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Development and characterization of transparent metal/ceramic and ceramic/ceramic nanomultilayers

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Development and characterization of transparent metal/ceramic and ceramic/ceramic nanomultilayers

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Content DEVELOPMENT AND CHARACTERIZATION OF TRANSPARENT METAL/CERAMIC AND CERAMIC/CERAMIC NANOMULTILAYERS by Chelsea Appleget A Dissertation Presented to the FACULTY OF THE USC GRADUATE SCHOOL UNIVERSITY OF SOUTHERN CALIFORNIA In Partial Fulfillment of the Requirements for the Degree DOCTOR OF PHILOSOPHY (AEROSPACE ENGINEERING) August 2020 ii To John Appleget and Jack Appleget iii Acknowledgments First and foremost, I would like to thank my advisor Professor Andrea Hodge, for the guidance and opportunity to work in her group. She has trained me as researcher, but has also invested in me as a person, and it has been an honor to work with someone that I respect so much. Thank you for believing in me and also for pushing me. You have shown me what strength is—in research, character, in every facet, and I am grateful to have worked with you. I would also like to acknowledge the Core Center of Excellence in Nano Imaging (CNI) and Matthew Mecklenberg for always answering questions and providing advice. I am grateful for the staff and faculty at USC that have helped me along the way, especially Kim Klotz, Jennifer Gerson, Natalie Guevara, and Samantha Graves, among many others, your time and advice has always been appreciated. Thank you also to my other wonderful mentors along the way, including but not limited to, Professor Roy Hartfield, Professor David Cicci, Professor Michael Kassner, Dr. Patrick Johnston, and Tim Carlton. Thank you to Professor Paulo Branicio, Professor Mitul Luhar, and Dr. Jonathan Salem from NASA GRC for agreeing to serve on my dissertation committee, as well as for encouraging me and providing their expertise during my entire PhD journey. Funding for this project was provided through Office of Naval Research Grant N00014- 18-1-2263. I would also like to thank NASA for support through the NASA Space Technology Research Fellowship (80NSSC17K0160) as well as the USC PhD Provost Fellowship. This work was also made possible through the generous hosting of distinguished researchers at NASA Glenn Research Center, including Dr. Jonathan Salem for mentorship and guidance, Dr. Anita Garg, Dr. Wayne Jennings, Dr. Tim Smith, and Pete Bonacuse for microscopy training as well as helpful technical discussions, and Brian Vyhnalek for ellipsometry. iv There is no possible way to individually thank each friend and family member for their love and encouragement, but I want to thank my family and friends for their unconditional support—both those that are here with me and those that we wish could be here. Thank you to my parents, sisters, and entire family for always encouraging me to follow big dreams, I love you to the moon and back. Thank you also to my many supportive, unwavering friends, Dr. Becky Kimball, Hayley Cox, and so many more, how lucky I am for each of you. To Allyssa VelDink and Joel Bahena, you two have shown me what it means to be truly good and kind, have cared for me during every up and every down and I hope you know how grateful I am. I want to also thank all members of the Hodge group, thank you for your support and for the fun times every step of the way, I am surrounded by brilliant colleagues but also good friends, including: Dr. Leonardo Velasco, Dr. Mikhail Polyakov, Dr. Nathan Heckman, Dr. Teri Juarez, Dr. Angelica Saenz-Trevizo, Dr. Sebastián Riaño, and Dr. Joel Bahena, each of whom have patiently taught and helped me. Alina Garcia-Taormina, Daniel Goodelman, Marion Wood, Karina Hemmendinger, Natalie Reck, Adie Alwen, Danielle White, and Roya Ermagan—thank you for this fun ride. I am proud of each of you. I’d be remiss if I didn’t thank www.Thesaurus.com, for supporting me, encouraging me, and helping me through this PhD—it never failed to help me find the right, suitable, and appropriate word anytime that I needed or required it. And to Stephen Samples, thank you for believing in me, your patience, your support, and for nearly ten years of friendship, love, and silliness. Thank you for being by my side for every adventure, including this one, I love you. v Table of Contents Dedication .................................................................................................................................. ii Acknowledgments ..................................................................................................................... iii List of Figures .......................................................................................................................... vii List of Tables .......................................................................................................................... xiv List of Abbreviations .............................................................................................................. xvi Abstract ................................................................................................................................. xviii Chapter 1 : Introduction ..............................................................................................................1 Chapter 2 : Background ..............................................................................................................4 2.1. Nanomultilayers ..............................................................................................................4 2.2. Transparent and Optical NMs .........................................................................................8 2.3. Optical Behavior of NMs ..............................................................................................14 2.4. Mechanical Behavior of NMs .......................................................................................23 Chapter 3 : Experimental Methods and Materials ....................................................................39 3.1. Magnetron Sputtering ....................................................................................................39 3.2. Microstructural Characterization Methods ....................................................................46 3.3. Optical Characterization Methods .................................................................................56 3.4. Mechanical Characterization .........................................................................................63 Chapter 4 : Optical and Mechanical Characterization of Metal/Ceramic NMs ........................68 4.1. Introduction ...................................................................................................................68 4.2. MBI Predictions ............................................................................................................70 4.3. Synthesis of AlN/Ag Multilayers ..................................................................................71 4.4. XRD and Optical Characterization of AlN/Ag Multilayers ..........................................72 4.5. Microstructural and Mechanical Characterization of AlN/Ag Multilayers ..................75 4.6. Conclusions ...................................................................................................................80 Chapter 5 : Microstructural Variations in Ceramic/Ceramic AlN/SiO2 NMs ...........................81 5.1. Introduction ...................................................................................................................81 5.2. Experimental Methods ..................................................................................................83 5.3. Results and Discussion ..................................................................................................84 5.4. Conclusion .....................................................................................................................93 Chapter 6 : High Transparency Ceramic Crystalline/Amorphous and Amorphous/Amorphous NMs ...................................................................................................95 6.1. Introduction ...................................................................................................................95 6.2. Experimental Methods ..................................................................................................97 6.3. Results and Discussion ..................................................................................................98 vi 6.4. Outlook and Conclusion ..............................................................................................110 Chapter 7 : Conclusions and Future Work ..............................................................................112 7.1. Conclusions .................................................................................................................112 7.2. Future Work ................................................................................................................114 References ...............................................................................................................................116 Appendix A: Summary of Sputtered Monolithic Samples .....................................................131 Appendix B: Summary of Sputtered Multilayers ...................................................................136 Appendix C: Spectroscopic Ellipsometry Measurements and Optical Properties of Monolithic Sputtered Films ....................................................................................................140 Appendix D: XRD Studies and Results ..................................................................................144 Appendix E: In-House MBI MATLAB Code .........................................................................150 vii List of Figures Figure 1: Schematic of a multilayer ............................................................................................4 Figure 2: Strain-free coherent interfaces: (a) coherent interface with two crystals of different chemical compositions with the same crystal structure and (b) two phases of different lattices [23]. .............................................................................................................................................6 Figure 3: Overview of different interface types: (a) a coherent interface with a slight mismatch in lattice parameters, leading to coherency strains, (b) a semi-coherent interface where the dislocations relieve the coherency strain and (c) an incoherent interface with fully overlapped dislocations[23]. .......................................................................................................7 Figure 4: Overview of amorphous interface types: (a) a crystalline/amorphous interface and (b) an amorphous/amorphous multilayer, modified from [30]. ..................................................8 Figure 5: SEM cross section of an antireflection coating comprised of SiOx, TiO2 and SiOx- TiO2 nano-multilayers (NML) [37]. ............................................................................................9 Figure 6: Structure-zone diagram for metal films deposited via magnetron sputtering [85]. ...12 Figure 7: Effect of oxygen flow rate on refractive index n and extinction coefficient k [88] ..13 Figure 8: Changes in refractive index n of TiO2 films deposited at ambient temperature (N26), 350°C (N24) and 500°C (N21) [89] ..........................................................................................13 Figure 9: An example of an optical multilayer showing nomenclature used for indicating thickness, layer number (m), refraction index (n), and angle of refraction (α) [93]. ................17 Figure 10: Reflectance vs. wavelength in antireflection coatings [94]. ....................................18 Figure 11: Reflectance for high-reflection multilayers [94]. ....................................................18 Figure 12: UV/Vis and NIR wavelength spectrum (in nm) [97]. .............................................19 Figure 13: a) Measurement and theoretical MBI calculations in SiOx/TiO2 NMs with inset showing the sample structure and b) process window simulation showing the defect of thickness deviation on overall transmittance for a four layer ACR coating, where ideal thicknesses are SiOx (84 nm)/TiO2 (110 nm)/SiOx (35 nm)/TiO2 (11 nm) corresponding to an optical transmittance of 99.34% [37]. .......................................................................................22 Figure 14: Schematic of the MBI recursive method for a multilayer system of m layers, with the calculation starting from the a) m th layer reflection amplitude coefficient to produce the b) effective reflection amplitude coefficient of the m th layer, reff,m, the process continues until the viii c) effective reflection amplitude coefficient of the 1 st layer is calculated, then d) the total reflection and transmission of the m-layer multilayer system is calculated. ............................23 Figure 15: Hardness values for metallic NMs systems as a function of the inverse square of the layer thickness [61]. ............................................................................................................24 Figure 16: Yield strength as a function of the inverse square root of the grain size, for nanocrystalline materials. There are three regimes for flow stress as a function of the grain size, in the > 100 nm regime, there is Hall-Petch strengthening, with a slope of k. At grain sizes of 100 nm and 10 nm the slope changes, which is indicative of the dislocation mechanism changing. ................................................................................................................26 Figure 17: Illustration of the dislocation mechanisms of metallic NM strength contributions operating at different length scales under an applied force normal to the layer interfaces [3]. 27 Figure 18: Deformation morphology of a crystalline/amorphous Cu40/PdSi10 micro pillar [103]. .........................................................................................................................................28 Figure 19: A schematic illustrating the deformation mechanism proposed in crystalline/amorphous Cu/CuNb multilayers [30] ....................................................................29 Figure 20: Schematic illustrating the competing governing mechanisms, localized shearing and shear band interactions, in amorphous/amorphous NMs [147]. .........................................30 Figure 21: (a) FIB cross-section of an indentation of a Al/SiC nanomultilayer, showing shear bands and SiC cracking, (b) higher magnification of SiC flexibility and microcracking, and (c) delamination of the multilayers/Si interface and fracture within the Si substrate. ...................31 Figure 22: (a) SEM cross-section showing the deformation mechanisms of Al and TiN multilayers, where the multilayer has been indented with a Berkovich tip. (b) Higher magnification shows that Al layers carry out plastic deformation by dislocation glide, while the TiN layers have cracking. These cracks are filled by the Al layer [157]. (c) Transmission electron microscopy (TEM) bright field cross-section showing nanoindentation W/NbN multilayers showing plastic deformation in the W layers with twin-like features (straight lines) in the NbN layers [5]. ......................................................................................................32 Figure 23: The effect of bilayer period on hardness for TiN/W multilayers [112]. ..................33 Figure 24: The effect of bilayer period on nanoindentation hardness and dissipation modulus on CN/BN multilayer films [122]. ............................................................................................36 Figure 25: Experimental hardness vs. superlattice (bilayer) period compared with model predictions for dislocation glide across and within layers in single and polycrystalline TiN/NbN multilayers [171]. ......................................................................................................37 Figure 26: Mechanisms of toughness enhancement in hard ceramic/ceramic multilayers [176]. ...................................................................................................................................................38 ix Figure 27: Schematic showing the setup and process of nonreactive magnetron sputtering. A neutral gas (Ar) is introduced and ionized by the negatively biased target, the Ar ions then strike the target and cause the target atoms to be ejected and coat the substrate. .....................41 Figure 28: Schematic of sputtering setup for depositing multilayer systems. The multilayer thickness is determined by the sputtering rate and ‘on time’ of each source, as controlled by the pneumatic shutter, (a) sputtering from source 1 while the shutter blocks source 2, and (b) sputtering from source 2 while the shutter blocks source 1. .....................................................41 Figure 29: Hysteresis plot for the reactive sputtering of Ti in an Ar/O2 atmosphere with flow control of the reactive gas [188]. ..............................................................................................43 Figure 30: Schematic of a stylus profilometer [190]. ...............................................................44 Figure 31: Substrate profile showing radius of curvature before and after film deposition, measured using profilometry. ...................................................................................................45 Figure 32: The incident beam (left) interacts with the atoms of spacing d and scatters. ..........47 Figure 33: Normalized XRD patterns of sputtered Al, Al-Mg, and Ni films [195]. .................47 Figure 34: (a) Electron and photon signals emanating from the tear-shaped interaction volume of the specimen, (b) energy spectrum of electrons emitted from the specimen surface and (c) the effect of the surface topography of the sample on electron emission [197]. ......................48 Figure 35: SEM image of a Vickers indent on a ceramic nanomultilayer sample, synthesized by C. Appleget. .........................................................................................................................49 Figure 36: Schematic of a dual beam FIB and SEM instrument with the inset showing the electron and ion beam interaction with the sample [198]. ........................................................51 Figure 37: Overview of FIB cross-section showing (a) top view of Vickers indentation sample with dotted line indicating the milling region and (b) cross-section of the indent produced via FIB milling (synthesized by C. Appleget). ...............................................................................51 Figure 38: Overview of FIB lift out technique beginning with a) carbon protective layer on top of sample and trenches milled around region of interest, b) preparation of peninsula with bright layers of multilayer sample visible, c) attachment of OmniProbe needle for lift out, and d) thinning of lift out specimen on TEM grid until desired thickness (prepared by C. Appleget). ..................................................................................................................................52 Figure 39: TEM images showing the depth of amorphization of Si prepared by Ga + and Xe + FIB [199]. ..................................................................................................................................52 Figure 40: Schematics of (a) bright field, (b) dark field, and (c) selected area diffraction TEM operating modes [200]. .............................................................................................................53 x Figure 41: (a) Bright field and (b) dark field cross-sectional TEM micrographs of an AlN/SiO2 multilayer sample, synthesized by Appleget. ...........................................................54 Figure 42: (a) Cross-sectional HAADF STEM of an AlN/SiOx multilayer sample and STEM EDS maps showing the distribution of (b) Al, (c) N, (d) Si, and (e) O (synthesized by C. Appleget). ..................................................................................................................................55 Figure 43: A summary of the interaction between light and matter, with the approximate energies of fundamental excitations [203]. ...............................................................................56 Figure 44: Schematic of a UV-Vis Spectrophotometer ............................................................58 Figure 45: Interaction of polarized light with a sample [206]. .................................................59 Figure 46: Characterization of physical properties by spectroscopic ellipsometry [208]. ........60 Figure 47: Ellipsometry curve fittings for TiO2 sample prepared in 10% oxygen environment [212]. .........................................................................................................................................61 Figure 48: Variation of refractive index (n) with % oxygen content in a TiO2 thin film [212]. ...................................................................................................................................................62 Figure 49: Optical constants (refractive index and extinction coefficient) of an 87 nm AlN thin film deposited by plasma-enhanced atomic layer deposition (PEALD) [213]. .................62 Figure 50: a) Schematic showing the indent geometry and dimensions, and b) typical nanoindentation load-displacement curve [215]. ......................................................................64 Figure 51: Load-displacement curves from nanoindentation testing of an AlN/SiO2 sample, 10 indents of a 10x10 (100 indent) array are shown (by C. Appleget). .........................................65 Figure 52: Overview of Vickers Indenter geometry [217]. .......................................................66 Figure 53: Comparison of several hardness scales [218]. .........................................................67 Figure 54: a) SEM isometric view of Vickers indent (with Pt coating to prevent charging of sample) and b) FIB top view of the same Vickers indent (by C. Appleget). ............................67 Figure 55: SEM cross-sections showing multilayers and corresponding top views indicating appearance of samples: (a) AlN(47)/Ag(21), (b) AlN(20)/Ag(20), (c) AlN(127)/Ag(10), (d) AlN/Ag MBI #1, and (e) AlN/Ag MBI #2 where the layer thicknesses are indicated in parentheses in nanometers. .......................................................................................................73 Figure 56: (a) Normalized X-Ray Diffraction (XRD) patterns of AlN/Ag multilayer system from 2θ 10-100⁰ and corresponding crystal structure where insets show detail of low intensity peaks and (b) Experimental transmittance curves of the AlN/Ag multilayers performed by spectrophotometry. ....................................................................................................................74 xi Figure 57: Representative top-view optical microscope images of indents after Vickers indentation (performed at 2.94 N with a 10 s dwell time) for multilayers samples: (a) AlN(47)/Ag(21), (b) AlN(20)/Ag(20), (c) AlN(127)/Ag(10), (d) AlN/Ag MBI #1, and (e) AlN/Ag MBI #2. .......................................................................................................................77 Figure 58: FIB and TEM cross-sectional views of AlN/Ag MBI #1 (a-c) and MBI #2 (d-f). (a) MBI #1 FIB cross-section of Vickers indent post-deformation and (b) bright field STEM with SAED inset of the as-sputtered sample including (c) bright field TEM showing a representative cross-section of an Ag layer breakthrough in MBI #1. (d) MBI #2 FIB cross- section of Vickers indent post-deformation, (e) bright field STEM with SAED inset of the as- sputtered sample including (f) bright field TEM showing a representative cross-section of a continuous Ag layer in MBI #2. In FIB cross-sections, the protective top layer of carbon is labeled; white arrows highlight regions of delamination and black arrows indicate through- layer cracking. ...........................................................................................................................78 Figure 59: Overview of AlN/SiO2 nanomultilayers, where (a) is a schematic of alternating and (b) representative cross-sectional TEM nanocrystalline AlN and amorphous SiO2, as indicated with the inset SAED patterns. (c) Experimental % transmittance curves, where layer thicknesses of repeated bilayer samples are indicated in parentheses in nanometers. Top views of as-deposited samples (d-g): (d) AlN(100 nm)/SiO2(100 nm), (e) AlN(50 nm)/SiO2(50 nm), (f) MBI 10 Layers, and (g) MBI 20 Layers. .............................................................................86 Figure 60: Overview of AlN(50 nm)/SiO2(50 nm) repeated bilayer sample, (a) BF cross- sectional TEM and inset SAED pattern of the multilayer film with yellow box indicating the region of (b) BF HRTEM of the AlN-SiO2 interface region. Red dotted line indicates the path of the 1D compositional line profile shown in (d). (c) Integrated radial intensity profiles interpolated from the SAED pattern and (d) the 1D compositional line profile as measured by STEM EDS, showing atomic % fraction of elements across the multilayers. ..........................88 Figure 61: Overview of AlN(100 nm)/SiO2(100 nm) repeated bilayer sample, (a) BF cross- sectional TEM with inset SAED pattern and (b) BF HRTEM of the AlN-SiO2 interface region. (c) Integrated radial intensity profiles interpolated from the SAED pattern and (d) the 1D compositional line profile as measured by STEM EDS, showing atomic % fraction of elements across the multilayers. ................................................................................................89 Figure 62: Overview of AlN/SiO2 MBI 20 Layer sample, (a) BF cross-sectional TEM with inset SAED pattern and (b) BF HRTEM of the AlN-SiO2 interface region. (c) Integrated radial intensity profiles interpolated from the SAED pattern and (d) the 1D compositional line profile. .......................................................................................................................................90 Figure 63: Overview of AlN/SiO2 MBI 10 Layer sample, (a) BF cross-sectional TEM with inset SAED and (b) BF HRTEM of the AlN-SiO2 interface region. (c) Integrated radial intensity profiles interpolated from the SAED pattern and (d) the 1D compositional line profile. .......................................................................................................................................92 xii Figure 64: Overview of ceramic nanomultilayer systems, including AlN/SiO2, TiO2/SiO2 and AlN/Al2O3, where (a) is experimental % transmittance curves of repeated bilayer samples with layer thicknesses indicated in parenthesis in nanometers and (b) experimental % transmittance curves of the MBI optically optimized samples. Top views of as-deposited samples (b-h): repeated bilayer (c) AlN(50)/SiO2(50), (d) AlN(50)/Al2O3(50), and (e) TiO2(50)/SiO2(50) and optically optimized MBI samples (f) AlN/SiO2 MBI, (g) AlN/Al2O3 MBI, and (h) TiO2/SiO2 MBI. .................................................................................................102 Figure 65: Representative BF TEM cross-sections of multilayers, with repeated bilayer configurations on the left side (a-c) and optically optimized MBI (d-f) configurations on the right side. The corresponding compositions are labeled on the left: (a,d) AlN/SiO2, (b,e) AlN/Al2O3, and (c,f) TiO2/SiO2. Inset SAED patterns show the crystalline or amorphous nature of the samples and the red dotted arrows indicate the growth direction of the films. .105 Figure 66: Representative BF HRTEM cross-sections of multilayers, with repeated bilayer configurations on the left side (a-c) and optically optimized MBI (d-f) configurations on the right side. The corresponding compositions are labeled on the left: (a,d) AlN/SiO2, (b,e) AlN/Al2O3, where interlayer regions are indicated by white lines, and (c,f) TiO2/SiO2, where nanocrystalline regions are indicated by the yellow ellipses with the corresponding inset FFT. .................................................................................................................................................107 Figure 67: Plots of multifunctional properties of optically optimized samples as a function of multilayer characteristics including (a) experimental hardness and transmittance as a function of amorphous volume fraction and (b) as a function of the normalized # of interfaces per µm. .................................................................................................................................................110 Figure 68: Psi and delta values of a sputtered AlN single layer sample (from C. Appleget). 141 Figure 69: Index of refraction and extinction coefficient generated by model fitting of measurements from a sputtered AlN single layer sample (from C. Appleget). ......................141 Figure 70: AlN thin films prepared via magnetron sputtering, showing A) % Transmittance for the three samples, where average transmittance decreases slightly with increasing deposition power, and decreases significantly with a higher Ar flow rate and higher working pressure and B-D) top views of the as-sputtered samples (from C. Appleget). ......................142 Figure 71: Psi and delta values of a sputtered Ag single layer sample (from C. Appleget). ..142 Figure 72: Index of refraction and extinction coefficient generated by model fitting of measurements from a sputtered Ag single layer sample (from C. Appleget). ........................143 Figure 73: XRD spectra of as-sputtered AlN-SR (sputtering rate) samples deposited by reactive DC magnetron sputtering on Corning Eagle 2000 substrates with varying deposition conditions. Sputtering parameters varied include power, non-reactive Ar gas flow, and reactive N2 gas flow. ...............................................................................................................145 xiii Figure 74: XRD spectra of as-sputtered AlN-SR (sputtering rate) samples deposited by reactive RF magnetron sputtering on Corning Eagle 2000 substrates with varying deposition conditions. Sputtering parameters varied include power, non-reactive Ar gas flow, and reactive N2 gas flow. ...............................................................................................................146 Figure 75: XRD spectra of as-sputtered ITO-SR (sputtering rate) samples deposited by non- reactive RF magnetron sputtering on Corning Eagle 2000 substrates with varying deposition power. ......................................................................................................................................146 Figure 76: XRD spectra of as-sputtered SiOx-SR (sputtering rate) samples deposited by reactive RF magnetron sputtering on Corning Eagle 2000 substrates with varying deposition conditions. Sputtering parameters varied include power, non-reactive Ar gas flow, and reactive O2 gas flow. All SiOx spectra are amorphous with no crystalline peaks detected. ...147 Figure 77: XRD spectra of as-sputtered TiOx-SR (sputtering rate) samples deposited by reactive RF magnetron sputtering on Corning Eagle 2000 substrates with varying deposition conditions. Sputtering parameters varied include power, non-reactive Ar gas flow, and reactive O2 gas flow. All TiOx spectra are amorphous with no crystalline peaks detected. ...148 Figure 78: XRD spectra of as-sputtered ZnO-SR (sputtering rate) samples deposited by non- reactive RF magnetron sputtering on Corning Eagle 2000 substrates with varying deposition power. ......................................................................................................................................148 Figure 79: XRD spectra of as-sputtered Ag-SR (sputtering rate) samples deposited by non- reactive DC magnetron sputtering on Corning Eagle 2000 substrates with varying non- reactive Ar flow. .....................................................................................................................149 Figure 80: (left) theoretical prediction of a 10 layer AlN/SiO2 NM and (right) associated layer thicknesses ..............................................................................................................................153 Figure 81: Spectroscopic predicted (MBI) vs. experimental % transmittance for AlN/SiO2 MBI 20 Layer sample .............................................................................................................154 Figure 82: Spectroscopic predicted (MBI) vs. experimental % transmittance for repeated bilayer AlN(50)/SiO2(50) ........................................................................................................154 xiv List of Tables Table 1: Predicted layer thicknesses of optimized MBI samples .............................................71 Table 2: Results of Experimental and Theoretical Optical and Mechanical Properties. ..........75 Table 3. Summary of layer thicknesses of MBI 10 Layer and MBI 20 Layer Samples ...........85 Table 4: Results of experimental mechanical and optical properties, including total film thickness, residual stresses as measured by profilometry, nanoindentation results, and experimental % transmittance. ..................................................................................................93 Table 5: Summary of layer thicknesses in AlN/SiO2 MBI 20 Layer sample ............................99 Table 6: Summary of layer thicknesses in AlN/Al2O3 MBI 20 Layer sample ........................100 Table 7: Summary of layer thicknesses in TiO2/SiO2 MBI 20 Layer sample .........................101 Table 8: Experimental results of mechanical and optical properties ......................................102 Table 9: Key for Types of Characterization Techniques ........................................................131 Table 10: Summary of Types of Sputtering Used for Thin Films Synthesized ......................132 Table 11: Ag Sputtered Samples via Non-Reactive DC Sputtering .......................................132 Table 12: Al2O3 Sputtered Samples via Non-Reactive RF Sputtering ...................................132 Table 13: Al2O3 Sputtered Samples via Reactive RF Sputtered (KIT) ...................................132 Table 14: AlN Sputtered Samples via Reactive DC Sputtering ..............................................133 Table 15: AlN Sputtered Samples via Reactive RF Sputtering (KIT) ....................................133 Table 16: ITO Sputtered Samples via Non-Reactive RF Sputtering (KIT) ............................134 Table 17: SiO2 Sputtered Samples via Non-Reactive RF Sputtering .....................................134 Table 18: SiOx Sputtered Samples via Reactive RF Sputtering (KIT) ...................................134 Table 19: TiN Sputtered Samples via Reactive DC Sputtering ..............................................135 Table 20: TiO2 Sputtered Samples via Non-Reactive RF Sputtering .....................................135 Table 21: TiOx Sputtered Samples via Reactive RF Sputtering (KIT) ...................................135 Table 22: ZnO Sputtered Samples via Non-Reactive RF Sputtering (KIT) ...........................135 xv Table 23: Sputtered AlN/Ag Multilayer Samples ...................................................................136 Table 24: Sputtered AlN/Al2O3 Multilayer Samples ..............................................................137 Table 25: Sputtered AlN/SiO2 Multilayer Samples ................................................................137 Table 26: Sputtered AlN/SiOx Multilayer Samples (KIT) ......................................................137 Table 27: Sputtered AlN/TiN Multilayer Samples .................................................................138 Table 28: Sputtered TiO2/Al2O3 Multilayer Samples .............................................................138 Table 29:Sputtered TiO2/SiO2 Multilayer Samples ................................................................138 Table 30: Sputtered TiOx/SiOx Multilayer Sample (KIT) ......................................................139 Table 31:Sputtered ZnO/ITO Multilayer Sample (KIT) .........................................................139 xvi List of Abbreviations A/A Amorphous/amorphous ALD Atomic layer deposition ARB Accumulative roll bonding BF Bright field BKD Backscatter Kikuchi diffraction C/A Crystalline/amorphous CLS Confined-layer slip CTE Coefficient of thermal expansion CVD Chemical vapor deposition DC Direct current DF Dark field DP Diffraction Pattern EBSD Electron back-scattered diffraction EDS Energy dispersive x-ray spectroscopy FCC Face-centered cubic FIB Focused ion beam GIS Gas injection system HAADF High-angle annular dark field HCP Hexagonal close-packed HRTEM High-resolution transmission electron microscopy HWOT Half-wave optical thickness IBS Inter-boundary slip MBE Molecular beam epitaxy MBI Multiple beam interference MFC Mass flow controller MSE Mean squared error NIR Near infra-red NM Nanomultilayer PVD Physical vapor deposition QWOT Quarter-wave optical thickness R Reflectance RF Radio frequency SAED Selected area electron diffraction SEM Scanning electron microscopy xvii STEM Scanning transmission electron microscopy T Transmittance TCO Transparent conductive oxide TEM Transmission electron microscopy UV Ultraviolet VASE Variable angle spectroscopic ellipsometry XRD X-ray diffraction xviii Abstract Nanomultilayer (NM) systems have demonstrated desirable properties such as corrosion resistance, radiation tolerance, and high strength due to the interfaces and length scale effects. Additionally, multilayers have commonly been used in optical systems as the layer thicknesses can be tuned to be identical to interaction lengths of photons. Although NMs have been shown to have useful optical and mechanical properties independently, few studies have been performed to investigate the interplay between optics and mechanics. Thus, the possible novel combination of properties, such as transparency and strength, presents an extraordinary area of research with potential applications in optical windows, sensor protection, and numerous other functions which require light penetration for function and a robust barrier for protection. The studies described in this dissertation provide a foundation for understanding the relationship between length-scale effects, interface effects, transparency, and multifunctional properties in ceramic/ceramic and metal/ceramic NMs. Specifically, this relationship was explored through (1) the development of an approach to synthesize and study the optical properties of multilayers in a model AlN/Ag system, then (2) the investigation of microstructural effects of maximizing transparency in AlN/SiO2 multilayers and finally (3) the exploration of amorphous/amorphous and crystalline/amorphous interface contributions in highly transparent ceramic/ceramic multilayers. By identifying the dominant mechanisms in the relationship between optical performance and resultant film properties, transparent optical multilayer systems can be tailored with optimal layer thicknesses to aid in the design and synthesis of new high-performance, long-lasting optical materials. 