Aluminium Nitride Piezoelectric Thin Films

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ALUMINIUM NITRIDE PIEZOELECTRIC THIN FILMS REACTIVELY DEPOSITED IN CLOSED FIELD UNBALANCED MAGNETRON SPUTTERING FOR ELEVATED TEMPERATURE ’SMART’ TRIBOLOGICAL APPLICATIONS

Masood Hasheminiasari A thesis submitted to the Faculty and Board of Trustees of the Colorado School of Mines in partial fulfillment of the requirements for the degree of Doctor of Philosophy (Materials Science).

Golden, Colorado

Date ______

Signed: ______

Masood Hasheminiasari

Signed: ______

Dr. Jianliang Lin

Thesis Advisor

Signed: ______

Dr. John J. Moore

Thesis Advisor

Golden, Colorado

Date ______

Signed: ______

Dr. Michael Kaufman

Professor and Head

Department of Materials Science

ABSTRACT

―Smart‖ high temperature piezoelectric aluminum nitride (AlN) thin films were synthesized by reactive magnetron sputtering using DC; pulsed-DC, and deep oscillation modulated pulsed power (DOMPP) systems on variety of substrate materials. Process optimization was performed to obtain highly c-axis texture films with improved piezoelectric response via studying the interplay between process parameters, microstructure and properties.

AlN thin films were sputtered with DC and pulsed-DC systems to investigate the effect of various deposition parameters such as reactive gas ratio, working pressure, target power, pulsing frequency, substrate bias, substrate heating and seed layers on the properties and performance of the film device. The c-axis texture, orientation, microstructure, and chemical composition of AlN films were characterized by means of X-ray diffraction (XRD), field emission scanning electron microscopy (FESEM), transmission electron microscopy (TEM), and x-ray photoelectron spectroscopy (XPS).

A Michelson laser interferometer was designed and built to obtain the converse piezoelectric response of the deposited AlN thin films. Thin films with narrow AlN-(002) rocking curve of 2.5° were obtained with preliminary studies of DOMPP reactive sputtering. In-situ high temperature XRD showed excellent thermal stability and oxidation resistance of AlN films up to 1000 °C. AlN films with optimized processing parameters yielded an inverse piezoelectric coefficient, d33 of 4.9 pm/V close to 90 percent of its theoretical value.

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TABLE OF CONTENTS

ABSTRACT ...... iii TABLE OF CONTENTS ...... iv LIST OF FIGURES ...... vi LIST OF TABLES ...... xii ACKNOWLEGMENTS ...... xiii CHAPTER 1. INTRODUCTION ...... 1 1.1 Hypothesis ...... 2 CHAPTER 2. RESEARCH BACKGROUND...... 3 2.1 Aluminum Nitride Structure and Properties ...... 3 2.2 Sensor Materials Selection Criteria ...... 7 2.3 Piezoelectricity Fundamentals...... 9 2.3.1 Thermodynamic Approach to Constitutive Equations ...... 18 2.3.2 AlN c-Axis Texture and Piezoelectric Response Model ...... 20 2.3.3 Piezoelectric Materials ...... 23 2.3.4 Piezoelectric Measurements Methods ...... 30 2.4 Low Density Plasma...... 40 2.4.1 Sputtering Deposition Process ...... 42 2.4.2 Magnetron Sputtering, Balanced and Unbalanced ...... 45 2.4.3 Closed Field Unbalanced Magnetron Sputtering (CFUBMS) ...... 47 2.4.4 Reactive Sputtering ...... 49 2.4.5 High Power Pulsed Magnetron Sputtering (HPPMS) ...... 52 2.4.6 Modulated Pulsed Power (MPP) Magnetron Sputtering ...... 55 2.4.7 MPP Pulse Characteristics ...... 56 2.5 Structure Zone Models and Film Microstructure Evolution ...... 61 CHAPTER 3. EXPERIMENTAL PROCEDURES ...... 68 3.1 AlN Thin Film Deposition Process ...... 68 3.2 Film Microstructure and Composition Characterization ...... 72 3.2.1 X-Ray Diffraction (XRD) ...... 72 3.2.2 X-Ray Photoelectron Spectroscopy (XPS) ...... 76 3.2.3 Field Emission Scanning Electron Microscopy (FESEM) ...... 77 3.2.4 Scanning Transmission Electron Microscopy (STEM) ...... 78 3.3 Film Properties Characterization ...... 78

3.3.1 Wafer Flexure Technique (Direct Piezoelectric Measurement) ...... 79 3.3.2 Michelson Interferometry (Remote Inverse Piezoelectric Measurement) ...... 80 3.3.3 Mechanical Properties (Nanoindentation) ...... 82 3.3.4 Residual Stress Measurements ...... 85 3.3.5 Oxidation Resistance and High Temperature Stability ...... 88 CHAPTER 4. RESULTS AND DISCUSSION ...... 90 4.1 DC Reactive Sputtering of AlN Films ...... 90 4.1.1 Reactive Gas Ratio Effect ...... 90 4.1.2 Insulation Characteristics of AlN Films ...... 94 4.1.3 Buffer Layer (Seed Layer) Effect on c-Axis Orientation ...... 97 4.1.4 Substrate Biasing Effect, DC and P-DC...... 100 4.2 Pulsed-DC Reactive Sputtering of AlN Films ...... 104 4.2.1 Working Pressure Effect ...... 105 4.2.2 Target Pulsing Frequency Effect ...... 109 4.2.3 Effect of Nitrogen/Argon Ratio...... 115 4.2.4 Seed-Layer Effect (Mo and Pt) ...... 123 4.2.5 High Temperature Stability (DSC and In-Situ HTXRD) ...... 126 4.2.6 High Temperature Oxidation Resistance ...... 130 4.2.7 Substrate Heating Effect on (002) Orientation ...... 136 4.2.8 Transmission Electron Microscopy (TEM) Study ...... 141 4.3 Modulated Pulsed Power Sputtering of AlN Films ...... 143 CHAPTER 5. CONCLUSIONS ...... 149 REFERENCES ...... 153

LIST OF FIGURES

Figure 1.1‎ Optimized tribological coating‘s schematic with smart layer sandwiched between the substrate and top wear resistance coating [175]...... 2 Figure 2.1‎ Wurtzite hexagonal close-packed structure of AlN with nitrogen atom in white circles and aluminum atoms in black circles [126]...... 4 Figure 2.2‎ Schematic of the films with induced imperfection introducing (a) rough surface, (b) porous film and (c) films with cavities [175]...... 7 Figure 2.3‎ The relations between physical domains relating mechanical, electrical and thermal domains with their linear coefficients labeled on the lines connecting two [37]...... 11 Figure 2.4‎ Schematic representation of bound charges within a dielectric insulator upon electric field application [38]...... 11 Figure 2.5‎ Relation between stress and electric displacement in piezoelectric materials [38]...... 13 Figure 2.6‎ Relation between electric field and strain in piezoelectric materials [38]...... 14 Figure 2.7‎ Three dimensional representation of a piezoelectric cube with its coordinate axes [38]...... 16 Figure 2.8‎ The transformation from the reference frame of the substrate to the coordinates of a crystal. The black arrows represents the new coordinate axes formed from a rotation of the gray axes [39]...... 20

Figure 2.9‎ The effective piezoelectric constant e33 as a function of the rocking curve FWHM showing the abrupt decay of piezoelectric response as a function of increasing FWHM [39]...... 22

Figure 2.10‎ Crystal structure of the Perovskite ferroelectric BaTiO3 (A) high temperature paraelectric cubic phase (B & C) room temperature, ferroelectric, tetragonal phases, showing up and down variants [42]...... 23 Figure 2.11‎ Typical wurtzite materials having hexagonal closed-packed (HCP) structures with tetrahedral interstitial sites of anion lattice filled by cations [176]...... 24 Figure 2.12‎ Lead zirconate titanate (PZT) crystal structure showing both (1) cubic and (2) tetragonal phases [44]...... 25 Figure 2.13‎ Lead titanate Perovskite structure with lead atoms in tetragonal lattice arrangement titanium atoms in octahedral interstitial site [176]...... 26 Figure 2.14‎ ZnO nanostructures synthesized by thermal evaporation of solid powders (a) nanocombs, (b) tetraleg, (c) hexagonal disks, (d) nanopropellers, (f) nanospiral, (g) nanosprings, (h) single crystal nanoring and (i) combination of rods, bow and ring [46]...... 27 Figure 2.15‎ Zinc oxide (ZnO) wurtzite hexagonal close packed structure [176]...... 27 Figure 2.16‎ Piezoelectric and spontaneous polarization effect in Ga (Al) or N‐face AlGaN/GaN heterostructures [50]...... 28

Figure 2.17‎ Lithium Niobate's trigonal crystal structure with Niobium atoms in blue, Lithium atoms in green and Oxygen atoms in red [176]...... 29

Figure 2.18‎ Experimental setup for d33 measurements by the method of the pneumatic pressure rig [58]...... 33

Figure 2.19‎ Experimental setup for d33 measurement by the method of the metallic rod [57]. .... 34 Figure 2.20‎ Schematic of a setup for the measurement of the transverse piezoelectric coefficient by the method of the cantilever beam [57]...... 35 Figure 2.21‎ Schematic of the double-beam interferometer (top) and principle of elimination of bending motion of the substrate (bottom) [57]...... 37 Figure 2.22‎ Schematic representation of a Michelson Interferometer...... 40 Figure 2.23‎ A schematic representation of the physical sputtering processes [63]...... 43 Figure 2.24‎ Geometry and typical voltage profile of the diode configuration dc glow discharge showing the plasma potential and the cathode and anode sheath [66]...... 44 Figure 2.25‎ Schematic illustration of typical magnetron sputtering showing the confinement of electrons to the target region [66]...... 45 Figure 2.26‎ A comparison of the magnetic configuration and plasma confinement in (a) dc balanced magnetron sputtering and (b) dc unbalanced magnetron sputtering...... 46 Figure 2.27‎ diagram showing the mirrored (a) and closed field (b) magnetic configurations where in the mirrored configuration similar poles face each other and in closed field configuration the opposite poles face each other [78]...... 48 Figure 2.‎ 28 Schematics of four unbalanced magnetrons arranged in closed field configuration. An EQP mass spectrometer is also put inside the chamber to analyze the species in the plasma [79]...... 49 Figure 2.29‎ An illustration of the reactive sputtering process, where in this case aluminum reacts with the nitrogen to form aluminum nitride [80]...... 50 Figure 2.30‎ Target voltage waveforms when operated in (a) continuous DC, (b) unipolar pulsed mode, (c) asymmetric bipolar pulsed mode, and (d) symmetric bipolar pulsed mode [79]...... 51 Figure 2.31‎ SEM micrographs of fracture section of aluminum oxide film deposited by (a) DC reactive sputtering and (b) pulsed closed field unbalanced reactive sputtering [70]...... 52 Figure 2.32‎ SEM micrographs of (Ti,Al)N coating cross-sections from a) dcMS, b) mfMS, and c) HPPMS [89]...... 53 Figure 2.33‎ SEM micrograph of (Ti,Al)N coating cross-section on a cutting edge deposited by dc-MS (top) and HPPMS (bottom) [89]...... 54 Figure 2.34‎ Comparison of pulse shapes from (a) HPPMS and (b) MPP [89], [91]...... 55 Figure 2.35‎ Typical target waveforms and parameters for (a) HPPMS and (b) MPP [92]...... 56

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Figure 2.36‎ The target voltage, current, and power waveforms during one modulated pulse used for the deposition of tantalum coatings (pulse width = 1000 μs, frequency = 65 Hz, Power-average =2 kW). The numbers indicate separate steps of the modulated pulse [93]...... 57 Figure 2.37‎ Ion energy distributions from a HPPMS discharge at various time periods during the discharge for Ar+ ions and Ti+ ions [98]...... 59 Figure 2.38‎ Ion energy distributions of Ar+, Ta+, and Ta2+ ions during the MPP Ta discharge [99]...... 60 Figure 2.39‎ Three modes of thin film nucleation and growth showing 2-D layer by layer growth mode proposed by Frank-van der Merwe, 3-D island growth by Volmer-Weber, and Stranski-Krastanov mixed growth...... 62 Figure 2.40‎ Zone model for the grain structure of vapor deposited metal films modified at low temperatures to take into account the fine equiaxed grain growth [106]...... 63 Figure 2.41‎ Structure zone model for thick films showing the effect of both bombardment and thermal induced mobility [108]...... 64 Figure 2.42‎ Microstructural zone diagram for sputter deposition processes proposed by Thornton [107]...... 65 Figure 2.43‎ Structure zone model relating to the CFUBMS system. Solid circles mark the position of coatings with zone 2 structure and shaded circles mark the position of coatings with zone 3 structures [72]...... 67 Figure 3.1‎ Illustration of four cathode arrangement inside vacuum chamber including the position of the substrate holder and the closed magnetic field created by the four facing unbalanced magnetrons...... 69 Figure 3.2‎ Photographs of (a) Horizontal vacuum chamber (DC), (b) P-CFUBMS system (pulsed-DC), and (c) MPP power supply and two cathode systems...... 70 Figure 3.3‎ X-ray diffraction in (a) conventional configuration /2 with the incident angle being the same as Bragg angle, (b) and (c): change of the incident angle for a fixed Bragg angle to obtain rocking curve...... 75 Figure 3.4‎ A photograph of the X-Ray Photoelectron Spectrometer XPS (AXIS HS XPS, Kratos Analytical Ltd.)...... 77 Figure 3.5‎ A photograph showing the JOEL JSM-7000 Field Emission Scanning Electron Microscope (FESEM)...... 78 Figure 3.6‎ Schematic of a piezoelectric measurement setup using wafer flexure mode...... 79 Figure 3.7‎ Homemade Michelson interferometer for remote piezoelectric measurements and its schematic ray diagram on the right...... 81 Figure 3.8‎ Schematic representation of the nanoindentation load-displacement curve...... 83 Figure 3.‎ 9 Initial steps of film deposition when the film is too narrow for the substrate (top- left) and too wide (top-right) and resulting end moments and how the film and substrate react. The bottom-left film is under residual tensile stress and the bottom- right film is under residual compressive stress...... 86

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Figure 3.10‎ Hot stage module attached to Philips XRD (HTXRD)...... 89 Figure 4.1‎ XRD patterns of AlN thin films deposited using DC reactive sputtering on Ti/SS as a function of various Ar/N2 ratios...... 91

Figure 4.2‎ Grain size measurements of AlN/Ti /SS thin films deposited with various N2 to Ar ratios...... 92 Figure 4.3‎ XRD patterns of as-deposited and annealed AlN film at 600 C for 2 hours...... 93 Figure 4.4‎ The I-V characteristics of the annealed AlN thin films at 600 for two hours in pure nitrogen environment...... 94 Figure 4.5‎ Two-probe I-V characteristics of AlN thin film with two different thicknesses, one deposited with 350 nm and the other coating deposited with the same deposition condition with thickness of about 900 nm...... 95 Figure 4.6‎ Schematic crystal structure of Ti, Pt and TiN seed layers showing their atomic symmetry within the closed pack planes and their lattice mismatch with AlN (002) basal planes...... 98 Figure 4.7‎ XRD of AlN films deposited with different seed layers with relation of their crystal structure and symmetry...... 99 Figure 4.8‎ XRD peaks of AlN films deposited on Si substrates with pulsed bias (-50 V/20 kHz) in red and unbiased substrate in black...... 101 Figure 4.9‎ Effect of Substrate biasing: DC and pulsed DC on grain size and microstructure of DC-sputtered AlN films with cross-section and top view SEMs...... 103 Figure 4.10‎ SEM micrograph of a thin Pt electrode deposited on a rough film surface compared to non-platinized AlN film...... 104 Figure 4.11‎ The working pressure effect on residual stress (a), crystallite size (b), and XRD rocking curves (c)...... 107 Figure 4.12‎ Cross-sectional FESEM micrographs of pulsed DC AlN films deposited with working pressure ranges from 2 to 6.5 mTorr...... 108 Figure 4.13‎ Nitrogen ion energy distribution as a function of pulsing frequency and reverse time [147]...... 109 Figure 4.14‎ Normalized preferred c-axis growth of AlN films deposited at various pulsing frequency...... 110 Figure 4.15‎ Comprehensive Rocking curves for film AlN/Cr/glass deposited at various pulsing frequencies...... 112 Figure 4.16‎ Dependence of Residual stress on the deposition parameters such as various pulsing frequencies and substrate bias...... 113 Figure 4.17‎ FESEM Cross-sectional micrograph of AlN/Cr films on Si substrate deposited at different pulsing frequencies...... 114 Figure 4.18‎ Conventional XRD peaks of AlN/Cr films deposited on glass at different nitrogen gas ratio at 3 and 5 mTorr...... 117

Figure 4.19‎ (a) Al 2p and (b) N 1s core level spectra of one-hour argon-etched AlN films deposited with different nitrogen to argon gas ratio...... 118 Figure 4.20‎ Gaussian curve fitting of (a) Al 2p and (b) N 1s core level spectra of one-hour argon-etched as a function of nitrogen/argon ratio...... 119 Figure 4.21‎ Residual stress measurements of AlN films deposited at ―5mTorr, 100 kHz, 1microsecond, FB‖ (top), and the calculated (002) grain size (bottom)...... 121 Figure 4.22‎ FESEM cross-sectional micrograph of AlN/Cr films on Si substrate deposited at 5 mTorr, 1 kW, 100 kHz, and 1 microsecond reverse time with different nitrogen gas ratio...... 122 Figure 4.23‎ XRD graphs of deposited AlN on stainless steel (a) at 3 and 5 mTorr using Pt and (b) comparison of XRD of AlN deposited on Pt and Mo...... 124

Figure 4.24‎ Piezoelectric coefficient (d33eff) measured by interferometry for optimized AlN film deposited on Mo seed layer...... 125 Figure 4.25‎ (a) DSC data and its first derivative acquired in argon environment (55 sccm) and (b) the corresponding XRD data before and after DSC test...... 127 Figure 4.26‎ FESEM cross-sectional and plan view of AlN film before and after DSC 1000 °C in argon environment...... 128 Figure 4.27‎ In-situ XRD spectra of AlN film studied at temperatures of 600, 700, 800, 900, and 1000 °C...... 129 Figure 4.28‎ In-situ high temperature XRD scans for non-textured AlN/Cr film at temperatures of 600, 700, 800, 900, and 1000 degree C in air, respectively...... 131 Figure 4.29‎ In-situ XRD scans of AlN sputtered on TiN/Ti/Si at 600, 700, 800, 900, and 1000 degree C in air...... 132 Figure 4.30‎ Residual stress relaxation of sputtered AlN films at 600 °C to 1000 °C in He controlled environment...... 134 Figure 4.31‎ Residual stress evolution of AlN films at 800 °C for 24 hours in air...... 135 Figure 4.32‎ Substrate temperature variation effect on preferred (002) orientation of AlN thin films deposited with optimized sputtering parameters...... 137 Figure 4.33‎ AlN (002) conventional XRD FWHM of films deposited at no substrate heating and substrate temperatures of 200, 300, and 400 °C...... 138 Figure 4.34‎ Laser interferometry displacement obtained by the application of increasing applied voltages from 2 to 10 volts for AlN deposited at no substrate heating compared to the one at 200 °C substrate temperature (a), rocking curve FWHM of AlN films deposited at no substrate heating and 200 °C, respectively...... 139

Figure 4.35‎ Piezoelectric coefficient d33 measurements for optimized AlN sample at frequency of modulation ranging from 10 to 100 kHz...... 140 Figure 4.36‎ TEM cross-section of optimized AlN deposited by P-DC on silicon substrates...... 141 Figure 4.37‎ Cross-sectional HRTEM image of AlN films deposited at 5 mtorr with its electron diffraction pattern at bottom...... 142

Figure 4.38‎ Deep oscillation MPP voltage and current waveforms showing macro pulses (a) and micro pulses (b)...... 143 Figure 4.39‎ Convention XRD of AlN film deposited by DOMPP (a) and its rocking curve comparison with respect to P-DC sputtered AlN film...... 144 Figure 4.40‎ Cross-sectional TEM micrograph of DOMPP AlN deposited on Si and its selected area electron diffraction (SAED) on the top...... 145 Figure 4.41‎ Argon IEDs for AlN deposited under a DC discharge (a) and P-DC (b) [174]...... 146 + + + Figure 4.42‎ IEDs of (a) Cr , (b) Ar , (c) N2 ions and (d) the integrated ion fluxes of Cr target in Ar:N2=1:1 atmosphere under DC, P-DC, and two MPP conditions [91]...... 147 Figure 4.43‎ A comparative residual stress measurements for AlN sputtered using DC, P-DC and DOMPP...... 148

LIST OF TABLES

Table 2.1‎ Physical properties of major piezoelectric materials widely used in devices [26–30]. .... 8 Table 3.1‎ Deposition summary of AlN deposited by reactive DC and pulsed DC magnetron sputtering...... 72 Table 4.1‎ Piezoelectric coefficient measurements for AlN films deposited with different thicknesses...... 97 Table 4.2‎ Piezoelectric coefficient measurements for AlN films deposited with different seed layer systems...... 100 Table 4.3‎ DC-sputtered AlN films deposited on DC and pulsed-DC biased substrates with similar optimized deposition parameters...... 104 Table 4.4‎ Deposition parameters of AlN films deposited at different working pressures of 2- 6.5 mTorr...... 106 2 Table 4.5‎ Piezoelectric coefficient e31 (c/m ) for AlN film deposited in pulsing frequencies of 100, 200 and 300 kHz...... 111 Table 4.6‎ Deposition parameters of AlN films deposited at various nitrogen/argon ratios...... 115 Table 4.7‎ XPS compositional and atomic concentration of AlN films deposited at different nitrogen ratio and their Al/N atomic ratio after one hour ion etching...... 120 Table 4.8‎ Cathode current, cathode voltage, and nitrogen flow rate for AlN films deposited at 5 mTorr, 1 kW, 100 kHz, and 1 microsecond reverse time...... 122 Table 4.9‎ Deposition summary of seed layers (Pt and Mo) sputtered on stainless steel substrates...... 123 Table 4.10‎ Grain size calculation from XRD of AlN (002) oriented film from room temperature to 1000 °C in controlled He atmosphere...... 134 Table 4.11‎ AlN deposition summary using different substrate heating from 200 °C to 400 °C during sputtering process...... 136

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ACKNOWLEGMENTS

This PhD work was performed in the Advanced Coatings and Surface Engineering Laboratory (ACSEL) in the Metallurgical and Materials Engineering department at Colorado School of Mines (CSM). I own a great dept to many people for their generous and endless assistance and positive inputs during my thesis studies.

