THE GROWTH OF ORGANIC SMALL MOLECULE AND INORGANIC HALIDE PEROVSKITE CRYSTALLINE THIN FILMS By Pei Chen A DISSERTATION Submitted to Michigan State University in partial fulfillment of the requirements for the degree of Material s Science and Engineering Doctor of Philosophy 201 9 ABSTRACT THE GROWTH OF ORGANIC SMALL MOLECULE AND INORGANIC HALIDE PEROVSKITE CRYSTALLINE THIN FILMS By Pe i Chen Organic semiconductors have shown except ional opportunities for manipulating energy in a range of structures in light - emitting diodes, lasers, transistors, transparent photovoltaics, etc. w ith the presence of excitons at room temperature that distinguishes them from traditional semiconductors . The con trol over the crystalline order, orie ntation, layer - coupling as well as defect formation are the key to the fabricati on and optimization for improving the performance of organic electronics. In the first part of this thesis, we focus on understanding organ ic crystalline growth. O rganic homoep itaxy growth mode is mapped as a function of vapor phase growth conditions on high quality organic crystalline substrates. Organic - organic hetero - quasiepitaxy is then studied to explore the design rules for ordered alternating organic growth similar to ino rganic quantum well structure . A unique organic edge driven case is demonstrated provid ing new routes to controlling molecular orientation and multilayer ordering . The se results could enable entirely new opportunities for enhancing unique excitonic tunabil ity and could also be used as a platform to study organic exciton confinement and strong coupling . The second part of the thesis is focused on inorganic halide perovsk ite growth. Hybrid halide perovskites have attracted tremendous attention as an exception al new class of semiconductors for solar harvesting, light emission, lasing, quantum dots, thin film electronics, etc. However, the toxicity of lead devices and lead m anufacturing combined with the instability of organic components have been two key barrie rs to widespread application s . In this work, we demonstrate the first single - domain epitaxial growth of halide perovskites. This in situ growth study is enabled by the study of homoepitaxy and mixed - homoepitaxy of metal halide crystals that demonstrates th e capability of performing reflection high - energy electron diffraction (RHEED) on insulating surfaces. We then focus on tin - based inorganic hali de perovskites, CsSnX 3 (X = Cl, Br, and I) , on lattice - matched metal halide crystals via reactive vapor growth r oute that leads to single - domain epitaxial films with excellent crystalline order lacking in solution process ing . Exploiting this highly controllable epitaxial growth we demonstrate the first halide perovskite quantum wells that creates photoluminescent tu nability with different well width. These demonstrations could spark the exploration of a full range of epitaxial halide perovskites and lead to novel applications for metal - halide - perovskite based single - crystal epitaxial optoelectronics. C opyright by PEI CHEN 2019 v ACKNOWLEDGME NTS I want to first express my sincere gratitude to my advis o r , Dr. Richard Lunt, for his guidance not only on science but also on life throughout my graduate career in Michigan State University. I will remember every piece of advice. I also thank my dis sertation committee members, Dr. Thomas Bieler, Dr. Scott Barton and Dr. Pengpeng Zhang for their help with my research. I am so grateful to be part of the MOELab family where everyone is so kind. Yimu, you were the big sister in this group and g a ve me s o many great pieces of advice when I first join ed the lab. Chris, I will so many great ideas for fixing things. And I would love to continue discussing assembling a desktop with you. I also love the e n dless funny jokes you have. Peggy, you are so talented. I remember ed one night when I was working late in the office, and you start ed practicing your songs. It was so beautiful, and it immediately cleared my tiredness. Also, your politeness to everyone is an inspiration . Paddy, thank you so much for introducing me to many amaz ing Indian food s. Y ou are so good at pronouncing Chinese ph r ases , I am very impressed! Chenchen, you are the second night owl in our lab and I definitely do not know who the first is . Also, your funniness is on a whole different level. Lili, it was great working with you. Your passion towards everything affect ed me tremendously . Dianyi, you gave me so many useful pieces of advice . You are a model for me with your professional attitud e in research. And I would love to take more photos for your beautiful family. Matt, I admire your working efficiency and I really enjoyed your jokes and sarcasm . Chris, the second Chris, you help ed increase the average height of MOELab. And it was so plea s ant to have your table right next to mine. Isaac ( I will make sure not to spell your name as Issac this time ) I have enjoyed mentoring you and look forward to great discoveries from you . Thank you for your trust in me. vi Audrey and John , I wish you the best in your future studies. Also, thank you to my former colleagues Yunhua, Dhanashree, Luca s , John, Tyler, Nathan, Zach, Alex, Matthew, Juan, Crystal, Jorge, Brian, Mark, Adam, Joe, and Kevin. I would also like to thank my collaborators: Dr. Richard Staples f rom the Department of Chemistry, Dr. Per Askeland from the Composite Materials and Structures Center , Dr. Johannes Pollanen and Dr. Kostyantyn Nasyedkin from the Department of Physics and Astronomy and Dr. Kai Sun from the University of Michigan . Your help and guidance have been invaluable . I want to thank my friend s for all those wonderful moments that I will always treasure. Zhilin, thank you for your heartwarming support and encouragement. Dong and Di, you will late - night group chats. Yuxi, you are my best buddy for photography and playing Warframe. Zeyang, you are the best roommate in the world. Six years without a single argument is just a miracle. Mingmin, thank you fo r all those amazing travel experience s . Xinting, you are so kind to me and thank you for the help on the TA labs! L ijiang, Yuanchao, Jialin, Xing, Yifeng, Yifan, Quan, Chuanpeng, Yuelin, Deliang, Mengyin, I appreciate all those good memories with you. Fina lly, thank you, Mo m and Dad, you are my dearest people in this world. You never ask for anything and provide me with unconditionally support. Your trust in me since I was a little kid is always the source of confidence to me. I can never thank you enough. vii TABLE OF CONTENTS LIST OF TABLES ................................ ................................ ................................ ................................ ..... ix LIST OF FIGURES ................................ ................................ ................................ ................................ .... x KEY TO ABBREVIA TIONS ................................ ................................ ................................ .................. xiv Chapter 1 Introduction ................................ ................................ ................................ ............................ 1 1.1 Organic Semiconductors ................................ ................................ ................................ ....... 1 1.1.1 Motivation ................................ ................................ ................................ ...................... 1 1.1.2 Review on Sm all Organic Molecule Crystalline Thin Film Growth .............................. 2 1.2 Inorganic Halide Perovskites ................................ ................................ ................................ 5 1.2.1 Motivation ................................ ................................ ................................ ...................... 5 1.2.2 Review on Halide Perovskite Epitaxial Growth ................................ ............................. 7 1.3 Thes is Outline ................................ ................................ ................................ ....................... 8 Chapter 2 Background on Thin Film Growth ................................ ................................ ....................... 9 2.1 Growth Physics ................................ ................................ ................................ ..................... 9 2.1.1 Elementary Processes ................................ ................................ ................................ ..... 9 2.1.2 Growth Modes ................................ ................................ ................................ .............. 10 2.2 Epitaxy ................................ ................................ ................................ ................................ 11 2.3 Quasiepitaxy (QE) ................................ ................................ ................................ ............... 13 2.3.1 Energetic Considerations ................................ ................................ .............................. 13 2. 3.2 Kinetic Considerations ................................ ................................ ................................ . 17 2.4 Key Issues with Inorganic Halide Perovskite Growth ................................ ........................ 19 2.5 Key Issues with Organic Thin Film Growth ................................ ................................ ....... 19 Chapter 3 Experimental Techniques ................................ ................................ ................................ .... 21 3.1 Physical Vapor Deposition (PVD) ................................ ................................ ...................... 21 3.1.1 Thermal Evaporation ................................ ................................ ................................ .... 21 3.1.2 Electron - beam (E - beam) Evaporation ................................ ................................ .......... 24 3.2 In situ Diffraction Technique: Reflection High - energy Electron Diffraction (RHEED) .... 25 3.2.1 RHEED Introduction ................................ ................................ ................................ .... 25 3.2.2 Electron Beam Damage ................................ ................................ ................................ 28 3.3 Ex situ Diffraction Technique ................................ ................................ ............................. 29 3.3.1 X - ray Diffraction (XRD) ................................ ................................ .............................. 29 3.3. 2 Pole Figure ................................ ................................ ................................ .................... 30 3.3.3 Rocking Curve ................................ ................................ ................................ .............. 34 3.4 Thin Film Morphology Measurements ................................ ................................ ............... 35 3.4.1 Atomic Force Microscopy (AFM) ................................ ................................ ................ 35 3.4 .2 Scanning Electron Microscope (SEM) ................................ ................................ ......... 36 3.5 Thin Film Optical Measurements ................................ ................................ ........................ 37 3.5.1 Ultraviolet - visible Spectroscopy (UV - Vis) ................................ ................................ .. 37 3.5.2 Variable Angle Spectroscopic Ellipsometry (VASE) ................................ .................. 38 viii Chapter 4 Organic Homoepitaxial and Hetero - quasiepitaxial Growth ................................ ............ 40 4.1 Organic Homoepitaxial Growth ................................ ................................ .......................... 40 4.1.1 Organic Single Crystal Substrate Growth ................................ ................................ .... 40 4.1.2 Homoepitaxial Growth of 9,10 - Diphenylanthracene (DPA) ................................ ........ 45 4.2 Ordered Organic - organic Hetero - quasiepitaxial (QE) Gro wth ................................ ........... 49 4.3 Conclusion ................................ ................................ ................................ ........................... 54 Chapter 5 Uncovering a New Growth Mode: Edge - Driven Organic - on - Organic Quasiepitaxy .... 56 5.1 Experiment al ................................ ................................ ................................ ....................... 56 5.2 NTCDA and DIP ................................ ................................ ................................ ................. 58 5.3 Ordered Alternating Crystalline Growth ................................ ................................ ............. 59 5.4 Reversed Or der Growth ................................ ................................ ................................ ...... 62 5.5 In - plane Registry Modelling ................................ ................................ ............................... 63 5. 6 Edge - driven Mechanism ................................ ................................ ................................ ..... 66 5.7 Conclusion ................................ ................................ ................................ ........................... 67 Chapter 6 Homoepitaxial Growth of Alkali Halide Crystals ................................ ............................. 69 6.1 Introduction ................................ ................................ ................................ ......................... 69 6.2 Experimental ................................ ................................ ................................ ....................... 70 6.3 Growth Modes of NaCl Homoepitaxy ................................ ................................ ................ 72 6.4 NaCl Based OLEDs ................................ ................................ ................................ ............ 78 6.5 Conclusion ................................ ................................ ................................ ........................... 79 Chapter 7 Single - Domain Epitaxy of Cesium Tin Bromide (CsSnBr 3 ) ................................ ............. 80 7. 1 Experimental ................................ ................................ ................................ ....................... 80 7.2 Phase Study via Stoichiometry Control ................................ ................................ .............. 84 7.3 Epitaxial Lift - off ................................ ................................ ................................ ................. 93 7.4 Discussion ................................ ................................ ................................ ........................... 96 7.5 Conclusion ................................ ................................ ................................ ......................... 101 Chapter 8 Epita xial Stabilization of Tetragonal Cesium Tin Iodide (CsSnI 3 ) ............................... 103 8.1 Experimental ................................ ................................ ................................ ..................... 104 8.2 Epitaxial Ph ase Stabilization ................................ ................................ ............................. 107 8.3 Optical P roperties and Quantum Well Fabrication ................................ ........................... 118 8.4 Epitaxial CsSnI 3 Thin Film Based Photodetector Fabrication ................................ .......... 120 8.5 Conclusion ................................ ................................ ................................ ......................... 122 Chapter 9 Concl usions and Future Outlook ................................ ................................ ...................... 123 9.1 Molecular Dynamics (MD) Simulation of Organic Surface Diffusion ............................. 123 9.2 Organic - organic Charge Transfer (CT) Complexes ................................ .......................... 127 9.3 Expand Halide Perovskite Heteroepitaxial Growth Pairings ................................ ............ 128 9.4 Epitaxial Hali de Perovskite Doping ................................ ................................ .................. 129 9.5 Epitaxial Halide Perovskite Applications ................................ ................................ ......... 129 APPENDIX ................................ ................................ ................................ ................................ .............. 130 BIBLIOGRAPHY ................................ ................................ ................................ ................................ ... 134 ix LIST OF TABLES Table 4.1. List of organic materials for CMD. ................................ ................................ ............ 43 Table 4.2. Calculated surface energies (SE) of LES) for a range of organic crystals including NTCDA and material s for paired alternating growth. ................................ ..... 50 Table 7.1. Lattice constants, film orientation and misfit of two phases, CsS nBr 3 and CsSn 2 Br 5 . 90 Table 7.2. Elemental ratio of as - deposited films with d ifferent of CsBr:Sn Br 2 obtained from XPS data. ................................ ................................ ................................ ............... 91 Table 7.3. Calculated band gaps of CsSnBr 3 and CsSn 2 Br 5 using the DFT - PBE and DFT - HSE06 methods. ................................ ................................ ................................ ................................ ........ 97 Table 7.4. Emission energy of quantum well CsSnBr 3 /NaCl with various well width. ............. 101 Table 7.5. Emission energy of quantum well CsSnBr 3 /CsSn 2 Br 5 with various well width. ....... 101 Table 8.1. Elemental ratio of as - deposited CsSnI 3 film obtained from XPS data. ...................... 113 Table 8.2. Crystal structures, lattice parameters, experimental and simulated bandgaps of the CsSnI 3 cubic, tetragonal, orthorhombic and the epi - tetragonal phase from this work. ............... 119 x LIST OF FIGURES Figure 1.1. Carrier mobility for single crystals, thick films, and monolayer films for various common organic small molecule semiconductors. ................................ ................................ ......... 1 Figure 1.2. TEM for AlGaN/AlN quantum wells for UV laser applications and sc hematic of multilayer organic single crystalline structure. ................................ ................................ ............... 2 Figure 1.3. Examples of organic - inorganic QE growth. ................................ ................................ . 3 Figure 1.4. Examples of organic - organic QE growth. ................................ ................................ .... 4 Figure 2.1. The elementary molecule diffusion processes of epitaxial thin film growth. .............. 9 Figure 2.2. Schematic diagrams and transmission electron microscopy (TEM) examples of commensurate, pseudomorphic and incommensurate growth in epitaxy. ................................ .... 12 Figure 2.3. Schematic modes of epitaxy and quasiepitaxial overlayer alignments. ..................... 14 Figure 2.4. Example o f V / V 0 analysis for zinc phthalocyanine (ZnPc) on deactivated Si(111) - B surface. ................................ ................................ ................................ ................................ .......... 16 Figure 2.5. The plot of free energy change G as a function of r . ................................ ............... 18 Figure 3.1. Evaporator diagrams. ................................ ................................ ................................ .. 22 Figure 3.2. Inficon SQS - 242 deposition software interface during a two - source co - deposition (shown in manual mode). ................................ ................................ ................................ .............. 24 Figure 3.3. Schematic of the E - beam deposition process. ................................ ............................ 25 Figure 3.4. RHEED schematics. ................................ ................................ ................................ ... 26 Figure 3.5. Example RHEED patterns. ................................ ................................ ......................... 27 Figure 3.6. Growth mode with the RHEED intensity signal vs. time. ................................ .......... 28 Figure 3.7. XRD geometry. ................................ ................................ ................................ ........... 30 Figure 3.8. Sphere of fixed - length scattering vect or and stereographic projection. ..................... 31 Figure 3.9. Example of pole figures. ................................ ................................ ............................. 32 Figure 3.10. Rocking curve measurement geometry and an example rocking curve scan of (004) for Si(Ge) on Si. ................................ ................................ ................................ ............................ 34 Figure 3.11. Typical set - up of AFM. ................................ ................................ ............................ 36 xi Figure 3.12. Basic construction of a SEM system. ................................ ................................ ....... 37 Figure 3.13. Schematic of a UV - Vis spectrometer configured to measure transmission for a thin film sample. ................................ ................................ ................................ ................................ ... 38 Figure 4.1 . CMD process for fabricating single crystalline organic substrates. ........................... 41 Figure 4.2. XRD patterns and cross polarized optica l images of the organic crystalline substrates prepared using CMD. ................................ ................................ ................................ .................... 42 Figure 4.3. Vapor growth method for fabricating single crystalline organic substrates. .............. 44 Figure 4.4. Schematic of DPA homoepitaxial growth. ................................ ................................ . 45 Figure 4.5. Structural characterizatio n of the DPA single crystal substrate. ................................ 46 Figure 4.6. RHEED for DPA homoepitaxy of 20nm at 75°C. ................................ ...................... 47 Figure 4.7. AFM images of DPA homoepitaxy. ................................ ................................ ........... 48 Figure 4.8. In situ RHEED patterns of the alternating growth pairs. ................................ ............ 51 Figure 4.9. XRD patterns of the alternating growth pairs. ................................ ............................ 52 Figure 4.10. AFM images of crystalline alternating growth pairs. ................................ .............. 53 - - DIP. ................................ ...................... 58 Figure 5.2. In situ RHEED patterns of ordered NTCDA/DIP alternating growth. ....................... 59 Figure 5.3. XRD characterization and structure models. ................................ .............................. 60 Figure 5.4. Reverse order growth. ................................ ................................ ................................ 62 Figure 5.5. Surface potential modeling of the DIP/NTCDA quasiepitaxial registry. ................... 64 Figure 5.6. Calculated dependence of the normalized geometric potential, V/V 0 , as a function of the azimuthal angle for (020) DIP on (100) NTCDA. ................................ ............................... 65 Figure 5.7. AFM image of early - stage nucleation and structural model of the s tep edge growt h. 66 Figure 6.1. Temperature dependent in situ R HEED patterns of NaCl homoepitaxy. ................... 72 Figure 6.2. RHEED oscillations of NaCl homoepitaxial growth vs. growth rate a nd temperature. ................................ ................................ ................................ ................................ ....................... 73 Figure 6.3. NaCl temperature de pendent homoepitaxial growth mode and 3D morphology. ...... 74 Figure 6.4. NaCl temperature dependent homoepitaxial growth mode and 2D line scan. ........... 75 xii Figure 6.5. NaCl based OLEDs. ................................ ................................ ................................ ... 