1 Chapter 1 : Introduction A multilayer is a composite structure typically made of one or more constituent materials, with layers that can range in size from microns to nanometers depending on the types of processing and applications. The repeated distances in the multilayers, or the thickness of two adjacent layers, can be designed to be identical to the interaction lengths characteristic of important physical properties such as electromagnetic or optical interaction lengths. Overall, nanomultilayers (NMs) have been shown to possess many desirable properties, however, research on optical multilayers has been primarily focused on the development of transparent conductive oxides or x-ray optics without exploiting other capabilities of multilayers. Thus, the possible novel combinations of properties, such as transparency and strength, presents an extraordinary area of research with potential use in optical windows, sensor protection, and other applications which require light penetration for function and a robust barrier for protection. Optical multilayers are comprised of nanoscale layers arranged to alter the transmission and reflection of light. This is achieved by tuning the layer thickness and individual layer material properties such as index of refraction and film density. Transparent multilayers are typically comprised of either ceramic/ceramic or metal/ceramic layer combinations. Conventional studies on optical multilayers have focused on multifunctional properties including transmittance, spectral selectivity, or sheet resistivity. Despite the significant interest in implementing these materials in applications that require structural integrity under high loading environments, studies on mechanical properties of optical NMs have been minimal. To better leverage the promise of optical nanomaterials, further studies are necessary to understand the interplay of optics and other multifunctional properties. The key to understanding the augmentation of two or more properties 2 lies in elucidating the microstructural variations as a function of tuning layer thicknesses, aperiodicity, and material selection. The present study focused on the synthesis and characterization of metal/ceramic and ceramic/ceramic NMs. Specifically, the nanoscale features are leveraged to enhance optical properties while investigating the impact of introducing aperiodicity and optically optimized configurations on the microstructural features and mechanical performance, where the term ‘optically optimized’ refers to multilayer configurations with layer thicknesses designed for desired optical properties. The combination of studying the microstructure, optical performance, and mechanical properties is the first of its kind for understanding multifunctionality in optical NMs. To approach this, metal/ceramic AlN/Ag NMs were synthesized with a wide range of transmittance values, showing that optically optimized configurations exhibited significantly increased hardness and optical performance despite similar deformation behavior. This methodology was then applied to exploring underlying microstructural variations in highly transparent ceramic AlN/SiO2 NMs, where distinct changes in grain morphology and interface character were associated with high transparency (%T300-800≈95%) arrangements. Finally, the approach was expanded to include a series of highly transparent ceramic NMs, including AlN/SiO2, AlN/Al2O3, and TiO2/SiO2, which were synthesized in repeated bilayer and optically optimized configurations. The contributions of common optical NM interfaces, including crystalline/amorphous and amorphous/amorphous, were then explored and correlated with microstructural fluctuations and multifunctional film properties. Through these investigations, the data demonstrates that incorporating aperiodicity for improved transmittance greatly influences the microstructural and interfacial characteristics across a range of optical NM systems. Furthermore, these studies present a versatile route for exploring 3 the relationship between optics and multifunctional properties in metal/ceramic and ceramic optical NMs. 4 Chapter 2 : Background 2.1. Nanomultilayers A nanomultilayer (NM) consists of alternating layers of material, where the individual layer thickness of each material is on the order of nanometers. In NMs, the characteristic length scale is the layer thickness rather than the grain size of traditional nanocrystalline materials [1]. Figure 1 shows a schematic of a NM configuration illustrating alternating layers of compositions. NMs present a wide range of material properties and applications, including, but not limited to, desirable mechanical [2-5], optical [6-8], electrical [9-11], and thermal properties [12-14]. The behavior of NMs is typically attributed to layer thickness effects or interfacial character [3, 15, 16]. In NMs, interfaces contribute to a substantial portion of the sample volume as the layer thicknesses approach the order of a few nanometers, thus making the interfacial properties much more important than in their coarse-grained counterparts. Figure 1: Schematic of a multilayer 5 2.1.1. Synthesis of NMs The structures and interfaces of NMs depend on their fabrication methods, which can be grouped into either top-down or bottom-up approaches. Top-down approaches include accumulative roll bonding (ARB), a severe plastic deformation process wherein a bulk material is repeatedly rolled to a severe reduction ratio, deforming it such that the final structure is a NM. Bottom-up approaches include chemical vapor deposition (CVD), electrodeposition, and physical vapor deposition (PVD). Optical multilayers typically consist of either ceramic/ceramic or metal/ceramic multilayers; therefore, ARB is a limited option due to a lack of plasticity in the ceramic layers and electrodeposition chemistry is limited such that many common optical multilayer materials cannot be deposited [17]. NMs are most commonly synthesized by atomic layer deposition (ALD), CVD, Sol-Gel, or PVD techniques. PVD nanoscale multilayer synthesis has focused primarily on three techniques: sputter deposition, vacuum evaporation, and molecular beam epitaxy (MBE), all of which are used for thin film deposition. However, the film formation and deposition kinetics between each technique vary widely. A main benefit of sputtering over vacuum evaporation is that the deposited films have the same composition as the source (target) material and a virtually unlimited range of available deposition materials [18-21]. Furthermore, sputtering allows for switching between the deposition from one source at a time, creating the distinct multilayered morphology. Additionally, sputtering allows for the synthesis of NMs with bilayer thickness as low as 1 nm [22]. For these reasons, sputtering was chosen as the technique for synthesizing the optical multilayers in this work, and further details on the technique and motivation for employing magnetron sputtering are discussed in Sections 2.2.2 and 3.1. 6 2.1.2. Interfaces in NMs The interface boundary that separates two different phases (layers) have different compositions, crystal structure, and lattice parameters. The resulting layer microstructure in a polycrystalline NM is dependent on the interface types, which are characterized as coherent, semi- coherent, or incoherent [23]. A strain-free coherent interface is when both crystals match perfectly at the interface plane, such that the two lattices are continuous across the interface, as shown in Figure 2a. This can only be achieved if the interfacial plane has the same atomic configuration in both phases, regardless of chemical composition. Another way strain-free coherence can be achieved is when the crystals are at a near perfect match by changing the orientation of one layer, shown in Figure 2b. Figure 2: Strain-free coherent interfaces: (a) coherent interface with two crystals of different chemical compositions with the same crystal structure and (b) two phases of different lattices [23]. When the lattices at the interface are not identical, coherency can be maintained by a strain that distorts one or both of the lattices at the interface to compensate for the mismatch between the lattice parameters (Figure 3a). This is called a coherency strain. A semi-coherent interface, Figure 3b, occurs when there is a sufficiently large lattice misfit such that it becomes more energetically 7 favorable for misfit dislocations to accommodate the mismatch at the interface. Both coherent and semi-coherent interfaces can display long-range periodicity. When the lattice mismatch becomes large enough, the misfit dislocation cores overlap, and the interface becomes incoherent, Figure 3c. Incoherent interfaces lack the long-range periodicity of coherent interfaces, and are characterized by high energy (500-1000 mJ/m 2 ). The coherency stresses at coherent and semi-coherent interfaces may repel or attract dislocations. Additionally, the presence of misfit dislocations will affect the dislocation movement across the interface due to the interaction of the dislocation stress fields [3, 24, 25]. Coherent interfaces are strong in shear, while incoherent interface boundaries are weaker in shear [26, 27]. On the other hand, incoherent interfaces that are weak in shear can act as a strong trap for glide dislocations due to core spreading at the interface [16, 28] The effect of the interface structure and strength on NM deformation and mechanical performance will be discussed further in Section 2.4. Figure 3: Overview of different interface types: (a) a coherent interface with a slight mismatch in lattice parameters, leading to coherency strains, (b) a semi-coherent interface where the dislocations relieve the coherency strain and (c) an incoherent interface with fully overlapped dislocations[23]. 8 In addition to crystalline multilayers with varying levels of coherency, there are composite structures with crystalline/amorphous or amorphous/amorphous alternating layers, both of which are illustrated in Figure 4. Amorphous materials lack the long-range order of crystalline materials but do exhibit short- and medium-range order at the atomic scale [29]. In crystalline/amorphous multilayers, the boundaries of the disparate ordered and amorphous layers meet at the interface. In amorphous/amorphous structures, microstructural defects such as dislocations, grain boundaries, and twin boundaries are not possible. The implications of these types of interfaces are further discussed in Section 2.4. Figure 4: Overview of amorphous interface types: (a) a crystalline/amorphous interface and (b) an amorphous/amorphous multilayer, modified from [30]. 2.2. Transparent and Optical NMs Optical NMs are a subset of NMs which are typically composed of metal/ceramic or ceramic/ceramic layers with desirable optical properties. The field of optical coatings has origins in the eighteenth century, with substantial strides being made since the 1930s through the use of vacuum deposition techniques [31]. With the understanding of optical interference in thin films, the field of designing layered media for antireflection coatings, metal-dielectric filters, and high- reflectance mirrors has made substantial progress. Optical coatings are defined as a coating that is intended to modify the optical properties of the surface or substrate that supports it [31]. Optical NMs can be designed to have virtually any 9 reflectance or transmittance characteristics by tuning layer thicknesses and layer material properties such as index of refraction and film density [32]. Optical coatings must be optically smooth, meaning that the size of imperfections or roughness is orders of magnitude less than the wavelengths of light that it is intended to modify. Optically transparent multilayers are promising materials because they offer extraordinary strength, hardness, heat resistance, and most importantly, transparency in both the UV-Vis and NIR wavelengths, a property which bulk silicate glass, a material traditionally used in these applications, typically lacks [33]. Two of the most common types of optical multilayer coatings are antireflection coatings, with applications such as window and lens coatings, and highly reflective mirrored coatings, with applications such as telescope components [34-36], heat reflectors, filters, or for cosmetic purposes. Figure 5 shows a representative SEM image of an optical multilayer sample [37]. Figure 5: SEM cross section of an antireflection coating comprised of SiOx, TiO2 and SiOx-TiO2 nano-multilayers (NML) [37]. 2.2.1. Applications of Optical NMs Typical subsets of optical NMs include antireflection coatings, high-reflection (mirror) coatings, and transparent conductive coatings. Other types of optical multilayer coatings include 10 beamsplitter coatings, filters (shortpass, longpass, or notch), heat-reflecting and heat-transmitting mirrors, and extreme ultraviolet (UV) coatings [38-40]. As previously mentioned, antireflection multilayer coatings have windows and lens coating applications, and include systems such as SiO2/Ta2O5 [41], TiO2/SiO2 [7, 37, 42, 43], and TiO2/VO2/TiO2 [44]. When high-reflection multilayer coatings are employed as telescope components or heat reflectors, some common systems include ZnS/MgF2 [45], HfO2/SiO2 [46] and Mo2C/Si [47], as well as Ag- or Al-based mirrors with protective coatings [48-52]. Optical multilayers for transparent conductive applications, rather than structural applications, have been well-studied in literature with a focus on transmittance and conductive properties. Transparent conducting oxides (TCOs) play important roles in transparent electrode applications such as displays and lighting devices, because they possess both transparency and electrical conductivity simultaneously [53]. TCO systems include TiO2/Ag/TiO2 [54], ITO/Au/ITO [55], ZnS/Ag/ZnS [56], AZO/Ag/AZO [57], ZnO/Ag/ZnO [10], BaSnO3/Ag/BaSnO3 [58], and many more. Most TCO studies focus on optical properties in relation to conductive properties and generally do not include any investigation of mechanical properties. As previously stated, nanomultilayers have been shown to possess many desirable properties such as high strength, corrosion resistance, and radiation resistance [59-61]. To date, research on multilayer optics has been primarily focused on TCOs or x-ray optics without exploiting the many other capabilities of multilayers [10, 55, 62]. The primary objective of designing optical coatings is to achieve desired optical behavior—and traditional applications of optical multilayers include for lasers, astronomy, aerospace, advanced windows, telecommunications, and sensor protection. This vast application space for optical coatings can create demands for additional properties such as mechanical, environmental, thermal, and 11 chemical durability [31], which presents a strong motivation for leveraging the ability to achieve novel combinations of properties in multilayered materials. 2.2.2. Synthesis of Optical NMs The early days of optical NMs focused on synthesis by various thermal evaporation methods [63, 64] but were inundated with issues such as low abrasion resistance and unstable spectral characteristics [31]. Heating of the substrates during evaporation deposition mitigated some of these problems [65] and further experiments on the synthesis of metals and dielectrics for optical properties has examined the relationship between synthesis parameters and resultant structure [66-68]. Since the 1970s, deposition of optical NMs has focused on energetic processes, meaning processes that leverage the addition of kinetic energy to the condensation process [31]. Synthesis techniques for optical NMs involve many of the same techniques for general NMs, with the capability of deposition dielectric materials being the first concern. Optical NM synthesis processes include ion-assisted deposition (IAD) [69], ion plating [70], sputtering [71-73], high power impulse magnetron sputtering (HIPIMS), cathodic arc evaporation [74], pulsed laser deposition [75], chemical vapor deposition [76], atomic layer deposition [77, 78], sol-gel processes [79, 80], etching [81], and other subtypes of thermal evaporation including glancing angle deposition (GLAD) [82, 83]. The microstructure of a material depends on the synthesis parameters, which in turn affect the resultant properties, a relationship which can be leveraged for engineering desired microstructures. For general materials, this is well-illustrated for magnetron sputtering deposition, where the resulting microstructure depends on the “independent” deposition parameters such as the working pressure, target polarization (including bias voltage, current, or power), and the 12 working distance between the target and substrate. These independent variables consequently affect many dependent deposition variables, including the sputtering rate, deposition temperature, and film microstructure [21, 84]. The impact on resulting microstructure is famously illustrated in the Thornton structural zone diagram, shown in Figure 6 [85], which shows four microstructural zones that are dependent on substrate temperature and working pressure. Figure 6: Structure-zone diagram for metal films deposited via magnetron sputtering [85]. Accordingly, the changes in deposition parameters and the dependent microstructural variations have a significant impact on optical performance [86-89]. Changes in crystal structure, density, grain morphology, roughness, and voids due to thin film processing parameters all influence light-matter interactions in optical materials. For example, the work by Chen et al in Figure 7 shows that by varying reactive gas flow (N2, O2) during the deposition of TiNxOy films, the refractive index of the film varies from n500nm = 1.5 to 2.5 [88]. Similar effects on the refractive index can be seen when changing deposition temperature, as illustrated in Figure 8, with the TiO2 refractive index n varying from 2.4 to 2.5 at 550 nm [89]. These variations in optical properties of 13 the thin films, as well as layer thicknesses, significantly impact transmittance of the multilayer stack, and are discussed in further detail in Section 2.3.2. Figure 7: Effect of oxygen flow rate on refractive index n and extinction coefficient k [88] Figure 8: Changes in refractive index n of TiO2 films deposited at ambient temperature (N26), 350°C (N24) and 500°C (N21) [89] Given the wide range of available deposition methods for optical NMs, the selection of a deposition technique is typically driven by the demands of the end application. Sputtering is known to yield high quality, smooth optical films with fully dense microstructures [71, 90, 91], and as 14 such has been selected for this work, with further details on the techniques employed presented in Section 3.1. 2.3. Optical Behavior of NMs The interaction of light with structured materials depends on the feature sizes at different scales, where controlling features on optical behavior include surface structures and film crystallinity. When expanding into multilayer optics comprised of homogeneous thin films, characteristics such the layer thickness and the layer boundaries must be taken into account. In the following sections, the optical behavior of NMs and methods for approximating this behavior are discussed. 2.3.1. Multilayer Optics The propagation of light refers to the manner in which an electromagnetic wave transfers energy from one point to another. When a plane wave interacts with a boundary between two homogenous media with differing optical properties, it is split into two waves: a transmitted wave that continues to the second medium and a reflected wave that propagates back into the first medium. The propagation of electromagnetic radiation through a medium, is described by the complex index of refraction, Ν, where: 𝛮 = 𝑛+𝑖𝑘 (1) The index of refraction, n, describes the refraction of radiation, and the extinction coefficient, k, describes the damping of the radiation. Both optical constants are wavelength dependent. 15 When light travels from one medium to another, the light refracts, and this is described in Snell’s Law, or the law of refraction, which states [92]: sin𝜃 ! sin𝜃 " = 𝑛 " 𝑛 ! (2) This law describes the relationship between the angles of incidence and refraction between two isotropic media. Fresnel’s equations, developed in 1823, describe the reflection and transmission of electromagnetic waves at an interface, and are based on his elastic theory of light. They yield the reflection and transmission coefficient for waves parallel and perpendicular to a plane of incidence. The Fresnel equations are generally written as: 𝑇 ∥ = 2𝑛 " cos𝜃 $ 𝑛 ! cos𝜃 $ +𝑛 " cos𝜃 % 𝐴 ∥ = 2sin𝜃 $ cos𝜃 $ sin(𝜃 $ +𝜃 % )cos(𝜃 $ −𝜃 % ) 𝐴 ∥ (3) 𝑇 & = 2𝑛 " cos𝜃 $ 𝑛 " cos𝜃 $ +𝑛 ! cos𝜃 % 𝐴 & = 2sin𝜃 % cos𝜃 $ sin(𝜃 $ +𝜃 % ) 𝐴 & (4) 𝑅 ∥ = 𝑛 ! cos𝜃 $ −𝑛 " cos𝜃 % 𝑛 ! cos𝜃 $ +𝑛 " cos𝜃 % 𝐴 ∥ = tan(𝜃 $ −𝜃 % ) tan(𝜃 $ +𝜃 % ) 𝐴 ∥ (5) 𝑅 & = 𝑛 " cos𝜃 $ −𝑛 ! cos𝜃 % 𝑛 " cos𝜃 $ +𝑛 ! cos𝜃 % 𝐴 & = − sin(𝜃 $ −𝜃 % ) sin(𝜃 $ +𝜃 % ) 𝐴 & (6) Where T is the transmission coefficient, R is the reflection coefficient, as parallel and perpendicular (polarized) components, 𝜃 $ is the angle of incidence, and 𝜃 % is the angle of transmission. A is the amplitude of the electric vector of the incident field, where A is complex with its phase equal to the constant part of the argument of the wave function. As previously stated, optical multilayers control the reflection and transmission of light, and this is achieved through optical interference. When two beams propagate along the same path, they can either interact constructively or destructively. Constructive interference is when two beams travel along coincident paths with matching phases, and their wave peaks match, thus 16 combining to create a larger total amplitude. Destructive interference is when the two beams are out of phase, and the overlay of the beams results in the combined amplitude decreasing [92]. The performance of an optical multilayer is dependent upon the number of layers, the thickness of each layer, and the index of refraction of each layer. Optical multilayers are designed for a specific incidence angle and polarization of light, such as S-polarization, P-polarization, or random polarization. If the incident angle is altered, the optical paths and angles within each layer will be affected and this will modify the phase change. When a non-normal incident angle is used, S-polarized and P-polarized light will reflect differently at each interface, therefore the optical performance at each polarization will be different. This same phenomenon is leveraged in the design of polarizing beamsplitters. Additionally, there is a polarizing angle, or Brewster angle (𝜃 ' ), such that the angle between the reflected and refracted ray is 90°. The Brewster’s angle is the angle of incidence at which light with a particular polarization is entirely transmitted through a surface with no reflection: tan (𝜃 ' ) = 𝑛 ! 𝑛 " (7) This means that when unpolarized light interacts at this incident angle, the light reflected from the surface is perfectly polarized. A multilayer stack has m total layers deposited upon a substrate, as illustrated in Figure 9. Each layer (m) in the stack is specified by three parameters: the physical thickness (tm), the refractive index (nm), and the extinction coefficient (km). These quantities of each layer, along with the incident refractive index (no), the index of refraction of the substrate (ns), and the extinction coefficient of the substrate (ks) determine the principal optical properties of the multilayer. If the 17 layers are non-absorbing, then the layer boundaries are parallel and the angle of refraction in the m th layer (θm) is determined from Snell’s law [92]. Figure 9: An example of an optical multilayer showing nomenclature used for indicating thickness, layer number (m), refraction index (n), and angle of refraction (α) [93]. By using two layers, one with a high index of refraction and one with a low index of refraction, it is possible to obtain zero reflectance at a single wavelength [94]. The addition of more layers allows for greater flexibility in minimizing reflectance along a wavelength spectrum. A three-layer stack can be designed with zero reflectance at two wavelengths, with less than 1% reflectance across the wavelength spectrum, as illustrated in Figure 10. This same principle can be extended to develop Bragg mirrors by leveraging the layers of high and low index of refraction to induce the desired interference effects. In order to minimize or maximize interference, the layer thicknesses will either be quarter-wave optical thickness (QWOT) or half-wave optical thickness (HWOT) of the wavelength of interest. By continuing to add layers, the high reflectance (or transmittance) performance of the multilayer increases, as illustrated in Figure 11 for a mirrored coating. There is a long history of leveraging interference in multilayer coatings for desired optical properties such as minimizing reflection, creating wavelength filters, band pass filters, and mirrored coatings [95]. 18 Figure 10: Reflectance vs. wavelength in antireflection coatings [94]. Figure 11: Reflectance for high-reflection multilayers [94]. 2.3.2. Multiple Beam Interference (MBI) Recursive Method The multiple-beam-interference (MBI) method, based on the Fresnel equations, provides estimated reflectance and transmittance values at various incident angles for multilayer thin films [37, 92, 96]. This method may be used with an iterative solver to calculate optimal layer thicknesses of each layer within the system for maximum optical properties. The MBI equations are used to calculate the thicknesses of each layer for maximum average transmittance, based on 19 the index of refraction of the layer, the wavelength (λ) range, and the incident angle of light. Changes in each layer thickness lead to changes in the optical path difference for each layer, which then impacts the effective reflectance and transmission. By creating an optimization loop, the calculations are run for a wide range of potential thicknesses for each layer in the system (up to m layers), and the systems meeting the minimum criteria (>95% transmittance over 200-1100 nm) are identified. Figure 12 shows the wavelengths of interest. Figure 12: UV/Vis and NIR wavelength spectrum (in nm) [97]. The following equations build upon the Fresnel equations to calculate the amplitude reflection and transmittance coefficients [92]. The amplitude reflection coefficient of the (m+1) th interface, 𝑟 ()" is calculated by: 𝑟 ()" = 𝑛 ( cos𝜃 ( −𝑛 ()" cos𝜃 ()" 𝑛 ( cos𝜃 ( +𝑛 ()" cos𝜃 ()" (8) where 𝑛 ( is the index of refraction of the m th layer, 𝑛 ()" is the index of refraction of the (m+1) th layer, 𝜃 ( is the incident angle in the m th layer, and 𝜃 ()" is the incident angle in the (m+1) th layer. The amplitude reflection coefficient of the m th interface, 𝑟 ( , is calculated by: 𝑟 ( = 𝑛 (*" cos𝜃 (*" −𝑛 ( cos𝜃 ( 𝑛 (*" cos𝜃 (*" +𝑛 ( cos𝜃 ( (9) 20 where 𝑛 (*" is the index of refraction of the (m-1) th layer and 𝜃 (*" is the incident angle in the (m-1) th layer. The amplitude reflection coefficient of the interface between the air and top layer, 𝑟 " , is calculated by: 𝑟 " = 𝑛 + cos𝜃 + −𝑛 " cos𝜃 " 𝑛 + cos𝜃 + +𝑛 " cos𝜃 " (10) where 𝑛 + is the index of refraction of air and 𝜃 + is the incident angle in air to the top layer. The amplitude transmittance coefficient of the (m+1) th interface, 𝑡 ()" , is calculated by: 𝑡 ()" = 2𝑛 ( cos𝜃 ( 𝑛 ( cos𝜃 ( +𝑛 ()" cos𝜃 ()" (11) The amplitude transmittance coefficient of the m th interface, 𝑡 ( , is calculated by: 𝑡 ( = 2𝑛 (*" cos𝜃 (*" 𝑛 (*" cos𝜃 (*" +𝑛 ( cos𝜃 ( (12) The amplitude transmittance coefficient of the interface between the air and top layer, 𝑡 " , is calculated by: 𝑡 " = 2𝑛 + cos𝜃 + 𝑛 + cos𝜃 + +𝑛 " cos𝜃 " (13) The phase difference due to the optical path difference for each layer m in the film, 𝛿 ( , is: 𝛿 ( = 4𝜋𝑑 ( cos𝜃 ( 𝜆 (14) where 𝑑 ( is the film thickness and 𝜆 is the incident wavelength, both in nanometers. The effective amplitude reflection coefficient, 𝑟 ,-- ,( , of the m th layer is calculated by: 𝑟 ,-- ,( = 𝑟 ( +𝑟 ()" 𝑒 $/ ! 1+𝑟 ( 𝑟 ()" 𝑒 $/ ! (15) The effective amplitude reflection coefficient, 𝑟 ,-- ,(*" , of the (m-1) th layer is: 𝑟 ,-- ,(*" = 𝑟 (*" +𝑟 ,-- ,( 𝑒 $/ !"# 1+𝑟 (*" 𝑟 ,-- ,( 𝑒 $/ !"# (16) 21 The effective amplitude reflection coefficient, 𝑟 ,-- ," , of the top layer is: 𝑟 ,-- ," = 𝑟 " +𝑟 ,-- ,! 𝑒 $/ # 1+𝑟 " 𝑟 ,-- ,! 𝑒 $/ # (17) The effective amplitude transmittance coefficient, 𝑡 ,-- ,( , of the m th layer is calculated by: 𝑡 , --,( = 𝑡 ( 𝑡 ()" 𝑒 $/ ! 1+𝑟 ( 𝑟 ()" 𝑒 !$/ ! (18) The effective amplitude transmittance coefficient, 𝑡 ,-- ,(*" , of the (m-1) th layer is: 𝑡 ,-- ,(*" = 𝑡 (*" 𝑡 ( 𝑒 $/ !"# 1+𝑟 (*" 𝑟 ( 𝑒 !$/ !"# (19) The effective amplitude transmittance coefficient, 𝑡 ,-- ," , of the top layer is calculated by: 𝑡 ,-- ," = 𝑡 " 𝑡 ! 𝑒 $/ # 1+𝑟 " 𝑟 ! 𝑒 !$/ $ (20) These effective amplitude reflection and transmission coefficients can then be used to calculate the total reflectance and transmittance of the entire m-layer system: 𝑇 = 𝑛 0 cos𝜃 0 𝑛 + cos𝜃 + 𝑡 ,-- ," 𝑡 ,-- ," ∗ (21) 𝑅 = 𝑟 ,-- ," 𝑟 ,-- ," ∗ (22) The total transmittance (T) and reflectance (R) values are calculated for each λ value in the 𝛿 ( (Equation 14) calculation. The process of estimating the total transmittance or reflectance for a given multilayer stack is summarized in Figure 14. Studies by Lu et al using MBI calculations have shown good agreement between the predicted transmittance calculations and the experimental results [33, 96, 98, 99]. An example is shown for a four-layer SiOx/TiO2 ACR multilayer in Figure 13a illustrating the correlation of theoretical and experimental measurements, where the absorption of visible light due to high sputtering speeds was mitigated using a heat treatment at 600⁰C for 10 min [37]. Additionally, the 22 average transmittance of the SiOx/TiO2 multilayers is affected when the layers deviate from the optimal calculated thickness, as summarized in Figure 13b. Figure 13: a) Measurement and theoretical MBI calculations in SiOx/TiO2 NMs with inset showing the sample structure and b) process window simulation showing the defect of thickness deviation on overall transmittance for a four layer ACR coating, where ideal thicknesses are SiOx (84 nm)/TiO2 (110 nm)/SiOx (35 nm)/TiO2 (11 nm) corresponding to an optical transmittance of 99.34% [37]. It is important to note that the MBI calculations are based on ideal systems and basic optics calculations, and do not account for factors such as surface roughness, interface roughness, or potential processing side effects (e.g. film density, residual stresses, contaminants), and thus only serve as a guide for expected transmittance and reflectance values. These optics calculations can be used to filter out multilayer systems unable to reach high transmittance values. Other effects, such as a grain size, could play a large role in optical properties and have not been examined in literature for their effect on transmittance, and necessitate further study. A MATLAB optimization code based upon the MBI recursive method was developed for this work, and is available in Appendix E. 23 Figure 14: Schematic of the MBI recursive method for a multilayer system of m layers, with the calculation starting from the a) m th layer reflection amplitude coefficient to produce the b) effective reflection amplitude coefficient of the m th layer, reff,m, the process continues until the c) effective reflection amplitude coefficient of the 1 st layer is calculated, then d) the total reflection and transmission of the m-layer multilayer system is calculated. 2.4. Mechanical Behavior of NMs In general, studies on the mechanical behavior of NMs have focused on metallic systems [27, 61, 100-111] and metal/ceramic systems [5, 112-121] with some studies on ceramic multilayers [2, 122-128] and even fewer studies on optical multilayers [42, 129, 130]. For ceramic/ceramic and metal/ceramic NMs there are several explanations for their behavior, however, none are universally applicable. Thus, there is a great opportunity to bridge this knowledge gap by performing a comprehensive study which could lead to a broadly applicable understanding of non-metallic NMs. The following overview of mechanical behavior of various NM systems highlights the contrast between the current understanding of metallic and non- metallic NMs, further emphasizing the need for future studies on non-metallic NMs. 24 2.4.1. Metal/Metal NMs Metal/metal NMs have been investigated extensively, and studies have found that these metallic NMs can achieve extraordinary strength and ductility, beyond the rule-of-mixtures, when the layer thickness is on the order of a few nanometers [16, 131, 132]. The overview for strengthening behavior of metallic NMs is illustrated in Figure 15 (note that for the X-axis smaller layer thicknesses are on the right) [3]. Overall, metallic NMs show a Hall-Petch type strengthening behavior and the strengthening of metallic NMs is attributed to changes in the dislocation mechanisms with decreasing layer thickness. This trend holds true for a variety of material systems, which indicates that this behavior is a result of the layered structure, not just a particular material system. In addition to having high strengths, metallic NMs are also strong in fatigue loading [133, 134], holding promise for use as protective coatings and in structural applications. Some recent work has shown that multilayers can exhibit varying deformation behaviors under scratch testing based on the stability and coherency of the interface, such as vortex formation or mixed nanostructures [135]. This type of deformation would affect the optical properties of NMs and has yet to be explored. Figure 15: Hardness values for metallic NMs systems as a function of the inverse square of the layer thickness [61]. 25 Hall-Petch type strengthening has been observed in both equiaxed and multilayered materials, and is summarized as: 𝜎 2 = 𝜎 + +𝑘 2 𝐷 * " ! (23) Where 𝜎 2 is the yield strength, 𝜎 + is the single crystal yield strength, 𝑘 2 is the Hall-Petch constant, and D is the average grain diameter. In Hall-Petch strengthening of equiaxed materials, the dislocations in a material begin to pile-up, and this pile-up of multiple dislocations increases the stress on the leading dislocation. This increased stress forces the leading dislocation to interact with the grain boundary, either by transmission, absorption, or nucleation of other dislocations across the grain boundary [136]. As the grain size decreases, fewer dislocations can pile up in a single grain, which means that there will be less stress acting on the lead dislocation, and thus a larger applied stress is required to cause the interaction between the grain boundary and the dislocation. Figure 16 shows an overview of Hall-Petch strengthening, where effect of changing grain size on the yield strength in a nanocrystalline material exhibits three distinct regimes. The first regime is the Hall-Petch strengthening regime, for nanocrystalline grain sizes between 1 μm and 100 nm. Between 100 nm and 10 nm, there is a change in the slope, corresponding to the change in deformation mechanisms. Below 10 nm, there is some debate about the slope observed in these ultrafine grain materials, with some studies reporting a negative Hall-Petch slope [137]. 