I wish to extend my deep gratitude to my supervisors Dr. John J. Moore and Dr. Jianliang Lin for their guidance and crucial advice during my research work. They have influenced me beyond just my research field and showed me how to manage and collaborate in the multidisciplinary subjects such as thin film area.

I would also like to express my deep gratitude to Dr. John Scales for his valuable comments and insight especially with his help to setup the interferometry measurements and laser vibrometer data acquisitions. Dr. Scales always has been a pleasure to work with and learn from his comprehensive knowledge. I would also like to thank Dr. Olson, Dr. Sproul and Dr. Ahrenkeil for their precious comments and suggestions on my thesis. I would like to thank Scott Pawelka for his continuous help on the equipment maintenance and problem solving during my research.

I would also like to thank my colleagues in ACSEL for their friendship and assistance on this research: Sterling Myers, In-Wook Park, Ari Feldman, William Garrett, Fengli Wang, Sudipta Bhattacharyya, William Moerbe, Zhilli Wu, Huiyao Wang, Bo Wang and Ningyi Zhang.

This research was financially supported by the Department of Energy (DOE) and North American Die Casting Association (NADCA). Without their gerenrous funding, this research would not be possible to be conducted.

Finally I would like to express my deepest gratidute to my Almighty creator for his bellesing opportunity of existance also to my parrents and my wife for their endless support, encouragment and love.

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CHAPTER 1. INTRODUCTION

Thermal and mechanical stresses during Aluminum die-casting introduce severe wear and result in parts out of desired specifications. Any differences in coefficient of thermal expansions (CTE mismatch) will induce intrinsic stress in the dies and the mechanical stress is also incorporated within the die in every shot. A tribological coating system can be designed to minimize these two types of stresses and prolong the life cycle of the dies. A piezoelectric ―smart layer‖ can be incorporated under the protective layer to avoid any catastrophic failure of the die.

A piezoelectric smart layer can identify the stresses that develop during processing with an appropriate calibration. In order to minimize any leakage of current through the structure, an insulating high temperature coating needs to be deposited on the smart layer to maintain the piezoelectric signals within the coating system.

The layer deposited on the top of smart layer consists of a nano-composite layer with high hardness, low friction coefficient and high mechanical adhesion, as shown in Figure 1‎ .1. Even though this is not the main focus of this research, it is an essential part of the idea of smart protective coating system for aluminum die casting. These superlattice coatings were deposited by pulsed closed field unbalanced magnetron sputtering system (P-CFUBMS). A fairly new technique with higher plasma intensity (modulated pulsed power (MPP)) is under further investigation.

In-situ ―smart layers‖ can be deposited using reactive sputtering to measure any abrupt changes in thermo-mechanical stress during die casting. In order to optimize the smart coating system with its highest piezoelectric response, the fundamental interrelation between synthesis, microstructure and properties must be fully understood. These optimization processes can help us extend the usage of these smart coating systems in other applications beyond die casting such as high temperature aerospace applications.

Figure ‎1.1 Optimized tribological coating‘s schematic with smart layer sandwiched between the substrate and top wear resistance coating [175].

1.1 Hypothesis

A ‗smart‘ high temperature coating system can be developed by embedding a thin film AlN piezoelectric stress/pressure sensor into a tribological coating system. The AlN thin film sensor with good c-axis or (002) orientation and high piezoelectric response can be deposited by optimizing key deposition parameters and the seed layers using dc and pulsed dc closed field unbalanced magnetron sputtering (PCFUMS).

In addition, the (002) orientation and piezoelectricity of the AlN thin films can be further enhanced by using a highly ionized plasma generated by deep oscillation- modulated pulsed power magnetron sputtering (DOMPPMS). By the appropriate control of ion energies and ion fluxes impinging on the growing film, highly c-axis textures AlN films having minimal FWHM of rocking curve with optimum piezoelectric response can be deposited

CHAPTER 2. RESEARCH BACKGROUND

This following chapter reviews the current literature on aluminum nitride (AlN) thin films and the materials selection based on the objectives of the project. Different types of magnetron sputtering system are summarized. Piezoelectric materials background and its characterization techniques are also briefly introduced.

2.1 Aluminum Nitride Structure and Properties

Under ambient conditions, Aluminum nitride (AlN) has a thermodynamic stable hexagonal closed packed Wurtzite structure, as illustrated in Figure 2‎ .1, with lattice parameters ranging from 3.110 to 3.113 Å for the a-axis and from 4.978 to 4.982 Å for the c-axis.

In AlN crystal structure, each Al atom is surrounded by four N atoms forming a tetrahedron with two different bonding angles of 107.7° and 110.5°. The Al atom has three semi-full orbits and one empty orbit; the N atom has three semi-full orbits and one full orbit. The coupling of the empty Al orbit with the N full orbit forms the c-axis bonding between Al and N atoms and hence the ionic character of this bonding is greater than the other three bonding. The bond energy of the c-axis bonding is relatively smaller than of the other ones.

The c/a ratio varies between 1.600 and 1.602 deviating from the ideal Wurtzite structure probably due to iconicity and lattice instability, and at very high pressure it has a meta-stable cubic rock-salt structure [1]. Due to its covalent bonding structure, it has high hardness and metal-like thermal conductivity that it shares with the other III-V nitride semiconductors [2, 3]. Thus, AlN ceramics have been used as heat dissipation components in power supply industries due to their excellent thermal conductivity [4]. Because of its high corrosion resistance and hardness, AlN is also used as wear resistance and corrosion protective coatings [5].

AlN has high chemical and thermal stability, high hardness, high thermal conductivity exceeding pure aluminum and gold (320 W/mK), high electrical resistivity, a wide direct band gap (6.2 eV) [6], a low thermal expansion coefficient ( 4 106 / C ), good piezoelectric property and

high acoustic velocity. AlN has relatively high piezoelectric coefficients d33 and d31 permitting static and dynamic motion. AlN thin films are widely used in variety of piezoelectric applications, from bio-sensing, pressure sensing to the actuating layer in complex micro-actuator arrays [7–9].

Figure ‎2.1 Wurtzite hexagonal close-packed structure of AlN with nitrogen atom in white circles and aluminum atoms in black circles [126].

Thin film AlN have been extensively studied using the most common deposition techniques including plasma enhanced chemical vapor deposition (PECVD) [10], metal-organic chemical vapor deposition (MOCVD) [11], molecular beam epitaxy (MBE) [12], pulsed laser deposition (PLD) [13–15], DC magnetron sputtering [16], pulsed DC magnetron sputtering [17–19] and RF magnetron sputtering [16], [20–22]. The CVD techniques have excellent process controllability and suitability for mass production [23] but substrate temperature in conventional CVDs is extremely high (larger than 1000 °C ) and due to the high grain growth rate, smooth surface morphology is rarely obtained. Reactive magnetron sputtering is a more desirable technique due to its ease of deposition, low temperature growth, inexpensive equipment and compatibility with semiconductor industry. In this section, the fundamental background of AlN structure, piezoelectric phenomena, magnetron sputtering and interferometry will be discussed.

AlN exhibits very useful mechanical, thermal, chemical and piezoelectric properties. Thermal conductivity of AlN has been assumed to be about 320 W/mK. The thermal expansion coefficient in the c-plane and in the c-direction is as follows [24]:

a/ a   8.679  102  1.929  10  4 T  3.400  10  7 T 2  7.969  10  11 T 3 0

and Equation 2‎ .1

c/ c   7.006  102  1.583  10  4 T  2.719  10  7 T 2  5.837  10  11 T 3 0

Aluminum nitride is very hard material (~9 at Mohs scale) having a high melting temperature above 2000 ℃. Therefore, this material can be used at high temperature applications and could be deposited as the wear/corrosion resistant component. Due to its piezoelectricity, acoustic waves can be excited in AlN using electric fields, making it a common choice for RF signal processing and filtering; it is often the active material in surface acoustic wave (SAW) and bulk acoustic wave (BAW) devices.

The most common crystal structures for piezoelectric materials are in the form of Perovskite and Wurtzite. In Perovskite structures, an oxygen anion bonds with two metal cations. A net polarization of the structure can be achieved by metal cation shifting at room temperatures. The Wurtzite structure is hexagonal usually containing one anion and one cation.

These mentioned piezoelectric systems have some key differences beyond crystal structure that make them useful for various applications. Wurtzite piezoelectric materials have about two orders of magnitude inferior piezo-response compared to Perovskite materials but can maintain hysteresis-free behavior [25], with high-temperature thermal stability and easier production techniques. Wurtzite materials usually have higher elastic modulus which makes them mechanically more viable choices. These Wurtzite films can be deposited by CVD, sputtering and PLD. In order to compensate their lower piezoelectric constants compared to Perovskite structures, a highly textured film needed to be deposited through the process optimization.

At room temperature, Perovskite materials exhibit tetragonal structure with ferroelectric properties which diminishes above the Curie temperatures having a paraelectric orthorhombic/cubic structures which show no sign of polarization upon field removal. Some of the Wurtzite structures also behave in the same manner and will lose their piezo-response at higher temperatures. Therefore a material of the choice needs to be stable at higher temperatures without going to phase change during process and would be able to maintain its piezoelectric properties at high temperatures.

AlN exhibits high chemical and mechanical strength with high temperature stability and can maintain the piezoelectric signal at elevated temperatures [26-30]. The strain in the materials is usually increased before the initiation of the crack to a critical value and then propagates through the material releasing its strain energy. Therefore by embedding an appropriate high temperature piezoelectric material, one can predict the onset of failure and prevent it by change or repair of the die.

Any surface inhomogeneity or severe roughness increases the probability of electrical breakdown by local electrical field exceeding the breakdown strength of samples at these irregular points. Also it is well known that during the sputtering process, due to the difference of ion and sputtered atoms, the angle of incidence of adatoms at the surface could be changed producing shadowing effect which at the end the film would exhibit porous structure where these pores could be the source of inhomogeneity within the structure. A schematic representation of these potential electrical failure mechanisms is depicted in Figure 2‎ .2. Any rough surface or porous structure increases the probability of formation of high electric field region which could result in the formation of breakdown and failure of the smart material of choice.

Figure ‎2.2 Schematic of the films with induced imperfection introducing (a) rough surface, (b) porous film and (c) films with cavities [175].

2.2 Sensor Materials Selection Criteria

In order to monitor any changes of stress state within the die, a piezoelectric material with high piezoelectric constant needs to be deposited with highly textured structure. The orientation and texture of the films is a key pint in these piezoelectric materials since the response of the coating is very directional. The piezoelectric material needs to be an insulator with a non-conducting top coating to prevent and minimize any current leakage during processes. One other requirement for these piezoelectric materials is the purity of the films since any source of imperfection could result in altered resistivity and hence piezoelectric response. As mentioned before, any rough surface area or any cavity/void will generate the regions with higher electric field which would exceed the material breakdown strength and hence degrading the piezoelectric properties of the material. The ―smart‖ sensor materials are required to have high Curie temperatures above the operational temperature and be stable at the elevated temperatures. In the case of Al pressure die-casting a sensor material needs to be stable and responsive up to about 700 °C. Table 2‎ .1 lists the potential piezoelectric materials that might be the candidates for the die casting applications based on the above criteria.

Table ‎2.1 Physical properties of major piezoelectric materials widely used in devices [26–30].

Figures of Merit PZT AlN ZnO LiNbO3

Current response: e (C m-2) 31,f -14.7 -1.0 -0.7 -5.8

-1.2 -10.3 -7.2 N/A Voltage response: (GV m-1)

Coupling Coefficient (k )2 on Si p,f 0.2 0.11 0.06 0.02

Curie Temperature T (C) c ~300 ~1100 N/A 1210

CTE α (10-6 K-1) 7.2 4 5 11

Electrical Resistivity (Ω.cm) 109 1011-1016 108 1011

-6 -1 Note: CTE α (10 K ): 13 for H13 steel, 6 for Al2O3, 6.5 for TiAlN, and 10.1 for Ti.

High Curie temperatures and piezoelectric properties were selected as the main criteria to choose the candidates as well as the compatibility of the film with other existing layers. Based on these requirements, AlN and LiNbO3 were selected from Table 2.1 to have high thermal stability beyond aluminum die-casting process temperatures (above about 700°C).

Near metal thermal conductivity of AlN and its high mechanical, thermal, chemical and piezoelectric properties, make this material one of the best candidates for die-casting dies. In contrast to its high thermal conductivity its electrical band gap is about (6.2 eV) which makes it amongst insulators [31] and would be able to hold its piezoelectric signals without breakdown of voltage. Very high electrical resistivity of (1016 Ω.cm) [32] is reported in the literature to put this material as one of the candidates with high electrical breakdown resistance.

Based on the above mentioned properties such as high thermal stability, reasonable piezoelectric response and high voltage breakdown resistance, AlN would be an outstanding candidate for elevated temperature applications. In AlN atomic nature of the bonds are mainly

covalent with minimal ionic conduction at high temperatures which is detrimental for the smart sensor.

The mismatch of thermal coefficient of expansion with the substrate materials such as tool steel is large hence creating thermal stresses at the interface of the sensor layer and the substrate that would be a drawback for its application especially at higher operational temperatures. In order to resolve this issue, an intermediate layer with an appropriate thermal expansion coefficient could minimize the thermal stresses generated at the interface.

The piezoelectric activity in this material is found to be restricted to the perpendicular direction to the hexagonal arrays or (002) orientation, also known as the c-axis [33–35]. It is also the natural (energetically stable) growth direction of AlN film. C-axis closed pack planes in the AlN (HCP) structure have the minimum surface energy without any unsaturated bonds and hence these planes and directions are the natural direction of the growth and the flake-like structure is the common configurations for AlN crystallites.

2.3 Piezoelectricity Fundamentals

Piezoelectricity is a widely used physical effect that is derived from the Greek word of piezein, meaning to apply pressure. Limited classes of materials have the particular property that would produce electric field when subjected to external force. This effect is also reversible and materials of this class would change dimension upon electrical stimulation.

The Curie brothers (Pierre and Jacques) first discovered the piezoelectric effect in α-quartz crystal in 1880 [36]. They observed an electrical signal when applied a mechanical strain to the quartz crystal. The inverse piezoelectric behavior of this material was not discovered until a year later Gabriel Lippmann discovered this reverse phenomenon mathematically in 1881. They are 32 typical crystalline classes of which 21 are non-symmetric and 20 of these exhibit the piezoelectric effect.

The physical origin of piezoelectricity is explained by charge asymmetry within primitive unit cell resulting in the formation of net electric dipole. Adding up these individual dipoles over

the entire crystal produces a net polarization and generated electric field across the material. This process, as mentioned above, is reversible and a mechanical strain would be created upon an electric field application. Only the crystals that are non-centro-symmetric can have the piezoelectric properties. Crystals with center of symmetry like cubic crystal would not show piezoelectric response and the net electric dipole within the unit cell is always zero.

In general, an electric field always causes some amount of mechanical distortion in materials, since matter is composed of positively charged nuclei surrounded by the negatively charged electron cloud as in metals. The polarization induced by an electric field will alter the charge distribution and hence mechanical distortion. This phenomenon is called ―electrostriction‖, which is not reversible [37].

Electrostriction occurs in all materials, but its effect is extremely small and negligible in practical cases. In order to understand the piezoelectric behavior of materials, fundamental interrelation of different physical domains need to be explored as depicted in Figure 2‎ .3. Every physical response is connected to the other physical domains with a linear coefficient such as thermal expansion for connecting the mechanical strain with thermal stress.

The piezoelectric behavior is also explained based on this graph, as interplay between mechanical stress domain and electrical response for both direct and converse situations. In direct piezoelectric effect, mechanical strain/stress would cause the electrical signal while in the converse situations; mechanical strain/deformation is generated by the application of electrical modulations.

As was explained before, this piezoelectric behavior is only seen in a very limited number of materials which have crystal structures without center of symmetry in order to produce charge imbalance within the unit cell and therefore permanent dipole moments to hold the charge separated in the non-conductive media.

Insulators, like AlN with the band gap of 6.2 eV, do not have a large portion of mobile carriers but instead they form an atomic structure that contains large portion of bound charges which will not conduct through the material but can be reoriented in the presence of an applied

electric field as shown in Figure 2‎ .4 [38]. Bound charge (dipole) is a pair of point charges of equal and opposite charge Q separated by distance d that can pivot around its center point.

Figure ‎2.3 The relations between physical domains relating mechanical, electrical and thermal domains with their linear coefficients labeled on the lines connecting two [37].

Figure ‎2.4 Schematic representation of bound charges within a dielectric insulator upon electric field application [38].

Placing a dipole in electric field will generate rotation of the dipole around its central axis but will not result in charge conduction since the charges are bound in an insulating structure. The dipole moment, p associated with a pair of point charges is defined as: p Qd Equation 2‎ .2 where the vector d is from the negative to positive charge and if there are n dipole moments per unit volume, the polarization of the material is described as:

1 n Pp lim Equation 2‎ .3 Vol i1 i

The polarization in materials can be thought of an additional term called electric displacement, D which can be written mathematically as:

DEP0 Equation 2‎ .4

The relation between polarization and electric field is assumed to be linear. In this case, the polarization is written as the product of a constant and the electric field vector as follows:

PE(R 1) 0 Equation 2‎ .5

where  R is the relative permittivity which is unitless and larger than one. If the equation 2.5 is substituted in equation 2.4, the electric displacement can be written as:

DEEE0 ( RR  1)  0   0  Equation 2‎ .6 where the permittivity of a material is defined at the product of the permittivity of free space and the relative permittivity. So, the relation between electric displacement and electric field can be arranged as follows:

DEER 0  Equation ‎2.7

Considering an elastic material that is under uniaxial tension, at low values of applied stress (T), the material‘s response is linear before it begins to yield. The slope of stress-strain curve is defined as modulus of the material and the constitutive relation is described as:

S sT Equation ‎2.8

where s is called the mechanical compliance and S is the strain produced in the material. Then, a piezoelectric material being subjected to the same uniaxial tension can be considered. In addition to the above mentioned elongation, a piezoelectric material will produce a charge flow on metal electrodes which are placed at the two end of the piezoelectric specimen. This charge flow is caused by the motion and reorientation of electric dipoles within the piezo-material. The charge produced divided by the electrode area is called the electric displacement, D . Over a certain range of applied mechanical stress, there is a linear relation between electric displacement and applied stress, as illustrated in Figure 2‎ .5:

Figure ‎2.5 Relation between stress and electric displacement in piezoelectric materials [38].

The piezoelectric strain coefficient d is defined as the slope of this curve. At the high level of applies stress, the piezoelectric material becomes nonlinear due to saturation of electric dipole motion. But for the purpose of this project, the linear response of the materials is only considered. Piezoelectric materials also exhibit an inverse effect in which an applied electric field will produce a mechanical deformation. Upon electric field application, the dipole rotation will occur and the mechanical displacement is measured at the electrodes. At low values of applied electric field, the relation between E and D is linear. The constant of proportionality, called the dielectric permittivity (  ) is linearly relating the applied electric field to the electric displacement through the equation 2.7. Similar to the direct case, the application of electric field in piezoelectric material will result in the rotation of the dipoles and hence the production of strain in the material, as shown in Figure 2‎ .6. This linear portion can be expressed in the following equation:

S d E Equation ‎2.9

The properties of piezoelectric materials can be expressed mathematically as a relation between two mechanical variables, strain and stress, and two electrical variables, electric displacement and electric field as mentioned in equations 2.8 and 2.9.

Figure ‎2.6 Relation between electric field and strain in piezoelectric materials [38].

These expressions can be combined to a matrix form by writing the relation between strain and electric displacement (dependent variables) as a function of applied stress and applied field (independent variables):

S   s d   T       Equation ‎2.10 D   d   E 

The top row of equation 2.10, defines an equation for the converse piezoelectric effect, while the bottom one expresses the direct effect. By taking the inverse of above matrix, the relation between stress and electric field as dependent variables with respect to the independent variables, strain and electric displacement will produce the following relation: s1  d/ s T 1   d   S 22  S  1d / s 1 d / s Equation ‎2.11  2      E sd   d s   D ds/ 1  D  1d22 / s 1 d / s

The square root of the term ds2 /  is called piezoelectric coupling coefficient ( k d/ s ) which is the representation of energy conversion between mechanical and electrical domains and hence is always positive and bounded between 0 and 1. So the equation 2.11 can be written in a simpler form as follows:

T 1  s1 d 1 k 2  S   Equation ‎2.12  2 1 2 1   ED 1 k dk   

The expansion of the above constitutive expression in all three dimensions can be done by the introduction of coordinate system for a cube of piezoelectric material, as illustrated in Figure 2‎ .7. The above equations can be written in a three dimensional representation in a compact form based on the fact that the stress and strain tensors are symmetric. Therefore, constitutive equation can be expressed in a compact three dimensional forms depending on the variable choice, as defined below:

Figure ‎2.7 Three dimensional representation of a piezoelectric cube with its coordinate axes [38].

E Si s ij T j d ik E k Equation ‎2.13 T Dm d mj T j mk E k

E Ti c ij S j e ik E k Equation ‎2.14 S Dm e mj S j mk E k

D Si s ij T j g ik D k Equation ‎2.15 T Em  g mj T j   mk E k

D Ti c ij S j h ik D k Equation ‎2.16 S Em  h mj S j   mk E k

where scij,,, ij ij ij, are elastic compliance constant, elastic stiffness constant, dielectric permittivity constant, and dielectric impermittivity constants. dij,, e ij g ij and h ij are piezoelectric coefficients which can be converted to each other, also i and j have the values between 1 and 6 and m and n take the values between 1 and 3. The superscripts T,,, S D and E mean that the stress, strain, dielectric displacement and electric field are held constant. Depending on which constants are known or measurable, the particular pair of equations will be considered. When evaluating direct and converse piezoelectric effects, the constant most typically used and measured is the piezoelectric strain coefficient dij .