78 Figure 7.1. I n situ RHEED patterns of the epitaxial growth of CsSnBr 3 on NaCl. ...................... 84 Figure 7.2. Rotation dependent RHEED patter ns of CsSnBr 3 and CsSn 2 Br 5 epitaxial film on NaCl . ................................ ................................ ................................ ................................ ....................... 85 Figure 7.3. RHEED oscillations of CsSnBr 3 growth on NaCl. ................................ ..................... 86 Figure 7.4. Ordering of the halide perovskite at the halide salt interface. ................................ .... 87 Figure 7.5. In situ real - time RHEED monitoring of the phase transition. ................................ .... 88 Figure 7.6. RHEED oscillations monitored during the phase transition. ................................ ...... 89 Figure 7.7. Crystal structure characterization of the two epitaxial phases, CsSnBr 3 and CsSn 2 Br 5 . ................................ ................................ ................................ ................................ ....................... 90 Figure 7.8. Epitaxial growth of CsSnBr 3 on pseudomorphic interlayer s of NaCl - NaBr. ............. 92 Figure 7.9. XRD patterns of alloyed NaCl - NaBr pseudomorphic interlayers and CsSnBr 3 . ....... 93 Figure 7.10. Film adhes ion of the epitaxial growth of CsSnBr 3 on NaCl and epitaxy lift - off/regrowth. ................................ ................................ ................................ ................................ . 94 Figure 7.1 1 . RHEED patterns of growth of CsSnBr 3 on Ge and InP . ................................ .......... 95 Figure 7.12. Absorption spectra of CsSnBr 3 , CsSn 2 Br 5 and NaCl substrate . ............................... 95 Figure 7.1 3 . Electronic band structures of CsSnBr 3 and CsSn 2 Br 5 . ................................ ............. 96 Figure 7.1 4 . Solar cell device from epitaxial lift - off of CsSnBr 3 . ................................ ................ 97 Figure 7.1 5 . Epitaxial CsSnBr 3 based 2D Quantum well fabrication. ................................ .......... 99 Figure 8.1. Schematic of the epitaxial growth. ................................ ................................ ........... 106 Figure 8.2. In situ RHEED patterns of the epit axial growth of CsSnI 3 . ................................ ..... 107 Figure 8.3. RHEED pattern of the bare KCl substrate and surface reconstruction of Cs SnI 3. ... 108 Figure 8.4. XRD and synchrotron XRD scans of epitaxial CsSnI 3 ................................ ............ 110 Figure 8.5. Pole figure scans of CsSnI 3 on KCl. ................................ ................................ ......... 111 Figure 8.6. TEM and SEM image of CsSnI 3 on KCl. ................................ ................................ . 112 Figure 8.7. Simulated Cs 2 SnI 6 diffraction patterns. ................................ ................................ .... 114 xiii Figure 8.8. Lifetime test for the epitaxial CsSnI 3 on KCl. ................................ .......................... 115 Figure 8.9. Growth of CsSnI 3 on KBr. ................................ ................................ ....................... 116 Figure 8.10. Epitaxial growth of CsSnI 3 on pseudomorphic interlayers of KCl - NaCl. .............. 117 Figure 8.11. Absorption of the epitaxial CsSnI 3 film and PL of the quantum wells. ................. 118 Figure 8.12. Epitaxial CsSnI 3 based photodetector fabrication and characterization. ................ 121 Figure 9.1. The set - up of the admolecule - substrate system 4 Figure 9.2. E m and D 0 plots as a function of L/W and molecular weight (M) 5 Figur e 9.3. (a) Fractional localized time (% T Localized ) vs. L/W . (b) Fractional time of a molecule travel along <01> (% T <01> ) vs. L/W 6 Figure 9.4. Preliminary results on alternating growt h of charge transfer salts 7 Figure A.1. Cover Picture for Advanced Materials Interfaces Volume 3, Issue 17 (September 2016). ................................ ................................ ................................ ................................ .......... 132 Figure A.2. Cover Picture for Advanced Materials Interfaces Volume 4, Issue 22 (November 2017). ................................ ................................ ................................ ................................ .......... 133 xiv KEY TO ABBREVIATIONS 1D one - dimensional 2D two - dimensional 3D three - dimensional - 4T - quaterthiophene - NPD - bis[N - (1 - naphthyl) - N - phe nyl - amino]biphenyl AFM atomic force microscopy Alq 3 tris - (8 - hydroxyquinolinato) aluminum AM1.5 G global standard spectrum APS advanced photon source BL bilayer (full unit cell thickness) BSE backscattered electrons CBP - - di carbazole - biphenyl CuPc copper(II) phthalocyanine CsBr cesium bromide C s I cesium iodide CsSnX 3 cesium tin chloride (X=Cl), cesium tin bromide (X=Br), cesium tin bromide (X=I) CT charge transfer DBP 5,10,15,20 - tetraphenylbisbenz[5,6]indeno[1,2,3 - - lm]perylene DB - TCNQ dibenzotetrathiafulvalene - tetracyanoquinodimethane DFT density functional theory DIP diindenoperylene DOS densities of states xv DPA 9,10 - diphenylanthracene DB - TCNQ dibenzotetrathiafulvalene - tetracyanoquinodimethane EL electrolu minescence ELO epitaxial lift - off F 16 CuPc copper hexadecafluorophthalocyanine FIB focused ion beam FWHM full width at half maximum GGA generalized gradient approximations H 2 Pc phthalocyanine HBC hexa - peri - hexabenzocoronene HOPG highly oriented pyroly tic graphite HSE06 Heyd - Scuseria - Ernzerhof hybrid functional IR infrared Ir(ppy) 3 fac tris(2 - phenylpyridine)iridium ITO indium tin oxide KCl potassium chloride L length of a molecule L1 - L4 layer 1 to 4 in alternating growth structure LES lowest en ergy surface LiQ 8 - hydroxyquinoline lithium M molecular weight MBE molecular beam epitaxy MeOH methanol xvi MD molecular dynamics ML monolayer (half unit cell thickness) NaCl sodium chloride NIR near - infrared NREL National Renewable Energy Laborator y NTCDA 1,4,5,8 - naphthalenetetracarboxylic dianhydride NVE microcanonical ensemble NVT canonical ensemble OLED organic light - emitting diode OPV organic photovoltaics p - 6P para - sexiphenyl PL photoluminescence PTCDA 3,4,9,10 - perylenetetracarboxylic d ianhydri de PTCDI - Ph N,N' - diphenyl - 3,4,9,10 - perylenedicarboximide PV photovoltaics PVD physical vapor deposition QCM quartz crystal monitor QE quasiepitaxial QT quaterrylene RAPD rapid automated processing of X - ray data RHEED reflection high - energy electron diffraction SAED selected area electron diffraction SE secondary electrons or surface energy xvii SEM scanning electron microscopy SnBr 2 tin bromide SnI 2 tin iodide STEM scanning transmission electron microscopy STO strontium titanate TC NQ tetracyanoquinodimethane TEM transmission electron microscopy TTF - TCNQ tetrathiafulvalene 7,7,8,8 - tetracyanoquinodimethane UV ultraviolet UV - Vis ultraviolet - visible spectroscopy VASE variable angle spectroscopic ellipsometry VASP Vienna Ab initio s imul ation package VIS visible light W width of a molecule XRD X - ray diffraction XPS X - ray photoelectron spectroscopy ZnPc zinc phthalocyanine xviii S ymbols Å Angstrom, equivalent to 0.1 nm or 10 - 10 m A absorbance , active area in calculating respons ivity of a photodetector A ij potential depth a lattice constant a B exciton Bohr radius a L lattice constant of the adlayer a S lattice constant of the substrate real lattice vectors of the adlayer real lattice vecto rs of the substrate B ij atomic interaction distance c speed of light C dose critical dose d thickness of film d width of the lattice rods D self (or tracer) diffusion coefficient D* specific detectivity D c chemical (or collective) diffusion coeff icient D 0 pre - exponential factor in Arrhenius equation for diffusivity E ES Ehrlich Schwoebel barrier E g bandgap E g 0 bulk band gap E m intra - terrace diffusion barrier xix Es diffusion barrier at the step edge EQE external quantum efficiency G free ener gy change during nucleus formation h Planck constant reduced Planck constant I intensity of the sample beam in UV - Vis measurement, electron beam current in RHEED I light current measured with light illumination I dark current measured without light ill umination I 0 intensity of the reference beam J current density J dar k extinction coefficient k 0 diameter of the Ewald sphere k B Boltzmann constant L luminescence current density L z thickness of the quantum well M transforma tion matrix between adlayer and substrate m e mass of electron m e * effective mass of the electron m h * effective mass of the hole n N M multiples of unit cell lattice vector ( ) of the adlayer xx N N multiples of unit cell lattice vector ( ) of the adlayer P incident incident light power density p v actual vapor pressure during deposition equilibrium vapor pressure q charge of the elect ron r separation distance between atoms in Lennard - Jones potential, radius of the nucleus R reflection; responsivity of a photodetector r p reflectivity of p - polarized light r s reflectivity of s - polarized light T transmittance , temperature %T Localized fractional localized time %T <01> fractional time of a molecule travel along [01] or [10] direction t time t damage time - to - damage t i initial (guessed) initial thickness in calculating tooling factor t f final (calculated) thickness in calculating tool ing factor TF i initial (guessed) tooling factor TF f final (calculated) tooling factor V voltage; simple periodic potential in registry modeling V 0 interatomic potential xxi tilt angle from sample surface normal direction in pole figure measurement rotation angle around sample surface normal direction in pole figure measurement free energy of the substrate f free energy of the interface free energy o f the epilayer phase component of ellipsometric angle permittivity r rel ative permittivity, also called dielectric constant 0 vacuum permittivity angle of the incident beam in XRD geometry, unit cell rotation angle between adlayer and substrat e complex reflectance wavelength intralayer p otential interlayer potential amplitude component of ellipsometric angle 1 Chapter 1 Introduction The goal of this chapter is to introduce small organic molecules and inorganic halide perovskites crystalline thin film growth, which are the two major themes d iscussed in this work. 1. 1 Organic Semiconductors 1. 1 .1 Motivation The presence of excitons, coulombically bound ed electron - hole pair s, in organic semiconductors at ro om temperature distinguishes them from traditional semiconductors, providing exceptional opportunities for manipulating energy in a range of structures in light emitting diodes, strongly coupled cavities, lasers, anisotropic transistors, transparent photov oltaics and excitonic switches . 2 - 7 The degree of crystal ordering for these organic layers has been linked to the efficiency of carrier transport with record mobilities of >30 cm 2 V - 1 s - 1 reported for single crysta ls ( Figure 1.1 ) , exciton migration , 8 and emerging photophysics such as singlet exciton fission and polariton lasing . 9 - 10 Defects have a large impact on charge transport in organic Figure 1 . 1 . Carrier mobility for single crystals, thick films, and monolayer films for various common organic small molecule semiconduct ors. Figure reprinted with permission. 1 2 electronic devices . 1 1 High defect density in the film will generally lead to large concentrati on of trap states that hinders the device performance . 12 Control over crystalline ord er, orientation and layer coupling are therefore critical to optimization and exploitation of exciton dynamics in these materials for next generation electronics. For inorganic semiconductors , quantum wells and quantum dots (i.e. that include many alternat ing layers) have played important roles in applications such as photodetector s, LEDs, and especially solid - state laser diode . 13 - 15 The organic equivalent ( Figure 1. 2 b ) could have simila rly substantial impact but has not been accessible before. Figure 1 . 2 . TEM for AlGaN/AlN quantum wells for UV laser applications and schematic of multilayer organic single crystalline structure . Figure rep roduced with permission . 13 1.1.2 Review on Small Organic Molecule Crystalline Thin Film Growth O rganic y from inorgan ic epitaxy that is based on lattice matching. T he quasiepitaxy (QE) energetically ordered growth (explained in Section 2.3 ) QE of organic lay ers can be achieved on both inorganic and organic substrates. Various high - qualit y inorganic crystalline substrates are readily available for organic thin film growth. Many studies have reported on organic - on - inorganic growth systems, including tetracene on SiO 2 , 16 - 17 Ag ; 18 pentacene on SiO 2 , 19 - 21 Au , 22 - 23 Ag , 24 - 25 Cu ; 26 - 27 diindenoperylene(DIP) 3 on SiO 2 , 28 - 30 Au ; 31 - 32 3,4,9,10 perylenetetracarboxylic dianhydride(PTCDA) on graphite , 32 - 3 3 SiO 2 , 34 GaAs , 35 Ag , 36 - 41 Au ; 42 - 44 copper phthalocyanine(CuPc) on graphite , 45 - 47 Si , 48 - 50 SiO 2 , 51 - 52 Au , 51 Ag 53 and metal halide . 45, 54 - 55 Example of organic - inorganic QE are shown in Figure 1.3 . One of the obstacles in understanding organi c homoepitaxial and organic - organic hetero - QE growth is the lack of high quality organic cr ystalline substrates. There are reports on the preparation of organic single crystals(substrates), but the sizes are relatively small. Single crystal naphthalene and anthracene sheets(<1cm 2 has been reported . 57 Single crystal naphthalene and anthracene a re also prepared from the liquid phase . Kikuchi patterns (explained in Section 3.2.1 ) have been observed by r eflection high - energy electron diffraction (RHEED) and were fitted to obtain the mean inner potential . 58 Needle - like rubrene single crystals (mm 2 size) were grown using physical vapor transport and used for homoepitaxial growth study . 59 - - 4T) single crystals grown fro m solution are used to demonstrate homoepitaxial growth 60 . Figure 1 . 3 . Examples of organic - inorganic QE growth . High - resolutio n STM data of (a) pentacene monolayer on HOPG and (b) PTCDA monolayer on the Ag. The unit cell of the adlayers are indicated. Figure reproduced with permission . 27, 56 4 Besides the difficulties in fabricating high quality organic crystalline substrates for subsequent layer growth, obstacles in understanding the organic - o rganic quasiepitaxy also include the incomplete understanding in the materials selection rules for organic quas iepitaxial growth as well as the limited use of effective in situ real time monitoring techniques. Cases of organic - organic hetero - quasiepitaxial growth systems are infrequent but includ e quaterrylene(QT) on hexa - peri - hexabenzocoronene(HBC) on Au(111) , 62 - 63 perylene - tetracarboxylic - dianhydride(PTCDA) - hexa - peribenzocoronene(HBC) - highly ordered pyrolytic gra phite (HOPG) , 64 rubrene on tetracene , 65 - 4T , 66 - 4T on Rubrene , 67 - 4T on tetracene sin gle crystal , 68 etc. Example of orga nic - organic Q E are shown in Figure 1.4. Case s of ordered alternating crystalline growth are extremely rare and only include PTCDA/3,4,7,9 naphthalenetetracarboxylic dianhydride(NTCDA) on highly ordered pyrolytic graphite(HOPG) , 69 dibenzotetrathiafulvalene - tetracyanoquinodimethane(DB - TCNQ)/ NTCDA on KBr , 70 pht halocyanine (H 2 Pc)/copper hexadecafluorophthalocyanine (F 16 CuPc) 71 and zinc phthalocyanine (ZnPc)/ N,N' - diphenyl - 3,4,9,10 - perylenedicarbox imide (PTCDI - Ph) on para - Figure 1 . 4 . Examples of organic - organic QE growth . - 4T on p - 6P and (b) C 60 on single crystal pentacene. Figure reproduced with permission . 61 - 62 5 sexiphenyl ( p - 6P)on Si/SiO 2 . 72 Clear QE relationships were discovered only in DB - TCNQ/NTCDA and PTCDI - Ph/ZnPc growth systems. It has been suggested that the matching of surfac e energy is essential in inducing ordered crystalline growth , 70 which is one of the predictive material selection criteria in pairi ng ordered organic growth that is more thoroughly investigated in Section 4. 2 . 1.2 Inorganic H alide P erovskite s 1.2.1 Motivation Hybrid organic - inorganic halide perovskite has the ABX 3 structure where A is an organic cation such as methylammonium (CH 3 NH 3 + ) and formamidinium ( HC( NH 2 ) 2 + ) , B is an inorganic cation which is usually Pb 2+ and X is halogen anion which can be Cl - , Br - , and I - . This type of perovskites ha s attracted tremendous attention as an exceptional new class of semiconductors for solar harves ting, 73 - 75 light emission, 76 lasing, 77 quantum dots, 78 water splitting 79 and thin film electronics. 80 Ever since the first report of organic - inorganic perovskite based solar cells with high power conversion efficiency of over 10%, 81 efficiencies of solar cells based on these materials h ave exceeded 22% only after five years , 82 however, the toxicity of lead devices and l ead manufacturing 83 - 85 combined with the instability of organic components 84, 86 - 90 have been two key barriers to widespread application. Other barriers include structura l defects , 91 - 92 trap states, 93 etc. Tin - based inorganic halide perovskites, such as CsSnX 3 (X = Cl, Br, and I), have been considered promising substitutes for their lead analogues since Sn is over 100 times less toxic than Pb and Cs has similar toxicity to Na or K. Ho wever, current research on photovoltaic and electronic applications of CsSnBr 3 and CsSnI 3 has, to date, been less encouraging, with solar cell efficiencies of < 5% for solution - processed thin film devices 94 - 95 that are likely limited by the low degree of crystalline ordering. 96 Indeed, structural ordering has been linked in traditional semiconductors to 6 a) carrier transport, where mobilities increase from amorphous - Si (1 cm 2 V - 1 s - 1 ) 97 to single crystalline Si (1,400 cm 2 V - 1 s - 1 ), 98 b) recombination rates, where unpa ssivated grain boundaries act as quenching sites for charge carriers and excited states and c) quantum confinement, which can make even Si an excellent NIR emitter with luminescent efficiency > 60%. 99 Th ese factors, among others, have motivated the recent interest in halide perovskite single crystal growth . 1 00 - 104 Thus, one of the main challenges for enhancing the properties of halide perovskites for high end optoelectronic applications is to o btain epitaxial crystalline films that can also be integrated into hetero epitaxial and quantum well structures. The epitaxial growth is a key step towards realizing 2 - dimensional (2D) electron gasses, 105 interface superconductivity, 106 ultra - high mobility strained transistors, 107 magnetoelectric multiferroics, 108 and the observation of fractional quantum hall effects. 109 As shown in previous research on oxide perovskites, numerous phases can be derived from t he perovskite structure with even minor changes in the elemental compositions. For example, by removing one - sixth of the oxygen atoms, phase transitions can occur from perovskite to brownmillerite structure s . 110 Therefore, it is key to gain precise control over the crystal phase, crystalline order, orientation, and interfaces for the optimization of halide perovskite based optoelectronics. Epitaxial growth has long been utilized to achieve the lowest defect densities and highest performance for applications in lasers , thin film transistors, 2D electron gases, sensors, and high - power devices. The development of inorganic materials such as III - V semiconductors has been extensively accelerated by epitaxial growth. Moreover, studies of epitaxial oxide perovskites have shown that unique properties can occur at the interfaces of different materials including superconductivity, ferroelectricity, and magnetism. These functional properties can be tun ed by engineering the symmetries and degrees of freedom of correlated electrons at the interfaces of 7 oxide perovskites, which is associated with the atomic arrangement at the interface . 111 - 114 Thus, epitaxial growth provides precise interface control and correspondingly yields the potential for interesting new systems such as magnetic superconductors, non - centrosymmetric superconductors and multiferroics . 115 - 119 These successes have drive n a new focus to control film quality, defect density, strain, and phases of new semiconducting materials that should be applicable to halide perovskites. 1.2.2 Review on Halide Perovskite Ep itaxial Growth While there has been significant research into th e epitaxial growth of oxide perovskites, such as the BaTiO 3 , SrRuO 3 and LaAlO 3 growth on SrTiO 3 (100), 112, 120 - 123 single - domain epitaxy has yet t o be explored for halide perovskites. This has likely been hindered in large part due to a number of challenges associated with the epitaxial growth of perovskites on dissimilar single crystal substrates including matching of lattice constants, lattice sym metry , coordination, wettability, thermal expansion differences, and bonding character (ionic versus covalent) . 124 There have been only a few reports of even metal halide salt epitaxy , 125 and only several recent studies of incommensurate v an der Waals epitaxy ( Q E ) of halide perovskite micro - sheets , 126 nanorods , 127 nanowires , 128 and nanocrystals 129 - 130 that are accompanied with a high degree of rotational disorder. Limited examples for halide perovskite epitaxy include the growth of hybrid perovskite MAPbBr 3 or MAPbB x I 3 - x on mica , 127, 131 but not none reported for pure ly inorganic perovskite s . In our work, l ow - cost alkali h alide salts such as NaCl, KBr, KCl, etc. are explored ( Chapter 7 and 8 ) as promising substrates for purely inorganic hal ide perovskite growth which provides the first demonstrat ion of the single - domain thin film epitaxy. These alkali halide substrates provide an ideal range of lattice constants (5.4 - 6.6 Å) closely matched to those of the halide perovskites (5.5 - 6.2 Å), with suitable wettability, and congruent ionic bonding . These 8 examples have confirmed the feasibility of vapor - phase epitaxial growth methods for t his group of materials. 1 . 3 Thesis Outline The rest of th is thesis is arranged as follows: In Chapter 2 , the theo ry behind thin film growth is discussed for both epitaxy and quasiepitaxy. Chapter 3 covers experimental techniques used in this study including diffraction techniques for analyzing crystal structure in situ and ex situ , thin film morphology , and optical m easurements . Chapter 4 discusses organic homo - epitaxy and hetero - quasiepitaxy (QE). Two methods for fabricating large area organic sing le crystals are demonstrated. A n example of organic homoepitaxy is studied and the growth mode is mapped as a function of growth rate and temperature. The rules for designing ordered alternating hetero - quasiepitaxial growth are then explored . Chapter 5 dis cusses a unique organic edge driven growth mechanism discovered in organic hetero - QE. Chapter 6 discusses the homoepitaxi al growth mode study for alkali metal halide crystals which demonstrates the capabilities to perform RHEED on insulating substrates tha t facilitates the demonstration o f inorganic halide epitaxy in Chapter 7 for CsSnBr 3 and Chapter 8 for CsSnI 3 where a new pseudomorphic phase is discovered in the latter . Chapter 9 provides conclusions and outlook for these areas. 9 Chapter 2 Background on Thin Film Growth In this chapter, the growth dynamics for thin film epitaxy are introduced . Conventional ino rganic epitaxy is compared with quasiepitaxy of small molecule semiconductor s . 2.1 Growth Physics 2. 1 .1 Elementary Processes When molecules are deposited onto a substrate, they experience multiple elementary diffusion processes : intra - terrace diffusion, inter - terrace diffusion, nucleation and aggregation as shown in Figure 2.1 . The most basic proces s is t he intra - terra c e diffusion ( b arrier E m ) when the terrace width is smaller than the mean free diffusion path of the admolecules. This process terminates when the admolecule hits a structural or chemical defect or with other admolecules. The molecule a pproaching a n edge on the substrate encounters a barrier ( E s ), which is usually large than E m . To descend from this step and achieve inter - terrace diffusion , extra diffusion barrier E ES needs to be overcome , known as Ehrlich Schwoebel barrier. 133 When surface diffusivity is Figure 2 . 