26 Figure 16: Yield strength as a function of the inverse square root of the grain size, for nanocrystalline materials. There are three regimes for flow stress as a function of the grain size, in the > 100 nm regime, there is Hall-Petch strengthening, with a slope of k. At grain sizes of 100 nm and 10 nm the slope changes, which is indicative of the dislocation mechanism changing. In metallic NMs, similar behavior has been observed, and an overview is shown in Figure 17 (note that for the X-axis smaller layer thicknesses are now on the left). As previously mentioned, in NMs the layer thickness is the characteristic length scale (D, in this case) rather than the grain size [136]. In the sub-micron to micron layer thickness regime, Hall-Petch strengthening is observed in metallic NMs, where the deformation is assisted by the effect of dislocation pile-ups. The deformation mechanisms for metallic NMs can be directly correlated with changes in the dislocation mechanisms; for example, when the layer thicknesses decrease below 50-100 nm, the layer thicknesses are too small to allow for dislocation pile-up and instead, a confined-layer slip (CLS) mechanism is observed, which operates by the bowing of dislocations within a single layer [138, 139]. With a further decrease in layer thickness below about 5 nm, there is another change in the deformation mechanism as the decrease in layer thickness forces the bowing 27 dislocations to assume very small bowing radii, requiring higher stresses. With these increasing stresses required to bow the dislocations, the mechanism switches from CLS to inter-boundary slip (IBS). In this regime, the necessary stress for a dislocation to cross a boundary is lower than that required for further dislocation bowing. Therefore, dislocations are no longer confined to one layer, and the strength of the samples is generally observed to either stay the same with decreasing layer thickness, or even decrease. Effects such as dislocation core spreading across the interface, interface stresses, and interface dislocation array behavior also affect the behavior [3]. Figure 17: Illustration of the dislocation mechanisms of metallic NM strength contributions operating at different length scales under an applied force normal to the layer interfaces [3]. An additional area of study in metallic NMs explores the contribution of crystalline/amorphous and amorphous/amorphous interfaces. These studies aim to leverage the beneficial properties of amorphous metallic glasses, which are known to have exceptionally high yield strength and wear resistance, while mitigating the drawbacks of metallic glasses including 28 poor ductility and catastrophic failure due to rapid shear band propagation [30]. Crystalline/amorphous NM studies have reported co-deformation without delamination, rupture, or cracking, up to plastic strains as large as 85% [103]. Both size-dependent strengthening and hardness plateaus have been reported in crystalline/amorphous systems. Previous studies have shown that when the thickness of amorphous layers is sufficiently thin, the crystalline layers can accommodate plasticity and constrain the formation and propagation of shear bands in the amorphous layer [140, 141]. Additional studies have observed transitions in plastic deformation modes—from shear banding to co-deformation— as a function of layer thickness reduction [29, 142, 143]. Load sharing between amorphous and crystalline layers, as well as inhibition of strain localization within individual layers, leads to higher failure strains and homogeneous deformation until catastrophic propagation of one of the shear bands, an example of which is shown Figure 18 [103]. Interfacial features such as medium range order (MRO) at the amorphous interface in Cu/CuNb has also been reported to impact the deformation mechanisms, illustrated in Figure 19, by lowering the resistance to shear transformation zones (STZs) in the amorphous layers [30]. Figure 18: Deformation morphology of a crystalline/amorphous Cu40/PdSi10 micro pillar [103]. 29 Figure 19: A schematic illustrating the deformation mechanism proposed in crystalline/amorphous Cu/CuNb multilayers [30] Introducing amorphous/amorphous interfaces in NMs has been explored as an additional approach to combat the drawbacks of amorphous metallic glasses, which expands upon the method of fighting the inferior fracture properties of brittle materials by deliberate introduction of a high- density of defects such as heterogeneous interfaces [144]. As discussed in Section 2.1.2, the lack of microstructural defects in amorphous materials results in the inability to restrict shear band motion, culminating in strain softening and catastrophic failure in monolithic amorphous materials [145, 146]. By introducing heterogeneous, sharp interfaces, they can act as a barrier to microcrack propagation and obstacles to the nucleation and propagation of shear bands. Transitions from inhomogeneous deformation modes to a more homogeneous co-deformation mode has been observed by stacking amorphous metallic layers with differing elastic moduli [144]. Amorphous/amorphous interface obstruction to shear bands has been observed, but there exists a critical layer thickness for the localized to homogeneous deformation transition. These competing deformation mechanisms are shown in Figure 20, where below this critical value multiple shear bands transmit through the interface, and above this the shear bands localize in the thicker layers [147]. 30 Figure 20: Schematic illustrating the competing governing mechanisms, localized shearing and shear band interactions, in amorphous/amorphous NMs [147]. 2.4.2. Metal/Ceramic NMs In metal/ceramic NMs, constituent materials can have large differences in mechanical strength and elastic modulus, and thus have been shown to exhibit a combination of high strength and toughness [148]. For example, the stress required to move a dislocation (e.g. Peierls stress) is generally much higher in a ceramic than in a metal [149]. The vastly different properties in the metal and ceramic layers show promise when combined into a multilayer system to achieve new deformation mechanisms. Most metal/ceramic NM studies on mechanical behavior involve nanoindentation testing [118, 150-153], with some studies incorporating tensile [118] and micropillar compression testing [116, 154]. The deformation mechanisms of metal/ceramic NMs are commonly characterized by cracks first forming in the brittle layer, then growing to neighboring ductile layers [113, 155]. A representative study by Chawla et al showed that in Al/SiC NMs, with decreasing layer thickness (to 25 nm), the hardness increased due to the greater plastic constraint of the metallic layers [155]. In this study, shown in Figure 21, the NMs exhibited significant pile-up due to ceramic layer 31 bending and plasticity in the metallic layers. The damage to the multilayer was primarily cracking of the ceramic layer with plasticity, void nucleation and growth in the metallic layers. The bilayer thickness and thickness ratio between ceramic and metal layers strongly affect the hardness and ductility of these multilayered materials [114]. In the thicker layer regime (> 10 nm thickness), there are two primary failure modes: failure via cracks caused by a combination of elastic-plastic deformation in the metallic layer and elastic deformation in the ceramic layer, or failure via localized shear deformation due to localized interface stresses [5, 115-118]. For example, Figure 22a and Figure 22b show plastic deformation in the metal layers of a TiN/Al multilayer, with cracking in the TiN layer layers that fills in with the ductile Al layer [156]. Figure Figure 21: (a) FIB cross-section of an indentation of a Al/SiC nanomultilayer, showing shear bands and SiC cracking, (b) higher magnification of SiC flexibility and microcracking, and (c) delamination of the multilayers/Si interface and fracture within the Si substrate. 32 22c shows NbN/W multilayers with plastic deformation in W layers with twin-like features in the NbN layers [5]. With nanoscale layer thicknesses (< 10 nm), experimental results have shown the potential for ceramics to plastically codeform with metal layers when the layer thickness reduces to a few nanometers [119]. An example of the effect of bilayer period on hardness is shown in Figure 23 for TIN/W multilayers (note that the smaller layer thicknesses are on the left) [112]. Studies have found that peaks in hardness properties of these metal/ceramic materials depend on bilayer period of the layers. It has also been shown that hardness decreases with increasing metal/ceramic layer thickness ratios. Figure 22: (a) SEM cross-section showing the deformation mechanisms of Al and TiN multilayers, where the multilayer has been indented with a Berkovich tip. (b) Higher magnification shows that Al layers carry out plastic deformation by dislocation glide, while the TiN layers have cracking. These cracks are filled by the Al layer [157]. (c) Transmission electron microscopy (TEM) bright field cross-section showing nanoindentation W/NbN multilayers showing plastic deformation in the W layers with twin-like features (straight lines) in the NbN layers [5]. 33 Figure 23: The effect of bilayer period on hardness for TiN/W multilayers [112]. The strength of metal/ceramic NMs has been shown to have a strong dependence on the combination of constituent materials, while the improvement of ductility and fracture toughness is observed in most systems [151]. While some studies observed improvements in hardness as a function of decreasing layer thickness [112, 116, 156, 157], others have seen no improvement in hardness [5, 117, 158]. A study by Abadias et al linked the lack of improvement in hardness in ZrN/W NMs to easier dislocation glide from the metallic layer to the ceramic layer due to similar shear moduli between ZrN and W, as well as enhanced delamination due to compressive stresses occurring during nanoindentation [117]. However, this theory did not extend to other systems, where another study by Abadias et al on TiN/W multilayers showed that despite similar shear moduli between TiN and W, the hardness increased with decreasing layer thickness [112]. This increase was attributed to the sharp interface barriers observed between the metallic and ceramic layers, but despite the many studies on metal/ceramic NM deformation, there is no overarching explanation for transitions in behavior. The deformation mechanisms for metal/ceramic NMs also correlate with changes in the dislocation mechanisms. In metal/ceramic NMs, dislocations are typically generated in the metal 34 layer and transmitted to the ceramic layer [114]. Metal/ceramic NMs also exhibit the confined layer slip deformation mechanisms of metallic NMs, where strain hardening occurs when dislocations glide and interact in a single layer, therefore higher stresses are needed to transmit dislocations into another layer. This gives rise to ductility in the NMs. Material selection and composition, bilayer thickness, and layer thickness ratio all affect strength and ductility of the NM systems by affecting the propagation and movement of dislocations during deformation [119]. Interfaces play a strong role in plastic deformation of metal/ceramic NMs, acting as sources, sinks, and barriers for defects as well as sites for nucleation and gliding of lattice dislocations [120, 121, 159]. The interface structure is determined by the crystallographic structure of the metal and ceramic layers. Like metallic NMs, metal/ceramic NMs that have incoherent interface structures hinder the transmission of dislocations, while weaker shear resistance characteristic of coherent (or semicoherent) interfaces enables core spreading of dislocations at the interface. During plastic deformation, dislocations originating in the metal layers may migrate to the interface, thereby increasing local stresses at the interface and facilitating the nucleation of lattice dislocations in the neighboring layer. 2.4.3. Ceramic/Ceramic NMs For monolithic ceramic thin films, dislocations are not as numerous and do not move as easily as they do in metals [160]. Though there are fewer dislocations, they still play an important role in nanostructured ceramics, including in crystalline ceramic thin films on crystalline substrates [160]. In ceramics, dislocations typically have complex and large unit cells, and the additional complexities of both charge and directional bonding must be considered [161]. The structure of the dislocation core in ceramics depends on the charge of ions, the size of ions, and the presence 35 of directional bonds. The interaction of defects with dislocations in ceramics differs from metals, primarily due to the dislocation core being a more open region of the ceramic crystal structure, resulting in a lower number of bonds. This means that defects such as vacancies and interstitials will more easily interact with the dislocation core. Dislocations do not glide easily in ceramics; climb dissociation is more common, but only when temperatures are very high [162]. During nanoindentation, deformation at the crack tip can lead to the generation of dislocations in the plastically deformed region, thus blunting the crack tip and toughening the ceramic [163]. Additionally, these dislocations formed under indentation can act as sites for local amorphization of the ceramic. In scratch testing, dislocations can be emitted from the region below the indenter tip and will tend to move out on the glide planes [160]. For ceramic NMs, the research is somewhat limited compared to metallic NMs and overall trends and mechanisms have not been well determined. However, some key observations can be made; for example, there seems to be a critical bilayer thickness for maximized mechanical properties (both hardness and wear resistance), as shown in Figure 24, but it is not necessarily the smallest bilayer period [122, 164]. In the CN/BN nanoindentation studies shown in Figure 24, the fracture begins in the weaker ceramic layer and decreasing layer thickness can lead to increased hardness and wear resistance due to the increase in interfacial content. In a TiN/AlN study by Boutos et al, it was found that decreasing the bilayer thickness resulted in an increased nanoindentation hardness, which was attributed to the increased volume fraction of interfaces, and a change in texture and average grain size of the layers was also observed [164]. A Hall-Petch- type mechanical relationship in TiC/TiB2 ceramic multilayers has been reported [165], but is not a generalized trend in ceramic NMs. In crystalline ceramic/ceramic NMs, high hardness can be 36 achieved through sharp interfaces and periodicity in the 5-10 nm range and are referred to as superlattice coatings [166-168]. Figure 24: The effect of bilayer period on nanoindentation hardness and dissipation modulus on CN/BN multilayer films [122]. Proposed explanations for hardness enhancement in ceramic NMs include shear modulus mismatch, lattice mismatch, and Hall-Petch type strengthening [169]. It has been suggested that a difference in the modulus between the two ceramic layers is required to increase the hardness, and that the coherency strain between layers has a minimal effect [170]. Other studies have shown reorientation and relocation of columnar grains and sliding along columnar grain boundaries [124] as well as intergranular shear sliding with crack formation due to lack of deformability [125]. A study by Chu and Barnett [171] proposed a model for yield stress and hardness enhancements in superlattice thin films, where the stress required for dislocations to glide across layers with varying shear moduli was calculated using core effects and interfacial effects. They found that the predicted strength and hardness enhancements increased based on the layer period, until reaching a saturation point that depended on the interface widths. This saturation point is illustrated in Figure 37 25 for TiN/NbN multilayers, where a peak hardness occurred at bilayer thicknesses of 5-10 nm, depending on the interface width. Figure 25: Experimental hardness vs. superlattice (bilayer) period compared with model predictions for dislocation glide across and within layers in single and polycrystalline TiN/NbN multilayers [171]. A key parameter in mechanical behavior of ceramic NMs is the interface strength, where the high density of interfaces provides barriers to limit crack propagation [126] and weak interfaces can lead to delamination [2]. Inserting weak interfaces within rigid ceramic layers can change the fracture behavior by changing the path of crack propagation, resulting in cracks becoming non- catastrophic [172]. Studies of ceramic NMs with strong interfaces have been shown to exhibit significant improvement in properties such as work of fracture and fracture toughness [127, 128, 173-175]. Toughening of ceramic materials through multilayer architectures has been a lively 38 field, where crack energy and deflection has been achieved through mechanisms illustrated in Figure 26 including: crack splitting at grain boundaries, crack deflection at layer interfaces, reduction of local stress concentrations via interface opening, and nanoplasticity in the form of plastic deformation at the interface [166, 176-178]. Figure 26: Mechanisms of toughness enhancement in hard ceramic/ceramic multilayers [176]. It is important to note that though there are studies on ceramic/ceramic NMs with repeated bilayer thicknesses, there are far fewer for optical ceramic/ceramic NMs with non-repeated bilayer thickness. The limited studies available typically report on a single coating system or layer configuration but have not been expanded to general trends in the mechanical behavior of optical coatings. For example, a study by Mazur et al reported an antireflection TiO2/SiO2 coating with a nanoindentation film hardness of 9.4 GPa and the corresponding 2D depth-of-penetration profile, but only for a single multilayer film. Hardness and wear behavior reported in other optical NM studies [33, 98, 179] are similar in scope. Thus, general trends of opto-mechanical behavior of optical NMs as a function of non-repeated bilayer thicknesses, microstructure, and interface characteristics are undetermined. 39 Chapter 3 : Experimental Methods and Materials The following section includes an overview of the synthesis of NMs and the deposition configurations for this study. In addition, characterization methods for the multilayer systems are discussed, including microstructural, optical, and mechanical characterization. 3.1. Magnetron Sputtering PVD, more specifically magnetron sputtering, was chosen for the synthesis of optical multilayers in this work due to the flexibility and control highlighted in Section 2.1.1. Since optical NMs are ceramic/ceramic and metal/ceramic systems, three types of sputtering techniques will be used: reactive sputtering (DC and RF) and traditional RF sputtering. By using these three combinations, a wide range of compositions and microstructures can be synthesized. Though there are many variants of sputtering (DC, RF, reactive, magnetron), as well as hybrids (reactive magnetron), all sputtering processes involve planar diodes with facing anode and cathode electrodes, and the deposited thin film originates from the target cathodes which play a role in the plasma [21]. “Sputtering” refers to the act of particles of a solid material ejecting from a source material after bombardment of energic particles, or incident ions. Sputter deposition typically uses an argon plasma because the noble gas will not react with the target material. 3.1.1. RF/DC Magnetron Sputtering A representative multilayer sputtering configuration is shown in Figure 27. The microstructure produced by sputtering depends upon the deposition parameters, such as the working pressure, control of target polarization (whether by voltage, current, or power), and the working distance between the target and the substrate [85, 180, 181]. These “independent” 40 variables determine dependent parameters such as the deposition temperature, deposition rate, and resulting microstructure. An additional variable is the type of sputtering power supplied to the target, which can be either direct current (DC) or radio frequency (RF). In magnetron sputtering, the magnetron behind the target causes the electrons to circulate on a closed path near the target surface, thus creating a high density of plasma from which ions can be extracted to sputter the target material [182]. In DC sputtering, the cathode electrode is the sputtering target with an applied direct current and the substrate is placed on the anode, typically at ground potential [183]. As a result of the constant potential on the cathode, DC magnetron sputtering allows for a higher deposition rate at lower applied potentials, and relatively low working gas pressure (< 5 mTorr), and also yields better coating uniformity [84, 184, 185]. However, with DC sputtering, the caveat is that an electrically conductive (metallic) target must be used. In RF sputtering, a large peak-to-peak voltage is used to generate an alternating positive/negative potential on the cathode surface. RF sputter deposition typically uses frequencies from 0.5-30 MHz, with 13.56 MHz being the most common. The advantages of RF sputtering include that lower gas pressures (< 1 mTorr) can be used, as well as the flexibility that the target can either be electrically conductive (metallic) or insulating. The tradeoff for increased target material versatility, especially in the case of insulating materials is that RF sputtering typically yields lower sputtering rates. During RF sputtering of dielectric targets, care must be taken to avoid thermal shock since most insulating materials have poor thermal conductivity, high coefficients of thermal expansion (CTEs), and are typically brittle. Thus, a proper ramp-up and ramp-down procedure must be used when introducing a voltage to the target, and the maximum power density for the target must not be exceeded. 41 Figure 27: Schematic showing the setup and process of nonreactive magnetron sputtering. A neutral gas (Ar) is introduced and ionized by the negatively biased target, the Ar ions then strike the target and cause the target atoms to be ejected and coat the substrate. The geometry of the sputtering configuration is shown in Figure 28, where two sources are focused on the sample substrate, and the power supplied to the sources (DC or RF) and pneumatic shutter is turned on and off intermittently to deposit the multilayers. Figure 28: Schematic of sputtering setup for depositing multilayer systems. The multilayer thickness is determined by the sputtering rate and ‘on time’ of each source, as controlled by the pneumatic shutter, (a) sputtering from source 1 while the shutter blocks source 2, and (b) sputtering from source 2 while the shutter blocks source 1. 3.1.2. Reactive/Non-Reactive Magnetron Sputtering 42 Furthermore, the type of working gas that flows into the sputtering chamber during deposition can be controlled, where films can be deposited under either reactive or non-reactive gas flow. In the schematic of the non-reactive sputtering process shown in Figure 27, only an inert working gas (Ar) is used during deposition. With non-reactive sputtering, the film deposited is the same composition as the source target, for example, sputtering a conductive Ti target with DC power under an Ar working gas flow results in a Ti film. During reactive sputtering, a film is deposited in the presence of a reactive gas (O2, N2) mixed with an inert working gas (Ar) and requires reaction between the target materials and reactive gas species [186]. For example, a Ti target may be sputtered, either with an applied DC or RF power, in the presence of O2 and Ar to deposit a TiOx film. A heavy inert gas, such as Ar, is required to facilitate effective sputtering, due to the fact that the typical reactive gases, O2 and N2, have low atomic masses which effectively limits their ability to sputter a material since sputtering is a momentum-driven process. Reactive sputtering can alleviate some of the limitations of using RF power and dielectric targets mentioned in Section 3.1.1, including the thermal shock and discharging constraints. Additionally, deposition of a compound via high purity reactive working gas and a metallic target can yield higher purity films as compared to using a sintered dielectric target [31]. Although non- reactive RF sputtering can be used to sputter oxides and nitrides, it is restricted to low sputtering rates due to the low maximum power densities that may be used and the low sputtering yields of insulating compounds [187]. Reactive magnetron sputtering is a flexible technique in the deposition of common materials in optical coatings, particularly oxides and nitrides, with the exception of flourides. In reactive DC sputtering, a common concern is to prevent the “poisoning”, or formation of a compound layer on the surface, of the target. When poisoning occurs, there is a drastic decrease in sputtering rate and efficiency. Since the sputtering rate decreases, there is less 43 consumption of the reactive gas and the partial pressure jumps rapidly from point A to point B as shown in Figure 29 and thus must be avoided by controlling the flow of the reactive gas. A sufficient flow must be supplied such that there is enough reactive gas to deposit the desired compound, but not so much that poisoning of the surface occurs [84]. Gas mixtures during reactive deposition are controlled by using mass flow controllers (MFCs) to achieve the desired composition. Overall, reactive sputtering can be used with either DC or RF power to deposit common dielectric materials, thereby offering an expanded range of materials for deposition. Figure 29: Hysteresis plot for the reactive sputtering of Ti in an Ar/O 2 atmosphere with flow control of the reactive gas [188]. 3.1.3. Residual Stresses and Profilometry In thin films deposited by PVD, there are two types of residual stresses: intrinsic and extrinsic. Intrinsic stresses arise from defects such as dislocations in the film. Extrinsic stresses arise mainly from adhesion to the substrate. Stress can be introduced during deposition due to a difference in thermal expansion between the deposited film and substrate, lattice mismatch with 44 the substrate, or a chemical reaction with the substrate. Residual stresses in the films can be calculated using the radius of curvature as measured by profilometry techniques. There are two main types of profilometry: optical profilometry and stylus profilometry. Optical profilometry uses light instead of a physical probe to detect the 3D features of the surface. Stylus profilometry is a common tool to measure surface film characteristics and is based on contact measurement of the sample. In this technique, a 20 nm diameter diamond stylus tip moves in the lateral direction over the sample and measures the vertical displacement of the stylus, an example of which is shown in Figure 30 [189]. In these profiles, vertical features ranging from 10 nm to 1 mm are detected. Figure 30: Schematic of a stylus profilometer [190]. Using these profiles, the radius of curvature of the substrate before and after deposition of a film can be estimated, as shown in Figure 31. This value is then used in Stoney’s equation to determine the residual stresses in the films [191, 192]: 45 𝜎 3%45,2 = −𝐸 0 𝑡 0 ! 6𝑡 6 (1−𝜐 0 )𝑅 (18) Where the radius of curvature, R, is: 𝑅 = 𝐿 ! 8𝛿 (78 (19) Where L is the length of the coating, δmax is the maximum deflection, ts is the substrate thickness, tc is the coating thickness, Es is the elastic modulus, and υs is the Poisson’s ratio. Stoney’s equation is valid up to a thickness ratio, tc/ts, of less than 0.05, which is the case for the 1 μm thick films deposited on 500 μm thick glass substrates [193]. Figure 31: Substrate profile showing radius of curvature before and after film deposition, measured using profilometry. 46 3.2. Microstructural Characterization Methods Due to the nanoscale features of NMs, including layer thicknesses, grain size, and interface character, multiple characterization techniques are leveraged to investigate the features of the NMs from the micro- to nanoscale. These various techniques are discussed in the following section. 3.2.1. X-Ray Diffraction X-ray diffraction (XRD) is a nondestructive characterization technique which uses x-rays to identify crystallographic orientations (texture), material compositions, grain size, and sample strain [194]. XRD peaks originate from x-ray scattering intensity upon interaction with the lattice planes in a sample. The XRD spectra produced are the scattered intensity as a function of 2θ. The direction of the scattered x-rays are governed by the wavelength of the incident radiation and the crystalline nature of the sample, and the relationship between the wavelength of the x-rays to the spacing of the atomic planes is formulated by Bragg’s law: 𝑛𝜆 = 2𝑑sin𝜃 (20) Where n is an integer, λ is the characteristic incident wavelengths, d is the interplanar spacing between the atoms, and θ is the angle of the x-ray with respect to these planes, as seen in Figure 32. When this relationship is satisfied, the x-rays scattered by the atoms are in phase, and the diffraction that occurs is defined by angle θ. A diffraction pattern, such as that shown in Figure 33, can be used be used as a “chemical fingerprint” to identify chemicals in the sample by comparing to a database of known diffraction patterns as well as to identify textures of the films. ‘Texture’ throughout this manuscript refers to the statistics of the crystallographic orientations in a polycrystalline material, meaning the randomness of orientations or presence of preferential orientations. 47 Figure 32: The incident beam (left) interacts with the atoms of spacing d and scatters. Figure 33: Normalized XRD patterns of sputtered Al, Al-Mg, and Ni films [195]. 3.2.2. Scanning Electron Microscopy In scanning electron microscopy (SEM), a beam of electrons is produced using a field emission gun, and is then condensed by using multiple electromagnetic lenses and apertures to focus the beam [196]. The focused electron beam scans over a surface and interacts with the surface of the sample to generate secondary electrons, backscattered electrons and characteristic x-rays, an overview of which is shown in Figure 34. This interaction produces images and the 48 various signals can be used to obtain information about the surface topology, crystal orientation, and elemental composition of the specimen. When the interaction is inelastic, the electrons collide with the top atomic layers causing emission of secondary electrons and x-ray. When the interaction is elastic, the energy of the electrons is almost totally conserved, the emitted electrons are back- scattered electrons, where the colliding electrons are reflected at high angles. Figure 34: (a) Electron and photon signals emanating from the tear-shaped interaction volume of the specimen, (b) energy spectrum of electrons emitted from the specimen surface and (c) the effect of the surface topography of the sample on electron emission [197]. Secondary electrons are generated from the top layers of the specimen, and are used to form images of the surface, and example of which is shown in Figure 35. These electrons are low energy, which means that they are originating from a subsurface depth of less than a few angstroms. The contrast in the image formed depends on the depth of the sample, where the closer features will appear brighter, as is the case with the edges of a sample. 49 Figure 35: SEM image of a Vickers indent on a ceramic nanomultilayer sample, synthesized by C. Appleget. Backscattered electrons are high energy electrons which are elastically scattered from the sample and possess nearly the same energy as the incident electrons. These electrons can be used in techniques such as electron backscattered diffraction (EBSD), also known as backscatter Kikuchi diffraction (BKD) to gather information about the crystallographic texture of a polycrystalline sample. X-rays of varying energies are emitted when the electrons interact inelastically with the sample surface. These x-rays are characteristic of the atoms from which they are emitted, and their associated energies can be identified and counted to determine the concentration of atomic species in the specimen. This is the basis for the energy dispersive x-ray spectroscopy (EDS) technique commonly used to identify the chemical composition of a sample. This work used cross-section SEM micrographs to measure film thickness and layer thickness, as well as EDS scans to measure the global composition of the samples. 50 3.2.3. Focused Ion Beam Focused ion beam (FIB) is a technique similar to SEM, but rasters a beam of ions over the sample rather than an electron beam [198]. Secondary electrons are generated by the interaction of the beam with the sample to produce images of the sample surface. In most FIB instrumentation, a dual beam system is used with both SEM and FIB, an example of which is shown in Figure 36. Most commonly, Ga + ions are used, and the FIB beam can be used to mill the specimen surface via sputtering at a specified current, resulting in excavations with nanometer precision. FIB techniques can be used for cross-section imaging of a sample, shown in Figure 37, or for TEM sample preparation, shown in Figure 38. For TEM sample preparation via FIB lift-out, the FIB and SEM beams are combined with micromanipulators and gas injection systems (GIS) for deposition of carbon or platinum to weld the sample. In the FIB lift-out technique, a cross-sectional lamella of the sample is prepared and lifted from the sample, welded to a TEM grid using a gas injection system (GIS), and thinned using FIB milling until electron transparency is achieved. In addition to conventional FIB systems using Ga + ions, FIB using a Xe + ion plasma source has been shown to have advantages in certain applications. Xe + plasma FIB (PFIB) techniques have a much higher maximum current which allows for high throughput milling of large sample volumes and also eliminates Ga + contamination and implantation. The depth of amorphization after TEM sample preparation, highlighted in Figure 39, has also been reported to be lower using PFIB preparation [199]. 51 Figure 36: Schematic of a dual beam FIB and SEM instrument with the inset showing the electron and ion beam interaction with the sample [198]. Figure 37: Overview of FIB cross-section showing (a) top view of Vickers indentation sample with dotted line indicating the milling region and (b) cross-section of the indent produced via FIB milling (synthesized by C. Appleget). 52 Figure 38: Overview of FIB lift out technique beginning with a) carbon protective layer on top of sample and trenches milled around region of interest, b) preparation of peninsula with bright layers of multilayer sample visible, c) attachment of OmniProbe needle for lift out, and d) thinning of lift out specimen on TEM grid until desired thickness (prepared by C. Appleget). Figure 39: TEM images showing the depth of amorphization of Si prepared by Ga + and Xe + FIB [199]. 53 3.2.4. Transmission Electron Microscopy (TEM) Transmission electron microscopy (TEM) is a characterization technique in which a high energy electron beam passes through a sufficiently thin (<100 nm) sample. The electrons are generated by a field emission gun and are condensed using a condenser lens [200]. The interactions between the electrons and atoms can be used to form an image and observe features such as the crystal structure, dislocations, and grain boundaries. To characterize these optical NMs, a variety of TEM imaging modes are used: including bright field (BF) TEM, dark field (DF) TEM, and selected area electron diffraction (SAED). An overview of these TEM operation modes is shown in Figure 40. Figure 40: Schematics of (a) bright field, (b) dark field, and (c) selected area diffraction TEM operating modes [200]. BF TEM images show mass-thickness and diffraction contrast based on the scattering and absorption of electrons, therefore showing elemental contrast, an example is shown in Figure 41a. 54 DF TEM images are formed by blocking the direct beam with an aperture, allowing specific diffracted beams to pass through the objective aperture, an example of which is shown in Figure 41b. DF TEM is used to observe planar defects, stacking faults, particle size, etc. In SAED, a parallel beam interacts with the sample, and an aperture is inserted to define an area from which the diffraction pattern is observed. SAED diffraction patterns (DPs) are typically spot patterns that correspond to a single-crystal diffraction or appear as ring patterns that correspond to multiple grains of differing orientations. SAED is used for determination of lattice parameters, phase identification and to determine growth direction. 3.2.5. Scanning Transmission Electron Microscopy (STEM) Scanning transmission electron microscopy (STEM) combines the principles of TEM with the scanning (rastering) electron microscopy employed in other techniques [201]. In STEM, the microscope lenses are aligned to converge the beam and signals are collected across the sample in a point-by-point fashion using a very small spot size (1 nm). The transmitted electrons form two Figure 41: (a) Bright field and (b) dark field cross-sectional TEM micrographs of an AlN/SiO2 multilayer sample, synthesized by Appleget. 