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Wear and microstructural integrity of ceramic plasma sprayed coatings Erickson, Lynn C. 1999

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WEAR AND MICROSTRUCTURAL INTEGRITY OF CERAMIC PLASMA SPRAYED COATINGS by L Y N N C. E R I C K S O N B.Sc, Chemical Engineering, Rensselaer Polytechnic Institute, 1977 M.Sc. Engineering/Management, Washington State University, 1987 Technical Licentiate, Materials Science, Uppsala University, 1993 A THESIS SUBMITTED IN PARTIAL FULFILLMENT OF THE REQUIREMENTS FOR THE DEGREE OF DOCTOR OF PHILOSOPHY in THE F A C U L T Y OF G R A D U A T E STUDIES (Department of Metals and Materials Engineering) We accept this thesis as conforming to the-*$quired standard THE UNIVERSITY OF BRITISH C O L U M B I A November 1998 © Lynn C. Erickson, 1998 In presenting this thesis in partial fulfilment of the requirements for an advanced degree at the University of British Columbia, I agree that the Library shall make it freely available for reference and study. I further agree that permission for extensive copying of this thesis for scholarly purposes may be granted by the head of my department or by his or her representatives. It is understood that copying or publication of this thesis for financial gain shall not be allowed without my written permission. Department of AU4^U 4. Aa\e^JLc The University of British Columbia Vancouver, Canada D a t e ;<? r r ^ DE-6 (2/88) ABSTRACT In this work a series of ceramic plasma sprayed (PS) coatings, both alumina- and chromia-based, were sprayed according to a matrix of deposition parameters in order to produce a broad range of microstructures. To investigate the effect of splat size on the coating response, a series of mono-crystalline a-alumina powders with very narrow particle size ranges, nominally 5, 10 and 18 microns in diameter, was sprayed. The coatings were extensively characterized for a variety of microstructural features, including porosity, the angular distribution and density of microcracks as well as the lamellar, or splat, dimensions, using techniques of metallurgical analysis and electron microscopy. The coatings were then evaluated using a series of micromechanical techniques, including indentation, controlled scratch testing, abrasion and dry particle erosion, to investigate their response to different contact situations. It was found that the microstructural features with the most influence on the behaviour of ceramic PS coatings during contact, or wear, by hard particles include, in order of importance: 1) macro-porosity, 2) horizontal crack density, 3) degree of flattening of the splats and 4) volume of unmelted particles, which are all linked to the level and strength of interlamellar bonding in the coating. The major effect of the inter-lamellar bonding in ceramic PS coatings was seen in the wear mechanism transitions. As the level of inter-splat bonding in the coating decreases, the contact load at which the transition from plastic deformation to splat fracture and debonding occurs does as well. However, the load at which catastrophic brittle fracture and spalling occur is increased. All of the micromechanical and wear methods evaluated in the present work were sensitive to differences in the coating microstructures to varying degrees. The low load abrasion results showed the most sensitivity to the microstructural differences of the coatings, followed by controlled scratching. ii TABLE OF CONTENTS Page Abstract i i Table of Contents i i i List of Tables viii List of Figures ix Acknowledgements xix A. BACKGROUND AND LITERATURE REVIEW 1 CHAPTER 1 BACKGROUND 1 1.1 Introduction 1 1.2 Objectives and Scope of the Thesis 5 CHAPTER 2 PLASMA SPRAYED COATINGS 7 2.1 Review of Basic Deposition Process 7 2.1.1 Plasma Spraying 8 2.1.2 Deposition Parameters 11 2.2 Microstructure of Ceramic Plasma Spray Coatings 12 2.2.1 Introduction 12 2.2.2 Lamellar Structure 12 2.2.3 Phase Composition and Structure 14 2.2.4 Porosity 15 2.2.5 Unmelted Particles 16 2.2.6 Impurities 16 2.2.7 Residual Stresses and Solidification Cracking 17 2.2.8 Adhesion to Substrate 19 2.2.9 Post-Deposition Sealing Treatments 20 2.3 Coating Properties 21 2.3.1 Effect of Microstructure on Coating Properties 21 2.3.2 Elastic Properties 22 2.3.3 Hardness 25 2.3.4 Adhesion Strength and Fracture Properties 25 2.3.5 Thermal Properties 27 2.3.6 Wear Resistance 28 2.3.6.1 Abrasive Wear 29 i i i 2.3.6.2 Erosive Wear 30 2.3.6.3 Sliding Wear 31 CHAPTER 3 METHODS FOR EVALUATION OF THE MICROSTRUCTURAL INTEGRITY OF 33 CERAMIC PLASMA SPRAYED COATINGS 3.1 Damage by Hard Particles 33 3.2 Indentation Methods 35 3.2.1 Introduction 35 3.2.2 Depth Sensing Indentation 38 3.2.3 Microindentation 41 3.2.4 Macroindentation and Fracture Toughness Testing 41 3.3 Controlled Scratch Testing 43 3.3.1 Scratch Testing of Ceramics 43 3.3.2 Scratching of Ceramic Plasma Sprayed Coatings 46 3.4 Abrasive Wear Testing 47 3.5 Erosive Wear Testing 50 3.5.1 Erosion of Brittle Materials 51 3.5.2 Erosion of Ceramic Plasma Sprayed Coatings 52 B. EXPERIMENTAL METHODS AND RESULTS 54 CHAPTER 4 MATERIALS 54 4.1 Introduction 54 4.2 Plasma Spray Powders 54 4.3 Coating Preparation 55 4.3.1 Spray Parameters for the Sapphire Alumina Coatings 56 4.3.2 Spray Parameters for Remaining Coatings 58 4.4 Reference Alumina 59 CHAPTER 5 COATING CHARACTERIZATION: 60 MICROSTRUCTURE AND MORPHOLOGY 5.1 Preparation for Coating Microstructural Examination 61 5.2 Porosity Determination 62 5.2.1 A A-Alumina Coatings 62 5.2.2 Other Coatings 66 5.3 Splat Morphology 69 iv 5.3.1 AA-Alumina Coatings 69 5.3.2 Other Coatings 73 5.4 Crack Distribution Analysis 76 5.4.1 AA-Alumina Coatings 77 5.4.2 Other Coatings 79 5.5 Elemental and Phase Composition 80 5.5.1 AA-Alumina Matrix 80 5.5.2 Other Coatings 81 CHAPTER 6 MICROMECHANICAL METHODS 83 6.1 Introduction 83 6.2 Indentation Techniques 83 6.2.1 Depth Sensing Indentation (DSI) 84 6.2.2 Microindentation 90 6.2.2.1 Vickers Microhardness - 90 6.2.2.2 Knoop Indentation 92 6.2.3 Macroindentation and Fracture Toughness Testing 98 6.2.4 Summary 101 6.3 Controlled Scratch Testing 102 6.3.1 Planar Surface Tests 102 6.3.1.1 Scratch Results 103 6.3.1.2 Scratch Morphology 106 6.3.1.3 Subsurface Cracking 111 6.3.1.4 Summary 113 6.3.2 Cross-Sectional Surface Scratch Tests 114 6.3.2.1 Scratch Morphology 116 6.3.2.2 Fracture cone lengths 117 CHAPTER 7 WEAR METHODS 121 7.1 Introduction 121 7.2 Abrasion Tests 121 7.2.1 AA-Alumina Coatings 122 7.2.2 Reference Materials 125 7.2.3 Effects of Particle Hardness 127 V 7.2.4 Abrasive Wear Mechanisms 128 7.3 Erosion Tests 133 7.3.1 AA-Alumina Coatings 135 7.3.2 Reference Materials 141 7.3.3 Erosive Wear Mechanisms 145 7.3.3.1 Erosion at a 90° impingement angle 146 7.3.3.2 Erosion at 45° impingement angle 151 7.3.4 Cross-Sectional Studies / Subsurface Cracking 155 • 7.3.4.1 90 Degree Erosion 155 7.3.4.2 45 Degree Erosion 157 7.3.5 Descriptive Model for Erosive Wear of Ceramic PS Coatings 158 C. DISCUSSION AND CONCLUSIONS 161 Chapter 8 Correlations Between the Microstructural Parameters, 161 Micromechanical Properties, and Wear Performance 8.1 Introduction 161 8.2 Porosity 162 8.2.1 Effect of Porosity on Micromechanical Properties 162 8.2.2 Effect of Porosity on Wear Properties of Coatings 167 8.2.3 Summary 168 8.3 Inter-lamellar (Horizontal) Microcracking 168 8.3.1 Effect of Inter-lamellar Microcracking on Micromechanical Properties 169 8.3.2 Effect of Inter-lamellar Microcracking on Wear Properties 171 8.3.3 Summary 172 8.4 Intra-lamellar (Vertical) Microcracking 173 8.4.1 Effect of Intra-lamellar Microcracking on Micromechanical Properties 175 8.4.2 Effect of Intra-lamellar Microcracking on Wear Properties 175 8.4.3 Summary 178 8.5 Correlations Between the Micromechanical Properties and Wear Rates 178 8.5.1 Effect of Material Hardness on Wear Rates 178 8.5.2 Correlation of Scratch Hardness with Wear Rates 180 8.5.3 Effect of Elastic Modulus on Wear Rates 181 8.5.4 Effect of Fracture Toughness on Wear Rates 182 vi 8.5.5 Summary 186 8.6 Sensitivity Comparison of Evaluation Methods 187 Chapter 9 Summary and Conclusions 190 9.1 Summary of Major Results for Each Method 190 9.1.1 Indentation 190 9.1.2 Controlled Scratch Testing 191 9.1.2.1 Planar Surface Tests 191 9.1.2.2 Cross-Sectional Surface Tests 193 9.1.3 Abrasion 193 9.1.3 Erosion 194 9.2 Contact Zone Size Range and Effects of the Microstructural Features 196 Found in PS Coatings 9.3 Conclusions 198 9.4 Recommendations for Future Work 200 9.5 New Knowledge and Major Contributions 201 References 203 Appendix A Procedure for Mounting and Polishing of Ceramic PS Coatings 211 Appendix B Scratch Wear Mechanism Maps 213 Appendix C Correlations Between the Microstructural Parameters, 214 Micromechanical Properties, and Wear Performance vii LIST OF TABLES Page Table 1.1 Wear Applications and types of Coatings [3, 4] 2 Table 1.2 Scale of microstructural parameters & defects found in ceramic plasma 4 sprayed coatings. Table 1.3 Techniques for applying different types of contact force to the coating. 6 Table 2.1 Comparison of the Basic Thermal Spray Processes [1,9] 8 Table 4.1 Details of Powders Used for Spraying 56 Table 4.2 Processing Conditions for Alumina PS Coatings 57 , Table 4.3 Details of Matrix Spray Parameters 58 Table 4.4 Details of Spray Parameters for Alumina-Titania and Chromia 59 PS Coatings Table 5.1 Analytical Techniques Used in Evaluating Coatings 60 Table 5.2 Splat Diameters (in pm) for AL30 and AT coatings 75 Table 6.1 Relative size of indentation footprint (diagonal length) to pass thickness and 87 splat diameter. Table 6.2 Ratios of elastic moduli measured in different directions. 93 Table 8.1 Microstructural Summary of Evaluated Plasma Spray Coatings 161 Table 8.2 Summary of Relationships between Microstructure and Properties of 186 Ceramic PS Coatings Table 9.1 Scratch Wear Characteristics (Load 0-60N with 100 pm radius indenter) 192 Table 9.2 Abrasive wear characteristics of the evaluated materials. 193 Table 9.3 Erosive wear characteristics of the evaluated materials. 195 Table 9.4 Ranking of the Effect of Microstructural Parameters on Wear by Hard 197 Particles. viii LIST OF FIGURES Page Fig. 1.1 Schematic showing typical microstructure of a thermal sprayed co 3 Fig. 2.1 Schematic of plasma spray torch with (a) radial and (b) axial powder 9 injection system. Fig. 2.2 Graph of temperature vs. heat content for typical plasma gases [10]. 10 Fig. 2.3 Schematic of idealized structure of flattened ceramic splats, or lamellae. 13 (After [15].) Fig. 2.4 SEM photo of PS alumina coating fracture surface showing columnar grain 14 structure within splats (this work). Fig. 2.5 SEM photo of the top surface of an alumina PS coating, showing a network 18 of fine cracks [this work]. Fig. 2.6 Idealized model of coating structure: (a) planar and (b) cross-sectional 23 views. (Adapted from [35].) Fig. 2.7 Stress-strain curves for PS alumina parallel and perpendicular to the coating 24 plane [36]. Fig. 3.1 Schematic of: (a) two-body and (b) 3-body abrasion and (c) erosion 34 processes. (Adapted from [71].) Fig. 3.2 Model for elastic-plastic indentation. (Adapted from [72].) 35 Fig. 3.3 Evolution of Hertzian cone cracks during spherical indentor loading / 36 unloading. (Adapted from [75].) Fig. 3.4 Progression of the crack systems during sharp indentor loading and 37 unloading. (Adapted from [75].) Fig. 3.5 Comparison of Vickers and Knoop indentors for hardness measurement. 38 (Adapted from [41].The load, P, is in kg and diagonal, D, is in mm.) Fig. 3.6 Typical Loading-unloading curves: (a) elastic, (b) plastic, and (c) elasto- 40 plastic materials, with the plastic depth (hp). Fig. 3.7 Schematic of indentation fracture pattern for Vickers geometry [75]. 42 Fig. 3.8 Schematic illustration of the crack network caused by a passing abrasive 44 particle across a ceramic material. Fig. 3.9 Schematic view of the half-cone fracture upon cross-sectional scratching of 47 a ceramic PS coating (Adapted from [94].) ix Fig. 3.10 Effect of the angle of impingement on the erosion rates of ductile aluminum 50 vs. brittle alumina. The curves display the general erosion response and not the actual magnitudes of the wear rates (After [100].) Fig. 4.1 Alumina Powders: (a) 18 pm monocrystalline (Sumicorundum) and 55 (b) 25-40 pm polycrystalline (Plasmalloy). Fig. 5.1 Comparison of average porosity of AA-alumina coatings as a function of 63 powder type and processing condition. (Sample symbols and numbers refer to Tables 4.3-4.4.) Fig. 5.2 O M micrographs of polished coating cross-sections: (a) AA5-1, (b) AA5-4, 64 (c) AA10-1, (d) AA10-4, (e) AA18-1, (f) AA18-4, (g) APS. Fig. 5.3 70% BSE / 30% SE micrographs of polished coating cross-sections: 65 (a) AA5-1, (b) AA5-4, (c) AA10-1, (d) AA10-4, (e) AA18-1, (f) AA18-4, (g) APS. Fig. 5.4 Comparison of average porosity of reference coatings. 66 Fig. 5.5 O M micrographs of polished coating cross-sections: (a) AL30-1, 67 (b) AL30-4, (c) A T I , (d) AT2, (e) CR1, (f) CR2-1, (g) CR2-2. Fig. 5.6 70% BSE / 30% SE micrographs of polished coating cross-sections: 68 (a) AL30-1, (b) AL30-4, (c) A T I , (d) AT2, (e) CR1, (f) CR2-1, (g) CR2-2. Fig. 5.7 Average dimensions: (a) splat diameter (as measured on coating planar 70 surfaces) and (b) pass thickness, as a function of powder type and processing conditions. Fig. 5.8 S E M photomicrographs of changes in splat diameter with increased particle 71 size: (a) AA5-4, (b) AA10-4, and change in heat input: (c) AA18-1 and (d) AA18-4. Fig. 5.9 Surface Roughness (Ra values) of the AA-alumina coatings as a function of 72 powder type and processing, conditions. Fig. 5.10 S E M photomicrographs of combined fracture/top surfaces: (a) AA5-4, (b) 73 AA10-4, (c) AA18-4 and (d) APS (radial injection sprayed AA18 powder). Fig. 5.11 S E M photomicrographs of as-deposited surfaces: (a) AL30-1, (b) AL30-2, 74 (c) AL30-4, (d) A T I , (e) AT2, (f) CR1, (g) CR2-1 and (h) CR2-2. Fig. 5.12 Average splat diameters of the polycrystalline alumina (AL30) and 74 alumina-titania (AT) coatings. Fig. 5.13 Surface Roughness (Ra values) of the reference coatings as a function of 75 powder type and processing conditions. Fig. 5.14 S E M photos of typical fracture surfaces: (a) PL30, (b) A T I , (c) CR1 and 76 (d) CR2 coatings. x Fig. 5.15 The ratio of horizontal to vertical cracks for the AA-alumina coatings, as a 77 function of powder type and processing conditions. Fig. 5.16 Typical angular crack distribution for the 3 powder sizes (AA5-4, AA10-4 78 and AA18-4). Fig. 5.17 Comparison of average linear crack densities [mm"1] for the AA-alumina 79 coatings: (a) vertical microcracks and (b) horizontal delaminations. Fig. 5.18 The ratio of horizontal to vertical cracks for the reference coatings. 80 Fig. 5.19 Comparison of average crack densities [mm 1] for the reference coatings: (a) 80 vertical microcracks and (b) horizontal delaminations. Fig. 5.20 X R D spectra for the alumina coatings deposited in the four runs: (a) A A 5 , 81 (b) AA10 and (c) AA18 (includes the APS deposition of AA18 sapphire powder using commercial radial powder injection system). Fig. 5.21 X R D spectra for: (a) commercial alumina (PL30), (b) alumina titania (ATI. 82 AT2) and (c) chromia (CRI, CR2-1, CR2-2) PS coatings. Fig. 6.1 Schematic of the scale of different indents and microstructural defects, or 84 parameters, on a ceramic PS coating. Fig. 6.2 Comparison of: (a) hardness (H), (b) elastic modulus (E), and (c) H/E, from 85 DSI at 0.1 N load obtained on cross-sectional and planar surfaces for the AA-series alumina coatings. (Sample symbols and numbers refer to Tables 4.2- 4.3.) Fig. 6.3 Comparison of: (a) hardness (H), (b) elastic modulus (E), and (c) H/E, from 86 DSI at 1.0 N load obtained on cross-sectional and planar surfaces for the AA-series alumina coatings. (Sample symbols and numbers refer to Tables 4.2-4.3.) Fig. 6.4 S E M micrographs of indentations on (a) AA10-4 (1.0 N) and (b) APS (1.0 89 N), (c) APS (0.1 N) alumina coatings showing the effects of microstructural variations or defects. Fig. 6.5 Comparison of Vickers microhardness values obtained on: (a) cross- 91 sectional and (b) planar surfaces of the AA-series alumina coatings, at a load of 3.0N. (Sample symbols and numbers refer to Tables 4.2- 4.3.) Fig. 6.6 Comparison of Vickers microhardness obtained on: (a) cross-sectional and 92 (b) planar surfaces for the reference coatings and bulk sintered alumina (SA). (Sample symbols and numbers refer to Table 4.4.) Fig. 6.7 Comparison of material properties: (a) Knoop hardness (H), (b) elastic 94 modulus (E), and (c) H/E, at 10.0 N load obtained on cross-sectional and planar surfaces for the AA-series alumina coatings. (Sample symbols and numbers refer to Tables 4.2- 4.3.) XI Fig. 6.8 Comparison of material properties: (a) Knoop hardness (H), (b) elastic 95 modulus (£), and (c) H/E, at 10.0 N load obtained on cross-sectional and planar surfaces for the reference coatings and sintered alumina (SA). (Sample symbols and numbers refer to Table 4.4.) Fig. 6.9 S E M micrographs (70% BSE, 30% SE) of Knoop indentations both 96 (a) perpendicular and (b) parallel to the coating cross-section, as well as (c) on the top surface of the AL30-2 coating, illustrating the relation of position to the lamellar structure. Fig. 6.10 S E M micrographs (70% BSE, 30% SE) showing variations in indentation 97 morphology due to differences in microstructure of: (a) AA5-4, (b) AA18-4, (c) AL30-2, (d) AT2, (e) CR2-1 coatings and (f) sintered alumina (SA). Fig. 6.11 Comparison of Vickers microhardness at 100 N obtained on the top surfaces 98 of: (a) AA-alumina coatings and (b) reference coatings and bulk sintered alumina (SA). (Sample symbols and numbers refer to Tables 4.3-4.4.) Fig. 6.12 Type L indentation fracture toughness values for the PS coatings and the 99 bulk sintered alumina (SA). (Eqn. 3.2) Fig. 6.13 Type P indentation fracture toughness values for the PS coatings and the 99 bulk sintered alumina (SA). (Eqn. 2.5) Fig. 6.14 S E M micrographs of 200 N indent on AA18-4: (a) overview and (b) close- 100 up of splat cracking. Fig. 6.15 (K/L/H)2 values for the PS coatings and the bulk sintered alumina (SA). 101 Fig. 6.16 Typical (a) acoustic emission signal and (b) tangential force vs. load plots 104 for an alumina PS coating at an linearly increasing load of 0-60 N (100 pm indenter). Fig. 6.17 Comparison of (a) scratch hardness and (b) scratch depth at 60N normal 105 load using a 100 pm radius indenter for the AA-alumina coatings and reference materials. Fig. 6.18 (a) Scratch cross-sectional area, Ap, as a function of indenter load (F„) for a 106 variety of coating microstructures, compared to the bulk alumina (SA). Fig. 6.19 S E M micrographs showing primarily ductile (Regime 1) behaviour: 107 (a) CR1 coating at 20N and (b) AA10-1 at 10N. Fig. 6.20 S E M micrographs showing Regime 2 behaviour, still mostly ductile with 107 some small scale cracking: (a) AA5-1 at 10N and (b) A T at 30N. Fig. 6.21 S E M micrographs showing Regime 3 behaviour, plastic deformation in the 108 groove bottom, with some brittle fracture at the edges of the groove: (a) AA10-4 coating and (b) SA at 40N. xn Fig. 6.22 S E M micrographs showing Regime 4 behaviour, with large scale-like 108 features in groove bottom, and the deformed layer cracking and buckling: (a) AL30-2 and (b) AA5-1 coatings at 60N. Fig. 6.23 S E M micrographs showing Regime 5 behaviour, exhibiting primarily brittle 109 fracture: (a) AA10-4 and (b) SA at 60 N . Fig. 6.24 S E M micrograph of wear groove after repeat scratching (n = 5, Fn = 20N): 109 (a) APS, (b) AA10 and (c) SA. Fig. 6.25 Examples of wear debris from: (a) AA10-4, (b) AL30 and (c) CRI coatings 110 along with (d) bulk alumina (SA), illustrating differences in wear debris morphology. Fig. 6.26 S E M micrographs of single scratch cross-sections: (a) AL30-2, (b) AA18-4 111 and (c) AT2 at 50 N , illustrating the effects of different microstructures. Fig. 6.27 S E M micrographs of AA5-1 repeat scratches (n = 5,Fn = 20N): (a) top 112 surface, (b) cross-section, showing the compressed layer, and (c) close-up showing densification and cracking under the scratch groove. Fig. 6.28 Cross-sections of scratches (n = 5, F„ = 20N): AA18-4 (a) overview and 113 (b) close-up of vertical macrocracks and (c) AT2 overview and (d) close-up of horizontal macrocracks. Fig. 6.29 Photographs of low-load scratch test equipment, mounted inside of a SEM: 115 (a) outside view and (b) close-up of scratch indenter setup. Fig. 6.30 (a) Normal and tangential forces, F n and Ff, plotted against scratch length, 115 with line A ' illustrating the boundary between substrate and coating and (b) S E M micrograph of 3.5 N scratch with fracture cone on cross-section of a typical alumina coating. Fig. 6.31 S E M micrographs of scratch tracks at a load of 3.5N illustrating different 116 damage mechanism regimes: (a) AT coating, showing ductile plowing, (b) APS coating, showing ductile plowing, tensile cracking and edge fracture, (c) CR and (d) AA10 coatings, showing brittle fracture. Fig. 6.32 S E M micrographs of fracture cones on cross-sections of (a) CR coating, 117 showing smooth brittle fracture with no clear sign of a lamellar structure and (b) AA10 coating, showing fracture occurring in shear steps revealing the coating lamellar structure. Fig. 6.33 Comparison of fracture cone lengths at a load of 3.5N for the: (a) A A - 118 alumina coatings and (b) reference materials. Fig. 6.34 Critical distance or cone length, Lc, raised to the 3/2 power as a function of 119 indenter load for representative coatings and bulk alumina (SA). Fig. 6.35 Fracture toughness values estimated using the fracture cone method. 119 Xl l l Fig. 7.1 Schematic of block-on-drum test equipment (Adapted from [115].) 122 Fig. 7.2 Comparison of specific wear rates for (a) 20 pm and (b) 200 pm abrasion of 123 the AA-alumina coatings as a function of powder type and processing conditions. (Sample symbols and numbers refer to Tables 4.2- 4.3.) Fig. 7.3 Factorial graph for 20 pm abrasion of: (a) A 'A5 , (b) AA10 and (c) AA18 124 alumina coatings. (Numbers correspond to coating runs as listed in Tables 4.2-4.3.) Fig. 7.4 Factorial graph for 200 pm abrasion of: (a) A A 5 , (b) AA10 and (c) AA18 125 alumina coatings. (Numbers correspond to coating runs as listed in Tables 4.2-4.3.) Fig. 7.5 Comparison of specific wear rates for (a) 20 pm and (b) 200 pm abrasion of 126 the reference coatings plus sintered alumina (SA) as a function of powder type and processing conditions. (Sample symbols and numbers refer to Tables 4.3- 4.4.) Fig. 7.6 Relative abrasive wear rates for: (a) 20 pm and (b) 200 pm abrasion, plotted 128 against the ratio of the abrasive hardness, Ha, to the surface hardness, Hs, for the evaluated materials. Fig. 7.7 S E M micrographs of 20 pm abraded surfaces: (a) AA5-1, (b) AA5-4, 130 (c) AA10/18* (same mechanism for both AA10 and AA18 coatings), (d) APS, (e) AL30-2 (f) A T and (g) CR coatings, showing plastic deformation, some brittle fracture and splat debonding, or pullouts. Fig. 7.8 S E M micrographs of 200 pm abraded surfaces: (a) A A 5 , (b) AA10/18* 131 (same mechanism for both AA10 and AA18 coatings), (c) AL30, (d) AT2 and (f) CRI coatings, showing plastic deformation, brittle (splat) fracture and splat debonding. Fig. 7.9 Higher magnification S E M micrographs of 200 pm abraded surfaces 132 showing the mechanisms of splat debonding and wear debris formation: (a) AA5 and (b) AL30 coatings. Fig. 7.10 (a) Outside and (b) inside views and (c) schematic of the test equipment 134 used to simulate solid particle erosion. The particles are fed into the center of the rapidly rotating disk where they are accelerated through four radial channels to impact the specimens (Adapted from [117].) Fig. 7.11 Determination of the erosion rate as the slope of the line through the 135 specimen weight loss vs. weight of the erodent data points. Fig. 7.12 Comparison of wear rates for (a) 75 pm, (b) 200 pm and (c) 600 pm erosion 136 using a 45 degree angle of incidence of the AA-alumina coatings as a function of powder type and processing conditions. (Sample symbols and numbers refer to Tables 4.2 and 4.3.) X I V Fig. 7.13 (a) Mass and (b) volume erosive wear rates, at a 45 degree impingement 137 angle, plotted as a function of erodent particle size for several AA-alumina coatings. Fig. 7.14 Comparison of wear rates for (a) 75 pm, (b) 200 pm and (c) 600 pm erosion 138 using a 90 degree angle of incidence of the AA-alumina coatings as a function of powder type and processing conditions. (Sample symbols and numbers refer to Tables 4.2 and 4.3.) Fig. 7.15 (a) Mass and (b) volume erosive wear rates, at a 90 degree impingement 139 angle, plotted as a function of erodent particle size for several AA-alumina coatings. Fig. 7.16 Factorial graph for 75 pm erosion, at 45 degrees, of: (a) A A 5 , (b) AA10 and 140 (c) AA18 alumina coatings. (Numbers correspond to coating runs as listed in Tables 4.2 - 4.3.) Fig. 7.17 Factorial graph for 75 pm erosion, at 90 degrees, of: (a) A A 5 , (b) AA10 and 141 (c) AA18 alumina coatings. (Numbers correspond to coating runs as listed in Tables 4.2 - 4.3.) Fig. 7.18 Comparison of wear rates for (a) 75 pm, (b) 200 pm and (c) 600 pm erosion 142 using a 45 degree angle of incidence of the reference coatings plus sintered alumina (SA) as a function of powder type and processing conditions. (Sample symbols and numbers refer to Tables 4.3-4.4.) Fig. 7.19 (a) Mass and (b) volume erosive wear rates, at a 45 degree impingement 143 angle, plotted as a function of erodent particle size for the reference materials. Fig. 7.20 Comparison of wear rates for (a) 75 pm, (b) 200 pm and (c) 600 pm erosion 144 using a 90 degree angle of incidence of the reference coatings plus sintered alumina (SA) as a function of powder type and processing conditions. (Sample symbols and numbers refer to Tables 4.3- 4.4.) Fig. 7.21 (a) Mass and (b) volume erosive wear rates, at a 90 degree impingement 145 angle, plotted as a function of erodent particle size for the reference materials. Fig. 7.22 S E M micrographs of 75 pm single impact erosion craters at 90° angle of 146 impingement: (a) AA18 (b) AL30 and (c) CR1, showing a combination of plastic deformation and brittle microfracture mechanisms. Fig. 7.23 S E M micrographs of typical 75 pm eroded surfaces at 90° angle of 147 impingement: (a) AT2 and (b) CR1, showing primarily plastic deformation to different penetration depths, along with some microfracture. X V Fig. 7.24 S E M micrographs of 600 pm single impact erosion craters at a 90 degrees 148 impingement angle on: (a) AA18 (b) AL30, (c) AT and (d) CR1, showing the mechanisms of plastic indentation and deformation, and brittle fracture. Fig. 7.25 S E M micrographs of 600 pm single impact erosion craters at an 149 impingement angle of 90 degrees showing details of the splat fracture and wear debris formation process: (a), (b) AL30 and (c), (d) AT coatings. Fig. 7.26 S E M micrographs (250x, lOOOx) of 600 pm eroded surfaces at 90° angle of 150 impingement: (a)AA5, (b) AA10/18, (c) AL30, (d) AT2 and (e) CR1, showing primarily brittle fracture and some splat debonding/delamination. Fig. 7.27 S E M micrographs of typical 75 pm single impact erosion craters at an 151 impact angle of 45 degrees, showing predominantly plastic deformation caused by: (a) blunt and (b) sharp edged particles. Fig. 7.28 S E M micrographs (250x, lOOOx) of 75 pm eroded surfaces at 45° angle of 152 impingement: (a) A A 5 , (b) AA10/18, (c) AL30 and (d) CR, showing primarily plastic deformation and some splat fracture and debonding mechanisms. Fig. 7.29 S E M micrographs of 600 pm single impact erosion craters at an 153 impingement angle of 45 degrees on: (a) AA10 (b) AL30, (c) CR and (d) SA, showing the mechanisms of plastic deformation, and splat fracture. Fig. 7.30 S E M micrographs (250x, lOOOx) of 600 pm eroded surfaces at 45° angle of 154 impingement: (a) A A 5 , (b) AA10/18, (c) AL30, (d) A T and (e) CR, showing plastic deformation, brittle fracture and some splat debonding. Fig. 7.31 S E M micrograph of bulk sintered alumina 600 pm eroded surfaces at 45° 155 angle of impingement showing plastic deformation and some brittle fracture. Fig. 7.32 Cross-sections of 600 pm eroded alumina coatings surface at a 90 degree 156 impingement angle: (a) AA18 and (b) AL30 coatings, showing crack linking and propagation from pre-existing cracks, i.e inter-splat boundaries and vertical solidification cracks; (c) A T and (d) CR coatings, showing evidence of horizontal (lateral) and vertical (median) macrocracks. Fig. 7.33 Cross-sections of 600 pm eroded surfaces at a 45 degree impingement 157 angle: (a) AA18, (b) AL30 and (c) CR coatings, showing micro- and macro-crack propagation mechanisms. Fig. 7.34 Wear mechanisms occurring in erosion: (a) in bulk ceramic materials and 159 (b) in ceramic PS coating with poor interlamellar bonding and a high density of interlamellar cracks. Fig. 8.1 Effect of porosity on measured hardness (Hw) at (a) 0.1N and (b) 1.0N on 163 planar and cross-sectional surfaces for the AA-alumina coatings. xvi Fig. 8.2 Effect of porosity on microhardness (Hv) at 3N for all evaluated materials. 163 Fig. 8.3 Effect of porosity on (planar) Knoop microhardness (Hk) at ION for: (a) all 164 materials and (b) alumina-based coatings only. Fig. 8.4 Relation between scratch hardness measured at 60 N and porosity of the 165 evaluated materials. Fig. 8.5 Effect of porosity on elastic modulus (£) at ION on coating cross-sectional 166 surfaces: (a) perpendicular and (b) parallel to the coating surface. Fig. 8.6 The "index of brittleness", H/Klc< as a function of porosity for the materials 166 with a measured Kic value. Fig. 8.7 Correlation between porosity and: (a) abrasive wear rate using 20 pm 167 particles and (b) the scratch cross-sectional area at a load of 60N, for all of the evaluated materials. Fig. 8.8 Correlation between coating porosity and the linear density of horizontal 169 microcracks. Fig. 8.9 Effect of the linear horizontal crack density on the Vickers hardness (Hv) 170 measured on the cross-sectional surfaces at 3N. Fig. 8.10 The effect of the horizontal crack density on: (a) 20 pm and (b) 200 pm 171 particle abrasive wear rates. Fig. 8.11 The effect of the horizontal, or inter-lamellar, crack density on the cross- 172 sectional scratch area, of the 60 N scratch groove, on the evaluated coatings. Fig. 8.12 Inverse relationship between the horizontal and vertical microcrack 174 densities. Fig. 8.13 Inverse relationship between coating porosity and the linear density of 174 vertical microcracks. Fig. 8.14 Effect of vertical crack density on coating hardness (Hv) measured at 3N. 175 Fig. 8.15 Relation between scratch groove width and distance between vertical 176 microcracks for Group 2 and 3 coatings. Fig. 8.16 Correlation between the scratch groove cross-sectional area at 60N and the 177 linear density of vertical microcracks. Fig. 8.17 Reation between the 600 pm erosion rate at a 45 degree impingement angle 177 and the linear density of vertical microcracks. Fig. 8.18 The 20 pm particle abrasion rate as a function of Knoop hardness, measured 179 at 10N on planar surfaces. xvii Fig. 8.19 Comparison of the small particle abrasion rate as a function of hardness (Hv) for bulk ceramics vs. ceramic PS coatings. The distances a and b illustrate the contributions of different microstructural factors. See text for details. 180 Fig. 8.20 Inverse correlation between the scratch hardness (at 60N) and the abrasive wear rate using 20 pm particles. 181 Fig. 8.21 Effect of elastic modulus (E) on 90 degree impact erosion results with 600 pm particles for all evaluated materials. 182 Fig. 8.22 The specific abrasive wear rate using 20 pm particles as a function of the fracture/ hardness parameter: (a) Klc~I/2H~1/2, and (b) combined with the elastic modulus as E4/5KlcinHin. 184 Fig. 8.23 The specific abrasive wear rate using 20 pm particles as a function of the toughness-hardness parameter Klcl/2Hia times a microstructural factor, M = P", with n = 0.4. 185 Fig. 8.24 Normalized comparison of abrasion evaluation methods, including 20 and 200 pm abrasion and the 60 N scratch cross-sectional area. 187 Fig. 8.25 Normalized comparison of all wear-related evaluation methods. 188 Fig. 8.26 Normalized comparison of 90 degree erosion evaluation methods. 189 Fig. 8.27 Normalized comparison of 45 degree erosion evaluation methods. 189 Fig. 9.1 Comparison of size ranges involved in evaluated forms of contact to the size ranges of the different microstructural features. 196 xviii A C K N O W L E D G E M E N T S This work was carried out at the Metals and Materials Engineering Department at the University of British Columbia, the National Research Council Innovation Center in Vancouver and the Materials Science Division, Uppsala University and was supported by an NRC CRD Grant (NSERC) and Northwest Mettech Corporation. I would like to express my sincere gratitude to my advisors, Associate Professor Tom Troczynski of the Metals and Materials Department and Adjunct Professor Howard Hawthorne ofthe National Research Council, as well as Professor Sture Hogmark of Uppsala University in Sweden, for their collaboration and invaluable guidance and encouragement throughout the course of this work, as well as for allowing me access to the various research facilities. I would also like to express my gratitude to Doug Ross and Herb Tai* at Northwest Mettech Corporation for the spraying of the coatings, along with many helpful technical discussions and recommendations. I am very grateful to my collaborators at Uppsala University in Sweden, including Richard Westergdrd and Docent Niklas Axen, for their technical assistance and recommendations, along with the other members ofthe department for providing a pleasant atmosphere to work in the two months ofthe two years that I was working over there. My thanks are extended to all the members of the Ceramics Group at MMAT for helpful discussions and pleasant diversions, along with a special thanks to Mary Magerfor her cheerful help in the running of the SEM, as well as the members of the NRC Tribology Group and Machine Shop for their technical assistance and recommendations, along with a number of very enjoyable ski days at Whistler/Blackcomb! I would also like to express my gratitude to the operators of the Beanery coffee shop in Fairview, for providing a perfect atmosphere for the writing and brainstorming of this thesis, along with great lattes and homemade soups! Finally, I would like to express my sincere thanks to my Mom for her encouragement and especially Gunnar for his unfailing support and understanding throughout this long process! *(no longer employed at NW Mettech) x i x A. BACKGROUND AND LITERATURE REVIEW C H A P T E R 1 B A C K G R O U N D 1.1 Introduction Surface engineering involves the use of techniques to either modify the surface of the component itself or cover it with a protective wear, heat or corrosion resistant coating. Thermal spraying (TS), one of the principal methods used in surface engineering, refers to a family of material deposition techniques. These techniques have some aspects in common. The coating material in the form of a rod or powder, is heated to a molten state. The heated material is propelled in the form of particles or droplets at high velocities to impact the substrate, where they rapidly solidify while spreading out to overlap and interlock with each other and with the substrate to form a protective coating. These versatile techniques can be used to produce coatings for improved wear, thermal and corrosion resistance, at a relatively rapid rate and with practically no limits on the size of the coated area. Most types of materials can be deposited with TS techniques, i.e. metals, ceramics and polymers [1]. Thermal spraying has been in use for more than 50 years with most improvements over the last 20 years arising from the design of new types of spray equipment and better understanding of the spraying process. Plasma spraying (PS) is one of the most versatile of the newer methods owing to the high plasma temperatures obtained (typically above 11,000°C), high enough that ceramics can be also used for coating surfaces. The size of the U.S. market for all types of thermal spray coatings is forecast to reach $2 billion by the end of the century [2]. Applications of ceramic PS coatings include: thermal barriers (aerospace and automotive applications), wear protection (chemical, textile and paper & printing industries) and even as orthopaedic implants in the biomedical field. Wear protection is one of the major uses of thermal spray coatings in industry today. Table 1.1 lists some of the 1 typical wear applications along with several commonly used types of ceramic and cermet coating materials. Table 1.1 - Wear Applications and Types of Coatings [3, 4] Wear Type Typical Coating Used Applications Abrasion A1 2 0 3 C r 2 0 3 Z r 0 2 WC-Co Cr 3 C 2 -NiCr - piston rods & rod liners - concrete mixer screw conveyors - fuel-rod mandrels - grinding hammers - cutting tools - mill liners & chutes - pump seals Erosion Cr 3 C 2 -NiCr WC-Co C r 2 0 3 A1 2 0 3 - exhaust fans - hydroelectric valves - exhaust valve seats - gas turbine components Adhesion (Sliding) A l 2 0 3 - T i 0 2 WC-Co - impeller shafts - piston rings (internal combustion) & cam followers - fuel pump rotors Dry Sliding WC-Co C r 2 0 3 Cr 3 C 2 -NiCr - paper rolls - piston rings - wire drawing capstans As result of the material deposition process, TS coatings have a lamellar microstructure as illustrated in Fig. 1.1. This lamellar structure is built up of layers of solidified droplets (or splats) combined with the presence of unmelted particles, different material phases, micro- and macro-porosity along with thermally induced microcracks. A TS coating thus has a unique microstructure and behaves differently from a comparable monolithic material. For example, numerous thermally induced microcracks of various lengths form when the splats shrink during rapid solidification [5, 6]. The number and length of such microcracks are 2 affected by the size of the splats. These cracks can propagate partially or all of the way through the splat thickness, which typically varies between 1-3 pm, see Fig. 1.1. Porosity and phase variations that result from incomplete melting during the deposition also reduce the strength of the individual splats. These types of intra-splat defects may be of the order of 0.1 to 10 pm. The size of the individual splats can vary, depending on the size of the precurser powder used, from 10 - 100 pm in diameter. Fig. 1.1 Schematic showing typical microstructure of a thermal sprayed coating. Defects of larger scale occur between the splats. These include macroporosity or voids (10-30 pm), often occuring at splat ends. Poor bonding between splats or lamellae, as well as unmelted particles or unwanted oxide inclusions may also strongly reduce the integrity of TS coatings. These coarse microstructural features (10 to 100 pm) are most characteristic of coatings deposited by the basic flame spray method. The emphasis in this thesis will focus on ceramic coatings produced using two different types of plasma spray (PS) techniques. The microstructural integrity of PS coatings includes all the factors that determine how well the coating 'holds together' and performs under various contact conditions. These include the whole size range of microstructural defects as listed in the table, in order of increasing scale of events. Unmelted Particle • Pores/Voids 1-3 pm Inclusions Intralamellar Cracking 3 Table 1.2 Scale of Microstructural Parameters and Defects Found in Ceramic Plasma Sprayed Coatings. Microstructural Parameters Size Range (pm) % (Typical Volume or Area) Intra-and inter-lamellar porosity 0.1-1 1-5% Vol Columnar grain width 0.2-1 -Thermal Microcracks 1-10 -Splat Thickness 1-3 -Pass Thickness 3-10 -Coarse porosity 1-30 0-20% Vol Unmelted Particles 5-50 0-20% Vol Splat Diameter 10-100 -Inter-splat contact - 24-33% Area This wide range of defects ultimately determines the mechanical performance of PS coatings. Any investigation into the cohesive properties or wear response of such coatings would have to consider the effects of defects of different types and sizes. Depending on the size of the surface damage or stress fields caused by the test method, different aspects of the coating microstructure will affect the material performance in that test. Due to the complexity of the wear processes in most materials, the modelling of wear mechanisms is a controversial area. A large number of wear models and equations have been published over the years [7], dealing with large numbers of variables and a variety of contact types and applications. Theoretical models dealing with the wear mechanisms that occur even in well characterized metals are still not fully evolved. In the area of ceramic materials, current wear models have failed to correlate with results [8], due to the use of macro-mechanical material properties to describe/study the micro-mechanical wear mechanisms that occur in engineering materials, i.e. in non-ideal, brittle materials like ceramics. The microstructure of 4 ceramic PS coatings are even more complex than those of bulk ceramics, and therefore, further study into the effects of the different microstructural defects is essential. 1.2 Objectives and Scope of the Thesis Although, in recent years, a number of studies concerning the many complex variables relating the deposition process and the coating microstructure have been performed, the microstructural integrity of PS coatings and its relation to their wear behaviour is still only poorly understood. A thorough investigation of the damage mechanisms that occur as a result of different contact forces can help determine which microstructural defects influence the material response to different types of contact forces. There is, therefore, also a need for methods that specifically evaluate the different types and sizes of microstructural defects. The aim of this study is to provide a scientific base for understanding the effects of different microstructural defects on the performance of PS coatings under different mechanical contact situations and wear modes. For this purpose, a series of ceramic coatings, both alumina- and chromia-based, was sprayed according to a matrix of deposition parameters and then extensively characterized for the types and sizes of microstructural defects occurring in each coating. A series of tests chosen to provide a range of contact disturbances was then performed on the coatings, as well as on a fine-grained sintered alumina as a reference material. The various tests and contact conditions included in this work are listed in Table 1.3. A thorough investigation of the damage mechanisms that occurred was then performed in order to evaluate which tests were the most useful for evaluating different types and sizes of defects, and the cohesive integrity of the coatings. Descriptive models of the mechanisms that occur during different micromechanical contact stress situations involving ceramic plasma 5 sprayed coatings are also evolved. The correlations between microstructural parameters, micromechanical properties and the wear performance are presented and discussed. Table 1.3 Techniques for Applying Different Types of Contact Force to the Coating. Technique Applied Force Size of Affected Zone Type of Disturbance Nanoindentation 0.01-0.1N 1-5 pm Hydrostatic Microindentation 0.1-10.0 N 3-30 pm Hydrostatic Macroindentation >10N 30-100 pm Hydrostatic Scratch Testing 1-200N 25-200 pm Normal + Shear Erosion: 90 0 v=70m/s, 75-600pm particles >100 mm 2 (Individual impact -10-500 pm) Normal Impact Erosion: < 90 ° v=70m/s, 75-600pm particles >100 mm 2 (Individual impact -10-500 pm) Normal Impact + Shear Abrasion 10-15 N , 0.2-0.3m/s, 20-200 pm particles >100 mm 2 (Individual scratch -50-300 pm) Normal + Shear To summarize, the objectives of this thesis are as follows: • To evaluate, both qualitatively and quantitatively, the effect of the microstructural features of TS coatings, such as porosity, microcrack density and splat morphology, on micromechanical properties and wear performance. • To characterize the damage micromechanisms of the coatings and present descriptive models of the mechanisms involved in "wear by hard particles". • To evaluate the different methods of examining the microstructural defects and cohesive integrity of ceramic PS coatings. 6 CHAPTER 2 PLASMA SPRAYED COATINGS 2.1 Review of Basic Deposition Process Thermal spraying is the generic name for a family of material deposition techniques involving the heating of a coating material to a molten or plastic state and subsequently propelling the particles at high velocities to impact the substrate material where they solidify and build up to form a protective coating. It is one of the most versatile techniques to produce coatings for improved wear, thermal and corrosion resistance, at a relatively rapid rate and with practically no limits on the size of the coated area. There are two basic methods of thermal spraying based on the way the coating material is heated and propelled towards the target. The first and oldest method is called the combustion method and depends on the combustion of a hydrocarbon fuel to provide the required energy. It was originally developed by Schoop in Switzerland who designed a compressed air system for melting and spraying tin or zinc in 1910 [1]. The second process relies on electric power, wire arc or thermal plasma, for the energy necessary for spraying. Details of the five most common variations of the two methods are compared in Table 2.1 below. 7 Table 2.1 Comparison of the Basic Thermal Spray Processes [1,9] Method Heat Source Plasma Temperature CQ Particle Impact Velocity (m/sec) Max. Spray Rate (kg/hr) Materials Deposited Wire Flame Spray Oxyfuel 2800 200 10 Metals Powder Flame Spray Oxyfuel 2800 30 7 Metals, ceramics, plastics High-Velocity Oxyfuel (HVOF) Oxyfuel 3100 600-1100 14 Metals, carbides Wire Arc Electric Arc 6000 240 !6 Metals Plasma Spray (PS) Electric Arc 6000-17000 240-1200 4.5- 23 Metals, ceramics, plastics, composites 2.1.1 Plasma Spraying Plasma spraying (PS) is one of the most versatile and sophisticated of the thermal spray methods. The gases forming the plasma serve as both the heat source and the propelling agent for the coating material, as shown in Fig. 2.1. The diameter of the spray nozzle, which is also the anode, constricts the electric arc producing high current densities (i.e. 1000 A at 80V) and temperatures [1]. The plasma jet is produced as the gases pass through the arc and are dissociated and ionized. The plasma is ejected from the nozzle at extremely high velocities (up to 1200 m/sec). 8 Powder Fig. 2.1 Schematic of plasma spray torch with (a) radial and (b) axial powder injection system. Monoatomic gases attain higher temperatures through ionization when more current is passed through the plasma, while polyatomic gases release more energy through the process of dissociation in addition to ionization, as shown in Fig. 2.2. The combined ionization and dissociation energy, i.e. the heat content of the gas, provides the necessary energy to the particles for melting. In practice, pure nitrogen or argon is typically used as the primary plasma gas together with 5-25% of a secondary gas (hydrogen or helium). The use of a secondary gas raises the ionization potential and thermal conductivity of the plasma, allowing higher temperatures to be produced at lower power levels [9]. 9 0 0 4 8 12 16 20 24 Temperature (x IO5 K) Fig. 2.2 Graph of temperature vs. heat content for typical plasma gases [10]. Due to the high temperatures attained almost any type of material can be sprayed, provided it can be made in powder form and will not decompose in the plasma. This includes metals, ceramics, carbides, cermets and plastics. The process can be done in air (APS) or vacuum (VPS), and the last few years have seen an increase in the number of speciality processes. The powder is typically injected into the side of the plasma where it melts and is propelled towards the target, see Fig. 2.1 (a). Usually, during spraying with radial-feed torches, powder particles experience variable processing conditions because of both their large size distribution and non-homogeneous exposure to the plasma. Particles injected normal to the plasma stream are prone to segregation (classification). Thus, while optimum sized particles are entrained, melted and deposited, smaller ones tend to be either deflected by the jet or overheated and partly reduced in the plasma. Larger particles tend to go through the jet or, if entrained, may be only partly melted. A l l of these contribute to low deposition efficiency at best and coatings with poor microstructure, and hence poor properties, at worst. In recent years, new designs of APS have been developed that utilize an axial injection torch, including the Northwest Mettech Axial III™ System [11]. In this system, the powder particles are injected into the center of the plasma stream, as shown in Fig. 2.1 (b), providing 10 better deposition efficiency along with improved uniformity of particle melting, entrainment and acceleration to the target. The high particle velocities achieved using this system have been shown to be comparable to those seen in H V O F systems [12]. The majority of the coatings investigated in this thesis were sprayed using this type of PS system. One set of coatings was sprayed using a radial-injection PS torch, for comparison. 2.1.2 Deposition Parameters In addition to the above mentioned system variables, such as power, gas composition and flow rates, parameters that affect the quality of the deposited coatings include the size and shape of the powder particles. The morphology of the particles affects the flowability of the powder, and thus the powder feed rate. The size of the particles affects the particle residence time in the plasma and the temperature that is reached [13]. In the case of ceramics, conventional powders are typically made by sieving fractions from fused and crushed material that are angular in shape. Methods of producing spherical powders that have recently been developed include sintering and agglomeration techniques. Commercial powders are available with wide or narrow size ranges and particles from 5-50 pm for oxides and 50-100 pm for metal powders. The residence time of the particles in the high temperature zone of the plasma jet is a key factor in achieving complete melting of the particles. This is determined by the mass of the particle since the drag coefficient is mainly a function of the gas properties. The residence time 110 will be proportional to p " , where p is density, for a given particle size [14]. It is also affected by the size of the injection nozzle, as this affects the flow rate of the plasma gas jet. The preparation of the surface to be sprayed is also important, especially for high bond strength. The substrate surface is typically cleaned and 'activated', i.e. energy is imparted to it, by grit-blasting with hard particles, typically aluminum oxide. This increases the surface area and the number of surface asperities, i.e. higher points on a rough surface. These are important in the bonding of splats to the substrate through mechanical interlocking with the asperities [4]. 11 2.2 Microstructure of Ceramic Plasma Sprayed Coatings 2.2.1 Introduction The microstructure of ceramic plasma sprayed (PS) coatings, shown in Fig. 1.1, even though made up of ceramic particles, is unlike that of any monolithic ceramic, in that is is built up of layers of rapidly solidified droplets or splats. Deviations from ideal behaviour occur in PS ceramic coatings because of the incorporation of partly melted particles, splashing of the droplets on impact, and defects arising from faults in the spraying process such as fluctuations in powder feed rate. Local temperature variations during spraying and differences in thermal expansion coefficients also result in microscopic and sometimes macroscopic cracking. The inter-splat bonding, or cohesion, within this lamellar structure is a key factor affecting the performance of these coatings. Each of the above factors, or microstructural features, affect the coating response to various forms of contact. In order to study these effects it is first necessary to fully understand these parameters. In the last 10 years the importance of these microstructural parameters as the link between the process conditions and coating quality has begun to be appreciated and studied. However, much work remains to be done. 2.2.2 Lamellar Structure The microstructure of the ideal plasma-sprayed coating is related to the coating formation process through two predominant factors: 1) the nature of the interface between randomly stacked pancake-shaped splats and 2) the internal structure within these splats, produced by their rapid solidification. The structure has been described as "lamellar" since it is built up of these flattened splats that form as the molten particles impinge on the substrate. Succeeding particles flow on top of those already solidified, forming an anisotropic lamellar structure parallel to the substrate [1]. Splat sizes for ceramic coatings typically vary from 1-3 pm in thickness, and from 50-100 pm in diameter depending on the size of the precursor powder, as shown in Fig. 1.1. 12 The lamellar structure has been studied using a copper electroplating technique on alumina PS coatings [15, 16]. This made it possible to examine the amount of interlamellar contact in the coatings, and clearly showed the network of horizontal (i.e. delaminations between the splats) and vertical cracks. The extent of bonding between splats was calculated using estimates of the actual contact area divided by the mean splat length. A maximum of 32% was found in one study [15]. A schematic of the structure of a typical splat in an alumina coating, as theorized based on the results obtained using the copper electroplating technique, is shown in Fig. 2.3 below. This agrees with other models put forth [14]. This model helps explain the reduced values of physical properties measured for PS coatings. Bonded Area Fig. 2.3 Schematic of idealized structure of flattened ceramic splats, or lamellae [after 15]. The amount of interlamellar contact is the major factor affecting the mechanical properties of PS coatings. The low bonding rate observed in PS coatings is a function of the short contact time of the impinging splats prior to solidification and the relatively low temperature of the already deposited splats. The temperature of the particles in the plasma is thus an important factor affecting the extent of bonding [15]. The interlamellar, or intersplat, contact or bonding has been described in many different ways. As it is a major factor affecting the porosity in a coating and consists of voids, it has been called interlamellar porosity. When discussing the 13 crack distribution in a coating it is also referred to as horizontal delaminations or horizontal cracking. 2.2.3 Phase Composition and Structure The crystal size and morphology within lamellae depend upon their relative nucleation and growth rates as a function of temperature. In general, high cooling rates (up to 106 K/sec) are caused by the large difference in temperature between the incoming, partially molten, small particles and the cooler metallic substrate with its much higher thermal conductivity. This gives rise to high nucleation rates, and polycrystalline lamellae with crystal sizes much smaller than conventionally processed materials (30-300 nm) [5, 17]. Rapid nucleation occurs at the cooler surface of the flattened droplet at large undercooling and crystals grow rapidly to form what has been called a columnar grain structure [6,14], Fig. 2.4. These columnar-looking cleavage steps indicate that crystal growth has taken place primarily in a direction perpendicular to the plane of the splats and the substrate. Fig. 2.4 SEM photo of PS alumina coating fracture surface showing columnar grain structure within splats (this work). The high cooling rates involved as the lamellae solidify (106 K/s) are also reponsible for the formation of metastable or nonstoichiometric phases in the coating. For example, pure zirconia undergoes a variety of phase transformations when cooling, from a cubic crystal structure at its melting point around 2700 °C, transforming to the tetragonal t-phase at about 14 2360°C and finally undergoing a martensitic transformation to the monoclinic phase at around 1200°C. The addition of various cubic oxides such as MgO, Y2O3 or CeC>2 allow the stabilization of the cubic or tetragonal phase at room temperature. Partially-stabilized zirconia (PSZ) has a two-phase structure consisting of cubic grains with tetragonal and/or monoclinic precipitates, while tetragonal zirconia polycrystal (TZP) is made by the addition of 6-8% Y2O3 or CeC>2 to stabilize the tetragonal phase. TZP is typically used for thermal barrier coatings as the tetragonal phase will not transform to monoclinic under stress [5]. Another well-known example is plasma sprayed aluminum oxide, where the metastable y-phase is preferentially formed at high cooling rates, with substantial droplet undercooling, since the energy barrier to homogenous nucleation is less than for the a-phase. Thus the percentage of the y-phase present increases as the energy of the spraying process increases and the particle size of the feed powder decreases. The presence of the a-phase in the as-sprayed coating is attributed to the presence of unmelted or partially melted particles. As the feed powder is typically a-phase alumina, if any part of a particle is not totally molten, that part will form the nucleus for the growth of an a-phase particle in the deposited coating [18]. 2.2.4 Porosity Porosity in ceramic PS coatings can vary from less than 2% to more than 20%, depending on the type of powder characteristics and deposition conditions used. The impact velocity of the particles has a large effect on the amount of porosity, along with the temperature of the particles. The porosity within PS coatings can be characterized as open, with pores connected to the surface, or closed, typically microporosity made up of spherical pores formed by entrapped gases. Most of the larger pores are due to the formation of cavities at the boundaries of overlapping splats, caused by incomplete filling in of the rapidly solidifying droplets. This effect is magnified the larger the particle size range of the powder being sprayed. The rapid 15 solidification that occurs when the melted particles impact the substrate induces stresses, and this can also lead to inter-splat porosity. The pore size distribution as measured by mercury porosimetry in some alumina coatings has been described as bimodal, with the fine pores around 0.01-0.4 pm and coarse pores ranging from 1-10 pm in size. The larger pores can be reduced or eliminated by increasing the temperature and energy of the plasma, producing only a network of monodispersed fine planar pores, roughly 0.1 pm in size. This microporosity has been interpreted as incomplete contact, due to nonwetting, between splats or lamellae. Transmission electron microscopy (TEM) of transverse sections of alumina coatings has revealed the presence of narrow planar pores (~ 0.1 pm) between lamellae [19]. 2.2.5 Unmelted Particles Unmelted particles are produced when the precurser powder particles do not melt completely in the thermal plasma, or do melt but subsequently resolidfy before impact. For a given size distribution of the starting powder, if the heat transferred to the particles in the plasma is not sufficient to produce a fully molten droplet, the resulting droplets will not completely flatten out and bond well with the underlying surface. The percentage of unmelted particles in the coating increases with increasing precurser particle size, as larger particles require more energy to melt completely. If the particles are only partially molten, they can still become part of the coating and negatively affect the microstructure and resulting properties [5]. 2.2.6 Impurities When spraying metals or cermets with a PS system in air, oxides can easily build up between the splats and affect the bonding. The use of a shroud of inert gas surrounding the part being sprayed effectively reduces the amount of oxidation, while spraying in low pressure (LPPS) or under vacumn (VPS) can eliminate it. This is not as much of a problem with oxide 16 ceramics. However, possible impurities include copper and tungsten from the torch electrodes, and sand from the grit blasting operation [17]. 2.2.7 Residual Stresses and Solidification Cracking Residual stresses in PS coatings have a large effect on their adhesion and operational performance [20]. There are two main sources of residual stress generation in plasma sprayed coatings. The majority of the residual stresses that build up come from the difference in thermal expansion coefficients (CTE's) between the substrate (15-18 x 10"6 m/m°C for metals) and the coating material (1-10 x 10"6 m/m°C for ceramics). These cooling stresses arise from the differential thermal contractions of the substrate and the coating that occur as the system cools from the final spray temperature down to room temperature [21], i.e. 6-c = £ c Ae , (2.1) and Ae=\T\as(T)-ac(T))dT (2.2) JT 2 where o c is the coating (or deposit) stress, Ec is the elastic modulus of the coating, Ae is the misfit strain caused by the cooling of the system from Ti to T2, a is the CTE for the substrate (s) and coating (c) and T is the temperature. As expected, coating stresses are reduced when the coating is less stiff, as stiffness is directly related to the elastic modulus. Thermal conductivity (K) is also important as it affects the actual rate of cooling. For example, an alumina coating has Kc of 25 W/mK(a) and thus will cool more slowly compared to a mild steel substrate with ^ s of 60 [20]. The temperature of the incoming particles is much higher than that of the substrate or the previously solidified droplets, thus the rate at which the coating is deposited will affect the heat input into the coating and the temperature profile of the system which ultimately affects the stress profile. The mismatch in CTE's is also the main cause of operational stresses, i.e. the stresses that occur when the coatings are being used in an operating environment. 17 The other major stress is known as the intrinsic, or quenching stress [22]. This is an in-plane tensile stress that is caused by the constraint or bonding of the impinging splat as it undergoes thermal contraction during solidification. It is relatively small in ceramics, less than 50 MPa, due to the stress relief obtained through microcracking perpendicular to the plane of the splat [5], as shown in Fig. 2.5. This is one of the stress relaxation mechanisms that can occur during the rapid cooling after solidification. Plastic yielding and creep take place in metals, while surface relaxation due to edge effects can occur to any type of coating material. Interfacial sliding can also occur due to the imperfect bonding between splats [20, 22]. Fig. 2.5 Top surface of an alumina PS coating, showing a network of fine cracks [this work]. The density of both vertical and horizontal microcracks in zirconia thermal barrier coatings has been studied to determine the connection to process parameters such as substrate temperature and particle temperature and velocity. In one study [6] the vertical microcrack density was found to increase with a decrease in substrate temperature, due to the high stress build-up, and subsequent relaxation, caused when the incoming splats cool to the low temperature of the underlying material, while spraying at a high substrate temperature reduced the density of both vertical and horizontal cracking. The reduced amount of horizontal cracking, or delaminations, was due to extended grain growth between splats. In another study [23], higher particle spraying temperatures were found to increase the amount of vertical cracking, while 18 increased particle velocities increased the horizontal cracking but not the vertical. In general, the greater the interlamellar bonding, the more constrained the splat surface and the greater the number of vertical microcracks that form to relieve the built-up stresses. Ways to reduce the stress buildup include preheating the substrate to minimize the temperature differences. The standard preheat temperature is 100-150 °C. Temperatures greater than 200 °C lead to increased oxidation of the metallic substrate which can then affect the adhesion. Selection of materials that have better matched CTE's and the use of bond coats between the substrate and the outer coating to gradually change the CTE's can also reduce the amount of residual stresses that build up. In applications such as thermal barrier coatings where a high strain tolerance is required, a high microcrack density is advantageous. 2.2.8 Adhesion to Substrate The strength of the bond between the coating and the substrate has a strong effect on its operational performance. Therefore, an understanding of this bonding is extremely important. Considerable debate has been ongoing as to the precise mechanisms involved in adhesion. It is now considered that in most cases it derives from a mixture of physical and chemical forces.The basic bonding mechanisms are grouped as physical, chemical and mechanical. Physical bonding is quite limited and is attributed to secondary and Van der Waal forces, while metallic and chemical bonding occur in regions of intimate contact between two compounds. In the case of metallic PS coatings, the oxide film present on the surface of the molten particles can prevent good metallic contact. N i A l coatings have very high adhesion to the substrate, along with good inter-splat cohesion due to in-flight deoxidation of the molten droplets by dissolved A l on their way towards the substrate [24]. The resulting splats are very dense and have an increased intersplat contact area. Ionic bonding occurs in regions of intimate contact such as occur in the columnar growth through freshly deposited splats in ionic ceramic coatings 19 like alumina and zirconia, see section 2.2.3 for further details. A chemical mechanism dominates when the fast moving near-molten particle interacts with the prepared substrate. Mechanical interlocking, where the impinging droplets become interlocked with surface asperities or deformities, has long been assumed to be the primary mechanism since most materials do not adhere well to polished surfaces [1]. This led to the standard procedure of grit blasting substrates in preparation for spraying. The process of grit blasting increases the surface energy of the substrate as well as breaking through the oxide layer on metallic substrates. Reduction of the cooling rate by preheating the substrate improves the bonding between lamellae and the substrate as well as between lamellae. The longer the temperature is maintained at elevated temperatures, the more inter-lamellar and coating-substrate bonding will occur. 2.2.9 Post-Deposition Sealing Treatments Many applications require sealing the pores of the PS coating before the final finishing operation, i.e. machining or grinding, to prevent the infiltration of corrosive fluids. The seal materials vary depending on the expected operating temperature [9]. Wax sealers are useful at low operating temperatures, with resin and silicone-based sealers useful in the ranges of 90-260 °C, and up to 480 °C, respectively. Epoxy and phenolic sealers are effective on coatings with higher porosity. There have been a number of reported attempts to seal the surface of plasma sprayed (PS) thermal barrier coatings by laser melting-solidification [25-27] in an effort to reduce the porosity and thus improve the protection of substrate corrosion. These attempts repeatedly resulted in a heavily cracked surface layer due to the solidification contraction of the rather porous zirconia, with its low thermal conductivity (1.1 W/mK) and high coefficient of thermal expansion (7.5-10.5 xlO" 6/K). In one report [28], preheating the substrate of an alumina plasma sprayed coating prior to laser-sealing prevented surface cracking. However, even if the "glazed" coatings could be obtained in uncracked form, additional drawbacks of the coating sealing through laser 20 melting-solidification include an uncontrolled resultant microstructure of the solidified layer, e.g. coarse-grained ceramic. The very high processing temperatures (the melting point for zirconia T m = 2400 °C) could adversely affect the metallic substrate. In addition, very dense coatings are no longer strain tolerant, as discussed in section 2.2.7 on solidification cracking. Post-deposition surface modification treatments that do not include additional melting of the ceramic, have also been investigated for ceramic plasma sprayed coatings [29-32]. The problems with solidification cracks, coarse microstructure and substrate alteration can then be avoided. The objective is to densify and seal the coating to substantially improve its integrity and wear resistance through controlled and selective deposition of active chemicals (for example, sol-gel processed ceramic) within the microcracks and pores which separate the individual splats within the ceramic coating. The active sub-micron grains of ceramic additive are subsequently sintered during dynamic or static heat treatment at temperatures well below the melting point of the coating. High intensity arc lamp or conventional furnace or the plasma torch itself, instead of a laser beam in a scanning mode, can be used to provide the required surface thermal treatments in a more uniform fashion and at much lower cost. Thus, the surface, to a depth of 50 to 200 pm depending on the porosity of the coating, of the as-deposited and additionally chemically modified porous ceramic coating can be sealed. 2.3 Coating Properties 2.3.1 Effect of Microstructure on Coating Properties Due to the variety of imperfections in the microstructure as described above, the mechanical and thermal properties of PS coatings are all inferior to those of a sintered material. It has also been found by many researchers that the wear resistance of plasma-sprayed ceramic coatings is significantly lower than that of the corresponding bulk ceramic. On the other hand, the defective microstructure of PS coatings results in increased strain tolerance and resistance to thermal shock compared to a very dense and defect free microstructure. A better understanding 21 of the effects of these defects will help in the prediction of coating properties. These will be discussed in the following sections. The physical properties of a coating can be directly related to its microstructure and porosity. The anisotropic lamellar structure, with limited contact between the lamellae and interfaces parallel to the load direction, together with the presence of pores, decreases the effective cross-sectional area of the coating. This causes a reduction in the physical and mechanical properties of the coating to considerably lower values as compared to similar monolithic materials [19, 33]. 2.3.2 Elastic Properties The modulus of elasticity (£) is the proportionality constant between elastic stress and elastic strain, and can be thought of as the amount of stress a required to produce a unit elastic strain s, i.e. E= c/e . The stronger the bonding within a material, the greater the stress required to pull it apart, or extend it. Typical polycrystalline ceramics have fairly high E values, on the order of 250-450 GPa [34]. However, due to their unique defect-filled, lamellar structure, ceramic PS coatings have measured E values in the range of 30-90 GPa [35, 36]. The model of the coating microstructure shown earlier in Fig. 2.3, has been used to study the effect of microstructure on the mechanical properties using flat circular plate theory [35]. An idealized model of the coating structure is shown in Fig. 2.6 for use in explaining its effect on the elastic modulus. When a load is applied to the coating perpendicular to the coating plane, the bonded regions between lamellae, with a radius of a, transfer the force from one splat layer to the next. In this model, it is assumed that the lamellae deflect elastically in the unbonded regions, while there is localized strain at the bonded regions. Based on this model, the deflection bending of the lamellae must be considered when the mean bonding ratio, a is less than 40 percent. When a is greater than or equal to 40 %, this term can be neglected and EJE is found to be equal to a. Thus the elastic modulus of a ceramic PS coating perpendicular to the plane of the coating can be 22 related to the localized elastic deformation of the bonded areas between splats and to the bending of the lamellae between the bonded areas. The extensive microcrack network typically seen within ceramic splats was not taken into account in this model, however the values predicted were quite close to those measured. The estimated value of Young's modulus for a typical alumina PS coating was roughly 10-30% that of a fully dense material, equivalent to the percentage of actual bonding contact, i.e. 30-90 GPa compared with 300 GPa for the equivalent sintered ceramic [35]. o o o Bonded Area O Interlamellar gap 3 C r n — Lamellar Interface Bonded Interface Fig. 2.6 Idealized model of coating structure: (a) planar and (b) cross-sectional views (Adapted from [35]). The lamellar structure of the coatings has a strong effect on their behaviour. The plate-like structure parallel to the coating plane, with the low level of actual inter-splat bonding and interlamellar porosity, can lead to nonlinear elastic behaviour and lower E values. The network of vertical cracks throughout the splats could account for the higher elastic modulus (89 GPa) found in the coating plane, while perpendicular to this plane, E was much lower (29 GPa ) [36]. The elastic and stress-strain behaviour of the coatings have both linear and nonlinear components, as shown in Fig. 2.7. Nonlinear elastic behaviour was observed in tension, while linear elastic behaviour was seen perpendicular to the coating plane. However the behaviour within the plane was found to be linear elastic in both tension and compression. 23 Knoop indentation can be used to measure the elastic modulus, as well as hardness, of a material, using methods based on the measurement of elastic recovery of the in-surface dimensions of indentations [37, 38]. The elastic recovery reduces the length of the shorter indentation diagonal (d), whereas the change in long diagonal length (D) is negligible. The following equation is used [37]: 0.45//. E = (2.3) D •0.1406 where E is the elastic modulus [GPa] and is the Knoop hardness [GPa] at, typically. ION. The value 0.1406 equals the ratio of the shorter to the longer identer axis. The elastic response of thermal spray deposits has recently been investigated using this method [39]. Values of E for alumina PS coatings were found to range from 50 to 175 GPa compared to 460 GPa for bulk alumina. 16 -14 -12 -(0 10 -Q. s in 8 -in o (A 6 -4 -2 -0 -Compression, / Tension, Perpendicul a r _ / ^ / Parallel ^ Tension, "" Perpendicular 0 50 100 150 200 250 300 350 400 450 Strain, (xlO" 6 ) Fig. 2.7 Stress-strain curves for PS alumina parallel and perpendicular to the coating plane [36]. Substantially higher values were found using nanoindentation, or depth sensing indentation (DSI) at low (50 mN) loads [40], e.g. 175 vs. 75 GPa. Low load indentation samples the "intrinsic" properties of the coating, or individual splats, as the volume of material sampled is smaller than the splat size. The ratio H/E, a measure of elastic response that provides a material 24 index of elastic-plastic behaviour, was also measured using both nano- and micro-indentation methods. H/E is small for rigid/plastic materials and becomes larger for more elastic materials. 2.3.3 Hardness Hardness testing is often used for preliminary assessment of coating quality due to its ease of use. The microhardness of ceramic PS coatings is strongly affected by the variety of defects found in the microstructure as well as by the size of the indentation load. The larger the load in relation to the size of the microstructural defects, the lower but more consistent the values measured become, as the increased test volume renders the readings less sensitive to variations in the microstructure. As the load and the resulting impression are reduced, the measured hardness of brittle materials increases. This is called the indentation size effect [41]. Smaller loads give hardness values with high scatter as well due to the effects of microstructural defects on the indentation impression. Porosity, unmelted particles as well as splat boundaries and the degree of inter-splat bonding all affect the measured microhardness, giving values for PS coatings that are 50% lower than are found for similar bulk materials, i.e. HV(3N) of 11 versus 20 GPa respectively for alumina [this work]. 2.3.4 Adhesion Strength and Fracture Properties The adhesion strength of ceramic PS coatings is commonly measured using a tensile adhesion test (TAT), such as A S T M C 633-79. In this test the coating on one end of a 25 mm diameter steel stub is attached to a similar stub by a high strength adhesive, typically an epoxy resin. The combined specimen is mounted in a uniaxial testing machine and loaded in tension to failure normal to the coating surface. There are four ways that fracture may occur: (1) interfacial or adhesive failure, occurring along the coating/substrate interface, (2) cohesive failure within the coating itself, (3) through the adhesive layer alone and (4) mixed-mode failure, a combination of the other three types [36, 42]. The tensile strength of the bond is determined by dividing the failure load by the cross-sectional area of the steel stub, provided the failure occurs 25 at the interface. If the sample fails in the adhesive then it is only known that the bond strength is higher than the measured stress. If cohesive failure occurs it can only be assumed that the tensile adhesive strength is greater than the cohesive strength. The cohesive strength can be related to the defects in the microstructure through the use of the Griffith equation for fracture strength of a brittle material. In this equation, 'lyE yTTC j (2.4) Of is the fracture strength, E is Young's modulus, y is the fracture surface energy and C is the length of the critical defect, i.e that leading to fracture [43]. Thus the strength is related to microstructural effects such as interlamellar porosity that have a pronounced effect on the elastic modulus, as discussed earlier, and also the fracture surface energy. The largest pore in the coating can often be the critical, fracture controlling defect, and as such strongly influences the strength and fracture toughness of the coating. The high density of microcracks substantially decreases the coating stiffness even if the porosity is low. Effectively, the stiffness of a PS coating is only 10 to 50% of that of the respective monolithic material. In contrast, the fracture toughness of the ceramic PS coatings is typically higher than that of sintered ceramics, because of the lamellar grain structure and the presence of microcracks induced by local shrinkage stresses. This microcrack toughening effect, hindering crack propagation, is not unlike that found in ceramic composites [33, 44]. The fracture toughness, Kc, and critical strain energy release rate, Gc, of ceramic PS coatings have been measured using indentation crack patterns. In one study [45] measurements were made on polished cross-sections of the coatings, and it was found that accurate values could not be obtained due to the various microstructural parameters, i.e. porosity, defects and residual stresses, found in the coatings. The method was deemed useful in providing a relative measure of Gc however. In a more recent study [46], indentations were made on polished planar surfaces of 26 ceramic PS coatings, and the KC values obtained were higher than those for the comparable monolithic ceramic. The spinel coating tested gave a value for KC of from 1.9-3.4 MPaVm as compared to 1.2-1.9 MPaVm for the bulk spinel. A unique method of determining the indentation fracture toughness (IFT) of cutting tool materials [47] has been used to study the properties of thermally sprayed WC-Co coatings [48, 49]. The relationship: , F T = 0 - 1 1 3 1 ^ ( 2 , 5 ) H I was used, where H is the Knoop microhardness [GPa], D is the diagonal of the Vickers indentation [pm] and CL is the total crack length, defined as Q = i X (2-6) where C „ is the length of the cracks emanating from each of the four corners of the Vickers indentations. This equation is based on the formation of Palmqvist rather than lateral cracks [48]. In this study Knoop indentations were made on polished cross-sections both perpendicular and parallel to the coating substrate. The resultant IFT values demonstrated the anisotropy of the thermal sprayed coatings, with larger values measured when the Knoop indentation on the cross-section was made perpendicular rather than parallel to the coating surface. The low toughness in the parallel direction was considered to be a direct consequence of the weak bonding between splats. 2.3.5 Thermal Properties The thermal conductivity and diffusivity of PS coatings are about one fourth the value of similar bulk materials [36], i.e. 0.45- 0.55 W/mK for CaO stabilized zirconia PS coatings tested in vacuum as compared to 1.8- 2.2 W/mK for the sintered form. This property, along with a 27 coefficient of thermal expansion nearly equal to that of a metal, is the reason that zirconia PS coatings are typically used as thermal barriers, to reduce the temperature seen by the metallic substrate. The reason for these low values is again due to the limited amount of actual contact between lamellae, since heat conduction in a vacuum takes place across areas of solid contact. In the presence of a gas, the conductivity is affected by the size of the pores in the coating. The very small width of some of the (interlamellar) pores in these coatings ( -0.1 pm) is comparable in size to the mean free path of the gas molecules and this reduces the amount of conduction, refer to Fig. 2.6. A microstructural model of the poor interlamellar contact in the coatings has confirmed the low values of thermal conductivity [50]. 2.3.6 Wear Resistance The microstructure of plasma sprayed coatings strongly affects their wear performance. The complexity of the lamellar microstructure is well known, as are the importance of specific details of the structure. However, the impact of these complexities on the material and system properties, such as wear, and the details of the various interactions have been studied very little, especially when compared to metals and ceramics. The wear behaviour of metals has been extensively studied and plastic deformation is the major wear mechanism in the majority of metal wear modes. Ceramics experience brittle fracture to a much larger extent although plastic deformation does occur, due to the high temperatures and pressures that occur within the contact zone [51], as well as the inhibition of fracture due to the hydrostatic pressure caused by the indentation of abrasive asperities [52]. Macroscopically, the "quasi-plasticity" seen in ceramics resembles the conventional plasticity seen in metals, as it is shear-driven [53, 54]. However, microscopically the source of the deformation is dissipative slip at microstructurally discrete "shear faults", e.g. weak particle-matrix interfaces, augmented by internal residual stresses rather than by dislocation slip. 28 The microstructure of bulk ceramics, i.e. grain size, phase distribution, etc., has a significant effect on the wear [55, 56]. Abrasive wear tests on different oxide ceramics have shown a type of Hall-Petch relation (D' ) between the wear resistance and the average grain size D [57, 58]. The volumetric wear intensity of alumina was found to increase linearly with increasing average grain size [59], while in another study the abrasive wear resistance was found to increase exponentially with increasing porosity [60]. Some results from the comparatively small number of articles relating to the wear of ceramic and cermet thermal spray coatings will be summarized in the remainder of this section. 2.3.6.1 Abrasive Wear A number of studies have been done on the abrasion resistance of thermal sprayed (TS) coatings. The main parameters influencing the abrasive wear of TS coatings include: - microstructural variables such as pore size and volume %, phase and impurity distribution, and intra- and inter-splat cohesive strength - size, shape and hardness of abrading particles - material hardness and fracture toughness - surface loading Abrasive wear testing performed by Barbezat et al [61] showed that at moderate loads oxide ceramic coatings with a more homogeneous microstructure exhibited better abrasion resistance than those with a less homogeneous microstructure. In this case Cr 2 0 3 and AI2O3 coatings performed better than Al 2 O 3 -40% Z r 0 2 and Cr 2 0 3 - ( 10-30%) T i 0 2 . At higher loads however, the Al2 O3-40% Z r 0 2 showed better wear resistance than the A12C>3 coatings, due to the increased fracture toughness caused by the addition of 40 wt.% Zr0 2 . Toughness and cohesion between the individual particles, or splats, of the coating were found to be extremely important for good abrasive wear resistance. The wear mechanisms involved in the abrasion process in tungsten carbide/cobalt are plastic deformation and fracture, characterized by selective removal of the binder, then cracking 29 and removal of carbide grains. The microstructure of WC/Co, including porosity, carbide grain diameter and mean free path within the binder, controls the fracture behavior and thus the wear behaviour of both sintered and plasma sprayed materials [49, 62]. The lamellar microstructures created by plasma spraying lead to anisotropic fracture toughness and variations in the wear resistance compared to the response of sintered WC/Co. Attempts to correlate the abrasion wear resistance to a structure/property factor revealed large scatter [49], due to problems with the variations in measured microhardness and fracture toughness caused by microstructural differences, such as W C grain size, porosity, carbon content and the mean free path of the cobalt, and anisotropy. The best abrasive wear performance for both sintered and plasma coated WC/Co was obtained using a fine dispersion of carbides in a cobalt binder matrix [62]. 2.3.6.2 Erosive Wear Plasma spray coatings are used in many applications requiring good erosion resistance, such as blades and nozzles in gas and hydraulic turbines as well as exhaust fan blades. In solid particle erosion tests on Zr02 and WC-Co based coatings [63], it was found that erosion rates were higher than similar bulk materials due to the presence of voids and lamellae in the coatings that had a tendency to crack and break apart. The WC-Co coatings were superior to the ZrCV based coatings, which had a higher porosity. As is typical of ceramic materials, wear losses were greater at an impact angle of 90° rather than 30°. The main wear mechanism in the erosion process is brittle fracture. The erosion resistance of a coating can be increased through microstructural optimization, i.e. the balancing of finer lamellae, with lower porosity and higher toughness. For carbides, it was found that factors such as coating porosity, intersplat strength and carbide/matrix bond strength may dominate over cobalt content or W C grain size [43]. A comparison of the erosive wear behaviour of vacuum PS alumina coatings and bulk alumina [64] found that the main wear mechanisms for vacuum PS coatings, which had good intersplat bonding, was brittle fracture, especially at high angles of impingement. As the angle 30 decreased, the mechanism changed more and more to that of ductile plowing. The erosion resistance was also found to be comparable to the bulk alumina when tested at high angles. Coatings with slight compressive stresses showed better wear resistance than those with tensile stresses. Porosity has a strong effect on the erosion resistance of PS coatings as well. Studies performed on porous zirconia PS coatings [65] found that for small sizes of eroding; particles, 27 pm in this case, the strength of the microstructure and the distribution of the porosity between discrete pores and intersplat boundaries at a given porosity level were useful in predicting the erosion rate. The erosion rate was found to be higher when the porosity was mainly at intersplat boundaries. In another study [62], it was found that an even distribution of round pores throughout the plasma spray coatings decreased their erosion resistance more than the long, flat pores found between splats in thermal spray coatings. This illustrates the need for a better understanding of the relationships between microstructure and wear performance. Throughout the studies done on the wear of ceramic TS coatings, the importance of inter-splat bonding, or contact, stands out. In a series of studies [66, 67], it was found that the particle erosion rate of ceramic coatings appeared to be controlled by bonding at the interfaces between splats. By measuring the erosion rate at 90 degrees, which occurs by a debonding, or interfacial separation, mechanism, they were able to obtain a statistical measure of coating cohesion. This will be discussed further in the next chapter. 2.3.6.3 Sliding Wear Ceramics typically exhibit low thermal conductivities, generating high local temperatures that lead to high thermal stresses at the points of contact during sliding. This results in spalling and the formation of abrasive grains. In ceramic PS coatings, the frictional and thermal stresses are centered at any pores and microcracks on the coating surface, leading to spalling of the coating at those sites. In one steel block-on-coated roll study of a number of oxide ceramic PS 31 coatings [68], a correlation was found between thermal diffusivity and wear resistance. The best coating tested was Cr203, which had the highest thermal diffusivity along with the lowest percent porosity. The effect of water on the wear of ceramic PS coatings is similar to that on monolithic ceramics. In a study of several oxide PS coatings versus steel in both oil and water, the wear rate increased by a factor of 10 when run in water [69]. This was due to the reaction of the water with ceramics to produce soluble tribochemical films. