trapping of surface charges. When this trapping effect is combined with the absorption spectrum, a peak is expected at the onset of the absorption, not at the absorption peak. The wavelength of the peak, which is essentially equal to the bandgap energy, is the wavelength at which sufficiently intense light propagates into the bulk region. At higher energies (shorter wavelengths), the higher absorption coefficient will limit the absorption to the near-surface regions. A third feature that differentiates compounds I and III 46, 50, 96 form APbX3 perovskite single crystals is the long-wavelength photocurrent tail at energies well below the bandgap; this tail is most pronounced for compound I. This region may be attributed to the photocurrent involving the mid-gap trap states. Interestingly, the charges produced in the single crystal volume (sharp peak and shoulder at longer wavelengths) are delayed compared to charges from the surface, corresponding to a decrease in the phase of the photoresponse (Figure 4.14 b). This may be due to the moderate charge mobility in these compounds.

Luminescent and photoconductive layered lead halide perovskite compounds comprising mixtures of caesium and guanidinium cations | Olga Nazarenko

Figure 4.14. (a) Room-temperature photoconductivity of Cs[C(NH2)3]PbI4 (I) and Cs2[C(NH2)3]Pb2Br7 (III) single crystals. (b) Photoresponse phase for single crystals of Cs[C(NH2)3]PbI4 (I) (red) and Cs2[C(NH2)3]Pb2Br7 (III) (blue). The decrease in phase indicates the delay with charge extraction for the photocarriers generated in the bulk of the crystal in comparison to carriers generated at the surface (spectral range of high absorbance). Electronic structure. Compounds I, II, and III have qualitatively similar electronic structures. We chose Cs[C(NH2)3]PbBr4 (II) as an example for the following detailed discussion. The quantitative differences such as the bandgaps and their relation to composition are addressed below. The density of states (DoS) per primitive cell is plotted in Figure 4.15. The solid black line represents the sum of all valence states, and the contributions of selected species, i.e., the partial density of states (PDoS), is represented by colored areas. The interaction of bromine with the other atoms is rather ionic, as can be inferred from the nearly 100% coverage (magenta areas) of the DoS in the regions where bromine contributes. The main contribution of lead to the occupied states originates from the electrons in the 6s orbitals at approximately 8 eV below the Fermi level, represented by the dashed gray area. The remaining valence electrons belong to guanidinium cations. The bonding, non-bonding and antibonding combinations of the CN3 π-system are labeled in Figure 4.15. These states are not stabilized by the interaction with the protons and appear higher in energy. The non-bonding section has a negligible contribution from carbon (cyan area) as expected for the nearly ideal D3h symmetry of guanidinium within this structure. The π contribution to the C–N bonds is also visible in the electron localization function (ELF, shown in Figure 4.16),159 where the bc cut (left side) and the ac cut (right side) show different widths of the bonding attractor; i.e., the C–N bond cross- section has an oval shape. The ELF also confirms the dominant ionic interaction between all Pb–Br, and Cs–Br pairs, with the formation of nearly spherical closed shells. In the present case, the guanidinium–Br hydrogen bond is the proton-bridge type of bond. Due to the deficit of one electron in the cation, the electron cloud is pulled by nitrogen and carbon toward the center, leaving the positive protons as an outer layer to interact with the bromine atoms. Accordingly, the very small contribution of hydrogen to the DoS (not highlighted in Figure

Luminescent and photoconductive layered lead halide perovskite compounds comprising mixtures of caesium and guanidinium cations | Olga Nazarenko

4.15) shows negligible mixing with bromine states. Note that the closed shell-like ELF attractor of Br is distorted in a way that “avoids” the protons, which is typical for ionic interactions. All three compounds feature guanidinium cation stacking in alternating orientations along a crystallographic axis, which suggests a stabilizing interaction between neighboring cations. It was found that this stacking allows van der Waals interactions involving the easily polarizable π electron cloud of guanidinium cations. This conclusion is corroborated by the small contribution of the carbon states to the non-bonding peak just below -2.5 eV (see Figure

4.15). The latter can appear only if the D3h symmetry is broken and the π electron cloud is polarized. At the same time, the π-bonding band appearing at -6.8 eV (see Figure 4.17) disperses along the W-R segment in reciprocal space. This band dispersion of approximately 0.7 eV is significant, especially when compared to the generally flat band structure of such systems. The alternate stacking allows the small energy cost for polarizing the shells to be overcompensated by the interaction between the induced dipoles, resulting in higher stability. Compounds I, II, and III are all semiconductors. DFT predicts that the highest occupied and lowest empty bands predominantly consist of halide p-orbital and Pb p-orbital contributions, respectively. The calculated bandgaps are approximately 2.7 eV for Cs[C(NH2)3]PbBr4, 2.2 eV for

Cs2[C(NH2)3]Pb2Br7, and 2.0 eV for Cs[C(NH2)3]PbI4. This trend is consistent with the measured absorption spectra (Figure 4.7) and can be understood in terms of basic quantum mechanical principles. The bandgap increases with the confinement, i.e., with the decreasing lead halide slab thickness (from n = 2 to n = 1) or with moving from heavier iodide to lighter bromide. Bromine is more electronegative than iodine, and therefore, the center of the valence and conduction bands should be closer in energy for the iodides (compound I), resulting in a smaller gap in this case.

Figure 4.15. Electronic density of states (DoS) per primitive cell per eV in Cs[C(NH2)3]PbBr4 (II).

Luminescent and photoconductive layered lead halide perovskite compounds comprising mixtures of caesium and guanidinium cations | Olga Nazarenko

Figure 4.16. ELF planar cut at x = 0 (left) and y = 0 (right), each covering 2  2 unit cells of Cs[C(NH2)3]PbBr4. Crystallographic sites are labeled with the name of the corresponding elements.

Figure 4.17. Electronic band structure of Cs[C(NH2)3]PbBr4 (II) from DFT calculations. The flatness of the majority of the occupied bands indicates the localized character of the states. Note the exceptional dispersion along the W-R segment of the π-bonding guanidinium states (in cyan at approximately -6.8 eV).

Luminescent and photoconductive layered lead halide perovskite compounds comprising mixtures of caesium and guanidinium cations | Olga Nazarenko

4.4. Conclusions In conclusion, in this Chapter the synthesis and characterization of three layered LHP compounds, containing a mixture of two stable interlayer cations in their structure (cesium and guanidinium): Cs[C(NH2)3]PbI4 (I), Cs[C(NH2)3]PbBr4 (II), Cs2[C(NH2)3]Pb2Br7 (III) was presented. The two interlayer cations act synergistically to enable stable 2D perovskites through the van der Waals interactions between guanidinium cations and through hydrogen–halogen bonds between guanidinium cations and halides ions of different layers. Cs+ ions fill the small voids within the Pb–Br layer in III, in addition to also occupying the interlayer space. All three compounds were shown to possess high thermal stability, up to above 300 C. The compounds exhibit clear signatures of strong quantum confinement. Narrow-band photoconductive response was observed, indicating the balance between the optical absorption effects and the varying charge transport efficiency in the bulk and at the surface of the material. These findings have strong relevance to perovskite thin-film photovoltaic research. In particular, recent studies have increasingly departed from the archetypical APbX3-like 260 compound, CH3NH3PbI3, to multi-cation and multi-ionic formulations such as Ref. , wherein + + + + four A-site cations are used (CH3NH3 , CH(NH2)2 , Cs , Rb ). Moreover, adding guanidinium ions has been reported to improve solar cell characteristics.261-262 Clearly, structural studies within this compositional space, as shown here for mixed-cation Cs–guanidinium systems, might unravel the complexities of such optoelectronic systems.

Reproduced with modifications from: O. Nazarenko, M. R. Kotyrba, M. Wörle, E. Cuervo-Reyes, S. Yakunin, M. V. Kovalenko. Luminescent and photoconductive layered lead halide perovskite compounds comprising mixtures of cesium and guanidinium cations. Inorg. Chem., 2017, 56, 11552-11564. Copyright 2017, American Chemical Society. https://pubs.acs.org/doi/abs/10.1021/acs.inorgchem.7b01204

Guanidinium-formamidinium lead iodide: a layered perovskite-related compound with red luminescence at room temperature | Olga Nazarenko

Chapter 5. Guanidinium-formamidinium lead iodide: a layered perovskite-related compound with red luminescence at room temperature

5.1. Introduction Layered LHPs are intensely studied as unique structurally soft and electronically quantum-well-like systems, exhibiting enormous structural diversity2, 9, 79, 82, 91, 168, 258, 263 as well as high potential for use in light-emitting diodes,91, 226, 264 photodetectors83, 265 or solar cells.244, 247 Most of 2D LHPs can be imagined as structural derivatives from the ideal cubic perovskite + lattice of an APbX3 composition (A – monovalent cation, primarily Cs , MA or FA, X – halogen anion) with 3D interconnection of corner-shared PbX6-octahedra and with A-cations filling the large voids in-between these octahedra. Slicing of this lattice along (100), (110) and (111) crystallographic planes (Figure 5.1) with the elimination of the octahedra lying in the slicing plane (gray-shaded octahedra in Figure 5.1) leads to common 2D-compounds with single- octahedra or thicker slabs. To induce the formation of 2D LHPs, A-site cations, which are typically used in 3D LHPs (Cs+, MA, FA), must be fully or partially replaced with bulkier cations; for instance, long-chain alkylammonium ions. Generally, (100)-layered 2D perovskites are most common and adopt K2NiF4 or RbAlF4 crystal structure. As to (110)-2D LHPs (Ca2Nb2O7 related phases), some known examples of interlayer cations include N-(3- 266 1 aminopropyl)imidazolium [in (C6H13N3)PbBr4], N -methylethane-1,2-diammonium [N- 264 + 4, 253 MEDA, in (N-MEDA)PbBr4], as well as guanidinium [C(NH2)3 , G, in G2PbI4]. An example of (111)-layered 2D LHP is the one comprising 3-(aminomethyl)pyridinium [(H23- 2+ 267 AMP) ], such as in (H23-AMP)2PbBr6 with K2PtCl6 type structure. (100)-2D LHPs are most commonly built with primary mono and diammonium cations.79 The heavier the halide ion and the thicker the slab the smaller is the optical bandgap energy (Eg). As an example, in iodides, 87 the absorption edge and PL maxima can be tuned from 2.57 eV in (CH3(CH2)3NH3)2PbI4 to 81 1.83 eV in (CH3(CH2)3NH3)2(CH3NH3)4Pb5I16.

Guanidinium-formamidinium lead iodide: a layered perovskite-related compound with red luminescence at room temperature | Olga Nazarenko

Figure 5.1. Derivation of 2D LHPs from the parental cubic perovskite lattice of 3D LHPs by cutting the latter along typical crystallographic planes - (100), (110) and (111). Large Stokes shifts and a broadband emission were recently reported for some representatives of (100) and (110) perovskites.91, 226, 259, 264 These phenomena were attributed to emission either by STEx or from charge carriers trapped on defects.268 The guanidinium cation is an interesting example of a small and highly symmetric organic molecule that forms a corrugated G2PbI4 structure with layers parallel to the (110) plane, with a bandgap energy of 2.40 eV.253 A similar but smaller cation, FA, forms instead a 200 3D LHP - FAPbI3 (α-phase, cubic symmetry) – with the smallest bandgap energy amongst all known LHPs, exhibiting PL maximum at 1.48 eV, corresponding to 840 nm in the near- o 61 infrared. α-FAPbI3 is thermodynamically stable only above 185 C, whereas at RT, α-FAPbI3 undergoes a phase-transition into a wide-bandgap, non-luminescent 1D-polymorph (δ-FAPbI3, 61, 200 hexagonal, P63mc space group). The difference between the crystal structures of α-FAPbI3 and G2PbI4 results from different ionic radii of FA and G cations - 253 pm and 278 pm, respectively.39 The tolerance factor describes a probability of a 3D perovskite lattice formation.39 Thus, when α = 0.9-1.0, formation of a cubic perovskite structure is favored and deviations from these values lead to distorted perovskites or non-perovskite, usually lower- dimensionality structures. While for FAPbI3 α = 0.987, for GPbI3 it increases to 1.039, causing crystallization in a different crystal structure made of 2D corrugated inorganic layers, with a general formula of the compound G2PbI4. Both FA and G ions are excellent donors of hydrogen bonds with the lead-halide framework. It is thus of high interest to explore the possibility of concomitant incorporation of FA and G for the formation of previously unknown compounds. This would lead to new compounds with the reduced volume of the interlayer region as compared to G-only compounds or those employing more traditional, larger spacer cations such as alkylamines. Such compactness of G has been shown to enable smaller bandgaps (for a given slab thickness).269-270 G and similar ions are increasingly used as additives to MA/FA/Cs-based 271-272 APbX3 compounds in perovskite photovoltaics. Hence, the rational synthesis and search for structural motifs within this compositional space is highly relevant for understanding and controlling the performance of perovskite solar cells as well as other perovskite optoelectronic devices. In this part of the Ph.D. project, a new member of hybrid 2D perovskites with a crystal structure composed of corrugated inorganic layers - [CH(NH2)2][C(NH2)3]PbI4 (Figure 5.2a-c),

Guanidinium-formamidinium lead iodide: a layered perovskite-related compound with red luminescence at room temperature | Olga Nazarenko further denoted also as FAGPbI4, was identified in the FA-G-PbI2 system. Such a layered network can be visualized as stair-like shaped slabs of corner-sharing Pb-I octahedra (Figure 5.2). It shows broad emission in the red region at RT (Figure 5.2d) that originates from both free and STEs.

5.2. Experimental section

Materials. Lead (II) acetate (Pb(OAc)2×3H2O, ≥99.99%) was purchased from Sigma Aldrich. Guanidinium carbonate (G2CO3, 99+%) was purchased from Acros. Hydrogen iodide

(HI, 57%, stabilized with 1.5% hypophosphorous acid, H3PO2) and formamidinium acetate [FA(OAc), 99%] were received from ABCR. Diethyl ether was used from different sources. Toluene (> 99.7%) was purchased from Fisher. All chemicals were used as received without further purification.

Synthesis of FAGPbI4, method 1. FAGPbI4 was crystallized from hot hydroiodic acid solution. FA(OAc) (4 mmol), G2CO3 (4 mmol) and Pb(OAc)2×3H2O (4 mmol) were loaded into a 25-ml round-bottom flask equipped with a magnetic stirring bar and a septum. The flask was then cooled with an ice bath, followed by the addition of HI (18 mL). Gas evolution was observed, while a mixture of brown and yellow precipitates formed. This mixture was heated with a glycerol bath (50-55 °C) with magnetic stirring until a clear yellow solution formed. Then the solution was cooled to RT, the stirring was stopped and the stirring bar removed. After a few hours, small brown crystals appeared. Crystals were allowed to crystallize for an additional 12 h and then were separated by vacuum filtration and washed with diethyl ether and toluene. Finally, the product was dried in a vacuum oven (10-50 mbar) overnight at 55 °C. A yield of 14 % was estimated relative to the initially loaded Pb(OAc)2×3H2O. To grow larger crystals (few mm), the initially obtained smaller crystals were carefully introduced into another vessel with a fresh reaction mixture (at the point of cooling it to RT), and left undisturbed for 1-2 days, followed by washing with diethylether and toluene.

Method 2 (higher yield). FA(OAc) (6.75 mmol) G2CO3, (4.5 mmol) and Pb(OAc)2×3H2O (4.5 mmol) were loaded into a 50-ml round-bottom flask equipped with a magnetic stirring bar and a septum. The mixture was cooled with an ice bath, and 20 ml of HI was added. A mixture of yellow and brown precipitates formed immediately, which fully dissolved upon further heating with a glycerol bath (60-70 °C) forming a clear yellow solution. During cooling to RT a yellow precipitate crystallized out. At RT, yellow needles began to redissolve and dark brown needles appeared. After ca. 1-2 h, the mixture was shaken to promote complete dissolution of the yellow precipitate (that remains below brown crystals). After 5 hours, the brown product was separated by vacuum filtration and washed with diethylether and toluene. A yield of 40 % was estimated relative to the initially loaded Pb(OAc)2×3H2O. 207 Reference compounds for characterization of FAGPbI4 by Pb SSNMR. G2PbI4 was obtained as yellow platelets from a hot acidic solution using [C(NH2)3]2CO3 (3 mmol), PbI2 (2 mmol) and 7.5 ml of HI (57%). The crystals were separated by vacuum filtration, washed with diethyl ether and dried in a vacuum oven at 60 C. A yield of 30 % was estimated relative to the initially loaded PbI2.

Guanidinium-formamidinium lead iodide: a layered perovskite-related compound with red luminescence at room temperature | Olga Nazarenko

148 α-FAPbI3. The synthetic procedure was adapted from Saidaminov et al.. The concentration of [Pb] and [FA] were adjusted to 1 M. Black crystals of α-FAPbI3 were taken out from the hot gamma-butyrolactone solution, dried with a filter paper and annealed under inert condition (Ar atmosphere) at 180 °C for 2 h. Crystals were then kept at 120 °C (overnight) until the 207Pb NMR measurement was conducted.

δ-FAPbI3 was crystallized from hot HI acid. For this FA(OAc) (1 mmol) and

Pb(OAc)2×3H2O (1 mmol) were dissolved in 3 ml of hot HI acid. A clear yellow solution resulted, and the solution was cooled to RT. A yellow product crystallized out upon cooling to RT. This was isolated by vacuum filtration, rinsed with diethyl ether and dried in a vacuum oven at 50 °C. A yield of 30 % was estimated relative to the initially loaded Pb(OAc)2×3H2O. Powder XRD patterns were collected in transmission mode with a STADI P diffractometer (STOE&Cie GmbH), equipped with a curved Ge (111)-Monochromator

(CuKα1=1.54056Å) and a silicon strip MYTHEN 1K Detector (Fa. DECTRIS). For the measurement, the ground powder was placed between the adhesive tape. Single crystal XRD measurements were conducted on Bruker Smart Platform diffractometer equipped with an Apex I CCD detector and molybdenum (MoKα=0.71073 Å) sealed tube as an X-ray source. Crystals were tip–mounted on a micromount with paraffin oil. The data was processed with APEX3 (Bruker software)273, structure solution and indexing were performed with SHELXS154 and SHELXL155 respectively, embedded in the Olex2 package.156 For twin indexing the cell_now algorithm (a part of the Bruker APEX3 software package) was used. The crystal structure of the synthesized compound was solved with direct methods, light elements (C, N) were located in the difference Fourier map, most of the positions of the cations were refined as rigid groups, and hydrogen atoms were placed at calculated positions. UV-Vis absorbance spectra of the microcrystalline powders were collected using a Jasco V670 spectrophotometer equipped with deuterium (D2) lamp (190 – 350 nm) for use in UV, a halogen lamp (330 – 2700 nm) for use in UV/NIR, and an integrating sphere (ILN-725) with a working wavelength range of 220 – 2200 nm. The absorbance spectra were estimated from diffuse reflectance measured on the powdered crystals transformed into the Kubelka-Munk function. PL spectra were measured in a Joule– Thomson cryostat (MMR Technologies) operated in the temperature range of 78-300 K. PL emission was recorded with a heating rate of approximately 5 K/min. A 405 nm diode laser with a wavelength of 405 nm was used for excitation. Scattered laser emission was filtered out using dielectric long-pass filters with edges at 400 and 450 nm, respectively. The emission from the samples was collimated to an optical fiber and recorded at 1 K intervals with a CCD spectrometer LR1 (Aseq Instruments). PL spectra were corrected to the spectral sensitivity of the setup using Planck irradiation from a calibrated halogen lamp. TR-PL measurements were performed using a time-correlated single photon counting (TCSPC) setup, equipped with an SPC-130-EM counting module (Becker & Hickl GmbH) and an IDQ-ID-100-20-ULN avalanche photodiode (Quantique) for recording the decay traces. The emission was excited by a BDL-488-SMN laser (Becker & Hickl) with a pulse duration of 50 ps and wavelength of 486 nm, the CW power equivalent of ~0.5 mW, externally triggered at a 1 MHz repetition rate. PL emission from the samples passed through a long-pass optical filter with an edge at 500 nm in order to reject the excitation laser line. Average radiative lifetimes were determined as: 휏푎푣푔 =

Guanidinium-formamidinium lead iodide: a layered perovskite-related compound with red luminescence at room temperature | Olga Nazarenko

2 2 ∑푖=1 휏푖 ∙퐴푖 2 , where Ai and τi are corresponding amplitudes and exponential decay parameters in the ∑푖=1 휏푖∙퐴푖 biexponential analysis. Absolute values of photoluminescence quantum yield (PLQY) was measured at Quantaurus-QY spectrometer from Hamamatsu in powder mode. Micro- photoluminescence spectra were recorded in a home-built setup. The sample, i.e., a single crystal of FAGPbI4, was mounted on xyz positioning stages (SmarAct, SLS 32:32; 1 nm resolution). A 405 nm laser diode (PicoQuant, LDH- D-C-405; repetition rate=2.5 MHz; excitation power density=90µJcm-2) was focused on the sample by an oil immersion objective (Olympus, 100X; NA=1.4). PL was recorded by the same objective and analyzed by a monochromator (500 mm focal length; 300 lines/mm grating) coupled to an EMCCD camera (Princeton Instruments ProEM: 1600x400B_eXcelon3). For time-resolved PL, a single photon avalanche diode (PicoQuant PDM, 50 ps time resolution) was coupled at the exit port of the monochromator. TR-PL traces were recorded by using a time-correlated single photon counting (TCSPC) module (PicoQuant, PicoHarp 300). Photoconductivity measurements were performed with illumination from a tungsten lamp dispersed by an Acton SP2150 (Roper Scientific) spectrograph/monochromator. The light was modulated by a mechanical chopper at a frequency of 9 Hz. The sample was contacted by soft conductive rubber from opposite sides along the shorter axis of the crystal and biased at 100 V through these contacts using a Keithley 236 SMU. The signal, amplitude, and phase were measured across series of resistance by of a Stanford Research 830 lock-in amplifier. The light intensity was controlled by a calibrated power detector (UM9B-BL, Gentec-EO). Thermal analysis (TG and DSC) was performed using a Netzsch Simultaneous Thermal Analyzer (STA 449 F5 Jupiter). A powdered sample (15.8 mg) was placed in an alumina crucible and heated under Ar gas flow (50 ml/min) to 800 C (10 °C min-1). 207Pb SSNMR. All experiments were performed on a Bruker spectrometer (16.4 T) equipped with a 2.5 mm triple-channel solid-state probe head and an Avance III console. A 2.5 mm zirconia rotor was used. The spinning frequency was set to 20 kHz except for α-FAPbI3 which was measured in static mode. Chemical shifts were referenced to PbMe4. A Hahn echo pulse sequence was used. The 90-degree pulse had a length of 5 µs. The echo delay was set to 1 rotor cycle and the recycle delay to 0.5-2 seconds. For every spectrum, between 5’000 and 786’432 scans were recorded.

5.3. Results and discussion

Synthesis and a crystal structure of FAGPbI4. FAGPbI4 was crystallized from hydroiodic acid in the form of 0.5-3 mm long needles (Figure 5.2c). The purity of the synthesized compound was estimated by powder XRD, and the thermal stability was studied with TG as well as DSC (Figure 5.3). FAGPbI4 crystallizes in a monoclinic crystal system (space group C2/m, Tables 5.1, S5.1, S5.2). The crystal structure is composed of corrugated Pb- I layers and G and FA cations situated in the interlayer space (Figure 5.2a, b). Interestingly, in comparison with a related compound G2PbI4 (included in Figure 5.4), composed of zig-zag Pb- 226 I layers, FAGPbI4 shows stair-like corrugation of inorganic layers. Concomitant use of FA and G ions in a lead iodide system results in a different crystal structure as compared to MA- 269-270 G-PbI2 and Cs-G-PbI2 systems, where Ruddlesden-Popper-like phases form (Figure 5.5).

Guanidinium-formamidinium lead iodide: a layered perovskite-related compound with red luminescence at room temperature | Olga Nazarenko

Figure 5.2. (a) Crystal structure of FAGPbI4. (b) Octahedral coordination of lead ions (ellipsoids shown at 50 % probability). (c) A photograph of typical crystals. (d) Photographs of a FAGPbI4 crystal taken under daylight and under UV illumination in dark. (e) Kubelka– 2 Munk function F(R∞)=(1-R∞) /2R∞ (R∞ - diffusive reflectance) and (f) temperature-dependent PL spectra measurements performed on powdered crystals.