55 types of images, bright field (BF) and high-angle annular dark field (HAADF). In BF STEM, the electrons transmitted along the axis of the beam are gathered using a BF detector. The HAADF STEM mode, also called the Z contrast mode, is based on high angle incoherent scattered electrons collected using a HAADF detector, which forms annular diffraction micrographs [202]. HAADF STEM is a valuable method with strong compositional sensitivity, detecting differences of the element irradiated based on the atomic number, an example is shown in Figure 42a. As previously mentioned in Section 3.2.2, when the electron beam interacts with the sample, x-rays are emitted, and can be used in energy dispersive x-ray spectroscopy (EDS) to measure sample composition [201]. When coupling STEM and EDS, it is possible to create a spatial compositional map of the sample (such as Figure 42b-e) with the resolution close to the size of the electron beam. Additionally, EDS point scans and EDS line scans can be used to calculate compositions of regions inside the sample. Figure 42: (a) Cross-sectional HAADF STEM of an AlN/SiOx multilayer sample and STEM EDS maps showing the distribution of (b) Al, (c) N, (d) Si, and (e) O (synthesized by C. Appleget). 56 3.3. Optical Characterization Methods Optical characterization techniques are typically non-destructive, and explore the change of intensity, energy, phase, direction, or polarization of the light wave after interaction with sample being studied [203]. When light interacts with matter, it exchanges energy with the material due to the oscillating electromagnetic field. An overview of the interaction between light and matter is shown in Figure 43. When light interacts with matter, the light can be scattered (either elastically or inelastically), absorbed, or transmitted. This interaction is governed by the intensity and energy of the incident photons, as well as physical, chemical, and structural properties of the matter. In this work, spectrophotometry and ellipsometry are the two main optical characterization methods and are detailed in the following sections. Figure 43: A summary of the interaction between light and matter, with the approximate energies of fundamental excitations [203]. 57 3.3.1. Spectrophotometry Spectrophotometry is one of the most common methods to investigate the interaction of light with a material, and it is where the quantitative measurement of light reflection or transmission properties of a material is measured as a function of wavelength [203, 204]. Spectrophotometry measures how much a material or chemical absorbs light by measuring the intensity of light as a beam passes through the material. Spectrophotometry is widely used for quantitative analysis in fields such as chemistry, physics, biochemistry, materials science, and for clinical applications. Typical spectrophotometry instrumentation covers the spectral range from near ultra-violet (UV) to near infra-red (NIR). The spectrophotometry follows the Beer Lambert law of light absorption, meaning that when a monochromatic light passes through a medium, the amount of light that is absorbed is directly proportional to the number of light absorbing molecules in that medium. Typical instrumentation consists of light sources, a diffraction grating based monochromator, a sample holder, and one or more detectors, a schematic of which is shown in Figure 44. A routine experimental mode in spectrophotometry is to perform transmittance measurements, and it consists of measuring how much a known incident light passes through the sample, expressed as a percentage: 𝑇 = 𝐼 𝐼 + ×100 (21) where I0 is the light transmitted through a reference blank, I is the light transmitted through the sample, and T is the transmittance. 58 Figure 44: Schematic of a UV-Vis Spectrophotometer 3.3.2. Variable Angle Spectroscopic Ellipsometry (VASE) Ellipsometry is a high precision thin film measurement technique that measures the interaction of polarized light with a sample, as shown in Figure 45 [205]. In ellipsometry, the change in the polarization state of light reflection from the surface of a sample is measured and expressed as psi (Ψ) and delta (Δ). Psi is the ratio of the amplitude diminutions and delta is the phase difference induced by the reflection: tan𝛹 = O𝑅 ' O |𝑅 0 | (22) where 𝑅 ' and 𝑅 0 are the total reflection coefficients which correspond to Fresnel coefficients. 𝛥 = 𝛿 " −𝛿 ! (23) where 𝛿 " is the phase difference before and 𝛿 ! is the phase difference after reflection. 59 Figure 45: Interaction of polarized light with a sample [206]. These values relate to the ratio of Fresnel reflection coefficients for p- and s- polarized light: 𝜌 = 𝑅 ' 𝑅 0 = tan(𝛹)𝑒 $9 (24) Fundamentally, ellipsometry only measures Δ and Ψ values, and these measurements are highly sensitive and reproducible. From these values, layer thicknesses and optical constants of the films can be calculated by constructing a model. The model contains known parameters such as the wavelength of the incident light, the incident beam polarization state, and the angle of incidence. The model is then used to calculate unknown optical physical properties such as film thickness, refractive index, surface roughness, interfacial regions, composition, crystallinity, anisotropy, and uniformity [207]. A summary of the physical properties that can be characterized by spectroscopic ellipsometry is shown in Figure 46. From this model, the best fit is calculated, and results in a generated data set. An example is shown in Figure 47, where the Δ and Ψ values are measured using VASE techniques, and the data is fitted to minimize the mean squared error (MSE) for the index of reaction and extinction coefficient, Figure 48 and Figure 49, given a measured film thickness. 60 Figure 46: Characterization of physical properties by spectroscopic ellipsometry [208]. The variable angle spectroscopic ellipsometry technique refers to the combination of varying the angle of incidence and spectroscopic (across a range of wavelengths) measurements. This generates a large data set about the optical behavior of the sample. The incident angles are chosen to maximize the sensitivity of the measurement, and the choice of incident angles depends on the optical constants of the samples [208]. For semiconductor characterization, the incidence angle is typically around 70⁰, and it is important to note that when measurements are performed at normal incidence, the p- and s-polarizations can no longer be distinguished so ellipsometry measurements are impossible. The benefit of these ellipsometry measurements is twofold. First, incorporating experimental optical data rather than literature estimates, such as index of refraction at a single 61 wavelength, yields higher fidelity MBI calculations. Secondly, this is important because index of refraction can vary with synthesis method (sol-gel, magnetron sputtering) and with sputtering deposition parameters, as previously discussed in Section 2.2.2, and these variations are important to capture quantitatively [89, 209]. Effects such as the oxygen partial pressure, substrate temperature, and working pressure can yield variations in refractive index [209-211]. This is shown as the refractive index varies with oxygen concentration in the deposition into TiO2 thin films in Figure 49 [212]. For example, a study by Meng et al found that varying the substrate temperature from 150⁰C to 500⁰C during magnetron sputtering of a TiO2 film, the index of refraction at λ = 500 nm varied from 2.38 to 2.47 [89]. Figure 47: Ellipsometry curve fittings for TiO2 sample prepared in 10% oxygen environment [212]. 62 Figure 48: Variation of refractive index (n) with % oxygen content in a TiO2 thin film [212]. Figure 49: Optical constants (refractive index and extinction coefficient) of an 87 nm AlN thin film deposited by plasma-enhanced atomic layer deposition (PEALD) [213]. 63 3.4. Mechanical Characterization This work centers on a comprehensive study of the mechanical behavior of transparent optical NMs, therefore it is important to investigate the deformation mechanisms of the multilayers on multiple length scales. Given the sizes of the samples (1-2 μm total thickness), standard ASTM testing techniques are not directly applicable for some thin film properties, particularly for brittle ceramic thin films, where accounting for size effects is imperative. Thus, testing designed for mechanical behavior of nanostructured materials and thin films must be utilized. Additionally, due to the length scales of NMs, as detailed in Section 2.1, the mechanical properties will likely be significantly different than bulk material behavior due to the influence of the interfaces, surfaces, and nanoscale features. Thus, due to the thickness of the films, nanoindentation is used for measuring hardness and modulus, while Vickers is used to investigate qualitative fracture behavior. 3.4.1. Nanoindentation Nanoindentation is a mechanical testing technique which consists of indenting a surface with a very small (100-200 nm tip radius) tip. The method of indenting the surface of a sample is similar to other hardness testing techniques such as Rockwell and Vickers testing, but the difference is in the resolution of the test due to the small tip size. In hardness testing, it is generally agreed upon that the indentation depth must be less than 10% of the film thickness to avoid substrate effects [214]. This test can achieve measurements at indent depths as small as 100 nm, which is important for the study of thin films. Most nanoindentation testing is concerned with measuring the elastic modulus and hardness of a material. In nanoindentation, the penetration depth of a diamond indenter is pressed into a sample with a prescribed load. The output of the test is a load-displacement curve, an 64 example is shown in Figure 50b, where there is elastic-plastic loading followed by elastic unloading. The elastic equations of contact are used with the unloading data to then determine the hardness and elastic modulus of the material. The three-sided Berkovich indenter geometry is the most common tip used in nanoindentation, a schematic of which is shown in Figure 50a. For this geometry, the contact depth, as related to the total depth of penetration during testing, is calculated as [215]: ℎ 6 = ℎ (78 −𝜀 𝑃 (78 𝑑𝑃 𝑑ℎ (25) Where the factor ε depends upon the indenter shape, and once the contact depth (hc) is known, the area of contact A is found from the indenter geometry. Figure 50: a) Schematic showing the indent geometry and dimensions, and b) typical nanoindentation load- displacement curve [215]. The reduced elastic modulus, E, and the hardness, H, can then be calculated using the Oliver and Pharr method [215]: 𝐸 = 𝑑𝑃 𝑑ℎ 1 2 √𝜋 √𝐴 (26) 65 𝐻 = 𝑃 (78 𝐴 (27) Due to the small tip geometry and penetration depth, care should be taken for avoiding the effects of thermal drift, surface roughness, substrate effects, creep, inaccurate tip area calculations, mounting of the sample, and minimizing the initial penetration depth [214]. An advantage of nanoindentation testing is that arrays of indents can be performed on a single sample, yielding statistically significant data about the physical properties. Multiple curves from a single sample can then be analyzed, such as the indents shown in Figure 51, to estimate hardness and modulus of the sample. Figure 51: Load-displacement curves from nanoindentation testing of an AlN/SiO2 sample, 10 indents of a 10x10 (100 indent) array are shown (by C. Appleget). 3.4.2. Vickers Hardness The Vickers hardness test method is a type of microhardness test that uses a diamond indenter with a pyramidal geometry at a range of loads. The diagonals of the indents are then measured, shown in Figure 52, and converted to a hardness value. The Vickers hardness numbers 66 are designated as the Vickers Pyramid Numbers (HV). Vickers is one of many microhardness testing methods, and a comparison of common hardness testing methods and their respective hardness scales is presented in Figure 53. Vickers testing is well suited for measuring the hardness of small, selected specimen regions with proper specimen surface preparation [203, 216], and an example of Vickers indents performed on a NM sample is shown in Figure 54. The Vickers hardness is calculated as: 𝑉𝐷𝐻 = 2𝑃 𝑑 ! sin 136° 2 = 1.8544 𝐹 𝑑 ! (28) Where the square pyramid indenter has opposite faces at an angle of 136° and edges at 148°. The average diagonal measured from corner to corner of the residual impression is d and F is the load applied during indentation. Figure 52: Overview of Vickers Indenter geometry [217]. 67 Figure 53: Comparison of several hardness scales [218]. Figure 54: a) SEM isometric view of Vickers indent (with Pt coating to prevent charging of sample) and b) FIB top view of the same Vickers indent (by C. Appleget). 68 Chapter 4 : Optical and Mechanical Characterization of Metal/Ceramic NMs The background and techniques discussed to this point were used to understand the optical and mechanical characterization of metal/ceramic NMs. The investigations resulted in the publication of Optical and Mechanical Characterization of Sputtered AlN/Ag Multilayer Films that can be found in the journal Advanced Engineering Materials in Volume 21, Issue 5 (DOI: 10.1002/adem.201801268). The study and this section explore the synthesis of metal/ceramic multilayers with a wide range of transmittance values while examining their resulting microstructure and mechanical behavior. In this work, AlN/Ag multilayers serve as a model optical NM system that is capable of tuned transmittance and also builds upon previous studies on the mechanical behavior of metal/ceramic NMs as discussed in Chapter 2. AlN/Ag nanomultilayers were synthesized by reactive DC magnetron sputtering with either repeated bilayer thicknesses or layer thicknesses optimized for transmittance. Five structures were investigated, and all showed agreement with predicted transmittance values. The measured hardness values ranged from 4.9 ± 0.0 to 19.4 ± 0.9 GPa, while the failure was brittle-dominated fracture, independent of the Ag layer thickness. The optically optimized samples exhibited a fourfold increase in hardness and increased optical performance over the repeated bilayer samples, despite the similar deformation behavior. 4.1. Introduction Optical multilayers, which are a subset of nanomultilayers (NMs), are comprised of nanoscale layers arranged to alter the transmission and reflection of light by tuning the layer thickness and individual layer material properties [32]. Optical multilayers are typically composed of either ceramic/ceramic layers, with a focus on maximizing transmittance or reflectance, or 69 metal/ceramic layer combinations, with a focus on spectral selectivity or sheet resistivity [40, 54, 57, 219]. Despite the significant interest in implementing these materials in applications that require structural integrity under high loading environments, studies on mechanical properties of optical NMs have been minimal [7, 33]. To better leverage the promising optical properties, further studies are necessary to understand the mechanical behavior of optical NMs by investigating optical and mechanical properties concurrently. Furthermore, a comprehensive understanding of the relationship between deformation behavior and optical properties also depends on thorough characterization of nanoscale features and microstructure. However, the microstructural characterization of multifunctional optical coatings is typically limited [10, 57, 220-222] even though these features have been shown to have a significant influence on the non-optical [12, 60] and optical [221] properties of metal/ceramic multilayer films. In general, the characteristic length scale of NMs is the layer thickness rather than the grain size, and since the interfaces contribute to a substantial portion of the sample volume, interfacial properties are more significant in NMs than in their coarse-grained counterparts [1, 3, 15, 16]. Therefore, material selection, synthesis, and resultant interface characteristics are crucial factors in the properties exhibited in these materials. To date, most efforts on characterizing the mechanical behavior of NMs have focused primarily on non-optical metal/metal [27, 60, 61, 100- 103, 106-111, 223] and metal/ceramic multilayer combinations [5, 112-118, 120, 121]. For non- optical metal/ceramic NMs, several explanations have been proposed for the effect of layer thickness on hardness. For example, some studies have shown that mechanical properties are primarily dependent on the bilayer period, such that a reduction in bilayer period correlates to an increase in hardness [110, 112], while other studies have shown that the mechanical response is 70 also dependent on the metal/ceramic layer thickness ratio (component fraction) and material selection [119, 158]. In this work, AlN/Ag multilayers were synthesized with a wide range of layer thicknesses and optical performance in order to explore the variation in mechanical response for samples with optimized and non-optimized transmittance. The mechanical behavior of these samples was examined by performing indentation at different length scales, while TEM was employed to determine nanoscale features and layer interfaces. Combining these characterization techniques with optical measurements is an important first step in studying the relationship between optical and mechanical properties, which is imperative for applications requiring both light transmittance and robust structural performance. 4.2. MBI Predictions Theoretical design predictions for the AlN/Ag samples optimized for maximum transmittance were made using the Multiple-Beam-Interference (MBI) Recursive Method detailed in studies by Lu et al [33, 96, 99]. An optimization scheme based on this method was written in MATLAB to design AlN/Ag multilayer systems with maximized transmittance in the UV/Vis wavelengths (300-800 nm). Design parameters set in the optimization loop include: (1) minimum Ag layer thickness of 5 nm to ensure layer uniformity [8], (2) five layer minimum in the AlN/Ag multilayer, (3) > 1 μm total thickness of the multilayer film for nanoindentation and mechanical testing [224]. Recent studies on Ag-based NMs have suggested that seed layers or minimum Ag layer thickness improve Ag nanolayer quality [8, 225, 226], hence the implementation of a 5 nm minimum layer thickness theoretical design constraint. Complex index of refraction values of AlN and Ag were used as inputs and were derived from model fitting of variable angle spectroscopic 71 ellipsometry (VASE) measurements (J.A. Woollam). The layer thicknesses and predicted % transmittance of two multilayer samples (MBI #1 and MBI #2) calculated using the MBI method are detailed in Table 1. Table 1: Predicted layer thicknesses of optimized MBI samples Layer MBI #1 MBI #2 AlN Layer #3 391 nm 574 nm Ag Layer #2 5 nm 5 nm AlN Layer #2 385 nm 451 nm Ag Layer #1 5 nm 5 nm AlN Layer #1 339 nm 287 nm Predicted % T 300-800nm 61.8% 62.7% 4.3. Synthesis of AlN/Ag Multilayers Samples were synthesized using DC reactive magnetron sputtering, with two sources used to deposit AlN and Ag onto Corning Eagle 2000 glass substrates. Each AlN layer was deposited under an Ar flow of 37.5 sccm and a reactive N2 flow of 12.5 sccm (1.2 Pa working pressure) at 150 W. All Ag layers were deposited under an Ar flow of 37.5 sccm (1.0 Pa working pressure) at 100 W. The on-times for deposition were controlled with a pneumatic shutter, allowing for deposition from one source at a time and control over layer thickness. An XP-2 profilometer (AMBiOS) was used to measure surface profiles of the substrate before and after deposition to calculate sputtering rates as well as the residual stresses of the as-deposited films using Stoney’s equation [192, 227]. Five total types of structures were synthesized with differing bilayer thicknesses, either optimized or not optimized for optical transmittance. The three non-optimized samples with repeated bilayer thicknesses span a range of layer thicknesses (10 nm – 125 nm) and 72 a range of ceramic: metal layer ratios (~1:1 to 10:1). The ceramic:metal layer thicknesses are (AlN(47)/Ag(21)), (AlN(20)/Ag(20)), and (AlN(127)/Ag(10)), respectively, where the layer thicknesses are indicated in parenthesis. Throughout this paper, the term “repeated bilayers” refers to samples where the individual thickness of the AlN and Ag layers are constant throughout the sample. These repeated bilayer samples provide both an optical and mechanical baseline of properties to compare to the optically optimized samples. In addition, samples MBI #1 and MBI #2 were synthesized with layer thicknesses optimized for optical properties from 300 to 800 nm wavelengths using the aforementioned MBI method. Sample MBI #1 consisted of five layers: AlN(339)/Ag(5)/AlN(385)/Ag(5)/AlN(391) with a total predicted % transmittance of 61.8%. Sample MBI #2 also consisted of five layers: AlN(287)/Ag(5)/AlN(451)/Ag(5)/AlN(574) with a total predicted % transmittance of 62.7%. SEM cross-sections of the five multilayer samples in this study are shown in Figure 55a-e, where the bright layers are Ag and the dark layers are AlN, with a vertical growth direction. Figure 55a-c shows SEM cross-sections of the repeated bilayer thickness samples while Figure 55d-e shows MBI #1 and MBI #2. 4.4. XRD and Optical Characterization of AlN/Ag Multilayers X-ray diffraction (XRD) spectra of the deposited AlN/Ag multilayers, shown in Figure 56a, were obtained using an Ultima IV diffractometer (Rigaku). The AlN(20)/Ag(20) and AlN(47)/Ag(21) samples exhibited AlN (0 0 2) and Ag (1 1 1) textures. Spectra for the MBI #1, MBI #2 and AlN(127)/Ag(10) samples showed low intensity peaks that are not visible after globally normalizing intensity; therefore insets are included to show detail. The % transmittance of the deposited AlN/Ag multilayers was measured by a Cary UV-Vis 60 Spectrophotometer with a solid sample holder, the profiles for each sample are presented in Figure 56b. The average measured transmittance values of the five samples from 300 to 800 nm are summarized in Table 73 2. From these results, it is clear that the repeated bilayer thicknesses without optical optimization yield low experimental average transmittance. The AlN(47)/Ag(21) and AlN(20)/Ag(20) samples have nearly 0% light transmission in the UV/Vis wavelength spectrum due to the presence of thick Ag layers (> 10 nm) [228]. AlN(127)/Ag(10) has a 16% transmission due to the thinner Ag layers. Thus, for these systems, minimizing the Ag layers appears to improve the transmittance values. It can be observed that the two MBI-based samples, MBI #1 and MBI #2, have higher average transmittance than those with the repeated bilayers. Close agreement between average predicted and experimental % transmittance values was achieved for both optimized samples, for MBI #1 (TMBI = 61.8% vs. Texp = 58.1%) and for MBI #2 (TMBI = 62.7% vs. Texp = 61.3%). Figure 55: SEM cross-sections showing multilayers and corresponding top views indicating appearance of samples: (a) AlN(47)/Ag(21), (b) AlN(20)/Ag(20), (c) AlN(127)/Ag(10), (d) AlN/Ag MBI #1, and (e) AlN/Ag MBI #2 where the layer thicknesses are indicated in parentheses in nanometers. 74 Figure 56: (a) Normalized X-Ray Diffraction (XRD) patterns of AlN/Ag multilayer system from 2θ 10-100⁰ and corresponding crystal structure where insets show detail of low intensity peaks and (b) Experimental transmittance curves of the AlN/Ag multilayers performed by spectrophotometry. 75 Table 2: Results of Experimental and Theoretical Optical and Mechanical Properties. Sample Total Thickness Exp. Hardness (GPa) Rule-of- Mixtures Hardness (GPa) Exp. Elastic Modulus (GPa) Ave Residual Stress (MPa) Exp. %T (300-800 nm) Predicted %T (300- 800 nm) AlN(47)/Ag(21) 1.0 μm 5.3 ± 0.1 9.05 109.6 ± 0.9 444 0.0% 0.0% AlN(20)/Ag(20) 1.1 μm 4.9 ± 0.0 6.85 95.7 ± 0.6 271 0.6% 0.0% AlN(127)/Ag(10) 1.0 μm 13.3 ± 0.3 11.76 150.2 ± 1.4 290 15.5% 3.6% AlN/Ag MBI #1 1.1 μm 18.3 ± 1.0 12.42 176.1 ± 6.3 338 58.1% 61.8% AlN/Ag MBI #2 1.3 μm 19.4 ± 0.9 12.51 158.0 ± 4.1 250 61.3% 62.7% 4.5. Microstructural and Mechanical Characterization of AlN/Ag Multilayers To investigate the effects of the layer thickness and AlN:Ag layer ratio on the deformation behavior at the nano and macroscale, both Vickers and nanoindentation tests were performed on all samples. Nanoindentation was performed using a force-controlled, constant loading rate load function with a 100 nm Berkovich tip, using a Hysitron Triboindenter. Table 2 shows the means and standard deviations of the nanoindentation results and residual stresses of the films. The nanoindentation hardness values are within the range from calculated rule-of-mixtures hardness values, which span from 9.05-12.51 GPa [229]. For reference, bulk AlN hardness and elastic modulus values are approximately 12.6 GPa and 308 GPa, respectively, and bulk Ag hardness and modulus values are approximately 1.1 GPa and 74 GPa. A minimum of 100 indents were performed per multilayer sample, yielding hardness values in the range of 4.9 to 19.4 GPa and reduced moduli values in the range of 95.7 to 176.1 GPa. The MBI-calculated samples exhibited an increase in hardness as compared to monolithic AlN and the repeated bilayer AlN/Ag multilayers, while the deformation behavior remained predominantly brittle despite the presence of Ag layers. The presence of 5 nm Ag layers in the MBI-calculated samples resulted in increased hardness as compared to monolithic AlN, but the deformation behavior remained predominantly 76 brittle. From the non-optimized (repeated bilayer) samples, the AlN(127)/Ag(10) with a ~10:1 ceramic:metal ratio and the smallest Ag layer thickness was the only sample to exceed the rule-of- mixtures calculations for hardness. Previous studies on multilayers with nanoscale layer thicknesses have shown measured hardness values exceeding the rule-of-mixtures calculations [230, 231], and others have shown mixed results [111, 117, 158, 232], similar to the results of the present study. Due to the thickness of the films (~1 μm), nanoindentation is used for measuring hardness and modulus, while Vickers is used to qualtitatively investigate fracture behavior. Representative Vickers indents performed at 2.94 N with a 10 s dwell time are presented in Figure 57. Ceramics, such as AlN, have high hardness but are also brittle and susceptible to cracking, and cracks emanating radially from the corners of the Vickers indents are typical. All samples, with either repeated bilayer or optimized MBI, show brittle failure under Vickers indentation. The samples with thicker Ag layers, AlN(47)/Ag(21) and AlN(20)/Ag(20) show less circumferential cracking, likely due to the plastic deformation occurring in the metallic layers. The indents of MBI #1 and MBI #2, Figure 57d-e, exhibit radial cracking, with some circumferential cracking. In the Vickers indents of the MBI samples, the light fringe distortions visible in the indents reveal uplift and curvature of the indent sides, suggesting that the multilayers are delaminating and peeling upward.[233] FIB cross-sections of the MBI #1 and MBI #2 Vickers indents were prepared to further investigate the fracture behavior of the optically optimized samples and are presented in Figure 58a and Figure 58d. Before performing the cross-section milling, a 15 nm Pt coating was deposited onto the sample to mitigate charging, and a 1 µm carbon coating was deposited to protect the 77 region of interest during ion milling. The indent cross-section of MBI #1 in Figure 58a shows multiple areas of delamination indicated by the white arrows, occurring between the first AlN layer and the substrate as well as between AlN and Ag layers. The delaminations indicate transverse tensile stresses during indentation [234]. This is in agreement with studies showing that in metal/ceramic NMs, the deformation mechanisms are commonly characterized by cracks first forming in the brittle layer, then growing to neighboring ductile layers [113, 235, 236]. Black arrows highlight the vertical cracking spanning the brittle AlN layers and terminating at the AlN/Ag interface. Subsurface cracking was also visible in the MBI #1 indent, which suggests high concentrations of stress within the substrate during indentation. MBI #2, shown in Figure 58c, also exhibits delamination between the first AlN layer and the substrate, but no delamination at the AlN/Ag interface and less vertical cracking as compared to MBI #1. Previous studies have shown that in NMs with thin metallic layers (< 10 nm) the fracture behavior remains dominated by the Figure 57: Representative top-view optical microscope images of indents after Vickers indentation (performed at 2.94 N with a 10 s dwell time) for multilayers samples: (a) AlN(47)/Ag(21), (b) AlN(20)/Ag(20), (c) AlN(127)/Ag(10), (d) AlN/Ag MBI #1, and (e) AlN/Ag MBI #2. Circumferential Cracking 78 ceramic layer deformation, which is in agreement with the observations for both MBI samples [113]. Figure 58: FIB and TEM cross-sectional views of AlN/Ag MBI #1 (a-c) and MBI #2 (d-f). (a) MBI #1 FIB cross- section of Vickers indent post-deformation and (b) bright field STEM with SAED inset of the as-sputtered sample including (c) bright field TEM showing a representative cross-section of an Ag layer breakthrough in MBI #1. (d) MBI #2 FIB cross-section of Vickers indent post-deformation, (e) bright field STEM with SAED inset of the as- sputtered sample including (f) bright field TEM showing a representative cross-section of a continuous Ag layer in MBI #2. In FIB cross-sections, the protective top layer of carbon is labeled; white arrows highlight regions of delamination and black arrows indicate through-layer cracking. Substrate Substrate 79 To further understand the complex relationship between mechanical and optical properties of metal/ceramic multilayer films, it is important to further investigate the microstructure. For example, in metal/ceramic systems it has been shown that nanoscale interface and layer roughening noticeably decreases transmittance in optical multilayer films [237, 238]. Slight changes in layer thicknesses (<10%) have also been shown to yield measurable changes in multilayer properties such as transparency [37], spectral filter behavior [239], and plastic behavior [158]. Despite previous work indicating that nanoscale features influence properties, studies of multifunctional optical coatings typically involve only limited microstructural characterization, with most studies utilizing SEM [10, 37, 57, 96, 220, 240], and with few studies examining layer interfaces and grain structures through transmission electron microscopy (TEM) [221, 222]. Therefore, in order to investigate these nanoscale effects in the MBI-optimized samples, TEM and STEM were performed to analyze the microstructure and interfaces. Figure 58 shows representative bright field TEM, prepared via FIB lift out preparation, of MBI #1 (Figure 58b) and MBI #2 (Figure 58e), where the indicated growth direction is vertical. The inset selected area electron diffraction (SAED) patterns of the multilayer regions of MBI #1 and #2 indicate a nanocrystalline microstructure with different orientations; this is in agreement with the acquired XRD spectra. The cross-sectional STEM Figure 58b and Figure 58e, highlight the multilayered structure in more detail. The Ag layers are visible as thin dark layers and show a sharp contrast compared to the AlN layers with distinct layer interfaces, likely attributed to the low miscibility of noble metals with nitrides [241, 242]. MBI #1 has a slightly higher occurrence of Ag layer breakthroughs, which corresponds to more AlN/Ag interface delamination observed in the indentation cross-sections, suggesting that the breakthroughs act as preferential paths for crack propagation. Over 50 cross-sectional BF TEM images at 600kx magnification of each MBI sample 80 were analyzed, corresponding to imaging over 8.5 µm linear length of the Ag layers for each sample. From these images, instances of Ag layer breakthroughs where the Ag has not formed a continuous layer were counted. Representative cross-sections of breakthroughs versus continuous Ag layers are shown in Figure 58c and f. The results showed 31 Ag layer breakthroughs in MBI #1 and 13 Ag layer breakthroughs in MBI #2, corresponding to approximately 3.6 breakthroughs/µm and 1.5 breakthroughs/µm, respectively. The increase in the number of breakthroughs in MBI #1 corresponds to more AlN/Ag interface delamination and through-layer cracking observed in the indentation cross-sections. Overall, MBI #1 and #2 exhibit similar microstructures, which is to be expected given the similar layer thicknesses and equivalent sputtering conditions. Both samples show layer uniformity with a slight progressive roughening through the increasing film thickness, in agreement with previous studies [243]. 4.6. Conclusions In summary, repeated and non-repeated bilayer AlN/Ag systems were synthesized via reactive DC magnetron sputtering and characterized to explore how microstructure and mechanical behavior relate to optical performance. AlN/Ag samples with layer thicknesses calculated for higher transmittance from 300-800 nm showed close agreement with predicted optical values. The optically optimized samples displayed a 4x increase in hardness and increased optical performance over the repeated bilayer samples, although the fracture behavior was similar across all samples. Moreover, it was observed that in repeated bilayer samples, decreasing the AlN/Ag bilayer thicknesses does not result in improved mechanical behavior and drastically decreases optical performance. Overall, this combinatorial approach for studying optical multilayers highlights a promising path forward for optimizing multifunctional properties in transparent materials. 81 Chapter 5 : Microstructural Variations in Ceramic/Ceramic AlN/SiO2 NMs The following work is published as a journal article titled Exploring Microstructural Variations in Highly Transparent AlN/SiO2 Nano Multilayers and is published in the journal Optical Materials Express in Volume 10, Issue 4, pages 850-959 (DOI: 10.1364/OME.389156). This work builds upon the methodology developed in Chapter 4 and extends the findings into ceramic NMs capable of achieving high (>95%) transmittance in the UV/Vis to NIR wavelengths. In this study, the microstructure of optically optimized transparent AlN/SiO2 nanomultilayers were investigated and compared with baseline repeated bilayer configurations. The multilayered films were synthesized by magnetron sputtering and characterized by transmission electron microscopy and spectrophotometry with multifunctional behavior evaluated by nanoindentation and residual stress analysis. The optically optimized AlN/SiO2 multilayers exhibit higher transmittance (%T300-800nm≈95%), distinct crystalline/amorphous interfaces, and changes in the grain morphology as compared to the periodic baseline samples (%T300-800nm≈70- 80%). Varying both layer thickness and layer ratio to maximize transparency showed a significant impact on microstructure and interface character. 