ET DS    d     Equation ‎2.17 TE   

The equation 2.13 can be simplified more depending on the boundary conditions of any given system. For example, in the case of thin films where one dimension (3rd dimension) is much smaller than the other two dimensions, assuming the situation like plane stress in the material, the state of piezoelectric material in the 3 direction can be simplified as follows:

S sE T d E 3 3 3 33 3 Equation ‎2.18 T D3 d 33 T 3 33 E 3

The free strain is defined as the strain produced in 3-direction when there is no resistance stress on the material. Under this mechanical boundary condition, T3 = 0 and the strain produced is:

S1 d 13 E 3 T3 0

S2 d 23 E 3 Equation ‎2.19 T3 0

S3 d 33 E 3 T3 0

The free strain (T3 = 0) and blocked stress ( S3 =0) are two common ways of characterizing the performance of piezoelectric materials. A general representation of the relation between stress and strain in the material as a function of the electric filed in the 3-direction is defined below:

EE T3 s 3 S 3 s 3 d 33 E 3 Equation ‎2.20

2.3.1 Thermodynamic Approach to Constitutive Equations

In this section, the basic thermodynamic principle for a reversible system is used to derive the constitutive equations of piezoelectric materials, as were expressed in previous section. From the first law of thermodynamics: dU dQ dW Equation ‎2.21

which states that the change in internal energy of a system is equal to the heat added and the work done by the system. For a reversible systems, the second law of thermodynamics states that the infinitesimal change in the heat can be expressed in terms of absolute temperature,  , and small change in the entropy, d which has the form of: dQ d . The external work in the case of piezoelectric materials is composed of mechanical and electrical work,

dW Ti dS i E k dD k Equation ‎2.22

and hence the internal energy of this system can be written as follows:

dU d  Ti dS i  E k dD k Equation ‎2.23

Assuming the linear electromechanical coupling in the piezoelectric system, the expression for Gibbs free energy can be used to reproduce the piezoelectric constitutive equations. The Gibbs free

energy of the above system has the form of, GUTSED  i i  k k and its derivative is obtained after the substitution of equation 2.23 into the complete derivative of Gibbs free energy:

dG  d  Si dT i  D k dE k Equation ‎2.24

GGG   then these following relations can be derived:   ,,SDik     . Substituting  TEik  these equations in the first part of equation 2.13 yields:

G E  sij T j  d ik E k Equation ‎2.25 Ti

Integrating the above equation 2.25 and define the constant of the integration by the application of the second part of equation 2.13, the Gibbs free energy of the system is derived,

11 G  sET TT  d T E   E E Equation ‎2.26 22ijij ikik klkl

Taking the derivative of equation 2.26 and compare the result with equation 2.24, dG  Si dT i  D k dE k (at constant temperature) will result in the following expression:

SDii ET dG ()() sTdEdTiji  ikk i  klk E  dTdE ikk i Equation ‎2.27

Compared to equation 2.13, the exact same piezoelectric constitutive equations are obtained with simple thermodynamic principles that were expressed previously

E T ( Si s ij T j d ik E k and Di d ik T k ik E k ).

2.3.2 AlN c-Axis Texture and Piezoelectric Response Model

The (002) plane of AlN has the highest growth rate of the crystal faces, so as deposition continues these grain can overtake crystallites with other crystal orientation. The effective piezoelectric response of an oriented film can be modeled considering coordinate reference frame (x,y,z) such that the z direction is perpendicular to the substrate; x and y lie in the plane of substrate

[39]. Using the reference frame R as a basis, the coordinate system Cgr (,,) x gr y gr z gr matching the orientation of individual crystallite is shown in Figure 2‎ .8. The coordination of any individual crystallite can be defined as a transformation of the reference frame by three successive angular rotation, ϕ, an in-plane rotation parallel to the substrate; θ, a tilt off the crystallite off of the substrate normal z, and ψ, a rotation about the c axis of the crystallite. The entire transformation between C (reference frame) and Cgr (individual grain frame) described above is contained within the matrix

cos cos cos  sin sin sin  cos cos  cos sin sin  sin  n  cos sin  cos  sin cos  sin  cos  cos  cos cos sin  cos ij  Equation ‎2.28  sin sin  sin  cos co s 

Figure ‎2.8 The transformation from the reference frame of the substrate to the coordinates of a crystal. The black arrows represents the new coordinate axes formed from a rotation of the gray axes [39].

By combining the transformation matrix (equation 2.28) with the piezoelectric tensor eij, the effective piezoelectric constant of an individual crystallite reduces to the below equation:

23 e33,crystallite ( ) (2 e 15  e 31 )(cos  sin  )  e 33 cos  Equation ‎2.29

The effective piezoelectric coefficient of an individual crystallite is a function of only tilt angle and a substitution for the strain piezoelectric constants dij can also be derived with the same procedure. In order to find the total piezoelectric response of the thin film, the contribution of each crystallite in the film must be integrated. If the distribution according to XRD results assumed to be Gaussian, then the total effective piezo-response would have the form of:

 e p(,)()  e  d  33,eff 33, crystallite   2  1 2 where,(,) p  e 2 Equation ‎2.30  2 Erf   2  2 and   FWHM 8ln 2

FWHM is the full-width half-maximum (FWHM) of an XRD rocking curve centered at the (002) peak. These equations above show that the piezoelectric response of the entire film is only a function of tilt angle. The above integration in equation 2.30 is relatively complex to solve and an exact analytical solution to the above integral can be found using computational software such as

Mathematica. The function e33,eff is plotted over the entire range of FWHM in Figure 2‎ .9.

Figure ‎2.9 The effective piezoelectric constant e33 as a function of the rocking curve FWHM showing the abrupt decay of piezoelectric response as a function of increasing FWHM [39].

The above model shows that the piezoelectric response of the film is strongly dependent on the c-axis orientation of individual crystallites to the substrate normal. However, there are some limitations to this model. The calculations assume no interaction among grains or any formation of voids during deformation. Assuming that the material of interest is highly aligned and have a true columnar dense structure, this model would be predicting with less error.

Also the modified Gaussian in the model does not completely represent true random orientation of the film. Nevertheless, this simplified mathematical model illustrates the importance of c-axis texture in AlN thin films to obtain the maximum piezo-response and also this model agrees with experimental data especially at low values of rocking curve full width half maximum

(FWHM); agreements with this model have been shown in the case of BaTiO3 [40]. Results for the AlN thin films on silicon also show that a rocking curve FWHM better than about 6 is necessary for good piezoelectric characteristics [41].

2.3.3 Piezoelectric Materials

A non-centrosymmetric lattice with charge separation and unbalanced distribution is needed to generate piezoelectric response. The most common structures that exhibit this non central symmetry are Perovskite and Wurtzite. In the case of Pervoskite crystals, oxygen is an anion bonded by two metal cations.

In Pervoskite tetragonal structures at room temperatures, the oxygen anions will create a polarization by moving the metal cation from its position as shown in Figure 2‎ .10. Barrium titanate (BaTiO3) is a very common Pervoskite piezoelectric material which at room temperature exhibits tetragonal structure with Ti cation being shifted due to the presence of oxygen anion. But at temperatures higher than Curie point, it transforms to the cubic lattice which is centro-symmetric with no net dipole moment. Wurtzite structures are hexagonal closed packed at room temperature (HCP) and consist of binary anion and a metal cation as depicted in Figure 2‎ .10

Figure ‎2.10 Crystal structure of the Perovskite ferroelectric BaTiO3 (A) high temperature paraelectric cubic phase (B & C) room temperature, ferroelectric, tetragonal phases, showing up and down variants [42].

techniques [43]. Wurtzite materials usually have higher elastic modulus which makes them more mechanically viable structure.

Figure ‎2.11 Typical wurtzite materials having hexagonal closed-packed (HCP) structures with tetrahedral interstitial sites of anion lattice filled by cations [176].

At room temperature, Pervoskite materials exhibit tetragonal structure with ferroelectric properties which diminishes above the Curie temperatures having a paraelectric orthorhombic/cubic structures which show no sign of polarization upon field removal. Some of the Wurtzite structures also behave in the same manner and will lose their piezoe-response at higher temperatures. Therefore a material of the choice needs to be stable at higher temperatures without going to phase change during process and would be able to maintain its piezoelectric properties at high temperatures.

2.3.3.1 Lead Zirconate Titanate (PZT)

Lead zirconate titanate (PZT) has been used widely since 1950 in thin film and bulk applications especially to make Surface Acoustic Wave (SAW) devices and sensors both in bulk and thin film form. PZT has been prepared with various methods from CVD to PVD with very

large dielectric and piezoelectric constants. Even though the melting temperature of PZT is fairly high and above 1650 C but its Curie temperature is only about 330 and will transfer to cubic non-ferroelectric above this temperature.

Both tetragonal and cubic structures of PZT, shown in Figure 2‎ .12, have extremely high piezoelectric coefficients at about -14.7 C/m2 and also coupling coefficient of 0.73 [44] which is more than an order of magnitude higher that Wurtzite piezoelectric materials. Sputtering of PZT is difficult because it requires the control of the stoichiometry of lead, zirconium, and titanium accurately. The further limitation of PZT is its low Curie temperature.

Figure ‎2.12 Lead zirconate titanate (PZT) crystal structure showing both (1) cubic and (2) tetragonal phases [44].

2.3.3.2 Lead Titanate (PbTiO3)

Lead Titanate (PbTiO3) has similar structure to PZT and is Perovskite in nature and used in ferroelectric applications and high frequency transducers. The Curie temperature of this material is slightly higher than PZT (490 C ) but with lower coupling coefficient. Figure 2‎ .13 shows the typical PbTiO3 Perovskite structure.

Figure ‎2.13 Lead titanate Perovskite structure with lead atoms in tetragonal lattice arrangement titanium atoms in octahedral interstitial site [176].

Similarly to PZT, PbTiO3 is not an appropriate candidate for die casting application due to its incompatibility with the rest of coating system and more importantly its low Curie temperature much below the die casting operational temperatures.

2.3.3.3 Zinc Oxide (ZnO)

Zinc oxide is widely used in diverse applications from biomedical to semiconductor industry and also as a pigment. This piezoelectric material has very similar structure to AlN and is used as bulk acoustic wave (BAW)resonators in telecommunications [45]. The direct band gap of 3.37 eV at room temperature makes ZnO one of the best candidates for band gap engineering applications. Different type and structures of ZnO have been produced on silicon nitride for micro-machined devices [46] as depicted in Figure 2‎ .14. As mentioned

before, ZnO has Wurtzite structure with Curie temperatures of 350 °C and coupling coefficient of 0.17 shown in Figure 2‎ .15. ZnO can be deposited by PVD, CVD and many other methods [47–49], with high quality.

Figure ‎2.14 ZnO nanostructures synthesized by thermal evaporation of solid powders (a) nanocombs, (b) tetraleg, (c) hexagonal disks, (d) nanopropellers, (f) nanospiral, (g) nanosprings, (h) single crystal nanoring and (i) combination of rods, bow and ring [46].

Figure ‎2.15 Zinc oxide (ZnO) wurtzite hexagonal close packed structure [176].

ZnO cannot be a suitable candidate for high temperature applications due to its low Curie temperature and also low coupling coefficients.

2.3.3.4 Gallium Nitride (GaN)

Gallium Nitride like its competitor ZnO has been studied widely in recent years for applications ranging from semiconductors to transistors. Due to its unprompted piezoelectric response can be utilized to cultivate extremely sensitive gas sensors, polar liquids and pressure sensors [50].

GaN has Wurtzite crystal structure similar to ZnO and AlN with lower piezoelectric constants compare to its rivals as shown in Figure 2.16. The Curie temperature is very low at about 180 C with coupling coefficient of 0.13 which limits its high temperature applications as piezoelectric materials. Several methods have been used to deposit GaN which ranges from molecular beam epitaxy (MBE) to PLD and sputtering [51, 52].

Figure ‎2.16 Piezoelectric and spontaneous polarization effect in Ga (Al) or N‐face AlGaN/GaN heterostructures [50].

2.3.3.5 Lithium Niobate (LiNbO3)

LiNbO3 is one of the candidates for high temperature application since its Curie temperature is about 1210 C with electromechanical coupling coefficient of 0.23 to be used for SAW transducers. Lithium Niobate has hexagonal closed packed (HCP) crystal structure below its Curie temperature but its trigonal structure also been made synthetically [53]. The structure is formed by the array of oxygen atoms in the HCP lattice with 2/3 of octahedral sites being occupied evenly with lithium and niobium cations. At temperatures above Curie point, lithium atoms transfer to oxygen position and niobium atoms will center between anion layers to make the crystal non-polar.

Figure ‎2.17 Lithium Niobate's trigonal crystal structure with Niobium atoms in blue, Lithium atoms in green and Oxygen atoms in red [176]. .

One of the drawbacks of this material is its complexity in bulk and thin film production. Single crystals of lithium niobate can be made with Czochralski method from the melt with proper seeding and rotation, and then it could be sliced similar to silicon wafers [54]. In thin film

form it has been deposited by MOCVD, low pressure CVD and sputtering in rf mode. As mentioned before, its high temperature stability could make this as a good candidate for our application in die-casting of aluminum but its complex processing methods and incompatibility with the rest of coating system and also its low thermal conductivity of 5.2 W/mK would limit its application [55].

2.3.3.6 Aluminum Nitride (AlN)

Aluminum nitride is one of the high temperature materials that is been developed recently especially as acoustic filters in GHz frequency range. This material has been synthesized since 1980 for mostly optical and electronic applications. It has been used in optoelectronic, piezoelectric filters in SAW devices, and also due to its high thermal conductivity, it is one of the heat sink material in semiconductor industry.

AlN has very high Curie temperature of 1150 C and coupling coefficient of 0.35 with reasonable piezoelectric constant that would be suitable for our specific application at elevated temperatures. However, the CTE mismatch is large with the tool steel substrate that can be compromised by the application of intermediate layer or adhesion film. The oxidation resistance of this material is about 800 [56] which is above the temperature of aluminum die casting process.

2.3.4 Piezoelectric Measurements Methods

Piezoelectric properties of thin films can be obtained by two main types of techniques: quasi-static and dynamic measurements [57]. The quasi-static methods of measurements are typically used for non-resonant applications such as sensors and actuators. In contrast, the dynamic measurements are usually for the piezoelectric thin films used in the generation or detection of high frequency bulk acoustic waves (BAW) or surface acoustic waves (SAW).

Since the topic of this project is focused on the high temperature piezoelectric sensor, only the quasi-static measurement techniques are summarized. These quasi-static measurements are based on either the direct or the converse piezoelectric effect. In the case of direct piezoelectric measurements, the voltage produced by an applied stress is measured and in the converse piezoelectric measurements, the displacement produced by an applied electric field is obtained.

2.3.4.1 Direct Piezoelectric Measurements Techniques

If a piezoelectric film with Wurtzite structure is assumed, the constitutive equations of state (equation 2.13), can describe the direct piezoelectric effect. If the stress is applied only in the c-axis direction, without any external electric field, the electric displacement is given by:

D1 dT15 5 0 D d T 0 Equation ‎2.31 2 15 4   D3 d31 T 1 d 31 T 2 d 33 T 3 dT33 3

Then the piezoelectric coefficient d33 can be calculated knowing the applied stress T3 and measuring the electric displacement D3 . However, when the stress is applied perpendicular to the c-axis direction, the electric displacement is given by:

D1 dT15 5 0  D d T 0 Equation ‎2.32 2 15 4   D3 d31 T 1 d 31 T 2 d 33 T 3 dT31 2

Therefore, no electric displacement should be obtained in direction perpendicular to the c-axis direction. In the following sections, direct piezoelectric techniques are summarized.

a) Pneumatic Pressure Rig Method

In this method, a calibrated perpendicular force to the film surface is applied and the generated charge is collected. The film needs to be sandwiched between the bottom and top electrodes. This method does not represent the piezoelectric coefficients of the free sample, but an effective coefficient since the film is clamped to the substrate [58]. The major problem with this method is believed to be the simultaneous bending of the sample and even a small substrate bending can generate very large biaxial stress in the piezoelectric film which produces large amount of additional electric charge collected in total. Moreover, the small thickness of the sample makes stress alignment a difficult task.

To solve these issues, a method using pneumatic pressure rig is designed to apply a uniaxial stress to the piezoelectric thin film as illustrated in Figure 2‎ .18. The samples are placed between two fixtures with cavities both above and beneath it and O-rings are used on both sides to seal the sample. By introducing high pressure gas inside, force is imposed on both sides of the sample uniformly introducing uniaxial compressive stress. The application of pneumatic pressure would not introduce bending even if the initial curvature exists within the sample.

The induced charge is collected using a charge integrator which converts the collected charge into a variation of voltage on a capacitor of known size placed in series with the sample. The effective piezoelectric coefficient of the film is calculated using the equation below:

D3 Q d33  Equation ‎2.33 TPA3 .

where D3 and T3 are the electrical displacement and mechanical stress in thickness direction,

P is the change of cavity pressure and Q is the measured electric charge induced by the pressure change on the area of the electrode A . Furthermore, since the film is clamped to the substrate, an effective piezoelectric coefficient can be obtained using some assumptions [59].

Figure ‎2.18 Experimental setup for d33 measurements by the method of the pneumatic pressure rig [58].

b) Metallic Rod Method

This method does not require the deposition of electrodes on the film surface and has been designed for quantitative check of the piezoelectric film [57], [60]. A low frequency transducer bonded on one end of the metallic rod produces longitudinal wave propagating in the rod. The other end of the rod is in contact with the film using a droplet of water to ensure a good mechanical energy transfer from the rod to the film, as shown in Figure 2‎ .19. In order to evaluate the piezoelectric coefficient, the measured voltage is compared with a known piezoelectric plate in the same conditions.

Figure ‎2.19 Experimental setup for d33 measurement by the method of the metallic rod [57].

c) Cantilever Beam Method

This method allows the measurement of transverse piezoelectric coefficient based on the direct piezoelectric effect. The film is deposited between two electrodes on a cantilever beam. The free part of the cantilever is displaced and then is suddenly released. The corresponding damping oscillation of the cantilever at its natural frequency is appeared between the electrodes. A variation of this technique consists of measuring the electrode charge at a frequency imposed by an external exciter moving the free edge of the beam as depicted in Figure 2‎ .20.

Figure ‎2.20 Schematic of a setup for the measurement of the transverse piezoelectric coefficient by the method of the cantilever beam [57].

2.3.4.2 Converse Piezoelectric Measurement Methods

Constitutive equation of state in the case of inverse piezoelectric effect for a wurtzite structure was derived previously (equation 2.13). Then if the electric field has only c-axis component with no external stresses, the strain is derived as:

S1 dE31 3 dE31 3  S dE dE 2 31 3 31 3

S3 dE33 3 dE 33 3 Equation ‎2.34 S4 dE15 2 0 S dE 0 5 15 1  S6 0 0

Therefore, the piezoelectric coefficient d33 can be calculated knowing the electric field E3 and by measuring the strain in the same direction ( S3 ). In the converse piezoelectric measurement techniques, a defined voltage is applied across the film surface and the strain is measured. As the magnitude of the displacement to be measured is in the range of angstroms or even picometers, very sensitive methods, such as interferometry or piezoresponse force microscopy (PFM), capable of measuring such small deformation must be used. a) Double-Beam Interferometry

This reliable and widely used interferometry technique is the double-beam Mach-Zender type, in which the displacement difference of the two faces of the sample is measured as illustrated in Figure 2‎ .21, eliminating the contribution of substrate bending motion. This method has been the reference method in the development of the other techniques.

Figure ‎2.21 Schematic of the double-beam interferometer (top) and principle of elimination of bending motion of the substrate (bottom) [57].

b) Michelson Interferometry

A single-beam interferometry can also be used in the same manner as mentioned above but with shining single coherent beam source. A Michelson interferometer can be used to measure the piezoelectric coefficients upon application of electric field across the bottom and top electrodes. In this part a brief overview of the optics and the interferometry background is covered. Wave optics is based on Maxwell‘s equations and can be expressed as follows:

 .Em  0

.0Bm B Equation 2‎ .35 E   m t E BJ  m   0 0t 0

These equations represent the Coulomb‘s law, the nonexistence of magnetic monopoles,

Faraday‘s law and Ampere‘s law with extra term introduced by Maxwell, respectively. For a light in a vacuum these equations become simplified and they have the form as follows:

.0E .0B B E   Equation 2‎ .36 t E B   00t

Therefore one can derive the wave equation from this simplified version of Maxwell‘s equation in vacuum. To begin, a curl operation needs to be applied to the Faraday‘s law in vacuum:

B 2E curl()()() E   curlm   curlB   Equation ‎2.37 t  t00  t 2

In the case of vacuum, the divergence of electric field is zero and the wave equation becomes:

2E E   0 Equation ‎2.38 00t 2

and in the one dimensional situation, like a scalar wave propagating in z-direction, the equation simplifies to:

22EE1 0 Equation ‎2.39 z2 c 2 t 2

One solution to the above partial differential equation has the following form:

E acos[ t k z ] Equation ‎2.40

where a is the amplitude,  is circular frequency and k is the wave vector. Now, if two monochromatic waves propagating in the same direction interfere at a given point, then the total electric field at this point is: EEE12 and if the two waves have the same frequency, like in the case of Michelson interferometer, the intensity at this point is:

2 IAA12where A1 a 1exp( i 1 ) , A2 a 2exp( i 2 ) are the complex amplitudes of the two waves and  2/  zc is the phase. So the intensity has the form of:

22  IAAAAAA1  2  1 2  1 2 1 Equation ‎2.41 2 IIIII1  2 2( 1 2 ) cos 

where   12   is the phase difference between two waves. If p is the difference in the optical path, then the order of interference is Np /  . The intensity has the maximum when N m,  p  m ,    2 m  , m is an integer, and its minimum value exist when, N(2 m  1)/2,  p  (2 m  1)/2,    (2 m  1)  . One of the convenient measures of contrast of the interference pattern is the visibility of the fringes, which is defined as follows:

II 2(II )1/2 K max min 12 Equation ‎2.42 IIIImax min() 1 2

In order to measure the displacement of the surface of AlN thin film, which is clamped to the substrate, an alternating electric field is applied and the output signal is calibrated based on a reference actuator. Michelson Interferometry is one of the sensitive techniques that can resolve the displacements down to 10-3 angstrom as schematically shown in Figure 2‎ .22. The phase shift is given by:  2kd  where d is the optical path-length difference of the two beams. The interference light intensity according to equation 2.43 can be expressed as follows:

Figure ‎2.22 Schematic representation of a Michelson Interferometer.