The inorganic framework in FAGPbI4 is rather similar to that of G2PbI4. This can be explained by relatively large size of FA cation as compared to MA [ri (MA) = 217 pm vs. 253 pm for FA]; also causing thermodynamic instability of α-FAPbI3 at RT. The crystal structure of FAGPbI4 had been deposited as a CIF file into the CCDC database with the number 1814744. In general, the bandgap energy is determined by such factors as the dimensionality, connectivity of the structure and distortions of the coordination polyhedra (lengths of the Pb-X bonds and Pb-X- Pb angles), as well as it is inversely proportional to the atomic number of the 207, 224 halide. In FAGPbI4, Pb-I distances lie in the range of 3.080(1)-3.361(1) Å and are 4 comparable to those in G2PbI4 [d(Pb-I) = 3.070(14)-3.362(14) Å]. I-Pb-I bond angles in FAGPbI4 deviate from 90⁰ and are in the range of 86.36(3) – 95.66(2)°; Pb-I-Pb angles range from 161.98° to 180°. In comparison, Pb-I-Pb angles in G2PbI4 are in the range of 156.9-177.5°. The optical measurement. The optical absorption spectrum of powdered FAGPbI4 is 2 represented here by the Kubelka-Munk function F(R∞)=/S=(1-R∞) /2R∞. The absorption transitions appear at 506 nm and 459 nm (Figure 5.2e, S5.1), with only a weak absorbance below 620 nm. The compound exhibits red PL under UV excitation at RT (Figure 5.2d), which is a spectrally broad convolution of several bands (Figure 5.2f). At RT, the PLQY reaches 3.5 %, when excited in the wavelength range of 550-600 nm and drops to about 1 % for the excitation at wavelengths shorter than 500 nm (Figure S5.2). Solid-state NMR, in standard magic-angle spinning version had been used to delineate whether the PL originates from amorphous/nanocrystalline contaminants, not detectable by 207 powder XRD. The high purity of FAGPbI4 was confirmed by comparing its Pb SSNMR spectrum with the spectra of plausible contaminants - ternary phases of G2PbI4, α-FAPbI3, and

δ-FAPbI3 (see NMR data in Figure 5.4, S5.3a, and S5.6, as well as Table S5.3 and, for structural

Guanidinium-formamidinium lead iodide: a layered perovskite-related compound with red luminescence at room temperature | Olga Nazarenko characterization of all compounds by XRD - Figures S5.3b,c, S5.4, S5.5). FAGPbI4 exhibits two peaks, at 1160 and 1465 ppm (referenced to PbMe4), with a FWHM of 22.4 and 20.2 kHz, which are typical values for inorganic lead (II)-iodide compounds274-276 and LHPs.277-278 The ratio of the peak areas of 2:1 corresponds to the staggered crystal structure. 207Pb SSNMR spectra of G2PbI4 and α-FAPbI3 exhibit single peaks at 1515 ppm. The narrower signal of α- FAPbI3 (FWHM = 22.2 kHz), as compared to G2PbI4 (FWHM = 24.5 kHz), can be attributed

Figure 5.3. (a) Powder XRD pattern of FAGPbI4 compared to the calculated pattern based on single crystal XRD measurements. A preferred orientation of the microcrystalline powder is observed through the difference in intensity of certain reflections in the experimental pattern comparing to the calculated data due to the fact that FAGPbI4 crystallizes in the shape of thin needles, giving rise to preferred orientation on a flat sample holder. (b) Thermal analysis of FAGPbI4: TG, DSC and corresponding derivatives (DTG, DDSC). The first step of decomposition of FAGPbI4 in the temperature range of 255 - 300 ℃ with a mass loss of about 5.5 % could correspond to the loss of formamidine. A weight percent (w) of formamidine (CN2H4) in FAGPbI4 is 5.4 %. Next steps of decomposition are not clearly defined, as a further loss of organic part [guanidine, products of its fragmentation, w(guanidine, CN3H5) = 7.2 %] and HI, as well as a sublimation of lead iodide take place. to the higher symmetry and rigidity of the former structure. In contrast, δ-FAPbI3 shows a complex spectrum with spinning side bands. The isotropic chemical shift was determined by varying the MAS spinning frequency and, after deconvolution, was found to be - 1175 ppm with a FWHM of 21.8 kHz (Figure S5.6). This spectrum indicates a larger asymmetry of the Pb environment in δ-FAPbI3. None of the other plausible impurities such as precursors (e.g. lead acetate) or side products (PbI2, PbCO3, PbO or Pb(OH)I, etc.) could be detected by sampling a broader range of chemical shifts (Figures S5.7). Understanding the broadband emission from layered LHPs is nontrivial due to effects of the increased structural dynamics in their soft lattices. Temperature-dependent PL measurements were taken to further investigate the complexity of the PL from FAGPbI4 (Figure 5.2f). With cooling, higher-energy emission shifts to a shorter wavelength, up to around 600 nm at 78 K and notably grows in intensity, as one would expect from the excitonic emission.

Guanidinium-formamidinium lead iodide: a layered perovskite-related compound with red luminescence at room temperature | Olga Nazarenko

Table 5.1. Crystallographic data for FAGPbI4.

[C(NH2)3][CH(NH2)2]PbI4 Formula weight 819.95 Temperature (K) 298 Crystal system monoclinic Space group C2/m Color brown a (Å) 26.948(2) b (Å) 12.8189(11) c (Å) 14.4080(12) α (°) 90 β (°) 109.9408(14) γ (°) 90 Volume (Å3) 4678.7(7) Z 12 3 ρcalc (g/cm ) 3.492 μ (mm-1) 18.710 F(000) 4224.0 Crystal size (mm3) 0.32×0.06×0.04 Radiation MoKα (λ = 0.71073) 2Θ range for data collection (°) 3.006 to 52.738 Index ranges -33≤h≤33, -16≤k≤16, -18≤l≤18 Reflections collected 21674

Independent reflections 5009 [Rint = 0.0386, Rsigma = 0.0319] Data/restraints/para-meters 5009/0/144 Goodness-of-fit on F2 1.074

Final R indexes [I>=2σ (I)] R1 = 0.0409, wR2 = 0.0945

Final R indexes [all data] R1 = 0.0552, wR2 = 0.1016 Largest diff. peak/hole (e Å-3) 1.87/-1.27

Guanidinium-formamidinium lead iodide: a layered perovskite-related compound with red luminescence at room temperature | Olga Nazarenko

207 Figure 5.4. Pb SSNMR spectra of FAGPbI4, G2PbI4, α-FAPbI3, and δ-FAPbI3. All spectra were acquired under 20 kHz magic-angle spinning, except for α-FAPbI3, which was recorded in static mode. The longer-wavelength side of the PL envelope is rather minor altered by the temperature. Time-resolved measurements also point to distinctly different nature of lower and higher energy emission bands (Figure 5.7, two spectral regions were probed; these were defined through the use of short- and long-pass optical filters, as indicated by shaded areas). Typically, structurally rigid 2D LHP structures exhibit narrow-band PL with short lifetimes for excitonic transitions, from a few to tens of ns, whereas more complex structures with STEx are

269 Figure 5.5. Crystal structures of FAGPbI4, MAGPbI4 (Ref. ) and CsGPbI4. MAGPbI4, as well as CsGPbI4, belong to Ruddlesden-Popper-related phases. Ruddlesden-Popper phases are composed of (100) perovskite slabs with a general formula of An+1BnX3n+1 and the simplest member (n=1) adopting the K2NiF4 structure. The slabs in Ruddlesden-Popper phases are displaced by a vector (a+b)/2 with respect to each other. In the case of MAGPbI4 as well as CsGPbI4 the displacement is a/2. characterized by broader emission with one-to-two orders of magnitude longer lifetimes.84, 90, 226, 279 Minor changes in an average lifetime for 500-650 nm spectral region as a function of the excitation intensity was observed, as opposed to the region of 700-750 nm, where a much

Guanidinium-formamidinium lead iodide: a layered perovskite-related compound with red luminescence at room temperature | Olga Nazarenko stronger power dependence was observed (Figure 5.7). Thus, for 700-750 nm spectral region the carrier lifetime decreases with increasing of the excitation power. Obtained results underline a different nature of the emission mechanisms of these bands. While still speculative, the PL band at ca. 645 nm can be attributed to (near) band-edge excitons, whereas lower-energy emission (at around 685 nm) might be STEx-related. However, for the lower-energy emission band, the emission can also be related to defects such as color centers.280All PL decay traces are multi-exponential and are compared by their average lifetimes (τavg, extracted from biexponential fits). The lower-energy PL is characterized by long

τavg on the order of hundreds of ns, which increases with increasing the temperature. Considering that the PLQYs are on the order of a few %, it is impossible to unambiguously assign these lifetimes to exclusively radiative processes. Additional measurements, such as spectral dependences of PL excitation and PLQYs, excitation power dependent time-resolved PL (TR-PL), and TR-PL with micrometer spatial resolution (see Figures 5.7, S5.2, S5.8), do not allow for a more accurate assignment of the complex and overlapping emission bands in such 2D LHPs. However, the extended 2D electronic structure of FAGPbI4 allows for the observation of photoconductivity (Figure 5.6d). The dark specific resistivity values were ca. 11010 Ω∙cm, indicating rather low intrinsic carrier concentration and/or mobility. The photoconductivity spectrum of a single FAGPbI4 crystal peaks around 510 nm, essentially

Figure 5.6. (a) RT TR-PL traces recorded in two spectral ranges, as defined by the combination of commercial short- and long-path filters (FEL/FES, from Thorlabs): 500-650 nm (green) and 700-750 nm (red). Inset: temperature-dependent average PL lifetime. (b) Regions of PL spectra selected for time-resolved measurements. (c) Configuration coordinate diagram for free excitons (Ex) and STEx. (d) Photoconductivity spectrum from a single FAGPbI4 crystal.

Guanidinium-formamidinium lead iodide: a layered perovskite-related compound with red luminescence at room temperature | Olga Nazarenko

Figure 5.7. (a, b) Time-resolved PL traces for 500-650 nm and 700-750 nm spectral regions recorded at different excitation intensities on powdered crystals synthesized with method 1. (c) Corresponding average lifetime calculated from TR-PL traces for 500-650 nm and 700-750 nm spectral regions from (a, b).

2 Figure 5.8. (a) Kubelka–Munk function F(R∞) = (1-R∞) /2R∞ of FAGPbI4 on powdered crystals prepared with method 2 and a temperature dependent PL spectra of the same sample (b). (c) A photo of crystals prepared with a higher-yield method 2 under daylight (top) and UV illumination (bottom). The absorption and PL spectra showed different intensities of certain transitions, compared to the spectra obtained on powdered crystals synthesized with method 1. For instance, the Kubelka–Munk function shows a large tail that goes roughly from 2.44 to

Guanidinium-formamidinium lead iodide: a layered perovskite-related compound with red luminescence at room temperature | Olga Nazarenko around 1.7 eV. The transition at 2.11 eV is clearly observable, while for larger crystals (Figure 5.2e) it was less pronounced. It is hard to make a definite conclusion as to the origin of the transition at 2.11 eV. Supposedly it originates from defects from the surface (microcrystalline powder possesses higher surface areas). coinciding with the absorption peak seen in Kubelka-Munk function. The tail of the weak photoconductivity signal extends further to 700 nm, thus covering the range in which PL bands are observed. Again, as in the case of PL, it remains impossible to disentangle two possible origins of this photoconductivity – from trap states or from STEx. Furthermore, the fine spectral details of the PL vary also with the synthesis method (see method 1 and 2 in for more details), such as reaction scale or crystal size, thus with the method 2 higher yields were obtained with smaller crystals’ size (Figure 5.8).

5.4. Conclusions

In conclusion, in this study FAGPbI4 was synthesized; it is thermally stable until 255 °C, exhibits PL at RT and pronounced photoconductivity. This study highlights a plethora of structures, other than traditional ABX3 perovskites, that can form in the compositional space comprising Cs+, FA, MA, G as A-site cations (interlayer cations), Pb(II), Sn(II), Ag(I), Bi(III) as B-site cations (octahedra-forming cations), and various halide anions.

Reproduced with modifications from: O. Nazarenko, M. R. Kotyrba, S. Yakunin, M. Aebli, G. Rainò, B. M. Benin, M. Wörle, M. V. Kovalenko. Guanidinium-formamidinium lead iodide: a layered perovskite-related compound with red luminescence at room temperature. J. Am. Chem. Soc., 2018, 140, 3850- 3853. Copyright 2018, American Chemical Society. https://pubs.acs.org/doi/abs/10.1021/jacs.8b00194

Guanidinium and mixed cesium-guanidinium tin(II) bromides: effects of quantum confinement and out-of-plane octahedral tilting | Olga Nazarenko

Chapter 6. Guanidinium and mixed cesium- guanidinium tin(II) bromides: effects of quantum confinement and out-of-plane octahedral tilting

6.1. Introduction Mixed organic-inorganic ternary and higher tin (II) halide compounds are intensely researched as to their photophysics and electronic properties, which are all governed by the structural diversity of Sn(II)-halide anionic networks.281-286 Such compounds are promising for application in light-emitting devices,286-288 solar cells289-290, and photodetectors.291 The crystal structure of hybrid tin (II) halide compounds consists of Sn(II)-X (X = Cl, Br, I) coordination anionic networks and organic cations. Sn(II)-X units can arrange themselves into a vast variety of structures: from extended 3D structures composed of Sn(II)-X coordination polyhedra connected by corners, edges and faces to isolated polyhedral units (0D-compounds) with a plethora of structures in-between.81, 126, 128, 281, 283, 292-293 The Sn(II)-X units thus far observed experimentally include trigonal pyramids (CN = 3),294-295 seesaws (CN = 4),286 square pyramids (CN = 5),285, 296-297 and, most commonly, octahedral units (CN = 6).281, 298 The coordination geometry of the Sn(II) ion is a complex interplay that involves the effects of the stereoactivity of the 5s-lone pair as well as the geometry of the organic cation and the interactions of the cation with inorganic units such as hydrogen bonding. Satisfactory rationalization (and prediction) of these effects on the atomistic and electronic structure of the resulting compounds remains a formidable challenge. The electronic band structure of hybrid Sn(II) halide compounds at the band gap is formed mainly by Sn 5p and 5s orbitals, and halide np orbitals. Non-conjugated organic amines do not significantly contribute to frontier orbitals.89, 299 In hybrid tin (II) halide perovskite 2 structures, such as (C6H5CH2CH2NH3)2SnI4, the ns pair of Sn defines an antibonding character

Guanidinium and mixed cesium-guanidinium tin(II) bromides: effects of quantum confinement and out-of-plane octahedral tilting | Olga Nazarenko of the valence band maximum by coupling with the p orbitals of halogen atoms.299 The bandgap depends on the dimensionality and connectivity of the Sn(II)-X units, the atomic number of the halogen, and distortions of the metal-halide polyhedra (governed by the organic moiety).281-282, 300 For instance, the optical bandgap of the 3D CH3NH3SnI3 is 1.26 eV, whereas the 2D 301 [C6H5(CH2)2NH3]2SnI4 has an optical absorption edge at 1.97 eV. In the context of real-world applications, low-dimensional hybrid Sn(II) halides might eventually be of use as bright light emitters for lighting and other applications. For instance, 0D

(C4N2H14Br)4SnBr3I3 (organic cation is N,N’-dimethylethylene-1,2-diammonium) demonstrates bright, yellow, broad-band PL with a QY of ∼85 % and a large Stokes shift (187 287 nm, or 1 eV) and FWHM - 126 nm. The corresponding compound, (C4N2H14Br)4SnBr6, containing exclusively bromide, exhibits a QY close to 100 %.288 The structure of these 4- compounds comprises SnX6 octahedra isolated by organic cations. Such emission characteristics are attributed to the emission from STEx, which are characterized by high binding energies and a low probability of interaction with intrinsic defects.302 STEx form due to a combination of electronic localization of an electron-hole pair (exciton) and structural distortion that occurs fast upon photoexcitation (Figure 6.1f), which lead to a Jahn-Teller-like distortion in the excited state.303 Hybrid 0D main-group metal-halides are characterized by a soft lattice, strongly favoring the formation of STEx due to facile structural distortion (exciton- phonon coupling). Radiative relaxation of STEs leads to broad-band emission with a large Stokes shift as reported recently for a wide range of Sn(II) and Sb(III) halides.283, 286-288, 304-305

In comparison, the emission from 3D CH3NH3SnI3 results from free-excitonic states and the PL band is characterized by a small Stokes shift, a narrow FWHM, and a fast radiative lifetime.81, 306

Figure 6.1. (a, b) Crystal structure of guanidinium tin bromide - G2SnBr4 (XI). (c) Kubelka– Munk function of CsSnBr3 and XI. (d) PL and PL excitation spectra measured on powdered crystals of CsSnBr3 and G2SnBr4 (XI). (e) TR-PL traces of CsSnBr3 and XI at RT. (f) A schematic of the formation of STEs.

Guanidinium and mixed cesium-guanidinium tin(II) bromides: effects of quantum confinement and out-of-plane octahedral tilting | Olga Nazarenko

Both organic and inorganic constituents determine the resulting crystal structure. Thus, + organic cations such as CH(NH2) facilitate the formation of the 3D perovskite network of 5a + FASnI3. However, the use of large cations, such as phenethylammonium [C6H5(CH2)2NH3 ], typically results in a layered structure.299 A cation that is only slightly larger than FA – guanidinium – forms a layered G2SnI4 perovskite crystal structure with corrugated Sn-I 253 layers. Guanidinium is a highly symmetric molecule (D3h) and a potential contributor of six hydrogen bonds. The G cation, due to its larger-than-FA ionic radius, destabilizes the 3D 38 perovskite network of corner-shared Sn-X octahedra. Instead, GSnI3 forms a 3D network composed of Sn-I octahedra with mixed connectivity: corner- and face-sharing .281 The appealing attributes of the G-cation are its high thermodynamic stability, strong hydrogen bonding capabilities, and high basicity.248, 250 In a G-Sn-Cl system, there are two known 257, 297 compounds: 1D G2SnCl4 and the 3D perovskite-like GSnCl3. Unlike G2SnI4, where the building blocks are Sn-I octahedra, G2SnCl4 has square pyramidal Sn(II) surrounded by Cl ions. So far, no compounds were reported in the G-Sn-Br system. Herein, the G-Sn-Br and the mixed-cation Cs-G-Sn-X (X – Br, I) systems were investigated, wherein the effect of the additional smaller cation on the obtained structure can be probed. In this study, the synthesis of quaternary mixed Cs-G-Sn(II) iodides was not 253, 281 307 successful, observing instead known ternary G2SnI4 and CsSnI3 phases. Therefore, the focus has been on the bromide systems. G2SnBr4 (XI) was synthesized (Figure 6.1a) and compared it to CsSnBr3, and subsequently explored mixed Cs/G phases. G2SnBr4 is a 1D, luminescent compound with broadband PL and a PLQY of 2 % at RT. It is thermally stable up to 300 ℃. Two layered perovskite compounds CsGSnBr4 (XII) and Cs2GSn2Br7 (XIII) were obtained and characterized in terms of their crystal structure, electronic structure, and optical properties. They reveal significant quantum confinement and that the planar shape of the G cation induces significant structural and electronic anisotropy. These structural effects were then compared to those in homologous lead-based compounds.

6.2. Experimental section

Chemicals, reagents and synthetic procedures. Guanidinium carbonate (G2CO3, 99+ %), hydrobromic acid (HBr, 48 % water solution), Sn powder (99.8 %, ~325 mesh) were purchased from Acros. Cesium bromide (CsBr, 99 %), hydroiodic acid (HI, 57% water solution stabilized with 1.5 % hypophosphorous acid, H3PO2) was obtained from ABCR. Tin(II) bromide (SnBr2, 99.2 %) was obtained from Alfa Aesar. All chemicals were used as received without further purification. All syntheses and further manipulations with the products were performed under inert conditions (argon or nitrogen atmosphere) using Schlenk technique. Temperatures of the syntheses, stated below, correspond to the temperatures of the glycerol bath used as a source of heat. The syntheses were conducted in 10-20 ml Schlenk vessels equipped with a stirring bar.

[C(NH2)3]2SnBr4 (XI) was crystallized from HBr. Briefly, 0.250 g (2.1 mmol) of Sn powder was dissolved in 3 ml of HBr (degassed under stirring in Ar atmosphere for ~20 min beforehand); this mixture was stirred for 10 min at RT and then heated in the glycerol bath at

85-90 °C. When all Sn dissolved, 0.378 g (2.1 mmol) of G2CO3 were slowly added. A strong

Guanidinium and mixed cesium-guanidinium tin(II) bromides: effects of quantum confinement and out-of-plane octahedral tilting | Olga Nazarenko evolution of gas was observed (CO2). The reaction mixture was then stirred for an additional 5 minutes resulting in a clear colorless solution, whereupon the stirring and the heating were discontinued. Colorless transparent needles crystallized upon cooling and were separated by vacuum filtration under Ar flow. The product was dried under vacuum at 65-70 °C for 6 h. Yield - 41 % (with regard to the initial [Sn]).

CsSnBr3 was synthesized from stoichiometric quantities of Sn and CsBr precursors in HBr. Briefly, 0.180 g (1.5 mmol) of Sn were dissolved in 3-ml of HBr as described above. Next, 0.320 g (1.5 mmol) CsBr was added, and a black precipitate formed immediately. A mixture was stirred for an additional 5 min at 105 °C. Further manipulations with the product were identical to those with G2SnBr4. Cs2[C(NH2)3]Sn2Br7 (XIII). 0.107 g (0.9 mmol) of Sn powder was dissolved in 3 ml of HBr, stirred for 10 min at RT and then heated to 85-90 °C. When all Sn dissolved, 0.486 g (2.7 mmol) of G2CO3 and 0.1915 g (0.9 mmol) of CsBr were slowly added. A strong evolution of gas (CO2) was observed and a black precipitate formed. The reaction mixture was stirred and kept at 85-90 °C for about 5 minutes until a slightly yellowish solution resulted. Next, the stirring and heating were discontinued, and the mixture was cooled down under a stream of cold tap-water. Initially, some black precipitate of CsSnBr3 crystallized out (Figure 6.2). In a few minutes, red crystals started appearing and the black powder re-dissolved. The mixture was left for one week for further crystallization. Finally, the red product was separated by vacuum filtration under an inert atmosphere and dried under vacuum at 65-70 °C for 6 h.

Figure 6.2. Identification of the black phase appearing upon cooling of the solution of

Cs2GSn2Br7 (XIII) to RT. The powder XRD pattern of the powdered sample (for the measurement the sample was deposited between the adhesive tape that also contributed to the background) obtained during the synthesis of XIII. The calculated pattern for Cs2GSn2Br7

Guanidinium and mixed cesium-guanidinium tin(II) bromides: effects of quantum confinement and out-of-plane octahedral tilting | Olga Nazarenko

(XIII) is based on the experimental single-crystal XRD data obtained in this work (CCDC code

1854833). The calculated pattern of CsSnBr3 is based on the experimental single-crystal XRD data reported by Donaldson J. D. et. al. [J. Chem. Soc., Dalton Trans., 1975, 0, 1500-1506; cif file was downloaded from ICSD database, ICSD card 4071]. When the synthetic solution of XIII was cooled to RT, a black phase rapidly appeared together with some red crystals. The microcrystalline product at this synthetic stage was quickly separated by vacuum filtration and dried under vacuum with a sand bath at 65 oC to identify the black phase. According to the powder XRD pattern, the product is a mixture of CsSnBr3 and XIII (red crystals). Thus, the black phase that forms upon cooling of the reaction mixture to RT is CsSnBr3.

Cs8Sn6Br13I7 was crystallized from a solution of hot HI and HBr acids under N2. CsBr (0.500 g, 2.350 mmol) was first dissolved in HBr (4.0 mL) and HI (1.0 mL) solution in a 10 mL Schlenk reaction tube and sparged with N2 while stirring at RT for 30 min. SnBr2 (0.111 g, 0.399 mmol) was then added and a black precipitate formed. This mixture was then heated with a glycerol bath until an orange solution resulted (110 °C). The stirring and heating were discontinued, and the solution was cooled to RT. Within an hour orange needles crystallized out. The solution was left overnight for further crystallization. The product was separated by vacuum filtration under N2 and dried overnight under vacuum (70-80 °C). A yield of 50.8 % was estimated relative to initial [Sn]. Characterization. Powder XRD patterns were collected in transmission mode (Debye- Scherrer-Geometry) with a STADI P diffractometer (STOE& Cie GmbH) equipped with a curved Ge (111)-monochromator (CuKα1 = 1.54056Å) and a silicon strip MYTHEN 1K Detector (Fa. DECTRIS) with. For the measurement, a ground powder was placed between an adhesive tape (for CsGSnBr4 and Cs2GSn2Br7) or sealed in a 0.3 mm glass capillary (for G2SnBr4 and CsSnBr3). Single crystal XRD measurements were conducted on Bruker Smart Platform diffractometer equipped with an Apex I CCD detector and molybdenum (MoKα = 0.71073 Å) sealed tube as an X-ray source. Crystals were tip – mounted on a micromount with paraffin oil. The data was processed with the APEX3 software package,153 and the structure solution and refinement were performed with SHELXS154 and SHELXL,155 respectively, which are embedded in Olex2.156 The crystal structures were solved with direct methods, light elements (C, N) were located in the difference Fourier map, and hydrogen atoms were placed at calculated positions. The crystallographic data for the reported tin halide compounds were deposited at the Cambridge Crystallographic Data Centre (CCDC) under the codes 1854819

(G2SnBr4), 1854833 (Cs2GSn2Br7), 1854838 (CsGSnBr4), as well as at the Inorganic Crystal Structure Database (ICSD), i.e., card number 434800 for Cs8Sn6Br13I7. UV-Vis diffuse reflectance spectra of the microcrystalline powders were collected using a Jasco V670 spectrophotometer equipped with deuterium (D2) lamp (190 – 350 nm) for use in UV and halogen lamp (330 – 2700 nm) for use in UV/NIR, and an integrating sphere detector (ILN- 725) with a working wavelength range of 220 – 2200 nm. The diffuse reflectance data were transformed with Kubelka-Munk model into the absorption/scattering ratio spectrum. PL spectra were measured with a CCD fiber spectrometer (LR1, Aseq Instruments) with a 355 nm excitation source (frequency-tripled, picosecond Nd:YAG laser, Duetto from Time-

Guanidinium and mixed cesium-guanidinium tin(II) bromides: effects of quantum confinement and out-of-plane octahedral tilting | Olga Nazarenko

Bandwidth). PL emission from the samples passed through a long-pass optical filter with an edge at 400 nm in order to reject the excitation laser line. PL spectra were corrected to the spectral sensitivity of the detection system. For temperature-dependent PL measurements, a sample of CsSnBr3 was placed atop of a 4-stage Peltier cooling/heating element in an evacuated chamber with a quartz window. The sample temperature was adjusted and stabilized with an accuracy of 0.25 °C by a home-made electronic scheme based on an Arduino microcontroller and thermocouple sensor. The current through the Peltier was reversible, thus the setup provided a wide working temperature range of -40 – 120 °C. This is an open-source project by the authors, deposited and described in details at https://www.researchgate.net/project/High- power-thermoelectric-cooler-TEC-controller-with-4-stage-Peltier-refrigerator-heater. TR-PL measurements were performed using a time-correlated single photon counting (TCSPC) setup, equipped with an SPC-130-EM counting module (Becker & Hickl GmbH) and an IDQ-ID-100- 20-ULN avalanche photodiode (Quantique) for recording the decay traces. The average 2 2 ∑푖=1 휏푖 ∙퐴푖 radiative lifetimes were determined as: 휏푎푣푔 = 2 , where Ai and τi are the corresponding ∑푖=1 휏푖∙퐴푖 amplitudes and exponential decay parameters in the biexponential analysis. For measurements of PL and PLE spectra at low temperature (77 K), the sample was encapsulated in a quartz tube filled with Ar gas and placed in a homemade cryostat. Low temperature (77 K) and RT absolute PLQY were measured with excitation at 340 nm using Quantaurus-QY spectrometer from Hamamatsu. The sample was encapsulated in a quartz tube filled with Ar gas. PL measurements at lower temperatures (in the case of Cs2GSn2Br7), down to 5 K, were conducted in a helium cryostat and PL spectroscopy was performed by exciting the sample with a frequency-doubled Ti:Sapphire mode-locked laser delivering pulses of about 150 fs duration at 400 nm and a repetition rate of 80 MHz. The time-integrated PL was analyzed using a CCD-coupled grating spectrometer, whereas TR-PL traces were recorded with a streak camera. Thermal analysis (TG and DSC) was performed using a Netzsch Simultaneous Thermal Analyzer (STA 449 F5

Jupiter). A powdered sample of G2SnBr4 (21.9 mg) was placed in an alumina crucible (without a lid) and heated under Ar gas flow (50 ml/min) to 850 ℃ (10 °C min-1). SEM was carried out on a Quanta 200F microscope (Thermo Fisher Scientific) operated at an acceleration voltage of 20 kV. EDXS was performed with an Octane SDD detector (EDAX, Ametec) attached to the microscope column. For spectra recording and quantification (ZAF correction), the software Gemini (EDAX) was used. Computational details. First-principles calculations were performed with experimental crystal structures using density functional theory (DFT) as implemented in the SIESTA package.162-163 Calculations have been carried out on experimental structures with the GGA functional in the PBE form.164 Core electrons were described with Troullier-Martins pseudopotentials,165 while valence wavefunctions were developed over double-ζ polarized basis set of finite-range numerical pseudoatomic orbitals.166 Spin-orbit coupling was taken into account through the on-site approximation as proposed by Fernández-Seivane et al..167 In all cases, an energy cutoff of 150 Ry for real-space mesh size was used.