5.1. Introduction Nanomultilayers (NMs) offer a wide range of tunability, as they demonstrate the ability to achieve desired properties through modulation of layer thicknesses, interfaces, and composition [1, 3]. Optical coatings for both single layer and multilayer systems have been widely studied, mostly focusing on the achievable optical behavior such as transmittance or reflectance, where highly transparent NMs are typically composed of layers with alternating high and low indices of refraction [35, 94]. These types of transparent NMs have attracted interest due to the possibility of introducing additional functionalities, such as low sheet resistance [57, 244] or favorable 82 mechanical properties [33] while retaining transparency in the UV-Vis and NIR wavelengths. However, before designing additional functionalities, it is imperative to better understand the juxtaposition between microstructure and interfaces to the optical behavior. Currently, methodologies do exist for optimizing transmittance in optical NMs [33], but the impact of tuning layer thickness for maximized transparency on the microstructural and interfacial character against the backdrop of multifunctionality remains relatively unexplored. In transparent NMs, the relationship between interfaces and optical performance has been investigated by focusing on minimizing interfacial roughness to avoid scattering effects [245]; however, the impact of nanoscale microstructural features on optics has been studied to a lesser extent, particularly focusing on the effects of varying layer thicknesses. For example, crystalline- crystalline-type interfaces with varying degrees of coherency in non-transparent NMs have been investigated and shown to follow classical interface interactions, where the interfaces act as defect sources and sinks, as well as barriers to defect mobility, thus dominating plastic deformation behavior [105, 246]. In contrast, amorphous-crystalline interfaces, which are common in transparent NMs, have been subject to more limited studies, and are expected to play essential roles in optical performance and functionality [247, 248]. Therefore, in order to expand upon current studies beyond solely optics and into functionality, this study focuses on the synthesis and subsequent characterization of a model AlN/SiO2 multilayer amorphous-crystalline system. A comprehensive exploration into the effect of varying layer thicknesses and layer ratios in optically optimized (aperiodic) and baseline repeated bilayer (periodic) multilayers uncovers the role of microstructure and interface character in highly transparent NMs. 83 5.2. Experimental Methods AlN/SiO2 multilayers were deposited on Corning Eagle 2000 glass substrates by DC and RF magnetron sputtering using two targets. AlN was deposited from a 99.999% Al target under an Ar flow of 37.5 sccm and a reactive N2 flow of 12.5 sccm (1.2 Pa working pressure) at 150 W of DC power. SiO2 was deposited from a 99.99% SiO2 target under an Ar flow of 0.6 Pa working pressure at 40 W of RF power. The on-times for deposition were controlled with a pneumatic shutter, allowing for precise control of layer thickness and deposition from one source material at time. An XP-2 profilometer (AMBiOS) was used to measure surface profiles for calculation of sputtering rates and an estimation of residual stresses of the films using Stoney’s equation. Two different types of multilayer samples were synthesized in this study: samples designed for maximum optical properties and those with repeated bilayer thicknesses as a baseline. The repeated bilayer samples were AlN(50nm)/SiO2(50nm) and AlN(100nm)/SiO2(100nm), where the parenthesis indicate the layer thicknesses held constant throughout the total 1 µm film thickness. From this point forward, the ‘nm’ designation will be dropped for concision. These constant bilayer thickness samples have the same total number of layers, and therefore the same number of interfaces, as the optically optimized samples, and provide a comparative standard for optical, microstructural, and mechanical properties. The optically optimized samples synthesized for maximized optical transmittance were designated “MBI 10 Layers” and “MBI 20 Layers,” indicating the total number of layers throughout the same 1 µm total film thickness. Optimized AlN/SiO2 layer configurations were calculated for maximized optical transmittance in the UV/Vis wavelengths (300-800 nm) using the Multiple Beam Interference (MBI) Recursive method [33, 249]. MBI formulas were implemented in an in-house MATLAB optimization scheme [250] with experimental optical 84 constants measured from the as-sputtered AlN and SiO2 films. The spectroscopic index of refraction data as measured from ellipsometry for the monolithic AlN and SiO2 films synthesized at the same deposition conditions. The refractive indices were approximately n550nm = 2.18 and n550nm = 1.43 for AlN and SiO2, respectively. The calculated layer thicknesses required non- constant bilayer thickness, and the 10 Layer and 20 Layer sample configurations are detailed in supplemental Table 3. The optically optimized samples exhibited good agreement between MBI- calculated predictions, within 2% for the 10 Layer sample (94.0% experimental vs. 96.1% predicted) and within 1% for the 20 Layer sample (95.3% vs. 96.2% predicted). The high transparency confirms good agreement to the optical predictions and precise deposition of layer thicknesses across all multilayered samples in this study. Sample microstructural characterization, including energy dispersive spectroscopy (EDS), was performed using a FEI Talos 200s (Thermo Scientific) transmission electron microscope (TEM). The TEM foils were prepared via focused ion beam (FIB) lift-out preparation using a Zeiss Auriga Dual Beam FIB. Hardness and reduced elastic moduli measurements were performed by nanoindentation using a Hysitron Triboindeter (Buehler) and a force-controlled, constant loading rate function with a 100 nm Berkovich tip. 5.3. Results and Discussion In general, studies in the field focus in non-optical NMs with repeated bilayers [118, 251] or optical behavior with varying layer thicknesses [57], both of which do not provide a direct link between interfaces and microstructural properties in high transparency NMs. Thus, this study provides a connection beyond optics by integrating a comprehensive study of the microstructural variations in both repeated bilayer and optically optimized NMs. Figure 59 presents a summary of the overall characteristics of the synthesized multilayer films. Figure 59a illustrates a schematic of 85 Table 3. Summary of layer thicknesses of MBI 10 Layer and MBI 20 Layer Samples AlN/SiO2 10 Layers AlN/SiO2 20 Layers Material Thickness (nm) Material Thickness (nm) SiO2 72 SiO2 128 AlN 5 AlN 5 SiO2 226 AlN 5 SiO2 249 SiO2 29 AlN 5 AlN 5 SiO2 22 AlN 5 SiO2 155 SiO2 19 AlN 5 AlN 5 SiO2 27 AlN 5 SiO2 250 SiO2 224 AlN 5 AlN 5 SiO2 34 AlN 5 SiO2 249 SiO2 211 AlN 5 AlN 5 SiO2 31 AlN 5 Corning Eagle 2000 Substrate a nano multilayered sample configuration, where Figure 59b highlights a representative TEM cross-section of the as-sputtered AlN/SiO2 multilayers noting the clear interface between the layers. The accompanying selected area electron diffraction (SAED) pattern (Figure 59b) indicates nanocrystalline spots in the AlN layer and a diffused amorphous pattern in the SiO2 layer for an overall amorphous/crystalline structure. The experimental % transmittance plot, as presented in Figure 59c, confirms much higher optical transmittance in the MBI samples than the repeated bilayer samples. The corresponding top views of the as-sputtered samples (Figure 59d-g) indicate the high transparency in the UV/Vis region of the optically optimized samples. Moreover, the experimental transmittance of the MBI 10 Layer and 20 Layer samples are 94.0% and 95.3% from 86 300 to 800 nm, respectively. This is a marked increase as compared to the repeated bilayer samples, which both have experimental transmittance values below 80%. Figure 59: Overview of AlN/SiO2 nanomultilayers, where (a) is a schematic of alternating and (b) representative cross-sectional TEM nanocrystalline AlN and amorphous SiO2, as indicated with the inset SAED patterns. (c) Experimental % transmittance curves, where layer thicknesses of repeated bilayer samples are indicated in parentheses in nanometers. Top views of as-deposited samples (d-g): (d) AlN(100 nm)/SiO2(100 nm), (e) AlN(50 nm)/SiO2(50 nm), (f) MBI 10 Layers, and (g) MBI 20 Layers. 87 To investigate the repeated bilayer microstructures as a baseline and to examine the implications of varying layer thickness, an AlN(50)/SiO2(50) sample is displayed in Figure 60a. Uniform layers were observed through the entire film thickness with columnar grains extending through the 50 nm AlN layers, a common morphology observed in sputter-deposited nanocrystalline films. This columnar grain structure can affect optical behavior of dielectric materials, such as increasing spectral selectivity and refractive index anisotropy [252, 253]. The inset SAED pattern shows nanocrystalline rings with some preferential texturing as indicated by the brighter spots on the rings. At higher magnifications, Figure 60b features a crystalline- amorphous AlN-SiO2 interface with the highly columnar AlN grains. Within the layers, the amorphous nature of the SiO2 is confirmed by the lack of SiO2 peaks in the radial intensity profile (Figure 60c) which indicates hcp and fcc AlN peaks. Previous studies have found that a formation of mixed phase hexagonal (wurtzite) and cubic (rock salt and zinc blende) crystal structure AlN is possible by reactive sputtering, especially when deposited at low discharge powers, which were employed in this study [254]. Additionally, some studies have demonstrated that with increasing AlN layer thickness, a rock-salt (cubic) to hexagonal transformation or zinc blende (cubic) to hexagonal transformation is observed [255, 256]. Beyond layer microstructure, the role of the interface character must be examined through both interface roughness and composition. The amorphous-crystalline interface of AlN(50)/SiO2(50) is highlighted in Figure 60b, where there is a sharp interface with clear delineation between the AlN and SiO2 layers. Upon closer examination, there appears to be varying termination of AlN grains, suggesting non-isotropic columnar growth in the AlN layer a type of interfacial deviations which could affect multilayer response, such as increasing resistance to deformation [247]. 88 Figure 60: Overview of AlN(50 nm)/SiO2(50 nm) repeated bilayer sample, (a) BF cross-sectional TEM and inset SAED pattern of the multilayer film with yellow box indicating the region of (b) BF HRTEM of the AlN-SiO2 interface region. Red dotted line indicates the path of the 1D compositional line profile shown in (d). (c) Integrated radial intensity profiles interpolated from the SAED pattern and (d) the 1D compositional line profile as measured by STEM EDS, showing atomic % fraction of elements across the multilayers. Furthermore, the AlN-SiO2 interfaces are compositionally discrete, as evidenced by the EDS line profiles in Figure 60d where the average slope of the compositional gradient [225] is calculated from the Al line profile and is ≈3.48 atomic fraction/nm. Thus, the repeated AlN(50)/SiO2(50) configuration exhibited distinct layer interfaces in addition to a strong columnar AlN grain morphology. Increasing the bilayer thickness in a nanolayered structure has implications on the microstructure and interface character, and it is expected that the columnar grain size will scale with layer thickness [9], thus leading to an increase in the anisotropic columnar growth in the AlN layers. Therefore, a periodic AlN(100)/SiO2(100) multilayer was investigated as an additional 89 baseline configuration, and it presents uniform layers (Figure 61a) with more distinct bright spots in the inset SAED patterns. The AlN columnar grains observed in the HRTEM (Figure 61b) have higher grain widths than in the AlN(50)/SiO2(50) multilayers, increasing from around 10 nm to 20 nm. The same mixed crystalline phase AlN is observed (Figure 61c) with a higher peak-to-valley ratio, which indicates more preferential texturing in the increased AlN layer thickness. Additionally, the average slope of the EDS line profile decreases to ≈3.05 atomic fraction/nm, which suggests more mixing in the thicker bilayer sample. The repeated bilayer configurations limit the achievable maximum transmittance values, but in view of these results, the optically optimized samples allow further investigation of the nanoscale features that potentially lead to improved transmittance in non-constant bilayer Figure 61: Overview of AlN(100 nm)/SiO2(100 nm) repeated bilayer sample, (a) BF cross-sectional TEM with inset SAED pattern and (b) BF HRTEM of the AlN-SiO2 interface region. (c) Integrated radial intensity profiles interpolated from the SAED pattern and (d) the 1D compositional line profile as measured by STEM EDS, showing atomic % fraction of elements across the multilayers. 90 thickness samples. Figure 62 highlights the MBI 20 Layer sample with transmittance ≥ 95%, which has the same total number of layers and interfaces as the baseline AlN(50)/SiO2(50) sample. The cross-section TEM shown in Figure 62a suggests uniform layers with concentric diffused rings in the SAED pattern, indicating an overall amorphous morphology. HRTEM in Figure 62b features an AlN layer within the amorphous SiO2. At 5 nm, the AlN layers remained nanocrystalline, with a smaller grain size than in the repeated bilayer samples. However, the radial intensity plot in Figure 62c illustrates the loss of an observed mixed crystalline phase AlN structure due to the thin layers. In nitride/SiO2 NMs it has been reported that amorphous SiO2 can only crystallize and grow epitaxially with sufficiently small (SiO2 < 0.6 nm) layer thicknesses due to templating effects from the hcp nitride layers, but beyond this the layers break coherent growth and transform back into Figure 62: Overview of AlN/SiO2 MBI 20 Layer sample, (a) BF cross-sectional TEM with inset SAED pattern and (b) BF HRTEM of the AlN-SiO2 interface region. (c) Integrated radial intensity profiles interpolated from the SAED pattern and (d) the 1D compositional line profile. 91 the amorphous state [257]. This agrees with what was observed in the present study, where the amorphous structure in SiO2 was maintained across all samples due to the sufficiently thick SiO2 layers. Additionally, due to the thin AlN layers, the grain structure confined to the nanocrystalline region does not develop into a highly aligned columnar structure as other studies have reported [258]. As for the effect on the interface due to varying layer thickness, there is less evidence of variation in grain termination as a result of reduced columnar growth in these thinner AlN layers, leading to decreased grain morphology-induced structural roughness [247]. However, there is increased compositional mixing of the AlN and SiO2 layers at the interface, as evidenced by the slope of the EDS line plot in Figure 62d, where the average slope is ≈1.20 atomic fraction/nm, a notable decrease from the repeated bilayer samples. Furthermore, the O2 atomic fraction remains above ~30% within the AlN layers, where this trace oxygen from the SiO2 leading to formation of compositional transition regions has been observed in AlN deposited on amorphous SiO2 substrates [258]. This increased interdiffusion and transition region at the AlN-SiO2 interface could lead to interfacial scattering, but the agreement between the MBI calculations and experimental transmittance suggests a minimal effect on achieving high transparency. Decreasing the total number of layers in a constant overall thickness could have implications on optical properties such as decreasing the incidents of scattering at interfaces. Thus, an overview of the optically optimized MBI 10 Layer sample with half the total layers as the MBI 20 is presented in Figure 63. Comparing the MBI 10 Layer sample to the MBI 20 Layer, the cross- sectional TEM in Figure 63a reveals more uniform layers with an absence of the previously observed layer waviness and fewer nanocrystalline spots in the inset SAED pattern. Figure 63b presents sharp interfaces similar to the previous sample and ~5 nm nanocrystalline grains. The 92 microstructure within the AlN layers and at the AlN:SiO2 interface are consistent with what was observed in the MBI 20 Layer sample. However, as the total number of layers decrease, there is less interdiffusion at the AlN-SiO2 interface, which is confirmed by the EDS slope increasing to ≈1.67 atomic fraction/nm. The previous results confirm that as the layer thicknesses are tuned for maximized optical properties, there are changes in the layer structure and interface properties, as confirmed by deviations in crystallinity and interface discreteness. These fluctuations also lead to changes in multifunctional behavior and two methodologies were employed to verify these effects: residual stress analysis and nanoindentation measurements which are summarized in Table 4. From these Figure 63: Overview of AlN/SiO2 MBI 10 Layer sample, (a) BF cross-sectional TEM with inset SAED and (b) BF HRTEM of the AlN-SiO2 interface region. (c) Integrated radial intensity profiles interpolated from the SAED pattern and (d) the 1D compositional line profile. 93 results, it is evident that varying the layer thicknesses for optimized transmittance impacts the residual stresses in the films, with repeated bilayer films having moderate tensile residual stresses (~300-500 MPa). Dominating compressive residual stresses of the SiO2 layer and reduced AlN crystallinity lead to the shift from tensile to compressive stresses in the optically optimized samples [259]. Similarly, microstructural variations due to tuning bilayer thickness affects hardness of the multilayer films. The measured nanoindentation hardness values span from ~7 GPa to 9.8 GPa. For reference, AlN hardness is approximately 12 GPa [260] and SiO2 hardness is approximately 4.5-9.5 GPa [261, 262], depending on crystallographic structure. In the MBI 10 and 20 Layer samples, the fraction of SiO2 is much higher, leading to lower hardness values. Overall, the changes in multifunctional properties were measurable, and these differences are in part due to composition, but are also driven by changes in interfacial character. Table 4: Results of experimental mechanical and optical properties, including total film thickness, residual stresses as measured by profilometry, nanoindentation results, and experimental % transmittance. Sample Details Total Film Thickness Residual Stress (MPa) Nanoindentation Pred. %T 300-800 nm Exp. %T 300- 800 nm Hardness (GPa) Modulus (GPa) AlN(50)/SiO 2(50) 1.0 µm 342 9.8 ± 0.2 94.5 ± 1.0 51.1% 70.5% AlN(100)/SiO 2(100) 1.0 µm 538 10.7 ± 0.5 106.4 ± 2.7 54.6% 78.6% MBI 10 Layers 1.0 µm -261 7.9 ± 0.1 68.8 ± 0.7 96.1% 94.0% MBI 20 Layers 1.0 µm -249 8.2 ± 0.2 69.3 ± 1.3 96.2% 95.3% 5.4. Conclusion In summary, microstructural evaluation of AlN/SiO2 NMs have shown that optimizing layer configurations for high optical transparency (94-96%) in the UV/Vis region induces nanoscale variations within the layers and at the interfaces. Based on the current study, designing multilayer configurations for maximized transparency leads to changes in the microstructure such as the observed columnar to nanocrystalline grain morphology in the AlN layer as well as reduction 94 of sharp compositional transitions at the interfaces. Furthermore, nanoindentation and residual stress results for the varying NM configurations, highlight a correlation between optimization for high optical transparency and physical attributes. An understanding of these interfacial and microstructural changes in relation to optical behavior facilitates eventual introduction of additional functionalities to optical NMs. 95 Chapter 6 : High Transparency Ceramic Crystalline/Amorphous and Amorphous/Amorphous NMs The work in this chapter has been published as a journal article titled Synthesis and Characterization of Optically Transparent Crystalline/Amorphous and Amorphous/Amorphous Multilayers and is and is published in the journal Scripta Materialia in Volume 187, October 2020 issue, pages 157-162 (DOI: 10.106/j.scriptamat.2020.06.013). While optical ceramic/ceramic NMs are promising candidates for durable transparent multifunctional coatings, the contributions of interfaces and high transparency configurations, particularly across a range of material systems, have yet to be explored. To address this, high transparency NMs (%T380-1100nm≈94-99%) were synthesized, and the effects of crystalline/amorphous (AlN/SiO2, AlN/Al2O3) and amorphous/amorphous (TiO2/SiO2) interfaces were characterized by spectrophotometry, transmission electron microscopy, and nanoindentation. This work demonstrates that tuning layer configurations for improved transmittance resulted in substantial variations in microstructure and multifunctional film properties and further evaluates the structure-properties relationship in optical NMs in the context of interface character and layer microstructure. 6.1. Introduction As emphasized throughout this work, nano multilayered (NM) materials are highly tailorable for a wide range of desirable characteristics including, but not limited to, mechanical [2- 5], optical [6-8], electrical [9-11], and thermal properties [12-14]. As such, NM structures present attractive routes towards multifunctional nanomaterials since the combination of distinct components at the nanoscale enables the augmentation of two or more properties. Optical NMs are those capable of achieving tuned optical properties (e.g. maximizing or minimizing reflectance, 96 absorption, etc.), including high UV/Vis to NIR transmittance [263], with ceramic optical NMs achieving the highest broadband transmittance performance. However, maximizing transparency requires the presence of aperiodic layers, meaning the layer thicknesses and layer ratios vary throughout the film. This aperiodicity can lead to substantial microstructural variations [250, 264] which in turn can impact multilayer multifunctional properties, such as mechanical performance, an area that has remained relatively unexplored in optical NMs. Regarding transparent optical NMs, the layer thicknesses and material selection are both critical parameters for tuning transmittance [265], where those same variables have also been shown to impact the functional performance in non-optical NMs [5, 176]. Prevalent base materials in optical NMs include both amorphous (SiO2, TiO2, Al2O3) and crystalline (AlN, TiO2, MgF2) structures, thus, common multilayered configurations can include crystalline/amorphous (C/A) and amorphous/amorphous (A/A) interfaces. To date, only non-optical NMs with C/A and A/A interfaces have been examined as a function of layer thickness, showing high strength and large plastic strain behavior due to the interfaces acting as preferential sites for dislocation nucleation and emission in C/A multilayers [266] as well as shear band propagation obstruction in A/A multilayers [147]. The preceding works in non-optical NMs demonstrate that C/A and A/A interface character shows a strong size effect which is expected to also influence the optical behavior. However, the contribution of common types of interfaces in optical NMs has not been previously explored, particularly regarding aperiodicity, which is crucial for understanding the complex relationship between transparency and functionality. The present study provides, to the authors' knowledge, the first study on the effects of optical optimization and aperiodicity, along with the impact of C/A and A/A interfaces, on resultant microstructure and behavior in highly transparent NMs. 97 6.2. Experimental Methods Three systems of ceramic nanomultilayers capable of achieving high optical transmittance were explored, including AlN/SiO2, TiO2/SiO2, and AlN/Al2O3. The NM films (all 1 µm total thickness) were deposited on Corning Eagle 2000 glass substrates by DC/RF magnetron sputtering from 3.3 cm (1.3 in) sources under a base pressure of 1.1 x 10 -4 Pa. During the deposition, a bias voltage was applied to the source targets under an inert Ar working gas (6.7 x 10 -1 Pa ) to sputter individual layers of SiO2, TiO2, and Al2O3, each at 40 W of RF power. For the AlN layers, an Al source with 150 W of DC power was sputtered in the presence of inert Ar and 2.0 x 10 -1 Pa of reactive N2 working gas to deposit the AlN compound. A pneumatic shutter allowed for control of on-times of deposition, enabling fine control over the layer thicknesses and deposition from a single target at a time. Under these conditions at low deposition temperatures, the AlN layers were nanocrystalline and the SiO2, TiO2, and Al2O3 layers were amorphous. Two C/A systems (AlN/SiO2, AlN/Al2O3) and one A/A system (TiO2/SiO2) each had a common base component to another system which allowed for direct comparison across the three compositions. To calculate the residual stresses using Stoney’s equation, an XP-2 profilometer (AMBiOS) was used to measure surface profiles before and after deposition. In-line transmittance of the samples was measured using spectrophotometry (Cary UV/Vis 60 with a solid sample holder) from 200-1100 nm. Complex refractive indices were model fit from variable angle spectroscopic ellipsometry (VASE) (J.A. Woollam) measurements of the as-sputtered monolithic thin films. Cross-sectional transmission electron microscopy (TEM) samples were prepared using focused ion beam (FIB) lift-out techniques (Zeiss Auriga Dual Beam, Thermo Fischer Helios G4 DualBeam PFIB) and examined using a JEOL 2100F TEM. Nanoindentation was performed using a Hysitron 98 Triboindenter (Buehler) with a 100 nm Berkovich tip and a force-controlled, constant loading rate function. Over 100 indents were measured per sample for the presented hardness and reduced elastic moduli values. 6.3. Results and Discussion For each multilayer system, two configurations were synthesized: (1) the ‘repeated bilayer’ configuration with periodic layer thicknesses held constant, and (2) the ‘optically optimized’ configuration with aperiodic layer thicknesses calculated for maximum transmittance. Both repeated bilayer and optically optimized configurations across all C/A and A/A material systems consisted of 20 total layers in the films for a constant number of interfaces. The repeated bilayer systems serve as non-optimized, baseline samples but include the same type of C/A and A/A interfaces as the optically optimized samples. The periodic samples consisted of AlN(50 nm)/SiO2(50 nm), TiO2(50 nm)/SiO2(50 nm), and AlN(50 nm)/Al2O3(50 nm), with the repeated layer thicknesses indicated in parentheses. The ‘nm’ layer thickness designation will be dropped hereafter for concision. Optically optimized multilayers were calculated for maximum transmittance in the UV/Vis to NIR wavelengths (380-1100 nm) using the Multiple Beam Interference (MBI) Recursive method [33] in an in-house MATLAB optimization scheme as detailed in previous studies [250, 264]. The experimental spectroscopic refractive indices for AlN (n550 nm ≈ 2.18), SiO2 (n550 nm ≈ 1.43), TiO2 (n550 nm ≈ 2.58), and Al2O3 (n550 nm ≈ 1.77) were used in the MBI calculations. All optically optimized configurations required non-constant bilayer thicknesses and the layer thicknesses synthesized are detailed in Tables 5-7. The optically optimized samples synthesized for maximized optical transmittance were designated “AlN/SiO2 MBI”, “AlN/Al2O3 MBI”, and “TiO2/SiO2 MBI” indicating the repeated base constituent layers repeated throughout the 1 µm total film thickness. Optimized layer configurations were calculated 99 for maximized optical transmittance in the UV/Vis wavelengths (380-1100 nm) for each system using the Multiple Beam Interference (MBI) Recursive method [33, 96]. MBI formulas were implemented in an in-house MATLAB optimization scheme [250] with experimental optical constants measured from the as-sputtered monolithic films. Please note that throughout this manuscript, ‘repeated bilayer’ and ‘periodic’ will be used interchangeably to identify the non- optimized multilayers. Additionally, the terms ‘optically optimized’, ‘MBI’, and ‘aperiodic’ are equivalent to each other and will be used interchangeably. Table 5: Summary of layer thicknesses in AlN/SiO2 MBI 20 Layer sample AlN/SiO2 20 Layers Material Thickness (nm) SiO2 128 AlN 5 SiO2 226 AlN 5 SiO2 29 AlN 5 SiO2 22 AlN 5 SiO2 19 AlN 5 SiO2 27 AlN 5 SiO2 224 AlN 5 SiO2 34 AlN 5 SiO2 211 AlN 5 SiO2 31 AlN 5 Corning Eagle 2000 100 Table 6: Summary of layer thicknesses in AlN/Al2O3 MBI 20 Layer sample AlN/Al2O3 Sample Material Thickness (nm) Al2O3 65 AlN 6 Al2O3 10 AlN 6 Al2O3 6 AlN 6 Al2O3 5 AlN 175 Al2O3 6 AlN 34 Al2O3 7 AlN 38 Al2O3 5 AlN 197 Al2O3 15 AlN 12 Al2O3 199 AlN 5 Al2O3 197 AlN 5 Corning Eagle 2000 101 Table 7: Summary of layer thicknesses in TiO2/SiO2 MBI 20 Layer sample TiO2/SiO2 20 Layers Material Thickness (nm) SiO2 149 TiO2 5 SiO2 34 TiO2 5 SiO2 133 TiO2 5 SiO2 38 TiO2 5 SiO2 148 TiO2 5 SiO2 137 TiO2 5 SiO2 150 TiO2 5 SiO2 120 TiO2 5 SiO2 20 TiO2 5 SiO2 20 TiO2 5 Corning Eagle 2000 Figure 64 shows the experimental % transmittance spectra of the as-sputtered multilayer and confirms that the repeated bilayer curves (Figure 64a) exhibit the expected lower broadband transmittance and the optically optimized curves (Figure 64b) exhibit high optical transmittance. Moreover, all optically optimized multilayers achieve an average %T380-1100nm > 94% and demonstrate good agreement (within 3% of predicted) between the MBI-calculated predictions and experimental measurements, as summarized in Table 8. The corresponding top views of the as-deposited samples (Figure 64c-h) confirm the improved optical transparency of the MBI- calculated layer configurations in contrast to the repeated bilayer samples. 102 Figure 64: Overview of ceramic nanomultilayer systems, including AlN/SiO2, TiO2/SiO2 and AlN/Al2O3, where (a) is experimental % transmittance curves of repeated bilayer samples with layer thicknesses indicated in parenthesis in nanometers and (b) experimental % transmittance curves of the MBI optically optimized samples. Top views of as- deposited samples (b-h): repeated bilayer (c) AlN(50)/SiO2(50), (d) AlN(50)/Al2O3(50), and (e) TiO2(50)/SiO2(50) and optically optimized MBI samples (f) AlN/SiO2 MBI, (g) AlN/Al2O3 MBI, and (h) TiO2/SiO2 MBI. Table 8: Experimental results of mechanical and optical properties Sample Details Residual Stress (MPa) Nanoindentation Predicted Ave %T 380-1100nm Experimental Ave %T 380-1100nm Hardness (GPa) Modulus (GPa) AlN(50)/SiO 2(50) 342 9.8 ± 0.2 94.5 ± 1.0 70.2% 51.1% AlN/SiO 2 MBI -249 8.2 ± 0.2 69.3 ± 1.3 95.3% 98.7% AlN(50)/Al 2O 3(50) 381 13.9 ± 0.2 139.3 ± 1.6 70.9% 63.3% AlN/Al 2O 3 MBI 414 15.5 ± 0.2 148.9 ± 1.3 93.2% 93.8% TiO 2(50)/SiO 2(50) -302 7.1 ± 0.1 75.6 ± 0.6 64.8% 44.3% TiO 2/SiO 2 MBI -250 7.1 ± 0.1 76.7 ± 0.4 94.4% 94.7% 103 6.3.1. Microstructural Overview Microstructural characterization was performed to investigate the effect of varying layer thicknesses from a periodic configuration to an aperiodic multilayer stack as well as the crystalline/amorphous and amorphous/amorphous interfaces. Figure 65 presents the bright field (BF) TEM cross-sectional overviews of each repeated bilayer (Figure 65a-c) and MBI (Figure 65d-f) configuration. Each overview includes an inset selected area diffraction pattern (SADP) with the growth direction of the film indicated by a red arrow. The cross-sections show that the AlN/SiO2 and AlN/Al2O3 NMs have crystalline/amorphous alternating layers and the TiO2/SiO2 NMs have amorphous/amorphous layers. Figure 65a and d present the repeated and MBI AlN/SiO2 system, noting a C/A configuration. For the periodic AlN(50)/SiO2(50), uniform layers are maintained throughout the film thickness, with columnar grains spanning the 50 nm AlN layers. This columnar AlN grain morphology is commonly observed in sputter-deposited nanocrystalline films [21, 84] and is confirmed by the inset SADP pattern where there is evidence of some preferential texturing as indicated by the brighter spots on the nanocrystalline rings. The nanocrystalline AlN layers are a mixed-phase hexagonal (wurtzite) and cubic (rock salt and zinc blende) crystal structure [264] due to the low discharge powers and low deposition temperature during DC reactive magnetron sputtering [256, 267, 268]. This preferential columnar grain structure can lead to higher resistance to deformation as well as increased spectra selectivity and optical anisotropy [252, 253]. In contrast, the MBI sample, with a high experimental transmittance of %T380-1100nm > 98%, has significant aperiodicity of the bilayers, as shown in the overall TEM cross-section in Figure 65d. The BF cross-section indicates uniform layers but due to a larger volume fraction of amorphous SiO2, the multilayer SADP shows more diffused concentric rings. 104 The AlN layers have a nanocrystalline morphology typical of the early stages of film growth rather than the highly aligned columnar grains seen in thicker layers [258]. To further expand on the formation and characterization of C/A interfaces and aperiodicity, Figure 65b and e present a periodic and an optically optimized multilayer of AlN/Al2O3. The BF cross-sections of both AlN/Al2O3 multilayers show uniform layers with clear interfaces between the crystalline and amorphous components. The AlN layers retain a columnar structure as seen in AlN(50)/SiO2(50) (Figure 65a), and the inset SADPs for both AlN/Al2O3 samples suggest a similar microstructure with preferential texturing. However, a transition region within the AlN layers, indicated by yellow lines in Figure 65b and e, is observed in both AlN/Al2O3 multilayers and is characterized by an equiaxed nanocrystalline AlN structure between the amorphous interface and columnar grain structure. The presence of this transition region could have implications on multilayer performance such as changes in index of refraction due to layer microstructure [209, 269] or variations in deformation behavior [270]. Overall the two MBI samples with C/A interfaces have sharp, uniform layers but dissimilarities emerge in the AlN columnar structure where thicker layers in the AlN/Al2O3 system allow for the growth and development columnar grains, rather than the equiaxed nanocrystalline structures seen in the AlN/SiO2 MBI layers. With regards to the optical performance, the % transmittance is higher for AlN/SiO2 MBI than the AlN/Al2O3 counterpart by 5% which is expected due to the combined effects of the lower SiO2 index of refraction and contributions of the equiaxed AlN microstructure. Discussion of the contribution of the interfaces will be expanded with Figure 66. Optically optimizing an amorphous/amorphous configuration is expected to alter microstructural characteristics and to investigate this the TiO2/SiO2 multilayer system is evaluated, with its respective microstructures presented in Figure 65c and f. The repeated bilayer arrangement 105 Figure 65: Representative BF TEM cross-sections of multilayers, with repeated bilayer configurations on the left side (a-c) and optically optimized MBI (d-f) configurations on the right side. The corresponding compositions are labeled on the left: (a,d) AlN/SiO2, (b,e) AlN/Al2O3, and (c,f) TiO2/SiO2. Inset SAED patterns show the crystalline or amorphous nature of the samples and the red dotted arrows indicate the growth direction of the films. 106 TiO2(50)/SiO2(50) (Figure 65c) indicates the presence of well-defined amorphous/amorphous layers, as confirmed by the changes in chemical contrast between layers and the diffused halo pattern in the inset SADP. Introducing aperiodic layer thicknesses for high transparency in the MBI sample (Figure 65f) maintains the overall amorphous nature of the microstructure (see inset). However, an increased ‘waviness’ is observed for layers with increasing film thickness along the growth direction (red arrow) but is minimal as compared to previous studies on line-of-sight deposition of multilayers [271, 272]. Such microstructural features could potentially lead to enhanced scattering of light at the interfaces, but the experimental transmittance (%T380-1100 nm = 94.7%) confirms that these layer structures yield high optical performance. 6.3.2. High Resolution Interface Evaluation Figure 66 provides a detailed look at the interfaces across all samples via high-resolution TEM (HRTEM) micrographs. Figure 66a-c highlights the periodic samples (left) and Figure 66 d- f highlight the optically optimized aperiodic samples (right). The HRTEM of the C/A interfaces (AlN/SiO2, AlN/Al2O3) in Figure 66a,b,d,e reveal clear boundaries between the crystalline and amorphous layers, where these compositionally sharp interfaces with disparate lattice structures can lead to a strong capacity for absorbing dislocations [266]. All C/A interfaces exhibit an interlayer region where the crystalline AlN grains terminate at the amorphous (SiO2, Al2O3) layer, as indicated with white lines in the HRTEM. For the repeated bilayer samples shown in Figure 66a and b, the interfacial regions are ~3-4 nm in thickness and exhibit columnar grain termination at the C-A interface. Meanwhile, the optically optimized AlN/SiO2 and AlN/Al2O3 interfaces (Figure 66d,e) appear smooth and show a ~2 nm thick interfacial region. The presence of these interlayers is in agreement with previous studies on non-optical C/A interfaces, which reported similar thin (~1-2 nm) interfacial regions with semi-crystalline structure [30]. 107 Figure 66: Representative BF HRTEM cross-sections of multilayers, with repeated bilayer configurations on the left side (a-c) and optically optimized MBI (d-f) configurations on the right side. The corresponding compositions are labeled on the left: (a,d) AlN/SiO2, (b,e) AlN/Al2O3, where interlayer regions are indicated by white lines, and (c,f) TiO2/SiO2, where nanocrystalline regions are indicated by the yellow ellipses with the corresponding inset FFT. 108 In the A/A system, TiO2/SiO2, HRTEM indicates layers with predominantly amorphous microstructures in both repeated bilayer and optically optimized samples. The interfaces are well defined and close inspection reveals ordered nanocrystalline regions (indicated in yellow) embedded at the interface, as shown in TiO2(50)/SiO2(50) in Figure 66c, where the inset FFT confirms the presence of crystallites along the entire interface. Similar nanocrystalline regions have been observed in amorphous/amorphous structures synthesized by room temperature sputtering [42, 273-275], however, in the current study, the crystallization was restricted to the interface rather than being embedded within the TiO2 layer matrix as previously reported [273]. In the TiO2/SiO2 MBI sample, fewer nanocrystallites formed along the interface, which is likely due to the thinner TiO2 layers [43, 276]. 