1 2 I IBABA  I 2( I I ) cos(2 k  d ) Equation ‎2.43

where I A and I B are the light intensities for the beams reflecting from movable (sample) mirror and a fixed mirror, respectively. After maximizing the equation 2.43, the piezoelectric coefficients can be calculated with following equation, where Vapp is the voltage applied across the film:

Iout d33  Equation ‎2.44 k() Imax I min Vapp

2.4 Low Density Plasma

Plasma is partially ionized gas with equal numbers of positive and negative charges. The degree of ionization is typically very small for low density plasmas in the order of 10-4 to 10-6. To

maintain the steady state of electron and ion densities, an ionization process is required to balance the recombination process. The positive particles are mainly atoms or molecules that have lost one or more electrons and the majority of negatively charged particles are free electrons. At steady state the free electrons obtain sufficient energy from the electric field potential difference to produce impact ionization of the gas molecules. In the case of collisions between electrons and neutral atom having masses men and m , respectively, by conservation of momentum and energy, the ratio of energy transferred has the following form:

En4 m e m n 2  2 cos  Equation ‎2.45 Ee() m e m n

where  is the incident angle of collision with respect to the line joining the center of mass of two colliding species. Since mn is much larger than me , and assuming head-on collision, then the energy transferred from a collision of an electron to a neutral atom is in the order of104 [61]. Therefore, electrons just change their angle of incident and not enough energy is transferred for neutral atoms to be moved. A typical electron temperature is in the range of 25 eV whereas in the case of ions and neutral atom, is only about room temperature (0.025eV ). These high energy or hot electrons have enough energy to excite electron-molecule reactions. In order to produce same reactive species without plasma would require temperatures exceeding1000 C .

As mentioned previously, the average speed of electrons is enormous compared to those of the ions and neutrals, due to both the high temperature and low mass of the electrons. If an electrically isolated substrate emerges into the plasma, initially it would be struck by electrons which have much higher speed compared to neutrals and ions. Therefore, the substrate immediately starts to build a negative charge and hence negative potential with respect to the plasma which is called floating potential and is defined by equation 2.46:

kTei m Vf  ln( ) Equation ‎2.46 22em e

where me and mi are the electron and ion masses respectively and Te is the electron temperature. As explained before, the ion mass is about 4 orders of magnitude higher than that of an electron, the floating potential is always a negative value several times the electron temperature in volts. Then the quasi-random motions of the ions and electrons in the vicinity of the object are disturbed. Since the substrate charges negatively, electrons are repelled and ions are attracted.

Thus, the floating potential of substrate is always more negative that the plasma potential (VVfp

). Since electrons are repelled by the potential differenceVVpf , it follows that the isolated substrate will acquire a net positive charge around it known as space charge or sheath. Since the electron density is lower in the sheath, it does not glow as much and it appears as a dark space.

2.4.1 Sputtering Deposition Process

Plasma in sputtering method consists of partially ionized Ar+ (~10‐4 ionization, mostly neutral species) due to potential drop, accelerates the inert gas ions toward a cathode. Upon the ion collisions with target material solid particles are ejected from the cathode and then deposited on substrate and chamber wall. Grove first observed sputtering in a DC discharge tube in 1852. Since then, sputtering has been widely used for surface treatments, etching, thin film deposition and surface analysis. The incident particles are usually inert gas ions but any ion, neutral atom, molecule or even photon can be injected if there is sufficient energy [62]. In sputter deposition as schematically illustrated in Figure 2‎ .23, the impact of the energetic particles on the target surface kinetically knocks one or more of the surface or near-surface atoms from their equilibrium positions.

Figure ‎2.23 A schematic representation of the physical sputtering processes [63].

These ejected atoms, which have received enough kinetic energy from the initial particles, would be attracted to the target material and hence causing more ejection of atoms from the cathode surface [64]. The interaction of energetic ions with surfaces would create a variety of particles such as secondary electrons, neutrals, photons, x-ray, and implantation of atoms into the substrate. The emission of secondary electron is essential to maintain the discharged plasma. The typical kinetic energy of sputtering deposition process is in the order of1 10eV , whereas, in the case of evaporated particles is about an order of magnitude lower. This added energy for the species in sputtering process can modify the nucleation and growth processes, improve film adhesion, increase film density and ease the film texturing. The most common modes of sputtering include the DC discharge, pulsed DC, RF, microwave and dielectric barrier discharge [65]. One of the simplest configurations of the sputtering is the diode sputtering where two counter-electrodes are separated in vacuum chamber and a DC voltage

potential difference is applied across the electrodes to break down the gas as shown in Figure 2‎ .24.

Figure ‎2.24 Geometry and typical voltage profile of the diode configuration dc glow discharge showing the plasma potential and the cathode and anode sheath [66].

The plasma between the cathode and anode consists of different distributions of potential, space charge and current density [67], [68]. The negative column region is the brightest part of the discharge where the electric field is close to zero and most of ionization collisions take place. In sputter deposition process, the electrode separation needs to be small in order to increase deposition rate and the anode (substrate) is usually located in the negative glow region and the positive glow is avoided [69]. In simple diode discharge, the ions move only relatively small fraction of the cathode sheath before they experience a charge exchange collision and are no longer accelerated by the field. As a result, the average energy of ions incident at the target is much less than 1 eV and the glow discharge sputter yields are less than expected.

2.4.2 Magnetron Sputtering, Balanced and Unbalanced

Even though diode sputtering is still used today due to its simplicity and relative ease of fabrication of target materials, it has several disadvantages. In the diode sputtering, some of the electrons coming out of the target material are not participating in glow discharge process and hence result in lower deposition rate. These wasted electrons are accelerated towards the chamber wall and substrate resulting in overheating and radiation. With the introduction of magnetic assistant, the sputtering achieved a widespread use in research and industrial applications.

In magnetron sputtering magnetic field lines would guide and focus the charged particles such as electron and ions as shown in Figure 2‎ .25.

Figure ‎2.25 Schematic illustration of typical magnetron sputtering showing the confinement of electrons to the target region [66].

The use of magnetic field behind the target material would address the electron loss in diode sputtering by capturing the escaping electrons and confine them to the vicinity of the target. The ion current density in the magnetron sputtering is increased by an order of magnitude

over conventional diode sputtering systems and resulting in higher deposition rates at lower pressures. Since the magnetron sputtering can work in lower pressures, it could create cleaner films at lower substrate temperatures. The magnetic field behind the target is designed to trap the electrons in the vicinity of the target surface. This could be done by using traditional balanced magnetron sputtering; where the north and south poles of the magnets are balanced meaning have equal strength [70]. One of the disadvantages of balanced magnetron sputtering is the plasma is mostly in the vicinity of the target resulting in low plasma density over the substrate area as depicted in Figure 2‎ .26.

Figure ‎2.26 A comparison of the magnetic configuration and plasma confinement in (a) dc balanced magnetron sputtering and (b) dc unbalanced magnetron sputtering.

The ion current density drawn at the substrate in balanced magnetron sputtering is less than 1 mA/ cm2 [70]. The invention of unbalanced magnetron sputtering by Window and

Savvides in 1986 [71] enhanced the plasma density further. In the unbalanced magnetron configuration, the outside magnets are stronger than the one in the center resulting in the expansion of magnetic field lines and pulling the plasma away from the vicinity of target surface towards the substrate. The effect of unbalanced magnetic field is to trap the hot secondary electrons from escaping the target surface and further enhance the ionization of neutral atoms to produce a greater number of ions and electrons near the substrate region. The ion current density in unbalanced magnetron sputtering is increased to 2-10 mA/ cm2 which is much higher than the balanced magnetron sputtering systems [72]. In order to deposit compound, reactive sputtering can be used to deposit most nitrides, carbides, and oxide thin films with controlled composition and texture employing appropriate deposition parameters [73–76].

2.4.3 Closed Field Unbalanced Magnetron Sputtering (CFUBMS)

It has been proposed by Sproul (1990) and Tominaga (1990) that the ion current density can be even further increased to 5-20 compared to unbalanced magnetron sputtering (2-10 ) [77], [78]. A comparison of a mirrored and closed field magnetron configuration is shown in Figure 2‎ .27. Using the closed field magnetron configuration, the deposition zone in which the substrates are located is surrounded by linking magnetic field lines. This traps the plasma region, prevents loss of escaping electrons hitting the chamber wall and resulting in higher plasma density (ion current density) at the substrates.

The closed field unbalanced magnetron sputtering (CFUBMS) configuration is a complex system where the neighboring magnets are opposite of polarity and arranged so that the magnetic field between the magnetrons is totally closed. This is only possible if one uses more than two targets in the chamber where the outer magnets are stronger than the inner magnets. In unbalanced configuration, electrons are less confined to the target surface and hence can escape towards the substrate are to extend the plasma of reactively sputtered species. A schematic of this unbalanced closed field configuration is shown in Figure 2‎ .28.

Figure ‎2.27 diagram showing the mirrored (a) and closed field (b) magnetic configurations where in the mirrored configuration similar poles face each other and in closed field configuration the opposite poles face each other [78].

Figure ‎2.28 Schematics of four unbalanced magnetrons arranged in closed field configuration. An EQP mass spectrometer is also put inside the chamber to analyze the species in the plasma [79].

2.4.4 Reactive Sputtering

In reactive sputtering, the metallic target is sputtered in a reactive atmosphere as illustrated in Figure 2‎ .29. As the sputtered atoms are deposited onto the substrate, they react with the reactive gas molecules forming the desired compound. There are several advantages for reactive sputtering of metal target over conventional sputtering of compound targets in inert gas environment. Metal targets can easily be machined and due to the high thermal conductivity of metal targets, the cooling is more efficient and higher power densities can be applied.

In the region, where the partial pressure of the reactive gas is rather low, the target is only marginally covered with compound and process behavior is very identical to the sputtering of pure metal target in inert environment with high sputtering yield. As the reactive gas flow is

slowly increased, the target surfaces becomes partially covered with the insulating compound (or becomes poisoned) and since the sputtering yield of the compound is usually much lower than that of the metal, the overall deposition rate decreases.

Figure ‎2.29 An illustration of the reactive sputtering process, where in this case aluminum reacts with the nitrogen to form aluminum nitride [80].

Normally a decrease in the discharge voltage is also observed because the secondary electron emission coefficient for compounds is usually higher than that of pure metal surfaces. This would lead to arcing problem in which the micro-particles will be ejected from the target and incorporated into the deposited films and generate non-uniform film with inhomogeneity and defects as well as reduced deposition rate [81]. To resolve these issues, the use of radio frequency (RF) to sputter the materials was developed in 1960‘s. However RF sputtering is not commercially feasible due to its lower deposition rate, high cost and the complexity of network matching. An alternative technique using pulsed DC plasma has been developed. The pulsed potential neutralizes the positive charge on the target surface and eliminating the arcing by

controlling the pulsing parameters. Figure 2‎ .30 exhibits some typical target voltage potential waveforms used in dc and pulsed dc magnetron sputtering.

Figure ‎2.30 Target voltage waveforms when operated in (a) continuous DC, (b) unipolar pulsed mode, (c) asymmetric bipolar pulsed mode, and (d) symmetric bipolar pulsed mode [79].

The continuous target voltage in DC mode is either turned off periodically in the unipolar mode or more commonly, switched to a positive voltage in the pulsed DC magnetron sputtering.

During the normal pulse-on period ( on ), the negative voltage is applied to the target. This negative voltage application is periodically interrupted by a positive pulse voltage with a period of ( rev ). This positive voltage is either reversed to a smaller positive voltage than nominal negative target voltage in the asymmetric bipolar mode or to the similar magnitude of negative voltage in symmetric bipolar systems. The duty cycle is also defined as: on/()  on  rev and the full range of frequencies may not be available due to the power supply limitations.

During the reversed positive pulsed DC sputtering process, the charge built up on the insulating material is discharged therefore breakdown and arcing is eliminated. With precise control of partial pressure of reactive gas, high deposition rate can be achieved in depositing insulating films. A comparison of microstructure of alumina films deposited by pulsed closed

field unbalanced magnetron sputtering (P-CFUBMS) and normal DC magnetron is shown in Figure 2‎ .31. A fully dense and smooth alumina film was produced by P-CFUBMS compared to a porous structure with micro-particles covering the film surface [70], [82], [83].

Figure ‎2.31 SEM micrographs of fracture section of aluminum oxide film deposited by (a) DC reactive sputtering and (b) pulsed closed field unbalanced reactive sputtering [70].

2.4.5 High Power Pulsed Magnetron Sputtering (HPPMS)

High Power Pulsed Magnetron Sputtering (HPPMS) or High Power Impulse Magnetron Sputtering (HIPIMS) is a relatively new magnetron sputtering technique. It was first introduced by Kouznetsov in 1999, by generating a high power pulse at a low duty cycle to produce peak power densities as high as several kW/cm2 [84], [85]. Generally, the maximum power used in dcMS is limited to 3 W/cm2 due to the thermal load on the target during deposition, but decreasing the duty cycle eliminates this concern [86]. Several authors have shown that a variety of materials can be used with HPPMS to produce dense films superior to films deposited by cathodic arc evaporation [87], [88].

Since Kouznetsov [84], [85] developed HPPMS, there have been numerous studies that compare properties of films deposited by HPPMS to films from other PVD techniques. The

general consensus is that HPPMS films are superior in many aspects. Bobzin et al. [89] evaluated films deposited at the same average power from dcMS, middle frequency magnetron sputtering (mfMS), and HPPMS. SEM micrographs of the cross-sections are shown in Figure 2‎ .32. Both films deposited from dcMS and mfMS exhibit a strong columnar microstructure and increased surface roughness, while the HPPMS film is much denser.

Figure ‎2.32 SEM micrographs of (Ti,Al)N coating cross-sections from a) dcMS, b) mfMS, and c) HPPMS [89].

Bobzin also deposited films on cutting inserts that had perpendicular edges to determine deposition rate on surfaces orthogonal to the target as shown in Figure 2‎ .33. It has been found that the coating on the perpendicular face was 56% less thick than that parallel to the target for dcMS whereas, it was only 28% less thick for HPPMS. This is most likely due to the high metal ion flux and energies of the deposited ions. They performed the same experiment on cutting

tools with edges perpendicular and parallel to the target using (Ti, Al, Si)N films and determined that the silicon content was lower on the perpendicular side for dcMS and mfMS films, but not HPPMS. The HPPMS films also provided a higher hardness value after oxidation tests.

Figure ‎2.33 SEM micrograph of (Ti,Al)N coating cross-section on a cutting edge deposited by dc-MS (top) and HPPMS (bottom) [89].

2.4.6 Modulated Pulsed Power (MPP) Magnetron Sputtering

Zpulser LLC from Mansfield, MA has recently developed MPP, a novel sputtering technique that applies the basic principles of HPPMS, but also has the ability to create and control stages within a single pulse [90]. These stages can be tailored to generate a combination of weakly-ionized and strongly-ionized plasmas for portions of each pulse. The pulse duration, voltage rise time, and voltage fall time can be controlled for each stage to create ―micropulses‖. The peak power, peak current and plasma density rival that from a HPPMS power supply. The resulting pulse shape is very different than that of HPPMS as shown in Figure 2‎ .34. MPP has shown to increase deposition rates above that of dcMS while maintaining optimum properties by optimizing the micropulses within the main pulse [98]. The MPP power supply can also deposit ―multilayer‖ coatings of the same species by having two different pulse shapes. Figure 2‎ .35 displays the differences between typical HPPMS and MPP pulses and common parameters. ( ( (a) (b)

Figure ‎2.34 Comparison of pulse shapes from (a) HPPMS and (b) MPP [89], [91].

Figure ‎2.35 Typical target waveforms and parameters for (a) HPPMS and (b) MPP [92].

2.4.7 MPP Pulse Characteristics

A single MPP pulse typically has four characteristic steps indicated in a typical pulse waveform as used in this thesis and shown in Figure 2‎ .36. The first two steps are used to initiate and maintain a plasma without generating arcs that may damage the power supply whereas, the final two steps use higher voltage and current to create high levels of ionization. Any of these steps may be modified to tailor coating properties.

The first step is the ignition of low power discharge which is indicated by the sharp spike in voltage. In this particular waveform, the step is about 300 μs long. This step is commonly observed in all magnetron sputtering techniques when trying to first generate plasma, although the voltage spike is much faster for dc magnetron sputtering and HPPMS than MPP. The slope and shape of the voltage spike is dependent on the working gas and target materials. The power and current remain near zero until the last portion of the step after which, the spike in current indicates that plasma has been generated and the discharge is maintained.

The second step begins after the spike in current between the ~300 – 500 μs portion of the waveform. This is a low power discharge with the voltage, current, and power remaining relatively low. This step is similar to a DC discharge.

The third step is a transient stage from a low power discharge to a high power discharge. This step is between 500 – 750 μs for this particular waveform. The current, power, and voltage rise because the voltage ―on‖ and ―off‖ time, or the ―micropulse‖ is modified. This waveform has an ―on‖ time of 14 μs and an ―off‖ time of 6 μs. During the ―off‖ time, the electrons diffuse to the target and chamber walls and the electron density decays with time. The ―off‖ time is considerably less than the previous step, which indicates less time for electrons to diffuse. This results in a critical point when the electron density is so large that it can ionize atoms in the sheath region. The fourth and final step of the MPP pulse is a high power steady state discharge i.e., the current, power, and voltage remain at steady state. This step is between 900 – 1000 μs for this particular waveform, in which the high ionization has taken place.

Figure ‎2.36 The target voltage, current, and power waveforms during one modulated pulse used for the deposition of tantalum coatings (pulse width = 1000 μs, frequency = 65 Hz, Power- average =2 kW). The numbers indicate separate steps of the modulated pulse [93].

Meng et al. [94] investigated plasma dynamics in a MPP magnetron discharge using a time-resolved Langmuir probe. They measured the electron density, which directly correlates to

plasma density, and electron temperature, which directly correlates to the degree of plasma ionization. The MPP can set certain parameters constant. For example, keeping the repetition rate (duty cycle) constant implies the average target power will increase or decrease when either the voltage or current change.

Meng et al. observed the electron density/temperature behavior directly mimics the voltage/current behavior when keeping any parameter constant. This implies when the current rises, the density of the plasma increases and the degree of plasma ionization rises when the voltage rises. Therefore, the ―micropulses‖ in steps three and four in the MPP pulse in Figure 2‎ .36 are very important when trying to maintain a balance between the density of the plasma and the degree of ionization within the plasma. They also observed the effect of several modulated pulse parameters on the electron density and electron temperature. They found that the rise in current increases the electron density because higher electron densities produce even more ions, which continually increase the current. The voltage drops with the electron temperature when the current is increased because of one of two reasons. First, when the current increases, more target material is sputtered that immediately ―cools‖ the plasma by electron collisions; thereby decreasing the electron temperature. The second reason is because as the voltage drops, secondary electrons aren‘t accelerated into the plasma as forcefully; thereby decreasing the electron temperature.

When increasing the working pressure, the voltage decreases because higher working pressures require lower voltages necessary to maintain plasma equilibrium. This causes the sheath potential to be lower, which decreases the electron temperature. Increasing the pressure also decreases the mean free path, which decreases the degree of plasma ionization and the electron temperature. The current increases when the voltage decreases in order to maintain constant power. This causes the plasma density and the electron density to rise.

The balance between the density of the plasma and the degree of ionization within the plasma is very important. If the electron temperature is too high, the degree of ionization will also be high. Thus, the ions arriving at the substrate will have an enormous amount of energy.

Typically, ions with ion energies greater than 30 eV can produce film defects [95], [96] and can translate into unwanted residual stress [97].

Bohlmark et al. [86], [98] measured the ion energy distributions from a HPPMS discharge during various times of the pulse as shown in Figure 2‎ .37. The ion flux is very large and the majority of metal and gas ions (107 – 108 ions) during each period of the pulse exhibit ion energies of ~5 – 10 eV. However, there are a significant amount of ions that are greater than 30 eV and some surpassing 100 eV. Thus, HPPMS exhibits a high ion flux, but at the expense of very high ion energies that can damage films and generate residual stress.

Figure ‎2.37 Ion energy distributions from a HPPMS discharge at various time periods during the discharge for Ar+ ions and Ti+ ions [98].

In contrast, with MPP, the residual stresses are minimized because the arriving ionized species exhibit much lower mean and maximum ion energies when compared to HPPMS. Lin et al. [99] demonstrated this by measuring the ion energy distribution (IED) of Ar+, Ta+, and Ta2+

using electrostatic quadrupole plasma diagnostics (EQP) for the deposition of tantalum with MPP (Figure 2‎ .38). All ion species were shown to exhibit the mean ion energy of 3 eV whereas, the maximum tail energy for Ar+ and Ta+ was 12 eV, and Ta2+ was 8 eV.

Similar observations of a high degree of ionization of the target materials in MPP plasmas have also been confirmed by other plasma diagnostic work for sputtering Cr and CrN films [91], [100]. These values are much lower than those found in the HPPMS technique, confirming that MPP processes provide additional flexible control of the IED than HPPMS.

This advantage is especially important when creating low defect density films for electronic applications. The IED also revealed that a majority of the ions were either Ta+ or Ta2+, which is much higher than the few percent of metal ion species produced by dc magnetron sputtering (dcMS) [101].

Figure ‎2.38 Ion energy distributions of Ar+, Ta+, and Ta2+ ions during the MPP Ta discharge [99].

Sproul et al. [102] confirmed these findings when measuring the IED as a function of working pressure and found that the metal ion intensity peaks between 0.67-0.93 Pa. Therefore, the EQP results confirm that MPP can generate plasmas containing a large amount of ionized target species while maintaining low ion energy. This plasma characteristic can be usefully utilized to provide enhanced ion bombardment to densify the coating and at the same time minimize the defect densities and residual stress in the growing film.

2.5 Structure Zone Models and Film Microstructure Evolution

In thin film processes, atoms absorb onto the substrate, produce nucleation sites and grow to form final film structure. Without the plasma aid, thin film process is thermo-chemically controlled and depends on the substrate chemistry, film material and temperature of the system [103]. There are three steps involved in a thin film deposition process: first is the production of the atomic, ionic and neutral species and secondly their transport to the substrate through the medium and the condensation on the substrate to from a final thin film structure. The film nucleation and growth can be described with three growth types as illustrated in Figure 2‎ .39 [104]. a) Island growth 3-D (Volmer-Weber type): In this growth process, the binding strength between the adatoms is greater than that between adatom and substrate and hence the clusters grow three dimensionally to form islands. This type of growth is widely observed for metal films especially on oxide substrates. b) Layered 2-D (Frank-van der Merwe type): When the bond strength of film adatoms to each other is equal or less than to that of the atoms in the substrate, a 2-D layer by layer growth is promoted. This type of growth is predicted for strongly attractive substrates and found in single crystal epitaxial systems with low deposition rates and high substrate temperatures. c) Mixed growth (Stranski-Krastanov type): A mixed growth mode combines layer growth and island growth. In this case, after few monolayer growths, the subsequent layer growth becomes

unfavorable and islands are formed. This growth mode is commonly observed in many kinds of coating systems.

An external energy is usually needed to facilitate the formation of the thin film and to improve the film quality and structure. This modification can be done through excessive thermal energy or kinetic energy. The thermal energy comes from the heating the substrate and the kinetic energy originate from the ion bombardment, the momentum transfer of ion species and neutral adatom during the thin film deposition process.