1 . The elementary molecule diffusion processes of epitaxial thin film growth. Figure reprinted with permission. 132 10 relatively low, a large E ES will hinder the admolecules from descending and cause more molecules to attach to the upper terra ce which prom otes 3D growth. At high diffusivity, 2D growth is possibly achieved even with high E ES. In thin film epitaxy, E ES is important for controlling the inter - terrace mass transport and Em for intra - terrace mass transport. 2. 1 .2 Growth Modes Depending on the st rength of adlayer - substrate interactions, four main growth epitaxial modes are commonly observed: Vollmer - Weber(VW ) growth ( island growth , 3D morphology ), Stranski - Krastinov (SK) growth ( layer - plus - island growth , change from 2D to 3D morphology after criti cal thickness ), Frank - van - der - Merwe(FM) growth ( layer - by - layer growth , 2 D morphology ) , and s tep g low growth . T he schematics for these growth modes along with r eflection h igh - e nergy e lectron d iffraction (RHEED) oscillation profiles are shown in Figure 3. 6 . 132 The classifi c ation of the first three growth modes are consider ed by three macroscopic surface energy : (free energy of the substrate), f ( free energy of the interface), ( free energy of the epilayer). 132 For an epitaxial film consist ing of n layer s, the FM growth mode occurs when : , ( 2 . 1 ) where f (n) is the interfacial energy for n monolayers . This suggests there is a surface energy gain when the ad layer convers the substrate. In this mode , ad - molecules can diffuse freely in an attempt to completely cover the high energy surface. For homoepita xy, that is = , this condition can be writ ten as f 0. The VW mode results when : , ( 2 . 2 ) 11 2.2 Epitaxy t - E pitaxial growth is achieved for lattice - matched materials and near - equilibrium growth conditio ns. Because of the former requirement, there is a significant limitation to possible materia l combinations. A key parameter for determining the success of a hetero - epitaxial system is the lattice misfit (or lateral strain) , f , defined as : , ( 2 . 3 ) where a S is the lattice constant of the substrate and a L is the lattice constant of the unstrained adlayer. If misfit values are too high (e.g. greater than several percent) , this leads to poor epitaxial growth with a high degree of peeling, cracking, or 3D islanding past the cr itical thickness and a high dislocation density below the critical thickness. 12 At the highest level, e pitaxy can be classified into two categories: homoepitaxy, when the film is deposited on a substrate of the same composition; and heteroepitaxy, when the substrate composition is different from the film. Based on the extent of lattice mismatch, epitaxy can be further classified as commensurate, pseudomorphic , coin cide nt, and incommensurate growth ( Figure 2.2 ) . Commensurate growth is based on ideal lattice matching between film and substrate . With lattice mismatch, the film s can exhibit pseudomorphic or incommensurate growth. In pseudomorphic growth, the film is str aine d (or structurally distorted) to achieve commensurism . With even larger lattice mismatch, the distortion of pseudomorphic growth causes increasing strain in the film to an extent that relaxation occurs (from monolayers to hundreds of nm) with misfit di slocations generated at the interface . Figure 2 . 2 . Schematic diagrams and transmission electron microscopy (TEM) examples of commensurate, pseudomorphic and incommensu rate growth in epitaxy. The s ubstrate is shown in blue and the adlayer is shown in grey. Figure reproduced wi th permission . 137 - 139 13 2. 3 Quasiepitaxy ( QE) , ( 2 . 4 ) 2. 3 . 1 Energetic Considerations Energy minimization is used for describing optimal q uasiepitaxial configurations at equilibrium conditions. Due to the rotational degrees of freedom and shallow potential landscape, this organic adlayer will rotate in - plane to minimize the total energy of this adlayer - substrate system. T he adlayer lattice v ectors are related to the substrate lattice, generally, via the tra nsformation matrix: 14 , ( 2 . 5 ) Figure 2 . 3 . Schematic modes of epitaxy and quasiepit axial overlayer alignments . (a) Schematic of molecular overlayer alignment showing real lattice vectors of the adlayer ( ) and substrate ( ) . Schematic modes of overlayer alignments showing (b) commensurate growth (epitaxy ), (c) in commensurate growth (quasiepitaxy) and (d ) point - on - line coincidence (quasiepitaxy). 15 , ( 2 . 6 ) ( 2 . 7 ) 16 ( 2 . 8 ) Figure 2 . 4 . Example of V / V 0 analysis for zinc phthalocyanine (ZnPc) on deactivated Si(111) - B s urface . The geometry potential model shows the fi r st minima at = 28±0.5° and gives V / V 0 value of ~ 0.7 indicating a QE relationship. The simulated azimuthal rotation angle matches well with experimental value of = 27±2 ° shown in the STM topography image . Note that the presence of two - dimensional Moiré pattern hi ghlights the incommensurate nature of the substrate and adlayer common to quasiepitaxial systems. Figure repr inte d with permission . 48 17 2. 3 . 2 Kinetic Considerations In the previous section, the adlayer - substrate interaction is only discussed in terms of energetic minimization, but there are also kinetic aspects involved. M aterials deposited onto a substrate experience several diffusion processes. There are two different typ es of diffusion coefficients : the chemical (or collective ) diffusion coefficient ( D c ) and self ( or tracer ) diffusion coefficient ( D ) . Chemical diffusion happens when there is a concentration gradient in the system while self - diffusion happens spontaneously without a gradient and can take place under equilibrium conditions 142 . Thu s, i n a very low molecular density case, D c =D. D is defined with the Arrhenius equation below: , ( 2 . 9 ) where D 0 is a pre - exponenti al factor , E m is the diffusion barrier. Typical values for E m are in the range of 1 - 50meV. Classical diffusion define s the mean square displacement from the starting position as being proportional to time as: , ( 2 . 10 ) Where the angular brackets denote an ensemble average over the equilibrium state in the system. For the 2D surface diffusion case, v = 2. ( v = 1 for 1D diffusion, v = 3 for 3D diffusion). This equation is used for simul atin g surface diffusivity discussed in Section 9.1 . The change in the free energy G during nucleus formation is described as 143 , ( 2 . 11 ) 18 w here is the change in the chemical free energy per unit volume for condensation (gas to soli d) , r is the radius of the nucleus and is the interfac ial energy per unit area . The relationship between G v and r is plotted in Figure 2. 5 . The e quilibrium radius of the nu cleus or critical radius is described as . At critical radius, the critical free energy change is G * = (16/3) 3 /( ) 2 , which is the barrier to nucleation. This critical radius represents the stability point where the cluster becomes stable by adding more molecules. For a 2D cluster, Equation 2.1 1 can be written as , ( 2 . 12 ) where a is the lattice parameter , and . is the vapor pressure at equilibrium related to growth temperature and is the actual vapor pressure during deposition as function of deposition rate . Thus, the 2D nucleation rate r n can be expressed as 144 Figure 2 . 5 . The plot of free energy change G as a function of r . G (red line) is balanced by the volume fre e energy change (green line) and interfacial energy change (blue line) due to formation of solid phase . 19 , ( 2 . 13 ) This equation can give insight into when nucleus form ation is favorable. When 2D nucleation rate is small enough (close to zero), new nucleus will stop to form, and the existing nuclei will grow larger from nuclei and step edge s . 2.4 Key Issues with Inorganic Halide Perovskite Growth 2. 5 Key I ssues with O rganic T hin F ilm G rowth inorganic 20 21 Chapter 3 Experimental Techniques In this chapter, variou s experimental techniques for thin film preparation and characterization used throughout this thesis are described in detail. C haracterization techniques are discussed in cluding diffraction methods , morphology , and optical measurement. 3. 1 Physical Va por Deposition (PVD) 3.1.1 Thermal Evaporation Thermal evaporation is a common method of physical vapor deposition (PVD). A simplified evaporator diagram is shown in Figure 3.1 . It is one of the simp lest forms of PVD and typically uses a resistive heat sou rce to evaporate a solid material in a vacuum environment to form a thin film. The materials are heated in baffled tungsten boats producing a high vapor pressure stream . The evaporated material traverses the vacuum chamber and coats the substrate . The requ ired deposition pressure to assure ballistic (line of sight) deposition is usually in the range of 10 - 6 torr , which is achieved by a two - stage pumping process: a rough mechanical pump to 10 - 2 torr and a cryopump further to10 - 6 torr . The deposition rate (usually 0. 01 1 nm /s) is monitored using quartz crystal monitors (QCM) above each source. Quartz crystals resonate at specific frequencies that depend on the total volume of thin film deposited on them. QCMs are calibrated to relate the thickness arrivin g on the QCM to the thickness on the subst rate by the ratio called the tooling factor (TF) . The equation below is used to calculate the calibrated tooling factor ( TF f ) : , ( 3 . 1 ) 22 w here t i is the initial measured thickness based on the initial guess of the tooling factor TF i and t f is the thickness measured on a reference substrate (usually undoped single - crystalline Si) using Figure 3 . 1 . Evaporator diagrams. (a) A simplified diagram of thermal evaporation chamber with one s ourc e (left). The labelled parts are: (1) source boat, (2) copper contacts , (3) source shutter, (4) quartz crystal monitor (QCM) , (5) su bstrate shutters, (6) substrate, and (7) rotation stage. During deposition, source and substrate shutters are open (righ t). (b) Schematic showing the positioning of the six evaporation sources and three QCM sensors located at the bottom in side the custom A ngstrom Engineering EvoVac deposition system. The substrate stage is above these sources. (E - beam and sputtering compone nts are also e quipped in the system.) 23 ellipsometry or atomic force microscopy (AFM) . Each QCM c an be us ed until the lifetime drops below 80% - 90% range. C o - deposition of two or three materials can be achieved by assigning each to a differen t QCM . Every source is tooled separately on a Si wafer using ellipsometry prior to the co - deposition. Note that the real thickness of the film grown on Si can be different from growth on other substrates ( particularly for alkali halide crystal s , for example) due to different wetting conditions . To achieve the desired stoichiometry (molar ratio), the rate (volumetri c) ratio for each precursor needs to be calculated from the molar mass and density. For reactive co - deposition, t he thickness of the reacted film needs to be measured by cross - sectional SEM as ellipsometry is not feasible on the optically transparent cryst alline substrates utilized in this work . Deposition control is accomplished using the Inficon SQS - 242 deposition software ( Figure 3.2 ) . The deposition process starts with preconditioning ( Ramp 1/Soak 1/Ramp 2/Soak) followed by a shutter delay for the grow th rate to stabilize at the set rate. Then the substrate shutter is open ed , and the substrate stage rotation is enabled to achieve uniform growth across the substrate . The software will con trol the output power to the source using closed loop PID control w ith feedback from the assigned crystal sensor. 24 3.1.2 E lectron - beam (E - beam) Evaporation Electron - beam (E - beam) is another PVD technique used in this work. In E - beam deposition, t he target materials are bombarded with electron beam created by a charged tu ngsten filament. The electrons are accelerated by the high voltage that is applied between the hearth and the filament. The focusing and the alignment of the electron beam is controlled by a strong magnetic field. The schematic of the E - beam dep osition pro cess is shown in Figure 3.3. The chamber pressure requirement is 10 - 5 torr and 10 - 3 for some modern systems. Compared with thermal evaporation, E - beam has many advantages. First, the E - beam material utilization efficiency is high which reduces cost. For ex ample, gold is deposited using E - beam for fabricating cavities for cavity - melt Figure 3 . 2 . Inficon SQS - 242 deposition software interface during a two - source co - deposition ( shown in manual mode). 25 deposition (CMD) discussed in Section 4.1.1.1 . Second, much higher heating temperature can be reached so that h igher deposition rates (up to 1 - 2nm/s) can be achieved. Third, bec ause the electron beam is only confined at the source material , the contamination from nearby components are largely reduced leading to higher purity of the deposited materials. 3. 2 In situ D iffraction T echnique: Reflection H igh - energy E lectron D iffractio n (RHEED) 3.2.1 RHEED Introduction Reflection high - energy electron diffraction (RHEED) is a powerful technique to study the evolution of surface structures. It has been successfully used in monitoring the growth of semiconductor thin films, 146 metals , 147 - 148 and complex oxides . 14 9 - 150 A schematic of the RHEED technique is shown in Figu re 3. 4 . A high - energy electron beam, typically 10 30 keV, is directed at the film surface at grazing inciden t angle (1 ~3 °), and the resulting diffraction pattern is capture d by a CCD camera. Because of the small incident angles, the electron penetration de pth is at the monolayer scale (0.5 - 1.0nm), making it a particularly surface sensitive technique. With such low penetration depth, the electron wave density in the normal direction is not sampled fo r perfectly smooth film s or crystals . In this case, the 3D reciprocal lattice devolves into Figure 3 . 3 . Schematic of the E - beam deposition process . 26 a 2D array of infinitely long lattice rods. The intercept of th e Ewald sphere with these rods corresponds to an allowed diffraction condition and streak ed patterns are formed. However , for rough crystalline sample s , the normal - direction electron density is sampled because of multiple small features on the sample surfa ce. Reciprocal lattice points emerge and lead to spotty features of the RHEED patterns. Typical diffr action patterns are shown in Figure 3. 5 correspond ing to different surface crystallinity . Because high energy electrons are used in RHEED, the diameter of the Ewald sphere ( k 0 ) is orders of magnitude larger than most reciprocal lattices (i. e. 20keV electro ns, k 0 = 785nm - 1 ), which means we can see the Ewald sphere as a plane cutting the reciprocal lattice. Kikuchi patterns in RHEED images can be seen in the case of thick single crystals . The formation of a Kikuchi lines is a two - step scattering process. First, electrons from the incident beam collide elastically (losing a small amount of energy) with the bulk sam ple and randomize the directions of their wavevectors. Then, if the energy los s is small, the resulting electron distribution Figure 3 . 4 . RHEED schematics. (a) Typical RHEED geometry. (b) Photo of the RHEED system installed in an evaporation vacuum chamber (Angstrom Engineering EvoVac Evaporator). (c) Ewald sphere construction with RHEED. Th k 0 d due to the structure factor width and crystalline defects . Figure reproduced with permission. 145 27 can then be rescattered inelastically and isot r opically from crystal planes leading to Kikuchi diffraction lines. These Kikuchi lines move in a continuous and coincident manner when the crystal is rotated , wh ich is a feature often used to align the azimuthal direct ion of the crystal with respect to the incident electron beam. Figure 3 . 5 . Example RHEED patterns. A smooth crystalline surface generally shows streak ed patterns while a rough crystalline surface generally shows spotty patterns. Ring patterns indicate 3D - disorder (or crystalline powder) and a diffuse halo pattern indicates the surface is amorphous . Figure reproduced with permission. 145 28 Growth modes can be deduced from RHEED oscillations represented by the intensity of RHEED specular beam versus growth time. Figure 3. 6 shows an illustration of the various growth modes with RHEED oscillations . In the 2D growth mode, the periodic formation and coalescence of 2D nuclei lead to the intensity oscillation ( Figure 3.6 c ) . Generally, one period of RHEED oscillation corresponds to one mol ecular layer or atomic layer 152 . In reality, the second layer w ould start to form before the growth of previous layer is fully complete, thus an overall decreasing in oscillation intensity is often observed. 3D growth will show a rapid damping in amplitude ( Figure 3.6 a ) . For step flow mode, the growth is viewed as steps traveling across the surface. With a relatively stable step density, the intensity will eventually reach a constant ( Figure 3.6 e ) . 3.2. 2 Electron Beam Damage Materials suffer damage a fter long exposure to high - energy electro ns . 153 The time - to - damage ( t damage ) of a material is defined as: t damage = C dose A /( qI) ( 3 . 2 ) Figure 3 . 6 . Growth mode with the RHEED intensity signal vs. time. The typical RHEED oscillation profile s are plotted belo w each growth mode schematic. Figure reproduced with permission. 151 29 where C dose is the critical dose in units of C/c m 2 , A is the electron beam area, q is the charge of the electron, and I is the electron beam current. In general, a minimum for the critical dose exists for a certain beam energy ( around 0.1 - 5keV ) for many organic molecules . From the equation, a direct way to increase the time before the film is damaged is to decrease the electron beam current. In the following chapters, in situ RHEED is demonstrated using ultralow current only in the nA range to capture images via a detector system with low f - number lense s and sensitive detectors , eliminating damage/charging even on organic and insulating layers over the measurement time ( over several hours) . This enabl es the monitoring of organic small molecule growth and perovski te growth on insulating halide crystals. 47 3. 3 Ex situ D iffraction T echnique 3. 3 .1 X - ray Diffraction ( XRD) X - ray diffraction (XRD) is one of the most important tools for characterizing structures, phases, texture, and other structural parameters such as average grain size, crystallinity, strain, and crystal defects. XRD peaks are produced by c onstructive interference of a monochromatic beam of X - rays scattered at certain angles from each set of lattice planes within the material. The typical Bragg - Brentano XRD geometry is shown in Figure 3. 7 . describes the diffraction condition for constructive interference , which is expressed as : ( 3 . 3 ) 30 where d is the interplanar spacing between parallel planes of atoms in the family (hkl) , is the angle of the incident beam with respect to these planes, n is an integer, is the characteristic wavelength of the X - rays . C u ( = 1.5406 Å) is frequently used o n labscale X - ra y instruments but monochromatic synchrotron sources were also utilized in this work ( : 0.0619870 nm) at the Argonne National Laboratory Beamline 24 - ID - C . The Advanced Photon Source (APS) at Argonne National Laboratory used a hot cathode heated to ~ to pro duce electrons and then accelerated to relativistic speeds in a linear accelerator. Further accel eration is performed in booster synchrotron which is a racetrack - shaped ring of electromagnets and 7 GeV can be reached within half a second . 154 - 155 3. 3 .2 Pole F igure Pole figure measurements are used to study the orientation distribution of crystallographic lattice planes and texture analysis via stereographic projections . To collect pole figures, the crystal structure of the sample needs to be well known so that the diffraction angle (2 ) corresponding to the hkl pole can be set for the scan . T he diffracted intensity is collected by varying two geometrical Figure 3 . 7 . XRD geometry. Two beams with identical wavelength and phase approach a crystalline solid and are scattered from separate ato mic planes . There is constructive interference (in - . 31 parameters : the tilt angle from sample surface normal direction ( ) and rotation angle around sample surface normal direction ( ). Th e collected data is plotted as a function of and . ( F igure 3.8 ) Figure 3 . 8 . Sphere of fixed - length scattering vector and stereographic projection. T he tilt angle from sample surface normal direction is the rotation angle around sample surface normal direction is . 32 In - plane pole figure measurement is a pole figure measurement p erformed using the in - plane arm. This method has two major advantages over the conventional pole figure measurement. F irst, diffraction from lattice planes perpendicular to the sample surface can be detected , which allows a complete mapping from =0° to 90 ° (rather than being limited to about 70 ° ) . Second, the sample is maintained horizontal in the case of a horizontal sa mple goniometer meaning this method Figure 3 . 9 . Example of pole figures. (a) Example pole figures for a randomly oriented sam ple , a textured polycrystalline sample that is (111) - fiber oriented and a single crystalline sample with a (111) surface. (b) Measured and simulated KCl (111) pole figure for KCl (100) substrate. 33 does not require the sample to be tilted, which eliminates the trouble of clipping the sample to the s tage to prevent them from falling. Examples of pole figures are shown in Figure 3. 9a . The p ole figure for a randomly oriented sample only has a uniform intensity of distribution due to the lack of preferred orientation. A textured polycrystalline sample that is (111) - fiber oriented will show strong distribution for {111} and {220} in a ring shape. The single crystalline sample with a (111) surface orienta tion shows strong intensity of distribution from {111} and {220}. If the 2 position for the pole figure scan is set the same as the surface orientation of the crystalline sample, a strong distribution will be located at the center ( =0°). Typically, distr ibution s at the center or edge ( =90°) are not ideal for pole figure scans because the s ymmetry obtained from non - surface orientations are more interesting compared to the know n distribution from the surface plane and th e signal on the edge is usually not strong . A n allowed 2 position other than the surface orientation with strong peak intensity and clear symmetry is usually chosen for pole figure scan. For example, KCl (111) at 2 =24.49° is chosen for the pole fig ure sc an of KCl ( 100) substrate. A fourfold symmetry is shown in Figure 3. 9b . It is also emphasized that pole figure measurements can only be obtained when the crystal structure of the film is well known so that the 2 position for a pole can be precisely set fo r the scan. 34 3. 3 .3 Rocking Curve Rocking curves are primarily used to study defects such as dislocation density, mosaic spread, curvature, misorientation, and inhomogeneity . This method is performed by rocking the thin - film sampl e while the detector is kept a fixed 2 angle to record dif fraction intensities from the preferentially oriented lattice planes. 157 The degree (or distribution) of preferred orientation is estimated from the full width at half maximum (FWHM) of the rocking curve profile. The FWH M f or a single crystalline sample are at 0.001° range, such as 0.003° for Si(220) 158 and 0.0082° for Strontium titanate (STO) (002) 159 . Organic crystalline samples suc h as pen tacene(001) show FWHM values between 0.08° to 0.09° for the thin film phase and 0.22°for the bulk phase. 160 An example of a rocking curve scan is shown in Figure 3 .10 b. Figure 3 . 10 . Rocking curve measurement geometry and an example rocking curve scan of (004) for Si (Ge) on Si. Figure reprinted with permission. 156 35 3.4 Thin F ilm M orphology M easurements 3. 