6.3.3. Mechanical Assessment The aforementioned changes in layer structure and interface characteristics as a function of tuned layer thicknesses can generate shifts in other facets of multilayer performance. To evaluate the potential impact of microstructural variations on multifunctionality in optical NMs, two techniques, residual stress analysis and nanoindentation measurements, were performed and results are summarized in Table 8. In the TiO2/SiO2 (compressive) and AlN/Al2O3 (tensile) systems, varying the layer thicknesses for optimized transmittance does not lead to a noticeable change in residual stresses. However, in AlN/SiO2, there is a shift from tensile to compressive residual stresses in the optically optimized samples that is ascribed to a build-up of compressive stresses caused by the thicker SiO2 layers. These changes in interfacial character and layer ratios can also impact film hardness, as is the case in the C/A systems, although the effect varies with composition. For instance, in the AlN/SiO2 system there is a decrease in hardness with optical optimization, which can be attributed to the increased amorphous SiO2 fraction. Meanwhile, in the 109 AlN/Al2O3 system, the amorphous volume fraction remained relatively constant, indicating that the rise in hardness is predominantly attributed to the microstructural variations. The C/A system observations are in contrast with the A/A system, where optical optimization did not show an effect on hardness, which remained at 7.1 GPa in both the repeated bilayer and optically optimized arrangements despite considerable changes in layer thickness and composition. Overall, introducing aperiodicity for optical optimization in C/A multilayers resulted in microstructural variations which in turn led to measurable changes in nanoindentation hardness. Conversely, aperiodic layers for improved transmittance in A/A multilayers did not show any effect on the nanoindentation hardness. To visualize the current observations, the contributions of interfaces and aperiodicity in optimized optical NMs are presented in Figure 67, where the trends between mechanical and optical performance are compared. Two plots are presented to illustrate the effect of varying multilayer parameters such as amorphous volume fraction and the average number of interfaces on experimental % transmittance and nanoindentation hardness measurements in a variety of material systems from present work and related studies [7, 33, 98]. When the amorphous volume fraction in a multilayer structure increases, the potential for higher experimental transmittance increases, as shown in Figure 67a. In these systems, TiO2, SiO2, and Al2O3 are all high transparency materials with an amorphous as-sputtered microstructure. Incorporating a higher fraction of these constituents leads to higher average transmittance, however, this is at the expense of overall multilayer hardness because most amorphous components have lower associated hardness values [261, 262, 277, 278]. As the average number of interfaces increases from ~4/µm to 20/µm, the results suggest that there is an increase in achievable average % transmittance but no clear effect on the nanoindentation hardness of the films. The traditional increase in hardness 110 with decreasing bilayer thicknesses is expected in ceramic NMs [122, 164], and optical NM systems with a higher density of interfaces/µm should be examined to elucidate the potential tradeoff between improved hardness and lower achievable transmittance due to scattering at interfaces. Figure 67: Plots of multifunctional properties of optically optimized samples as a function of multilayer characteristics including (a) experimental hardness and transmittance as a function of amorphous volume fraction and (b) as a function of the normalized # of interfaces per µm. 6.4. Outlook and Conclusion In this work, we explored the effect of maximizing optical properties on microstructural and multifunctional properties in both crystalline/amorphous and amorphous/amorphous ceramic optical NM systems capable of achieving high transparency. Optically optimized configurations across three ceramic nanomultilayer compositions demonstrated high experimental transmittance ranging from 94-99% at the 380-1100 nm wavelengths. Tuning layer thicknesses and introducing aperiodicity for optimized transmittance led to microstructural variations both within the layers and at the interfaces across all systems. Tuning layer thicknesses for higher optical performance in the C/A multilayers produced measurable changes in residual stresses and nanoindentation hardness. This is in contrast with the A/A TiO2/SiO2 system, where changes in crystalline features at the amorphous-amorphous interface were observed as a function of optimizing for high transmittance, but these variations did not give rise to fluctuations in nanoindentation hardness. This study demonstrates that incorporating aperiodicity for improved transmittance greatly 111 influences the microstructural and interfacial characteristics of C/A and A/A multilayers and highlights a versatile route for exploring the relationship between optics and multifunctionality in ceramic optical NMs. 112 Chapter 7 : Conclusions and Future Work 7.1. Conclusions Optical nanomultilayers (NMs) are promising candidates for materials that are simultaneously transparent and durable. However, additional understanding of the microstructural variations and deformation behavior as a function of optical optimization is needed before fully realizing their multifunctional potential. Thus, the work presented in Chapters 4-6 explored the role of microstructure, interfaces, and composition in optically optimized NMs and related these characteristics to multilayer optical and mechanical behavior as a crucial step towards discerning, then leveraging, this interplay of properties. In the first part of this study, a metal/ceramic AlN/Ag optical NM system was investigated to explore the relationship between transmittance and deformation behavior. It was revealed that optically optimized samples exhibited a significant increase in hardness in conjunction with improved optical performance over repeated bilayer samples. On the other hand, fracture behavior via Vickers indentation was similar across all samples. Furthermore, this work demonstrated the success of a combinatorial approach of synthesis, optical predictions, and mechanical behavior investigations in optical NMs. From there, this approach was extended to ceramic optical NMs capable of achieving high (>95%) transmittance, which elucidated the role of microstructural variations as a function of introducing aperiodic layer thicknesses for improved optical behavior. This was achieved by synthesizing AlN/SiO2 NMs with repeated bilayer and optically optimized layer thicknesses across a range of total interfaces. It was found that, in addition to achieving high experimental transparency (%T200-1100nm=94-96%), substantial shifts in grain morphology and compositional transitions occurred when designing configurations for maximized transmittance. Furthermore, 113 nanoindentation and residual stress results for the varying NM configurations highlight a correlation between high optical transparency configurations and the observed physical attributes. Such correlations laid the groundwork for an understanding of these interfacial and microstructural changes concerning optical behavior to facilitate the introduction of additional functionalities to optical NMs. However, global insight into optical and mechanical properties across multiple material compositions is necessary to elucidate the dominating parameters in optical NM functionality. A first undertaking of this was conducted in the final study presented in Chapter 6, which explored the effect of maximizing optical properties on microstructural and multifunctional properties across a range of common crystalline/amorphous (C/A) and amorphous/amorphous (A/A) optical NM systems. The three ceramic nanomultilayer compositions (AlN/SiO2, AlN/Al2O3, TiO2/SiO2) demonstrated high experimental transmittance ranging from 94-99% and verified that tuned layer thicknesses and aperiodicity generated microstructural variations both within the layers and at the interfaces across all material systems. Moreover, optically optimized designs in C/A (AlN/SiO2, AlN/Al2O3) multilayers led to deviations within each system in the residual stresses and nanoindentation hardness that was not observed in the A/A (TiO2/SiO2) multilayers. By capturing the changes that arise with optimized optical properties, these shifts can be correlated to other multilayer properties. Overall, these studies showed: 1) the success in predicting and synthesizing optical NMs with high transmittance in a variety of material systems via DC/RF magnetron sputtering, 2) the importance of examining microstructural variations both within nanoscale layers and at the interfaces as a function of aperiodicity for improved transmittance, and 3) a new avenue for exploring the relationship between optics and multifunctionality in optical NMs. Through these observations, the main characteristics and microstructural transitions in relation to optical 114 performance were identified and the resultant film properties were reported. These findings lay the groundwork for elucidating how transparent optical multilayer systems can be tailored with optimal layer thicknesses and the effect this has on multifunctional performance, all in the effort to enable the design and synthesis of new high-performance, long-lasting optical materials. 7.2. Future Work The studies to date have addressed multiple questions around understanding the effect of optically optimizing NMs, but work remains to explain the opto-mechanical relationship in optical NMs. Promising future routes building off of this work are twofold: (1) to expand the scope of deformation modes to different nanomechanical loading modes and (2) to further isolate the contributions of interfaces in optical NMs. As previously discussed, nanoindentation testing provides insight into the hardness and modulus of nanostructured materials and can be considered a type of compressive loading configuration. To perform a comprehensive study on the deformation behavior of optical NMs, it is important to test the NMs under a variety of loading conditions. Due to the size constraints of these samples, the tests must be specifically designed for micro and nanoscale testing. As such, tensile and fracture toughness tests for nanostructured materials are proposed to build upon results to date on the influence of surfaces, interfacial effects, and nanoscale microstructures in these optical NMs. The NMs will likely exhibit mechanical properties and deformation that are significantly different from bulk materials due to the influence of surfaces, interfacial volume, and nanoscale microstructure. Microtensile testing and post-mortem characterization allows for the study of localized deformation and can provide further insight into the combined effect of layer thickness, interfacial properties, and composition. Fracture toughness is another critical parameter needed for a comprehensive understanding of the mechanical behavior of these systems and will 115 be estimated using micropillar compression testing. Ceramic/ceramic NMs are predicted to show improved toughness through crack absorption and deflection, and these are important behaviors to examine. By combining nanoindentation, microtensile experiments, and microcompression testing, the distinct loading configurations will provide an outlook on the driving deformation mechanisms for these types of NMs. Additionally, there are still several challenges surrounding ceramic/ceramic NMs, as little is understood about the effects of thickness and composition on strength and toughness. The nanoindentation testing and residual stress analysis detailed in this work have provided important insight into mechanical properties of the optical NMs, but expanding the mechanical testing will further elucidate the controlling mechanisms for NMs that yield both high optical performance and improved mechanical behavior. The studies detailed in this dissertation performed optimization and testing across multiple material systems to determine the effect of different interface types. 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Table 9: Key for Types of Characterization Techniques Identifier Characterization Technique Type of Technique P Profilometry Residual Stress X X-Ray Diffraction (XRD) Crystallographic S Scanning Electron Microscopy (SEM) Microscopy F Focused Ion Beam (FIB) T Transmission Electron Microscopy (TEM) ST Scanning Transmission Electron Microscopy (STEM) Sp Spectrophotometry Optical E Ellipsometry V Vickers Mechanical N Nanoindentation 132 Table 10: Summary of Types of Sputtering Used for Thin Films Synthesized Type of Film Reactive/Non- Reactive DC/RF Targets Used Ag Non-Reactive DC 1.3” Ag 99.99% Al2O3 Non-Reactive RF 1.3” Al2O3 99.99% Al2O3 (KIT) Reactive DC 2” Al 99.999% AlN Reactive (N2) DC 1.3” Al 99.9999% AlN (KIT) Reactive (N2) RF 2” Al 99.999% ITO (KIT) Non-Reactive RF 2” Indium Tin Oxide (ITO) 99.99% SiO2 Non-Reactive RF 1.3” SiO2 99.995% SiOx (KIT) Reactive (O2) RF 2” Si 99.999% TiN Reactive (N2) DC 1.3” Ti 99.995% TiO2 Non-Reactive RF 1.3” TiO2 99.9% TiOx (KIT) Reactive (O2) RF 2” Ti 99.995% ZnO Non-Reactive RF 1.3” ZnO 99.99% ZnO (KIT) Non-Reactive RF 2” ZnO 99.99% Table 11: Ag Sputtered Samples via Non-Reactive DC Sputtering Sample Name Working Pressure (mTorr) Power Density (W/cm 2 ) Thickness (nm) Time (min) Deposition Rate (nm/min) Characterization/Comments Ag-SR-1a & 1b 3.8 11.7 2143 30 71.4 P, X Ag-SR-2a & 2b 3.8 11.7 1984 30 66.1 P, X Ag-SR-3a & 3b 7.9 11.7 1673 30 55.8 P, X Ag-SR-4a & 4b 7.9 11.7 59.8 1 59.8 P, E Ag-SR-5a & 5b 8.0 11.7 5.0 0.08 60.0 P Table 12: Al2O3 Sputtered Samples via Non-Reactive RF Sputtering Sample Name Working Pressure (mTorr) Power Density (W/cm 2 ) Thickness (nm) Time (min) Deposition Rate (nm/min) Characterization/Comments Al2O3-SR-1a & 1b 5.0 4.7 37 60 0.62 P, X Al2O3-SR-2a & 2b 5.0 4.7 4 6.5 0.62 P Table 13: Al2O3 Sputtered Samples via Reactive RF Sputtered (KIT) Sample Name Working Gas Flow (sccm) Power Density (W/cm 2 ) Thickness (nm) Time (min) Deposition Rate (nm/min) Characterization/Comments Al2O3-SR-KIT-1 Ar 46 O2 4 4.9 125 60 2.1 P, X 133 Table 14: AlN Sputtered Samples via Reactive DC Sputtering Sample Name Working Gas Flow (sccm) Power Density (W/cm 2 ) Thickness (nm) Time (min) Deposition Rate (nm/min) Characterization/ Comments AlN-SR-1 & 2 Ar 25 N2 7.5 8.8 141 30 4.7 P, X AlN-SR-3 & 4 Ar 25 N2 7.5 17.5 190 30 6.3 P, X, Sp AlN-SR-5 & 6 Ar 37.5 N2 7.5 17.5 - 30 - P, X, Sp AlN-SR-7 & 8 Ar 25 N2 7.5 23.4 - 21.7 - P, X AlN-SR-9a & 9b Ar 20 N2 12 17.5 165 30 5.5 P, X AlN-SR-10a & 10b Ar 20 N2 12 17.5 180 30 6.0 P, X AlN-SR-11a & 11b Ar 25 N2 5 17.5 395 30 13.2 P, X AlN-SR-12a & 12b Ar 25 N2 5 17.5 340 30 11.3 P, X AlN-SR-13a & 13b Ar 37.5 N2 12.5 17.5 1000 150 6.7 P AlN-SR-14a & 14b Ar 37.5 N2 12.5 17.5 100 15 6.7 P, X AlN-SR-15a & 15b Ar 37.5 N2 12.5 17.5 112 15 7.5 P AlN-SR-16a &16b Ar 37.5 N2 12.5 17.5 - 60 - - AlN-SR-17a & 17b Ar 37.5 N2 12.5 17.5 535 60 8.9 P Table 15: AlN Sputtered Samples via Reactive RF Sputtering (KIT) Sample Name Working Gas Flow (sccm) Power Density (W/cm 2 ) Thickness (nm) Time (min) Deposition Rate (nm/min) Characterization/Comments AlNRF-SR-1 Ar 50 N2 4 4.9 339 60 5.7 P, X AlNRF-SR-2 Ar 40 N2 10 4.9 155 60 2.6 P, X AlNRF-SR-3 Ar 40 N2 10 7.4 264 55 4.8 P, X, E AlNRF-SR-4 Ar 40 N2 10 7.4 - 60 - - AlNRF-SR-5 Ar 40 N2 10 7.4 246 60 4.1 P 134 Table 16: ITO Sputtered Samples via Non-Reactive RF Sputtering (KIT) Sample Name Working Gas Flow (sccm) Power Density (W/cm 2 ) Thickness (nm) Time (min) Deposition Rate (nm/min) Characterization/Comments ITO-SR-1 Ar 20 2.5 417 50 7.0 P, X, E ITO-SR-2 Ar 20 3.0 109 14 7.9 P, X Table 17: SiO2 Sputtered Samples via Non-Reactive RF Sputtering Sample Name Working Pressure (mTorr) Power Density (W/cm 2 ) Thickness (nm) Time (min) Deposition Rate (nm/min) Characterization/Comments SiO2-SR-1a & 1b 5 3.5 49 60 0.8 P, X SiO2-SR-2a & 2b 5 4.7 57 60 0.9 P, X, E SiO2-SR-3a & 3b 3 4.7 - 60 - X, E SiO2-SR-4a & 4b 5 4.7 - 60 - X Table 18: SiOx Sputtered Samples via Reactive RF Sputtering (KIT) Sample Name Working Gas Flow (sccm) Power Density (W/cm 2 ) Thickness (nm) Time (min) Deposition Rate (nm/min) Characterization/Comments SiOx-SR-1 Ar 50 O2 4 2.5 - 60 - P, X SiOx-SR-2 Ar 50 O2 2 2.5 - 60 - P, X SiOx-SR-3 Ar 50 O2 4 1.5 22 60 0.4 P, X SiOx-SR-4 Ar 50 O2 4 2.5 66 60 1.1 P, X, E SiOx-SR-5 Ar 75 O2 4 2.5 49 60 0.8 X SiOx-SR-6 Ar 35 O2 4 2.5 57.8 60 1.0 P, X SiOx-SR-7 Ar 50 O2 4 2.5 55 60 0.9 P 135 Table 19: TiN Sputtered Samples via Reactive DC Sputtering Sample Name Working Gas Flow (sccm) Power Density (W/cm 2 ) Thickness (nm) Time (min) Deposition Rate (nm/min) Characterization/ Comments TiN-SR-1 & 2 Ar 50 N2 4.5 7.0 136 60 4.5 P, X, Sp TiN-SR-3 & 4 Ar 50 N2 6 7.0 - 60 - P, X, Sp TiN-SR-5 & 6 Ar 25 N2 4.5 14.0 132 120 4.4 P, X TiN-7 & 8 Ar 60 N2 4.5 3.5 - 30 - P, X TiN-SR-9a & 9b Ar 37.5 N2 12.5 7.6 80 60 1.3 P TiN-SR-10a & 10b Ar 37.5 N2 4 7.6 - 60 - - Table 20: TiO2 Sputtered Samples via Non-Reactive RF Sputtering Sample Name Working Pressure (mTorr) Power Density (W/cm 2 ) Thickness (nm) Time (min) Deposition Rate (nm/min) Characterization/Comments TiO2-SR-1a & 1b 5.0 4.7 14 30 0.5 P TiO2-SR-2a & 2b 5.0 4.7 34 60 0.6 P, E Table 21: TiOx Sputtered Samples via Reactive RF Sputtering (KIT) Sample Name Working Gas Flow (sccm) Power Density (W/cm 2 ) Thickness (nm) Time (min) Deposition Rate (nm/min) Characterization/Comments TiOx-SR-1 Ar 50 O2 4 3.7 60 P, X, E TiOx-SR-2 Ar 50 O2 4 4.4 60 P, X TiOx-SR-3 Ar 50 O2 2 3.7 26 60 0.4 P, X TiOx-SR-4 Ar 50 O2 4 3.7 12 60 0.2 P Table 22: ZnO Sputtered Samples via Non-Reactive RF Sputtering (KIT) Sample Name Working Gas Flow (sccm) Power Density (W/cm 2 ) Thickness (nm) Time (min) Deposition Rate (nm/min) Characterization/Comments ZnO-SR-1 Ar 20 2.5 311 60 5.2 P, X, E ZnO-SR-2 Ar 20 3.0 368 60 6.1 P, X 136 Appendix B: Summary of Sputtered Multilayers Table 23: Sputtered AlN/Ag Multilayer Samples AlN Layers Ag Layers Sample Name Type of Sample Thickness (nm) Number of Layers Thickness (nm) Number of Layers Total Thickness (nm) Characterization AlN/Ag-1 & 2 Repeated Bilayer 167 6 10 5 1052 P, X, S, F, T, Sp, E, V, N AlN/Ag-3 & 4 Repeated Bilayer 167 6 10 5 1052 P, X, S, Sp, N AlN/Ag-5a & 5b Repeated Bilayer 50 10 50 10 1000 P, X, S, F, T, Sp, E, V, N AlN/Ag-6a & 6b Repeated Bilayer 100 6 100 5 1100 P, S, Sp, N AlN/Ag-7a & 7b Repeated Bilayer 100 5 100 5 1000 Failed P AlN/Ag-8a & 8b Repeated Bilayer 50 11 50 10 1050 P, S, Sp, N AlN/Ag-9a & 9b MBI Test Varies 3 Varies 2 166 P, S, Sp AlN/Ag-10a & 10b Repeated Bilayer 20 26 20 25 1020 P, X, S, F, T, Sp, E, V, N AlN/Ag-11a & 11b MBI 9 Layers Varies 5 Varies 4 1053 P, S, Sp, N AlN/Ag-12a & 12b MBI 25 Layers Varies 13 Varies 12 929 P, S, Sp AlN/Ag-13a & 13b MBI 9 Layers #2 Varies 5 Varies 4 1054 P, Sp AlN/Ag-14a & 14b MBI 5 Layers Varies 3 Varies 2 1322 P, Sp, E AlN/Ag-15a & 15b MBI 3 Layers Varies 2 5 1 1047 P, Sp, E AlN/Ag-16a & 16b MBI 5 Layers #2 Varies 3 5 2 1125 P, Sp AlN/Ag-17a & 17b MBI 7 Layers Varies 4 7 3 1137 P, Sp AlN/Ag-18a & 18b MBI 5 Layers #2 Varies 3 5 2 1125 P, X, S, F, T, ST, Sp, E, V, N AlN/Ag-19a & 19b MBI 9 Layers Varies 5 5 4 1054 P, Sp AlN/Ag-20a & 20b MBI 5 Layers #3 Varies 3 5 2 1122 P, X, S, F, T, ST, Sp, E, V, N AlN/Ag-21a & 21b Repeated Bilayer 100 6 100 5 1100 P, X, S, Sp AlN/Ag-22a & 22b MBI 5 Layers #4 Varies 3 Varies 2 1501 Failed P, X, Sp AlN/Ag-23a & 23b MBI 5 Layers #4 Varies 3 Varies 2 1501 P, X, S, Sp, N AlN/Ag-24a & 24b Repeated Bilayer 50 11 50 10 1050 Extra Sample P, N AlN/Ag-25a & 25b Repeated Bilayer 20 26 20 25 1020 Extra Sample P, N AlN/Ag-26a & 26b Repeated Bilayer 167 6 10 5 1052 Extra Sample P, N AlN/Ag-27a & 27b MBI 5 Layers #2 Varies 3 Varies 2 1322 Extra Sample AlN/Ag-28a & 28b MBI 5 Layers #3 Varies 3 Varies 2 1322 Extra Sample 137 Table 24: Sputtered AlN/Al2O3 Multilayer Samples TiO2 Layers Al2O3 Layers Sample Name Type of Sample Thickness (nm) Number of Layers Thickness (nm) Number of Layers Total Thickness (nm) Characterization AlN/Al2O3-1a & 1b MBI 20 Layers Varies 10 50 10 1000 P, F, T, ST, Sp, N AlN/Al2O3-2a & 2b Repeated Bilayer 50 10 Varies 10 1000 P, F, T, ST, Sp, N Table 25: Sputtered AlN/SiO2 Multilayer Samples AlN Layers SiO2 Layers Sample Name Type of Sample Thickness (nm) Number of Layers Thickness (nm) Number of Layers Total Thickness (nm) Characterization AlN/SiO2-1a & 1b Repeated Bilayer 50 10 50 10 1000 P, X, F, T, ST, Sp, E, V, N AlN/SiO2-2a & 2b Repeated Bilayer 50 10 50 10 1000 Test Sample P, Sp AlN/SiO2-3a & 3b Repeated Bilayer 50 10 50 10 1000 Test Sample P, Sp AlN/SiO2-4a & 4b Repeated Bilayer 50 10 50 10 1000 Failed P, Sp AlN/SiO2-5a & 5b Repeated Bilayer 50 10 50 10 1000 Failed P, Sp AlN/SiO2-6a & 6b Repeated Bilayer 100 5 100 5 1000 P, X, Sp, V, N AlN/SiO2-7a & 7b Repeated Bilayer 25 20 25 20 1000 P, X, F, T, ST, Sp, V, N AlN/SiO2-8a & 8b MBI 10 Layers Varies 5 Varies 5 1000 P, X, F, T, ST, Sp, V, N AlN/SiO2-9a & 9b MBI 20 Layers Varies 10 Varies 10 1000 P, X, F, T, ST, Sp, E, V, N AlN/SiO2-10a & 10b MBI 30 Layers Varies 15 Varies 15 1000 P, Sp AlN/SiO2-11a & 11b Repeated Bilayer 50 10 50 10 1000 Extra Sample P, Sp AlN/SiO2-12a & 12b MBI 20 Layers Varies 10 Varies 10 1000 Extra Sample P, Sp Table 26: Sputtered AlN/SiOx Multilayer Samples (KIT) AlN Layers SiOx Layers Sample Name Type of Sample Thickness (nm) Number of Layers Thickness (nm) Number of Layers Total Thickness (nm) Characterization AlN/SiOx-1 Repeated Bilayer 50 10 50 10 1000 P, X, S, F, T, Sp, N AlN/SiOx-2 MBI Test Varies 4 Varies 4 225 P, X, Sp 138 AlN/SiOx-3 MBI 18 Layers Varies 9 Varies 9 1000 P, X AlN/SiOx-4 MBI 18 Layers #2 Varies 9 Varies 9 1001 P, X AlN/SiOx-5 MBI 20 Layers Varies 10 Varies 10 1000 P Table 27: Sputtered AlN/TiN Multilayer Samples AlN Layers TiN Layers Sample Name Type of Sample Thickness (nm) Number of Layers Thickness (nm) Number of Layers Total Thickness (nm) Characterization AlN/TiN-1 & 2 Repeated Bilayer 158 6 23 5 1062 P, X, Sp, N AlN/TiN-3a & 3b Repeated Bilayer 50 10 50 10 1000 P, Sp AlN/TiN-4a & 4b MBI 20 Layers Varies 10 Varies 10 1000 P, Sp Table 28: Sputtered TiO2/Al2O3 Multilayer Samples TiO2 Layers Al2O3 Layers Sample Name Type of Sample Thickness (nm) Number of Layers Thickness (nm) Number of Layers Total Thickness (nm) Characterization TiO2/Al2O3-1a & 1b Repeated Bilayer 50 10 50 10 1000 P, Sp TiO2/Al2O3-2a & 2b MBI 20 Layers Varies 10 Varies 10 1000 P, Sp Table 29:Sputtered TiO2/SiO2 Multilayer Samples TiO2 Layers SiO2 Layers Sample Name Type of Sample Thickness (nm) Number of Layers Thickness (nm) Number of Layers Total Thickness (nm) Characterization TiO2/SiO2-1a & 1b Repeated Bilayer 50 10 50 10 1000 P, S, F, T, ST, Sp, N TiO2/SiO2-2a & 2b Repeated Bilayer 100 5 1000 5 1000 P, Sp, N TiO2/SiO2-3a & 3b Repeated Bilayer 25 20 25 20 1000 P, Sp, N TiO2/SiO2-4a & 4b MBI 10 Layers #1 Varies 5 Varies 5 1000 P, F, T, Sp, N TiO2/SiO2-5a & 5b MBI 20 Layers #1 Varies 10 Varies 10 1000 P, Sp, N TiO2/SiO2-6a & 6b MBI 10 Layers #2 Varies 5 Varies 5 1000 P, Sp TiO2/SiO2-7a & 7b MBI 20 Layers #2 Varies 10 Varies 10 1000 P, Sp TiO2/SiO2-8a & 8b MBI 20 Layers #3 Varies 10 Varies 10 1000 P, F, T, Sp TiO2/SiO2-9a & 9b MBI 20 Layers #4 Varies 10 Varies 10 1000 P, Sp 139 TiO2/SiO2-10a & 10b MBI 20 Layers #5 Varies 10 Varies 10 1000 P, Sp TiO2/SiO2-11a & 11b MBI 20 Layers #6 Varies 10 Varies 10 1000 P, F, T, ST, Sp, N TiO2/SiO2-12a & 12b MBI 20 Layers #6 Varies 10 Varies 10 1000 Extra Sample P, Sp TiO2/SiO2-13a & 13b Repeated Bilayer 50 10 50 10 1000 Extra Sample Failed TiO2/SiO2-14a & 14b Repeated Bilayer 50 10 50 10 1000 Extra Sample P, Sp Table 30: Sputtered TiOx/SiOx Multilayer Sample (KIT) TiOx Layers SiOx Layers Sample Name Type of Sample Thickness (nm) Number of Layers Thickness (nm) Number of Layers Total Thickness (nm) Characterization TiOx/SiOx-1 Repeated Bilayer 50 10 50 10 1000 P, X, S, F, T, ST, Sp, N TiOx/SiOx-2 MBI Test Varies 3 Varies 3 135 P, X, Sp TiOx/SiOx-3 MBI 6 Layers Varies 3 Varies 3 1057 P, X, F, T, ST, Sp TiOx/SiOx-4 MBI 18 Layers Varies 9 Varies 6 1001 P, X, F, T, ST, Sp Table 31:Sputtered ZnO/ITO Multilayer Sample (KIT) TiOx Layers SiOx Layers Sample Name Type of Sample Thickness (nm) Number of Layers Thickness (nm) Number of Layers Total Thickness (nm) Characterization ZnO/ITO-1 Repeated Bilayer 50 10 50 10 1000 P, X, S, Sp, N ZnO/ITO-2 MBI 8 Layers Varies 4 Varies 4 1004 P, X, S, Sp, N 140 Appendix C: Spectroscopic Ellipsometry Measurements and Optical Properties of Monolithic Sputtered Films In this work, optical characterization was performed by spectrophotometry and ellipsometry. Spectrophotometry was primarily used for spectroscopic % transmittance measurements and ellipsometry measurements were used to model the complex index of refraction values. Chapters 4-6 presented experimental % transmittance for multilayered samples, however, the same optical characterization was performed on monolithic thin films. These measurements were then used in MBI optical predictions and to examine how sputtering parameters affected the optical behavior of the monolithic films. The following figures show AlN and Ag thin film measurements via ellipsometry and spectrophotometry from the AlN/Ag study in Chapter 4. For the AlN thin film, the psi and delta values were measured at various incident angles from 300- 1700 nm, Figure 68, and then model fit (WVASE software) for the complex index of refraction and extinction coefficient values, Figure 69. Spectrophotometry % transmittance measurements for AlN monolithic films at varying deposition parameters, highlighted in Figure 70, shows a strong detrimental effect on broadband transmittance with increased inert Ar working pressure (thereby lowering the Ar:N2 working gas ratio), while increasing applied power yields only a small decrease in spectroscopic transmittance. 141 Figure 68: Psi and delta values of a sputtered AlN single layer sample (from C. Appleget). Figure 69: Index of refraction and extinction coefficient generated by model fitting of measurements from a sputtered AlN single layer sample (from C. Appleget). Generated and Experimental Wavelength (nm) 300 580 860 1140 1420 1700 Y in degrees D in degrees 0 20 40 60 80 0 30 60 90 120 150 180 Model Fit Exp Y-E 65° Exp Y-E 75° Model Fit Exp D-E 65° Exp D-E 75° 400 600 800 1000 1200 1400 1600 0 1 2 3 4 5 6 Exp <n>-E 65° <n> Model Fit Exp <k>-E 65° <k> Model Fit Wavelength (nm) <n> -2 -1 0 1 2 3 4 <k> 142 Figure 70: AlN thin films prepared via magnetron sputtering, showing A) % Transmittance for the three samples, where average transmittance decreases slightly with increasing deposition power, and decreases significantly with a higher Ar flow rate and higher working pressure and B-D) top views of the as-sputtered samples (from C. Appleget). For the Ag, the psi and delta values were measured from 300-1700 nm, Figure 71, and then model fit for the complex index of refraction and extinction coefficient values, Figure 72. Figure 71: Psi and delta values of a sputtered Ag single layer sample (from C. Appleget). Generated and Experimental Wavelength (nm) 250 540 830 1120 1410 1700 Y in degrees D in degrees 24 27 30 33 36 39 42 45 60 80 100 120 140 160 Model Fit Exp Y-E 65° Model Fit Exp D-E 65° 143 Figure 72: Index of refraction and extinction coefficient generated by model fitting of measurements from a sputtered Ag single layer sample (from C. Appleget). 400 600 800 1000 1200 1400 1600 0.5 1.0 1.5 2.0 Exp <n>-E65° <n> Model Exp <k>-E 65° <k> Model Wavelength (nm) <n> 0 2 4 6 8 10 <k> 144 Appendix D: XRD Studies and Results XRD was used to characterize crystallinity in the monolithic and multilayered films. The monolithic sputtering rate (SR) samples were deposited to measure sputtering rate using profilometry and optical constants using ellipsometry, and XRD was performed to determine which crystal structures or phases were present in the films as a function of deposition parameters. For example, when sputtering AlN via reactive DC magnetron sputtering, as shown in Figure 73, XRD was used to determine if the Ar:N2 working gas ratio resulted in deposition of stoichiometric AlN. All scans presented in Figures 73-79 were deposited on amorphous Corning Eagle 2000 glass, and the amorphous background has not been removed. In reactive DC sputtering of AlN, Figure 73 indicates that with increasing applied power (75 W, 150 W, 200 W) at a constant working gas ratio (Ar 25 sccm, N2 7.5 sccm), the AlN (1 1 1) peak intensity increases. However, in reactive RF sputtering of AlN as shown in Figure 74, when increasing the applied power from 100 W to 150 W at a constant working gas ratio (Ar 40 sccm, N2 10 sccm), there is no measurable increase in the AlN (1 1 1) peak intensity. 145 Figure 73: XRD spectra of as-sputtered AlN-SR (sputtering rate) samples deposited by reactive DC magnetron sputtering on Corning Eagle 2000 substrates with varying deposition conditions. Sputtering parameters varied include power, non-reactive Ar gas flow, and reactive N2 gas flow. 146 Figure 74: XRD spectra of as-sputtered AlN-SR (sputtering rate) samples deposited by reactive RF magnetron sputtering on Corning Eagle 2000 substrates with varying deposition conditions. Sputtering parameters varied include power, non-reactive Ar gas flow, and reactive N2 gas flow. Figure 75: XRD spectra of as-sputtered ITO-SR (sputtering rate) samples deposited by non-reactive RF magnetron sputtering on Corning Eagle 2000 substrates with varying deposition power. 147 Additionally, amorphous ceramic layers were explored in Chapters 5 and 6, and the XRD spectra of SiOx and TiOx are shown in Figure 76 and Figure 77, respectively. These plots confirm an overall amorphous structure irrespective of varying sputtering power or working gas ratios. However, it is important to note that the presence of nanocrystallites or short-range order must be determined by TEM/HRTEM techniques [279]. Figure 76: XRD spectra of as-sputtered SiOx-SR (sputtering rate) samples deposited by reactive RF magnetron sputtering on Corning Eagle 2000 substrates with varying deposition conditions. Sputtering parameters varied include power, non-reactive Ar gas flow, and reactive O2 gas flow. All SiOx spectra are amorphous with no crystalline peaks detected. 148 Figure 77: XRD spectra of as-sputtered TiOx-SR (sputtering rate) samples deposited by reactive RF magnetron sputtering on Corning Eagle 2000 substrates with varying deposition conditions. Sputtering parameters varied include power, non-reactive Ar gas flow, and reactive O2 gas flow. All TiOx spectra are amorphous with no crystalline peaks detected. Figure 78: XRD spectra of as-sputtered ZnO-SR (sputtering rate) samples deposited by non-reactive RF magnetron sputtering on Corning Eagle 2000 substrates with varying deposition power. 149 Figure 79: XRD spectra of as-sputtered Ag-SR (sputtering rate) samples deposited by non-reactive DC magnetron sputtering on Corning Eagle 2000 substrates with varying non-reactive Ar flow. As for Ag thin films, with increasing working gas pressure, preferential (1 1 1) FCC peaks increased. A similar trend was observed by Jung et al, where a peak in the intensity ratio was observed as a function of sputtering pressure [280]. Overall, the XRD spectra served to examine the effects of deposition parameters on the film structure on the monolayer films before incorporation into a multilayered film. 150 Appendix E: In-House MBI MATLAB Code As described in Chapters 4-6, an in-house MATLAB optimization code based off the MBI Recursive Method was developed to predict high transmittance layer configurations. There are many off-the-shelf options available, such as Essential Macleod and OpTaliX, which can be used for design, analysis, and optimization of dielectric multilayer thin films. These are robust options that can incorporate experimental properties such as index of refraction, however, the goal with this work was to have a highly configurable in-house program, both with adaptable user variables and outputs. Input variables include the wavelengths of interest, the lower and upper bounds within the multilayer stack (this is to account for both sputtering limitations and to keep layers on the nanoscale), total film thickness, and to easily load experimental optical data measured by ellipsometry. As noted in Section 2.2.2, this is essential, as the synthesis parameters and microstructure strongly impact optical behavior: %% ============ USER VARIABLES ============ lambda = [251:1199]*10^-9; % Wavelengths to optimize over in nanometers num_bilayers = 10; % Total number of bilayers in the system (10 bilayers = 20 total layers in film stack) % Number of repeated bilayers. Currently set up to do (Si Ti) layers. num_runs = 200; % Number of optimization runs to do (with randomized start, more of these reduces possibliity of getting trapped in a local minima) d_lb = ones(1,num_bilayers*2)* 5e-9; % Lower bound for individual layer thickness in meters. d_ub = ones(1,num_bilayers*2)*0.2e-6; % Upper bound for individual layer thickness in meters. %% ============ LOAD MATERIAL DATA ============ Mat1_n = load('Mat_TiO2-SR-2a.txt'); % index of refraction values from first material file Mat2_n = load('Mat_Al2O3.txt'); % index of refraction values from second material file Corning_n = load('Mat_Corning.txt'); % index of refraction values from Corning Glass material file By using a relatively simple, in-house code, additional features and constraints, such as layer roughness, can be efficiently incorporated in the future. The optimization scheme is then 151 performed which finds a local solution and refines it within the user-specified tolerances. Multiple runs and configurations are calculated to maximize optical outputs by varying initial guesses: %% ============ OPTIMIZATION SETUP ============ % Create random initial guesses between bounds. d_guess = [rand rand]; for i=1:num_bilayers-1 d_guess(end+1) = rand; d_guess(end+1) = rand; end d_guess = d_lb + d_guess.*0.5.*(d_ub-d_lb); % SET LINEAR EQUALITY CONSTRAINT % there is the constraint A_eq*d=b_eq A_eq = ones(1,num_bilayers*2); % this means add all the layer thicknesses together b_eq = 1.0e-6; % this is the max total thickness % SET LINEAR INEQUALITY CONSTRAINT % there is a constraint that is A*d<=b A = []; b = []; %% ============ RUN OPTIMIZATION ============ fun = @(d)FunctionToOptimize2(n_data,n0,d,lambda,theta_mp1,num_bilayers); options = optimoptions('fmincon','Display','iter','TolCon',1e-20,'TolFun',1e- 11,'StepTolerance',1e-19); for j=1:num_runs % Run the optimization multiple times with different initial guesses, pick the best result. d_out(j,:) = fmincon(fun,d_guess,A,b,A_eq,b_eq,d_lb,d_ub,[],options); for i=1:length(lambda) Ti_n = n_data(i,1); Si_n = n_data(i,2); Cor_n = n_data(i,3); n = [Si_n Ti_n]; for k = 1:num_bilayers-1 n(end+1) = Si_n; n(end+1) = Ti_n; end n(end+1) = Cor_n; [R_f(i),T_f(i)] = MBI_Recursive(n,n0,d_out(j,:),lambda(i),theta_mp1); % Idk why T_f comes out complex in some cases end % THIS IS THE VALUE CASES ARE JUDGED ON meanT(j) = mean(real(T_f(51:551))); % calculate mean reflectance for this initial guess d_tot(j) = sum(d_out(j,:)); % Redo initial guess! d_guess = [rand rand]; for i=1:num_bilayers-1 d_guess(end+1) = rand; d_guess(end+1) = rand; end d_guess = d_lb + d_guess.*(d_ub-d_lb).*0.5; end [val,IDX] = max(meanT); d_out = d_out(IDX,:); 152 Studies on global optimization of multilayer thin films have been performed to overcome this nonconvexity, however, they are traditional computationally intensive [281, 282]. Therefore, though this code is only guaranteed to find the local minima rather than a global solution, it has proved to be a simple yet accurate guide for predicting highly transparent multilayer stacks. The calculations for the MBI equations detailed in the 2011 study by Lu et al [249] are used: function [R_f,T_f] = MBI_Recursive(n,n0,d,lambda,theta_mp1) num_layers = length(n)-1; m=num_layers; r_eff = zeros(1,num_layers); theta(m+1) = theta_mp1; % incoming angle from glass %% Calculate intermediate values [r(m+1),t(m+1)] = Fresnel(n(m+1),n(m),theta(m+1),theta(m)); for i=m+1:-1:3 % Theta should be calculated from m -> 1 (theta(m+1) given, theta_0 calculated later) % delta is calculated from m -> 1 (for each layer) theta(i-1) = asin(sin(theta(i))*n(i)/n(i-1)); delta(i-1) = 4*pi*d(i-1)*cos(theta(i-1))/lambda; [r(i-1),t(i-1)] = Fresnel(n(i-1),n(i-2),theta(i-1),theta(i-2)); end theta(1) = asin(sin(theta(2))*n(2)/n(1)); % incidence angle in first layer theta_0 = asin(sin(theta(1))*n(1)/n0); % Light incidence angle in air above layers delta(1) = 4*pi*d(1)*cos(theta(1))/lambda; [r(1),t(1)] = Fresnel(n(1),n0,theta(1),theta_mp1); % for top interface between air and top layer %% Calculate effective values for second-to-bottom interface (r_eff_m, t_eff_m) r_eff(m) = (r(m) + r(m+1)*exp(2i*delta(m))) / (1+r(m)*r(m+1)*exp(2i*delta(m))); t_eff(m) = (t(m)*t(m+1)*exp(1i*delta(m)))/(1+r(m)*r(m+1)*exp(2i*delta(m))); %% Do recursive method up to top layer for i=m:-1:2 r_eff(i-1) = (r(i-1)+r_eff(i)*exp(1i*2*delta(i-1))) / (1+r(i- 1)*r_eff(i)*exp(2i*delta(i-1))); t_eff(i-1) = (t(i-1)*t_eff(i)*exp(1i*delta(i-1))) / (1+r(i- 1)*r_eff(i)*exp(2i*delta(i-1))); end %% Calculate effective total values R_f = r_eff(1)*conj(r_eff(1)); T_f = ((n(m+1)*cos(theta(m+1)))/(n0*cos(theta_0)))*t_eff(1)*conj(t_eff(1)); end 153 The optimization function is then used to minimize reflectance in the desired wavelengths. The current code outputs the estimated complex % transmittance, % reflectance, and associated layer thicknesses, thus serving as a guideline for subsequent optically optimized NM synthesis: Figure 80: (left) theoretical prediction of a 10 layer AlN/SiO2 NM and (right) associated layer thicknesses It is important to note that the MBI calculations serve as a guideline for synthesis, and do not account for complexities that can arise in multilayers such as anisotropy or homogeneities within the layers, interface roughness, or voids. The agreement between calculated MBI % transmittance predictions and experimental % transmittance was highlighted in Chapters 4-6, where average % transmittance values across the wavelength spectrum were presented. The good agreement between MBI-calculated values and experimental measurements, an example of which is highlighted in Table 8 (e.g. predicted %T380-1100nm≈95.3% vs. experimental %T380-1100nm≈98.7% for AlN/SiO2 MBI), suggests that the assumptions of homogeneous layers with smooth interfaces were acceptable. Figure 81 displays good spectroscopic agreement in the AlN/SiO2 MBI 20 Layer sample, with deviations only occurring below 400 nm, while Figure 82 shows a poor spectroscopic match. Possibly sources of these differentials could be deviations between desired and 154 experimental layer thicknesses, layer homogeneities, presence of optical imperfects, or absorbance of the substrate. Figure 81: Spectroscopic predicted (MBI) vs. experimental % transmittance for AlN/SiO2 MBI 20 Layer sample Figure 82: Spectroscopic predicted (MBI) vs. experimental % transmittance for repeated bilayer AlN(50)/SiO2(50) The current full code and example data are both available from the author.