Guanidinium and mixed cesium-guanidinium tin(II) bromides: effects of quantum confinement and out-of-plane octahedral tilting | Olga Nazarenko

Figure 6.3. (a) Electronic band diagram of G2SnBr4 (XI) computed by DFT. No dispersion can be observed except in the direction of the 1D chains [010] (Γ➝Y, R➝U, and Z➝T paths). The space group Pna21 is a polar space group. Therefore, a Rashba splitting of the band edge states is predicted both in the conduction and valence bands away from the Y, R and T critical points of the Brillouin zone. (b) Representation of the electronic densities of the conduction (red) and valence (blue) edge states.

6.3. Results and discussion All compounds in this project were crystallized from hydrobromic acid (HBr, 48 % water solution) under Argon atmosphere using Sn powder, CsBr, and G2CO3 as precursors. G2SnBr4 (XI), obtained from stoichiometric quantities of G and Sn(II), crystallizes as colorless transparent needles. Compound XI was then characterized by single crystal XRD (Figure 6.1a, b, S6.1, Tables S6.1-S6.4) and found to crystallize in the orthorhombic crystal system, space group Pna21. This structure consists of corrugated chains oriented along [010] (i.e., extended 2- in b direction), comprising corner-sharing square pyramids - [SnBr5] . Formation of such pyramids emphasizes stereoactivity of the lone pair of Sn(II). G cations are situated in the space between the chains and connect them with a net of hydrogen-halogen bonds into a 3D supramolecular structure (Table S6.2). Two slightly non-equivalent pyramids can be identified, connected in an alternating manner. In both kinds of pyramids, Sn atoms are displaced out of the basal planes (by 0.23 and 0.24 Å, for Br1Br2Br5Br4 and Br2Br5Br8Br7 planes, respectively, Figure 6.1b). The Sn-Br bond lengths are in the range of 2.724(2)–3.117 Å, which is comparable to the distances found in (NH4)SnBr3×H2O (2.623-3.047 Å), where Sn(II) has the same square-pyramidal coordination.308 The shortest Sn-Br bond is with the Br ion at the apex of the pyramid. G2SnBr4 (XI) has a similar crystal structure as G2PbBr4 containing an 2- 270 alternating orientation of [PbBr5] square pyramids in 1D chains. The 1D character is confirmed by the inspection of the DFT electronic structure (Figure 6.3). Indeed, the band diagram (Figure 6.3a) exhibits an indirect band gap of 3.12 eV and no dispersion of the edge

Guanidinium and mixed cesium-guanidinium tin(II) bromides: effects of quantum confinement and out-of-plane octahedral tilting | Olga Nazarenko state bands is observed except when following a path corresponding to the [010] crystal direction. Furthermore, the

Figure 6.4. Thermal stability analysis of G2SnBr4 (XI): TG, DSC, and corresponding derivatives (DTG, DDSC). computed electronic edge state densities exhibit sizeable overlaps, but only along the b crystallographic direction (Figure 6.3b) thus confirming the 1D character of the density of states close to the bandgap. It is also worth noticing that a Rashba splitting of the band edge states is observed both in the conduction and valence bands away from the Y, R and T critical points of the Brillouin zone as a result of the polar space group Pna21 assigned to XI. Sn-I square pyramidal units were also observed when using larger cations, such as 2+ 296 [C3H7N(C2H4)3NC3H7] . G2SnBr4 (XI) is thermally stable up to at least 300 °C (see TGA/DSC analysis in Figure 6.4). Two endothermic processes without mass loss occur before the decomposition at around 146 and 190 °C and indicate melting and/or other structural transitions.

Studies on a ternary 3D phase – the black CsSnBr3 – date back to as early as the 1970s.

CsSnBr3 crystallizes in a cubic perovskite crystal structure (space group Pm-3m, Figure 6.8a) and melts congruently at 450 °C.309-311 A phase transition to a tetragonally distorted lattice 311 appears upon cooling to below 19 °C. The yellow CsSnI3 crystallizes in the orthorhombic -1 312 crystal system. CsSnI3 is a 1D compound composed of edge-sharing [SnI6] octahedra, and 311 converts into a cubic 3D phase only at 152 °C. Similarly, CsPbI3 is known to crystallize in a 3D-lattice (cubic, space group Pm-3m) at higher temperatures and converts into 1D-phase, 36 analogous to 1D-CsSnI3, within the 329 - 290 °C temperature range.

Guanidinium and mixed cesium-guanidinium tin(II) bromides: effects of quantum confinement and out-of-plane octahedral tilting | Olga Nazarenko

The effect of Br-I mixing in the Cs-Sn-X system was investigated. Interestingly, the addition of iodide with [I]/[Sn]=19/1 (in concentrated HBr) leads to the precipitation of

Cs8Sn6Br13I7 in the form of orange needles. It crystallizes in the orthorhombic crystal system, space group Cmcm (Tables S6.5-S6.7) and consists of layers of the corner (through μ2-Br) and -1⅔ -1⅙ edge-sharing (through μ3-I ions) distorted [SnBr4I2] and [SnBr3I3] octahedra (Figure 6.5a, b). Iodide has one full and two partially occupied positions that are shared with bromide ions. The Cs+ cations are situated in-between, as well as within, Sn(II) halide layers. There are two types of the crystallographic surrounding of Sn(II) ions; one characterized by the presense of a -1⅙ very weak Sn-I bond of 3.723 Å – in [SnBr3I3] octahedra (see Sn1 atom in Figure 6.5b). The -1⅔ Sn2-I1 distances within [SnBr4I2] are 3.454 Å, which is in the range of d(Sn-I) in CsSnI3, 312 2.942 – 3.470 Å. Based on powder XRD analysis, Cs8Sn6Br13I7 was obtained as a pure phase (Figure 6.5c) with a homogeneous composition that was estimated with EDXS (Figure 6.6a-c). 313 The same structural motive was found in Cs2.38Rb1.62Sn3Cl8I2, recently reported by Li et al., obtained by solid-state reaction. The optical band edge of Cs8Sn6Br13I7 lies at 2.23 eV (Figure 313 6.7a, b), which is 0.58 eV smaller than the optical band gap of Cs2.38Rb1.62Sn3Cl8I2.

Figure 6.5. (a) Crystal structure of the 2D Cs8Sn6Br13I7 (ICSD card number 434800). (b) Coordination of Sn(II) ions with weak Sn(2)-I bonds (shown in dotted lines). (c) The powder XRD pattern of a powdered sample of Cs8Sn6Br13I7. The calculated pattern is based on the experimental single-crystal XRD data obtained in this work. A preferred orientation of the microcrystalline powder is observed through the difference in intensity of certain reflections in the experimental pattern comparing to the calculated data because Cs8Sn6Br13I7 crystallizes in the shape of thin needles, giving rise to a preferred orientation in a sample holder: powdered crystals were deposited between the adhesive tape.

Guanidinium and mixed cesium-guanidinium tin(II) bromides: effects of quantum confinement and out-of-plane octahedral tilting | Olga Nazarenko

Figure 6.6. (a) SEM image of Cs8Sn6Br13I7 crystals. (b) Results of EDXS analysis on single crystals of Cs8Sn6Br13I7 (wt % - weight percent). (c) SEM-EDXS spectrum.

2 Figure 6.7. (a) Kubelka–Munk function F(R∞)=(1-R∞) /2R∞ of Cs8Sn6Br13I7. (b) Estimation of 2 the bandgap of Cs8Sn6Br13I7 from an [F(R∞)·hν)] (where hν is the incident photon energy) versus energy (hν) graph.

Guanidinium and mixed cesium-guanidinium tin(II) bromides: effects of quantum confinement and out-of-plane octahedral tilting | Olga Nazarenko

Figure 6.8. (a) Crystal structure of cubic 3D-perovskite CsSnBr3. (b, c) Crystal structure of 2D perovskites: CsGSnBr4 (XII) and Cs2GSn2Br7 (XIII). (d) Kubelka–Munk function of CsSnBr3 and XIII at RT. (e) Selected low-temperature PL spectra of XIII. (f) Selected low-temperature time-resolved PL spectra of XIII.

By comparing the experimental powder XRD patterns with those simulated based on single-crystal XRD data obtained herein (except for CsSnBr3, where the ICSD card 4071 was used), XI and CsSnBr3 were determined to be phase-pure (Figure S6.2, S6.3). CsSnBr3 also appears to contain amorphous CsSnBr3 (Figure S6.3), which can explain the partially-resolved component at 650-700 nm in the PL spectrum (Figure 6.1d). Next two mixed guanidinium- cesium tin bromides were obtained: CsGSnBr4 (XII, orange-colored) and Cs2GSn2Br7 (XIII, red-colored), as a mixture of crystals. Both compounds are layered perovskites (Figure 6.8b, c). The formation of these phases was found to be in strong competition with the crystallization of compounds XI and CsSnBr3. Cs2GSn2Br7 (XIII) was obtained as the dominant phase, with CsGSnBr4 (XII) as a major impurity (below 10%, Figure S6.4) was obtained at a molar ratio of Cs/G/Sn = 1/6/1. Optical absorption analysis by a Kubelka-Munk function points to a possible trace quantity of rather amorphous CsSnBr3 thus giving rise to a small shoulder in the 650-700 nm region (Figure 6.8d). Both XII and XIII crystallize in the orthorhombic crystal system, namely in Imma (XII) and Cmmm (XIII) space groups, respectively (Figure 6.8b, c, Table S6.1). The crystal structures of XII and XIII consist of perovskite layers of corner-sharing 4- [SnBr6] octahedra that are separated by cesium and guanidinium cations. The derivation of these structures can be visualized by cutting the parental cubic CsSnBr3 perovskite lattice along the (100) crystallographic plane and adding the octahedral tilting (discussed later in detail). The importance of mixed-cation design for obtaining 2D-perovskites had been shown in Chapter 4

Guanidinium and mixed cesium-guanidinium tin(II) bromides: effects of quantum confinement and out-of-plane octahedral tilting | Olga Nazarenko for the analogous CsGPbBr(I)4 and Cs2GPb2Br7 compounds. Compound XII contains monolayers of corner-sharing octahedra (Figure 6.8b), whereas compound XIII features double layers, connected via corners (Figure 6.8c). To the best of our knowledge, this is the first tin bromide featuring multi-layers. The perovskite slabs are mutually shifted in-plane by a vector a/2 (a being an ideal cubic cell vector, Figure 6.8c right). Cs+ and G cations in XII and XIII are both situated in the interlayer space, and alternate with each other in a periodic manner. In

Cs2GSn2Br7 (XIII), cesium cations also fill the voids within the perovskite layers. G forms hydrogen bonds with bromide ions of the perovskite layers (Table S6.8, S6.9). G cations are closely packed with intermolecular d(C···N) distance of ca. 3 Å. Due to the difference in ionic radii of the Cs+ and G cations and the planar shape of the latter, Sn-Br octahedra are tilted in an alternating manner along the crystallographic axis b and c for XII (Figure 6.8b) and XIII (Figure 6.8c left), respectively. This generates smaller cavities for Cs+ and large voids for G cations. However, in the other in-plane direction (for instance along the crystallographic axis a for XIII, Figure 6.8c right) the planar shape of the G cation and small atomic radius of Cs+ lead to the absence of steric strain with almost perfectly collinear Sn and Br atoms and subsequently longer Sn-Sn distances (6.011 and 5.965 Å for XII and XIII, respectively). This might contribute to an observed disorder in atomic positions of bromide along this direction. In

CsGSnBr4 (XII), d(Sn-Br) = 2.909(2)-3.006(1) Å, are comparable to Sn-Br distances in a 309 CsSnBr3 cubic phase [d(Sn-Br) = 2.898 Å]. Sn-Br-Sn angles are in the range of 155.9(5)-

180.0 ° (Figure 6.9). CsGSnBr4 and CsGPbBr4 both crystallize in the space group Imma, and 270 270 are isotypic. XIII is isotypic to Cs2GPb2Br7, and crystallizes in the space group Cmmm. The average Sn-Br distances for XIII are in the range of 2.915-2.987(1) Å (for additional crystallographic data for XII and XIII please see Tables S6.10-S6.13). The interlayer Br···Br distances along the stacking direction (c and b for XII and XIII, respectively) are rather short: d(Br···Br) = 4.161 Å in CsGPbBr4 (XII) and d(Br···Br) = 4.070 Å in Cs2GSn2Br7 (XIII). In 3D perovskites, the critical point of the Brillouin zone where a direct bandgap is observed is either the R-point,314for the cubic Pm-3m phase, or the Z-point for the distorted tetragonal P4/mbm phase (Figure S6.5). In both cases, the critical wavevectors have the point symmetries of their respective lattices (m-3m and 4/mmm). Similar symmetry considerations lead to critical wavevectors located along the stacking directions at the X (0,0,1) and Y (0,1,0) high symmetry points for XII and XIII, respectively, in the Imma and Cmmm Brillouin zones (both with point groups mmm). However, DFT simulations (Figure 6.10a and c) show that the electronic band gaps are predicted to be direct at the R (0.5,0,0.5) and S (0.5,0.5,0) points of the Brillouin zones of XII and XIII, respectively. In both cases, the point symmetry of the critical wavevector is reduced from mmm to 2/m. This can be related to the anisotropy of the out-of-plane octahedral tilts, which is directly induced by the shape of the G cations. The octahedral tilting is clearly apparent only in the (b,c) crystallographic planes (Figure 6.8b and c). The difference between XII and XIII lies in the fact that the (b,c) plane contains the stacking axis along c for XII and along b for XIII. Furthermore, due to the lattice distortions, the periodicity of the electronic density is doubled along b (c) for XII (XIII), in contrast to the a direction. In turn, the valence band maxima and the conduction band minima show different hybridizations along all three directions (Figure 6.10b, d, and Figure S6.6) as well as different

Guanidinium and mixed cesium-guanidinium tin(II) bromides: effects of quantum confinement and out-of-plane octahedral tilting | Olga Nazarenko

Figure 6.9. Octahedral coordination of Sn(II) ions (ellipsoids shown with 50 % probability) and characteristic Sn-Br-Sn angles in XII (a) and XIII (b). band dispersions (Figure 6.10a and c). Both compounds XII and XIII exhibit flat dispersions along the stacking axis despite the short Br···Br distances, a classic feature of 2D systems. Analogously to CsSnBr3 (Figure 6.8) and contrarily to XI (Figure 6.3), the band gap remains direct in both XII and XIII (Figure 6.10a and c) with a systematic decrease of the bandgap values with the decrease of dimensionality and thickness of perovskite layers. These variations in DFT band gap energies (Table S6.14) are related to the equivalent quantum well thickness

(n = 1, 2 or ), which is roughly doubled when going from CsGSnBr4 (XII) to Cs2GSn2Br7 (XIII) as a result of quantum confinement. To experimentally probe the energy bandgaps in these hybrid tin bromide compounds, an optical absorption analysis by means of a Kubelka-Munk function as well as PL measurements were performed (Figure 6.1c, d, Figure 6.7, Figure 6.8d,e,). The Kubelka-Munk absorption representation was derived from the diffuse reflectance spectra of powdered 2 samples, and was estimated as F(R∞) = α/S = (1−R∞) /2R. The influence of excitonic effects on the absorption was neglected, because the excitonic band cannot be consistently resolved for all compounds and thus correctly de-convoluted from the bandgap absorption edge. Hence the bandgap can be slightly underestimated. The optical bandgap energies for XI, XIII and r CsSnBr3 were determined by plotting [F(R∞)·hν] versus (hν) (hν - the incident photon energy, r = 2 for direct bandgap semiconductor, and r = 1/2 for indirect bandgap semiconductor; Figure

6.11a,b). The highest absorption edge energy (ca. 3.11 eV) was found for G2SnBr4 (XI), which is in agreement with its lowest electronic dimensionality (1D). The 3D cubic CsSnBr3 has the

Guanidinium and mixed cesium-guanidinium tin(II) bromides: effects of quantum confinement and out-of-plane octahedral tilting | Olga Nazarenko

Figure 6.10. (a) Electronic band structure and (b) the representation of the electronic densities of the conduction (red) and valence (blue) edge states (at the R point of the Imma Brillouin zone) of CsGSnBr4 (XII). R➝E0 corresponds to the stacking direction. (b) and (d), same for Cs2GSn2Br7 (XIII). S➝J2 corresponds to the stacking direction. (e) Electronic band structures of CsGPbBr4 and (f) Cs2GPb2Br7. Electron densities are computed at the R point for XII (b) and S point for XIII (d).

2 Figure 6.11. (a) Tauc plots: [F(R∞)·hν)] (where hν is the incident photon energy) versus 1/2 energy (hν) for Cs2GSn2Br7 (XIII) and CsSnBr3. (b) Tauc plots: [F(R∞)·hν)] versus energy (hν) for indirect bandgap G2SnBr4(XI). smallest band gap energy (1.74 eV, in agreement with the literature),315 whereas 2D Cs2GSn2Br7 (XIII) has an intermediate band gap value of 2.16 eV (Figure 6.11a). Compound XII could not

Guanidinium and mixed cesium-guanidinium tin(II) bromides: effects of quantum confinement and out-of-plane octahedral tilting | Olga Nazarenko be obtained in the phase-pure form needed for diffuse reflectance (only individual small crystallites could be selected for a single-crystal XRD analysis). However, the visual appearance of XII pointed towards an expected higher bandgap in comparison to XIII (orange vs. red). Despite the well-known underestimation of band gaps, such a decrease is predicted from DFT calculations (Figure 6.10 and Table S6.14) with 1.04 eV for XII and 0.84 eV XIII.

Cs2GPb2Br7 has a larger optical experimental band gap than XIII, by 0.44 eV (2.60 eV 270 vs. 2.16 eV). A similar scenario was observed for the pair of (C6H9C2H4NH3)2SnBr4 (2.5 eV optical band gap) and its lead analog (2.9 eV).316 This might be a result of the difference in the

E3 ionization energies between tin and lead [E3(Pb) = 31.94 eV and E3(Sn) = 30.5 eV], the effect of the additional polarizability (presence of filled diffuse d and f orbitals below the valence electrons in Pb), and the distortions of the M-X (M = Sn, Pb) octahedra. A different degree of tilting of the MBr6 octahedra is observed for Cs2GPb2Br7 and XIII (Figure 6.12), with stronger deviations from the ideal 180° for the lead compound. Interestingly, despite the different tilts of the octahedra, the angles are at the end compensated in a manner as to give the required amount of space for guanidinium and cesium cations. Figure 6.10e and f report the computed band structures for the lead-based analogs of compound XII and XIII that share many similarities with the tin-based layered perovskites (Figure 6.10a and c). Both for n=1 (XII) and n = 2 (XIII) the predicted electronic bandgaps follow the same trends as the above mentioned experimental results. For instance, the bandgap increases from 1.04 eV (Sn) to 1.55 eV (Pb) for XII and from 0.84 eV (Sn) to 1.32 eV (Pb) for XIII.

Both CsSnBr3 and XI are luminescent at RT, unlike XII and XIII. The PL of G2SnBr4 (XI) at RT is characterized by a broad emission band (Figure 6.1d) with an FWHM of ca. 121 nm and 2 % PLQY and with a large Stokes shift of ca. 208 nm. At 77 K, PLQY reaches 75±5 % and the PL band narrows (FWHM ~72 nm) while retaining the same peak position at

Figure 6.12. (a, b) Octahedral coordination of Sn(II) and Pb(II) ions in Cs2GPb2Br7 and

Cs2GSn2Br7 (XIII) (ellipsoids shown with 50 % probability) and characteristic geometric parameters. The 2D metal-halide layers of both compounds are extending in (ca) plane. The crystallographic data (cif-file) for Cs2GPb2Br7 was downloaded from Nazarenko O. et. al., Inorg. Chem. 2017, 56, 11552-11564, and corresponding CCDC card 1552602.

Guanidinium and mixed cesium-guanidinium tin(II) bromides: effects of quantum confinement and out-of-plane octahedral tilting | Olga Nazarenko

Figure 6.13. Temperature dependent PL measurements of CsSnBr3. 557 nm with a Stokes shift of 233 nm. Additionally, the PL excitation spectrum at 77 K is blue- shifted comparing to that at RT. CsSnBr3 exhibits a PL maximum at 1.70 eV (Figure 6.1d) at 315 RT, which is consistent with the literature value. In PL spectrum of CsSnBr3, a shoulder was observed at the higher energy side, which seems to originate from an amorphous or 317 nanocrystalline CsSnBr3 impurity. A blue shift of the emission bands of CsSnBr3 was observed when the temperature increased from 228 to 383 K (Figure 6.13). This trend is consistent with previous temperature-dependent PL theoretical and experimental studies and is explained by the lattice expansion upon heating and local displacements of Sn(II) cation.315, 318- 319 Large differences in the width of the PL bands and Stokes shifts for G2SnBr4 (XI) and CsSnBr3 indicate different radiative processes in these compounds. TR-PL measurements (Figure 6.1e) yielded PL lifetime of 0.4 ns for CsSnBr3 and 18 ns in compound XI. The short lifetime, the small Stokes shift and the narrow PL band attest to the excitonic nature of the radiative transition in CsSnBr3. On the contrary, the longer PL lifetime and larger Stokes shift and broader FWHM point to emission via STEs in compound XI (Figure 6.1f). Importantly, the

PL lifetime in 1D G2SnBr4 is much faster (18 ns) than that commonly found in 0D Sn(II)- halides (102-103 ns for bromide) at RT.286-288, 304 This might eventually lead to a practical advantage, for instance, in the context of applications such as white-light sources,287-288, 303 as the achievable saturated emission brightness scales with the PL lifetime. For XII and XIII, which are structurally more rigid and exhibit a delocalized electronic structure that reduces the propensity to form STEs, the PL was measurable only at cryogenic temperatures (5 – 120 K, Figure 6.8e). At 5 K, the main band is peaked at 517 nm with an additional emission around 577 nm. A broad emission band above 700 nm is also visible. Due to the very low PLQY and presence of different phases within the same sample, it is difficult

Guanidinium and mixed cesium-guanidinium tin(II) bromides: effects of quantum confinement and out-of-plane octahedral tilting | Olga Nazarenko to draw solid conclusions. Yet one can attribute the main emission bands to free excitons or STExs with low stabilization energy (due to hindered structural distortion as compared to 0D and 1D-compounds). Finally, the lowest energy band might stem from surface defects. By increasing the temperature, the PL intensity quickly drops due to the activation of non-radiative processes. TR-PL measurements were performed with a streak camera system, and the results are displayed in Figure 6.8f. The measured sub-100-ps decay times and their dependence on temperature attest to the presence of non-radiative processes, which efficiently quench radiative recombination.

6.4. Conclusions

In conclusion, in this Chapter a luminescent 1D-compound G2SnBr4, as well as layered 2D-perovskites CsGSnBr4 and Cs2GSn2Br7 were presented. In G2SnBr4, the Sn(II) lone pair is stereoactive and influences the coordination environment of Sn(II). G2SnBr4 is a luminescent compound with broadband emission from self-trapped excitonic states that are enabled by the soft-lattice and electronic localization. This emission is characterized by PLQY of 2 % at RT, and 75±5 % at 77 K. G2SnBr4 does not thermally decompose until 300 ℃. The addition of cesium to the G-Sn-Br system promotes the formation of structurally more rigid, layered perovskite compounds, obscuring the formation of STEs but leading to a delocalized electronic structure. Cs2GSn2Br7 has a smaller bandgap than its lead analog, Cs2GPb2Br7, due to the difference in E3 ionization energies of tin and lead, higher electronegativity of Pb, as well as the different degree of structural distortion. Lastly, in-plane electronic coupling is anisotropic as a result of the planar shape of the guanidinium cation, which leads to significant octahedral out-of-plane tilting in only one of the two directions.

Reproduced with modifications from: O. Nazarenko, M. R. Kotyrba, S. Yakunin, M. Wörle, B. M. Benin, G. Rainò, F. Krumeich, M. Kepenekian, J. Even, C. Katan, M. V. Kovalenko. Guanidinium and mixed cesium-guanidinium tin(II) bromides: effects of quantum confinement and out-of-plane octahedral tilting. Chem. Mater. 2019. Copyright 2019, American Chemical Society. https://pubs.acs.org/doi/abs/10.1021/acs.chemmater.9b00038

Rubidium copper and silver iodobismuthates | Olga Nazarenko

Chapter 7. Rubidium copper and silver iodobismuthates

7.1. Introduction Halogenidobismuthates attract an increasing scientific interest owing to their structural diversity320-323 and interesting physical properties such as semiconductivity,324-327 PL,326, 328 thermochromism,329-331 photochromism.332-335 These compounds are studied for potential applications in a field of hard radiation detection,336 photodetection,337 and in solar cells.123, 140, 338 The boost in research on Bi(III) halide coordination compounds is a result of the increased interest to metal-halide complexes in general, motivated by a recent success of hybrid organic- inorganic Pb(II) and Sn(II) halides in a field of photovoltaics - a third generation solar cells, based on thin-film technology with LHPs as photon to energy converter reached efficiencies above 20 %.339 On the other hand, issues with the toxicity of lead and ease of oxidation of Sn2+ - that is considered a replacement of toxic lead and can form analogous mixed organic-inorganic halide perovskite compounds - to Sn4+, prompt an exploration of new semiconductors. A typical polyhedron of Bi3+ in halogenidobismuthate anionic networks is octahedron 3- [BiX6] (X – Cl, Br, I). These octahedral units can condense via corners, edges or octahedral faces into extended polymeric networks or isolated anionic species (bi- to octanuclear units), as well as form heterometallic complexes.320, 322-323 Comparing to Pb(II) halides, Bi(III) halide coordination compounds have somewhat a higher tendency to form lower dimensionality networks with edge and face sharing BiX6 octahedra. 3D networks were found among heterometallic structures, for instance, cesium (or MA) silver bismuth bromobismuthates.134, 326 Silver bromo- and chlorobismuthates with small Cs+ and MA cations form cubic 3D perovskite + 134, 326, 340 structures Cs2AgBiCl6, and A2AgBiBr3 (A = Cs or MA). When a larger cation than Cs+ or MA is used in A-Bi-Ag-X system, a layered structure might form. For instance, the crystal structure of [(C2H5)4N]2[Ag2Bi2I10] is composed of layers made of Bi-I octahedra linked by edge-sharing with Ag-I tetrahedral units.325

Rubidium copper and silver iodobismuthates | Olga Nazarenko

A number of known heterometallic bismuth halide complexes is rather small. They include such metals as Ag(I), Cu(I), and Hg(II) (Table 7.1) and have different dimensionality structures (from 0 to 3D). In the case of mercury iodobismuthates, use of tetraethylammonium cation, or n-butylammonium cation leads to 0D quaternary coordination compounds.341 Cu+ 2+ and Hg typically form halide coordination polyhedra with CN = 4 - tetrahedral MX4 units, while Ag+ can be in tetrahedral as well as the octahedral surrounding of halide ions. So far, no quaternary heterometallic bismuth iodide structure with a 3D crystal structure have been reported.