4 .1 Atomic Force Microscopy (AFM) Atomic force microscopy ( AFM ) is a versatile and powerful microscopy technology for studying surfaces at the nanoscale. It has the advantage of imaging almost any type of surface, including polymers, ceramics, composites, glass, and biological samples and overcome s the drawback of sc a nning tunneling microscope (STM) imaging which is limit ed to conducting or semiconducting surfaces . The typical AFM set - up is shown in Figure 3.11 . AFM uses a low spring constant cantilever to image the sample surface . At one end of the cantilever, a shar p tip is fabricated using semiconductor processing techniques. The cantilever moves backward and forward above and cross the sample surface. The force exerted on the tip varies with the difference in the surface height and thus leads to the bending of the cantilever. A laser beam gets constantly reflected from the top of the cantilever towards a position - sensitive photodetector . The deflection is tracked and used to calculate the actual position of the cantilever . In th is way , AFM records a three - dimensiona l image of the surface topography of the sample under a constant applied force (n N range ) without causing any damage to the sample surface. Most commercial AFM instruments can reach a vertical resolution as low as 0.01 nm for more rigid cantilevers and the lateral resolution is related to the tip condition, but it is usually lower. 161 Common scan modes are contact mode, non - contact mode and tapping mode. In contact mode, c ontamination to the tip o ften happen s and lead s to distortion to the image . N on - contact mode generally has lower resolution. Tapping mode is most widely used to achie ve higher resolution without causing destructive damage to the sample surface. 36 3. 4 .2 Scanning Electron Microscope (SEM) A s canning electron microscope (SEM) is a type of electron microscope that produces topography images of a sample by scanning the surface with a focused beam of electrons ( Figure 3.12 ) . The resolution of the SEM is can go to 1n m . Both backscattered electron s (BSE) and secondary electron s (SE) can be used for SEM imaging. In the most common SE i maging mode, SE emitted by atoms excited by the electron beam are detected using an Everhart - Thornley detector. The number of SE detected depends on the topogr aphy of the sample. When the incident e lectron enters the sample, SE are produced from the emission of the valence electrons of the constituent atoms in the specimen. Due to low energy (<50 eV) of the SE, those generated deeper are absorbed while those on the surface are emitted, thus SE are ve ry surface sensitive. The incident angle of the electron beam will affect the amount of emission and therefore creating brightness differences in an SEM image. Nonconductive specimens collect charge when scanned by th e electron beam, thus they are usually coated with gold, gold/palladium alloy, etc. via sputtering. Figure 3 . 11 . Typical set - up of AFM. The l aser beam is reflected off the back of the tip onto a quadrant photodiode. The deflection of the cantilever is measured as it tracks the surface to obtain the sample topograph y . 37 3. 5 Thin F ilm O ptical M easurements 3. 5 . 1 Ultraviolet - visible S pectroscopy (UV - Vis) Ultraviolet - visible (UV - V is ) spectroscopy can be used to measure the t ransmission in the UV visible - infrared region of the electromagnetic spectrum of thin films coated on substrates or solutions contained in cuvettes. The schematic of a dual - beam spectrometer is shown in Figure 3.1 3 . A beam of light from a visible and/or UV light source ( wavelength from 300nm to 19 00nm) is separated into its component wavelengths by a diffraction grating. Each monochromatic (single wavelength) beam in turn is split into two equal intensity beams by a half - mirrored device. The sample beam pas ses through the thin film sample or a cuvette containing a solution of the compound being studied in a transparent solvent. The reference beam passes through nothing for abs olute thin film measurements or an identical cuvette containing only the solvent. N ote that for thin film transmission measurements, reference samples (e.g., clear glass slides) should not be utilized since the reflection from the glass/air interface canno t be properly separated by simple Figure 3 . 12 . Basic construction of a SEM system. 38 subtraction. The intensities of these light beams are then measured by electronic detectors and compared . The intensity of the reference beam is defined as I 0 and t he intensity of the sample beam is defined as I . Absorptio n may be presented as transmittance ( T = I / I 0 ) or absorbance ( A = log I 0 / I ). 3. 5 .2 Variable A ngle S pectroscopic E llipsometry ( VASE) Variable angle spectroscopic ellipsometry (VASE) is a technique used to characterize thin film thicknesses and optical constants. It measures the change in polarization of reflected light. A monochromatic be am of light is polarized and sent through a fiber optic directed at a thin film sample on a reflective surface. The light reflected off the sample is collected at a detector, and, based on the phase difference and light attenuation of the reflected ligh t, the ellipsometric angles (amplitude component) and (phase component) can be calculated: 162 ( 3 . 4 ) Figure 3 . 13 . Schematic of a UV - Vis spectrometer configured to measure transmission for a thin film sample. 39 where is the complex reflectance, r p and r s are the reflectivity of p - and s - polarized light , respectively , and are a polar description of the ratio of reflectance for p - and s - polarized light . A Cauchy model is used to fit the d ( film thickness ) , optical constants n and k (refractive index and extinction coefficient) represented as: 162 ( 3 . 5 ) ( 3 . 6 ) To reduce the model to three fit parameters, only the spectroscopic region with no absorption is used ( k =0). Thus, from the measured and for a wavelength range, n and d can be fitted. The measured thickness can be used to calculate the tooling factor (discussed in Section 3.1 .1 ) of thermally evaporated material such as organic semiconductors, charge transfer (CT) compounds, alkali m etal halides, halide perovskites, etc. 40 Chapter 4 Organic Homoepitaxial and Hetero - quasiepitaxial Growth This chapter focuses on understanding the homoepitaxial and hetero - quasiepitaxial (QE) g rowth of classical organic semiconductors . For homoepitax ial growth, two methods for fabricating single crystalline substrates were developed and the homoepitax y of 9,10 - diphenylanthracene ( DPA ) was studied . This is the first demonstration of studying homoepitaxial growth of organic crystals using RHEED . For het ero - QE growth, v arious organic - organic growth pairs with low and high energy mismatch of lowest energy surface ( LES) were explored to understand the design rules for ordered organic crystalline structures. 4 . 1 Organic Homoepitaxial Growth 4 . 1 . 1 Organic Sin gle Crystal Substrate Growth 4 . 1 . 1 .1 Cavity - melt Deposition (CMD) Method C avity - melt deposition (CMD) was utilized to fabricate crystalline substrates of controlled thickness over several cm 2 substrates which can be used for subsequent homo - epitaxial or h etero - QE vapor growth of organic films. 41 Figure 4 . 1 . CMD process for fabricating single crystalline organic substrates. (a - d) Schematics for CMD process , (e) Photo of the assembled cavities , (f) Photo of the assembled cavities placed in molten organic material in a 3 - zone furnace . 42 Using this method, various organic crystalline substrates were successfully prepared . XRD patterns and cross polarized optica l images of the fabricated organic substrates are shown in Fig ure 4.2 . The cross polarized optical imag ing is a direct way to visualize the grains and grain boundaries . Anthracene (001), - - Di(1 - naphthyl) - - diphenyl - - biphenyl) - - - N PD ) (101), perylene (100), triphenylene (200) and DPA (020) organic substrates are successfully fabricated showing single crystalline growth with only one preferred orientation . The fluora nthene substrate is polycrystalline. The melting points and CMD resul ts of the materials explored are summarized in Table 4 .1 . V arious parameters (pressure, temperature, cooling rate, purity of the material, etc.) can be optimized in order to prepare true single crystalline organic substrates. Figure 4 . 2 . XRD patterns and cross polarized optical images of the organic crystalline substrates prepared using CMD. The black square marked the Au (111) peak (from the metal spacers) at 2 =37.18°. 43 4 . 1 . 1 .2 Vapor Growth Method The drawback of using CMD to fabricat e organic single crystalline substrates for organic homoepitaxy or hetero - quasiepitaxy lies in the difficulties of performing in situ and real - time RHEED characterization because t he glass substrates used for cavities c reate heavy charging when electron beam is directed onto the surface . Furthermore , this method is unable to provide thickness control < 500nm and is unable to support gro wth of more than one layer. To overcome the charging problem , a v apor growth method wa s also developed to create small bulk single - crystals as substrates that could be mounted on conductive surfaces . The vapor growth of DPA single crystals ( lateral length up to 1.5 cm ) wa s demonstrated as an example . DPA is a promising candidate for device applications as high mobilities have been reported up to 34 cm 2 V - 1 s - 1 . 163 The source boat containing DPA powder was placed in a long quartz sleeve at the high temperature zone. A short one - side opened quartz sleeve was placed on the low temperature zone for crystal collection. This separated short sleeve makes it convenient to remov e the crystals. Another one - side opened quartz sleeve w as placed at the end of high temperature zone to confine Table 4 . 1 . List of organic materials for CMD. p˚ 44 the organic vapor during heating. All the inner sleeves were placed in a long quartz tube that can be evacuated and filled with N 2. This long q uartz tube was then enclosed in a programmable PID controlled 3 - zone tube furnace. The high temperature zone is set to 10 - 15 below the melting point of the organic material . The melting point of DPA is 245 - temper ature zone wa s set to ~ 223 , which fully melts the mate r ials if maintained for 5 - 10min . The low temperature zone is set to 10 - the high temperature zone. When the DPA fully melt ed in the boat, the inner sleeve s were filled with DPA vapor and crystallization initiates on the inner walls of the sleeve in the low temperature zone. The temperature was maintained until enough crystals were formed or the boat containing the molten DPA became empty. Figure 4 . 3 . Vapor growth method for fabricating single crystalline organic substrates. (a) Sche matics for vapor growth of DPA crystals. The open ends of the inner sleeves are indicated by dashed lines. (b) Photo of the vapor growth in a 3 - zone furnace. 45 4 . 1.2 Homoepitaxial Growth of 9,10 - Diphenylanthracene (DPA) Homoepitaxial growth of DPA was performed in our custom thermal evaporation chamber (Angstrom Engineering) equipped with reflecti on high energy electron diffraction (RHEED) system (STAIB Instruments) s hown in Figure 4. 4 a . The DPA single crystal substrate wa s prepared using the vapor growth method discussed in Section 4. 1 . 1 .2 . Due to the fragility of the crystals, t hey were carefully attached to the growth stage using a tiny drop of ionic liquid (1 - e thyl - 3 - methylimidazolium dicyanamide, Sigma Aldrich) instead of conductive carbon tape used for thicker metal halide single crystal substrates used in Chapter 7 and 8 . The ionic liquid is important to reduce charging and to reduce stress on the crystal, an d it has a low vapor pressure to prevent outgassing. Figure 4 . 4 . Schematic of DPA homoepitaxial growth. (a) Schematics for homoepitaxi al growth of DPA crystals. (b) Crystal structure of DPA (monoclinic, a = 0.94 nm, b = 2.03 nm, c = 1.00 nm, = 90.00°, = 113.00°, = 90.00°. 164 (c) C ross polarized optical images for single crystal DPA substrate (area: ~15mm 2 ). 46 Cross polarized optical imag es were taken using a simple rotational polarize and analyzer kit (360° rotation) prior to loading in the evaporator . Out - of - plane XRD and rocking curve measurements were performed 40 kV and 44 mA and using a Ni filter. Atomic force microscopy (AFM) was performed in contact mode for ex situ film morphology characterization. A silicon tip coated in Ti/Ir was used for the AFM measurements (Asylum Research). Figure 4 . 5 . Structural characterization of the DPA single crystal substrate. (a) XRD, (b) rocking curve analysis and (c) Rotation dependent RHEED of the of DPA single crystal substrate prepared using vapor growth method. 47 Figure 4 . 6 . RHEED for DPA homoepitaxy of 20nm at 75 ° C . (a) RHEED pattern transition. (b) Specular intensity change vs growth time. 48 Figure 4 . 7 . AFM images of DPA homoepitaxy. (a) Bare DPA substrate and homoepitaxial growth (20 nm) performed at (b) room temperature, 0.05nm/s , (c) low temperature ( - 0.05nm/s and (d) room temperature, 1nm/s. 49 4 . 2 Ordered Organic - organic Hetero - quasiepitaxial (QE) Growth For a lternating hetero - QE growth , we utilize the same deposition and in situ RHEED set up shown in Figure 4.1a . We focus on looking at various organic molecules (listed in Table 4.2 ) paired with NTCDA to study the design rules for ordered alternating growth of organic crystalline layers . NTCDA and the paired material is deposi ted alternating onto a s ingle - crystal KBr (100) 50 substrate freshly cleaved prior to the growth. A n alternating growth structure containing t wo growth cycles is used in this study shown in Figure 4. 8 i . O ut - of - plane XRD and AFM characterization were uti lized to measure the stacking orientation and growth features respectively . A range of surface energy systems with different energy mismatch between their LES are studied t o understand the design rules for ordered organic - organic crystalline hetero - QE growth and the ro le of this mismatch plays in growth modes. Specific combination s for energy - 1) matched, 2) mismatched - low, and 3) mismatched - high are explored. Pair growth of NTCDA with the chosen small molecules are explored in the order of increasing energy difference of their LES with NTCDA(100): 5,10,15,20 - tetraphenylbisbenz[5,6]indeno[1,2,3 - - lm]perylene (DBP) (energy - matched), C 60 , tetracene, pentacene (energy mismatched - low ) , coronene, copper phthalocyanine (CuPc) and perylene (energy mismatched - high). XR D is performed after one growt h circle to determine the crystalline orientation of the film. Despite the small energy Table 4 . 2 . Calculated surface energies (SE) of LES) for a range of organic crystals including NTCDA and materials for paired alternating growth. Isolated surface energies of different planes were calculated with Materials Studio v7.0 (Morphology module ) using the equilibrium method with the universal force field . The materials are listed in the order of increasing S E difference with NTCDA (100). The values for acenes agree with reported values. 166 - 167 *Second lowest energy surface. 51 mismatch between LES of DBP and NTCDA , t he L3 for DBP growth pair is polycrystalline with ring - like RHEED pattern and L4 is completely amo rphous with diffuse halo patte rn (Figure 4. 8 b) . This suggests DBP cannot support multilayer ordering pairing with NTCDA. The energy mismatched - low set shows ordered crystalline growth for at least two pairs. After the second growth circle, the L4 for C 60 , tetracene, pentacene all show clear rotational depend ent RHEED patterns indicating good in - plane order . AFM shows that the L2 for C 60 set is smooth and the grain size as well as orientation is maintained after L4 ( Figure 4. 10 a, g ) . RHEED patterns suggest smooth crystalline growth for both L2 and L4 of C 60 set. Both the L2 and L4 for Figure 4 . 8 . In situ RHEED patterns of the alternating growth pairs . (a) the bare K Br substrate, (b - h) alternating growth of NTCDA with DBP, C 60 , tetracene, pentacene, coronene, CuPc and perylene, (i) schematic of bottom - up alternating growth. (L is short for layer). Crystal structure of the KBr substrate and molecular structures of the growth materials are shown (C: gray; H: white; O: red; N: blue, K: purple; Br: maroon and Cu: pink). 52 tetracene set show needle - like growth from the AFM, where the preferred alignment of grains is more obvious for L4( Figure 4. 10 b, h ). For pentacene, AFM shows th at L4 has two p referred Figure 4 . 9 . XRD patterns of the alternating growth pairs . X - ray diffraction characterization of the b are KBr substrate(K), NTCDA growth on KBr(K/N) and growth of DBP, C 60 , tetracene, pentacene, coronene, CuPc and perylene on NTCDA. All the materials show crystalline growth from the LESs on NTCDA (100) surface (except for DBP that grows amorphously). The p entacene (001) peak is not shown in the XRD scans, but this LES growth direction is confirmed by RHEED. 53 orientation perpendicular to each other ( Figure 4. 10 i ). Each of the tetracene and pentacene layer s show clear spotty RHEED patterns indicating the rough crystalline surface. Figure 4 . 10 . AFM images of crystalline alternating growth pairs. 20nm of (a,g) C 60 , ( b,h ) tet racene (c,i) pentacene (d,j) coronene (e, k) CuPc (f, l) perylene growth on NTCDA after the first (L1 - L2) and second (L3 - L4) growth circle. 54 The energy mismatched ( high ) set show little ordering beyond one pair suggested by the RHEED for L4 . F or coronene set, L2 shows needle - like growth but the orientation is random ( Figure 4. 10 d ) , which agrees with the ring patterns observed over the streaks in RHEED ( Figure 4. 8 f ). Non - needle - like grains are seen for L4 ( Figure 4. 10 j ) and crystalline features from RHEED faded into a diffusive halo pattern ( Figure 4. 8 f ). AFM shows one major grain a lignment for L2 of CuPc set and RHEED shows spotty patterns indicating a rough crystalline surface ( Figure 4. 10 e ) . However, the L4 of C uPc shows random grain orientation ( Figure 4. 10 k) and shows a corresponding diffuse ring feature from RHEED ( Figure 4. 8 g ) . For the perylene set, the L4 AFM features look similar to pentacene ( Figure 4. 10l ) but the RHEED pattern shows rings , indicating a po lycrystalline structure is developed ( Figure 4. 8 f ). In sum, these results indicate that multilayer order is indeed likely supported by surface energy matching , requiring a matching in the range of 10 - 20% . This is likely a good starting point for designin g organic - organic multilayer structures moving forward and strengthens the hypothesis that surface energy matching is a key criteria on order ed multilayer growth (in contrast to lattice matching in traditional epitaxy) . 70 4 . 3 Conclusion 55 56 Chapter 5 Uncovering a New Growth Mode: Edge - Driven Organic - on - Organic Quasiepitaxy In this chapter, we demonstrate the ordered alternating hetero - quasiepitaxial growth of 1,4,5,8 - naphthalenetetracarboxylic dianhydride (NTCDA) and diindenoperylene (DIP) crystalline films on single - crystal substrates from the botto m up exploiting a new edge - interaction driven mechanism. Each layer is shown to grow wit h a well - defined quasiepitaxial registry regardless of the incommensurate unit cells, where the multilayer pair growth is maintained for multiple cycles. While there is a close surface energy matching between the lowest energy surface ( LES ) of DIP and NTCD A, it is a non - lowest energy plane of DIP that is formed in the multilayer structure. While surface energy matching of the lowest energy surface is likely to play an im portant role in many other organic - organic systems, the ability to grow this multilayer system with sustained ordering provides an intriguing new mechanism in organic - organic quasiepitaxy. E pitaxial superlattice structure has had significant importance in inorganic quantum wells and quantum dots applications, such as photodetectors, light emi tting diodes, and solid state laser diodes . 14, 168 - 173 The organic equivalent could have similarly substantial impact but has not been widely accessible previously. Thus, this new growth mode for achieving quasiepitaxial multilayers could provide novel pathways to fabricate highly ordered organic superlattices with tunable orientation for enhanced excitonic electronic devices. 5 . 1 Experimental Vapor deposition of each material was performed in the multi - source thermal evaporator described in Section 3.1.1 . Growth was carried out under a base pressure less than 3×10 - 6 torr at a rate of 0.03 nm/s at room temperature on freshly cle aved single - cryst al KBr (100) substrates. The d eposition temperature was monitored with a thermocouple integrated into the substrate holder 57 that was calibrated using an in situ non - contact infrared - to - analog converter module (Omega, OSM101). After the depo sition of each la yer, growth is paused to collect diffraction patterns at various azimuthal rotation s and to allow for shutter switching. In - plane lattice constants of the organic layers were measured from the RHEED patterns with the initial KBr pattern as a reference. The o ut - of - plane lattice structure was confirmed by X - ray diffraction (XRD) in the Bragg - (AFM) was performed in contact mode to study the film morpholog y and early stage nucleation. A silicon tip coated in Ti/Ir was used for the AFM measurements (Asylum Research). Isolated surface energies of different planes were calculated with Materials Studio v 7.0 (Morphology) using the equilibrium method with the uni versal force fiel d, validated against standard systems discussed elsewhere 70 . To understand the predicted equilibrium align ment of the organic - inorganic and organic - organic lattices, full surface potentials between substrate - adlayers were calculated in Matlab by summing the non - bonding interactions (van der Waals) between every atom - atom combination in respective layers . 70 The film lattice constants used in the calculations were obtained directly from R HEED measurements. The equilibrium separation is first determined for a fixed adlayer lattice (e.g. 5 × 5) over an excess of substrate lattices (50 × 50). The full system energy is then calculated as a function of translation across an entire unit cell at every rotation angle, . Thus, a plot of the minimum energy at every rotation can identify the overall equilibrium orientation registry. For reference, the rotation angle is defined as the angle between one of the two real lattice vectors of the adlayer ( ) and substrate ( ) (inset in Figure 5.5 b ). Calculations were repeated for surface mesh sizes of 5 × 5 and 9 × 9 adlayer meshes over 50 × 50 substrate meshes, respectively, to avoid edge effects. This method has been used previously to successfully i dentify nearly all known 58 quasiepitaxial lattice registries . 70, 141, 165 5 . 2 NTCDA and DIP NTCDA is a wide bandgap semic onductor, which we have already demonstrated to grow quasiepitaxially on KBr crystals with excellent ordering and registry at room temperature 70 . The archetypal perylene derivation, DIP, was paired w ith NTCDA due to its 1) promising electron and hole mobilities in thin - film transistors >1cm 2 V - 1 s - 1 ; 176 2) good crystallization behavior, stability against decomposition and oxidation at high temperature, and moderate photoluminescence efficiency ; 176 - 178 and 3) close surface e n ergy matching of LES to that of NTCDA. A number of studies have investigated DIP growth on inorganic surfaces like SiO 2 179 - 183 showing an upright crystalline orientation with little in - plane ordering. The 3D growt h of upright DIP has also been observed on organic layers including F 16 CoPc and F 16 CuPc 184 - 185 with in - plane disorder. Bulk structures for NTCDA and DIP are shown in Figure 5 .1. Figure 5 . 1 - - DIP. (a)The bulk structure for NTCD A s - DIP (a = 1.166nm, b = 1.301nm, c = - DIP (a = °) . 174 - 175 59 5 . 3 Ordered Alternating Crystalline Growth A typical series of reflection high - energy electron diffraction (RHEED) patterns for the growth of NTCDA/DIP alternating layer structures grown at room temperature are shown in Figure 5 .2 . Although th ere is no lattice matching between lay ers, RHEED patterns reveal that each layer is highly crystalline with its own well - defined surface lattice and registry. Schematic Figure 5 . 2 . In situ RHEED patterns of ordered NTCDA/DIP alternating growth. In situ RHEED patterns for the (a) the bare KBr substrate, (b,c) first, (d,e) second and (f,g) third pair. (h,i) AFM images of the first two layers. Note that as the number of layers increases, the streak ed patterns for the NTCDA star t to become less streak - like indicating an accumulating surface roughness. Crystal structure of the KBr substrate and molecular structures of NTCDA and DIP are shown (C: gray; O: red; K: purple; Br: maroon and H: white). 60 structural models of the NTCDA and DIP multilayer heterostructure are shown in Figure 5 .3 , along with the schematic of the in dividual layer orientations for growth directly on glass. When NTCDA is grown on the KBr substrate, the RHEED patterns show continuous streak features (similar to the starting substrate), indicative of a crystalline flat surface that is consistent Figure 5 . 3 . XRD characterization and structure models. X - ray diffraction characterization of (a) NTCD A(20nm) on glass. (b) NTCDA powder. (c) DIP(20nm) on glass. (d) DIP powder. Growth on glass substrates exhibits the LES for NTCDA(10 - DIP(001). (e) DIP(30nm) on NTCDA(20nm) on KBr. (f) NTCDA(10nm) on KBr. (g) DIP(30nm) on KBr. (h) the bare KBr subst rate. The peaks at 2 =11.80° belongs to NTCDA(100). The peak at 2 =20.68° is the - DIP(020) peak and the other peaks between 22.5 - 25 ° are satellite peaks from the KBr(200) peak at 27.00°. Structural models of (i) NTCDA growth on glass. (j) DIP growth on gl ass. (k) NTCDA/DIP alternating growth on single crystal KBr. (N is for NTCDA and D is for DIP). 61 with m easurements with atomic force microscopy (AFM). The diffraction streak spacings vary with different azimuthal angles indicating the ordered single - crystalline - like growth. XRD patterns show that the LES of NTCDA, the (100) plane, is formed parallel to the KBr and glass substrates ( Figure 5.3a, f ). The bulk lattice of NTCDA with this orientation has a rectangular unit mesh with lattice constants b 1 =0.531nm and b 2 =1.257nm , while the measured lattice from RHEED is b 1 = 0. 534(±0.001) nm and b 2 =1.319 (±0.002) nm, ind icating a slight distortion as seen previously. With the subsequent growth of the DIP layer at room temperature, long unbroken NTCDA diffraction streaks are eventually replaced by spotty diffraction patte rns of DIP which indicates a roughening of the cry stalline surface. Nonetheless, clear azimuthal registry is still observed and the single - crystal - like rotational dependence of the RHEED patterns indicates that the DIP continues to remain aligned in - plane - DIP(001) - DIP(020) plane is formed on the NTCDA surface, which is confirmed by b oth RHEED and XRD data. From the RHEED patterns of the first DIP layer, we measure out - of - plane spacings of d (020) =0.868(±0.007)nm, which is close to the spacing of d (020) =0.855 nm for the bulk lattice. The (010) peak, which is not typically an allowed peak, is in fact observed under the dynamical conditions with e lectron diffraction from multiple scattering events in RHEED . 186 In contrast, the (020) se ries of peaks are observed in the XRD pattern where dynamical scattering is not appreciable, and the measured d (020) =0.859( ±0.003 ) nm from XRD is in good agreement within the error of the RHEED data. DIP is also grown directly on glass as a comparison showi ng the LES orientation (upright DIP molecules) with d (001) =1. 684 nm, which is in stark contrast to the growth on NTCDA. While typical Figure 5 .3 d ) , growth and the LES - DIP(001) is formed on glass ( Figure 5 .3 c ) . From the RHEED patterns of 62 the first DIP layer, measure d unit mesh dimensions ar e b 1 =0.712(±0.001)nm, b 2 =1.734(±0.002)nm and =94.1(±0.3) °, which shows only a very slight distortion compared to the bulk surface structure of b 1 = 0.717 nm, b 2 =1.680nm and =92.42 ° . 5. 4 Reversed Order Growth To further understand this multilayer orderin g, growths were performed with a reverse layer sequence, starting with the DIP layer. Interestingly, the RHEED patterns in Figure 5 .4 b no longer - DIP(120) plane is instead observed, Figure 5 . 4 . Reverse order growth. In situ RHEED patterns of (a) the bare KBr substrate, (b,d) the DIP layer (10nm) on KBr and (c,e) subsequent growth of NTCDA (10nm) on DIP. Note that the diffraction pattern of (b) for DIP on KBr is di stinct to the pattern for DIP grown first on NTCDA ( Figure 5.2 c ) indicating a significantly altered molecular stacking orientation, - DIP (120) , that subsequently leads to a 3D disordered powder with NTCDA deposited on top. 63 which is an entirely distinct growth orientation. AFM shows the formation of long nanowire features which is also distinct from the growth o n NTCDA ( Figure 5 .4 d ). Subsequent growth of NTCDA on DIP/KBr results in ring diffraction patterns ( Figure 5 .4 c ) which indicat es the evolution of three - dimensional crystalline disorder (i.e. crystalline powder) after just the first layer pair ( Figure 5 .4 e ). 5. 5 In - plane Registry Modelling The measured azimuthal registry of NTCDA films with the underlying KBr substrate is consist ent with both full - structure potentia l energy and geometric model of exposing the LES parallel to the substrate on weakly interacting surfaces . 70 Indeed, when N TCDA and DIP are deposited individually on glass (SiO 2 ) substrates ( Figure 5 .3 a,c ), a weakly interacting surface, they both grow with the LES parallel to the substrate; these surfaces are NTCDA(100)(12.1 kcal mol - 1 nm - 2 - DIP(001)(10.1 kcal mol - 1 nm - 2 ), respectively. Consistent with this picture, we have found that many common organic molecules ( Section 4. 2 ), grow with their LES on both glass and this highly ordered NTCDA layer without a preferential nucleation at e dges. In the case of DIP, howeve - DIP (020) on NTCDA (100), which is an unexpected 64 and non - LES orientation. In this growth orientation, DIP molecules lay flat on NTCDA, leading to the potential for higher transport in the nor mal direction and one of the in - plane directions, which could be ideal for solar cell or anisotropic transistor applications. The results of the full - structure potential energy calculations predict an equilibrium re gistry between (020) DIP and (100) Figure 5 . 5 . Surface potential modeling of the DIP/NTCDA quasiepitaxial registry. (a) Measured real ( b 1 , b 2 ), and reciprocal ( b 1 *, b 2 *) lattice vector alignments for NTCDA and DIP layers. (b) V an der Waals potential energy as a functi on of adlayer - substrate azimuthal angle ( ) for DIP 5×5 (9×9) unit cells over an NTCDA 50×50 mesh. The azimuthal angle ( ) was defined as the angle between b 1 of the DIP adlayer ( ) and NTCDA substrate ( ). The minimum in the potential energy gives an equilibrium rotation angle (registry) of 65.9° compared to the measured angle of 0±5° (red arrows). Model of (c) the measured and (d) predicted real - space overlayer alignment for DIP on NTCDA . 65 NTCDA at a rotation angle of 65.9°, independent of the number of DIP cells ( Figure 5 .5 b ). However, the measured orientation of DIP on NTCDA was determined to be 0±5 ° from the RHEED data , which is in stark disagreement to the predicted registry. Figure 5 . 5 c,d sho w the schematics of the measured and predicted alignments. In contrast to NTCDA on KBr, both the full - structure potential energy ( Figure 5 . 5 b ) and geometric ( Figure 5 . 6 ) models fail to correctly predict the registry of (020) DIP grown on NTCDA or KBr, even though full potential modeling 141 has successfully predicted the registry of hundreds of organic - inorganic quasiepitaxia l systems including the registry of NTCDA on KBr in this work . 69 - 70, 140 - 141, 165, 187 - 190 This failure, and the emergence of a non - lowest energy surface suggest that there are other aspects that are dominating the ordered crystalline multilayer growth. Figure 5 . 6 . Calculated dependence of the normalized geometric potential, V/V 0 , as a function of the azimuthal angle for (020) DIP on (100) NTCDA. 66 5. 6 Edge - driven Mechanism Considering the elemental simplicity of these organic systems, this suggests that the DIP adlayer interacts with and nucleates from the step edges of the NTCDA. The absence of such edge nucleation in other acene molecules (which predominately show intra - terrace nucleatio n) points to a key distinguishing interaction between the NTCDA and DIP directly. I ndeed, the DIP has a perylene core capable of near covalent - like bonding to other aromatic rings, metals , 191 - 193 an d can support hy drogen bonding interactions , 194 which has been p reviously seen in PTCDA and exploite d for molecular templating . 195 - 197 This can create stronger interactions between the DIP molecule and the flat facing edge of the NTCDA step ( Figure 5 . 7 b ). From the RHEED patterns of the multilayer growth, the LES (001) of the DIP is indeed growing from either the (002) or (102) plane of NTCDA and DIP(110) from the NTCDA(011), which orients the DIP cofacially with the NTCDA edges. The initial monolayer gro wth of DIP is shown in AFM images for DIP on NTCDA ( Figure 5 . 7 a ), where, indeed, we observe a clear evidence of DIP nucleating from the step edges of NTCDA and subsequently filling individual NTCDA terraces. Despite a coverage nominally approaching 1 ML (a nd a uniform deposition flux/coverage observed concurrently f or DIP on SiO 2 Figure 5 . 7 . AFM image of early - stage nucleation and structural model of the step edge growth. (a) AFM image of a sub - monolayer film of DIP (< 1nm) deposited on NTCDA (10nm). The DIP nuclei are marked with white arrows. (b) Schematic o f DIP molecules ly ing cofacially with the NTCDA molecules at terraces showing edge nucleation. Interactions between molecular cores are highlighted in blue. 67 here and as in ref. 198 ) we do not find DIP isla nds uniformly distributed on every NTCDA terrace, and find pa rticularly few DIP islands on the largest terraces. This indicates that there is both a small Ehrlich Schwöebel barrier 133, 199 for DIP molecules to reach different terraces and there is high diffusivity for DIP on NTCDA at room temperature. Th is high diffusivity also likely plays a role in the ability t o observe this growth mode, similarly to the role that diffusivity plays in the transition between layer - by - layer and step flow growth in inorganic systems. The fact there is a layer sequence de pendence (whether DIP or NTCDA is grown first) for the first layer, but not subsequent layers, further confirms the role of NTCDA edge - nucleation in guiding this quasiepitaxial growth NTCDA plays an important role in initiating a growth orientation capab le of supporting multilayer quasiepitaxy. This highlights the impact that the nucleation environment can play as a dominant mechanism in the structural evolution for a range of organic molecules designed to interact more strongly with molecular moieties. I t is likely that by targeting these properties, such as hydro gen bonding or aromatic interactions between molecules at step edges, other systems can be tailored to show a similar growth mode. 5. 7 Conclusion In this chapter , we demonstrate the ordered hete ro - quasiepitaxial superlattice growth of two distinct organic crystalline films on single - crystal substrates via deposition from the bottom up, reminiscent of molecular beam epitaxy of inorganic quantum wells and quantum dots. The crystalline alignment, mo nitored with in situ RHEED, can be maintained for many pairs, where DIP grows on NTCDA exposing its non - - DIP (020) plane. This surprising growth configuration is shown to derive from edge - nucleation driven quasiepitaxial dynamics confirmed by in situ diffraction, equilibrium energetic predictions, gro wth order reversal, and early growth studies. The understanding of this non - equilibrium quasiepitaxial growth of crystalline organic 68 multilayers could provide new routes to controlling molecular orientation in ordered organic superlattices via control over nucleation dynamics through the design of step - edge molecule i nteractions. Therefore, t hese results could enable entirely new opportunities for enhancing unique excitonic tunability and could also be used as a platform to study organic exciton confinement and strong coupling. 69 Chapter 6 Homoepitaxial Growth of Alkali Halide Crystals In this chapter, w e investigate the homoepitaxial growth of NaCl on NaCl (001) investigated with in situ reflection high - energy electron diffraction (RHEED) and ex situ atomic force microscopy (AFM). Epitaxial growth is explored as a function of temperature and growth rate which provides the first detailed report of RHEED oscillations for alkali halide growt h that uncovers previously unrecognized growt h modes. This chapter provides the foundation for the epitaxial growth of halide perovskites on metal halide substrates described in the subsequent chapters. 6 .1 Introduction Alkali halide crystals are widely used in many optical and optoelectronic appl ications. NaCl crystals are used as one of the most common infrared transmission windows for spectrophotometers and useful for energy detection in x - ray monochromators. 200 - 201 Do ped NaCl bulk crystals have been used for single crystal emitters and scintillation detectors 202 - 204 and thin film NaCl can be used as buffer layers for OLEDs. 205 - 206 NaCl crystal surfaces have been used for studying epitaxial growth of organic and inor ganic thin films. 207 - 211 The large range of lattice constants of alkali halide crystals could also enable growth of epitaxial metal - halide perovskites. Thus, understanding fundamental aspects of homoepitaxial metal halide growth could have a signif i cant impact. RHEED intensity oscillations have been routinely used during epitaxial growth of metals and semiconductors to monitor the growth rate, 13 6 , 212 understand crystal growth mechanisms and surface reconstructions, 213 - 216 estimate surface diffusion, 152, 217 - 218 and measure dopant incorporation. 219 - 220 A better understanding of the characteristics of the oscillations would provide insight into the microscop ic processes during epitaxial growth. Although significant progress has 70 been made in studying RHEED oscillations in epitaxial growth on metals 221 - 224 and semiconductor s, 225 - 228 there are no reports on halide insulators. Hetero - and homoepitaxial growths of alkali halides using molecular beam epitaxy (MBE) have been studied by both electron and He atom scattering, however with a limited growth parameter range in ultrahigh vacuum. 229 - 235 In the case of the studies with RHEED in MBE, electron beams had to be removed within seconds to avoid charging, therefore limiting the ability to monitor oscillations and to obtain real t ime growth mode information. In the case of the studies with He sca ttering very little variation was found as a function of growth conditions. In this work, NaCl homoepitaxial growth on NaCl (001) surfaces is performed by thermal evaporation and studied us ing both in situ ultra - low current RHEED and ex situ atomic force m icroscopy (AFM), which reveal strong variations in the growth modes as a function of growth conditions. Layer - by - layer growth is observed at room temperature accompanied by clear RHEED osci llations while the growth mode transitions to an island (3D) mode a indicate a transition to a step - flow growth mode. To show the importance of such metal halide growth, gree n organic light - emitting diodes (OLEDs) are demonstrated using a do ped NaCl film with a phosphorescent emitter as the emissive layer. This study demonstrates the ability to perform in situ and non - destructive RHEED monitoring even on insulating substrates that could enable doped single crystals and crystalline substrates for a range of optoelectronic applications. 6 . 2 Experimental Vapor deposition of NaCl was performed in the multisource thermal evaporator (Angstrom Engineering) equipped with real - time and in situ RHEED system (STAIB Instruments) operated with an ultra - lo w (< 10 nA) current and beam energy of 30.0keV. By using ultra - low current, the damage and charging of the film is negligible over the growth times investigated enabling the use 71 of this technique on organic and insulating surfaces. 236 RHEED oscillations are monitored with substrates fixed at various in - plane orientations (KSA400). Rotation dependent patterns were collected after growth w as halted. The NaCl single crystal subs trates were freshly cleaved in a glovebox prior to deposition with the (001) surface exposed and attached to the substrate holder using conductive tape. D eposition rates were measured in situ with a quartz crystal microbalance a nd calibrated ex situ by cro ss - sectional scanning electron microscopy (SEM) using single crystal potassium bromide as a substrate for layer contrast. The heating/cooling stage inside the chamber was used to pre - heat/cool the NaCl single crystal substrates to the required temperature before deposition. Substrate temperature was monitored and controlled (Watlow EZ - Zone) with a thermocouple integrated into the substrate heating/cooling plate that was calibrated to the surface temperature using an in situ non - c ontact infrared - to - analog co nverter module (Omega OSM101). Atomic force microscopy (AFM) was performed in contact mode for ex situ film morphology characterization. A silicon tip coated in Ti/Ir was used for the AFM measurements (Asylum Research). Electric al characterization was perf ormed using a digital source meter (Keithley 2420) and a picoammeter (Keithley 6487). Luminescence current was measured using a large area Si photodetector (Hamamatsu) and electroluminescence (EL) spectra were measured using a c alibrated spectrometer (Ocea n Optics USB 4000). 72 6 .3 Growth Modes of NaCl Homoepitaxy Experiments were performed to study homoepitaxial growth of NaCl on NaCl (001) surfaces as a function of growth rate and temperature. For NaCl, the mo nolayer (ML) and bilayer (BL) thicknesses are d efined as a /2 (2.8Å) and a (5.6 Å), respectively, consistent with other BCC and FCC crystals. 237 different rates to confirm RHEED oscillation interpretation: 0.1 BL/s, 0.3BL/s and 0.6BL/s. In the second set, low ( - re performed at a growth rate of 0.1 BL/s as co mparison to the room temperature growth to study the effect of temperature on the growth mode . RHEED patterns of NaCl homoepitaxial growth of 200Å at different te mperatures are compared in . A typica l RHEED pattern of the NaCl single crystal is shown in Figure 6 . 1 . Temperature dependent in situ RHEED patterns of NaCl homoepitaxy. Diffraction patterns for (a) bare single crystal NaCl cleaved along the (001) plane; homoepitaxial growth of 200Å NaCl at (b) - 120 , (c) 25 and (d) 120 . Note that the Kikuchi lines and continuous streaky patterns of the NaCl persist after growth at 25 a nd 120 , while at - 120 Kikuchi lines disappear, and streaky pattern becomes discontinuous after growth. 73 features of the patterns are still present after the growth with a clear single - crystal rotation dependence and Kik uchi patterns still observable. For growth at - observed with a corresponding disappearance of Kikuchi patter ns. Specular RHEED intensities are shown in ll show clear RHEED oscillations with a sudden specular inte nsity drop when the growth is initiated. The period of the RHEED oscillations is consistent with the deposition rate within the error of the deposition rate recorded by the quartz crystal microbal ance (±0.1 Å /s), where each period corresponds to a complete unit cell (or BL) similar to the oscillations observed for stoichiometric FeAl on AlAs. 238 As the growth proceeds there is an overall damping of t he specular intensity Figure 6 . 2 . RHEED oscillations of NaCl homoepitaxial growth vs. growth rate and t emperature. Room temperature growths as a function of growth rate: (a) 0.1 (±0.02)BL/s, (b) 0.3(±0.05)BL/s and (c) 0.6(±0.1)BL/s; fixed rate growths of 0.1BL/s as a function of growth temperature (d) - 120 , (e) 25 and (f) 120 . RHEED oscillations intens ities were recorded from the specular spot and correspond to one unit cell (or BL) per oscillation. Note that (a) and (e) are the same data. 74 shown in obs erved for > 25 BLs, along with a recovery in the specular intensity during longer deposition time s as shown in In contrast, the oscillations observed at low temperature gro wth at - . shows the contact mode AFM morphology of 200 Å of homoepitaxial NaCl growth for each growth temperature. The average step heights in the se AFM images correspond to either ML (~2.8Å ) or BL (~5.6 Å) steps where BL steps are only observed at low temperature (line scans shown in ) . At room temperature, intra - terrace nucleation is clearly evident by the emergence of isolated ML - thick circular islands speckled within terraces. W hen the growth temperature is reduced to - rougher terraces with greater terrace heights (closer to a BL) are found to form during growth Figure 6 . 3 . NaCl t emperature dependent homoepitaxial growth mode and 3D morphology. AFM images for ( a) bare single crystal NaCl and as a function of growth temperature for homoepitaxial growth of 200Å NaCl at ( b) - 120 , ( c) 25 , ( d) 120 . ( e) Crystal structure of NaCl. Line scans showing step height are i ncluded in the Supplementary Information. Two a verage step height s are observed: 2.5 Å (corresponding to the ML) and 5. 0 Å (corresponding to the BL), with the ML steps observed most frequently. 75 along with a large density of voids and no observable intr a - terrace nucleation. At room temperature, layer - by - layer (2D) growt h of NaCl is observed for both low and high rate deposition indicated by strong intensity oscillations. From the rate - dependence, it is clear that highly quality epitaxial growth ca n be maintained even for the highest rates measured here (0.6 BL/s), which allows the growth of 1 micron thick film in 50 min. Figure 6 . 4 . NaCl t emperature dependent homoepitaxial growth mode and 2D line scan. AFM images for (a) bare single crystal NaCl and as a function of growth temperature for homoepitaxial growth of 200Å NaCl at ( b) - (e - g ) Corresponding line scans sho wing step height. (h) Crystal structures for one NaCl unit cell (BL, 5.6 Å ) and half unit cell (ML, 2.8 Å ). 76 In general, RHEED oscillations stem from interference of the periodic formation and coalescence of 2D nuclei during the deposit ion of each layer. 151 This mechan ism is briefly described: when the deposition is initiated there is a formation of 2D nuclei that creates a disruption in the crystal potential and causes part of the diffraction be am to become out - of - phase. The coalescence of multiple islands into a smoot h surface returns the layer to a uniform potential and restores coherence and specular intensity. 239 In the case of NaC l with FCC symmetry, the full period of the RHEED oscillation corresponds to a full unit cell (or BL). This means that the bottom of the oscillation corresponds to a completed half unit cell (or ML) and each ML forms a complete layer prior to the next. Thi s picture is consistent with the top - layer interference model described above where a completed half unit cell has greater destructive interference analogous to the phase mechanism that forbids the (001) and allows (002) diffraction peaks. This also explai ns the observation of half unit cell (or ML) step heights in AFM images. In the case of perfect layer - by - layer (2D) growth on singular surfaces, each layer is completed before nucle i are formed on the next layer the resulting oscillations would then pers ist ideally with the same amplitude (undamped) indefinitely. However, in many cases adsorbates begin to nucleate on the previous layer before it is fully completed, resulting in eit her an accumulating degree of roughness or a lack of oscillation coherence that dampens the overall intensity of the specular spot. For the latter, this specular decay is also typically seen for layer - by - layer growth on vicinal surfaces such as with GaAs h omoepitaxy where growth on a singular surface (miscut < 1 mrad) results in a nearly constant oscillation amplitude while vicinal surfaces (miscut ~ 7 mrad) results in an intensity decay 218 . Given the vicinal nature of the cleaved NaCl (001) shown in Fig ure 6.4 a , this is the most 77 When the growth temperature is increas a smaller amplitude) for over 25 BLs and there is an overall recovery of the specular intensity (instead of a decay).This specular recovery is a hallmark of step - flow driven growth. 186 That is, as the temperature and the surface diffusivities of adsorbates increase, eventually only growth from the step edge is favorable and intra - terrace nucleation is su ppressed. When the diffusion length of adsorbates becomes large enough to allow adsorbates move freely between the boundaries of the terraces, a roughly constant step density is maintained that flow s uniformly across the substrate and results in a RHEED in tensity that recovers to a nearly constant value. This growth mode and specular recovery is consistent with the morphology in d where only terraces are observed (no ML islands). The contin ued observation of weak oscillations suggests that we are close to the transition between layer - by - layer (2D) and step - likely that we are observing a mixed - mode growth where growth is initiated at step edges but grows rad ially across the terrace from these nucleation points. Thi s can be clearly seen in the AFM as apparent triple junctions at the end of steps. Thus, it would be expected to see a superposition of both weak oscillations and an overall intensity recovery. For the low temperature growth at - of the base NaCl substrate disappear within a coverage of 200Å, which suggests the formation of an increasing number of grain boundaries and dislocation density. In this case, the fast drop of o scillation intensity within the growth of less than 5 BLs ind icates a transition from layer - by - layer (2D) to island (3D) growth. The AFM image in b shows non - uniform coverage at the terraces, larger terrace step height, and a rough er surface morp hology. Interestingly, it was predicted that for alkali halid e epitaxy, diffusivities would be large enough to see RHEED oscillations at growth temperatures above 0.1 x T M (absolute melting temperature), 232 which translates to - 78 surprisingly consistent with our observations (weak but still pres ent oscillations at - would then place the diffusivity in the 10 - 7 - 10 - 8 cm 2 /s range at low temperature ( - - and 10 - 4 to 10 - 5 cm 2 /s at room temperature. 232 6 . 4 NaCl Based OLEDs To further demonstrate the potential of NaCl in electronic devices, we show the incorporation of vapor deposited NaCl as the host in a green phosphorescent OLEDs structure . OLEDs utilizing NaCl as t he host were fabricated on degreased substrates coated wit h indium tin oxide(ITO) using the following structure: ITO/ MoO 3 (150 Å) / NaCl:Ir(ppy) 3 (6 vol.%, 200 Å)/ LiQ(8 Å)/ Al (800 Å) where Ir(ppy) 3 is fac tris(2 - phenylpyridine)iridium and LiQ is 8 - hydro xyquinoline lithium. The control device used the same stru cture without any emitter in the - NPD (400 Å) / CBP: Ir(ppy) 3 (6 vol.%, 200 Å)/ BCP (75 Å)/ Alq 3 - - bis[N - (1 - naphthyl) - N - phenyl - amino]biphenyl, CB - - dicarbazole - biphenyl and Alq 3 is tris - (8 - hydroxyquinolinato) aluminum. shows a comparison of current density ( J ) and luminescence current density ( L ) vs. voltage ( V ) of th e NaCl:Ir(ppy) 3 device and the control device Figure 6 . 5 . NaCl based OLEDs. (a) Current density ( J ) and l uminescence current density ( L ) vs. v oltage ( V ) of the NaCl , NaCl:Ir (ppy) 3 and CBP:Ir(ppy) 3 devices . ( b) Electroluminescence (EL) spectrum of the conventional CBP:Ir(ppy) 3 and NaCl:Ir(ppy) 3 devices . (Inset ) Photo graph of the NaCl:Ir(ppy) 3 device showing green electroluminescence . 79 with only Na Cl as the emissive layer. The absence of electroluminescence from the control devices confirms that the emission in the NaCl:Ir(ppy) 3 device originates from Ir(ppy) 3 . shows the EL spectr um of the device, which is consistent with the EL for the conventional host architecture. 240 While the peak external quantum efficiency (EQE) for the NaCl:Ir(ppy) 3 device is low at 0.005%, the performance could be improved substantially with additional thickness optimization, layer optimization (e.g. exciton blocki ng layers), or host codoping. Co nsidering NaCl can be doped to become conductive (at least at high temperatures) 241 or epit axially coupled with highly conductive halide perovskites, these demonstrations suggest new potential pathways to integrate epitaxial NaCl layer into optoelectronics. 6 . 5 Conclusion In this chapter, in situ and real time RH EED monitoring of homoepitaxial growth of NaCl on NaCl (001) is demonstrated. RHEED oscillations and AFM morphologies are observed and discussed to show the impact of growth rate and temperature on changes in growth mode which were previously unrecognized. This study demonstrates the cap ability of performing RHEED monitoring of epitaxial growth on insulating halide crystals in thermal evaporators. This understanding could be used to enable the mixed growth of alkali halides with tunable lattice constants an d enable epitaxial halide perovs kite growth for photovoltaic and optoelectronic applications discussed in Chapter 7 and 8 . We have also demonstrated NaCl as a potentially interesting component for OLEDs, further suggesting new opportunities for metal - halid e and metal - halide - perovskite ba sed single - crystal epitaxial optoelectronics. 80 Chapter 7 Single - Domain Epitaxy of Cesium Tin Bromide (CsSnBr 3 ) In this chapter, we focus on the crystalline growth of CsSnBr 3, which has been shown to be a promising and air - stabl e candidate in optoelectronics with a bandgap of 1.8 eV. We demonstrate single - domain epitaxy enabled by room - temperature reactive vapor phase deposition onto single crystal metal halide substrates with congrue nt ionic interactions and find this approach t o be generally applicable across the halide (X) series. The lattice constant of cubic CsSnBr 3 (5.80 Å ) is most closely lattice matched from all the MX alkali halide salts with NaCl (cubic lattice constant of 5. 64 Å). While the compressive misfit between Cs SnBr 3 and NaCl is - 2.8%, this provides one of the smallest misfits readily available. For th is archetypical halide perovskite, we uncover two epitaxial phases, a cubic phase and tetragonal phase, which emerge v ia stoichiometry control that are both stabili zed with vastly differing lattice constants and accommodated via epitaxial rotation. We exploit this epitaxial growth to demonstrate multilayer 2D quantum wells of a halide - perovskite system. This work ultimate ly unlocks new routes to push halide perovskit es to their full potential. 7 . 1 Experimental Reactive vapor deposition of halide perovskites was performed in a multisource custom thermal evaporator. The two precursors, CsBr and SnBr 2, were co - evaporated from separate tungsten boats to form the perovs kite layer. Prior to growth, pre - polished NaCl (100) single crystal substrates were prepared through cleaving in a glovebox or used as polished. RHEED oscillations were monitored with subst rates fixed at various in - plane orientations (KSA400). Rotation dep endent RHEED patterns were collected after each deposition was halted via source and substrate shutters. Quantum well multilayers were fabricated under similar growth conditions, where epit axial NaCl was vapor deposited from a NaCl powder source with a rat e of 0.02 Å/s and 81 a thickness of 1.5 nm. Cross - section transmission electron microscope (TEM) samples were prepared by focused ion beam (FIB), attached to a FEI Nova 200 Nanolab SEM/FIB, an d then investigated by a JEOL 3100R05 Double Cs Corrected TEM/STEM. A carbon top - layer was deposited on the cutting area to protect the epitaxial film. Scanning electron microscopy (SEM, Carl Zeiss Auriga Dual Column FIB SEM) was performed for ex situ film thickness calibration and morphology characterization. Photolumine scence spectra were measured using a PTI Quanta Master 40 spectroflurometer under nitrogen atmosphere and various excitation wavelengths. Dielectric long - pass filters were used during the P L measurement to prevent both wavelength doubling and light bleedin g. UV - VIS transmission spectra were taken using Perkin Elmer UV - VIS Spectrometer (Lambda 900). X - ray diffraction was characterized by using a Bruker D2 Phaser urce at 30 kV and 10 mA and a Ni filter in the Bragg - Brentano confi guration. X - ray photoelectron spectroscopy was performed in a separate chamber with a - ray source. Before collecting XPS data, the films were etched by Argon ions for 1.5 min to prevent t he interference of surface contamination. Epitaxial lift off device fabrication was performed by immersing the epitaxial film grown on the substrate into liquid nitrogen for about 30 s to fully cool the sample . The film and substrate were then quickly im mersed into diethyl ether. After warming to room temperature, the film and substrate were removed from the solvent and copper tape was pressed onto the halide perovskite film (with a pre - de posited gold layer of ~300 Å on top). The tape was slowly peeled to separate the halide perovskite epitaxial film from the substrate. C 60 (M. E. R. Corporation 99.9%), bathocuproine (BCP, Lumtec >99%), silver (Kurt Lesker, 99.99%), and tris - (8 - hydroxyquino linato) aluminum (Alq 3 , Lumtec, >99.5%) were then deposited onto th e surface of 82 halide perovskite film in sequence followed by Ag(4nm)/Alq 3 (60nm) as the transparent cathode for top illumination. P hotoconductive - atomic force microscope (AFM) measurements w ere carried out on an MFP - 3D - AFM from Asylum Research in a nitrogen filled cell. The illumination condition was established by a light fiber shining from the top of the sample. Pt/Ir coated tip with the spring constant of 0.2 N/m was used in the pc - AFM mea surements, while 20 nN force was applied between tip and sample. As the tip moved across the surface, the topology was measured. In the point I - V measurements, the AFM tip was fixed at different locations. While a bias connected to the bottom electrode was varied, the current between the AFM tip and bottom electrode was r ecorded. Electronic band structures and densities of states (DOS) of CsSnBr 3 and CsSn 2 Br 5 were calculated using density functional theory (DFT) implemented in the Vienna Ab initio Simulation Package (VASP). The exchange - correlation functional utilized were the Perdew - Burke - Ernzerhof (PBE) functional 242 , wh ich belongs to the generalized gradient approximations (GGA) class, and the screened Heyd - Scuseria - Ernzerhof (HSE06) 243 hybrid functional. Additional computational details can be found in the Supporting Information. Crystal structures were drawn using VESTA and selected area electron diffraction (SAED) patterns were calcul ated with CrystalMaker. The fitting method of the Bohr radius ( ) is explained below. The emission energy of a quantum well is described by the Brus equation as 244 - 245 , ( 7 . 1 ) where E g 0 is the bulk band gap, E is the confinement energy of both electrons and holes, is reduced Planck constant, L z is the thickness of the quantum well, and m * is the reduce d mass that can be obtained from the effective masses of the electron ( m e * ) and the hole ( m h * ) as 83 , is the charge of electron, is the relative permittiviy, and is the vacuum permittivity. Since the ex citon bindin g energy of CsSnBr 3 has been reported to be less than 1 meV 246 the size dependence of the quantum well bandgap ca n be express ed as 247 , ( 7 . 2 ) The value of obtained by fitting the PL data of quantum well PL data can be extracted and then used for calculation of Bohr radius of this material by using 248 , ( 7 . 3 ) where is the Bohr radiu s, is the electron charge, is the vacuum permittivity, and is the dielectric constant of the semiconductor which has been reported for CsSnBr 3 to be 32.4. 246 84 7 . 2 Phase Study via Stoichiometry Control RHEED patterns captured during the epitaxial growth of the perovskite at room temperature are shown in Figure 7 . 1 . The first row of Figure 7 . 1 shows the initial RHEED p atterns of the NaCl(100) crystal with the electron beam directed along the NaCl[110]. Th e impact of the precursor ratio on the crystal structure of epitaxial film is investigated with CsBr:SnBr 2 molar ratios ranging from 0.25:1 to 1.5:1. We confirm that th e epitaxially deposited halide perovskite films form single - domain epitaxial layers from the rotation dependent RHEED ( Figure 7 .2 a - d ) . Figu re 7 . 1 . In situ RHEED patterns of the epitaxial growth of CsSnBr 3 on NaCl . The epitaxial halide perovskite film is grown on single crystalline NaCl(100) substrates with various ratios of precursors, CsBr to SnBr 2 , where two distinct phases are observed depending on stoichiometry: cubic (1:1) and tetragonal (0.25:1 and 0.5:1). The epitaxial growth of stoichiometric CsSnBr 3 (cubic) is highlighted in the box. The uncertainty of film thickness is 1 - 1.5 MLs. 85 When depositing CsSnBr 3 using a molar ratio of 1:1 (CsBr:SnBr 2 ), the RHEED patterns remain streaky, indicating the form ation of a smooth crystalline layer. After the deposition of the first monolayer, the underlying substrate Kikuchi lines disappear as expected due to the shift in elemental composition and lattice type (face - centered to primitive). The geometry and spacing of these reciprocal lattice points obtained along the [110] and [100] direction indicate that the crystal structure of the perovskite is cubic with a calculated lattice constant of 5. 8±0.1 Å. The lattice constant is further confirmed with ex situ X - ray di ffraction (XRD) to obtain an out - of - plane lattice constant of 5.80±0.01 Å (discussed below). Several other bulk phases have been reported for CsSnBr 3 , including tetragonal and monoclin ic phases; however, only the cubic phase is stable at room temperature 249 - 250 . The high symmetry shown in all of the diffraction data clearly indicate the presence of the cubic phase ( Figur e 7 .2a - d ). Thus, we find the films are not pseudomorphic past the first monolayer, which is not surprising considering the level of compressive misfit. This implies there is a critical thickness and that there will be a considerable dislocation density to F igure 7 . 2 . Rotation dependent RHEED patterns of CsSnBr 3 and CsSn 2 Br 5 epitaxial film on NaCl . Streaky RHEED patterns of ( a - d ) cubic phase CsSnBr 3 taken from different rotation angles showing that single - doma in epitaxial layer is formed on the substrates. RHEED patterns of roughened ( e - h ) tetragonal phase CsSn 2 Br 5 taken from different rotation angles. 86 accommo date this misfit. However, because it is compressive misfit, this is less likely to lead to film cracking than if it were tensile misfit, 251 and no cracking was observed fo r thicknesses > 2000 Å. During the 1:1 growth, we also observe d clear RHEED oscillations that vary with deposition rate. Such oscillations are a hallmark of layer - by - layer growth ( Figure 7 .3 ), where the oscillation period typically corresponds to the grow th of a monolayer or bilayer, 252 but can also show complex bimodal periods. 253 Here, we find the oscillation period corresp onds to half a monolayer (two periods per monolayer), which suggests a more complex underlying reactive growth mechanism or an associated reconstruction during the reaction. This is also similar to RHEED oscillation Figure 7 . 3 . RHEED oscillations of CsSnBr 3 growth on NaCl. Specular RHEED intensity recorded during CsSnBr 3 epitaxial growth at 1:1 s toichiometry on NaCl at (a) 0.28 Å/s and (b) 0.14 Å/s. The oscillation period is 5 s and 10 s for (A) and (B), respectively, and corresponds to thickness of a half monolayer. (C) C ross - section SEM image used for calibration of the growth rate. 87 beating seen in ZnSe migration - enhanced epitaxial growth on GaAs where the oscillation period corresponded to a half monolayer. 254 Cross - section TEM images of the epitaxial CsSnBr 3 film are shown in Figure 7 .4 . Due to the inherent misfit, we do o bserve dislocations in the first several monolayers at the interface ( Figure 7 .4b ) wh en no pseudomorphic layer is present. Nonetheless, the atomic arrangement of the two materials are nearly indistinguishable, which is consistent with the observation of th e RHEED patterns. The cross - section SEM images shown in Figure 7 .4c further confirms the smooth surface of films prepared with 1:1 ratio of CsBr and SnBr 2 , indicating its suitability for the fabrication of thin - film optoelectronic devices. Figure 7 . 4 . Ordering of the halide perovskite at the halide salt interface. (a ) Enlarged cross - section TEM image (viewed along the [100] direction of NaCl) of sample prepared at 1:1 CsBr:SnBr 2 ratio with CsSnBr 3 film thickness of ~25 nm, where the black arrow shows the boundary between epitaxy and NaCl ; (b ) Enlarged image of the area marked by white frame; (c ) cross - section SEM image showin g the smooth surface o f epitaxial film. ( d ) and ( e ) show the film color prepared at different ratios 0.25:1 and 1:1, respectively. Schematics of the epitaxial structures : ( f ) Top and ( g ) side view of cubic CsSnBr 3 on NaCl. ( h ) Top and ( I ) side view of tetr agonal CsSn 2 Br 5 on NaC l. Green spheres are Cs; gray spheres are Sn; red spheres are Br; yellow spheres are Na; and light green spheres are Cl. 88 In contrast to the growth with a 1:1 stoichiometry, growth with a 0.5:1 stoichiometry results in a p hase transition from the cubic CsSnBr 3 to a stable tetragonal phase that takes place at the earliest stages of growth within the first two monolayers. Rotation dependent RHEED patterns for the tetragonal phase are shown in Figure 7 .2e - h . While tetragonal d istortions are common for large lattice misfits in pseudomorphic growth, this tetragonal phase is not a simple distortion, nor is it the low temperature tetragonal phase ; 249 - 250 it even appears to be a monoclinic phase when monitored along the NaCl [110] direction. This indicates that the growth with moderate Cs deficiency leads to a susc eptibility to transitioning to Cs Sn 2 Br 5 . This phase transition process is further elucidated by the RHEED data in Figure 7 .5 where only the first ML is cubic and then the subsequent layers transform to the tetragonal phase. The transition was monitored und er rough Figure 7 . 5 . In situ real - time RHEED monitoring of the phas e transition. A phase transition from the cubic to tetragonal phase occurs when the deposition ratio of CsBr to SnBr 2 is 0.5:1 after 1 - 2 monolayers. Note that while the pattern for the tetragonal phase appears monoclinic, it is actually a rotated tetrago nal phase and the diffraction spots are therefore not along primary axes. 89 growth conditions with s ubstrates fixed at NaCl [110] due to the symmetry and lattice matching of the a - b plane of the tetragonal phase. The RHEED intensity monitoring also allows the study of phase transitions from cubic to tetragonal because RHE ED pattern changes can be monitor ed as a change of particular diffraction peak locations ( Figure 7 .6 ) . At later stages of growth (>7 MLs) the spotty patterns of the tetragonal phase become streak ed , which is indicative of the crystalline film changing from rough to smooth while maintaini ng the initial tetragonal crystal structure. To further confirm the phases shown in the RHEED patterns, XRD was used to determine the out - of - plane lattice parameter for the epitaxial films. As the ratio of CsBr to SnBr 2 increases from 0.25:1 to 1:1, the p eaks at 11.57 o ( d = 7.64±0.01 Å) and 23.46 o are replaced by peaks at 15.31 o ( d = 5.80±0.01 Å) and 30.83 o as shown in Figure 7 .7 . The observed peaks are consistent with the d - spacings along the c - axis calculated from RHEED patterns and correspond to the (00 1)/(002) and (002)/(004) peaks of the cubic CsSnBr 3 and tetragonal CsSn 2 Br 5 phases respectively. Based on the RHEED and XRD dat a, the measured lattice constants and orientations of the two epitaxial phases are summarized in Table 7 . 1 . Surprisingly, we find that both the cubic CsSnBr 3 and tetragonal CsSn 2 Br 5 can grow epitaxially, even though the lattice constant of CsSn 2 Br 5 is much Figure 7 . 6 . RHEED oscillations monitored during the phase transition. (a) RHEED pattern with the monitored intensity area hig hlighted with the red circle and the corresponding (b) RHEED intensity profile with time. 90 larger and the mismatch between CsSn 2 Br 5 and NaCl is 4.9 %. This larger lattice is accommodated via the rotation of CsSn 2 Br 5 rel ative to the metal halide substrate. Schematics of the epitaxial growth of CsSn 2 Br 5 and CsSnBr 3 on NaCl substrates are shown in Figure 7 .4 f - i . Figure 7 . 7 . Crystal structure characterization of the two epitaxial phases, CsSnBr 3 and CsSn 2 Br 5 . XRD patterns of NaCl (blue) and samples grown at different ratios of CsBr:SnBr 2 : 0.25:1 (black) and 1:1 (red). Table 7 . 1 . Lattice constants, film orientation and misfit of two phases , CsSnBr 3 and CsSn 2 Br 5 . Bulk Crystal Structure CsSnBr 3 CsSn 2 Br 5 Pm3m, a = b = c =5.80 Å I4/mcm, a = b =8.48 Å, c =15.28 Å Precursor Ratio of CsBr to SnBr 2 0.25:1 - Observed 0.5:1 Observed < 2 ML Observed > 3 ML 1:1 Observed - 1.5:1 - - Orientation Along NaCl [110] [110] [210] Along NaCl [100] [100 ] [3 0] Along NaCl [001] [001] [002] Misfit - 2.8 % 4.9 % 91 Epitaxial growths with both a greater CsBr deficiency (0.25:1) and CsBr excess (1.5:1) were also investigated as shown in Figure 7 . 1 . At the CsBr deficient (0.25:1) ratio, the pure tetragonal phase is observed without first seeing the cubic structure. In c ontrast, the CsBr excess (1.5:1) ratio resulted in ring - like patterns, which indicates the film becomes a thr ee - dimensional polycrystalline powder. Epitaxial films were also characterized by X - ray photoelectron spectroscopy (XPS) to measure the elemental ratios in the films deposited using various ratios. By fitting the XPS peak, the elemental ratio of Cs to Sn can be extracted, and the results are su mmarized in Table 7 . 2 . The epitaxial film deposited with 1:1 ratio of CsBr to SnBr 2 is indeed stoichiometric CsSnBr 3 (1:1:3 for Cs:Sn:Br). The other two ratios of 0.25:1 and 0.5:1 both lead to films deficient in Cs. The combination of RHEED, XRD and XPS an alysis indicates that the growth of CsSnBr 3 is more Table 7 . 2 . Elemental ratio of as - deposited films with different of CsBr :SnBr 2 obtained from X PS data. of CsBr to SnBr 2 p> p> p> p> p> p> p> p> p> 92 favorable when Cs is stoichiometric or in slight excess, while CsSn 2 Br 5 dominates when there is a Cs deficiency. Both selective elemental vacancies and lattice mi sfit can ultimately play a role in initiating strain - driven phase transitions in these systems. While the elemental vacancies can be controlled by stoichiometry, the lattice misfit can be tuned through compositional alloying of the metal halide substrate, either in the bulk or as thin pseudomorphic interlayers. NaBr has larger lattice const ant Figure 7 . 8 . Epitaxial growth of CsSnBr 3 on pseudomorphic interlayers of NaCl - NaBr. ( a - e) RHEED patterns of growth of pseudomorphic interlayers of NaCl - NaBr on NaCl substrates. The ratio of NaBr: NaCl is gradually increased to reduce the misfit between substrates and interlayers. Note that whe n the ratio of NaBr: NaCl reaches 1:1, the lattice constant of interlayers is 5.81 Å perfectly matching that of CsSnBr 3 (5.80 Å). During the growth, the lattice constant calculated from RHEED patterns shows gradual increase along with increasing the ratio of NaBr: NaCl. ( f - i) RHEED patterns of growth of CsSnBr 3 on pseudomorphic interlayers of NaCl - NaBr. 93 (5.98 Å) than cubic CsSnBr 3 . Therefore, alloying NaBr and NaCl can provide near perfect lattice matching for epitaxial growth of cubic CsSnBr 3 ( Figure 7 .8 and 7 .9 ) , r ep resenting a general strategy for tuning the lattice misfit and dislocation density. 7 . 3 Epitaxial Lift - off The epitaxially deposited halide perovskite films are strongly bonded to the metal halide crystals (see Figure 7 .10 a - c ). The strong bond formed a t the interface of the halide perovskite and metal halide substrates is very likely due to the similar ionic characteristic between halide perovskite and metal halide salt. This is confirmed by the attempted growth of halide perovsk ite Figure 7 . 9 . XRD patterns of alloyed NaCl - NaBr pseudomorphic interlayers and CsSnBr 3 . ( a) Comparison of NaCl substrate alone and alloyed NaCl - NaBr layers. (b) Enlarged view of (a) shows a shoulder peak arises in the sample of alloyed NaCl - NaBr pseudomorphic interlayers on NaCl substrates. ( c) Enlarged view shows the peak splitting that refle cts the change o f lattice constant from NaCl substrate, alloyed NaCl - NaBr, and CsSnBr 3 . 94 on Ge and InP substr ates, which have similar degree of lattice misfit but are more covalent in their bonding character. According to the RHEED pattern shown in Figure 7 .11 , neither Ge nor InP provides readily suitable surfaces for epitaxial growth of halide perovskite, whi ch leads to the formation of polycrystalline films. With the successful growth of single - domain epitaxial films we further demonstrate an epitaxial lift - off (ELO) process to allow the separation of these epitaxial layers from single crystalline substrate s and enable substrate regrowth ( Figure 7 .10 d - g ). ELO has been shown to be an important processing method for making, for example, single - domain GaAs more economically viable for solar cells. 255 - 258 In our work this is achieved by flash cooling/heatin g with liquid Figure 7 . 10 . Film adhesion of the epitaxial growth of CsSnBr 3 on NaCl and epitaxy lift - off/regrowth. ( a) Epitaxial CsSnBr 3 sample b efore application of Scotch tape; (b) after attachment on the film surface; and (c) after peeling off. (d) Epitaxial lift - off procedure: the epitaxial CsSnBr 3 grown on NaCl with a gold layer (~300 Å) on the top is rapidly immersed into liquid nitr ogen and then rapidly transferred into diethyl ether; Cu tape is then pressed onto the surface and then slowly peeled which results in separation of the epitaxial film from the substrate; (e - f) RHEED patterns for the epitaxial re - growth of CsSnBr 3 after ep itaxial l ift - off. 95 nitrogen immersion followed by rapid immersion in diethyl ether to initiate cracking at the interface of the substrate and epitaxial film by differences in the thermal expansion coefficients. Figure 7.1 . RHEED patterns of growth of CsSnBr 3 on Ge and InP. (a) Ge single crystalline substrate along the [100]; (b) CsSnBr 3 grown on Ge; (c) Ge single crystalline substrate along the [100]pre-treated with HCl acid etching for 10 min; (d) CsSnBr 3 grown on Ge from (c); (e) InP single crystalline substrate along the [100]; (f) CsSnBr 3 grown on InP .Figure 7.1 . Absorption spectra of CsSnBr 3, CsSn 2Br5 and NaCl substrate. (a) CsSnBr3 of varying well thickness and (b) CsSn 2Br5 (black curve) and NaCl (blue curve). The spectra are converted from (1-Transmission) and shifted for clarity. 96 Subsequently , epitaxial films are peeled with conductive copper tape. Moreover, the substrates can be reused for further epitaxial growth ( Figure 7 .10 d - g ). 7 . 4 Discussion Experimental and theoretical studies are performed on the CsSn 2 Br 5 and CsSnBr 3 phases to underst and the properties for each. Absorption spectra of as - prepared epitaxial films are shown in Figure 7 .1 2 a and confirm that the band - gap of epitaxial CsSnBr 3 is 1.83±0.02 eV, which is Figure 7 . 13 . Electronic band structures of CsSnBr 3 and CsSn 2 Br 5 . PBE b and structure, density of states (DOS) and projected density of states (PDOS) of ( a ) CsSnBr 3 and ( b ) CsSn 2 Br 5 . HSE06 band structure, density of states (DOS) and projected density of states (PDOS) of ( c ) CsSnBr 3 and ( d ) CsSn 2 Br 5 . 97 consistent with both theoret ical 246 and experimental 249 results reported previously. For CsSn 2 Br 5 , we measure a bandgap of 3.34±0.04 eV, wh ich is clearly distinguishable from the NaCl bandgap of ~9 eV ( Figure 7 .1 3 b ). The calculated band structures, density of states (DOS) and projected Tab le 7 . 3 . Calculated band gaps of CsSnBr 3 and CsSn 2 Br 5 using the DFT - PBE and DFT - HSE06 methods. Figure 7 . 14 . Solar cell device from epitaxial lift - off of CsSnBr 3 . AFM image of ( a ) single - domain epitaxial lift - off film, and ( b ) amorphous film. ( c ) Device arch itecture of the photovoltaic cells. ( d ) I - V curve of devices fabricated with the single domain epitaxy film and amorphous film, respectively, showing nearly an order of magnitude more photocurrent and double the voltage for single domain films versus the a morphous film. 98 density of states (PDOS) of CsSnBr 3 and CsSn 2 Br 5 using the DFT with the HSE06 functional Figure 7 .1 3 c - d . A s ummary of the calculated band gap values can be found in Table 7 .3 . The res ulting HSE06 band gaps for CsSnBr 3 and CsSn 2 Br 5 are 0.84 eV and 3.12eV, respectively. Note that DFT methods underestimate the gap of most semiconductors. These values are in reasona ble agreement with the observed properties of CsSnBr 3 and CsSn 2 Br 5 . Had the perovskite been pseudomorphic on pure NaCl, we predict that the bandgap would decrease by around 0.5 eV, which is clearly not observed experimentally. For comparison, the PBE band structures, DOS and PDOS of CsSnBr 3 and CsSn 2 Br 5 are shown in Figure 7 .1 3 a - b . These calculations further confirm that the tetragonal phase is CsSn 2 Br 5 with a large bandgap. Given the suitable bandgap of CsSnBr 3 for solar cell applications we fabricate d p hotovoltaic devices to compare the optoelectr onic properties of single - domain epitaxial films and amorphous films of CsSnBr 3 . AFM images show that the surface of epitaxial films is much smoother than that of amorphous films ( Figure 7 .1 4 a - b ). Current - voltag e ( I - V) curves measured by conducting probe AFM show that the devices fabricated with the single - domain epitaxial film as the absorber layers have both higher I sc and V oc than the control devices fabricated with amorphous films by a factor of 10 and 2 resp ectivel y ( Figure 7 .1 4 c - d ). This clearly indicates that the high defect concentration present in amorphous and polycrystalline CsSnBr 3 is 99 Figure 7 . 15 . Epitaxial CsSnBr 3 based 2D Quantum well fabrication . RHEED patterns of ( a ) NaCl along the [110] direction, ( b ) NaCl/CsSnBr 3 (~40 nm), and ( c ) NaCl/Cs SnBr 3 (~40 nm)/NaCl (1.5 nm). ( d ) Schematic illustration of NaCl/CsSnBr 3 quantum well structure , green spheres are Cs; gray spheres are Sn; red spheres are Br; ye llow spheres are Na; and light green spheres are Cl; ( e ) PL spectra of quantum well samples wit h various well width: 5 nm (black curve), 10 nm (red curve), 20 nm (blue curve), 40 nm (magenta curve), 80 nm (orange curve), and 100 nm (violet curve). (f ) Emis sion energy of quantum wells with varying well width . The fitting is described in the Supportin g Information . The inset shows the photograph of samples illuminated under UV light. Samples from left to right are bare single crystal, quantum well of NaCl/CsS nBr 3 ( 40 nm), and quantum well of NaCl/CsSnBr 3 (~100 nm). (g) PL spectra of quantum well samples C sSnBr 3 /CsSn 2 Br 5 with various well widths: 20 nm (black curve), 40 nm (red curve), 80 nm (blue curve), 100 nm ( green curve). PL of quantum well samples CsSnBr 3 /Na Cl as a comparison: 80 nm with NaCl ( orange curve), and 100 nm with NaCl ( olive curve). 100 indeed a key limitation for enhancing device performance 96 that can be over come with epitaxial layers. Based on the control afforded b y this epitaxial halide growth, we further fabricate quantum wells with varying well thicknesses for CsSnBr 3 paired with both vapor - deposited NaCl and CsSn 2 Br 5 as the well barrier. Quantum wells are important in a range of optoelectronic devices and provid e critical insight into the physical properties of quantum confined charge carriers, two - dimensional electron gas, 259 and tunable luminescence. The growth process was investigated by RHEED to confirm the formation of epitaxial multilayers as shown in Figure 7 .1 5 a - c where NaCl was grown under similar conditions to homoepitaxial growth demonstrated previousl y. 125 The data in shows that no obvious change occurs after depositing the epi taxial barrier layer on the halide perovskite or after depositing multiple quantum well layers. That is, the NaCl epitaxial layers are pseudomorphic with the p erovskite film. The PL spectrum of CsSnBr 3 /NaCl quantum wells were studied by adjusting the well thickness shown schematically in Figure 7 .1 5 d . When the well thickness is reduced from 100 nm to 5 nm, the emission peak redshifts ( Figure 7 .1 5 e ) to a similar degree as seen with colloidal nanocrystals. 260 From fitting the size de pendence of the bandgap, we can estimate an effective reduced mass of m * = 0.30 m e , where m e is the rest mass of the electron, and the Bohr radius of CsSnBr 3 of ~5.6 nm). This is similar in magnitude to CdSe (5.6 nm) 261 and Si (~5nm) 262 , smaller than PbS (20 nm) 263 and larger than ZnS (2.5 nm). 264 In moving from the weak to strong confinement regime, we expect that the bandgap of CsSnBr 3 to reach up to 3.0 eV with the smallest well thickness around 1 nm. We note that quantum wells with CsSn 2 Br 5 barrier layers were also feasible through switching of the stoichiometry. As shown in Figure 7 .1 5 g , the PL spectra of CsSnBr 3 /CsSn 2 Br 5 quantum wells show consistent changes with those of CsSnBr 3 /NaCl quantum wells when varying well thickness. A summary of the emission energy of 101 CsSnBr 3 /NaCl and CsSnBr 3 /CsSn 2 Br 5 quantum well with various well width are shown in Table 7 .4 and 7 .5 . This provides a route to better index - match the two layers in multiple quantum wells with enhanced potential for el ectrical injection. 7 . 5 Conclusion In summary, we demonstrate a route to the room - temperature epitaxial growth of inorganic halide perovskites using low cost metal halide crystals and show the emergence of two epitaxial phases of cesium tin bromide (CsSn Br 3 and CsSn 2 Br 5 ) with vastly differing lattice constants and bandgaps based on stoichiometry control. The larger lattice of CsSn 2 Br 5 is accommodated via the rotation of crystal planes relative to the metal halide substrates. The phase transitions between the Table 7 . 4 . Emission energy of quantum well CsSnBr 3 /NaCl with various well width. Table 7 . 5 . Emission energy of quantum well CsSnBr 3 /CsSn 2 Br 5 with various well width. 102 cubic CsSnBr 3 and tetragonal CsSn 2 Br 5 phases were manipulated and observed in real - time. The lattice misfit between the ionic epi taxial film and the substrate is precisely tuned by applying a pseudomorphic buffer layer of alloyed alkali metal halide sa lts and an epitaxial lift - off method has been demonstrated for further device fabrication. The dominant performance of devices fabric ated with the epitaxial film confirms that the high crystallinity and low defect intensity are beneficial for halide perovs kite optoelectronic applications. We further exploit the epitaxial growth of CsSnBr 3 to demonstrate multilayer epitaxial quantum well s of a halide perovskite and extract the Bohr radius for CsSnBr 3 of 5.6 nm, which provides a guide for manipulating quantum confinement in this class of materials. These demonstrations are likely to spark the exploration of a full range of epitaxial halide perovskites and help enable their ultimate potential in many emerging applications. 103 Chapter 8 Epitaxial Stabilization of Tetragonal Cesium Tin Iodide (CsSnI 3 ) ( This chapter was reproduced with permission from Wang, L., Chen, P., et al., Epitaxial Stabilization of Tetragonal Cesium Tin Iodide. ACS Appl. Mater. Interfaces, 2019. DOI: 10.1021/acsami.9b05592 . Copyright 2019 American Chemical S ociety. ) In this chapter, we demonstrate the single - domain epitaxial growth of cesium tin iodide (CsSnI 3 ) on closely lattice matched single crystal potassium chloride (KCl) substrates. One of the critical chal lenges for improving the optoelectronic performance of CsSnX 3 - based (X = Cl, Br, and I) devices is controlling the phase durin g film growth. For cesium tin iodide (CsSnI 3 ) , four different phases can form depending on the temperature including cubic (B - tetragonal (B - orthorhombic (black, B - (black) and orthor hombic (yellow) phases are stable at room temperature. 164, 265 These phases have different optical and electronic properties which i mpact carrier transport. To date, most research on CsSnI 3 focuses on the orthorhombic (black, B - . 266 However, there has been little work on the cubic and tetragonal phases since they are both unstabl e at room temperature. In the previous chapter , the epitaxial growth of the cubic and tetragonal phases of CsSnBr 3 are achieved on NaCl substrates via stoichiometry control. Using similar growth techniques, we demonstrate the room - temperature epitaxial growth of CsSnI 3 on single crystalline alkali hali de salt substrate KCl. The substrate was chosen to closely match the high - temperature cubic phase lattice constants and maintain similar ionic bonding character of the film. However, we find that the epita xial growths lead to a new stable tetragonal phase at room temperature that has a lattice provides a route to the stabilization of a metastable phase. The impact of strain is explored by 104 tuning substrate lattice con stants via alloyed interlayer deposition. Exploiting the precise growth enabled by vapor phase epitaxy, we also explore the quantum confinement effect in this material and determine a surprisingly small ef fective Bohr radius. Finally, epitaxial films of Cs SnI 3 are integrated into a lateral photodetector application with good photoresponse. This work provides new insight into the stabilization of halide perovskite crystal phases and expands the selection of epitaxial halide perovskites. 8 . 1 Experimental The epitaxial growth of CsSnI 3 was carried out by using a customized multisource thermal evaporator (Angstrom Engineering) equipped with a reflection high energy electron diffraction (RHEED) system (STAIB Ins truments). The perovskite layer was formed by co - evaporating the precursors, CsI and SnI 2, from separate tantalum boats. Prior to growth, KCl single crystal substrates were cleaved to expose the fresh (100) surfaces in a glovebox. The base pressure during epitaxial growth was controlled to be less than 3×10 - 6 torr and the deposition rates were measured in situ with separate quartz crystal microbalances for each source . The crystal structure was monitored in real - time and in situ using RHEED (30.0 keV) optim ized with an ultra - low current (< 10 nA) to elim inate damage and charging of the film over the growth times investigated. The capability of performing low - current RHEED on insulating halide crystals has been demonstrated in our previous studies. 267 RHEED oscillations were monitored with substrates fixed at various in - plane orientations ( KSA400). Rotation dependent RHEED patterns were collected after the deposition was halted via substrate shutters. Quantum well multilayers were fabricated under similar growth conditions. Epitaxial KCl (barrier layer) was vapor deposited from a KCl powder source with a rate of 0.002 nm/s and a barrier t hickness of 1.5 nm. 105 Cross - section TEM samples were prepared by focused ion beam (FIB), attached to a FEI Nova 200 Nanolab SEM/FIB, and then investigated by a JEOL 3100R05 Double Cs Corrected TEM/STEM oper ated in TEM mode. A carbon top - layer was deposited on the cutting area to protect the epitaxial film. Scanning electron microscopy (SEM, Carl Zeiss Auriga Dual Column FIB - SE M ) was performed for ex situ film thickness calibration and morphology characteriza tion. PL spectra were measured using a PTI Quanta Master 40 spectrofluorometer under nitrogen atmosphere and various excitation wavelengths. Dielectric long - pass filters wer e used during the PL measurement to prevent both wavelength doubling and light blee ding. UV - Vis transmission spectra were taken using a Perkin Elmer UV - Vis Spectrometer (Lambda 900). Nitrogen gas was purged during both PL and UV - Vis measurements to protect the sample during the scan. Out - of - plane XRD and pole figure were characterized b y using a Rigaku SmartLab XRD mA and using a Ni filter in the Bragg - Brentano configuration. The sample was placed in a nitroge n sealed dome to protect from oxygen and moisture during the scan. Note that the in tensity of the film peaks is reduced due to the presence of the dome . Accelerated synchrotron XRD data were collected at 100 K at the NE - CAT beamline 24 - ID - C, Advanced Phot on Source (APS) (wavelength: 0.0619870 nm). The sample was prepared by carefully cl eaving the single crystalline epitaxial film from the substrate and mounting/folding the film onto a copper holder stored at liquid nitrogen during transferring. The data we re processed with XDS as implemented in RAPD ( https://github.com/RAPD/RAPD ). XPS was performed in a X - ray source. Befor e collecting XPS data, the films were etched by Argon ions for 1 min to prevent the interference of surface contamination. 106 Lateral photodetectors were fabricated by depositing a thin layer (70 nm) of Au film onto the epitaxial CsSnI 3 single crystalline film by e - beam deposition. The channel between the two Au electrode was fabr icated by 3 film as the mask to form the channel region. Silver paste (Ted Pella) was used to fix a gold wire onto each Au electrode for better electri cal conne ction. The gold wires of both electrodes were connected to a customized photocurrent set - up with a Keithley 2420 source meter. Simulated AM1.5 G solar illumination ( xenon arc lamp) was used as the light source, with NREL - calibrated Si reference ce ll with K G5 filter . Figure 8 . 1 . Schematic of the epitaxial growth. (A) The epitaxial halide perovs kite film is grown on a single crystalline K Cl (100) substrate with two precursors, Cs I and Sn I 2 (1:1). In situ RHEED is used to monitor the growth in real time . (B) A schematic of the epitaxial structure of the CsSnI 3 film on KCl ( 100 ) (Cs: orange ; Sn: grey; I: purple: K: blue ; Cl: light green). 107 8 . 2 Epitaxial Phase S tabiliz ation Thin film CsSnI 3 was grown epitaxially on KCl single crystalline substrates via reactive thermal deposition of CsI and SnI 2 with a molar ratio of 1:1, which was monitored in situ and real - time using reflectio n high - energy electron diffraction (RHEED ) . A schematic of the epitaxial growth is shown in Fi gure 8 . 1b . The freshly cleaved KCl substrate exhibits sharp streak ed peaks along the [100] direction along with Kikuchi lines indicative of the single crystalline (single domain) and thick nature of the KCl substrates. After growing one monolayer of CsSnI 3 , the Figure 8 . 2 . In situ RHEED patterns of the epitaxial growth of CsSnI 3 . ( a ) RHEED pattern of the bare KCl substrate and 20nm of CsSnI 3 . ( b ) Rotation dependent RHEED patterns of the smooth CsSnI 3 epitaxial film. ( c ) RHEED oscillations monitored by capt uring the peak intensity change of the central streak of the pattern as a function of time. The observation of 0.5ML/per iod is consistent with the epitaxial growth of CsSnI 3 . ( Note the growth rate 0.02nm/s is for the reacted film not the sources.) 108 pattern changes and the KCl Kikuchi patterns disappear. A new set of streak ed lines emerge overlapping with the (20) streaks of KCl, along with (10) streaks between the (20) streaks ( Figure 8 . 2a ). These additional streaks are expected as the structure changes from a face - centered cubic (FCC) lattice of the substrate to the primitive cell of the halide perovskite. These patterns are observed without any subst rate rotation, indicating good lattice matching and alignment between the film and the substrate. The streak ed RHEED patterns remain essentially unchanged up to and beyond thicknesses of ~20 nm. The streak ed pattern is an indication of a smooth crystalline layer which is consistent with the morphology shown in the scanning electron microscop e (SEM) ( Figure 8 . 6 c - d ) . To further understand the crystal structure of the epitaxial film, rotation dependent RHEED patterns were collected ( Figure 8 . 2b ). The patterns collected from different azimuthal angles vary as expected for a single crystal, indicating single - domain growth is achieved Figur e 8 . 3 . RHEED patter n of the bare KCl substrate and surface reconstruction of CsSnI 3 . H alf - order streaks are occasionally observed at high growth rate indicating a surface reconstruction of the perovskite mo nolayer. 109 across the KCl substrate. From th e d - spacing obtained at various azimuthal angles, we obtained the lattice constant of the epitaxi al film: a = b = 0.622±0.007 nm. During the epitaxial growth, low source deposition rates (<0.02 nm/s) are found to aid in the formation of smoother films and more streak ed patterns. If the sources are deposited with a much higher growth rate, spotty patt erns quickly replace the streak ed patterns from the KCl substrate, indicating that a rough crystalline film is formed . The growth mode was also studied by moni toring the intensity of the RHEED patterns ( Figure 8 . 2c ). When using low source deposition rates, RHEED oscillations are observed indicating Frank - van der Merwe (layer - by - layer) growth mode . Two full cycles of the oscillation are found to correspond to the growth of one monolayer (half of a complete unit cell) of the CsSnI 3 fi lm as confirmed from ex s itu thickness measurements ( Figure 8 . 2c ). This is similar to the observation for epitaxial CsSnBr 3 growth on KBr discussed in Chapter 7 . 268 110 Under certain g rowth conditions (e.g. high rate growth), half - order RHEED streaks sometimes appear as seen in Figure 8 . 3 . This indicates the formation of a surface reconstruction. In these cases, the half - order streaks fade as the film growth proceeded to higher thicknes s. Because these streaks are only observed under high growth rates, it is unlikely to be indicative of the underlying growth front mechanism. That is, the growth front does not typically proceed via a surface reconstruction. Figure 8 . 4 . XRD and synchrotron XRD scans of epitaxial CsSnI 3 ( a ) Out - of - plane XRD scan of 40nm of CsSnI 3 on KCl substrate and photograph of the film on a 1cm x 1cm substrate . The KCl (200) pe ak is at 28.35 o (CuK and WL peaks are seen at 25.50 o and 27.42 o ) and the CsSnI 3 (001) and (002) peaks are at 14. 4 6 o and 29. 1 6 o . ( b ) Synchrotron X - ray diffraction (XRD) patterns of 20nm of CsSnI3 removed from the substrate and formed into a folded (text ured pseudo - powder) film. An obvious peak split is seen at 11.28°and 11.39° (inset), which corresponds to the (200) and (002) planes of the epi - tetragonal phase respectivel y. 111 Figure 8 . 5 . Pole figure scans of CsSnI 3 on KCl. ( a ) The (111) p ole figure scan of KCl substrate. ( b - d ) The (112) (021) (112) p ole figure scan of 40nm of CsSnI 3 on angle s around the sample surface normal direction. (e - f) Simulated stereograp hic projections of KCl and epi - tetragonal CsSnI 3 , which are in good agreement with the experimental scans. Note that a couple of the (211) peaks are missing due to the slight t ilt of the crystal. (The (112) (021) (112) poles for epi - tetragonal CsSnI 3 are c learly distinguishable from the KCl substrate. 112 Out - of - plane X - ray d iffraction (XRD) data was collected to confirm the crystal structure of the epitaxial film ( Figure 8 .4 ) . The peaks observed at 14.46 o and 29.16 o are the (001) and (002) of the epitaxial phase respectively, which gives a lattice constant c = 0.612 ±0.002nm . Combined with the RHEED d - phase ( a = b = 0.622±0.007 n m , c = 0.612±0.002nm) stabilized by lattice matching with the KCl substrate ( a = 0.629nm) that has a significantly different structure and atomic arrangement from the bulk high - temperature tetragonal phase. 164 The d spacings are well matched with the substrate in the early growth stage (within error) and remain nearly unchanged as the growth proceeds to higher thickness in dicating the growth is pseudomorphic . The (112), (021), and (211) pole figure Figure 8 . 6 . TEM and SEM image of CsSnI 3 on KCl. ( a ) Cross - section HR TEM image at the inter face . The film is shaded in red on the left side of the image. N ote a substrate step can be observed just to the bottom right of the shading . ( b ) SAED pattern of the film clearly indicating the tetragonal nature of the film. ( c ) SEM and ( d ) cross - section SEM i mage of 40nm of CsSnI 3 on KCl substrate. 113 scans of the film are shown in Figure 8 .5 . These were chosen as they are clearly distinguishable from any substrate peaks and not allowed for KCl. The fixed diffraction angles fo r each pole figure scan are calculat ed from the lattice constant of the epi - tetragonal phase. The (112) and (021) pole figure show the expected 4 - fold symmetry while the (211) scan show the expected 8 - fold symmetry. The a / c ratio obtained from the transmis sion electron microscopy (TEM) high - resolution images of the epitaxial interface is 1.03 , which clearly emphasizes the tetragonal nature of the structure and is close to the 1.02 value obtained from diffraction data (Figure 8 .6 ) . The TEM also clearly confirms the pseudomorphic nature of the film. To further verify the crystal structure, the film was removed from the substrate and folded to obtain a pseudo - powder (textured powder) for synchrotron X - ray characterization. This data confirms that the epitaxial film is a new tetragonal phase as sh own in Figure 8 .4b . An obvious peak split is seen at 11.28°and 11.39° that correspond to the (200) and (002) planes of the tetragonal phase respectively. From the synchrotron XRD data fitting, we obtain lattice constants of the epitaxial film: a = b = 0.6 240nm, c = 0.6178nm, which are close to and consistent with the XRD and RHEED analysis. The slight difference is due to the temperature difference of the scan conditions. Extrapolating the lat tice constant from room temperature to 100K for comparison to th e synchrotron data, the lattice constant of the cubic KCl substrate would be 0.625±0.001nm, assuming a thermal expansion coefficient 3.2x10 - 5 /K, 269 which is within error of the in - plane lattice constant measured with the synchrotron ( a = b = Table 8 . 1 . Elemental ratio of as - deposited CsSnI 3 film obtained from XPS data. 114 0.6240nm). This further confirms the pseudomorphic nature of the film and is consistent with the findings from TEM. Whe n determining the crystal structure and phases of these compounds, it is important to consider the oxidation states. For example, Sn 2+ and Sn 4+ oxidation states have been observed in this class of materials, which leads to other analogs such as Cs 2 SnI 6 (cu bic, a = 1.165 nm 270 ). This phase is ruled out from the simulated diffraction patterns ( Figure 8 . 7 ). Furthermore, we con firm the chemical state of Sn and the ratio of these three elements with X - ray photoelectron spectroscopy ( XPS) . The ratio matches with the stoichiometry of CsSnI 3 and the Sn is found to have a +2 valence state in the epitaxial film ( Table 8 .1 ). We note th at we do obverse a phase transformation from CsSnI 3 to Cs 2 SnI 6 when epitaxial films are exposed to air for more t han 20hrs ( Figure 8 . 8 ). This is consistent with other reports for the B - of CsSnI 3. 271 However, we do not observe this transfo rmation for films kept in a N 2 environment under dark conditions or under illumination (~1 - Sun) ( Figure 8 . 8 ). Figure 8 . 7 . Simulated Cs 2 SnI 6 diffraction patterns. ( a ) Simulated SAED and (b) XRD patterns of Cs 2 SnI 6 . 115 Both the tetragonal phase and the cubic phase are unstable at room temperature. 272 Epitaxy has been commonly deployed to grow meta - stable phases at lower temperature. 273 The mechanism for stabilization can include pseudomorphism (strain accommodating lattice matching) or Figure 8 . 8 . Lifetime test for the epitaxial CsSnI 3 on KC l . ( a ) XRD patterns and ( b ) color change of an epitaxial CsSnI 3 film (40nm) on KCl as a function of time in the dark and exposed to air. ( c ) XRD patterns and ( d ) color change of an epitaxial CsSnI 3 film (40nm) on KCl as a function of time under light soaking and in N 2 . ( e ) Color chan ge of an epitaxial CsSnI 3 film (40nm) on KCl as a func tion of time in dark and in N 2 . 116 superstructures, while relaxation past a critical thickness would result in a bulk structure. In our case, pseudomorphism is the dominant mechanism (seen directly in the high resolution TEM and synchrotron data) where lattice matching constraints result in a phase that looks similar to the high temperature cubic phase but exhibit tetragonal distortion to accommodate the tensile strain and perovskite octahedra. We subsequently utilized the growth on other halide crystal salts with much larger lattice misfit (> 10%) to emphasize the control over the resulting phase. When changing the substrate to KBr ( a = 0.660nm), 164 growt h of the orthorhombic phase was observed from RHEED and XRD ( Figure 8 . 9 ). The RHEED pattern of the film grown on KBr substrate shows the formation of rotated microdomains. The increase in misfit strain to a value where the critical thickness is less than a monolayer results in unstrained incommensurate ( quasiepitaxial ) film growth that then exhibits the expected room temperature orthorhombic phase. Figure 8 . 9 . Growth of CsSnI 3 on KBr. ( a ) RHEE D and ( b ) XRD patterns of Cs SnI 3 growth on KBr. 117 As we have shown previously, lattice engineering by alloying epi taxial metal halide interlayer films is an effective approach to reduce the misfit strain between the film and the substrate. 268 For CsSnI 3 growth, a thin all oyed layer (6nm) of KCl and NaCl with a ratio of 6.7:1 is grown homoepitaxially on the KCl substrate to provide a near - ideal lattice match to the halide perovskite. Streak ed RHEED patterns of the alloyed layer and the epitaxial CsSnI 3 layer are shown in Fi gure 8 . 10 . By adding a tuned pseudomorphic interlayer of alloyed alkali halide salts, the misfit defect concentration is like ly to be reduced even further. Figure 8 . 10 . Epitaxial g r owth of CsSnI 3 on pseudomorphic interlayers of K Cl - Na Cl. RHEED pattern of the bare KCl substrate, the alloyed growth of 5 nm of KC l and NaCl ( 6. 7:1) and 5 nm of CsSnI 3 . 118 8 . 3 Optical Properties and Quantum Well Fabrication The optical properties of the CsSnI 3 epi - te tragonal phase on KCl were studied as shown in Figure 8 .1 1 . The band gap energies were extracted with a Tauc plot, giving a direct band gap of 1.85 eV ( Figure 8 .1 1 a inset). This value is consistent with the photoluminescence (PL) peak emission and onset em ission at 1.47eV and 1.60eV respectively. The bandgap simulation of all the relevant phases uses the B3LYP h ybrid functional using the CASTEP module in Material Studio 7.0 with the B3LYP functional. (300 eV cutoff energy with the k - points set to 2 × 2 × 2 for the Brillouin zone integration). The epi - tetragonal bandgap is calculated to be 2.07 eV. This overestima tion from simulation is consistent with slight overestimation for the orthorhombic phase (1.40 eV from simulation and 1.31 eV from experiment 274 ). This bandgap is notably distinct from the high - temperature tetragonal phase of 0.41eV. As discussed above, the epitaxial growth al lows the stabilization of a new epi - tetragonal CsSnI 3 phase. Thus, the observation of a larger bandgap is no t surprising. Figure 8 . 11 . Absorption of the epit axial CsSnI 3 film and PL of the quantum wells. ( a ) A bsorption spectr um of 40nm CsSnI 3 grown on KCl substrate . The absolute absorption from the film is plotted. The Tauc plot shows the extrapolation of a direct bandgap of 1.85eV (inset). (b) PL spectra of quantum wells with well width of 17nm, 34nm and 60nm showing peaks at 838nm , 836nm and 843nm, respectively. 119 Quantum wells were fabricated with the halide perovskite as the well and KCl was used as a barrier to study the quantum confinement e ffect. KCl was grown under similar conditions to the homoepitaxial growth demonstrated previously. 267 RHEED was used during growth to confirm that each layer was crystall ine and smooth as indicated by streak ed patterns. PL spectra of the quantum wells show that by varying well width from 60 nm to 34 nm, the photoluminescence peak shifts modestly from 843 nm to 836 nm indicating that the Bohr radius of CsSnI 3 is relatively small (< 10nm). A similarly small quantum confinement effect has been reported for CsSnBr 3 nanocrystals and CsSnBr 3 quantum wells. 268, 275 Table 8 . 2 . Crystal structures, lattice parameters, experimental and simulated bandgaps of the CsSnI 3 cubic, tetragonal, orthorhombic and the epi - tetragonal phase from this work. Crystal structur es, lattice parameters, experimental and simulated bandgaps of the cubic, tetragonal, orthorhombic a nd the epi - tetragonal phase from this work. 1 64 Band structure and bandgap calculations were performing using the CASTEP module in Material Studio 7.0 with the B3LYP functional. 0.612±0.002 120 8 . 4 Epitaxial CsSnI 3 Thin Film Based Photodetector Fabrication Sing le crystalline CsSnI 3 thin film - based photodetectors were fabricated with the architecture shown in Figure 8.12 . The CsSnI 3 film serves as the photosensitive material to absorb light and Au electrodes were vapor deposited on the epitaxial CsSnI 3 layer to c ollect photogenerated charge carriers. The symmetrical and linear characteristics of the current - voltage (I - V) curve indica tes the ohmic contact between the Au electrode and CsSnI 3 . When the channel width rrent is over two magnitudes higher of the depletion Figure 8 .12 , the photocurrent is shown to continuously increase with increasing illum ination intensity. It is noted that the photocurrent increases significantly faster at the lower range of illumination inte nsity. The responsivity ( R ) is a key parameter commonly utilized to evaluate the performance of a photodetector. The responsivity is calculated as: , ( 8 . 1 ) where I light and I dark represent the current measured with or without light illumination respectively, P incident is the incident light power density, and A is the active area. R is typically reported as a fu nction of wavelength but can also be reported for broadband light sourc es. For this proof - of - principle demonstration we utilize the average incident white light power since we were better able to characterize the incident power density of white light versu s monochromatic light in this configuration. As shown in Figure 8 .12 , t he responsivity exhibits different trends from the photocurrent. As the responsivity is closely correlated to 121 Figure 8 . 12 . Epitaxial CsSnI 3 based p hotodetector fabrication and characterization. ( a ) A rchitecture of single crystalline CsSnI 3 thin film - based photodetector. simulated AM1.5 G solar illumination . ( b ) I - V characteristics of the device with 50 - wide channel wi dth under different illuminations. The arrows indicate the increase in the current when increasing the light intensity . ( c ) I - V - wide channel width under different illuminations. ( d ) I - V characteri stics of the devic - wide channel width under different testing voltages. ( e ) Photocurrent, - wide channel at 1V. 122 quantum efficiency, the quantum efficiency then decreases with increasing intensity. This reduction is likely due to an accumulation of space charge that creates additional carrier scattering and has been shown to turn off halide perovskite solar cells. 276 Nonetheless , the photo current is significantly h igher than the solution - processed lead - based hybrid halide perovskite CH 3 NH 3 PbI 3 incorporated in similar device configurations . 277 For the photodetector with an applied bias, the specific detectivity ( ) is limited by Shot noise as: , ( 8 .2) Hz 1/2 8 . 5 Conclusion In this chapter , we demonstrate a route to the room - temperature epitaxial growth of another lead - free inorganic perov skite CsSnI 3 on low - cost metal halide crystals via reactive thermal deposition characterized by in situ and ex situ diffraction techniques. This growth is investigated on KCl (100), KBr (100), and alloyed (NaCl/KCl) interlayers on KCl. The growth on KCl (100) reveals a ne w room temperature epitaxial tetragonal phase of CsSnI 3 that is closely lattice matched to the KCl. In contrast, the growth observed on KBr (with large misfit) results in quasiepitaxial films with the orthorhombic phase (common room temperature phase) and rotated microdomains. We further exploit the epitaxial growth of CsSnI 3 to demonstrate multilayer epitaxial quantum wells and lateral photodetectors. This work provides insight into the control over phase and ordering during halide perovskite epitaxial gro wth and expands the selection of photoactive materials for growing epitaxial halide perovskites that can be exploited in high performa nce electronic applications. 123 Chapter 9 Conclusions and Future Outlook In the first part of this work, organic crystalline i s sy stematically studied. H omoepitaxy growth mode s were mapped as a function of growth rate and temperature on high quality organic crystalline substrates. Organic - organic hetero - quasiepitaxy was studied showing energy mismatch (<20%) of lowest energy surfaces (LES) is ideal for achieving ordered alternating growth. A unique organic edge driven nucleation case was demonstrated provid ing new ideas for controlling molecular orientation. These growth studies on organic multilayer structures could motivate the expl oration of organic exciton confinement phenomena and opens new opportunities for enhanced excitonic tunability. The control over the crystalline order and orientation demonstrated in this work are the key to managing energy transport for novel orga nic semi conductor devices. In the second part of this work, we demonstrate a low - cost route to single domain halide perovskite thin film growth with excellent crystalline order that is lack ing with solution processing. We utilize in situ RHEED on insulatin g alkali halide crystalline surfaces to investigate the homo epitaxial growth studies of alkali halide crystals . Phase studies including phase control by manipulating precursor ratio and phase stabilization were demonstrated . P hotoluminescent tunability was observe d for epitaxial perovskite layer - based quantum wells with different well width. These rsults could spark the exploration of a full range of epitaxial halide perovskites and lead to novel applications for metal - halide - perovskite based single - crystal epitaxi al optoelectronics. Below, future work on organic - organic quasiepitaxy and perovskite epitaxy are discussed. 9 .1 Molecular Dynamics (MD) Simulation of Organic Surface Diffusion 124 Based on the initial exploration of organic surface diffusion using classical forcefield - based molecular dynamics, we developed algorithms that we have applie d to a large variety of admolecule/substrate systems for greater mo lecular combinations . The effect of molecular shape, molecular weight, and conjugation on surface diffusion behavior was studied with eight different organic admolecules. The molecules were selected by considering molecules that only hav e C - C and C - H bonds with varied shape, molecular weight, and aspect ratio. Simulations were performed on the same crystalline substrate, KBr (100) using Berendsen thermostat under COMPASS forcefield with the following procedure: 1) Start with a canonical isothermal ensemble (NVT) to set T , 2) Use a microcanonical constant energy ensemble (NVE) to run the experiment under more realistic fixed system energies, and 3) repeat the simulation for each Figure 9 . 1 .The set - up of the admolecule - substrate system . The anthracene on KBr (100) is shown as an example. 125 temperature with 30 - 40 runs for good statistical averaging . The example model fo r anthracene is shown in Figure 9.1. The diffusion process in these cases follow s the Arrhenius behavior ( Equation 2. 9 ) T he diffusivity ( D ) is extracted as the slope from the mean square displacement ( MSD ) as a function of time ( t ) according to ( Equation 2 .1 0 ). The diffusion barrier E m can be obtained from the slope of a ln D vs. 1/ T plot. Since D 0 is mechanistically composed of two contributions, the attempt frequency and the average hopping distance, we looked to separate the trend in the attempt frequency . Therefore, the trajectory is imported and analyzed in Matlab to extract the localized time ( % L ocalized ) and fractional time of a molecule travel along [10]&[01] direction of the substrate ( %T <01> ).The former should be inversely proportional to the attemp t frequency. Figure 9 . 2 . E m and D 0 plots as a function of L/W and molecular weight (M). E m plotted as a function of (a) L/W and (b ) M for the eight molecules studied. Corresponding D 0 plotted as a function of (c) L/W and (d) M. 126 In general, we found a complex dependence of activation energy with length/width ( L/W) aspect ratio. Close to an aspect ratio of 1, the activation barrier starts off very low and increases with increasing L/W , reaches a maximum, and then dec reases with L/W . For highly symmetric molecules it is not surprising to observe bo th low E m and D 0 . There is little driving force for the molecule to preferentially hop in a particular crystallographic direction resulting in little barrier for motion in an y particular direction but also resulting in only short hops. In contrast, for highly anisotropic molecules there can be stronger barriers to find alignment along molecular axes of the substrate that can result in larger ener getic barriers for hopping but larger hops when overcome. For D 0 we do not see a correlation to aspect ratio but rather see more straightforward trend of decreasing magnitude with increasing molecular weight ( Figure 9 . 2 ). We find that larger L/W increases the localized time on the subst rate, which is observed at both low T and high T conditions. Thus, we suggest that L/W > 1 can promote diffusion along crystallographic directions of the substrate indicated by longer fractional time along these trajectories ( Figure 9 . 3 ), which is more obv ious at high T . This study suggests the activation energy and the attempt frequency are closely related to the anisotropy of the molecule. Further studies are warranted on the simulation of admolecules on inorganic substrate s Figure 9 . 3 . (a) Fractional localized time ( %T Localized ) vs. L/W . (b) Fractional time of a molecule trave l along <01> ( %T <01> ) vs. L/W . 127 with other surface orientation s such as KBr (111), (110), etc. and on organic substrate s with different orientations (different surface energy). Step edge s can also be introduced and could provide computational evidence to the organic edge driven mechanis m discussed in Chapter 5 . 9 . 2 O rganic - organic Charge Transfer (CT) Complexes In Chapter 4 and 5 , we demonstrated t he ability to grow ordered organic crystalline structures using bottom - up vapor - deposition routes . A key challenge moving forward is the ability to integrate conductive lay ers into organic quasiepitaxial multilayers to fabricate complete organic epitaxial devices. However, the growth of pure metals in these systems generally prevent the ability to maint ain organic crystalline ordering grown from the bottom up. O rganic CT com plexes provide an enticing solution to this problem . We have demonstrated sustained alternating quasi - epitaxial layers of two organic semiconductors 1,4,5,8 - naphthalene - tetracarboxyli c - dianhydride (NTCDA), and dibenzotetrathiafulvalene - tetracyanoquinodimet hane (DB - TCNQ) grown on single crystal substrates via vapor phase deposition. Both in - plane and out - of - plane ordering were Figure 9 . 4 . Preliminary results on alternating growth of charge transfer salts. (a) In situ RHEED patterns of alternating NTCDA/DBTTF - TCNQ growth on single crystal sub strates. Ordering for over 2 pairs has been achieved with this combination, where there is a close surface energy matching of the LESs. ( b - c ) RHEED patterns of TTF - TCNQ and TCNQ growth. ( d ) RHEED pat terns of alternating TCNQ/TTF growth . 128 preserved for each subsequent layer under optimized growth c onditions despite the vastly differing lattice constants and lattice symm etries ( Fig ure 9.4 a ). Other potential starting layers such as tetrathiafulvalene 7,7,8,8 - tetracyanoquinodimethane (TTF - TCNQ) and TCNQ were preliminarily investigated using in situ dif fraction ( Figure 9.4 b - c ) . Both materials show single crystalline like gro wth on KBr suggested by the rotation dependence of RHEED patterns. Paired growth of TCNQ/TTF have also been studied and show viable crystalline growth across the multilayer structure ( Figure 9.4 d ). Based on these results , more growth pairings should be stu died with greater growth optimization to identify and map quasi - epitaxial growth modes of select organic charge transfer single crystal thin films . 9 . 3 Expand Halide Perovskite Heteroepitaxial Growth Pairings In C hapter 7 and 8 , the heteroepitaxial growth of CsSnBr 3 and CsSnI 3 on metal halide single crystals are discussed. It is important to expand the se growth stud ies to other lead - free perovskite compositions including CsSnCl 3 ( to complete the CsSnX 3 se ries ) and double perovskites such as Cs 2 AgIn Cl 6 , Cs 2 AgBiX 6 , Cs 2 AuBi X 6 , Cs 2 CuBi X 6 (X=Cl, Br, and I), etc. 279 as well as lead based compounds. Such future studies can start with metal halide single crystals substrates that are most closely matched to each composition. Parameters such as growth temperature, pressure, stoichiometry and strain can be studied to clarify the specific accessible phase is at certain growth condition. Lattice enginee r ing (described in Chapter 7 and 8 ) can be used to modify the surface lattice constant of the substrate by adding a pseudomorphic alloyed layer based on the Vegard's rule, ( 10.1 ) 129 where the alloyed lattice constant ( a k ) is a lin ear function of the lattice constants from the two constituent materials with lattice constant as a i and a j . This method has been proved to help growing smooth crystalline growth based on the CsSnBr 3 growth on alloyed NaCl/NaBr and CsSnI 3 growth on alloyed KCl/NaCl that are discussed in Chapter 7 and 8 . 9 . 4 Epitaxial Halide Perovskite Doping To open up the possibilities for epitaxial perovskite film - based applications, controlled doping achieved by neat layer doping or modulation (delta) doping can be s tudied . 280 Neat layer dopin g is a method that can be achieved by co - evaporation of dopant and perovskite composition. The direct neat layer doping could affect the structural integrity of the quantum wells , thus modulation (delta) doping can be introduced . In modulation (delta) dopi ng, the dopants are spa tially separated from the well barrier layers so that the epitaxial layer is deposited on top of a doped single crystalline substrate . 9 . 5 Epitaxial Halide Perovskite Applications After we understand the effect of doping method s and concentration on these epitaxial halide perovskites layers, it is possible to incorporate them into multilayer structures to build a range of interesting electronic devices including photodetectors, light emitting diodes ( LED) , and high mobility transistors. Since tunable optoe lectronic properties were demonstrated with halide perovskite quantum well (QW) structures in Chapter 7 and 8 , transport measurements should be performed o to explore interesting low temperature and q uantum - based physics. 130 APPENDIX 131 List o f publications and patent from this thesis (1) (2) (3) (4) (5) Wang, L., Chen, P., Lunt, R. R. Method for Fabricating Epitaxial Halide Per ov skite Films and Devices, filed as U.S. provisional patent application on June 13, 2017. (1) Traverse, C. J.; Chen, P.; Lunt, R. R. Lifetime of Organic Salt Photovoltaics. Adv. Energy. Mater. 2018, 8 (21), 1703678. (2) Liu, D.; Traverse , C. J.; Chen, P.; Elinski, M.; Yang, C.; Wang, L.; Young, M.; Lunt, R. R. Adv. Sci. 2018, 5 (1), 1700484. (3) Liu, D., Yang, C., Chen, P., Bates, M., Han, S., Askeland, P.A. and Lun t, R.R. Lead Halide Ultraviolet - Harvesting Transparent Photovoltaics with an Efficiency Exceeding 1%. ACS Appl. Energy Mater. 2019 (4) Kuttipillai, P. S., Yang, C., Chen, P., Wang, L., Bates, M., Lunt, S. Y., & Lunt, R. R. Enhanced Electroluminescence Efficie ncy in Metal Halide Nanocluster Based Light Emitting Diodes through Apical Halide Exchange. ACS Appl. Energy Mater. 2018, 1(8), 3587 - 3592. (5) Liu, D.; Wang, Q.; Elinski, M.; Chen, P.; Traverse, C. J.; Yang, C.; Young, M.; Hamann, T . W.; Lunt, R. R. Ultrathin Hole Extraction Layer for Efficient Inverted Perovskite Solar Cells. ACS Omega 2018, 3 (6), 6339 - 6345. (6) Wang, L.; Moghe, D.; Hafezian, S.; Chen, P.; Young, M.; Elinski, M.; Martinu, L.; K na - Cohen, S. p.; Lunt, R. R. Alkali met al halide salts as interfac e additives to fabricate hysteresis - free hybrid perovskite - based photovoltaic devices. ACS Appl. Mater. Interfaces 2016, 8 (35), 23086 - 23094. (7) Traverse, C.J., Young, M., Suddard - Bangsund, J., Patrick, T., Bates, M., Chen, P., Win gate, B., Lunt, S.Y., Ancti l, A. and Lunt, R.R. Anions for near - infrared selective organic salt photovoltaics. Sci. Rep. 2017, 7(1), p.16399. 132 Figure A . 1 . Cover Pict ure for Advanced Materials Interfaces Volume 3 , Issue 1 7 ( September 2016 ) . Corresponding article: Organic Step Edge Driven Heteroquasiepitaxial Growth of Organic Multilayer Films . 133 Figure A. 2 . Cover Picture for Advanced M aterials Interfaces Volume 4, Issue 22 (November 2017). 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