Abstract (if available)

Abstract Nanomultilayer (NM) systems have demonstrated desirable properties such as corrosion resistance, radiation tolerance, and high strength due to the interfaces and length scale effects. Additionally, multilayers have commonly been used in optical systems as the layer thicknesses can be tuned to be identical to interaction lengths of photons. Although NMs have been shown to have useful optical and mechanical properties independently, few studies have been performed to investigate the interplay between optics and mechanics. Thus, the possible novel combination of properties, such as transparency and strength, presents an extraordinary area of research with potential applications in optical windows, sensor protection, and numerous other functions which require light penetration for function and a robust barrier for protection. ❧ The studies described in this dissertation provide a foundation for understanding the relationship between length-scale effects, interface effects, transparency, and multifunctional properties in ceramic/ceramic and metal/ceramic NMs. Specifically, this relationship was explored through (1) the development of an approach to synthesize and study the optical properties of multilayers in a model AlN/Ag system, then (2) the investigation of microstructural effects of maximizing transparency in AlN/SiO₂ multilayers and finally (3) the exploration of amorphous/amorphous and crystalline/amorphous interface contributions in highly transparent ceramic/ceramic multilayers. By identifying the dominant mechanisms in the relationship between optical performance and resultant film properties, transparent optical multilayer systems can be tailored with optimal layer thicknesses to aid in the design and synthesis of new high-performance, long-lasting optical materials.

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Asset Metadata

Creator Appleget, Chelsea Diane (author)

Core Title Development and characterization of transparent metal/ceramic and ceramic/ceramic nanomultilayers

School Viterbi School of Engineering

Degree Doctor of Philosophy

Degree Program Aerospace Engineering

Publication Date 07/31/2020

Defense Date 06/12/2020

Publisher University of Southern California (original), University of Southern California. Libraries (digital)

Tag multilayer thin films,nanoindentation,nanostructured materials,OAI-PMH Harvest,optical materials,sputtering,transmittance

Language English

Contributor Electronically uploaded by the author (provenance)

Advisor Hodge, Andrea (committee chair), Branicio, Paulo (committee member), Luhar, Mitul (committee member), Salem, Jonathan (committee member)

Creator Email capplege@usc.edu

Permanent Link (DOI) https://doi.org/10.25549/usctheses-c89-350955

Unique identifier UC11663956

Identifier etd-ApplegetCh-8834.pdf (filename),usctheses-c89-350955 (legacy record id)

Legacy Identifier etd-ApplegetCh-8834.pdf

Dmrecord 350955

Document Type Dissertation

Rights Appleget, Chelsea Diane

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Source University of Southern California (contributing entity), University of Southern California Dissertations and Theses (collection)

Access Conditions The author retains rights to his/her dissertation, thesis or other graduate work according to U.S. copyright law. Electronic access is being provided by the USC Libraries in agreement with the a...

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