Figure ‎2.39 Three modes of thin film nucleation and growth showing 2-D layer by layer growth mode proposed by Frank-van der Merwe, 3-D island growth by Volmer-Weber, and Stranski- Krastanov mixed growth.

In general, the plasma assisted depositions need moderate substrate temperatures to enhance adatom mobility and hence the thin film quality. However, this could be done at room temperature substrates with using high energy ion bombardment processes.

In past decades, different structure-zone diagrams have been proposed to correlate the microstructure of thin films with the deposition parameters. The melting temperature of the film

(Tm ) was assumed as a basic material parameter and the substrate temperature (Ts ) as the main process factor. Several zone models have been developed based on the TTsm/ ratio (the homologous temperature), where typical example is the Movchan-Demchishin model for vacuum vapor deposition [105]. Grovenor model for thermal evaporation [106], Thornton model for cathode and magnetron sputtering [107], and Messier model for ion beam deposition [108] have been developed thoroughly.

The cases of thermally evaporated films, the microstructures were classified to four zones according to their homologous temperatures as shown in Figure 2‎ .40. At the substrate temperatures below 0.15Tm , the film consists of porous columnar grains (zone 1). When the substrate temperature is between 0.15 and 0.3 , the film exhibits a transitional structure (zone T) between zone II having a dense columnar structure and zone III which is controlled by bulk diffusion due to the substrate temperatures higher than 0.5 .

Figure ‎2.40 Zone model for the grain structure of vapor deposited metal films modified at low temperatures to take into account the fine equiaxed grain growth [106].

For ion beam deposited coatings, Messier modified the Thornton model by replacing gas pressure in the sputtering zone with energy of ions reaching the substrate surface as depicted in Figure 2‎ .41. As the bombardment energy increases, the width of zone T increases at the expense of zone I.

In the most well-known zone model for sputtering films, Thornton revised the diagram by including gas pressure as an additional axis. His findings are based on 25 micron metal and alloy films sputtered at deposition rates 0f 100-200 nm/min. Three distinct zones and one transition zone was proposed in this model as illustrated in Figure 2‎ .42.

Figure ‎2.41 Structure zone model for thick films showing the effect of both bombardment and thermal induced mobility [108].

Zone 1 microstructures occur at low TTsm/ temperatures (≥ 0.2 – 0.3) that increase with rising pressure. The microstructure consists of cone-like crystals containing numerous voids and dislocations at grain boundaries. The columnar structure concludes in domes at the surface of the coating creating a rough surface that appears black resulting from trapped light in the pores. The consequential open columnar structure is caused by low adatom mobility caused by higher

deposition pressure and lower homologous temperatures. Higher pressures usually decrease the mean free path causing decreased energetic adatoms that result in shadowing.

The shadowing phenomena result from low angle impingement from depositing atoms that creates areas of higher coating flux resulting in formation of islands. Increasing the temperature can eliminate pores caused by shadowing due to the increased adatom movement. Zone 1 microstructures include poor mechanical properties due to low density, poor optical properties caused by the dome-like surface, and extreme anisotropic magnetization and electrical resistivity properties. Zone 1 structures rarely exhibit significant residual stress values caused by the porous network.

Figure ‎2.42 Microstructural zone diagram for sputter deposition processes proposed by Thornton [107].

Zone T of the Thornton Diagram is known as a transition region between zones 1 and 2.

The ratio of substrate temperature: thin film melting temperature TTsm/ range from 0.2 to 0.5 [66], noting that raising the pressure while maintaining a low will result in the formation of a Zone 1 structure. The microstructure consists of poorly defined fibrous grains with reduced columnar structure. Zone T exhibits much higher density than Zone 1 and is much smoother. The surface diffusion of adatoms is high enough to prevent some pore formation, but overhanging atoms can create self-shadowing effects that result in small void networks. Thus, the anticipated properties for Zone T microstructures are similar, yet not as discouraging as Zone 1 structures. However, the smooth surface enhances the optical properties. Residual stress becomes apparent caused by the increased density of the film.

The Zone 2 structure provides a significantly improved microstructure created by raising to 0.2 – 0.5. The increase in temperature enhances surface diffusion of adatoms, which were once subjected to shadowing. The porous columns of Zone T recrystallize and form grains that have fewer defects at higher densities. The overall structure consists of columns, which have tight grain boundaries and, in some cases, are broken into smaller sub-grains. The surface is rough due to the continuing columnar structure. The appearance of the film is matte or milky in nature. Since defects are at a minimum, the overall anticipated properties are significantly improved. The coatings exhibit high strength with low ductility. However, optical properties are poor because of the rough surface.

The Zone 3 transition begins at larger than of 0.5. The growth features are characterized by the addition of equiaxed grains depending on stress distribution and structure. The grains are extremely dense and often form twinning boundaries. All columnar grains are broken into sub-grains. Transport features include large diffusion among the bulk of the material. Hence, recrystallization and grain growth occur on a bulk scale. However, recrystallization is time dependent due to nucleation. The surface of Zone 3 is smoother than any of the previous zones due to increased bulk diffusion. Nevertheless, the grain boundaries have a tendency to develop grooves. Most residual stresses are removed by recrystallization. Zone 3 anticipated properties are similar to Zone 2, with the coating exhibiting high strength, but

low ductility. Yet, coatings deposited at higher Zone 3 temperatures demonstrate properties analogous to bulk annealed materials. Nevertheless, Zone 3 is often not observed in most materials because of the difficulty of using substrate temperatures > 0.5 Tm.

Obtaining a zone 2 or 3 microstructure without increasing deposition temperature has become a goal for most coating processes. Kelly et al. [72] developed zone T and 2 microstructures of several metal films below temperatures recommended by Thornton [107] by increasing the ion energy and ion flux ratio using CFUBMS with biased voltages as shown in Figure 2‎ .43. Kelly found that the high adatom movement caused by increased deposition temperature is replicated by higher energy ion bombardment.

Figure ‎2.43 Structure zone model relating to the CFUBMS system. Solid circles mark the position of coatings with zone 2 structure and shaded circles mark the position of coatings with zone 3 structures [72].

CHAPTER 3. EXPERIMENTAL PROCEDURES

In this section, the experimental work to deposit AlN thin films with DC, pulsed DC and MPP is summarized. The section is separated by the deposition parameters in which the films were produced and the characterization techniques implemented to determine various properties.

3.1 AlN Thin Film Deposition Process

Aluminum nitride thin films were deposited on 304 stainless steel, boro-silicate glass and (100) single crystal Si wafers by DC, pulsed DC and modulated pulsed power magnetron sputtering systems using a pure metal Al target (99.95%) in the high purity (99.999%) argon and nitrogen gas mixture.

In our chamber, four unbalanced magnetrons are configured in 90-degree arrangements and are closed completely. Advanced energy pinnacle plus power source has been used to deposit the coatings in mid-frequency and DC regions. In the positive portion of the voltage waveform, target is charged with about 10 percent of its nominal sputtering voltage attract electrons to the target surface and discharge any regions that were charged during the negative voltage application on the target and avoid target poisoning.

An illustration of the target configuration inside the four cathode closed field vacuum chamber is shown in Figure 3‎ .1. Only one of four cathodes was powered during AlN deposition and also one cathode was on while seed-layer (adhesion layer) is deposited, whereas, the other three unpowered cathodes were utilized to create a magnetically closed-field arrangement. The powered cathode was installed with a metal 99.95 % pure Al target and was positioned 240 mm from the other cathode. The target size was 228 mm x 127 mm x 6 mm (9 in. x 5 in. x 0.25 in.).

The rotating substrate holder was equidistantly placed in the middle of chamber. The substrate holder was machined such that a substrate-to-target distance of 60, 90, 140, and 165 mm could be used. The substrate holder was attached from the top of the chamber and could be

electrically connected to either the middle frequency pulsed dc power supply for plasma sputter etching or to the bias power supply to provide a negative substrate bias.

Ti, Ti/TiN, Mo, Pt

Figure ‎3.1 Illustration of four cathode arrangement inside vacuum chamber including the position of the substrate holder and the closed magnetic field created by the four facing unbalanced magnetrons.

AlN films were deposited using DC sputtering in a Horizontal chamber as shown in Figure 3‎ .2 (a). Further optimizations of AlN films were done in closed field system depicted in Figure 3‎ .2 (b). Preliminary investigations of AlN sputtered films were obtained using modulated pulsed power system as photographed in Figure 3‎ .2 (c).

Polished AISI 304 stainless steel coupons (2.5 x 3.75 cm), (100) doped silicon wafers (~1 x 2 cm) and borosilicate glass (2.5 x 7.5 cm) were used as substrates. The substrates ultrasonically cleaned using acetone and alcohol for 15 minutes each. Base pressure of 1×10-6 Torr (1.3 ×10-4 Pa) or less was maintained prior to film deposition. The substrates were also etched in argon environment by applying biasing voltages of about 300 to 400 volts in order to

clean the substrates from any contamination. Different thin layers (about 100-200 nm) of Cr, Pt, Mo, Ti and Ti/TiN were coated as the bottom seed layer/adhesion electrode.

(a)

(b)

(c)

Figure ‎3.2 Photographs of (a) Horizontal vacuum chamber (DC), (b) P-CFUBMS system (pulsed-DC), and (c) MPP power supply and two cathode systems.

Several different regions of the stainless steel substrates were lightly marked with a marker to create a region with poor film adhesion in order to determine film thickness post deposition. The prepared samples were then mounted on a substrate holder and loaded in the vacuum chamber.

The sample holder was rotated to a position perpendicular to the Al target prior to plasma sputter etching. The samples were sputtered etched prior to deposition to ensure clean sample surfaces. Sputter etching was achieved by a Pinnacle Plus Advanced Energy Inc. middle frequency pulsed dc power supply using Ar+ plasma ion bombardment at a pulsed bias voltage of 300-400 V, 100 kHz and a 90 % duty cycle for 30 minutes. The resulting power and current were typically 50 W and 0.1 A, respectively. High purity argon gas (99.999 %) was used for sputter etching and coating deposition. The working pressure for the sputter etching was 1.33 Pa (10 mTorr).

After samples were sputter-etched, the substrate holder was rotated such that the samples were facing away from the target which seed layer deposition would occur. Contaminations on the target were removed by applying low power (300 W) with either the pulsed DC or the DC power supply for at least five minutes. With the conclusion of this step, the substrate holder was rotated directly in front of the Cr, Ti, Ti/TiN, Pt or Mo targets to deposit an appropriate buffer layer that would further enhance c-axis texture of AlN films.

Several depositions parameters such as N2 to Ar gas ratio, working pressure, negative substrate bias, substrate heating, substrate-to-target distance, deposition time (thickness) and target pulsing frequency alteration were investigated for DC, pulsed DC and MPP systems. In reactive sputtering, the atmosphere is consists of the reactive gases (e.g. N2, O2 and etc.) and the inert gas mainly argon. Ar is used to initiate the plasma and generate ions in glow discharge accelerated in cathode sheath to sputter the target. The reactive gas is introduced through the chamber to deposit compounds such as nitrides and oxides on the substrates. Reactive gas ratio, working pressure and flow rate of the reactive gas have profound impact on the structure, composition and properties of compound films. In this study, the reactive gas (N2) percentage was varied from 10 to 100 at a given working pressure. The working pressure inside the

chambers was maintained in a range of (3 mtorr to 6.5 mtorr) with flow controllers. Substrate bias and temperatures were applied to optimize the coating texture. A brief summary of all the deposition parameters studied in this work are given in Table 3‎ .1.

Table ‎3.1 Deposition summary of AlN deposited by reactive DC and pulsed DC magnetron sputtering.

Base Operating Sputter gas frequency Power Time Pressure Pressure

≤ 10-6 Torr 3.0-6.5 Argon& 20-300 kHz 2-3 W/cm2 (Pulsed 2-3 hr. mTorr DC) Nitrogen 5-10 W/cm2 (DC) 1-1.5 hr.

3.2 Film Microstructure and Composition Characterization

The following section describes all of the characterization techniques used in this thesis to study the relation between the deposition process, film microstructure and properties of the AlN films.

3.2.1 X-Ray Diffraction (XRD)

X-ray diffraction (XRD) is a non-destructive technique for crystalline material characterization. The XRD peaks are obtained due to constructive interference of a monochromatic beam scattered from lattice planes at specific angles. Therefore, the x-ray

diffraction pattern is the representative of crystal structure of given sample. XRD can also provide information of preferred orientation (texture), average grain size, residual stress, and crystal defects. Three types of XRD configurations (conventional, glancing incident and rocking) were recorded to study the crystal structure, stress and c-axis texture of AlN films, respectively.

AlN films structure were characterized by Cu-Kα radiation (1.5418 angstrom) on a Philips X-ray diffractometer model PW1729 operated at 45kV and 40 mA. C-axis or (002) texture of the AlN thin films were analyzed by the rocking curve XRD method using Simens X- ray diffractometer (model KRISTALLOFLEX-810).

3.2.1.1 Conventional /2 Configuration

In conventional XRD scans, the diffraction angle was recorded from 30° to 60°, with the scanning step size of 0.05 and 4 seconds per step. In this arrangement, the incident and detection angles relative to substrate plane were equal during all the measurements. In this arrangement, the only information related to the planes parallel to the substrate surface is obtained. Data taken from conventional XRD scans can be used to determine the size of the crystallites in the sputtered AlN thin films.

Considering a single crystal film that is free of defects, grain boundary, or any other scattering centers, reflections from the planes would approach delta function. In a crystal that contains scattering centers, the reflection that occurs would be broadened over some range of angles. In polycrystalline materials such as AlN films, the presence of crystallites and grain boundaries reduce the periodicity of the crystal lattice thus broadening the reflection peak. The Scherrer method uses an analysis of x-ray data from a /2 scan to calculate the average crystallite size in a polycrystalline material [109]. For a polycrystalline thin film, the crystallite size is given as:

K D  Equation ‎3.1 2 cos( )

 is the x-ray wavelength (0.15418 nm) for Cu K ; 2 is the location of the peak 1 corresponding to the (002) orientation. 2 is the full with at half maximum (FWHM) of the peak and K is the appropriate shape factor which is assumed to be 0.9 [110], [111]. From experimental results, an empirical relation has been developed for the crystalline size as a function of crystalline texture of the films. The crystalline size is inversely proportional to the FWHM of the rocking curve [112].

3.2.1.2 Rocking Curve ( -scan)

In this XRD configuration, the detector is fixed at diffraction angle of interest and the sample is rocked through an angular range, bringing the plane of interest in and out of the Bragg condition. The width of the measured peak measured at its half-maximum (FWHM) contains information of the amount by which the measured plane is off the surface normal and is often referred to as degree of orientation or texturing of the specimen. A schematic of conventional and rocking XRD configuration is shown in Figure 3‎ .3.

3.2.1.3 Glancing-Incident X-Ray Diffraction (GIXRD)

Glancing incident angle XRD is used to study the diffraction pattern of AlN films deposited on various substrates with suppressed peaks interfering from the substrates. In this configuration, the incident angle  is fixed at grazing angles of (1 10 ) and the diffraction profile is recorded by detector scan only.

Since the penetration depth of the x-ray in this configuration is lower, the substrate effect is reduced and only the surface layers of the film are measured. The GIXRD method can also be used to measure the residual stress inside the film. Diffraction patterns obtained at conventional arrangement contain information averaged over a larger fraction of the sample. Whereas, analysis of diffraction patterns acquired at gradually increasing incident angles would allow quantitative residual stress in the films.

Figure ‎3.3 X-ray diffraction in (a) conventional configuration /2 with the incident angle being the same as Bragg angle, (b) and (c): change of the incident angle for a fixed Bragg angle to obtain rocking curve.

3.2.2 X-Ray Photoelectron Spectroscopy (XPS)

X-ray Photoelectron Spectroscopy (XPS) (AXIS HS XPS, Kratos Analytical Ltd.) using an Al K x-ray source (13 kV and 15 mA) was used to investigate the chemical composition of the surface of the coatings (instrument shown in Figure 3‎ .4). The incident x-rays cause the ejection of core-level electrons from sample atoms. The energy of a photoemitted core electron is a function of its binding energy and is characteristic of the element. When the core electron is ejected by the incident x-ray, an outer electron fills the core hole.

The energy of this transition is balanced by the emission of an Auger electron or a characteristic x-ray. Analysis of Auger electrons can be used in XPS, in addition to emitted photoelectrons. The photoelectrons and Auger electrons emitted from the sample are detected by an electron energy analyzer, and their energy is determined as a function of their velocity entering the detector. By counting the number of photoelectrons and Auger electrons as a function of their energy, a spectrum representing the surface composition is obtained. The energy corresponding to each peak is characteristic of an element present in the sampled volume. The area under a peak in the spectrum is a measure of the relative amount of the element represented by that peak.

The samples were ultrasonicated in acetone and rinsed with methanol prior to mounting in the vacuum chamber. A base pressure of at least 1 X 10-10 Torr was reached before sputter etching and scanning each sample. The samples were sputtered etched with high purity argon for thirty minutes in an effort to remove surface contamination. Once etched, the XPS completed a survey scan across the entire energy spectrum to determine the presence of elements. Energy peaks in the survey scan identify the elemental composition of the surface. After the survey scan, high resolution scans were completed for Al, N, O, and C energy spectra to obtain the chemical composition of the films. XPS is not sensitive to very light elements such as hydrogen or helium but can detect all other elements.

Figure ‎3.4 A photograph of the X-Ray Photoelectron Spectrometer XPS (AXIS HS XPS, Kratos Analytical Ltd.).

3.2.3 Field Emission Scanning Electron Microscopy (FESEM)

Film microstructures were examined using a JOEL JSM-7000F field-emission scanning electron microscope (FESEM) operating at an accelerating voltage of 2-5 kV as shown in Figure 3‎ .5. The top surface morphology and cross-section were observed for deposited AlN

films on silicon. Silicon wafers were carefully fractured after scratching the back of the silicon wafer with a diamond scribe to obtain fresh cross-section.

Figure ‎3.5 A photograph showing the JOEL JSM-7000 Field Emission Scanning Electron Microscope (FESEM).

3.2.4 Scanning Transmission Electron Microscopy (STEM)

Longitudinal cross-section TEM sample were prepared from the AlN film deposited on a silicon wafer using mechanical thinning, polishing and ion milling processes. A dimple was produced at the cross section of the film and was further thinned in ion beam miller to create a crater. A Philips/FEI CM 200 transmission electron microscope operated at 200 kV was used to study the film cross sectional microstructure, columnar size and defects.

3.3 Film Properties Characterization

In order to establish the processing, microstructure and property relation, several characterization techniques have been applied to analyze the piezoelectric coefficient, residual stress, mechanical, thermal stability and high temperature oxidation behavior of AlN films.

3.3.1 Wafer Flexure Technique (Direct Piezoelectric Measurement)

Figure ‎3.6 Schematic of a piezoelectric measurement setup using wafer flexure mode.

In this technique, the film was first attached firmly with a flexible silicon wafer (about 0.30-0.40 mm thick), which was used as a diaphragm between two chambers with good pressure sealing. The pressure in one of these chambers was varied periodically by injecting air in and out of the chamber using a vibrating diaphragm of an audio speaker placed in front of one opening in the chamber. This caused a periodic pressure difference across the diaphragm. This pressure difference in turn caused a bending of the diaphragm, and a resultant in-plane periodic strain.

The time varying strain caused transient voltage across the film (and hence a transient current flowing from the positively charged electrode to the negatively charged electrode), which was subsequently measured using a lock-in-amplifier, and the total charge collected on the electrodes (during the transient) was calculated by integrating the signal. The transverse piezoelectric coefficient e31 given as:

charge e  Equation ‎3.2 31 2  area strain

Knowing the amount of pressure difference across the silicon wafer and also the elastic constants of the wafer, one could calculate the strain. In order to minimize the error due to sample shape (which was non-circular) the strain was calculated by taking the average of its two values along two perpendicular in-plane directions.

3.3.2 Michelson Interferometry (Remote Inverse Piezoelectric Measurement)

The piezoelectric properties of AlN films needed to be characterized in-house therefore a homemade interferometer was designed and built in collaboration with Prof. John Scales in CSM physics department. One proposed idea is to use a non-contact method such as a laser beam to shine at the system of the AlN thin film sandwiched between two electrodes, and then with the application of cyclic voltage, the displacement variation within the piezoelectric AlN thin film can be measured remotely.

A Michelson interferometer was proposed to be used due to its high precision in displacements measurements down to 10-3 angstrom. In this setup, as shown in Figure 3‎ .7; a light source such as a laser source emits an electromagnetic wave, which is then divided at a half- transparent mirror into two beams with equal intensities. These beams are reflected back at two plane mirrors, one mirror is fixed and the other mirror is our sample, which will change dimensions upon voltage application, and return to the half-transparent mirror, where they are

combined before they emerge at the screen or detector. The different paths may be of different lengths or be composed of different materials to create alternating interference fringes on a back detector.

Figure ‎3.7 Homemade Michelson interferometer for remote piezoelectric measurements and its schematic ray diagram on the right.

By applying a voltage difference across the top and bottom electrodes in the AlN thin film system, one can measure the displacement change due to piezoelectric effect by using the mentioned Interferometer setup.

A panametric transducer was used as a reference sample to go between commercial laser vibrometer and our homemade laser interferometer. The vibrometer measures the surface modulation velocity produced by sinusoidal 10 Vp-p applications. The transducer was driven at 60 kHz which produced a measurable signal to calculate the surface motion speed. The same transducer was tested with our interferometer driven at the same peak to peak voltage and frequency. Therefore, the displacement variation was calculated and in-direct piezoelectric coefficient d33 was obtained in (pm/V).

3.3.3 Mechanical Properties (Nanoindentation)

Nanoindentation is defined as a tool to measure mechanical properties of materials at the nanoscale. This method was first developed in the early 1980s evolving from traditional Vicker hardness testing. Nanoindenter tips with various shapes are used in this technique to analyze the resistance of materials to an external force. A schematic representation of continuous load- displacement data is shown in Figure 3‎ .8 which consists of three different regions called loading, holding and unloading regions. Some important quantities that are indicated in this figure are:

Pmax (peak load), hmax (maximum displacement) and S dP/ dh (the slope of unloading curve at maximum indentation depth).

Hardness and elastic modulus measurements were completed by a nanoindenter (Nanoindenter XPTM, MTS Systems Inc.) equipped with a Berkovich diamond indenter. Figure 3‎ .8 displays a typical indentation cycle of loading and unloading the indenter from Pharr et al. [113]. The slope of the unloading curve is the reduced elastic modulus Er (all modulus contributions from the indenter and sample). This slope determines the elastic modulus of the film E with Equation 3.4 where Ei is the elastic modulus of the indenter; νi is the Poisson‘s ratio of the indenter, ν is the Poisson‘s ratio of the film which is about 0.25[114].

It is noticeable that S has the dimensions of force per unit length, which is known as elastic contact stiffness. The elastic modulus of materials can be derived using the following equations:

() S E  Equation ‎3.3 r 2 A

where  is a constant that depends only on the geometry of the indenter and A is the projected contact area. Er is called the reduced modulus, a parameter that considers the effect of a non- rigid indenter on the load-displacement behavior. The elastic modulus, E of the test material is calculated using the following expression:

Figure ‎3.8 Schematic representation of the nanoindentation load-displacement curve.

11 2 1 2 i Equation 3‎ .4 EEEri

where  is the Poisson‘s ratio of test material and Ei and  i are the elastic modulus and Poisson‘s ratio of the indenter, respectively. It may seem inappropriate that we have to know the material‘s Poisson‘s ratio in order to calculate its modulus. But, even a rough number, say 0.25 0.1, produces only about a 5% error in the calculation of elastic properties of most materials [115]. For different types of nanoindenter tip, a specific  constant is used for calculations. For indenters with square cross-sections such as the Vickers pyramid,  1.012 ; for triangular cross-sections such as the Berkovich and the cubic corner indenters,  1.034 . An additional mechanical property that is usually obtained from nanoindentation is the hardness, H , as defined below:

P H  Equation 3‎ .5 A where P is the applied load and A is the projected contact area of indentation at load P as a function of contact depth. Therefore, the elastic contact stiffness ( S ) and projected contact area ( A ) must be obtained in order to be able to derive the elastic modulus and hardness. There are two distinct methods that measure the stiffness and projected contact area, i.e., the continuous stiffness measurement (CSM) method, and the unloading stiffness measurement (USM) method.

In the CSM method, a small oscillating force is applied either to the sample or indenter during the indentation period and the contact stiffness value is calculated from the displacement response against the depth of indentation. Once the stiffness of contact, S is defined, the elastic modulus and hardness can be obtained using equations mentioned previously. In the USM method, there are several ways to calculate the contact stiffness. The method of Oliver and Pharr [116] is the most widely used. According to the method, data analysis procedure begins by fitting the load-displacement data acquired during unloading to the power-law relation:

m P B() h hf Equation 3‎ .6 where P is the applied load to the test surface and, h is the resulting penetration, B and m are empirically determined fitting parameters, and hf is the final displacement after complete unloading. The contact stiffness, S , is then obtained using:

dP S  Bm() h  h m1 Equation 3‎ .7 dh max f

It is worth noting that, in nanoindentation, the projected contact area is not obtained by optical

imaging. Rather it is computed as a function of the contact depth, hc , applying an empirical relationship given by:

2 A f() hc  C0 h c  C 1 h c  C 2 h c Equation 3‎ .8

the constants CC01, andC2 are determined prior to the experiment by indenting a sample of known properties such as fused silica. The contact depth, hc is different from the total penetration depth, h , and is estimated using:

P hh Equation 3‎ .9 c S where  is a constant that depends only on the geometry of the indenter. For cones,   0.72 and for spheres,  is 0.75. There is empirical justification for using   0.75 for Berkovich and Vickers tips as well.

3.3.4 Residual Stress Measurements

Residual stress is the resulting internal stress left in a material when all external forces are removed. They are present in all materials that are altered by mechanical, chemical, or thermal processes. For example, Figure 3‎ .9 displays two growing films on the top. The top-left film initially shrinks compared to the substrate and the top-right film expands. Ohring [97] states that surface tension and misfit accompanying epitaxial growth are some of the reasons for this phenomenon.

Figure ‎3.9 Initial steps of film deposition when the film is too narrow for the substrate (top-left) and too wide (top-right) and resulting end moments and how the film and substrate react. The bottom-left film is under residual tensile stress and the bottom- right film is under residual compressive stress.

The film and substrate either expand or contract in order to be the same ―equilibrium‖ length. Figure 3‎ .9 illustrates the resulting stresses as the left film is now in tension and the substrate is in compression and the opposite is true for the film on the right. The tensile forces are balanced by the compressive forces, but the entire sample is not in mechanical equilibrium because of the uncompensated end moments. As a result, the left film will buckle upwards and the right film will buckle downward. The substrate will also respond if it is thin and elastic. If the substrate is too thick and resists the moment, the tensile film will fracture and the compressive film will wrinkle and lose adhesion to the substrate. The built-up strain in the coating is a function of the thickness of the coating.

There are many factors that affect residual stress, such as the lattice mismatch between the film and substrate and thermal expansion coefficients. Both of these factors are difficult to

change without changing materials. However, one can modify residual stresses through changes in ion bombardment during the deposition process as ion energy has a square root dependence on stress. Some compressive stresses can be very beneficial to mechanical and corrosion resistant properties to counteract harmful tensile stresses, but high levels of stress can cause catastrophic failure [66]. These failures are amplified during thermal treatments.

There are several methods to determine residual stress, one being mechanical deflection with the Stoney formula [117–120] and another by determining peak shifts with XRD [121–124]. Mechanical deflection is typically easier than XRD and will determine stresses for non- crystalline films, but the accuracy is much lower.

The theory behind the XRD measurement is that residual stress in a material causes a uniform biaxial micro-strain that causes a broadening of the original diffraction peak. By measuring the peak shift, the strain can be calculated which, directly correlates to the stress by Hooke‘s Law. The common method for thin film residual stress measurement is sin2 ψ. This method uses Glancing Incidence X-ray Diffraction (GIXRD) where the x-rays are directed at a very low incident angle (γ = 1–10°). This method does not work for textured thin films and has been modified with two assumptions made in this case: (1) the stress state in the film is equi- biaxial and (2) the crystal is very small and each crystal has a random orientation around the c- axis with 2 rotational freedom. Under these assumptions, the relationship between the lattice strain 33 and the stress  is [125], [126]:

c c c2 c 33{(s 11  s 12  2 s 13 )sin   2 s 13 }  Equation 3‎ .10

c where, sij is the compliance of a single crystal and  is the angle between normal axis of the diffracting planes and the surface normal axis, which is identical to the c-axis. The lattice strains for each diffraction plane should lie on a straight line against sin2 . The residual stress value can be estimated by knowing the elastic compliance of the AlN. The values for an AlN crystal

c5 c  5 c  5  1 are: s110.35  10 , s 12   0.10  10 , s 13   0.08  10 MPa [125]. AlN powder diffraction data is used as a standard and the strain is then calculated by the following equation:

ddfp 33  Equation 3‎ .11 d p

where, d f and d p are measured interplanar spacing of the film and the powder, respectively. There is also another technique specifically for textured thin films to calculate residual stress inside the films. Assuming that residual stress is the only mechanism governing deviations from theoretical lattice constants, excluding the effect of impurities and interstitials, the strain 33 perpendicular to the film plane as observed by the shift of the (002) peak from its theoretical value, would be the result of the in-plane biaxial stress where r  12  . The latter influences c-axis strain through stiffness tensor cij as follows [127–129]:

()c  c   c  r 11 12 11 13 33 Equation ‎3.12 320cc 13  11  33  33 

The residual stress can thus be calculated from33()/d 002 dth d th :

33c 33() c 11 c 12  r {2c13 [ ]} Equation ‎3.13 2 c13

where, c11,,, c 12 c 13 and c33 are the stiffness constants, dth is the theoretical stress free c-axis single crystal AlN, and d002 is the fitted peak position.

3.3.5 Oxidation Resistance and High Temperature Stability

The AlN film in the die casting applications needs to be stable at high temperatures and being oxidation resistant. The dynamic thermal stability studies were conducted in a Netzsch

STA-409C differential scanning calorimeter (DSC). About 30 mg of film on silicon substrate was placed in an alumina crucible and heated to 1000 C using 20 K/min of heating rate in argon environment. An empty pure alumina crucible served as the reference material. XRD measurements were made before and after the DSC test to study their structural and chemical changes due to this thermal treatment. Cross sectional and plan view FESEM were also obtained to see any microstructural alterations.

Several AlN samples were sent to Prof. Dahan [130] to perform in-situ high temperature XRD measurements. This method examines the high temperature structural evolution of AlN thin films in air and inert gas environment. This was done in their customized module attached to Philips XRD system. Anton-Paar HTK-1200N hot stage was attached to their XRD system to study high temperature behavior of different thin films and powder samples, as shown in Figure 3‎ .10.

Figure ‎3.10 Hot stage module attached to Philips XRD (HTXRD).

CHAPTER 4. RESULTS AND DISCUSSION

AlN thin films deposited with DC, pulsed DC and MPP systems are summarized in this chapter. The interplay of structure, processing and properties of AlN coatings are discussed in order to maximize c-axis preferred orientation and hence optimize the devise final piezoelectric response.

4.1 DC Reactive Sputtering of AlN Films

AlN thin films were deposited on borosilicate glass, silicon and stainless steel (SS) substrates using pure Al target reactively sputtered in nitrogen and argon environment with the application of negative DC voltage on the cathode. Deposition process parameters were optimized to obtain highly c-axis textured films to achieve higher piezoelectric response.

Ti under-layer (seed layer) was deposited on the substrates to act as the adhesion layer for the AlN films and also a seed layer for the (002) texture growth of the functional AlN films. Ti has the hexagonal closed packed structure similar to the AlN films and is a well-known adhesion layer in sputtering processes. Ti seed layers were found to grow along its c-axis direction for small thicknesses which would become mixed orientation as the deposition time was increased.

4.1.1 Reactive Gas Ratio Effect

The crystalline structures of the DC reactively sputtered AlN films deposited at different nitrogen to argon ratios are shown in Figure 4‎ .1. It can be seen that the crystal phase starts forming from nitrogen content as low as 10% in the plasma, but there was a significant proportion of un-reacted aluminum in the deposited film. Pure AlN was obtained once the nitrogen content in the plasma exceeded 20%.

Irrespective of substrate material used, the obtained films were found to exhibit two Bragg reflections in the XRD, one at the (002) reflection angle (36.05 degrees) and the other one

(at 38.2-38.4 degrees) between the AlN (101) and Ti (002) peaks. Also, the second peak was evolving with deposition parameters, as seen from Figure 4‎ .1. This led to the conclusion that the second peak could be the convolution of AlN (101) and Ti (002) peaks [131]. Therefore, the second peak could be the (101) AlN, and hence, the sample had a mixed orientation. Although the behavior of the second peak intensity with respect to the main peak was quite irregular, there was a definite increase in the second peak intensity observed at very high nitrogen content in the plasma. It is known that the (002) have the lowest energy HCP plane in the AlN system.

Figure ‎4.1 XRD patterns of AlN thin films deposited using DC reactive sputtering on Ti/SS as a function of various Ar/N2 ratios.

This result indicated that the system is far from the lowest energy state and therefore the crystallization parameter could be changed with deposition parameters since the deposition was mainly a kinematically limited process instead of an energy limited one. To further characterize these films deposited with different reactive gas percentages, the crystallite sizes of the films were calculated using the Scherrer‘s equation as following [132]:

K D  Equation ‎4.1 Bcos

where D is the crystallite size, K is a constant number around 0.9,  is the x-ray wavelength, and B is the full width at half maximum of the (002) peak. AlN thin films deposited above 20 percent nitrogen have about 30 nm c-axis oriented grain size. The film deposited at 10 percent of nitrogen reactive gas percentage has very broad peak in the XRD pattern indicating the formation of either XRD-amorphous or fine nano-crystallites of about 5 nm in size as shown in Figure 4‎ .2.

Figure ‎4.2 Grain size measurements of AlN/Ti /SS thin films deposited with various N2 to Ar ratios.

4.1.1.1 Annealing Treatment Effect on Texture

When the films were exposed to post-annealing treatment of 600 C for two hours in nitrogen environment, the mixed orientation films changed to the pure c-axis texture as shown in Figure 4‎ .3. It also helped in improving the electrical properties of the films. The current-voltage (I-V) curve of the annealed film is shown in Figure 4‎ .4 which illustrates a current density of less than 10-6 A/cm2 at an applied voltage of 1 V. This indicated a possibility of off-stoichiometry in the as deposited films, which was compensated by the external nitrogen flow during annealing. It was also a possibility that the film was smoother after annealing the surface, which reduced the electric field inhomogeneity, and therefore less chances of breakdown occurred.

Figure ‎4.3 XRD patterns of as-deposited and annealed AlN film at 600 C for 2 hours.

1E-3

) 2 1E-4

1E-5

1E-6

1E-7 Current density(A/cm

0.1 1 Voltage (V)

Figure ‎4.4 The I-V characteristics of the annealed AlN thin films at 600 C for two hours in pure nitrogen environment.

4.1.2 Insulation Characteristics of AlN Films

The insulation characteristics of the AlN films were measured by a two-probe I-V measurement. The film being insulating, would not introduce much error due to contact resistance in two-probe technique. The films were already deposited on metal like TiN substrates; one part of the substrate was physically masked during the deposition of the AlN film, to keep that part of substrate being exposed to air acting as the bottom electrode.

On top of the film, smaller dot electrodes were deposited using another shadow mask. Each dot had a diameter of about 0.5 mm. Pt was sputtered mainly for top electrode even though the best electrode would be Al, due to its good electrical contact formation with AlN film, but to have sharp interface, platinum was chosen. It is seen that the film deposited with thinner thickness exhibits very poor insulation properties, whereas a film with the deposition time just double of this can reveal a good I-V curve with fairly low current at low voltages, as shown in Figure 4‎ .5.

Figure ‎4.5 Two-probe I-V characteristics of AlN thin film with two different thicknesses, one deposited with 350 nm and the other coating deposited with the same deposition condition with thickness of about 900 nm.

Electrical properties of the thicker film (900 nm) were still inferior to the reported results. There could be several reasons for this poor insulation of the as deposited films. If the film is not at its lowest energy orientation, the surface would not be smooth any longer; instead it would have a surface reconstruction provided the surface atoms had enough mobility just before the deposition ended. This corrugation of the surface will produce area of concentrated field that would exceed the breakdown strength in some places. Then an avalanche breakdown could be initiated all through the sample.

Also the scattering between sputtered atoms/ions might alter the angle of incident of arriving particles hitting the substrate and produce significant shadowing of the non-normal beam resulting in a porous film as was shown previously in Figure 2‎ .2.

4.1.2.1 Film Thickness Effect on Piezoelectric Coefficient

For the piezoelectric characterization of AlN thin films, research collaboration was established with the W. M. Keck Smart Materials Integration Laboratory headed by Professor S. T. McKinstry in Penn state university (PSU). The technique adopted was called the wafer flexure technique. In this technique, the film was first attached firmly with a flexible silicon wafer (about 0.30-0.40 mm thick), which was used as a diaphragm between two chambers with good pressure sealing.

The pressure in one of these chambers was varied periodically by injecting air in and out of the chamber using a vibrating diaphragm of an audio speaker placed in front of one opening in the chamber. This caused a periodic pressure difference across the diaphragm. This pressure difference in turn caused a bending of the diaphragm, and a resultant in-plane periodic strain. The time varying strain caused transient voltage across the film (and hence a transient current flowing from the positively charged electrode to the negatively charged electrode), which was subsequently measured using a lock-in-amplifier, and the total charge collected on the electrodes (during the transient) was calculated by integrating the signal. The transverse piezoelectric coefficient e31 given as:

charge e  Equation ‎4.2 31 2  area strain

Knowing the amount of pressure difference across the silicon wafer and also the elastic constants of the wafer, the strain can be calculated. To minimize the error due to sample shape (which was non-circular) the strain was calculated by taking the average of its two values along two perpendicular in-plane directions.

In order to study the effect of film thickness on the piezoelectric response of the above films, the transverse piezoelectric coefficients of the films are measured and reported in Table 4‎ .1.

Table ‎4.1 Piezoelectric coefficient measurements for AlN films deposited with different thicknesses.

Architecture Thickness Processing parameter Piezoelectric coeff.

2 (nm) (DC-sputtered AlN) [e31 (C/m )]

AlN/Ti/Si 700 250W, 3mT, Pure -0.74 Nitrogen, 3h

AlN/Ti/Si 400 250W, 3mT, Pure -0.14 Nitrogen, 3h

As explained before, the AlN films deposited at the same deposition parameters having different thicknesses produce significantly altered performance. The AlN film deposited at 400 nm thickness has a much lower piezoelectric coefficient compared to the same film just deposited with twice its thickness. This also confirms the previous difference that was seen with I-V characteristics of these films.

4.1.3 Buffer Layer (Seed Layer) Effect on c-Axis Orientation

Thickness increase above the critical range produced an encouraging effect on the insulation characteristics and piezoelectric coefficient of the AlN films. However, this effect is dependent on the bottom layer used with the AlN films deposited by DC magnetron sputtering. Three bottom layers (seed layers) were studied based on their crystal symmetry and lattice matching with AlN thin films. The c-axis oriented AlN thin films can be deposited on silicon (111) and (100) wafers but this c-axis texture can be further improved by the incorporation of seed layer having similar hexagonal atomic arrangement (3-fold symmetry) and minimal lattice mismatch.

Ti was deposited as a buffer layer on top of several substrates such as Si (100), stainless steel and amorphous glass. Ti seed layer has a similar HCP crystal structure as AlN and its lattice mismatch for (002) a plane is about 5 percent, as illustrated in Figure 4‎ .6.

Furthermore, TiN/Ti seed layer system was chosen since the lattice mismatch between TiN (111) and AlN (002) plane is only about 3.6%. TiN is also a well-known diffusion barrier material [133–135], which would minimize any interaction or inter-diffusion that might happen between AlN and Ti layers. In addition, the crystal symmetry between TiN (111) and AlN (002) planes both exhibits hexagonal atomic arrangement as depicted in Figure 4‎ .6.

Figure ‎4.6 Schematic crystal structure of Ti, Pt and TiN seed layers showing their atomic symmetry within the closed pack planes and their lattice mismatch with AlN (002) basal planes.

AlN films that were deposited using the Ti and TiN/Ti as seed layers were tested for their piezoelectric response. The piezoelectric coefficient e33 was improved when TiN was deposited between the AlN and Ti layers to act as a diffusion barrier preventing the Ti diffusion into AlN layer, as reported in Table 4‎ .2.

Thin layer of Pt was deposited to act as a seed layer due to its crystal symmetry with AlN (002) and well known for its sharp interface with AlN films [136–139]. The best bottom layer (seed layer) was found to be a TiN/Ti bi-layer system having the lowest lattice mismatch with (002) AlN films. The effect of crystalline with different seed layer systems is shown in Figure 4‎ .7 which shows the improvement of c-axis orientation changing the under-layers from Ti to Pt/Ti and ultimately to TiN/Ti bi-layer system.

Figure ‎4.7 XRD of AlN films deposited with different seed layers with relation of their crystal structure and symmetry.

Piezoelectric coefficient, e33 , of these AlN films deposited on various seed layers were characterized with the previously mentioned method. As was evident from the XRD data, the AlN film deposited using TiN as an inter-layer between AlN and Ti shown the highest response

among the other candidates due to its impeding effect of Ti diffusion to AlN layer. Pt (111) lattice mismatch with AlN (002) is higher than the other two candidates and hence has a lower piezoelectric response.

Table ‎4.2 Piezoelectric coefficient measurements for AlN films deposited with different seed layer systems.

Architecture Thickness Processing parameter Piezoelectric coeff.

2 (nm) (DC-sputtered AlN) [e31 (C/m )]

AlN/Ti/Si 700 250W, 3mT, Pure -0.74 Nitrogen, 3h, -50V bias

AlN/TiN/Ti/Si 700 250W, 3mT, Pure -0.90 Nitrogen, 3h, -50V bias

AlN/Pt/Ti/Si Not measured 250W, 3mT, Pure -0.49 Nitrogen, 3h, -50V bias

4.1.4 Substrate Biasing Effect, DC and P-DC

The fact that AlN is one of the best insulators in nature makes it difficult to apply a DC substrate bias during deposition. The deposition of only a couple of AlN layers on the substrate will make the whole thing insulating, and the subsequent depositing species will be practically unaffected by the bias applied well below the insulting layer. Therefore a pulsed bias was applied to the substrate and found a significant effect.

The crystallinty was significantly improved upon pulse biasing the substrate. The following graph, Figure 4‎ .8, compares the XRD of a pulsed biased sample with an unbiased one. The XRD was shown in logarithmic scale for the same set of samples, the unbiased substrate

100

initiated the growth of some other peaks than the main one, whereas, the pulsed biased substrate has proved to favor (002) orientation (there was only a single extra peak at 59.22 degrees which belonged to AlN (110) plane.

The following SEM graphs shown in Figure 4‎ .9, compares the cross sections and top view of an unbiased film with a pulsed biased film. One not only notices a thicker film for pulsed biased substrates, but also an increase in density.

Figure ‎4.8 XRD peaks of AlN films deposited on Si substrates with pulsed bias (-50 V/20 kHz) in red and unbiased substrate in black.

The comparison in the surface reveals an increase in grain size with pulsed substrate biasing, however, DC biasing also did have a positive effect in the grain size increase. This is not unnatural, since, the metallic substrate holder was behind the silicon substrate, and some amount of the electric field (generated from the charged substrate holder) would have penetrated the semiconducting silicon wafer and interfered with the plasma.

At a biasing frequency of 100 kHz, the structure became flaky, when looked at a larger scale. This structure, however, revealed as a conglomeration of finer rains along a plane. It was probably the effect of the increased surface mobility, which made the small nuclei migrate more along the surface and join with neighboring nuclei. The fact that the joining took place along a

101

plane non-parallel to the substrate was still not very clear. These planes were also not the preferred low energy (002) growth planes of AlN, since the XRD revealed the (002) planes parallel to the substrate. The two SEM micrographs depicted in Figure 4‎ .10 compared the film surface without and with a platinum electrode (<50 nm) on it. The platinized area looked very similar (except a little blurring of finer surface features) to the one without the platinum, and proved that the electrode layer, being very thin, followed the same surface morphology (texture), at least along the viewing direction, and it could be expected to have left no physical gap between the film and the electrode.

The film surface, consisting of planner (and non-parallel to substrate) agglomerates of finer grains indicated a rough surface, and hence a very irregular electric field distribution. Also, due to the surface roughness would cause significant shadowing effect to the sputtered platinum atoms during electroding process, which in turn would make a non-continuous electrode layer, and further increase the field inhomogeneity (possibly in a negative way). This could have been the result of a decreased piezoelectric constant in pulse biased AlN films as is recorded in Table 4‎ .3. Another possible degradation mechanism for piezoelectric response of pulsed bias substrate could be due to the fact that additional peak for AlN (110) existed in all of the pulsed biased samples.

The AlN films with excellent and near epitaxy (with narrow FWHM of rocking curve around 1 ) can provide a negligible piezoelectric response. Inversely, perfectly (002) oriented films with poor crystal quality (with wide rc-FWHM of up to 20 ) can still exhibit good piezoelectric coefficients [140].

These behaviors are explained by the presence of other than (002) oriented crystals in the structure of AlN films, that results in a presence of grains with opposite crystallographic polarities along the normal of the surface (inversion domains), which significantly reduce the net piezoelectric field [141].

102

Figure ‎4.9 Effect of Substrate biasing: DC and pulsed DC on grain size and microstructure of DC-sputtered AlN films with cross-section and top view SEMs.

103

Substrate Film

Figure ‎4.10 SEM micrograph of a thin Pt electrode deposited on a rough film surface compared to non-platinized AlN film.

Table ‎4.3 DC-sputtered AlN films deposited on DC and pulsed-DC biased substrates with similar optimized deposition parameters.

Architecture Processing parameter Piezoelectric coefficient, e31 (C/m2) (DC-sputtered AlN)

AlN/TiN/Ti/Si DC substrate bias: -50V -0.90

AlN/TiN/Ti/Si Pulsed-DC substrate bias: -0.63 -50V/20 kHz/2microSec

AlN/TiN/Ti/Si Pulsed-DC substrate bias: -0.69 -50V/100kHz/2microSec

4.2 Pulsed-DC Reactive Sputtering of AlN Films

AlN thin films were deposited with four-cathode closed field unbalanced magnetron sputtering and three-cathode Semicore units. To further study the effect of pulsing frequency on film properties, an Advanced Energy Pinnacle Plus with pulser unit was used as power supply.

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Here is the summary of pulsed-DC sputtered AlN films optimized based on the c-axis texture by changing one deposition parameter at a time while others are held constant.

4.2.1 Working Pressure Effect

AlN thin films grown at different working pressures were deposited using the P- CFUBMS chamber, the crystal size, residual stress, and c-axis texture / (002) preferred orientation were studied by conventional XRD and ω-scans.

According to the kinetic theory of molecular gas [66], the mean free path of a gas molecule at constant temperature has a value that is inversely proportional to the pressure. At low sputtering pressure, the mean free path of the species increases owing to the decrease of particle collision and thus the energetic particles in the plasma can easily transfer their energy to the adatom at the growing surface. This effect can promote the growth of highly c-axis oriented AlN films at room temperature [142]. On the other hand, at low working pressures due to larger mean free path of particles, highly energetic ions and fast neutrals could approach the substrate due to fewer collisions and introducing defects and residual stress inside the coating, and further instability of plasma ignition [143, 144].

Generally, the magnitude of compressive stress drastically increases with decreasing sputtering pressure. This could be understood in terms of internal stress accumulation by relatively high momentum transfer at the film surface in low sputtering pressures. In AlN crystalline cell, the bond in c-axis direction is formed by the coupling of the Al empty orbit and the N full orbit and hence the ionic character of this bonding is greater than other Al-N bond in [100] direction. Therefore sufficient energy of arriving particles is required for the formation of c-axis oriented film but excessive amount of impinging particle/ion would introduce defects and stress resulting in a mixed-oriented AlN film [126].

105

Table ‎4.4 Deposition parameters of AlN films deposited at different working pressures of 2-6.5 mTorr.

Architecture Reverse Power (W) N2 /Ar ratio, Frequency Pressure Time Bias (V) (KHz) (mTorr) (μs)

AlN/Cr/Si 1.0 1000 100/0, FB 100 2.0

AlN/Cr/Si 1.0 1000 100/0, FB 100 3.0

AlN/Cr/Si 1.0 1000 100/0, FB 100 4.0

AlN/Cr/Si 1.0 1000 100/0, FB 100 5.0

AlN/Cr/Si 1.0 1000 100/0, FB 100 6.5

The AlN films were deposited varying the working pressure ranging from 2 mTorr up to 6.5 mTorr to study the residual stress evolution, crystal size and c-axis orientation of the AlN thin films, as depicted in Figure 4‎ .11 (a-c), respectively.

According to the kinetics theory of gases, at higher gas pressures, more ion and particle collision would occur which would lower the energy of incoming particles arriving towards the substrate and growing layers thus producing films with lower compressive stress level due to shadowing of incoming particles with oblique angles. On the other hand, at low-working pressure situation the mean free path of particles is larger and the substrate is encountered by more energetic ions and neutrals creating a denser film but with excessive amount of residual stress and possible defects formation at extreme energies.

Therefore, there is an optimum working pressure that would produce highly c-axis oriented AlN film. The grain size (columnar diameter) of the growing film increases with increasing pressure and reaches a maximum at 5 mTorr, as depicted in Figure 4‎ .11 (b), and decreases at 6.5 mTorr. Films deposited at 5mTorr showed the highest grain size which is in agreement with XRD data and also exhibit approximately 98% (002) texture.

106

Figure ‎4.11 The working pressure effect on residual stress (a), crystallite size (b), and XRD rocking curves (c).

Highly (002)-oriented growth promote larger crystallite size due to the narrowing of (002) diffraction peak in conventional XRD. At lower working pressures, the neutral particles and ions approach the growing film surface with higher velocity and hence produce more nucleation sites which would reduce the grain size. This grain size would become larger as the working pressure increases but at excessive pressures, the shadowing effect would take over and crystalline size drops again [146].

AlN films deposited at 5mTorr of working pressure had less residual stress, as depicted in Figure 4‎ .11 (a and c). The amount of compressive stress inside the AlN films decreased

107

significantly from 3 to 4 mTorr and then became constant at 5 and 6.5 mTorr. The transition in residual stress from compressive to tensile is believed to happen at higher pressures but due to the significant damage that higher pressures would produce to the vacuum components, this was not studied. The smaller full with at half maximum (FWHM) and stronger peaks in rocking curve, would indicated the lower amount of residual stress with improved c-axis texture. Thus, the film deposited at 5mTorr exhibits the lowermost residual stress, larger crystal size and higher texture in preferred (002) orientation. Furthermore, cross-sectional FESEM micrographs of pulsed DC sputtered AlN films at different pressures are shown in Figure 4.12. Films all exhibited columnar growth at different working pressures. The columnar width increases as the pressure rises from 2 to 5 mTorr and eventually decreases at 6.5 mTorr. This is in good agreements with calculated grain size from XRD peak broadening using Scherrer‘s equation.

Figure ‎4.12 Cross-sectional FESEM micrographs of pulsed DC AlN films deposited with working pressure ranges from 2 to 6.5 mTorr.

108

4.2.2 Target Pulsing Frequency Effect

The power applied to target material can be in the form of rectangular shape pulses where a small portion of the voltage waveform is in the positive potentials to remove any charging of the target material. Depending on what polarity is applied to the cathode, ions and electron can be directed to the target removing material of interest. At lower pulsing frequencies (higher duty cycles), larger flux of atoms are sputtered but with less energy while during longer pulsing reverse time, more energetic atoms are injected. In order to control the texture and density of the films, one needs to choose an optimum ion energy that is sufficient for texture growth but at the same time not too excessive to introduce defects in the films with proper ion flux. This change in ion energy and ion flux has been studied previously [147] and indicated that in pulsed-dc unbalanced closed field systems, the ion energy and ion flux both increase with increasing the pulsing frequency. Higher energy tails were observed at higher pulsing frequency as shown in Figure 4‎ .13.

Figure ‎4.13 Nitrogen ion energy distribution as a function of pulsing frequency and reverse time [147].

109

C-axis or (002) texture is identified as the parameter to optimize in order to obtain higher piezoelectric signal. AlN deposited on the glass substrates were selected to omit any substrate peak overlapping. The degree of preferred (002) orientation increased from around 91.8% at 0 KHz to above 98.6% at 150~200 KHz, then decreased to 94.6% at 250KHz and 90% at 300KHz, as shown in Figure 4‎ .14. At DC sputtering (0 kHz), the normalized peak intensity was low (~92%). This could be due to a lower ion flux in DC sputtering [91] with several eV of energy that is not sufficient to produce highly (002) oriented films which has a higher bonding strength. But further increase in pulsing frequency in the range of (100-200 kHz) produced ions with appropriate energies and flux compared to its DC counterpart which promoted c-axis orientation of AlN films. Further increase in pulsing frequency to 300 kHz introduced ions with energies of about 100 eV or higher. This excessive energy tail produces defects and deteriorates the coating preferred orientation.

Figure ‎4.14 Normalized preferred c-axis growth of AlN films deposited at various pulsing frequency.

110

Table 4‎ .5, Summarizes the direct piezoelectric measurements that also proves the fact that at higher pulsing frequencies the ion energy and flux is greater than the lower frequency situation and enhance the (002) orientation and hence piezoelectric response.

2 Table ‎4.5 Piezoelectric coefficient e31 (c/m ) for AlN film deposited in pulsing frequencies of 100, 200 and 300 kHz.

Piezoelectric coefficient Architecture Processing parameter 2 e31 (C/m )

100kHz/1microSec/1kW/closed AlN/TiN/Ti/Si -0.69 field/3mT/3hr

200kHz/1microSec/1kW/closed AlN/TiN/Ti/Si -0.82 field/3mT/3hr

300kHz/1microSec/1kW/closed AlN/TiN/Ti/Si -0.92 field/3mT/3hr

When the pressure increased to 5 mTorr and slightly decreased the frequency to 100 KHz, a highly (002) growth (98.8%) was obtained. The plasma energy at lower pulsing frequency and higher pressure also provided appropriate ion bombardment for the preferred oriented growth as was explained previously. Additionally, it established the smallest FWHM or rocking curves as shown in Figure 4‎ .15.

XRD-rocking curve is one of the main methods to study the texture and residual stress of the films. Figure 4‎ .15 shows the rocking curve plots of the samples at various pulsing frequencies. The lowest FWHM is around 6° which was acquired for the AlN film deposited at 5 mTorr as was expected based on the prior findings.

111

Figure ‎4.15 Comprehensive Rocking curves for film AlN/Cr/glass deposited at various pulsing frequencies.

The residual stress of the AlN/Cr films was measured through XRD method. The influence of deposition parameters on the residual stress is shown in Figure 4‎ .16. Stress shifted to more compressive for all the samples deposited at (150-300 kHz) pulsing frequency with floating bias on the substrate. Substrate bias attracted more ions and caused more ion implantation, which brought ‗compressive‘ to the film.

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Figure ‎4.16 Dependence of Residual stress on the deposition parameters such as various pulsing frequencies and substrate bias.

When pulsing frequency changed from 0 to 100 kHz at floating bias, the residual stress was reduced significantly due to less arc and defect formation. Bombardment of ions and particles produced higher density and defects at this condition. AlN thin films sputtered at various pulsing frequency were examined in cross section mode using JEOL JSM-7000F Field Emission Scanning Electron Microscope as shown in Figure 4‎ .17. These Cr columns initiated from the substrates and grew to the AlN/Cr interface while the AlN initiated from the interface and grew continuously to the film surface.

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Figure ‎4.17 FESEM Cross-sectional micrograph of AlN/Cr films on Si substrate deposited at different pulsing frequencies.

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Seed layer and its orientation play a big role in the texture of AlN deposited on top. Cr has been grown in preferred orientation of (110) which is the closed packed planes in its BCC structure having the same hexagonal symmetry as the basal planes in AlN films with slight difference in lattice parameter.

4.2.3 Effect of Nitrogen/Argon Ratio

Previously optimized pulsed-DC AlN films with respect to working pressure and pulsing frequency were further investigated by changing the reactive gas ratio during sputtering while other parameters are held unchanged. Table 4‎ .6 summarizes the deposition parameters that were altered during this stage film optimization.

Table ‎4.6 Deposition parameters of AlN films deposited at various nitrogen/argon ratios.

Architecture Pressure Frequency Power Reverse N2 /Ar ratio, (mTorr) (KHz) (W) Time Bias (V) (μs)

AlN/Cr 5.0 100 1000 1.0 20/80, FB

AlN/Cr 5.0 100 1000 1.0 40/60, FB

AlN/Cr 5.0 100 1000 1.0 60/40, FB

AlN/Cr 5.0 100 1000 1.0 80/20, FB

AlN/Cr 5.0 100 1000 1.0 100/0, FB

As mentioned before, it is well known that the formation energy of the closed-packed (002) plane of wurtzite AlN is greater than that of the loosely packed (110) or (100) planes [148]. Therefore moderate energy condition will promote the formation of (002) plane. Ar+ ion with

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+ heavier atomic mass with respect to N2 and larger sputtering yield will transfer more energy to the target and generate more Al atoms/ions. Thus, at lower nitrogen percentage conditions, the ejected Al particles will have higher kinetic energies to rearrange atoms in (002) plane [149– 151].

Furthermore, the crystals with closed packed (002) plane parallel to the substrate will have higher probability of survival under the energetic bombardment because of the higher formation energy in (002) plane [152]. Other crystals with loosely packed planes like (110) and (100) parallel to the substrate will undergo higher etch rate than those with closed packed (002) plane parallel to the substrate surface [153]. Increase of nitrogen concentration will lower the kinetic energy of sputtered ions which undergo more scattering due to higher scattering probability and hence the (002) preferential orientation growth is deteriorated.

When the reactive gas such as nitrogen and inert gas like argon are introduced through the chamber, they will ionize due to the potential difference between the target and substrate then both types of these ions will accelerate towards the target and participate in sputtering process. The results shown in Figure 4‎ .18, illustrates that the films deposited at high nitrogen percentage did not have a good (002) orientation, as was explained in detail above, despite the fact that all the films were deposited with the same other parameters and thicknesses.

To study the composition, stoichiometry and atomic bonding nature of AlN samples as a function of increasing nitrogen concentration in plasma were examined by XPS. Figure 4‎ .19 (a) and (b) depicts the Al 2p and N 1s core level peaks of AlN films deposited at 20/80, 40/60, 80/20 and 100/0 nitrogen/argon ratios, respectively. Ar+ ion etching has been used to clean the samples from any contaminants. Despite the etching and cleaning step, the oxygen peak was still present in the films, which could indicate that oxygen had been diffused deeply into the film.

Figure 4‎ .19 (a)-(b) and Figure 4‎ .20 (a)-(b) show the Al 2p and N 1s peak and their Gaussian fit at 73.5 eV and 396.5 eV, respectively. The 73.5 eV and 396.5 eV binding energies are allocated to Al-N and N-Al atomic bonding respectively [154]. There are no shoulder peaks evident, such as those that would indicate the presence of Al-O peak at 75 eV or Al-Al peak at 72 eV, in the Al 2p spectra. This latter result indicates that aluminum atoms are just bonded to

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nitrogen atoms and aluminum oxide or metallic aluminum is not present in these films. Important information that can be obtained from XPS is the atomic composition of AlN thin films that can be quantified by the area under each peak. Table 4‎ .7 shows the atomic concentrations of every atom in the films with their Al to nitrogen atomic ratio for the formation of stoichiometric AlN.

Figure ‎4.18 Conventional XRD peaks of AlN/Cr films deposited on glass at different nitrogen gas ratio at 3 and 5 mTorr.

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Figure ‎4.19 (a) Al 2p and (b) N 1s core level spectra of one-hour argon-etched AlN films deposited with different nitrogen to argon gas ratio.

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C P S Figure ‎4.20 Gaussian curve fitting of (a) Al 2p and (b) N 1s core level spectra of one-hour argon- etched as a function of nitrogen/argon ratio.

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Table ‎4.7 XPS compositional and atomic concentration of AlN films deposited at different nitrogen ratio and their Al/N atomic ratio after one hour ion etching.

Reactive Nitrogen % Atomic Conc. % Atomic Conc. % Atomic Conc. Al/N Atomic Ratio ratio (O 1s) (Al 2p) (N 1s)

(20/80, N2/Ar) 12.04 39.42 48.53 0.81

(40/60, N2/Ar) 7.32 48.04 44.04 1.09

(80/20, N2/Ar) 6.46 48.33 44.52 1.08

(100/0, N2/Ar) 10.58 46.46 41.15 1.12

For the films deposited at 20 percent of nitrogen to argon ratio, the atomic ratio is below unity which indicates the formation of sub-stoichiometric films. But at higher nitrogrn to argon gas ratios, the atomic ratio of Al/N is close to unity which indicates the formation of stoichiometric AlN films. Since there was no shoulder peak indicating the formation of Al-O bonding, the oxygen is most likely to be free and being diffused inside the films after exposure to the air.

Moreover, the residual stress and (002) grain size variation, shown in Figure 4‎ .21, as a function of nitrogen percentage during AlN deposition were also calculated using Equation 3.13 and Scherrer‘s equation, respectively. AlN stiffness constants in this stress evaluation were as follows: ( c11395 GPa , c 12  137 GPa , c 13  108 GPa , c 33  373 GPa ) [155].

Under constant power deposition, increasing the N2 gas fraction would increase the degree of the target nitridation (poisoning effect) and hence higher cathode voltages with lower cathode current. This was apparent in all of the depositions done at a constant pressure and power. At pressure of 5 mTorr and constant target power of 1kW, the observed cathode current and voltage are summarized in Table 4‎ .8.

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Figure ‎4.21 Residual stress measurements of AlN films deposited at ―5mTorr, 100 kHz, 1microsecond, FB‖ (top), and the calculated (002) grain size (bottom).

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Table ‎4.8 Cathode current, cathode voltage, and nitrogen flow rate for AlN films deposited at 5 mTorr, 1 kW, 100 kHz, and 1 microsecond reverse time.

Nitrogen gas ratio Target Current Target Voltage Nitrogen flow (A) (V) rate (sccm)

AlN (20% N2) 4.59 218 13.25

AlN (40% N2) 4.05 247 23.80

AlN (60% N2) 3.88 258 33.76

AlN (80% N2) 3.72 269 42.40

AlN (100% N2) 3.57 280 51.30

Figure ‎4.22 FESEM cross-sectional micrograph of AlN/Cr films on Si substrate deposited at 5 mTorr, 1 kW, 100 kHz, and 1 microsecond reverse time with different nitrogen gas ratio.

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Initial particle energy increases proportional with cathode voltage and the flux is proportional to the discharge current [156]. Therefore at higher nitrogen gas ratio, the higher cathode voltage would result in a smaller flux of energetic particles with higher energies. These higher energy particles with lower flux are believed to be the source of c-axis deterioration and production of defects thus higher compressive stresses at high nitrogen concentrations. These higher compressive stresses at higher nitrogen concentration often produced delamination the film from the substrate as was observed in the AlN film deposited at 80:20 nitrogen to argon ratio in Figure 4‎ .22

4.2.4 Seed-Layer Effect (Mo and Pt)

AlN films with additional seed layer materials (Pt and Mo) were deposited on the stainless steel substrates as presented in Table 4‎ .9. The first pair of films was deposited at two different working pressures of 3 and 5 mTorr to study the effect of deposition pressure on the orientation of these films.

Table ‎4.9 Deposition summary of seed layers (Pt and Mo) sputtered on stainless steel substrates.

Seed-Layers Pressure Frequency Power Reverse N2 /Ar ratio, (mTorr) (KHz) (W) Time Bias (V) (μs)

AlN/Pt 3.0 100 1000 1.0 40/60, FB

AlN/Pt 5.0 100 1000 1.0 40/60, FB

AlN/Mo 5.0 100 1000 1.0 40/60, FB

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As it is shown in Figure 4‎ .23 (a), the AlN film that was deposited at 5 mtorr showed higher degree of preferred (002) orientation as was expected according to previous findings. The AlN was also deposited on Mo seed layer with preferred (110) texture.

Figure ‎4.23 XRD graphs of deposited AlN on stainless steel (a) at 3 and 5 mTorr using Pt and (b) comparison of XRD of AlN deposited on Pt and Mo.

Pt was also sputtered with the same deposition parameters to see the effect of seed layers (Mo and Pt) on the (002) texture of AlN thin films. As illustrated in Figure 4‎ .23 (b), both thin 124

films deposited on Pt and Mo showed AlN (002) peaks on the stainless steel substrates but the AlN film that was deposited on Mo seed layer illustrated higher preferential orientation (FWHM of conventional XRD peak decreased slightly from 0.28 to 0.20 degrees), because the (110) planes in body centered cubic molybdenum is a closed pack plane with hexagonal atomic symmetry and has a good lattice matching with AlN (002) planes hence an epitaxial film could be grown using this under-layer. The grain size of AlN film deposited on Mo seed layer, of about 42 nm, did not show much improvement over other buffer layers investigated previously, indicating that the texturing was only improved based on the hexagonal atomic symmetry of Mo (110) and lattice matching thus was not energetically enhanced.

The previously optimized film deposited on highly doped Si was electroded with sputtered gold using mask spot size of1mm2 . Then the film was wired and tested using the homemade laser interferometry calibrated by a transducer with a known displacement modulation. The result is shown in Figure 4‎ .24.

Figure ‎4.24 Piezoelectric coefficient (d33eff) measured by interferometry for optimized AlN film deposited on Mo seed layer.

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Since the piezoelectric film is rigidly clamped to the substrate, then the substrate constrains the in-plane contraction and expansion of the film. Therefore only an effective value of d33,eff is measured instead of d33 . Thus, a correction term involving the in-plane piezoelectric coefficient d31 and the elastic compliance s11,, s 12 s 13 is used in order to calculate the unclamped value of d33 [157–161]:

ds31 13 dd33 33,eff 2 Equation ‎4.3 ss11 12

The values for elastic compliance are taken from [155], and d31 is assumed to be d33,eff /2.

Calculated piezoelectric coefficient d33 of unclamped AlN is about 4.2 pm/V. This can be translated back to the direct piezoelectric coefficients by the equation given below [162]:

2 d31 e31(/) c m  Equation ‎4.4 ss11 12

2 Therefore, an effective e31 for this optimized film is about 0.93 C/m which is about 90 percent of theoretical piezoelectric value reposted in literature [162–164].

4.2.5 High Temperature Stability (DSC and In-Situ HTXRD)

To study the thermal stability of AlN thin films at elevated temperatures, DSC method is employed as shown in Figure 4‎ .25 (a), and the films were further evaluated by XRD to report any changes in phase or texture, as depicted in Figure 4‎ .25 (b).

The AlN deposited on silicon with Ti/TiN as seed layer/electrode were heated from room temperatures up to 1000°C with a heating rate of 20 K/min in a flowing of argon (55 sccm) environment. The DSC data and its derivative show no apparent phase change below 1000 °C,

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confirming the superior thermal stability of these AlN films for aluminum die-casting applications at about 700 °C.

Figure ‎4.25 (a) DSC data and its first derivative acquired in argon environment (55 sccm) and (b) the corresponding XRD data before and after DSC test.

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In order to recognize any texture or phase change of the films during the DSC test, XRD graphs were obtained for the films prior and after heat cycle. As it shown in Figure 4‎ .25 (b), there is not much change in AlN c-axis orientation and films before and after DSC test look very similar which is very promising result indicating the stability of AlN thin films up to 1000 °C without any phase change. The (002) peak of AlN is not changed in its intensity after DCS test compared to as-deposited film which indicates the excellent thermal stability of AlN films at high temperatures. The shift in (002) peak of AlN after DSC analysis is most likely due to the compressive stress relief during high temperature treatment. The thermal stability is one of important parameters to evaluate AlN in high temperature application. The DSC curve indicated that there was no apparently phase change or reaction below about 1000°C.

Figure ‎4.26 FESEM cross-sectional and plan view of AlN film before and after DSC 1000 °C in argon environment.

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The cross sectional and top view microstructure of AlN before and after DSC were investigated by FESEM, as shown in Figure 4.26. Both plan view and cross section microstructure indicated that the size and columnar structure were not affected by annealing. Above all, AlN exhibited strong stability upon high temperature annealing (1000°C) in inert environment.

Figure ‎4.27 In-situ XRD spectra of AlN film studied at temperatures of 600, 700, 800, 900, and 1000 °C.

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In-situ XRD measurements was also performed in He controlled environment on the optimized AlN sample which was deposited at 5 mtorr pressure, 100 kHz, 1kW, and 40 percent nitrogen, at 600, 700, 800, 900, and 1000 degree C, respectively and holding it for one hour to study its phase stability behavior at the above temperatures. The films that were examined in controlled environment of He did not show any sign of oxidation of any layers and the peak intensities also remained unchanged, as illustrated in Figure 4‎ .27.

4.2.6 High Temperature Oxidation Resistance

In-situ XRD measurements were conducted for a non-(002) and highly (002) oriented AlN films in air at temperatures of 600, 700, 800, 900, and 1000 degree C, respectively, holding it for one hour to study its oxidation and texturing behavior at the elevated temperatures. As shown in Figure 4‎ .28, the AlN film tested at mentioned temperatures did not show considerable crystalline aluminum oxide formation indicating that the film is stable against oxidation up to 1000 degree C.

AlN films held at 600 degree C showed the oxide peaks corresponding to chromium oxide where the under-layer was more prone to oxygen. The films exhibited the same XRD patterns up top 900 degree C which the oxidation of substrate silicon initiated. The AlN film did not show any oxidation until the temperature reached 1000 degree C. Also, the (002) preferred orientation of the weakly crystalline film at room temperature was enhanced and maintained at high temperatures indicating that the high temperature energy transferred the crystallinty of AlN films to its most thermodynamically stable orientation. The Cr peak also disappeared at temperatures above 900 degree C, signifying that chromium was totally oxidized above this temperature.

A highly (002) oriented AlN film was also investigated with in-situ high temperature XRD, as shown in Figure 4‎ .29, at the same temperature scans as prescribed before. The film was held at each temperature step for one hour before proceeding to the next step.

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Figure ‎4.28 In-situ high temperature XRD scans for non-textured AlN/Cr film at temperatures of 600, 700, 800, 900, and 1000 degree C in air, respectively.

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Figure ‎4.29 In-situ XRD scans of AlN sputtered on TiN/Ti/Si at 600, 700, 800, 900, and 1000 degree C in air.

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The AlN film was highly (002) oriented with a small (110) peak at room temperature. When the temperature reached 600 °C, the silicon substrate started to react with oxygen in air producing silica. Since the thickness of silicon (substrate) is more than two orders of magnitude higher compared to the thin AlN film and its under-layers, then the penetrated XRD will easily detect the formation of crystalline silica at lower temperatures. After temperature reached 800 °C, the crystalline titanium oxide peaks started to appear which was maintained up to 1000 °C.

The AlN film of interest, did not exhibit any crystalline aluminum oxide peak until the temperature approached 1000 °C. Multiple alumina peaks were presented as depicted in Figure 4‎ .29 with no crystalline AlN peak at this temperature. Therefore AlN showed good high temperature oxidation resistance up to 900 °C and the onset of it oxidation is observed to be about 1000 °C.

It is interesting to note that, at 1000 °C, all the AlN, TiN, Ti, and Si peaks are disappeared and only the oxide peaks composed of alumina, titanium oxide and silica exist and the whole coating system (AlN/TiN/Ti/Si) has been oxidized at 1000 °C. This onset of oxidation temperature of AlN film tested here is in good agreement with reported results in literature [56], [165–167]. Residual stress measurements were also obtained for the AlN films treated in temperature ranges of 600 °C to 1000 °C in controlled He atmosphere, as is shown in Figure 4‎ .30.

The (002) grain size is also calculated for these temperatures to study the stability of AlN at higher temperatures. Table 4‎ .10 presents the grain size calculated from XRD data which indicates that grain were slightly larger when the temperature reached the 600 °C from room temperature.

But, the grain size was maintained constant over the temperature range of 600 °C to 1000 °C. This also agrees with the previous cross sectional SEM images of AlN films after DSC treatments of 1400 °C which did not show considerable columnar width change compared to as deposited films (Figure 4‎ .26).

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Table ‎4.10 Grain size calculation from XRD of AlN (002) oriented film from room temperature to 1000 °C in controlled He atmosphere.

Temperature Range FWHM (°) (002) Grain Size (nm)

Room Temperature 0.192 43.2

600 °C 0.190 44.0

700 °C 0.195 43.0

800 °C 0.197 42.5

900 °C 0.199 42.0

1000 °C 0.200 41.8

Figure ‎4.30 Residual stress relaxation of sputtered AlN films at 600 °C to 1000 °C in He controlled environment. 134

The calculated residual compressive stress of (- 4.38 GPa) existed in the pulsed DC sputtered AlN film at room temperature which was further released as the temperature rises to 600 °C in the controlled He environment. This exponential drop in compressive stress was observed for temperatures up to 1000 °C, where the coating stress was minimal.

In contrary, when the films were heat treated at 800 °C for 24 hours long in air, it appears that the stress was released in the first fast ramp to 800 °C (3 minutes ramp) from room temperature and then gradually increased as the time past, as illustrated in Figure 4‎ .31. This development of stress in the film at 800 °C for longer period of holding is possibly due to the incorporation of oxygen atoms in AlN lattice which introduced higher stress in AlN films.

Figure ‎4.31 Residual stress evolution of AlN films at 800 °C for 24 hours in air.

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4.2.7 Substrate Heating Effect on (002) Orientation

Previously optimized AlN thin film were further investigated to study the effect of substrate heating during deposition on the (002) orientation of these films. Table 4‎ .11 shows the deposition parameters that were used to deposit these set of AlN films.

Table ‎4.11 AlN deposition summary using different substrate heating from 200 °C to 400 °C during sputtering process.

Architecture Reverse Power (W) N2 /Ar Frequency Pressure Substrate Time ratio (KHz) (mtorr) Temperature (μs) (°C)

AlN/TiN/Ti 1.0 1000 40/60 100 5.0 RT

AlN/TiN/Ti 1.0 1000 40/60 100 5.0 200

AlN/TiN/Ti 1.0 1000 40/60 100 5.0 300

AlN/TiN/Ti 1.0 1000 40/60 100 5.0 400

AlN films with substrate temperatures ranging from 200 °C to 400 °C were sputtered on the stainless steel and silicon (100) substrates. The gas ratio inside the chamber during deposition was kept constant at 40/60 percent of nitrogen to argon, a gas ratio which was the optimum concentration from prior results. As shown in Figure 4‎ .32, the (002) texture of the film is highly improved at substrate temperature of 200 °C possibly due to higher adatom mobility at higher substrate temperatures and hence enhancement of AlN (002) preferred orientation. The temperatures higher than 200 °C did not improve the texture significantly as observed in the figure below. This is probably due to the out gassing from the oxide ceramic (MgO-based) heater element at higher temperatures since the color of the coating also changed indicating the formation of possible oxides.

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Figure ‎4.32 Substrate temperature variation effect on preferred (002) orientation of AlN thin films deposited with optimized sputtering parameters.

The crystalline size (grain size) is inversely proportional to the FWHM of the conventional XRD peak as was shown in Scherrer‘s equation. It has been also seen from previous studies [112], that the FWHM of conventional and rocking curve have linear relation with respect to one another. At 200 °C the FWHM of conventional XRD peak was reduced significantly to a lower value compared to the film sputtered at no substrate heating for the pulsed DC deposited AlN. This indicates that the film structure was further optimized by higher substrate heating during deposition as depicted in Figure 4‎ .33. The FWHM was increased again

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after temperature reached 300 °C-400 °C which was possibly due to the out-gassing of ceramic heater and the deterioration of (002) texture.

Figure ‎4.33 AlN (002) conventional XRD FWHM of films deposited at no substrate heating and substrate temperatures of 200, 300, and 400 °C.

AlN film deposited at substrate temperature of 200 °C with optimized c-axis texture was analyzed with laser interferometry to obtain its piezoelectric coefficient d33 (/) pm V . The voltage applied across the AlN film was varied from 2 V to 10 V and the corresponding displacement was recorded to calculate the piezoelectric coefficient d33,eff . As it is shown from Figure 4‎ .34 (a), the piezoelectric coefficient of film deposited at 200 °C substrate temperature is enhanced compared to the one without the substrate heating. This finding follows the model that was explained earlier with respect to the rocking curve FWHM and piezoelectric response. As the FWHM of rocking curve decreases from 6° to 4.5 °, the crystallinty and (002) texture of AlN is improved leading to higher piezoelectric coefficient.

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Figure ‎4.34 Laser interferometry displacement obtained by the application of increasing applied voltages from 2 to 10 volts for AlN deposited at no substrate heating compared to the one at 200 °C substrate temperature (a), rocking curve FWHM of AlN films deposited at no substrate heating and 200 °C, respectively.

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The unclamped piezoelectric coefficient d33 is derived using equation 4.3. This turns out to be (d33  4.9 pm / V ) which is higher compared to the film deposited without substrate heating (d33  4.2 pm / V ) . This value is similar to the optimum piezoelectric response obtained in literature [168–172]. The inverse piezoelectric response, d33 (/) pm V , can also be back calculated

2 to the direct piezoelectric coefficient e31(/) C m using equation 4.4. This is calculated to be (0.98 C/m2), which is about 95% of its theoretical value reported elsewhere [157] [173]. By changing the applied frequency of modulation at a given voltage (VVpp 10 ), the piezoelectric response d33 of the sample was not affected as depicted in Figure 4‎ .35.

Figure ‎4.35 Piezoelectric coefficient d33 measurements for optimized AlN sample at frequency of modulation ranging from 10 to 100 kHz.

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4.2.8 Transmission Electron Microscopy (TEM) Study

TEM sample preparation was also carried out on the optimized P-DC deposited AlN film by cutting the cross section of the film on silicon substrate and then polishing and thinning them up to 60 micron. Then, the films were dimpled and ion milled to even make the sample thinner to be transparent to the electron beams. A Philips/FEI CM 200 transmission electron microscope (CM 200-TEM) is used to look at the transparent thin sections. Figure 4‎ .36 shows the cross sectional view of AlN thin film exhibiting columnar grains with the average size about 40-60 nm which is with good agreements with the grain size obtained from XRD by using the Scherrer‘s equation. High resolution TEM image of this film, as depicted in Figure 4‎ .37, showed the lattice fringes corresponding to (002) AlN planes with d-spacing of about 0.251 nm.

Figure ‎4.36 TEM cross-section of optimized AlN deposited by P-DC on silicon substrates.

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Figure ‎4.37 Cross-sectional HRTEM image of AlN films deposited at 5 mtorr with its electron diffraction pattern at bottom.

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The diffraction pattern also showed polycrystalline feature of this AlN thin film with mostly (002) planes presented there and with the small amount of (100) planes also presented in the electron diffraction patterns which were not evident in the XRD studies.

4.3 Modulated Pulsed Power Sputtering of AlN Films

Thin AlN films were sputtered previously using DC and pulsed-DC with optimized structure and properties. Further preliminary study was carried out using a new deep oscillation MPP (DOMPP) power supply to deposit AlN films. The preliminary results were very promising for further investigations. Films were sputtered using average power of 1.5 kW, peak voltage of 600 V, peak current of 62 A, in a 30% nitrogen reactive gas ratio as shown in Figure 4‎ .38.

Figure ‎4.38 Deep oscillation MPP voltage and current waveforms showing macro pulses (a) and micro pulses (b).

The AlN film deposited by this technique is purely c-axis oriented and exhibits no other peaks in the conventional XRD runs. To further understand this high c-axis texture, rocking curve XRD was also investigated. As shown in Figure 4‎ .39, the FWHM of this highly textured was greatly reduced from around 6 ° for an optimized P-DC films to about 2.5 ° in DOMPP

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sputtered AlN films. This much smaller rocking curve FWHM confirms the highly (002) texture of AlN film with similar thicknesses.

Figure ‎4.39 Convention XRD of AlN film deposited by DOMPP (a) and its rocking curve comparison with respect to P-DC sputtered AlN film.

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Cross-sectional TEM sample was prepared for above DOMPP sputtered sample to further study its microstructure. The films were dense and showed larger columnar width compared to P-DC sputtered films in general. The SAED pattern showed highly (002) electron diffraction spots which was possibly overlapped with silicon single crystal electron diffraction spots as shown in Figure 4‎ .40.

Figure ‎4.40 Cross-sectional TEM micrograph of DOMPP AlN deposited on Si and its selected area electron diffraction (SAED) on the top.

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Based on the time-averaged ion energy distributions (IEDs) taken with a Hiden analytical electrostatic quadruple plasma analyzer (EQP) in the past in Advanced coating and surface engineering lab (ACSEL) [174], [91] for DC, pulsed DC and MPP sputtered films, as shown in Figure 4‎ .41, the IEDs of AlN sputtered with pulsed-DC resulted in increased ion energy when compared to its DC sputtered AlN film where only low energy ions are observed.

Figure ‎4.41 Argon IEDs for AlN deposited under a DC discharge (a) and P-DC (b) [174].

Furthermore, a comparison between DC, P-DC and MPP sputtered film [91], illustrates the formation of low ion energy particles mainly several eV and maximum ion energy of about 15 eV for DC sputtered films. When the target was pulsed, a wide range of ion energies 30-50 eV were observed with a small numbers of ions with a wide energy distribution exceeded 100 eV. However, MPP plasma behaves like DC plasma with much higher ion numbers (ion flux) where the ion tail increases to 20 eV as shown in Figure 4‎ .42.

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+ + + Figure ‎4.42 IEDs of (a) Cr , (b) Ar , (c) N2 ions and (d) the integrated ion fluxes of Cr target in Ar:N2=1:1 atmosphere under DC, P-DC, and two MPP conditions [91].

High energy tails in the pulsed DC plasma could potentially create defects and deterioration from c-axis texture. By choosing higher ion fluxes without having too high of ion energy tails, the formation of (002) planes which have the highest bonding strength, could be favorable. A comparative residual stress analysis, depicted in Figure 4‎ .43, can also show that the high ion energy tail in the P-DC plasma introduces higher amount of defects and stress while high ion fluxes in the MPP sputtered AlN films with low ion energy distribution can enhance the adatom mobility on the substrate and hence improving the c-axis texture and piezoelectric response of AlN films.

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Figure ‎4.43 A comparative residual stress measurements for AlN sputtered using DC, P-DC and DOMPP.

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CHAPTER 5. CONCLUSIONS

DC Sputtering of AlN: The AlN films deposited by reactive DC sputtering were characterized and optimized to obtain c-axis preferred orientation and enhanced piezoelectric response. Deposition parameters studied in this case were the reactive gas ratio during deposition, seed layer and substrate biasing effect on the AlN films properties and performance. Nitrogen gas ratios above 10 percent produced AlN films with preferred (002) orientation with the average grain size of 30 nm. However, the films showed a second peak of (101) AlN film which indicated that the system is far from its lowest energy state thermodynamic equilibrium. Further annealing treatments at 600 C in nitrogen improved the c-axis texture and insulation characteristic of AlN films, significantly. The films that were deposited thinner than 350 nm showed very poor insulation properties due to a higher surface roughness which induces a strong electric field inhomogeneity that would exceed the breakdown field in some areas. This degraded insulating property of films deposited at lower thicknesses showed also deterioration in piezoelectric response. Seed layer has a crucial effect on the texture and response of hexagonal AlN films. An appropriate seed layer with similar lattice spacing and hexagonal atomic symmetry similar to (002) basal planes of AlN was observed to be beneficial to obtain c-axis

2 oriented films with optimized piezoelectric response (e31 0.90 C / m ). Pulsed substrate biasing improved the (002) orientation of the films but introduced an additional peak of (110) AlN. The AlN films with excellent and near epitaxy (rocking curve FWHM of around 1 °) can produce negligible piezoresponse due to the formation of inversion domains with opposite polarity along the normal of the surface reducing the net piezoelectric field.

Pulsed-DC sputtering of AlN: It should be pointed out that extensive arching was observed on the Al target during reactive sputtering while the target was operated in DC mode due to the formation of insulating AlN film on the target surface covering the Al cathode. In order to minimize the arching formation, the AlN films were deposited by pulsed DC reactive sputtering and the interrelation between processing parameters, microstructure and properties were investigated. Processing parameters altered in the pulsed DC deposition were: working

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pressure, target pulsing frequency, reactive gas ratio, seed layer, and substrate temperature. Various deposition parameters were studied to obtain a highly (002) oriented film with enhanced piezoelectric response.

Based on the kinetic theory of molecular gases, the mean free path of a gas molecule at constant temperature is inversely proportional to the pressure. At low sputtering pressure, the mean free path of the species increases owing to the decrease of particle collision and thus the energetic particles in the plasma can easily transfer their energy to the adatom at the growing surface. This effect can promote the growth of highly c-axis oriented AlN films at room temperature. On the other hand, if the working pressure becomes too low, ions and fast neutrals approach the substrate with excessive energies due to fewer collisions introducing defects and residual stress inside the coating and hence weakening the (002) texture and piezoelectric response.

In AlN crystalline cell, the bond in c-axis direction is formed by the coupling of the Al empty orbit and the N full orbit and hence the ionic character of this bonding is greater than other Al-N bond in [100] or [101] directions. Therefore sufficient energy of arriving particles is required for the formation of c-axis oriented film but excessive amount of impinging particle/ion energy would introduce defects and stress resulting in a mixed-oriented AlN film.

When the films were sputtered in closed field unbalanced magnetron sputtering configuration with pulsed-DC power input, (002) textured AlN films were obtained at lower power density since the ion energy and ion flux are higher compared to DC sputtering system. Cathode frequencies of 100 to 300 KHz improved the c-axis texture of AlN films due to this higher ion energies impinging growing film introducing higher adatom mobility and also produced optimum piezoelectric coefficients. But when the frequency of target exceeded the 300 KHz range, the c-axis texture was degraded due to the formation of very high energy ions (≥100 eV) introducing defects and deteriorating the coating (002) preferred orientation.

Nitrogen, aluminum and argon ions are the major ion species in the reactive sputtering of AlN films. Since argon is heavier gas compared to nitrogen, argon atoms have better sputtering yield than nitrogen atoms. The amount of metal atoms sputtered from the target by the argon

150

bombardment decreased with increase in N2 to argon flow ratio in the chamber. Thus, the amount of ions created from the ionization decreases due to the decrease in argon and metal atoms at higher nitrogen flow. As the nitrogen flow rate is increased in the chamber, the target surface became poisoned due to the formation of AlN insulating compound on the surface with much lower sputtering yield compared to metal target that reduced the deposition rates at higher nitrogen to argon ratios.

Initial particle energy increases proportional with cathode voltage and the flux is proportional to the discharge current. Therefore at higher nitrogen gas ratio, the higher cathode voltage would result in a smaller flux of energetic particles with higher energies. These higher energy particles with lower flux are believed to be the source of c-axis deterioration and production of defects thus higher compressive stresses at high nitrogen concentrations.

An optimized film with respect to c-axis texture was deposited with low substrate heating and showed very high piezoelectric coefficient (d33  4.9 pm / V ) which is comparable to the literature findings. High temperature stability of AlN films were studied in air and inert environment which showed thermal stability of up to 1400 C in controlled environment and oxidation resistance of about 900 , where the AlN film did not show any oxidation before 1000 in air.

MPP sputtering of AlN: The c-axis texture of AlN films were further enhanced exhibiting rocking curve FWHM of about 2.5 which was greatly improved compared to the optimized P- DC (r-FWHM of 4.5 ). Based on the time-averaged ion energy distributions (IEDs) taken with a Hiden analytical electrostatic quadruple plasma analyzer (EQP) for DC, pulsed DC and MPP sputtered films, a conclusion can be drawn as follows: The IEDs of AlN sputtered with pulsed- DC resulted in higher ion energies when compared to its DC sputtered AlN film where only low energy ions are observed.

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Furthermore, a comparison between DC, P-DC and MPP sputtered film [91], illustrates the formation of low ion energy particles mainly several eV and maximum ion energy of about 15 eV for DC sputtered films. When the target was pulsed, a wide range of ion energies 30-50 eV were observed with a small numbers of ions with a wide energy distribution exceeded 100 eV. However, MPP plasma behaved like DC plasma with much higher ion numbers (ion flux) where the ion tail increases to 20 eV.

High energy tails in the pulsed DC plasma could potentially create defects and deterioration from c-axis texture. By choosing higher ion fluxes without having too high of ion energy tails, the formation of (002) planes which have the highest bonding strength, could be favorable. A comparative residual stress analysis can also show that the high ion energy tail in the P-DC plasma introduces higher amount of defects and stress while high ion fluxes in the MPP sputtered AlN films with low ion energy distribution can enhance the adatom mobility on the substrate and hence improving the c-axis texture and piezoelectric response of AlN films.

Based on plasma diagnostics, it was found that the effect of the ion energy can be both detrimental and beneficial to the growing film. The higher ion energy and ion flux can facilitate the ion bombardment on the substrate, increase the adatom mobility resulting in denser film, and change the grain size as well as change the preferred orientation of the film. The excessive ion energies of impinging particles would result in increase in point and lattice defects, and introduced higher residual stress in the structure of the films hence deteriorating c-axis texture of AlN films.

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