Table 7.1. Heterometallic bismuth halide complexes

Compound Dimensionality 342 [Bi2(C4H8O3H)3(C4H8O3H2)][CuBi5I19] 1D 325 [Et4N]2[Ag2Bi4I16] 1D 325 [Et4N]2[Ag2Bi2I10] 2D + + 326, 340 A(Ag0.5Bi0.5)Br3, (A=Cs , CH3NH3 ) 3D 134 Cs(Ag0.5Bi0.5)Cl3 3D 343 [PPh4]4[M2Bi2I12] (M=Cu(I), Ag(I)) 0D 324 [n-Bu4N]2[Cu2(CH3CN)2Bi2I10] 0D 324 [Et4N]2[Cu2Bi2I10] 1D 324 [Cu(CH3CN)4]2[Cu2Bi2I10] 1D 341 (Et4N)4(Bi4Hg2I20) 0D 341 (n-Bu4N)2(Bi2HgI10) 0D 344 [Na4((CH3)2CO)15][PtBi2I12]3(CH3)2CO 0D

In this part of the Ph.D. project, a study on fully inorganic quaternary bismuth iodides: rubidium copper and silver iodobismuthates, was performed. Rb+ facilitates the formation of a 2D network in the Bi-I system.147 Therefore, the experiment to connect Bi-I anionic layers into a 3D network with metal-halide bridges was undertaken. The synthesis was performed in the hydriodic acid water solution. As a result, two compounds, XIV in the system Rb-Cu-Bi-I-O- H, and XV in the system Rb-Ag-Bi-I-O-H were obtained. These compounds have complex crystal structures that were studied with single crystal XRD, EDXS, and WD-XRF. The purity of the obtained compounds was estimated with powder XRD.

7.2. Experimental section Chemicals, reagents and synthesis procedures. Hydriodic acid (HI, 57%, water solution, stabilized with 1.5% hypophosphorous acid, H3PO2) was purchased from ABCR.

Rubidium carbonate (Rb2CO3, 99.8 %) was purchased from Acros and silver (I) iodide (AgI,

>99 %) - from Fluka. Bismuth (III) carbonate basic, (BiO)2CO3, was obtained from Siegfried Handel. Copper (I) iodide (CuI, 99.999 %) was bought from Sigma Aldrich.

Synthesis of XIV (the method I). 0.277 g of Rb2CO3 (1.2 mmol), 0.229 g of CuI (1.2 mmol) and 0.153 g of (BiO)2CO3 (0.3 mmol) in a 10 ml round bottom flask equipped with a

100

Rubidium copper and silver iodobismuthates | Olga Nazarenko stirring bar and a stopcock were cooled on ice, and 1.5 ml of HI were added slowly. A strong gas (CO2) evolution was observed, a dark red solution formed and a dark precipitate appeared. The mixture was next heated in the glycerol bath (~80 ℃) until a dark red solution result. The solution is next placed on an ice bath and left for about 15 min. During cooling, dark green prismatic crystals appeared. Next, the product was separated by vacuum filtration and rinsed with cold HI. Further, to remove the rests of HI the crystals were washed with 1-propanol. Finally, the powders were dried under ambient conditions — the mass of the product was 0.55 g.

Synthesis of XV (the method I’). 0.29 g of Rb2CO3 (1.26 mmol), 0.293 g of AgI (1.25 mmol) and 0.32g of (BiO)2CO3 (0.63 mmol) in a 10 ml round bottom flask equipped with a stirring bar and a stopcock were cooled on ice, and 1.75 ml of HI were added slowly. A strong gas (CO2) evolution was observed, a reddish solution formed and a red precipitate appeared. Next, the procedure was the same as described above. 0.46 g of red prismatic crystals were obtained. The growth of larger crystals XIV (method II). 0.165 g of Rb2CO3 (0.7 mmol), 0.136 g of CuI (0.7 mmol) and 0.182 g of (BiO)2CO3 (0.36 mmol) placed in an 8 ml vial, cooled on ice, and 2 ml of HI were added slowly. The formed dark precipitate was re-dissolved upon heating and a dark red solution result. Next, the heating and the stirring were discontinued, and the solution was cooled to RT and left undisturbed. In 3-4 days few crystals of 1-4 mm size grew. Crystals were further removed from the mother liquid and dried under ambient conditions.

XV (method II’). 0.174 g of Rb2CO3 (0.75 mmol), 0.175 g of AgI (0.75 mmol) and 0.192 g of (BiO)2CO3 (0.38 mmol) in an 8 ml vial were cooled on ice, and 2 ml of HI were added slowly. Further, the procedure was identical to the described above.

Figure 7.1. Photos of the crystals of XIV (a) and XV (b) grown with method 2. Powder XRD patterns were collected in transmission (Debye-Scherrer-Geometry) with STADI P diffractometer (STOE& Cie GmbH), equipped with a silicon strip MYTHEN 1K

Detector (Fa. DECTRIS) with a curved Ge (111)-Monochromator (CuKα1=1.54056Å). For the measurements, a ground powder was filled into 0.5 mm glass capillary. Single crystal XRD measurements were conducted on Bruker Smart Platform diffractometer equipped with an Apex

I CCD detector and molybdenum (MoKα=0.71073 Å) sealed tube as an X-ray source. Crystals

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Rubidium copper and silver iodobismuthates | Olga Nazarenko were a tip – mounted on a micro-mount with paraffin oil. The data were processed with CrysAlisPro, structure solution and refinement was performed with SHELXS and SHELXL respectively, embedded in the Olex2 package. UV-Vis diffuse reflectance spectra of the microcrystalline powders were collected using a Jasco V670 spectrophotometer equipped with such light sources, as deuterium (D2) lamp (190 – 350 nm) for use in UV and halogen lamp (330 – 2700 nm) for use in UV/NIR. An integrating sphere detector (ILN-725) designed to measure the transmittance (diffused transmittance) and reflectance (diffuse reflectance) with a working wavelength range of 220 – 2200 nm was used. A white standard is barium sulfate. The diffuse reflectance data were transformed into the Kubelka-Munk function: 2 F(R∞)=α/S=(1−R∞) /2R∞, where F(R∞) - Kubelka-Munk units, α - absorption coefficient, S - the scattering coefficient, and R∞ - the reflectance of an infinitely thick layer. Thermal analysis (TG and DSC) was performed using a Netzsch Simultaneous Thermal Analyzer (STA 449 F5

Jupiter). ~25 mg of sample was deposited into Al2O3 crucible and heated under Argon flow (40 ml/min) to 850 °C with a speed of 10 °C min-1. SEM was done on a Quanta 200F microscope (Thermo Fisher Scientific) operated at an acceleration voltage Vacc = 20 kV. EDXS was performed with an Octane SDD detector (EDAX (Ametec)) attached to the microscope column. For spectra recording and quantification (ZAF correction), the software Gemini (EDAX) was used. WD-XRF: the pellets were examined with Rigaku Primus IV WD-XRF instrument with a standardless calculation of the element mass content. The measuring time was 10 min, with an energy of a 4 kW X-ray tube and an adjusted aperture of 1 mm diameter was used. A relative error of 7-10 % must be expected for standardless semiquantitative WD-XRF analyzes. For resistivity measurement, the sample was contacted by a soft conductive rubber from opposite sides, and the current-voltage (IV) characteristics were measured using a Keithley 236 SMU. 1H SSNMR. All experiments were performed on a Bruker spectrometer (16.4 T) equipped with a 2.5 mm double-channel solid-state probe head and an Avance III console. A 2.5 mm zirconia rotor was used. Chemical shifts were referenced to tetramethylsilane (TMS). A one-pulse experiment was conducted. The 90-degree pulse had a length of 2.6 μs. The recycle delay was set to 4 seconds. For every spectrum, four scans were recorded. The spinning frequency was set to 11.8 kHz. 87Rb SSNMR. Experiments were performed on a 16.4 T Bruker spectrometer equipped with a 4 mm double-channel solid-state probe head and an Avance III console. A 4 mm zirconia rotor was used. Spectra were measured under static conditions and spinning at 5 kHz at the magic angle.

Chemical shifts were referenced to 0.01 M RbCl in D2O. One-pulse and Hahn echo pulse- sequences were used. 90° pulses lasted 8.25 µs and 180° pulses accordingly 16.5 µs. Echo delays varied from 0.087 – 0.788 ms. 256-5120 transients were acquired and recycle delays of 0.5 s were applied. Continuous wave EPR measurements were performed at X-band frequency (~9.7 GHz) using a commercial X-band Elexsys E500 spectrometer (Bruker) equipped with a super high Q resonator (Bruker). Spectra were acquired at 10 K using 2 mW incident microwave power and 0.3 mT magnetic field modulation amplitude. Single crystal samples were placed into 3 mm quartz EPR tubes. Per sample, several random orientations with respect to the external magnetic field were tested. Baseline measurements were performed using the same settings on the same empty 3 mm tube.

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7.3. Results and discussion The compounds were crystallized from hydroiodic acid water solution. XIV crystallizes as dark green colored prisms and XV as red-colored irregular polyhedrons. XIV crystallizes in the tetragonal crystal system, space group P4mm, whereas XV is isotypic, but crystallizes in the lower symmetry orthorhombic crystal system, space group Cmmm (Tables S7.1-S7.5). Assignment of Bi, Rb, and part of I atoms in the structures of XIV and XV, during the crystal structure determination, was straightforward, whereas the electron density within a certain region was challenging to assign due to a strong static disorder. This region is best described with a body-centered cube capped with I- ions. According to geometrical parameters, such as distances to the I- ions and the geometry of the surrounding, the corners of the cube are best described as Ag and Cu atoms with partial occupancies (Table 7.2). The electron density in the center of the cube was assigned to Cu [0.36(3) occupancy] and Ag [0.68(2) occupancy] for XIV and XV respectively. The rest of the disordered electron density in and around the M-I (M = Cu, Ag) cube was described as I- ions with partial occupancies. Thus, according to the single crystal XRD, the empirical formula of XIV and XV are Rb20Cu8.32Bi6I48.14O8 and

Rb20Ag8.419Bi6I47.482O8 correspondingly, where O originates from water molecules. The amounts of Cu and Ag, and I in both compounds are comparable. However, the electroneutrality + condition is not maintained. The electroneutrality might be reached if hydronium H3O ions are present in the crystal structure (compounds were crystallized from an acid - water solution). The problem with electroneutrality might also point towards a different oxidation state of Cu and Ag than 1+.

Table 7.2. Atomic occupancy for XIV.

Atom Occupancy Atom Occupancy Atom Occupancy I8 0.580(16) I4A 0.516(18) I5B 0.42(3) Rb1 0.5 Cu3 0.32(2) I4B 0.484(18) I5A 0.58(3) I9 0.157(12) Cu1 0.36(3) I11 0.34(2) Cu2 0.156(17) Cu4 0.47(2)

Crystal structure of XIV consists of a 3D network made of BiI6 octahedra linked by Cu- I complex polyhedral units (Figure 7.2). This 3D network has channels large enough to host 3- + BiI6 octahedra (shown on Figure 7.2a as polyhedra), Rb cations, and water (and possibly 3- hydronium) molecules. BiI6 octahedra within the channels have an ideal geometry with d(Bi- I) = 3.022 – 3.073 Å and I-Bi-I angles in the range of 89.69(1) – 90.31 º. Whereas BiI6 octahedra - within the network have a strong disorder of iodide ions. Thus, four of six I ions in BiI6 octahedra within the 3D anionic network are disordered on two positions. Cu-I unit, presented in Figure 7.2b, is made of a Cu [0.47(2) occupancy] cube with 4Cu ions disordered on two positions. All faces of the cube are capped with I- ions. In the center of a copper cube, an electron density was described as Cu1 with 0.36(3) occupancy. The other electron density within the cube and between two neighboring cubes was described as I11 [0.34(2)] and I9 [0.157(12)] ions. As to the copper ions of the cube, Cu3 has a tetrahedral surrounding of I- ions

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Rubidium copper and silver iodobismuthates | Olga Nazarenko with d(Cu3-I) = 2.628 – 2.767 Å, which is close to the range of Cu-I bonds in RbCu2I3, d(Cu- I) = 2.592 – 2.716 Å.345 Except, in XIV, due to a disorder of iodide ions, one of the bonds is longer - d(Cu3-I4B) = 2.767 Å than the maximum Cu-I distance in RbCu2I3. Cu2, on the other hand, is surrounded by 3I- in a trigonal-planar geometry with d(Cu2-I) = 2.478(1) – 2.556 Å, similar to Cu-I distances in Cs3Cu2I5, d(Cu-I) = 2.500(20) – 2.567 (7) Å for Cu(I) with the same trigonal-planar coordination surrounding.346 The closest I- neighbors form an octahedral surrounding around Cu1 with rather long d(Cu1-I) = 3.062 – 3.272 Å. Analogously to XIV, in XV crystal structure is composed of an extended network made of BiI6 octahedra linked via corner-sharing by Ag-I units (Figure 7.3a) with cavities are filled 3- + with BiI6 octahedra, Rb , and H2O (Figure 7.3b). Ag1 ions forming the cube have occupancies of 44.1(6) %. The electron density in the center of the cube (Figure 7.3c) was described with the Ag atom with 68(2) % occupancy. I4 and I6 are disordered on two positions (Figure 7.3c, d)

Figure 7.2. (a) A complex network of XIV: BiI6 octahedra (four of six iodide ions are disordered on two positions) are connected with Cu-I units via corner-sharing forming a 3D 3- + anionic network with channels filled with BiI6 octahedra (violet polyhedra), Rb cations and water molecules. (b) A copper iodide unit, which consists of Cu-I cube with all the faces capped with I- (ellipsoids are shown at 50 % probability). with a total occupancy of I6 + I6xiv ~ 90 %. The Ag-I complex unit is less disordered comparing to Cu-I unit. Ag1 ions are in the tetrahedral surrounding of the I- ions with d(Ag1-I) = 2.614 –

2.987 Å, which is similar to the Ag-I distances in C26H78N18O6Ag8Pb6I16 - d(Ag-I) = 2.767(4) – 2.9782(18) Å - that is composed of layers made of a complex Pb(II) based cation connected - 347 to Ag8I15 clusters and I ions situated in between the layers. Ag8I15 cluster is composed of Ag-I tetrahedral units connected via edges and faces. In XV there is a rather short Ag1-I6 distance of 2.614 Å present due to a disorder of I at this position (Figure 7.3c). The closest Ag2-

I distance is to 3.140(2) Å. The four I ions from six in BiI6 octahedra (those connected to Ag-I units) are disordered on two positions, which was also observed in XIV. In both XIV and XV crystal structures, there is a disorder of Rb1 (in XIV) and Rb2 (in XV) cations on two positions

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Rubidium copper and silver iodobismuthates | Olga Nazarenko with occupancies of 0.5 each (Table 7.2, Figure 7.3d). No analogous compounds to XIV and XV were found in the literature. Therefore, it is not completely clear if such a static disorder of Cu(Ag)-I unit can be described as a disordered cluster unit or rather a superposition of a few coordination units. The purity of XIV and XV compounds was estimated with a powder XRD (Figure 7.4a, b, collected in the transmission mode through the powders deposited into a 0.5 mm glass capillaries) by comparing the experimental powder XRD patterns with the calculated ones (based on the single crystal XRD data). No other crystalline phases were detected. Thermal analysis was performed in the inert atmosphere using an alumina crucible as a sample holder. A mass reduction before 200 ℃ for XIV (peak at 159 ℃, reduction of 2.13 % of mass) and XV

(peak at 142 ℃, a mass reduction of 2.05 %) is likely to result from H2O loss or H2O and HI in + case there are hydronium (H3O ) ions present in the structure. The further decomposition steps are connected with the sublimation of the metal iodides. Due to challenges encountered during the single crystal XRD analysis, other methods, such as WD-XRF and EDXS, were employed to study the composition of the obtained phases. The results are presented in Table 7.3. All three analyses give a different result. Although, according to EDXS (performed in SEM mode) the compounds have a homogenous composition (minimum deviations of the composition from crystal to crystal). The discrepancies in the compositions are largest when it comes to the amounts of Cu and Ag, and I in XIV. The reason for these inconsistencies might lie in the different composition of the bulk of the crystal comparing to its surface (as the resolution for EDXS, for instance, is ~1-2 μm in depth), and in the use of high vacuum necessary for XRF and EDXS measurements that might facilitate a decomposition of the samples. Thus, single crystal XRD, WD-XRF, and EDXS analyses did not elucidate the matter with the composition of the obtained phases. To perceive the problem with the charge neutrality in compounds, additional measurements, such as electron paramagnetic resonance [EPR, for paramagnetic Cu(II), Ag(II) species], and 63Cu solid-state nuclear magnetic resonance [SSNMR, for Cu(I)] were conducted on single crystals of XIV and XV. However, neither with EPR nor with 63Cu SSNMR any signals were detected, which indicates the complexity of the structures and the nature of M-I, M···M (M = Ag, Cu) interactions in XIV and XV. The 1H SSNMR (Figure 7.5a) analysis of XIV (Rb-Cu-Bi-I-OH system) shows a complex spectrum with numerous spinning sidebands. There are several types of protons with different local magnetic fields present in the structure, but it is not possible to disentangle the different species. Furthermore, the 87Rb SSNMR spectrum of XIV reveals a signal that, although, not easy to interpret (Figure 7.5b), is characterized by short relaxation times expected for this quadrupolar nucleus. Homogeneous broadening must be significant as almost no signal can be recovered for echo delays of not even

0.8 ms. This fast T2-relaxation could be potentially induced by fast dynamics and mobility of rubidium in the range of the Larmor frequency of 87Rb (229.131 MHz at 16.4 T), which is in agreement with the crystallographic data, as part of the Rb+ ions are disordered on two positions. The asymmetric shape of the NMR signal can originate either from larger order quadrupole interactions which are not averaged out by magic angle spinning (MAS), i.e., second-order perturbation broadening, or by numerous Rb species (inhomogeneous broadening, site disorder) – or the combination of both. Under static experimental conditions, a broad unresolved signal

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Rubidium copper and silver iodobismuthates | Olga Nazarenko is observed (black spectrum in Figure 7.5b). When spinning the sample at 5 kHz under MAS, spinning sidebands appear, which are poorly resolved/overlap because of the fast T2 relaxation

Figure 7.3. A crystal structure of XV, which consists of a 3D network of BiI6 octahedra bound with Ag-I complex coordination units via corner-sharing (a), with voids, filled with Rb+ cations, 3- solvent molecules, and isolated BiI6 octahedra (b). (c) Ag-I complex unit can be described as an Ag cube with I- ions capping all faces of this cube, and another Ag ion in the center (ellipsoids are shown at 50 % probability). (d) The occupancies of the selected atoms. Table 7.3. Compositional analysis for compounds XIV and XV according to different instrumental techniques.

Compound Single crystal XRD Wavelength dispersive The energy dispersive X-ray fluorescence X-ray spectroscopy

XIV Rb20Cu8.32Bi6I48.14O8 Rb21.5Cu6.1Bi6I48.8 Rb22.4Cu6.9Bi6I41.6

XV Rb20Ag8.42Bi6I47.49O8 Rb20.4Ag6.3Bi6I47.3 Rb20.7Ag6.2Bi6I47.4

leading to the broadening of the central transition and the sidebands (blue spectrum in Figure 7.5b). Notably, the dark green color of IV is a property of the surface, whereas, when crystals are broken or polished a red color of the bulk is observed. Such a difference in color insinuates a difference in the chemical composition of the surface comparing to the bulk of the crystals. The Kubelka-Munk function was used to estimate the difference in absorption edge between

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XIV, XV, and BiI3 (Figure 7.6). Thus, comparing to BiI3 - a 2D compound composed of Bi-I edge-sharing octahedra348 – XV and XIV have larger optical bandgaps. Interestingly, the second excitonic feature in XIV and XV has the same energy for both compounds. The single crystals of XIV and XV (grown with method II and II’, see the Experimental section) showed a specific resistivity of 108 Ω m.

Figure 7.4. (a, b) The powder XRD patterns of powdered samples of XIV and XV. The calculated patterns are based on the experimental single-crystal XRD data obtained in this project. A shift of the reflections in the experimental patterns (measured at RT) towards smaller 2Θ values comparing to the calculated patterns is because the single crystal XRD data was collected at 100 K (at such low T the lattice constants get smaller).The presence of small background in the powder XRD patterns of both samples might result from the amorphization of the compounds when mortaring, or a bad diffraction due to a strong static disorder. (c, d) The thermal analysis of XIV and XV.

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Figure 7.5. (a) 1H SSNMR spectra of XIV, spectrum acquired under 11.8 kHz spinning frequency. (b) 87Rb SSNMR spectra measured under static (black) conditions and while spinning the sample (blue) at 5 kHz at the magic-angle.

Figure 7.6. (a) Kubelka-Munk function of XIV and XV comparing to BiI3. (b, c) Photographs of XIV and XIV taken under the daylight.

7.4. Conclusions In conclusion, in this Chapter a study on the semiconductor compounds XIV and XV obtained in the Rb-Cu-Bi-I-OH and Rb-Ag-Bi-I-OH systems was presented. The synthesized compounds have complex crystal structures consisting of BiI6 octahedra connected via corner- sharing with Cu-I or Ag-I coordination polyhedra into networks with large channels filled with + 3- Rb , H2O as well as isolated BiI6 octahedra. The Cu and Ag iodide units are best described as metal cubes with all faces capped with I- ions and an additional metal ion in the center of the cube. According to the empirical formula obtained from single crystal XRD analysis, the charge neutrality condition is not fulfilled for both compounds. With additional EPR (on Ag, Cu), 1H, 63Cu SSNMR measurements it was, unfortunately, possible neither to ascertain the oxidation

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Rubidium copper and silver iodobismuthates | Olga Nazarenko states of Cu and Ag (no signal were observed with either of methods) nor to procure a clear + evidence of a presence of H3O cations (proton species were found, but it was not possible to resolve a complex signal). It is possible to grow a few mm large single crystals of XIV and XV. The XIV and XV compounds with fascinating crystal structures that present a perplexing problem of charge neutrality, an elemental composition, and Cu and Ag oxidation states, underline a variety of the anionic lattices that can be obtained within A+-B3+-B+-X- systems.

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Chapter 8. Conclusions and outlook

8.1. Conclusions This dissertation comprises five stories that include fundamental studies on the crystal structure, optical, and electronic properties of novel ternary, quaternary and quinary Pb(II), Sn(II), and Bi(III) halide coordination compounds. In particular, three projects presented in Chapters 3-5 concern 3D and 2D LHP compounds. The electronic properties of the 3D LHPs (Chapter 1), which are composed of high Z-value Pb and I atoms, were probed through their use as γ-radiation detectors. In Chapters 6-7, studies were conducted on the lead-free, cesium guanidinium Sn(II) bromide system and rubidium copper (silver) iodobismuthates. The obtained results are summarized below:

(1) The phase stability of the cubic FAPbI3 compound was drastically improved by a + - - substitutional doping of the FA cation with Cs , and I with Br ions. The CsxFA1−xPbI3−yBry (x = 0–0.1, y = 0–0.6) compounds were grown as single crystals (few mm in size) from solution utilizing an inversed temperature solubility method. According to the powder XRD, the compounds are pure phases, and the reflections are shifted (without broadening) towards higher 2Ɵ values when Cs+ and Br- are incorporated in the lattice. The PL maxima can be tuned by changing the amount of the dopant, and it is blue-shifted compared to the parental FAPbI3. High charge carriers’ mobility × lifetime products of up to 1.210-1 cm2 V-1 were obtained for these single crystals, underlining their good electronic quality. The resistivity of the single crystals drops with an increase in the Br concentration, and the dark current concomitantly increases. A compromise between phase stability and dark-current values was found for the

Cs0.1FA0.9PbI2.8Br0.2 composition, which also showed the best performance as a γ-radiation

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Conclusions and outlook | Olga Nazarenko detector. Additionally, the region for efficient γ-ray counting was calculated for

Cs0.1FA0.9PbI2.8Br0.2 single crystals, and it was found to fall in the region with energies < 400 keV.

(2) Layered Cs[C(NH2)3]PbI4 and Csn[C(NH2)3]PbnBr3n+1 (n = 1, 2) were crystallized from aqueous acidic solutions, and these compounds were found to be Ruddlesden-Popper-related + + phases. Fascinatingly, Cs and C(NH2)3 cations are observed to periodically alternate with one another in the interlayer space between the layers of corner-sharing lead-halide octahedra. + Moreover, C(NH2)3 cations are closely packed due to the Van der Waals interactions; they also connect the lead-halide layers through hydrogen-halogen bonds. The bandgap of

Cs[C(NH2)3]PbI4 is 0.1 eV smaller than the one of known (n-BuNH3)2PbI4, and its specific density is 1.54 times greater. Another very important factor is distortions of Pb-I octahedra. The in-plane distortions of Pb-I-Pb angles in (n-BuNH3)2PbI4 are larger than in Cs[C(NH2)3]PbI4, which is consistent with the difference in the bandgaps. Cs[C(NH2)3]PbI4 and Csn[C(NH2)3]PbnBr3n+1 (n = 1, 2) possess good thermal stability and decompose above 571 K. + In the case of Cs2[C(NH2)3]Pb2Br7, Cs facilitates the formation of thicker slabs by filling the voids within the lead-iodide layers. Cs2[C(NH2)3]Pb2Br7 is luminescent under UV light at RT in the blue region. Cs[C(NH2)3]PbBr4 is luminescent in blue and Cs[C(NH2)3]PbBr4 in green under UV light upon cooling. Cs2[C(NH2)3]Pb2Br7 and Cs[C(NH2)3]PbI4 are photoconductive with dark resistivities of 3108 Ω∙cm and 5109 Ω∙cm, respectively.

(3) In the FA-G-Pb-I system a common phase with a formula FAGPbI4 was obtained. FAGPbI4 is composed of corrugated Pb-I layers that are arranged in a stair-like configuration with FA and G cations situated in the interlayer space. It was possible to grow few mm crystals of 10 FAPbI4. The compound is photoconductive, and has a specific resistivity of 110 Ω∙cm. It is luminescent in the red region with a max PLQY of 3.5 %. The purity of the synthesized compound was studied with powder XRD and 207Pb SSNMR. In detail, the Pb207 NMR spectrum of the FAGPbI4 was compared to spectra of -FAPbI3, α-FAPbI3, G2PbI4 and possible side products, such as PbO, PbCO3, Pb(OH)I. According to the powder XRD and SSNMR FAGPbI4 was synthesized as a pure phase. Therefore, it was concluded, that the complex PL spectrum (resulting from several types of emissive transitions with different lifetimes) originated from near-band-edge excitonic states as well as STEs, and possibly defects in the FAGPbI4, and not from the impurities. According to the thermal analysis, FAGPbI4 is stable up to 528 K. This study underlined the importance of fundamental studies on the complex multi-cation and multi- anion systems used in solar cell research, while, within typically used compositional space – Cs+, Rb+, FA, MA, G as A-type cations, Pb(II), Sn(II), Ag(I), Bi(III) as B-site cations and Cl, Br, I anions, other than 3D perovskite phases can form.

(4) In the G-Cs-Sn-Br system, two common phases, Cs2[C(NH2)3]Sn2Br7 and Cs[C(NH2)3]SnBr4, were crystallized from an aqueous hydrobromic acid. The Cs[C(NH2)3]SnBr4 compound could only be obtained as an impurity in the Cs2[C(NH2)3]Sn2Br7 phase. The crystal structure of both compounds was determined, and Cs2[C(NH2)3]Sn2Br7 and

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Cs[C(NH2)3]SnBr4 were found to be Ruddlesden-Popper-related phases and are isostructural to the corresponding lead compounds - Csn[C(NH2)3]PbnBr3n+1 (n=1, 2). The pure [C(NH2)3]2SnBr4, on the other hand, is a 1D polymer, where the Sn-Br chains are composed of corner-sharing square pyramids, and the G cations are situated in between the chains forming hydrogen bonds. Based on the coordination surrounding of the Sn(II) ions, the tin 5s2 lone pair is stereoactive. [C(NH2)3]2SnBr4 has a broad (FWHM ~121 nm) emission band resulting from the STEs with a maximum at ~557 nm, a PLQY of 2 %, and a large Stokes shift of 1.3 eV. At

77 K the PLQY increases to 75±5 %. The optical bandgap of CsSnBr3 is 1.74 eV whereas it is about 3.11 eV for [C(NH2)3]2SnBr4, which is consistent with their respective dimensionality -

3D and 1D. Cs2[C(NH2)3]Sn2Br7 has a smaller optical bandgap than its lead analog, and this might be due to several factors: a lower ionization energy for Sn, E3(Sn) = 30.5 eV vs. E3(Pb) = 31.94 eV, differences in electronegativity between Sn and Pb, and different distortions of the metal-halide octahedra. Interestingly, although CsSnBr3 is a black 3D compound, CsSnI3 is - + yellow and has a 1D polymeric network. The addition of I ions to the HBr(aq.) containing Cs 2+ and Sn results in the formation of an orange product with the formula Cs8Sn6Br13I7. This orange compound is a 2D polymer with layers consisting of distorted, corner- (through μ2-Br) -1⅔ -1⅙ and edge-sharing (through μ3-I ions) [SnBr4I2] and [SnBr3I3] octahedra.

(5) Pursuing toxic-element-free semiconductor compounds, the Rb-Cu-Bi-I and Rb-Ag-Bi-I systems were investigated. The water and hydronium ions are a part of the crystal lattices, as the synthesis was conducted in the aqueous acidic solutions. Two compounds, XIV in Rb-Cu- Bi-I-OH system and XV in Rb-Ag-Bi-I-O-H system were obtained. Compounds XIV and XV have fascinating and complex crystal structures. Their empirical formulas, obtained from single crystal XRD were determined to be Rb20Cu8.32Bi6I48.14O8 for XIV and Rb20Ag8.42Bi6I47.49O8 for XV. This suggests that charge neutrality is not maintained. To understand the composition of the synthesized compounds and the oxidation states of Cu and Ag methods such as EPR, SSNMR, WD-XRF, and EDXS were used. Unfortunately, signal was detected neither in EPR for both Cu and Ag, nor in SSNMR for Cu. Therefore, no additional information on the oxidation states of Cu and Ag was obtained. The question with the charge neutrality can be answered assuming the presence of hydronium ions in the voids of these structures, but no direct proof was found. The 1H SSNMR gave a complex spectrum indicating several unresolvable proton species. The crystal structures of XIV and XV are made of BiI6 octahedra connected via corner-sharing by Cu and Ag iodide complex units that all-together form a carcass with large channels. Cu and Ag are in the tetrahedral surrounding of I- ions. The channels within the lattice 3- + + are filled with isolated BiI6 octahedra, water molecules, Rb , and H3O ions. The optical 348 bandgaps of XIV and XV are larger comparing to BiI3. The resistivity of the single crystals of XIV and XV are 108 Ω m. With SSNMR in XIV the dynamic movement of the Rb+ cations within the lattice was observed through the fast T2-relaxation time. This result was supported by single crystal XRD, as some Rb atoms is the crystal structure are disordered in two positions.

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8.2. Outlook Research on metal-halide perovskites is highly dynamic, and numerous efforts are dedicated to the understanding of the intrinsic optical and electronic properties of APbX3 (A = Cs+, MA, FA+; X=Cl-, Br-, I- or a mixture thereof) with the help of both experimental techniques and theoretical calculations. These compounds have been tested not only as a photon to energy converters in solar cells but also for use in LEDs, lasers, broadband and narrowband photodetectors operating in the ultraviolet-visible-near infrared regions, as well as soft X-ray and even γ-radiation detectors. LHPs are in general investigated in forms of single crystals, thin films, as well as nanoparticles. With regards to the potential use of the investigated here

Cs0.1FA0.9PbI2.8Br0.2 single crystals in γ-radiation detectors, it is possible using single crystals of sizes 0.3-1 cm and radiation energies < 400 keV in small personal devices. The growth of single crystals of sizes > 1 cm of high quality, however, needs further development. Such large crystals might be obtained by - multiple transfers of smaller crystals to fresh, warm solutions until a necessary crystal size has been reached (as was observed from experiments reported in literature on FAPbI3 and

MAPbI3, and as was performed on the quinary systems presented here by a visiting student, Oleh Hordiychuk, in the group of Prof. Kovalenko) - by developing a setup that would allow a slow-flow of a fresh solution to pass over the single crystals, and thus bring the necessary building blocks for the creation of larger crystals. There remain several important issues with LHPs that should be addressed:

- phase transitions, for instance, MAPbI3 undergoes a phase transition from a tetragonal crystal structure to a cubic one at ~ 327 K, which means a change in the volume and properties of the compound;

- the preparation of APbX3 is performed from toxic solvents like N,N- dimethylformamide, or dimethyl sulfoxide that can penetrate deep into the human tissues and intoxicate the body with the lead dissolved in it; - the toxicity of lead. As to the third point, it was discussed in the introduction of this work that the amount of lead is currently restricted in products and should constitute only 0.1 % of the homogeneous product by mass. While LHP nanoparticles can meet these RoHS restrictions in TV displays, the use of LHPs as thin polycrystalline films in solar cells exceeds the allowed Pb(II) limit. It is, therefore, a problematic question if there should be an exemption for the use of LHPs in solar cells, photodetectors, and other devices as well given that the stability of LHPs is further improved to meet industrial demands. Nowadays, a principal research direction is finding alternative semiconductor compounds with suitable optoelectronic properties to substitute LHPs. The substitution of Pb(II) for Ge(II) or Sn(II) was reported in the literature, but unfortunately, Sn(II), and especially Ge(II) easily oxidize to Sn(IV) and Ge(IV). When all work was conducted under inert conditions, S. Shao et al. could obtain reproducible FASnI3 based solar cells with an efficiency of 9 %.349 The investigation of Sn(II) hybrid halide perovskites also brings new insights into the structure-property relationships that exist in these materials, which were also investigated

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Conclusions and outlook | Olga Nazarenko here. Although, the actual application of tin(II) halide coordination compounds is hindered by the ease of oxidation of Sn(II) to Sn(IV).

Concerning quarternary A2B(I)B(III)X6 systems; they are more complex than LHPs, and challenges are to obtain a direct bandgap semiconductor that can be processed into uniform thin films of high purity. The successful substitution of Ag(I) and Bi(III) in Cs2AgBiBr6 for Tl(I) and Tl(III) was shown in 2018 that lead to a direct bandgap semiconductor. However, the toxicity of Tl is a substantial hinder to any practical use. Conceptually, though, this reported study demonstrated the possibility to achieve a direct bandgap semiconductor with a notably small bandgap (0.95 eV) in the case of Cs2AgTlBr6. Finding toxic-element-free, thermodynamically stable metal-halide semiconductors with excellent optoelectronic properties is a challenging task. The investigation of Sn(II), Bi(III)/B(I) halide complexes has brought some novel compounds to light that have high PLQYs, a semiconductor to metallic properties, or can detect soft X-rays. In general, 2D and 3D ternary and higher semiconductor metal halide compounds based on Pb(II), Sn(II), Bi(III), or Sb(III) are of interest for the application in solar cells, photodetectors, as for these applications, among others, the efficient charge transport is important. Whereas, 0D compounds - for instance, some of mixed organic-inorganic tin halides – have high exciton binding energies, which can result in high PLQYs. Such compounds, therefore, can be efficient lighting sources.

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Chapter 9. Appendix

Appendix to Chapter 3 Table S3.1. Crystal data and a structure refinement for formamidinium iodide.

Empirical formula CH5N2I Formula weight 171.97 Temperature (K) 300 Crystal system monoclinic

Space group P21/c a (Å) 4.8386(9) b (Å) 13.831(3) c (Å) 7.0496(13) α (°) 90 β (°) 97.196(3) γ (°) 90 Volume (Å3) 468.07(15) Z 4 3 ρcalc (g/cm ) 2.440 μ (mm-1) 6.652 F(000) 312.0 Crystal size (mm3) 0.58 × 0.04 × 0.04 Radiation MoKα (λ = 0.71073) 2Θ range for data collection (°) 5.89 to 57.394 Index ranges -6 ≤ h ≤ 6, -18 ≤ k ≤ 18, -9 ≤ l ≤ 9 Reflections collected 4978

Independent reflections 1213 [Rint = 0.0303, Rsigma = 0.0252] Data/restraints/parameters 1213/0/38 Goodness-of-fit on F2 1.193

Final R indexes [I>=2σ (I)] R1 = 0.0230, wR2 = 0.0526

Final R indexes [all data] R1 = 0.0268, wR2 = 0.0542 Largest diff. peak/hole (e Å-3) 0.70/-0.83

Table S3.2. Fractional atomic coordinates and isotropic or equivalent isotropic displacement parameters (Å2).

x y z Uiso*/Ueq I1 0.37747 (4) 0.88392 (2) 0.70654 (3) 0.04315 (12)

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N2 1.1108 (6) 0.5832 (2) 0.6610 (4) 0.0516 (7) H2A 1.1199 0.5374 0.7442 0.062* H2B 1.2198 0.5827 0.5736 0.062* N1 0.7629 (6) 0.6603 (2) 0.7900 (4) 0.0518 (7) H1A 0.7586 0.6176 0.8782 0.062* H1B 0.6514 0.7089 0.7842 0.062* C1 0.9358 (7) 0.6516 (2) 0.6673 (4) 0.0430 (7) H1 0.9333 0.6995 0.5745 0.052*

Table S3.3. Atomic displacement parameters (Å2).

U11 U22 U33 U12 U13 U23 I1 0.04721 (16) 0.04141 (15) 0.04223 (16) −0.00079 (7) 0.01107 (9) −0.00306 (8) N2 0.0506 (15) 0.0587 (17) 0.0473 (16) 0.0022 (14) 0.0130 (13) −0.0030 (14) N1 0.0509 (16) 0.0486 (16) 0.0563 (17) 0.0073 (13) 0.0081 (13) −0.0038 (14) C1 0.0462 (16) 0.0413 (16) 0.0396 (15) −0.0105 (14) −0.0018 (13) 0.0058 (14)

Table S3.4. Selected geometric parameters (Å, °).

N2—H2A 0.8600 N1—H1B 0.8600 N2—H2B 0.8600 N1—C1 1.282 (4) N2—C1 1.274 (5) C1—H1 0.9300

N1—H1A 0.8600 H2A—N2—H2B 120.0 C1—N1—H1B 120.0 C1—N2—H2A 120.0 N2—C1—N1 125.9 (3) C1—N2—H2B 120.0 N2—C1—H1 117.1 H1A—N1—H1B 120.0 N1—C1—H1 117.1 C1—N1—H1A 120.0

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Table S3.5. Comparison of the theoretical composition (i.e. expected from the composition of the mother solution) to the elemental analysis

Sample (nominal mixing Theoretical (%) Elemental analysis of ratio in the mother single crystals (%) solution)

Cs0.1FA0.9PbI3 [C] 1.68 [C] 1.61 [H] 0.71 [H] 0.68 [N] 3.93 [N] 3.98

Cs0.1FA0.9PbI2.9Br0.1 [C] 1.70 [C] 1.65 [H] 0.71 [H] 0.67 [N] 3.96 [N] 3.67 [Br] 1.25 [Br] 1.68 [I] 57.77 [I] 57.17

Cs0.1FA0.9PbI2.8Br0.2 [C] 1.71 [C] 1.65 [H] 0.72 [H] 0.65 [N] 3.99 [N] 3.78 [Br] 2.53 [Br] 3.11 [I] 56.19 [I] 55.33

Cs0.1FA0.9PbI2.6Br0.4 [C] 1.74 [C] 1.72 [H] 0.73 [H] 0.57 [N] 4.05 [N] 3.94 [Br] 5.13 [Br] 7.63 [I] 52.96 [I] 49.66

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Figure S3.1. Powder XRD of CsxFA1-xPbI3 (x = 0, 0.05, 0.1): (a) over a wide 2θ range and (b) with a focus on the (100) reflection [the black dashed lines are calculated as the centres of mass (i.e., the first mathematical moments) of the corresponding reflections].

Figure S3.2. Specific resistivity of single crystals of various compositions.

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Appendix to Chapter 4 Table S4.1. Crystallographic data for selected hybrid LHPs.

Cs[C(NH2)3]PbI4 Cs[C(NH2)3]PbBr4 Cs2[C(NH2)3]Pb2Br7 Cs[C(NH2)3]PbI4 (I) (II) (III) (I’) Formula weight 907.79 719.83 1299.66 907.79 Temperature (K) 300 300 300 250 Crystal system orthorhombic orthorhombic orthorhombic orthorhombic Space group Pnnm Imma Cmmm Pnnm Color red yellow yellow red a (Å) 12.7426(13) 6.0613(6) 6.0029(12) 12.6947(12) b (Å) 18.5776(19) 11.5868(12) 28.773(6) 18.6066(18) c (Å) 12.2390(13) 17.4591(18) 11.598(2) 12.1767(12) Volume (Å3) 2897.3(5) 1226.2(2) 2003.2(7) 2876.2(5) Z 8 4 4 8

3 4.162 3.899 4.309 4.193 ρcalc (g/cm )

μ (mm-1) 22.607 29.671 34.311 22.773 F(000) 3056.0 1240.0 2208.0 3056.0 Crystal size 0.3 × 0.07 × 0.02 0.2 × 0.08 × 0.02 0.46 × 0.08 × 0.04 0.28 × 0.1 × 0.02 (mm3) Radiation MoKα (λ = 0.71073) MoKα (λ = MoKα (λ = 0.71073) MoKα (λ = 0.71073) 0.71073) 2Θ range for data 3.876 to 63.122 4.218 to 62.904 3.512 to 62.718 3.884 to 60.052 collection (°) Index ranges -18≤h≤17, -8≤h≤8, -8≤h≤8, -17≤h≤17, -26≤k≤26, -17≤k≤16, -41≤k≤41, -26≤k≤26, -17≤l≤17 -25≤l≤24 -16≤l≤16 -17≤l≤17 Reflections 33,586 7142 11,917 32,295 collected

Independent 4767 [Rint = 0.0637, 1094 [Rint = 1814 [Rint = 0.0473, 4385 [Rint = reflections Rsigma = 0.0432] 0.0406, Rsigma = Rsigma = 0.0370] 0.0600, Rsigma = 0.0295] 0.0469] Data/restraints/ 4767/18/109 1094/0/40 1814/0/58 4385/18/87 parameters Goodness-of-fit 1.004 1.041 1.035 1.008 on F2

Final R indexes R1 = 0.0389, wR2 = R1 = 0.0223, wR2 R1 = 0.0321, wR2 = R1 = 0.0352, wR2 0.0722 = 0.0519 0.0772 = 0.0806 [I >= 2σ (I)]

Final R indexes R1 = 0.0835, wR2 = R1 = 0.0261, wR2 R1 = 0.0407, wR2 = R1 = 0.0693, wR2 [all data] 0.0888 = 0.0535 0.0823 = 0.0985 Largest diff. 1.36/-1.55 0.84/-0.88 2.21/-1.52 2.44/-1.54 peak/hole (e Å-3)

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Figure S4.1. The reciprocal space reconstructions of the hk0-layers at RT and 250K (a* to the right and b* up) clearly show the weak superstructure reflections (rows of superstructure reflections marked with arrows) requiring a doubling of the crystallographic a-axis. The superstructure reflections become stronger at lower temperature, indicating that the deviations with respect to the high symmetry substructure with a’=a/2 are even more pronounced.

Cs[C(NH2)3]PbI4 (I): When taking only strong reflections into account, indexing of the diffraction pattern is possible with a small body centered orthorhombic unit cell with a' = 6.37 Å, b' = 12.24 Å and c' = 18.58 Å. A careful examination of the reciprocal space reconstructions clearly reveals weak superstructure reflecions requiring a doubling of the a'-axis, finally

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Appendix | Olga Nazarenko leading to a primitive orthorhombic unit cell of a = 12.7426 Å (2 x a'), b = 18.5776 Å and c = 12.239 Å (according to the standard setting in the space group Pnnm). Since the superstructure reflections are only very weak the refinement needs a slight damping (DAMP 20) to avoid large shifts/esd ratios. For I and I’ in order to obtain a correct unit cell a twin indexing was necessary. While a twin domain in each case was negligibly small, an integration was done for the main domain only.

Table S4.2. C···N distance between neighboring guanidinium molecules.

Compound C - N d (Å) C1 - N1 3.188(1) C1 - N4 3.207(1) Cs[C(NH2)3]PbI4 C2 - N1 3.206 C2 – N4 3.179

Cs[C(NH2)3]PbBr4 C1 - N2 3.036(1)

Cs2[C(NH2)3]Pb2Br7 C1 - N1 3.014(1)

Table S4.3. Fractional atomic coordinates and isotropic or equivalent isotropic displacement 2 parameters (Å ) for Cs[C(NH2)3]PbI4 (I).

x y z Uiso*/Ueq Occ. (<1) Pb1 0.37509 (8) 0.74989 (2) 0.25057 (2) 0.02935 (8) I4 0.37240 (11) 0.91990 (2) 0.19690 (4) 0.04476 (13) I2 0.37771 (11) 0.58237 (2) 0.30614 (4) 0.04658 (13) Cs2 0.37518 (15) 0.38744 (4) 0.5000 0.05799 (18) I5B 0.33718 (15) 0.71836 (9) 0.0000 0.0543 (4) 0.872 (5) I3A 0.41452 (17) 0.78161 (9) 0.5000 0.0523 (4) 0.872 (5) I1 0.62513 (9) 0.75058 (2) 0.22422 (5) 0.05065 (14) Cs1 0.62532 (16) 0.90195 (4) 0.0000 0.05791 (18) C1 0.877 (2) 0.4569 (5) 0.0000 0.0400 (19) N1 0.8739 (17) 0.5281 (4) 0.0000 0.0479 (19) H1A 0.8699 0.5511 0.0608 0.057* 0.5 H1B 0.8759 0.5513 −0.0608 0.057* 0.5 C2 0.377 (2) 0.4589 (5) 1.0000 0.048 (2) N2 0.8737 (12) 0.4229 (3) 0.0933 (5) 0.0523 (15) H2A 0.8694 0.3767 0.0946 0.063* H2B 0.8759 0.4467 0.1536 0.063* N4 0.3741 (18) 0.5296 (4) 1.0000 0.055 (2) H4A 0.3714 0.5527 1.0608 0.066* 0.5

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H4B 0.3749 0.5527 0.9392 0.066* 0.5 N3 0.3757 (15) 0.4263 (4) 0.9074 (7) 0.087 (2) H3A 0.3738 0.3801 0.9051 0.104* H3B 0.3767 0.4507 0.8477 0.104* I5A 0.3765 (16) 0.7252 (7) 0.0000 0.0543 (4) 0.128 (5) I3B 0.3682 (14) 0.7844 (7) 0.5000 0.0523 (4) 0.128 (5)

Table S4.4. Atomic displacement parameters (Å2) for I.

U11 U22 U33 U12 U13 U23 Pb1 0.02713 0.03339 0.02754 0.00023 (8) −0.00045 −0.00131 (8) (12) (14) (13) (7) I4 0.0556 (3) 0.0342 (2) 0.0444 (3) 0.0001 (5) −0.0007 (6) 0.00038 (17) I2 0.0547 (3) 0.0347 (2) 0.0503 (3) 0.0006 (5) −0.0016 (7) −0.00062 (18) Cs2 0.0472 (4) 0.0721 (5) 0.0546 (4) 0.0022 (6) 0.000 0.000 I5B 0.0917 (11) 0.0486 (6) 0.0225 (3) −0.0137 (7) 0.000 0.000 I3A 0.0916 (11) 0.0424 (4) 0.0230 (3) −0.0031 (7) 0.000 0.000 I1 0.0240 (2) 0.0489 (3) 0.0791 (3) −0.00024 −0.0001 (5) −0.0015 (2) (19) Cs1 0.0626 (4) 0.0659 (5) 0.0452 (4) −0.0002 (6) 0.000 0.000 C1 0.034 (4) 0.039 (4) 0.047 (4) −0.004 (8) 0.000 0.000 N1 0.048 (4) 0.043 (4) 0.052 (5) −0.002 (7) 0.000 0.000 C2 0.045 (5) 0.041 (4) 0.059 (5) 0.004 (9) 0.000 0.000 N2 0.068 (4) 0.045 (3) 0.044 (3) −0.002 (6) 0.001 (5) 0.004 (2) N4 0.056 (5) 0.045 (4) 0.065 (6) 0.003 (8) 0.000 0.000 N3 0.094 (6) 0.085 (5) 0.081 (5) 0.003 (9) 0.000 (8) −0.032 (4) I5A 0.0917 (11) 0.0486 (6) 0.0225 (3) −0.0137 (7) 0.000 0.000 I3B 0.0916 (11) 0.0424 (4) 0.0230 (3) −0.0031 (7) 0.000 0.000

Table S4.5. Table of hydrogen bonds for I.

D-H d(D-H) d(H..A)

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N3-H3A 0.86 2.897 148.73 3.659 I1 [ -x+1, -y+1, -z+1 ] N3-H3B 0.86 3.086 159.46 3.903 I2 [ x, y, -z+1 ]

Table S4.6. Fractional atomic coordinates and isotropic or equivalent isotropic displacement 2 parameters (Å ) for Cs[C(NH2)3]PbBr4 (II).

x y z U */U Occ. (<1) iso eq Pb1 0.5000 0.5000 0.5000 0.02526 (9)

Cs2 0.5000 0.2500 0.85462 (3) 0.05236 (15)

Br2 0.5000 0.44450 (5) 0.66797 (3) 0.04505 (14)

Br1 1.0000 0.5000 0.5000 0.0512 (2)

C1 0.0000 0.7500 0.7066 (4) 0.0316 (12)

N2 0.0000 0.7500 0.7828 (4) 0.0517 (15) H2A 0.0000 0.6857 0.8074 0.062* 0.5 H2B 0.0000 0.8143 0.8074 0.062* 0.5

N1 0.0000 0.6509 (5) 0.6719 (3) 0.0657 (14)

H1A 0.0000 0.6481 0.6227 0.079*

H1B 0.0000 0.5881 0.6981 0.079* Br3 0.4496 (3) 0.7500 0.52800 (6) 0.0555 (6) 0.5

Table S4.7. Atomic displacement parameters (Å2) for II.

U11 U22 U33 U12 U13 U23 Pb1 0.02294 (13) 0.02264 (13) 0.03020 (14) 0.000 0.000 −0.00119 (7) Cs2 0.0504 (3) 0.0432 (3) 0.0635 (3) 0.000 0.000 0.000 Br2 0.0529 (3) 0.0475 (3) 0.0347 (2) 0.000 0.000 −0.0009 (2) Br1 0.0232 (3) 0.0879 (7) 0.0425 (4) 0.000 0.000 −0.0068 (3) C1 0.026 (3) 0.035 (3) 0.034 (3) 0.000 0.000 0.000 N2 0.039 (3) 0.074 (4) 0.041 (3) 0.000 0.000 0.000 N1 0.058 (3) 0.061 (3) 0.078 (4) 0.000 0.000 −0.025 (3) Br3 0.099 (2) 0.0214 (4) 0.0458 (5) 0.000 0.0032 (6) 0.000

Table S4.8. Table of hydrogen bonds for II.

D-H d(D-H) d(H..A)

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Table S4.9. Fractional atomic coordinates and isotropic or equivalent isotropic displacement 2 parameters (Å ) for Cs2[C(NH2)3]Pb2Br7 (III).

x y z Uiso*/Ueq Occ. (<1) Pb1 0.5000 0.60158 (2) 0.24823 (2) 0.03092 (11) Cs1 0.5000 0.81452 (3) 0.0000 0.0562 (2) Br4 0.5000 0.5000 0.29753 (15) 0.0661 (4) Br2 0.5000 0.70236 (3) 0.19305 (7) 0.0515 (2) Br5 0.5000 0.57983 (5) 0.0000 0.0638 (4) Cs3 0.0000 0.5000 0.0000 0.0820 (4) Cs2 0.0000 0.5000 0.5000 0.0843 (5) Br1 0.0000 0.59626 (4) 0.24229 (8) 0.0598 (3) N1 0.0000 0.7680 (4) 0.5000 0.063 (3) H1A 0.0000 0.7830 0.5642 0.075* 0.5 H1B 0.0000 0.7830 0.4358 0.075* 0.5 N2 0.0000 0.7003 (3) 0.4008 (7) 0.071 (2) H2A 0.0000 0.6704 0.3990 0.085* H2B 0.0000 0.7159 0.3374 0.085* C1 0.0000 0.7222 (4) 0.5000 0.0379 (19) Br3 0.5513 (6) 0.61620 (6) 0.5000 0.0669 (12) 0.5

Table S4.10. Atomic displacement parameters (Å2) for III.

U11 U22 U33 U12 U13 U23 Pb1 0.02974 (15) 0.03251 (16) 0.03050 (15) 0.000 0.000 −0.00042 (8) Cs1 0.0559 (4) 0.0646 (5) 0.0481 (4) 0.000 0.000 0.000 Br4 0.0841 (9) 0.0296 (5) 0.0846 (9) 0.000 0.000 0.000 Br2 0.0610 (5) 0.0359 (4) 0.0576 (5) 0.000 0.000 0.0029 (3) Br5 0.0887 (9) 0.0738 (9) 0.0289 (5) 0.000 0.000 0.000 Cs3 0.0751 (8) 0.0485 (7) 0.1223 (13) 0.000 0.000 0.000 Cs2 0.0716 (9) 0.1210 (14) 0.0603 (7) 0.000 0.000 0.000 Br1 0.0294 (4) 0.0632 (6) 0.0869 (8) 0.000 0.000 −0.0024 (4) N1 0.046 (5) 0.043 (6) 0.098 (8) 0.000 0.000 0.000 N2 0.067 (5) 0.080 (6) 0.065 (5) 0.000 0.000 −0.020 (5) C1 0.031 (4) 0.040 (5) 0.043 (5) 0.000 0.000 0.000 Br3 0.114 (4) 0.0579 (9) 0.0289 (5) −0.0114 (11) 0.000 0.000

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Table S4.11. Table of hydrogen bonds for III.

D-H d(D-H) d(H..A)

Table S4.12. Fractional atomic coordinates and isotropic or equivalent isotropic displacement 2 parameters (Å ) for Cs[C(NH2)3]PbI4 at 250K (I’).

x y z U */U Occ. (<1) iso eq Pb1 0.62490 (7) 0.25000 (2) 0.25065 (2) 0.02434 (8)

I2 0.58009 (7) 0.21818 (4) 0.5000 0.0446 (2)

I3 0.62187 (9) 0.41734 (3) 0.30616 (4) 0.03708 (13)

I4 0.66905 (8) 0.28003 (4) 0.0000 0.0464 (2)

Cs2 0.87423 (13) 0.40452 (4) 0.5000 0.04461 (17)

I1 0.62796 (9) 0.08012 (2) 0.19688 (4) 0.03553 (12)

Cs1 0.37535 (12) 0.38675 (5) 0.5000 0.04694 (18)

I5 0.87494 (5) 0.25047 (3) 0.28387 (6) 0.03974 (13)

N3 0.1254 (16) 0.4706 (5) 0.0000 0.0492 (11) H3A 0.1248 0.4475 0.0612 0.059* 0.5 H3B 0.1258 0.4475 −0.0612 0.059* 0.5

C2 0.1257 (17) 0.5415 (6) 0.0000 0.0338 (17)

N4 0.1250 (11) 0.5766 (4) 0.0948 (6) 0.0492 (11)

H4A 0.1244 0.5532 0.1557 0.059*

H4B 0.1252 0.6228 0.0954 0.059*

N2 0.3744 (16) 0.5319 (5) 1.0000 0.0492 (11) H2A 0.3751 0.5550 1.0612 0.059* 0.5 H2B 0.3735 0.5550 0.9388 0.059* 0.5

C1 0.3749 (19) 0.4619 (6) 1.0000 0.0368 (18)

N1 0.3764 (11) 0.4282 (4) 0.9044 (6) 0.0492 (11)

H1A 0.3770 0.4526 0.8443 0.059*

H1B 0.3767 0.3820 0.9022 0.059*

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Table S4.13. Atomic displacement parameters (Å2) for I’.

U11 U22 U33 U12 U13 U23 Pb1 0.02184 0.02747 0.02372 0.00014 0.00034 0.00098 (8) (14) (15) (14) (9) (7) I2 0.0787 (6) 0.0354 (4) 0.0198 (3) −0.0046 0.000 0.000 (4) I3 0.0417 (3) 0.0288 (2) 0.0408 (3) 0.0004 (4) 0.0010 (5) 0.00084 (18) I4 0.0788 (6) 0.0405 (4) 0.0199 (3) −0.0062 0.000 0.000 (4) Cs2 0.0485 (4) 0.0486 (4) 0.0368 (3) −0.0007 0.000 0.000 (5) I1 0.0430 (3) 0.0280 (2) 0.0356 (2) 0.0002 (4) 0.0008 (5) −0.00028 (18) Cs1 0.0348 (3) 0.0591 (5) 0.0469 (4) 0.0017 (5) 0.000 0.000 I5 0.0187 (2) 0.0411 (3) 0.0594 (3) −0.0001 0.0007 (4) −0.0006 (2) (2) N3 0.060 (3) 0.045 (2) 0.043 (2) 0.001 (5) 0.000 0.000 C2 0.019 (4) 0.042 (3) 0.040 (3) 0.002 (5) 0.000 0.000 N4 0.060 (3) 0.045 (2) 0.043 (2) 0.001 (5) 0.000 0.000 N2 0.060 (3) 0.045 (2) 0.043 (2) 0.001 (5) 0.000 0.000 C1 0.027 (4) 0.044 (3) 0.040 (3) 0.002 (5) 0.000 0.000 N1 0.060 (3) 0.045 (2) 0.043 (2) 0.001 (5) 0.000 0.000

Table S4.14. Table of hydrogen bonds for I’.

D-H d(D-H) d(H..A)

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Figure S4.2. The powder XRD pattern of (n-C4H9NH3)2PbI4. The simulated pattern is based on the experimental single-crystal XRD data reported by D. B. Mitzi [Chem. Mater. 1996, 8 (3), 791-800]. The difference in intensities of some reflections in the experimental data from the simulated pattern comes from the preferred orientation of crystallites in the mortared powder, because (n-C4H9NH3)2PbI4 crystallizes as thin orange plates.

Figure S4.3. The powder XRD pattern of [C(NH2)3]2PbI4. The simulated pattern is based on the experimental single-crystal XRD data reported by M. Szafranski and A. Katrusiak [Phys. Rev. 2000, B61 (2), 1026 and ICSD card 92045].

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Figure S4.4. The powder XRD pattern of CsPbBr3. The simulated pattern is based on the experimental single-crystal XRD data reported by M. Rodova [see J. Therm. Anal. Calorim. 2003, 71 (2), 667-673 and ICSD card 97851].

Figure S4.5. Powder XRD pattern of (n-C4H9NH3)2PbBr4. The simulated pattern is based on the experimental single-crystal XRD data provided by Prof. M. Koshimizu and reported in [J. Phys. Chem. C, 2014, 118 (17), 9101 – 9106]. The difference in intensities of some reflections in the experimental data from the simulated patterns comes from the preferred orientation of crystallites in the mortared powder, because (n-C4H9NH3)2PbBr4 crystallizes as thin colorless plates.

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Figure S4.6. The powder XRD pattern of CsPb2Br5. The simulated pattern is based on the experimental single-crystal XRD data obtained in this work (ICSD number 432533).

Figure S4.7. The powder XRD pattern of [C(NH2)3]2PbBr4. The simulated pattern is based on the experimental single-crystal XRD data obtained in this work (ICSD number 432534).

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Figure S4.8. Powder XRD data of CsPbI3. The simulated pattern is based on the experimental single-crystal XRD data reported by C. K. Möller [see Nature 1958, 182 (4647), 1436-1436 and ICSD card 27979].

Table S4.15. Crystallographic data for selected compounds.

[C(NH2)3]2PbBr4 CsPb2Br5 Formula weight 1294.01 946.84 Temperature (K) 300 300 Crystal system triclinic tetragonal Space group P-1 I4/mcm Colour colourless colourless a (Å) 8.2753(8) 8.4833(10) b (Å) 10.4787(11) 8.4833(10) c (Å) 15.9791(16) 15.1830 (18) α (°) 90.954(2) 90 β (°) 91.981(2) 90 γ (°) 90.883(2) 90 Volume (Å3) 1384.4(2) 1092.7(3) Z 2 4 3 ρcalc (g/cm ) 3.104 5.756 μ (mm-1) 23.706 52.269 F(000) 1152.0 1576.0 Crystal size (mm3) 0.16 × 0.04 × 0.02 0.3 × 0.16 × 0.02

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Radiation MoKα (λ = 0.71073) MoKα (λ = 0.71073) 2Θ range for data 2.55 to 57.034 5.366 to 52.708 collection (°) Index ranges -10 h 10, -13 k 13, -21 l 21 -10 h 10, -10 k 10, - 18 l 18 Reflections collected 14481 4496

Independent reflections 6797 [Rint = 0.0346, Rsigma = 327 [Rint = 0.0500, Rsigma 0.0578] = 0.0308] Data/restraints/parameters 6797/0/235 327/0/16 Goodness-of-fit on F2 1.036 1.142

Final R indexes [I>=2σ (I)] R1 = 0.0396, wR2 = 0.0749 R1 = 0.0255, wR2 = 0.0578

Final R indexes [all data] R1 = 0.0646, wR2 = 0.0817 R1 = 0.0259, wR2 = 0.0582 Largest diff. peak/hole (e 1.06/-0.79 1.34/-1.17 Å-3)

Table S4.16. Fractional atomic coordinates and isotropic or equivalent isotropic displacement 2 parameters (Å ) for [C(NH2)3]2PbBr4.

x y z Uiso*/Ueq Pb01 0.21920 (3) 0.83285 (3) 0.24662 (2) 0.03824 (9) Pb02 0.71757 (4) 0.68632 (3) 0.24481 (2) 0.03991 (9) Br3 0.45606 (10) 0.75353 (8) 0.11485 (5) 0.0459 (2) Br5 0.95513 (10) 0.74887 (8) 0.11321 (5) 0.0464 (2) Br2 0.18901 (10) 1.06906 (8) 0.16395 (6) 0.0487 (2) Br1 −0.03950 (10) 0.93046 (9) 0.35510 (5) 0.0516 (2) Br4 0.48918 (10) 0.96305 (9) 0.33816 (6) 0.0510 (2) Br6 0.99083 (11) 0.56573 (9) 0.33890 (6) 0.0535 (2) Br8 0.45587 (11) 0.60134 (9) 0.35556 (6) 0.0550 (2) Br7 0.68152 (10) 0.43737 (9) 0.17259 (6) 0.0544 (2) N9 0.8281 (9) 0.1063 (8) 0.0487 (4) 0.059 (2) H9A 0.8063 0.1322 −0.0011 0.071* H9B 0.9267 0.0927 0.0646 0.071* N8 0.5592 (8) 0.1070 (7) 0.0786 (5) 0.057 (2) H8A 0.5345 0.1330 0.0291 0.068* H8B 0.4844 0.0939 0.1137 0.068* N7 0.7476 (9) 0.0496 (7) 0.1771 (4) 0.0508 (18) H7A 0.8469 0.0378 0.1922 0.061* H7B 0.6722 0.0367 0.2118 0.061*

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N6 0.3187 (9) 0.3915 (8) 0.0533 (5) 0.060 (2) H6A 0.4183 0.4019 0.0696 0.072* H6B 0.2950 0.3641 0.0033 0.072* N4 0.2382 (9) 0.4598 (7) 0.1816 (4) 0.059 (2) H4A 0.3375 0.4703 0.1982 0.070* H4B 0.1618 0.4767 0.2150 0.070* N3 0.3039 (11) 1.1589 (7) 0.4767 (5) 0.073 (3) H3A 0.3509 1.0982 0.4506 0.088* H3B 0.2958 1.1565 0.5302 0.088* N1 0.1765 (11) 1.3503 (8) 0.4726 (5) 0.082 (3) H1A 0.1407 1.4137 0.4444 0.099* H1B 0.1672 1.3495 0.5261 0.099* N5 0.0501 (9) 0.3991 (7) 0.0802 (5) 0.061 (2) H5A 0.0265 0.3699 0.0306 0.073* H5B −0.0258 0.4161 0.1139 0.073* C1 0.2444 (11) 1.2558 (9) 0.4347 (5) 0.048 (2) C3 0.7110 (10) 0.0873 (8) 0.1007 (5) 0.044 (2) N10 0.7827 (13) 0.1752 (8) 0.4811 (6) 0.092 (3) H10A 0.7541 0.1870 0.5318 0.111* H10B 0.8349 0.1079 0.4678 0.111* N2 0.2578 (12) 1.2579 (8) 0.3541 (5) 0.079 (3) H2A 0.2217 1.3215 0.3262 0.095* H2B 0.3026 1.1956 0.3285 0.095* C4 0.7472 (11) 0.2589 (8) 0.4246 (6) 0.049 (2) N12 0.7979 (11) 0.2404 (8) 0.3491 (5) 0.072 (3) H12A 0.7794 0.2962 0.3112 0.086* H12B 0.8500 0.1725 0.3371 0.086* C2 0.2030 (11) 0.4176 (8) 0.1047 (5) 0.046 (2) N11 0.6687 (13) 0.3611 (8) 0.4431 (6) 0.096 (4) H11A 0.6392 0.3744 0.4936 0.115* H11B 0.6462 0.4156 0.4050 0.115*

2 Table S4.17. Atomic displacement parameters (Å ) for [C(NH2)3]2PbBr4.

U11 U22 U33 U12 U13 U23 Pb01 0.03524 0.04164 0.03789 0.00255 0.00191 −0.00051 (17) (18) (17) (13) (12) (13) Pb02 0.03681 0.04349 0.03963 0.00115 0.00201 0.00470 (13) (17) (19) (18) (14) (13)

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Br3 0.0447 (5) 0.0566 (5) 0.0373 (4) 0.0150 (4) 0.0041 (3) 0.0051 (4) Br5 0.0443 (5) 0.0577 (5) 0.0370 (4) −0.0124 (4) 0.0047 (3) −0.0043 (4) Br2 0.0445 (5) 0.0443 (5) 0.0576 (5) 0.0002 (4) −0.0008 (4) 0.0113 (4) Br1 0.0494 (5) 0.0601 (6) 0.0464 (5) 0.0165 (4) 0.0110 (4) 0.0036 (4) Br4 0.0408 (5) 0.0579 (6) 0.0533 (5) −0.0003 (4) −0.0087 (4) −0.0076 (4) Br6 0.0467 (5) 0.0626 (6) 0.0515 (5) 0.0102 (4) −0.0032 (4) 0.0082 (4) Br8 0.0542 (5) 0.0614 (6) 0.0502 (5) −0.0027 (5) 0.0123 (4) 0.0102 (4) Br7 0.0449 (5) 0.0439 (5) 0.0740 (6) 0.0045 (4) −0.0022 (4) −0.0111 (4) N9 0.046 (4) 0.091 (6) 0.041 (4) 0.001 (4) 0.008 (3) 0.015 (4) N8 0.041 (4) 0.077 (6) 0.052 (5) −0.007 (4) −0.008 (3) 0.014 (4) N7 0.060 (5) 0.061 (5) 0.032 (4) 0.014 (4) 0.005 (3) 0.005 (3) N6 0.048 (4) 0.086 (6) 0.047 (4) 0.005 (4) 0.006 (3) −0.015 (4) N4 0.063 (5) 0.067 (5) 0.044 (4) 0.000 (4) −0.005 (4) −0.009 (4) N3 0.120 (7) 0.058 (5) 0.043 (5) 0.031 (5) 0.010 (5) 0.015 (4) N1 0.127 (8) 0.079 (6) 0.041 (5) 0.051 (6) 0.001 (5) −0.011 (4) N5 0.046 (5) 0.066 (5) 0.069 (5) 0.011 (4) −0.007 (4) −0.015 (4) C1 0.059 (6) 0.051 (6) 0.032 (4) 0.006 (5) −0.004 (4) 0.001 (4) C3 0.045 (5) 0.047 (5) 0.041 (5) −0.008 (4) −0.008 (4) 0.008 (4) N10 0.142 (9) 0.067 (6) 0.067 (6) 0.016 (6) −0.021 (6) 0.017 (5) N2 0.145 (9) 0.056 (5) 0.035 (4) 0.012 (6) −0.002 (5) 0.004 (4) C4 0.062 (6) 0.034 (5) 0.049 (5) 0.004 (4) −0.005 (4) 0.005 (4) N12 0.114 (7) 0.051 (5) 0.051 (5) −0.006 (5) 0.030 (5) 0.003 (4) C2 0.059 (6) 0.043 (5) 0.037 (5) 0.009 (4) 0.002 (4) −0.002 (4) N11 0.147 (10) 0.061 (6) 0.085 (7) 0.039 (6) 0.052 (7) 0.009 (5)

Table S4.18. Table of hydrogen bonds for [C(NH2)3]2PbBr4

D-H d(D-H) d(H..A)

N9-H9A 0.86 3.055 144.19 3.786 Br3 [ -x+1, -y+1, -z ]

N9-H9A 0.86 2.983 125.03 3.548 Br5 [ -x+2, -y+1, -z ]

N9-H9B 0.86 2.662 160.33 3.484 Br2 [ x+1, y-1, z ]

N8-H8A 0.86 2.61 162.74 3.44 Br3 [ -x+1, -y+1, -z ]

N8-H8B 0.86 2.61 156.57 3.416 Br2 [ x, y-1, z ]

N7-H7A 0.86 2.897 150.32 3.669 Br2 [ x+1, y-1, z ]

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N7-H7A 0.86 2.98 125.45 3.549 Br1 [ x+1, y-1, z ]

N7-H7B 0.86 2.682 166.6 3.524 Br4 [ x, y-1, z ]

N6-H6A 0.86 2.702 160.04 3.522 Br7

N6-H6B 0.86 3.098 123.81 3.647 Br3 [ -x+1, -y+1, -z ]

N6-H6B 0.86 2.956 146.81 3.705 Br5 [ -x+1, -y+1, -z ]

N4-H4A 0.86 2.973 126.36 3.551 Br8

N4-H4A 0.86 2.915 150.39 3.687 Br7

N4-H4B 0.86 2.64 164.31 3.475 Br6 [ x-1, y, z ]

N3-H3A 0.86 2.579 163.44 3.412 Br4

N3-H3B 0.86 2.993 134.72 3.649 Br1 [ -x, -y+2, -z+1 ]

N3-H3B 0.86 3.01 130.48 3.628 Br4 [ -x+1, -y+2, -z+1 ]

N1-H1A 0.86 2.631 166.56 3.473 Br6 [ x-1, y+1, z ]

N1-H1B 0.86 2.706 147.75 3.464 Br6 [ -x+1, -y+2, -z+1 ]

N5-H5A 0.86 2.605 161.96 3.433 Br5 [ -x+1, -y+1, -z ]

N5-H5B 0.86 2.638 160.23 3.46 Br7 [ x-1, y, z ]

N10-H10B 0.86 2.803 154.38 3.597 Br1 [ x+1, y-1, z ]

N2-H2B 0.86 3.038 126.94 3.621 Br2

N2-H2B 0.86 2.908 148.64 3.67 Br4

N12-H12A 0.86 2.788 168.32 3.634 Br7

N12-H12B 0.86 2.725 157.72 3.536 Br1 [ x+1, y-1, z ]

N11-H11A 0.86 2.571 176.18 3.43 Br8 [ -x+1, -y+1, -z+1 ]

N11-H11B 0.86 2.637 146.71 3.389 Br8

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Table S4.19. Fractional atomic coordinates and isotropic or equivalent isotropic displacement 2 parameters (Å ) for CsPb2Br5.

x y z Uiso*/Ueq Pb1 0.33061 (3) 0.83061 (3) 0.5000 0.0321 (3) Cs1 0.5000 0.5000 0.7500 0.0317 (3) Br2 0.5000 0.5000 0.5000 0.0258 (4) Br1 0.15393 (6) 0.65393 (6) 0.62911 (6) 0.0296 (3)

2 Table S4.20. Atomic displacement parameters (Å ) for CsPb2Br5.

U11 U22 U33 U12 U13 U23 Pb1 0.0288 (3) 0.0288 (3) 0.0386 (4) −0.00253 (12) 0.000 0.000 Cs1 0.0350 (4) 0.0350 (4) 0.0251 (5) 0.000 0.000 0.000 Br2 0.0207 (4) 0.0207 (4) 0.0360 (8) 0.000 0.000 0.000 Br1 0.0298 (3) 0.0298 (3) 0.0292 (5) −0.0001 (2) 0.00206 (19) 0.00206 (19)

Figure S4.9. Absorbance spectra of I-X. Absorbance was calculated from reflectance (R) and transmittance (T), A = -log (T+R). R and T were measured from the thin layer of powder dispersed in a transparent Teflon grease on a quartz substrate.

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Appendix to Chapter 5 Table S5.1. Fractional atomic coordinates and isotropic or equivalent isotropic displacement parameters (Å2).

x y Z U */U Occ. (<1) iso eq Pb1 0.2500 0.2500 0.5000 0.03861 (13)

Pb2 0.36972 (2) 0.24941 (3) 0.17333 (3) 0.03893 (11)

I1 0.14363 (3) 0.27884 (7) 0.31678 (6) 0.0661 (2)

I2 0.22996 (5) 0.0000 0.48564 (10) 0.0679 (3)

I3 0.31523 (3) 0.23463 (7) 0.35085 (6) 0.0633 (2)

I4 0.36981 (5) 0.0000 0.13945 (9) 0.0706 (3)

I5 0.47971 (3) 0.21806 (6) 0.33598 (6) 0.0572 (2)

I6 0.41596 (3) 0.27989 (7) 0.00697 (6) 0.0656 (2)

I7 0.2500 0.2500 0.0000 0.0810 (4)

I8 0.37575 (5) 0.5000 0.20851 (8) 0.0604 (3) N9 0.4220 (6) 1.013 (2) 0.9076 (12) 0.094 (4)* 0.5 H9A 0.4115 0.9666 0.9400 0.113* 0.5 H9B 0.4085 1.0744 0.8998 0.113* 0.5 N7 0.4747 (8) 1.0602 (13) 0.8174 (15) 0.094 (4)* 0.5 H7A 0.4982 1.0445 0.7916 0.113* 0.5 H7B 0.4612 1.1217 0.8095 0.113* 0.5 N8 0.4802 (8) 0.8941 (14) 0.8815 (14) 0.094 (4)* 0.5 H8A 0.5037 0.8784 0.8557 0.113* 0.5 H8B 0.4703 0.8482 0.9151 0.113* 0.5 C4 0.4593 (4) 0.9892 (13) 0.8694 (9) 0.094 (4)* 0.5 N14 0.195 (2) 0.591 (4) 0.145 (3) 0.094 (8)* 0.27 (2) H14A 0.1950 0.6349 0.1902 0.113* 0.27 (2) H14B 0.1824 0.6077 0.0841 0.113* 0.27 (2) C6 0.216 (2) 0.500 (3) 0.170 (3) 0.107 (11)* 0.27 (2) H6 0.2161 0.4527 0.1216 0.128* 0.27 (2) N13 0.2352 (14) 0.474 (4) 0.261 (3) 0.094 (8)* 0.27 (2) H13A 0.2349 0.5179 0.3066 0.113* 0.27 (2) H13B 0.2487 0.4132 0.2776 0.113* 0.27 (2)

N5 0.4878 (5) 0.5000 0.6087 (9) 0.062 (4) H5A 0.5016 0.4422 0.6356 0.075* 0.5 H5B 0.5043 0.5578 0.6279 0.075* 0.5

N6 0.4161 (4) 0.4134 (8) 0.5081 (8) 0.081 (3)

H6A 0.4293 0.3549 0.5342 0.097*

H6B 0.3858 0.4149 0.4618 0.097*

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C3 0.4419 (6) 0.5000 0.5386 (11) 0.048 (3) N3 0.4250 (7) 0.989 (3) 0.4367 (12) 0.094 (7)* 0.5 H3A 0.4235 0.9216 0.4328 0.113* 0.5 H3B 0.4373 1.0237 0.3986 0.113* 0.5 N4 0.3901 (6) 0.984 (3) 0.5564 (11) 0.092 (7)* 0.5 H4A 0.3886 0.9174 0.5529 0.110* 0.5 H4B 0.3793 1.0167 0.5982 0.110* 0.5 C2 0.4087 (9) 1.0360 (18) 0.4993 (16) 0.072 (7)* 0.5 H2 0.4103 1.1085 0.5033 0.087* 0.5 N1 0.7385 (12) 0.469 (3) 0.220 (2) 0.176 (15)* 0.5 H1A 0.7415 0.4067 0.2011 0.211* 0.5 H1B 0.7557 0.4875 0.2801 0.211* 0.5 N2 0.6823 (9) 0.508 (4) 0.0711 (17) 0.173 (11)* 0.5 H2A 0.6851 0.4458 0.0512 0.208* 0.5 H2B 0.6620 0.5525 0.0310 0.208* 0.5 C1 0.7082 (10) 0.535 (2) 0.1602 (18) 0.081 (8)* 0.5 H1 0.7050 0.6021 0.1814 0.098* 0.5 N12 0.5242 (7) 0.4326 (12) 0.1525 (14) 0.084 (3)* 0.5 H12A 0.4923 0.4454 0.1152 0.101* 0.5 H12B 0.5364 0.3701 0.1572 0.101* 0.5 N10 0.6045 (5) 0.490 (2) 0.2595 (10) 0.084 (3)* 0.5 H10A 0.6244 0.5401 0.2912 0.101* 0.5 H10B 0.6167 0.4276 0.2641 0.101* 0.5 N11 0.5359 (8) 0.6064 (13) 0.1962 (14) 0.084 (3)* 0.5 H11A 0.5040 0.6193 0.1589 0.101* 0.5 H11B 0.5556 0.6561 0.2290 0.101* 0.5 C5 0.5547 (4) 0.5097 (13) 0.2033 (8) 0.084 (3)* 0.5 N15 0.224 (2) 0.569 (4) 0.204 (5) 0.094 (8)* 0.23 (2) H15A 0.2240 0.5440 0.2590 0.113* 0.23 (2) H15B 0.2260 0.6359 0.1971 0.113* 0.23 (2) N16 0.219 (3) 0.409 (4) 0.141 (4) 0.094 (8)* 0.23 (2) H16A 0.2186 0.3836 0.1964 0.113* 0.23 (2) H16B 0.2171 0.3689 0.0929 0.113* 0.23 (2) C7 0.222 (2) 0.508 (4) 0.132 (4) 0.107 (11)* 0.23 (2) H7 0.2218 0.5362 0.0723 0.128* 0.23 (2)

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Table S5.2. Atomic displacement parameters (Å2).

U11 U22 U33 U12 U13 U23 Pb1 0.0392 (3) 0.0382 (3) 0.0419 (3) −0.0003 (2) 0.0183 (2) 0.0007 (2) Pb2 0.0367 (2) 0.0404 (2) 0.0406 (2) −0.00008 (15) 0.01428 (16) 0.00074 (15) I1 0.0523 (4) 0.0724 (5) 0.0594 (5) 0.0004 (4) 0.0008 (4) 0.0041 (4) I2 0.0647 (7) 0.0298 (5) 0.1045 (9) 0.000 0.0228 (6) 0.000 I3 0.0643 (5) 0.0802 (5) 0.0607 (5) −0.0091 (4) 0.0412 (4) −0.0086 (4) I4 0.0891 (9) 0.0299 (5) 0.0812 (8) 0.000 0.0138 (7) 0.000 I5 0.0431 (4) 0.0581 (4) 0.0591 (4) −0.0016 (3) 0.0028 (3) 0.0007 (3) I6 0.0687 (5) 0.0858 (6) 0.0544 (4) −0.0052 (4) 0.0366 (4) −0.0036 (4) I7 0.0403 (6) 0.0951 (9) 0.0885 (9) 0.0006 (6) −0.0026 (6) 0.0145 (7) I8 0.0729 (7) 0.0315 (5) 0.0668 (7) 0.000 0.0108 (5) 0.000 N5 0.054 (8) 0.061 (8) 0.057 (8) 0.000 0.000 (7) 0.000 N6 0.071 (7) 0.057 (6) 0.104 (9) −0.012 (5) 0.015 (6) −0.006 (6) C3 0.044 (8) 0.050 (8) 0.046 (8) 0.000 0.010 (7) 0.000

2 Figure S5.1. Kubelka–Munk function F(R∞)=(1-R∞) /2R∞ (R∞ - diffusive reflectance), PL, PL excitation (PLexc) and photoconductivity (PC) spectra of FAGPbI4 (same as shown in Figure 5.6 of the Main Text). The measurements were performed on the powdered crystals synthesized with a method 1.

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Figure S5.2. Spectral dependence of PLQY for FAGPbI4. The measurements were performed on the powdered crystals synthesized with a method 1.

207 Figure S5.3. (a) Pb MAS NMR spectrum of α-FAPbI3 (black) showing two different signals. The spectra of α-FAPbI3 in static mode (blue) and of δ-FAPbI3 with MAS are plotted for comparison (red); (b,c) Powder XRD pattern of α-FAPbI3 before and after the MAS NMR measurements, compared to the calculated patterns of α-FAPbI3 (blue) and δ-FAPbI3 (red). The calculated pattern for α-FAPbI3 is based on the experimental single-crystal data reported by M. T. Weller et. al. [J. Phys. Chem. Lett., 2015, 6, 3209-3212]. Both SSNMR and powder XRD methods show the formation of δ-FAPbI3 when spinning the sample during NMR experiment. Hence only the static NMR data are truly relevant for comparison with other compounds.

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Table S5.3. 207Pb SSNMR chemical shifts (referenced to PbMe4) and FWHM of the measured compounds under MAS or, in the case of α-FAPbI3, static mode.

Compound Chemical shift FWHM Scans MAS (ppm) (kHz) (kHz)

FAGPbI4 1160, 1465 22.4, 20.2 786’432 20

α-FAPbI3 1515 22.2 51’454 0

δ-FAPbI3 1175 21.8 117’702 20

G2PbI4 1515 24.5 131’072 20

Figure S5.4. The powder XRD pattern of G2PbI4. The calculated pattern is based on the experimental single-crystal XRD data reported by M. Szafranski and A. Katrusiak [Phys. Rev. 2000, B61 (2), 1026 and ICSD card 92045]. A preferred orientation of the microcrystalline powder is observed through the difference in intensity of certain reflections in the experimental pattern compared to the calculated data due to the fact that G2PbI4 crystallizes in the shape of thin plates, giving rise to preferred orientation on a flat sample holder.

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Figure S5.5. The powder XRD pattern of δ-FAPbI3 ([CH(NH2)2]2PbI4, P63mc). The calculated pattern is based on the experimental single-crystal XRD data reported by C. Stoumpous et. al. [Inorg. Chem. 2013, 52, 9019-9038 and ICSD card 250741].

207 Figure S5.6. Pb MAS NMR spectrum of δ-FAPbI3 showing a peak at 1175 ppm with a FWHM of 21.8 kHz. The isotropic signal and spinning sidebands were identified by varying the MAS frequency. The peak remaining at the same chemical shift was identified to be the isotropic signal whereas the spinning sidebands (marked by asterisk) are shifting according to the spinning speed. The spectrum was acquired with 117 702 scans. Fitted spinning sidebands and

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Appendix | Olga Nazarenko their sum are shown in blue. For the fit the separation between the individual Gaussian peaks was set to 20 kHz.

207 Figure S5.7. Pb MAS NMR spectra of FAGPbI4 spectral regions around 0 and -1300 ppm. The number of scans acquired were 5000 and 12 536, respectively. The recycle delay was set to 2 seconds to detect impurities with longer relaxation times. No signal could be detected in this spectral range.

Figure S5.8. Analysis of PL and TR-PL of a single crystal with micrometer spatial resolution.

In order to clarify the complexity of the emission spectrum, and, in particular, to check whether different emission bands originate from spatially-separated emitting centers, PL and TR-PL were performed on a home–built micro-photoluminescence setup. In these experiments, a FAGPbI4 single crystal was placed on thin (150 µm) glass coverslip and excited by a laser diode at 405 nm (repetition rate = 2.5 MHz; excitation power density = 90 µJ/cm2). In most studied crystal areas, the PL spectrum is composed by different bands and the PL line shape

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Appendix | Olga Nazarenko does not strongly depend on the particular spatial location (Figure S5.8 a). TR-PL experiments (detection wavelength sets at 700 nm) reveal a modest change of the average lifetime for the different spatial locations (Figure S5.8 b). Given our experimental spatial resolution of a few µm, it is difficult to draw final conclusions concerning the origin of the different emission bands: whether they are associated with some local crystal distortions, free and/or STEs, or color centers.

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Appendix to Chapter 6 Table S6.1. Crystallographic data of the synthesized compounds.

[C(NH2)3]2SnBr4 Cs[C(NH2)3]SnBr4 Cs2[C(NH2)3]Sn2Br7 Formula weight 558.51 631.33 1122.66 Temperature (K) 298 250 300 Crystal system orthorhombic orthorhombic orthorhombic

Space group Pna21 Imma Cmmm Colour colorless orange red a (Å) 32.082(3) 6.0112(10) 5.9653(6) b (Å) 8.3226(6) 11.5363(18) 28.625(3) c (Å) 10.4238(8) 17.342(3) 11.5503(11) α=β=γ (°) 90 90 90 Volume (Å3) 2783.2(4) 1202.6(3) 1972.3(3) Z 8 4 4 3 ρcalc (g/cm ) 2.666 3.487 3.781 μ (mm-1) 13.296 18.348 20.333 F(000) 2048.0 1112.0 1952.0 Crystal size (mm3) 0.34 × 0.34 × 0.3 0.18 × 0.06 × 0.02 0.38 × 0.04 × 0.04 Radiation MoKα (λ = 0.71073 MoKα (λ = 0.71073 MoKα (λ = 0.71073 Å) Å) Å) 2Θ range for data 4.66 to 58.052 4.24 to 62.872 3.526 to 61.014 collection (°) Index ranges -43 ≤ h ≤ 43, -11 ≤ -8 ≤ h ≤ 8, -16 ≤ k ≤ -8 ≤ h ≤ 8, -39 ≤ k ≤ k ≤ 11, -14 ≤ l ≤ 14 16, -25 ≤ l ≤ 25 40, -16 ≤ l ≤ 16 Reflections collected 29756 7135 11738

Independent reflections 7393 [Rint = 0.0477, 1079 [Rint = 0.0361, 1734 [Rint = 0.0341, Rsigma = 0.0444] Rsigma = 0.0247] Rsigma = 0.0243] Data/restraints/parameters 7393/1/237 1079/0/46 1734/0/63 Goodness-of-fit on F2 1.013 1.032 1.045

Final R indexes [I>=2σ R1 = 0.0348, wR2 = R1 = 0.0227, wR2 = R1 = 0.0258, wR2 = (I)] 0.0674 0.0511 0.0584

Final R indexes [all data] R1 = 0.0574, wR2 = R1 = 0.0279, wR2 = R1 = 0.0336, wR2 = 0.0737 0.0529 0.0616 Largest diff. peak/hole 0.52/-0.55 0.67/-1.01 0.74/-0.93 (e Å-3)

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Figure S6.1. An examination of the reconstructed reciprocal layer hk0 (a* to the right, b* up) of compound XI shows that only very few, weak reflections are responsible for the violation of a b-glide plane perpendicular to the c-axis in the crystal structure of G2SnBr4. The corresponding lines in the diffraction pattern containing the weak reflections are marked with black arrows.

Accordingly, during the refinement in the space group Pna21 (no. 33), large correlation matrix elements arise for parameters of atoms related by the pseudo-b-glide plane perpendicular to the c-axis (larger than 0.8 between displacement parameters of heavy atoms). Therefore, a refinement in the higher symmetric space group Pbcn (no. 60, which is the standard setting of Pnab, obtained when maintaining the setting given by the space group Pna21) has been tested. However, the second parameter of the empirical weighting scheme refines to an unusually high value of more than 20 and one of the two crystallographically independent guanidinium molecules had to be described as disordered over two positions. Therefore, for the final refinement, the space group Pna21 has been chosen. A refinement as an inversion twin yields a contribution of 10(2) % of an inverted domain, which is just slightly above the significance criterion of 3*sigma. There are also very few reflections originating from a small, randomly oriented crystal. The weak reflections 4 1 0, 3 1 0 were larger observed than calculated with an error/esd larger than 8 times the standard uncertainty. It was assumed that these reflections are affected by a contribution from reflections of the small randomly oriented domain, and were subsequently excluded from the refinement. The reflection 2 0 0, has also been omitted from the refinement since its intensity is biased by the beamstop.

Table S6.2. Table of hydrogen bonds for G2SnBr4.

D-H d(D-H) d(H···A) D-H···A d(D···A) A N1-H1A 0.86 2.599 162.24 3.428 Br3 [x, y+1, z] N1-H1B 0.86 2.61 163.61 3.443 Br5 [-x+1, -y+1, z+1/2] N2-H2A 0.86 2.67 165.46 3.509 Br3 N2-H2B 0.86 2.981 127.18 3.567 Br2 [-x+1, -y, z+1/2]

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N2-H2B 0.86 3.122 143.53 3.847 Br5 [-x+1, -y+1, z+1/2] N3-H3A 0.86 3.051 148.07 3.808 Br3 N3-H3A 0.86 2.893 131.9 3.526 Br4 N3-H3B 0.86 2.715 163.71 3.549 Br1 [x, y+1, z] N4-H4A 0.86 2.573 163.74 3.407 Br4 [-x+3/2, y+1/2, z-1/2] N4-H4B 0.86 2.64 153.18 3.43 Br7 N5-H5A 0.86 2.645 163.34 3.478 Br7 [-x+3/2, y-1/2, z+1/2] N5-H5B 0.86 3.138 143.74 3.865 Br1 [x, y+1, z] N5-H5B 0.86 2.895 120.51 3.413 Br4 N6-H6A 0.86 2.866 153.3 3.655 Br7 [-x+3/2, y-1/2, z+1/2] N6-H6B 0.86 2.909 158.03 3.72 Br3 [-x+3/2, y+1/2, z-1/2] N7-H7A 0.86 2.83 135.56 3.496 Br8 N7-H7B 0.86 2.623 162.74 3.454 Br1 [-x+3/2, y+1/2, z-1/2] N8-H8A 0.86 2.841 141.15 3.552 Br1 N8-H8B 0.86 2.586 165.59 3.426 Br8 [-x+3/2, y-1/2, z+1/2] N9-H9A 0.86 2.931 148.38 3.691 Br1 [-x+3/2, y+1/2, z-1/2] N9-H9B 0.86 3.003 135.33 3.665 Br6 [-x+3/2, y-1/2, z+1/2] N9-H9B 0.86 3.016 146.75 3.765 Br8 [-x+3/2, y-1/2, z+1/2] N10- 0.86 2.659 170.17 3.509 Br8 H10A N10- 0.86 2.938 148.9 3.701 Br6 [x, y-1, z] H10B N10- 0.86 2.908 127.4 3.498 Br7 [x, y-1, z] H10B N11- 0.86 2.561 164.73 3.398 Br2 [-x+1, -y, z-1/2] H11A N11- 0.86 2.667 156.93 3.475 Br6 H11B N12- 0.86 2.991 145.96 3.735 Br2 [-x+1, -y, z-1/2] H12A N12- 0.86 2.697 159.77 3.516 Br6 [x, y-1, z] H12B

Table S6.3. Fractional atomic coordinates and isotropic or equivalent isotropic displacement 2 parameters (Å ) for G2SnBr4.

x y z Uiso*/Ueq Sn1 0.62067 (2) 0.04517 (8) 0.41578 (9) 0.0385 (3) Sn2 0.62186 (2) 0.54671 (7) 0.25611 (9) 0.0380 (3) Br1 0.66976 (4) −0.21157 (13) 0.53744 (17) 0.0513 (4) Br2 0.55581 (3) −0.20011 (16) 0.34115 (18) 0.0479 (3)

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Br3 0.58258 (4) 0.06514 (12) 0.64880 (16) 0.0486 (4) Br4 0.67370 (4) 0.29787 (14) 0.5100 (2) 0.0575 (4) Br5 0.55849 (3) 0.30095 (16) 0.33406 (18) 0.0468 (3) Br6 0.58170 (5) 0.57547 (13) 0.02709 (16) 0.0473 (4) Br7 0.67658 (4) 0.80806 (13) 0.16236 (18) 0.0534 (4) Br8 0.66722 (4) 0.29759 (13) 0.13147 (18) 0.0530 (4) N1 0.5373 (3) 0.6949 (11) 0.6849 (13) 0.066 (4) H1A 0.5535 0.7748 0.6700 0.079* H1B 0.5121 0.7115 0.7101 0.079* N2 0.5271 (3) 0.4231 (11) 0.6990 (16) 0.059 (4) H2A 0.5367 0.3268 0.6935 0.071* H2B 0.5020 0.4393 0.7242 0.071* N3 0.5890 (3) 0.5215 (13) 0.6307 (13) 0.055 (3) H3A 0.5986 0.4253 0.6251 0.066* H3B 0.6046 0.6019 0.6110 0.066* C1 0.5510 (3) 0.5453 (12) 0.6688 (16) 0.043 (4) N4 0.7680 (4) 0.6207 (15) 0.2437 (11) 0.099 (4) H4A 0.7834 0.6449 0.1788 0.119* H4B 0.7420 0.6443 0.2434 0.119* N5 0.7616 (3) 0.5126 (14) 0.4422 (11) 0.087 (4) H5A 0.7724 0.4659 0.5079 0.105* H5B 0.7355 0.5364 0.4418 0.105* N6 0.8229 (3) 0.5090 (14) 0.3506 (15) 0.084 (4) H6A 0.8324 0.4624 0.4181 0.101* H6B 0.8393 0.5298 0.2874 0.101* C2 0.7846 (4) 0.5471 (15) 0.3441 (18) 0.050 (4) N7 0.7592 (3) 0.1298 (14) 0.2517 (10) 0.081 (3) H7A 0.7324 0.1259 0.2526 0.097* H7B 0.7719 0.1778 0.1900 0.097* N8 0.7620 (3) −0.0105 (14) 0.4389 (10) 0.086 (3) H8A 0.7353 −0.0157 0.4416 0.103* H8B 0.7766 −0.0537 0.4988 0.103* N9 0.8213 (3) 0.0635 (12) 0.3414 (13) 0.062 (3) H9A 0.8343 0.1080 0.2786 0.075* H9B 0.8351 0.0192 0.4026 0.075* C3 0.7805 (4) 0.0632 (14) 0.3438 (16) 0.045 (3) N10 0.5901 (3) 0.0184 (12) 0.0463 (15) 0.054 (4) H10A 0.6072 0.0956 0.0611 0.065* H10B 0.5973 −0.0793 0.0613 0.065*

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N11 0.5405 (3) 0.1945 (10) −0.0232 (11) 0.052 (3) H11A 0.5154 0.2128 −0.0484 0.063* H11B 0.5577 0.2729 −0.0138 0.063* N12 0.5256 (3) −0.0655 (10) −0.0150 (18) 0.064 (4) H12A 0.5006 −0.0443 −0.0402 0.077* H12B 0.5328 −0.1632 0.0001 0.077* C4 0.5523 (4) 0.0501 (11) 0.0002 (14) 0.041 (3)

2 Table S6.4. Atomic displacement parameters (Å ) for G2SnBr4.

U11 U22 U33 U12 U13 U23 Sn1 0.0375 (5) 0.0415 (5) 0.0366 (7) −0.0018 (3) −0.0004 (4) 0.0019 (4) Sn2 0.0391 (5) 0.0391 (5) 0.0358 (7) −0.0013 (3) −0.0008 (5) 0.0004 (4) Br1 0.0486 (7) 0.0495 (7) 0.0557 (11) 0.0053 (5) −0.0065 (7) 0.0091 (7) Br2 0.0412 (6) 0.0502 (7) 0.0523 (9) −0.0040 (6) 0.0064 (8) −0.0125 (6) Br3 0.0595 (9) 0.0478 (7) 0.0385 (10) −0.0008 (5) 0.0089 (8) 0.0021 (6) Br4 0.0448 (7) 0.0559 (7) 0.0720 (13) −0.0115 (5) −0.0053 (7) −0.0017 (8) Br5 0.0389 (6) 0.0484 (7) 0.0531 (9) −0.0035 (5) −0.0049 (7) 0.0123 (6) Br6 0.0609 (9) 0.0420 (6) 0.0391 (9) 0.0040 (5) −0.0124 (8) −0.0006 (7) Br7 0.0470 (7) 0.0543 (7) 0.0588 (11) −0.0070 (5) 0.0040 (7) 0.0129 (7) Br8 0.0527 (8) 0.0473 (7) 0.0591 (12) 0.0126 (5) 0.0114 (7) −0.0017 (7) N1 0.058 (6) 0.048 (6) 0.091 (11) 0.007 (5) 0.026 (7) 0.013 (6) N2 0.047 (6) 0.049 (5) 0.082 (11) −0.016 (4) 0.007 (7) 0.003 (6) N3 0.043 (6) 0.077 (7) 0.045 (9) −0.005 (5) 0.009 (6) −0.005 (7) C1 0.033 (6) 0.057 (8) 0.040 (9) 0.001 (5) −0.019 (6) 0.005 (6) N4 0.086 (8) 0.149 (10) 0.062 (7) 0.043 (8) 0.006 (6) 0.033 (8) N5 0.061 (6) 0.144 (10) 0.058 (7) −0.019 (7) 0.008 (5) 0.022 (7) N6 0.059 (7) 0.146 (10) 0.046 (7) 0.027 (6) −0.004 (6) −0.001 (8) C2 0.043 (7) 0.067 (9) 0.039 (8) −0.004 (5) −0.013 (6) 0.007 (6) N7 0.048 (5) 0.133 (9) 0.061 (7) −0.004 (6) −0.008 (5) 0.039 (7) N8 0.067 (6) 0.132 (9) 0.060 (7) 0.010 (7) 0.018 (5) 0.036 (7) N9 0.038 (5) 0.113 (9) 0.036 (6) −0.005 (5) −0.003 (4) 0.013 (5) C3 0.041 (6) 0.058 (8) 0.037 (8) 0.002 (5) −0.006 (6) 0.001 (6) N10 0.034 (6) 0.056 (6) 0.073 (11) 0.001 (5) −0.008 (7) 0.000 (7) N11 0.061 (6) 0.043 (5) 0.053 (8) 0.000 (4) −0.009 (6) −0.003 (5) N12 0.047 (6) 0.049 (6) 0.096 (12) 0.003 (4) −0.015 (7) −0.007 (7) C4 0.049 (7) 0.043 (7) 0.030 (8) −0.001 (5) −0.018 (6) 0.000 (6)

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Table S6.5. Crystallographic data of Cs8Sn6Br13I7.

Cs8Sn6Br13I7 Formula weight 3702.55 Temperature (K) 250 Crystal system orthorhombic Space group Cmcm Colour orange a (Å) 5.9758(8) b (Å) 26.947(4) c (Å) 17.164(2) α=β=γ (°) 90 Volume (Å3) 2763.9(6) Z 2 3 ρcalc (g/cm ) 4.449 μ (mm-1) 21.174 F(000) 3132.0 Crystal size (mm3) 0.44 × 0.12 × 0.02 Radiation MoKα (λ = 0.71073 Å) 2Θ range for data collection (°) 3.844 to 63.176 Index ranges -8 ≤ h ≤ 8, -39 ≤ k ≤ 38, -24 ≤ l ≤ 25 Reflections collected 16164

Independent reflections 2428 [Rint = 0.0436, Rsigma = 0.0335] Data/restraints/parameters 2428/0/58 Goodness-of-fit on F2 1.092

Final R indexes [I>=2σ (I)] R1 = 0.0302, wR2 = 0.0684

Final R indexes [all data] R1 = 0.0365, wR2 = 0.0713 Largest diff. peak/hole (e Å-3) 1.45/-1.64

Table S6.6. Fractional atomic coordinates and isotropic or equivalent isotropic displacement 2 parameters (Å ) for Cs8Sn6Br13I7.

x y z Uiso*/Ueq Occ. (<1) Cs1 0.5000 0.52941 (2) 0.7500 0.04154 (14) Cs2 0.5000 0.67763 (2) 0.95953 (2) 0.04094 (11) Cs3 1.0000 0.84818 (2) 0.7500 0.04312 (15) I1 1.0000 0.57633 (2) 0.60665 (2) 0.03999 (12) I2 1.0000 0.66472 (2) 0.40086 (3) 0.03593 (12) 0.5 I3 0.5000 0.55974 (3) 0.41881 (4) 0.05076 (17) 0.25 Sn1 1.0000 0.56050 (2) 0.42345 (2) 0.02954 (10)

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Sn2 1.0000 0.66629 (2) 0.7500 0.02340 (11) Br1 1.0000 0.55311 (3) 0.2500 0.03254 (18) Br2 0.5000 0.55974 (3) 0.41881 (4) 0.05076 (17) 0.75 Br3 1.0000 0.66472 (2) 0.40086 (3) 0.03593 (12) 0.5 Br4 0.5000 0.66403 (4) 0.7500 0.0409 (2) Br5 1.0000 0.73517 (2) 0.86826 (4) 0.03332 (14)

2 Table S6.7. Atomic displacement parameters (Å ) for Cs8Sn6Br13I7.

U11 U22 U33 U12 U13 U23 Cs1 0.0395 (3) 0.0341 (3) 0.0510 (3) 0.000 0.000 0.000 Cs2 0.0380 (2) 0.0542 (3) 0.03063 (19) 0.000 0.000 −0.00028 (18) Cs3 0.0602 (4) 0.0349 (3) 0.0343 (3) 0.000 0.000 0.000 I1 0.0539 (3) 0.0362 (2) 0.0299 (2) 0.000 0.000 0.00027 (16) I2 0.0334 (2) 0.0347 (2) 0.0397 (3) 0.000 0.000 −0.0037 (2) I3 0.0363 (3) 0.0591 (4) 0.0569 (4) 0.000 0.000 −0.0158 (3) Sn1 0.0308 (2) 0.0283 (2) 0.0295 (2) 0.000 0.000 −0.00040 (15) Sn2 0.0220 (2) 0.0246 (2) 0.0236 (2) 0.000 0.000 0.000 Br1 0.0361 (4) 0.0383 (4) 0.0231 (4) 0.000 0.000 0.000 Br2 0.0363 (3) 0.0591 (4) 0.0569 (4) 0.000 0.000 −0.0158 (3) Br3 0.0334 (2) 0.0347 (2) 0.0397 (3) 0.000 0.000 −0.0037 (2) Br4 0.0454 (5) 0.0478 (5) 0.0296 (4) 0.000 0.000 0.000 Br5 0.0424 (3) 0.0274 (3) 0.0302 (3) 0.000 0.000 −0.0037 (2)

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Figure S6.2. The powder XRD pattern of a powdered sample of guanidinium tin bromide -

G2SnBr4 (XI). The calculated pattern is based on the experimental single-crystal XRD data obtained in this work (CCDC number 1854819). A preferred orientation of the microcrystalline powder is observed through the difference in intensity of certain reflections in the experimental pattern comparing to the calculated data due to the fact that XI crystallize in the shape of thin needles, giving rise to a preferred orientation in a sample holder: 0.3 mm glass capillary.

Figure S6.3. The powder XRD pattern of CsSnBr3. The calculated pattern is based on the experimental single-crystal XRD data reported by Donaldson J. D. et. al. [J. Chem. Soc., Dalton Trans., 1975, 0, 1500-1506; cif file was downloaded from ICSD database, ICSD card 4071]. A background, obtained during the experimental measurement of powder XRD of

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CsSnBr3 sample is attributed to a presence of the amorphous CsSnBr3 phase that is also observed as a shoulder in the PL spectra of CsSnBr3 (Figure 6.1d, Main Text). A preferred orientation of the microcrystalline powder is observed through the difference in intensity of certain reflections in the experimental pattern comparing to the calculated data due to preferred orientation of microcrystalline powder in a sample holder: 0.3 mm glass capillary.

Figure S6.4. The powder XRD pattern of a powdered sample of cesium guanidinium tin bromide - Cs2GSn2Br7 (XIII). For the measurement the sample was deposited between adhesive tape that contributed to the background. In the inset a selected area of 2Ɵ angles is presented to demonstrate the reflections assigned to CsGSnBr4 (XII) phase. The calculated patterns of

XIII as well as of XII are based on the experimental single-crystal XRD data obtained in this work (CCDC codes 1854833 and 1854838 correspondingly). The reflection appointed with an asterisk (*) seemingly results from the Sn (II) hydrolyzation product. The reflection matches (being shifted by ~ 0.1 deg to lower angles) to the most intense reflection in the powder XRD pattern of red modification of SnO (reported by Donaldson J. D. et. al., Acta Crystallogr., 1963, 16, 22; ICSD card 60619).

Table S6.8. Table of hydrogen bonds for Cs2GSn2Br7.

D-H d(D-H) d(H···A) D-H···A d(D···A) A N1-H1A 0.860 2.794 140.60 3.502 Br5 [ -x+1, -y+1, z ] N1-H1B 0.860 2.888 155.99 3.690 Br2 [ x-1/2, y+1/2, z ] N2-H2A 0.860 2.819 158.88 3.634 Br2 [ x-1/2, y+1/2, z ] N2-H2B 0.860 2.819 158.88 3.634 Br2 [ x-1/2, y+1/2, -z ]

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Table S6.9. Table of hydrogen bonds for CsGSnBr4.

D-H d(D-H) d(H···A) D-H···A d(D···A) A N1-H1A 0.860 2.800 158.77 3.616 Br1 [ x-1/2, y, -z+3/2 ] N1-H1B 0.860 2.800 158.77 3.616 Br1 [ -x+3/2, -y+1/2, - z+3/2 ] N2-H2A 0.860 2.721 142.41 3.443 Br2 N2-H2B 0.860 2.859 156.13 3.662 Br1 [ x-1/2, y, -z+3/2 ]

Table S6.10. Fractional atomic coordinates and isotropic or equivalent isotropic displacement 2 parameters (Å ) for Cs2GSn2Br7.

x y z U */U Occ. (<1) iso eq Cs1 0.0000 0.31517 (2) 0.5000 0.05061 (13)

Cs2 0.5000 0.0000 0.5000 0.0701 (2)

Cs3 0.5000 0.0000 0.0000 0.0754 (2)

Sn1 0.0000 0.10131 (2) 0.25200 (2) 0.02658 (9)

Br1 0.0000 0.08243 (3) 0.5000 0.0558 (2)

Br2 0.0000 0.20163 (2) 0.30561 (4) 0.04836 (13) Br3A 0.064 (9) 0.1137 (5) 0.0000 0.051 (6) 0.13 (7) Br3B 0.0000 0.11446 (18) 0.0000 0.076 (4) 0.74 (13)

Br4 0.0000 0.0000 0.21000 (9) 0.0622 (2)

Br5 0.0000 0.40416 (2) 0.25728 (5) 0.05620 (16)

N1 0.0000 0.29988 (17) 0.0988 (4) 0.0585 (12)

H1A 0.0000 0.3299 0.1002 0.070*

H1B 0.0000 0.2844 0.1626 0.070*

N2 0.0000 0.2319 (2) 0.0000 0.0560 (16) H2A 0.0000 0.2168 0.0645 0.067* 0.5 H2B 0.0000 0.2168 −0.0645 0.067* 0.5

C1 0.0000 0.2779 (2) 0.0000 0.0326 (12)

2 Table S6.11. Atomic displacement parameters (Å ) for Cs2GSn2Br7.

U11 U22 U33 U12 U13 U23 Cs1 0.0514 (3) 0.0576 (3) 0.0428 (2) 0.000 0.000 0.000 Cs2 0.0706 (5) 0.0411 (4) 0.0986 (6) 0.000 0.000 0.000 Cs3 0.0734 (5) 0.0992 (7) 0.0537 (4) 0.000 0.000 0.000 Sn1 0.02649 (14) 0.02852 (15) 0.02473 (14) 0.000 0.000 0.00078 (9) Br1 0.0713 (5) 0.0633 (5) 0.0327 (3) 0.000 0.000 0.000 Br2 0.0581 (3) 0.0376 (2) 0.0493 (3) 0.000 0.000 0.0009 (2)

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Br3A 0.086 (12) 0.027 (8) 0.040 (6) −0.012 (4) 0.000 0.000 Br3B 0.130 (14) 0.063 (5) 0.0341 (18) 0.000 0.000 0.000 Br4 0.0769 (5) 0.0327 (4) 0.0769 (5) 0.000 0.000 0.000 Br5 0.0471 (3) 0.0527 (3) 0.0688 (4) 0.000 0.000 −0.0029 (3) N1 0.058 (3) 0.066 (3) 0.052 (2) 0.000 0.000 −0.020 (2) N2 0.040 (3) 0.036 (3) 0.092 (5) 0.000 0.000 0.000 C1 0.022 (2) 0.040 (3) 0.036 (3) 0.000 0.000 0.000

Table S6.12. Fractional atomic coordinates and isotropic or equivalent isotropic displacement 2 parameters (Å ) for CsGSnBr4.

x y z U */U Occ. (<1) iso eq Cs1 0.0000 0.2500 0.64389 (2) 0.04059 (13)

Sn1 0.0000 0.0000 0.0000 0.01998 (10)

Br1 0.0000 0.55447 (4) 0.16707 (2) 0.03745 (12)

Br2 0.0000 0.0000 0.5000 0.04440 (17) Br3A 0.0000 0.2500 0.0219 (9) 0.048 (5) 0.56 (10) Br3B 0.054 (5) 0.2500 0.0302 (12) 0.0448 (14) 0.22 (5)

N1 0.0000 0.2500 0.2165 (3) 0.0381 (10) H1A 0.0000 0.1854 0.1917 0.046* 0.5 H1B 0.0000 0.3146 0.1917 0.046* 0.5

N2 0.0000 0.1503 (4) 0.3285 (2) 0.0503 (10)

H2A 0.0000 0.1485 0.3780 0.060*

H2B 0.0000 0.0867 0.3026 0.060*

C1 0.0000 0.2500 0.2923 (3) 0.0251 (9)

2 Table S6.13. Atomic displacement parameters (Å ) for CsGSnBr4.

U11 U22 U33 U12 U13 U23

Cs1 0.0401 (2) 0.0344 (2) 0.0473 (2) 0.000 0.000 0.000

Sn1 0.01912 (16) 0.01761 (15) 0.02319 (17) 0.000 0.000 −0.00134 (10)

Br1 0.0424 (3) 0.0371 (2) 0.0329 (2) 0.000 0.000 0.00420 (15)

Br2 0.0448 (4) 0.0585 (4) 0.0300 (3) 0.000 0.000 0.0046 (3)

Br3A 0.089 (15) 0.0299 (14) 0.026 (2) 0.000 0.000 0.000

Br3B 0.067 (4) 0.038 (2) 0.029 (3) 0.000 −0.010 (3) 0.000

N1 0.030 (2) 0.057 (3) 0.027 (2) 0.000 0.000 0.000

N2 0.048 (2) 0.049 (2) 0.054 (2) 0.000 0.000 0.0234 (17)

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C1 0.020 (2) 0.031 (2) 0.024 (2) 0.000 0.000 0.000

Figure S6.5. An electronic band diagram of CsSnBr3 computed by DFT with SOC for the tetragonal phase.

Figure S6.6. (a) Representation of the electronic densities of the conduction (red) and valence

(blue) edge states of (a) CsGSnBr4 (XII) and (b) Cs2GSn2Br7 (XIII).

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Table S6.14. Computed band gaps at the DFT+SOC level of theory.

G2SnBr4 CsGSnBr4 Cs2GSn2Br7 CsSnBr3 CsGPbBr4 Cs2GPb2Br7

n=1 n=2 n=∞ n=1 n=2 (XI) (XII) (XIII) Dimensionality 1D 2D 2D 3D 2D 2D Space group Pna2 Imma Cmmm P4/mbm Imma Cmmm 1 Band gaps 3.12 1.04 0.84 0.20 1.55 1.32 (eV)

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Appendix to Chapter 7 Table S7.1. Crystallographic data of XIV and XV.

XIV XV

Empirical formula Rb20Cu8.32Bi6I48.14O8 Rb20Ag8.42Bi6I47.49O8 Formula weight 9728.90 10026.03 Temperature (K) 100 100 Crystal system tetragonal orthorhombic Space group P4mm Cmmm Colour dark green red a (Å) 13.0717(3) 18.5360(7) b (Å) 13.0717(3) 18.7551(7) c (Å) 10.2949(5) 10.3272(4) α=β=γ (°) 90 90 Volume (Å3) 1759.08(12) 3590.2(2) Z 0.5 1 3 ρcalc (g/cm ) 4.592 4.637 μ (mm-1) 26.158 25.409 F(000) 2047.0 4215.0 Crystal size (mm3) 0.1 × 0.04 × 0.02 0.28 × 0.14 × 0.04 Radiation MoKα (λ = 0.71073 Å) MoKα (λ = 0.71073 Å) 2Θ range for data collection 3.116 to 62.698 3.09 to 63.058 (°) Index ranges -19 ≤ h ≤ 19, -18 ≤ k ≤ 18, -26 ≤ h ≤ 27, -27 ≤ k ≤ 26, -14 ≤ l ≤ 14 -14 ≤ l ≤ 15 Reflections collected 20445 20563

Independent reflections 3096 [Rint = 0.0510, Rsigma = 3142 [Rint = 0.0842, Rsigma = 0.0253] 0.0447] Data/restraints/parameters 3096/1/111 3142/0/100 Goodness-of-fit on F2 1.220 1.079

Final R indexes [I>=2σ (I)] R1 = 0.0435, wR2 = 0.0804 R1 = 0.0645, wR2 = 0.1318

Final R indexes [all data] R1 = 0.0518, wR2 = 0.0842 R1 = 0.0835, wR2 = 0.1433 Largest diff. peak/hole 2.80/-1.70 4.71/-3.12 (e Å-3) Flack parameter 0.50(3)

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Table S7.2. Fractional atomic coordinates and isotropic or equivalent isotropic displacement parameters (Å2) for XIV.

X y z U */U Occ. (<1) iso eq Bi0A 0.5000 0.5000 0.8549 (4) 0.0125 (2)

Bi2 0.5000 1.0000 0.3563 (3) 0.01627 (18)

I6 0.5000 0.5000 0.1484 (4) 0.0101 (8)

I3 0.5000 0.76683 (7) 0.3570 (4) 0.0168 (2)

I2 0.5000 0.5000 0.5602 (5) 0.0187 (9)

I1 0.5000 0.73508 (7) 0.8565 (4) 0.0174 (2)

I7 0.16560 (6) 0.83440 (6) −0.1457 (4) 0.0230 (2) I8 0.0000 1.0000 0.1621 (7) 0.0108 (14)* 0.580 (16)

Rb3 0.2925 (3) 0.7075 (3) 0.1245 (4) 0.0289 (9)

I10 0.0000 1.0000 0.5433 (8) 0.061 (2) I4A 0.3185 (5) 1.0000 0.1667 (7) 0.0185 (19) 0.516 (18)

Rb2 0.7061 (2) 0.7061 (2) 0.5860 (3) 0.0194 (7) I5B 0.6935 (11) 1.0000 0.5374 (11) 0.025 (2) 0.42 (3) Rb1 0.5643 (4) 1.0000 0.8577 (12) 0.0411 (11) 0.5 Cu3 0.1636 (8) 1.0000 −0.0010 (10) 0.014 (3)* 0.32 (2) I4B 0.3557 (6) 1.0000 0.1118 (6) 0.0206 (17) 0.484 (18) I5A 0.6494 (11) 1.0000 0.5846 (11) 0.036 (4) 0.58 (3) I9 0.0000 1.0000 0.351 (7) 0.044 (7)* 0.157 (12) Cu1 0.0000 1.0000 −0.139 (8) 0.063 (8) 0.36 (3) I11 0.0000 1.0000 0.050 (2) 0.072 (7) 0.34 (2) Cu2 0.1079 (16) 1.0000 −0.0359 (18) 0.019 (6)* 0.156 (17)

O1 0.8027 (10) 0.8027 (10) 0.366 (5) 0.042 (6) Cu4 0.0000 1.1449 (12) −0.2885 (13) 0.050* 0.47 (2)

Table S7.3. Atomic displacement parameters (Å2) for XIV.

U11 U22 U33 U12 U13 U23 Bi0A 0.0106 (3) 0.0106 (3) 0.0164 (5) 0.000 0.000 0.000 Bi2 0.0162 (4) 0.0102 (3) 0.0224 (4) 0.000 0.000 0.000 I6 0.0106 (12) 0.0106 (12) 0.0092 (16) 0.000 0.000 0.000 I3 0.0160 (4) 0.0093 (4) 0.0250 (5) 0.000 0.000 0.0029 (11) I2 0.0175 (14) 0.0175 (14) 0.021 (2) 0.000 0.000 0.000 I1 0.0176 (4) 0.0101 (4) 0.0245 (5) 0.000 0.000 −0.0106 (11) I7 0.0223 (3) 0.0223 (3) 0.0242 (5) 0.0031 (4) −0.0077 (9) 0.0077 (9) Rb3 0.0242 (13) 0.0242 (13) 0.039 (2) 0.0089 (16) −0.0016 (11) 0.0016 (11) I10 0.067 (3) 0.067 (3) 0.051 (3) 0.000 0.000 0.000

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I4A 0.012 (2) 0.018 (2) 0.025 (4) 0.000 −0.008 (2) 0.000 Rb2 0.0184 (11) 0.0184 (11) 0.0214 (16) 0.0054 (13) −0.0028 (10) −0.0028 (10) I5B 0.024 (4) 0.019 (3) 0.033 (4) 0.000 −0.006 (3) 0.000 Rb1 0.091 (3) 0.0109 (14) 0.0219 (16) 0.000 0.007 (6) 0.000 I4B 0.026 (3) 0.018 (2) 0.017 (2) 0.000 0.006 (3) 0.000 I5A 0.044 (7) 0.015 (2) 0.050 (6) 0.000 −0.031 (5) 0.000 Cu1 0.058 (9) 0.058 (9) 0.073 (18) 0.000 0.000 0.000 I11 0.056 (7) 0.056 (7) 0.103 (17) 0.000 0.000 0.000 O1 0.037 (5) 0.037 (5) 0.051 (16) 0.008 (7) 0.017 (11) 0.017 (11)

Table S7.4. Fractional atomic coordinates and isotropic or equivalent isotropic displacement parameters (Å2) for XV.

X y z U */U Occ. (<1) iso eq Bi1 0.5000 0.5000 0.5000 0.0201 (3)

Bi2 0.2500 0.2500 0.0000 0.0240 (2)

I1 0.38212 (7) 0.38440 (7) 0.5000 0.0240 (2)

I2 0.5000 0.5000 0.79381 (14) 0.0214 (3)

I3 0.13506 (6) 0.13372 (6) 0.0000 0.0245 (2) I4A 0.15312 (10) 0.34062 (10) 0.17671 (16) 0.0234 (5) 0.497 (4) I4B 0.18424 (10) 0.31782 (10) 0.24619 (17) 0.0282 (6) 0.497 (4)

I5 0.0000 0.32623 (13) 0.5000 0.0396 (5) I6 0.0000 0.5121 (2) 0.1758 (5) 0.0587 (19) 0.448 (6)

I7 0.16940 (13) 0.5000 0.5000 0.0527 (7) Ag1 0.0764 (4) 0.4212 (2) 0.3508 (4) 0.095 (2) 0.441 (6) Ag2 0.0000 0.5000 0.5000 0.099 (4) 0.68 (2)

Rb1 0.0000 0.20304 (12) 0.2294 (3) 0.0375 (5) Rb2 0.2134 (2) 0.2880 (2) 0.5000 0.0297 (7) 0.5

Rb3 0.29387 (11) 0.5000 0.7721 (2) 0.0326 (4)

O1 0.0000 0.2977 (16) 0.0000 0.093 (15)

O2 0.1985 (11) 0.5000 1.0000 0.035 (5)

Table S7.5. Atomic displacement parameters (Å2) for XV.

U11 U22 U33 U12 U13 U23 Bi1 0.0178 (6) 0.0261 (7) 0.0164 (5) 0.000 0.000 0.000 Bi2 0.0228 (5) 0.0217 (5) 0.0276 (5) −0.0055 (3) 0.000 0.000 I1 0.0213 (5) 0.0243 (6) 0.0262 (5) −0.0050 (4) 0.000 0.000 I2 0.0221 (7) 0.0273 (8) 0.0148 (6) 0.000 0.000 0.000

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I3 0.0202 (5) 0.0185 (5) 0.0348 (6) −0.0019 (4) 0.000 0.000 I4A 0.0269 (9) 0.0216 (8) 0.0217 (9) −0.0015 (6) 0.0054 (7) −0.0028 (6) I4B 0.0324 (10) 0.0284 (9) 0.0238 (9) −0.0050 (7) 0.0083 (8) −0.0083 (7) I5 0.0411 (11) 0.0455 (12) 0.0322 (10) 0.000 0.000 0.000 I6 0.0368 (16) 0.031 (4) 0.109 (3) 0.000 0.000 −0.019 (2) I7 0.0303 (10) 0.102 (2) 0.0254 (10) 0.000 0.000 0.000 Ag1 0.165 (6) 0.068 (3) 0.053 (2) −0.060 (3) −0.008 (3) 0.010 (2) Ag2 0.116 (9) 0.128 (10) 0.052 (5) 0.000 0.000 0.000 Rb1 0.0279 (9) 0.0325 (10) 0.0520 (14) 0.000 0.000 −0.0122 (10) Rb2 0.038 (2) 0.0321 (19) 0.0195 (15) −0.0166 (15) 0.000 0.000 Rb3 0.0308 (9) 0.0458 (11) 0.0212 (9) 0.000 −0.0027 (7) 0.000 O1 0.039 (15) 0.039 (15) 0.20 (5) 0.000 0.000 0.000 O2 0.026 (10) 0.028 (10) 0.050 (14) 0.000 0.000 0.000

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343. Dehnhardt, N.; Borkowski, H.; Schepp, J.; Tonner, R.; Heine, J., Ternary iodido bismuthates and the special role of copper. Inorg. Chem. 2018, 57 (2), 633-640. 2- 344. Adonin, S. A.; Sokolov, M. N.; Smolentsev, A. I.; Kozlova, S. G.; Fedin, V. P., [PtBi2I12] : the first polyiodobismuthate containing an octahedral heterometallic unit. Dalton Trans. 2013, 42 (27), 9818-9821. 345. Bigalke, K. P.; Hans, A.; Hartl, H., Synthese und Strukturuntersuchungen von Iodocupraten(I). IX. Synthese und Kristallstrukturen von Cs3Cu2I5 und RbCu2I3. Z. Anorg. Allg. Chem. 1988, 563 (1), 96-104. 346. Hull, S.; Berastegui, P., Crystal structures and ionic conductivities of ternary derivatives of the silver and copper monohalides—II: ordered phases within the (AgX)x–(MX)1−x and (CuX)x–(MX)1−x (M=K, Rb and Cs; X=Cl, Br and I) systems. J. Solid State Chem. 2004, 177 (9), 3156-3173. 347. Shen, Y.; Lu, J.; Tang, C.; Fang, W.; Jia, D.; Zhang, Y., Syntheses and properties of 2-D and 3- D Pb–Ag heterometallic iodides decorated with ethylene polyamines at the Pb(ii) center. Dalton Trans. 2014, 43 (24), 9116-9125. 348. Keller, L.; Nason, D., Review of X-ray powder diffraction data of rhombohedral bismuth tri- iodide. Powder Diffr. 2013, 11 (2), 91-96. 349. Shao, S.; Liu, J.; Portale, G.; Fang, H.-H.; Blake, G. R.; ten Brink, G. H.; Koster, L. J. A.; Loi, M. A., Highly reproducible Sn-based hybrid perovskite solar cells with 9% efficiency. Adv. Energy Mater. 2017, 8 (4), 1702019.

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A creative page | Olga Nazarenko

A creative page

The images presented below are made by the Ph.D. candidate on the transmission electron microscope CM12 that was provided by the Scientific Center for Optical and Electron Microscopy (ScopeM) at ETH Zurich. Each Image is followed by a short history of the sample.

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Curriculum Vitae| Olga Nazarenko

Curriculum Vitae

Name Olga Nazarenko

Nationality Ukrainian

Education: Since September 2013 Doctoral studies, under the supervision of Prof. Dr. Maksym V. Kovalenko, Department of Chemistry and Applied Biosciences, ETH Zürich, Switzerland and Empa - Swiss Federal Laboratories for Materials Science and Technology - Dübendorf, Switzerland.

September 2011 - June M.Sc. studies in Chemistry, Department of Chemistry, Taras 2013 Shevchenko National University of Kyiv, Ukraine.

Awards 2017 SCNAT_SCS Chemistry Travel Award Professional experience Teaching assistant for Laboratory Course General Chemistry, 529-0011-04L (fall semesters 2014 - 2016)

List of publications: Doctoral publications 12. O. Nazarenko, M. R. Kotyrba, S. Yakunin, M. Wörle, B. M. Benin, G. Rainò, F. Krumeich, M. Kepenekian, J. Even, C. Katan, M. V. Kovalenko Guanidinium and mixed cesium-guanidinium tin(II) bromides: effects of quantum confinement and out-of-plane octahedral tilting Chem. Mater. 2019

11. B. M. Benin, D. N. Dirin, V. Morad, M. Wörle, S. Yakunin, G. Rainò, O. Nazarenko, M. Fischer, I. Infante, M. V. Kovalenko Highly emissive self‐trapped excitons in fully inorganic zero‐ dimensional tin halides Angew. Chem. Int. Ed. 2018, 57, 11329 10. L. Protesescu, S. Yakunin, O. Nazarenko, D. N. Dirin, M. V. Kovalenko Low-cost synthesis of highly luminescent colloidal lead halide perovskite nanocrystals by wet ball milling ACS Appl. Nano Mater., 2018, 1, 3, 1300-1308

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9. O. Nazarenko, M. R Kotyrba, S. Yakunin, M. Aebli, G. Rainò, B. M. Benin, M. Wörle, M. V. Kovalenko Guanidinium-formamidinium lead iodide: a layered perovskite- related compound with red luminescence at room temperature J. Am. Chem. Soc., 2018, 11, 140, 3850-3853

8. F. Krieg, S. Ochsenbein, S. Yakunin, S. Brinck, P. Aellen, A. Süess, B. Clerc, D. Guggisberg, O. Nazarenko, Y. Shynkarenko, S. Kumar, C.-J. Shih, I. Infante, and M.V. Kovalenko Colloidal CsPbX3 (X=Cl, Br, I) nanocrystals 2.0: zwitterionic capping ligands for improved durability and stability ACS Energy Lett., 2018, 3 (3), 641-646

7. O. Nazarenko, M.R. Kotyrba, M. Wörle, E. Cuervo-Reyes, S. Yakunin, and M.V. Kovalenko Luminescent and photoconductive layered lead halide perovskite compounds comprising mixtures of cesium and guanidinium cations Inorg. Chem., 2017, 56, 11552–11564

6. M. Ibáñez, R. Hasler, Y. Liu, O. Dobrozhan, O. Nazarenko, D. Cadavid, A. Cabot, and M.V. Kovalenko Tuning p-type transport in bottom-up-engineered nanocrystalline Pb chalcogenides using alkali metal chalcogenides as capping ligands Chem. Mater., 2017, 29, 7093–7097

5. C. Mittag, M. Karalic, S. Mueller, T. Tschirky, W. Wegscheider, O. Nazarenko, M.V. Kovalenko, T. Ihn, and K. Ensslin Passivation of edge states in etched InAs sidewalls Appl. Phys. Lett., 2017, 111, 082101

4. O. Nazarenko, S. Yakunin, V. Morad, I. Cherniukh, M. V. Kovalenko Single crystals of caesium-formamidinium lead halide perovskites: solution growth and gamma dosimetry NPG Asia Mater., 2017, 9, e373

3. S. Yakunin, D. N. Dirin, Y. Shynkarenko, V. Morad, I. Cherniukh, O. Nazarenko, D. Kreil, T. Nauser, M. V. Kovalenko Detection of gamma photons using solution-grown single crystals of hybrid lead halide perovskites Nat. Photon. 2016, 10 (9), 585-589

Pre-doctoral 2. O. M. Nazarenko, E. B. Rusanov; A. N. Chernega, K. V. publications Domasevitch

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Curriculum Vitae| Olga Nazarenko

Cobalt(II) and cadmium(II) square grids supported with 4,4’- bipyrazole and accommodating 3-carboxyadamantane-1- carboxylate Acta Cryst. Sect. C, 2013, C69, p. 232-236

1. O. M. Nazarenko, J.A. Rusanova, H. Krautscheid, K.V. Domasevitch Catena-Poly[thorium(IV)-tetrakis-(2-3-carboxyadamantane-1- carboxylato)]: a quadruple helical strand driven by a synergy of coordination and hydrogen bonding Acta Cryst. Sect. C, 2010, C66, p.276-279

List of presentations: Oral presentations 6. Olga Nazarenko, Martin R. Kotyrba, Sergii Yakunin, Marcel Aebli, Gabriele Rainò, Bogdan M. Benin, Michael Wörle, and Maksym V. Kovalenko Guanidinium-formamidinium lead iodide: a layered perovskite- related compound with red luminescence at room temperature 2018 – Swiss Chemical Society (SCS) fall meeting, Lausanne, Switzerland

5. O. Nazarenko, M. R. Kotyrba, S. Yakunin, M. Wörle, E. Cuervo-Reyes, M. Aebli, G. Rainò, B. M. Benin, M. V. Kovalenko Layered cesium - guanidinium Sn(II) and Pb(II) halide perovskites 2018 – Journées Pérovskites Halogénées 2018 (JPH2018), Grenoble-Autrans, France

4. O. Nazarenko, M. R. Kotyrba, S. Yakunin, M. Wörle, E. Cuervo-Reyes, M. Aebli, G. Rainò, B. M. Benin, M. V. Kovalenko Layered lead halide perovskite compounds comprising both cesium and guanidinium cations 2018 – Materials Research Society (MRS), Arizona, Phoenix, USA

3. O. Nazarenko, S. Yakunin, V. Morad, I. Cherniukh, M. V. Kovalenko Solution grown caesium formamidinium lead halide perovskites for detection of gamma photons 2017 – SCS fall meeting, Bern, Switzerland

2. O. Nazarenko, M. R. Kotyrba, M. Wörle, E. Cuervo-Reyes, S. Yakunin, M. V. Kovalenko Structure and optical properties of organic – inorganic lead halide compounds with guanidinium cation

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2016 – LAC (Laboratorium für Anorganische Chemie) Christmas symposium, Zürich, Switzerland

1. O. Nazarenko, K. V. Domasevitch Adamantancarboxylates: building blocks for structural design of coordination polymers 2011 – XII Ukrainian conference of students and post-graduate students, Taras Shevchenko National University of Kyiv, Ukraine

Poster presentations: 4. O. Nazarenko, S. Yakunin, V. Morad, I. Cherniukh, M. V. Kovalenko Single crystals of caesium formamidinium lead halide perovskites: solution growth and gamma dosimetry 2017 – PSCO (perovskites solar cells and optoelectronics), Oxford, UK

3. O. Nazarenko, S. Yakunin, V. Morad, I. Cherniukh, M. V. Kovalenko Single crystals of caesium formamidinium lead halide perovskites: phase stability and gamma dosimetry 2016 – MAP (Materials and Processes) symposium, Zürich, Switzerland

2. O. Nazarenko, M. R. Kotyrba, M. Wörle, E. Cuervo-Reyes, S. Yakunin, M. V. Kovalenko Structure and optical properties of organic-inorganic lead halide compounds with guanidinium cation 2016 – International Workshop on Supramolecular Chemistry & Functional Materials, Japan

1. O. Nazarenko, M. R. Kotyrba, S. Yakunin, M. Wörle, M. V. Kovalenko Hybrid lead and tin halide perovskites with guanidinium cation 2016 – SCS fall meeting, Zürich, Switzerland

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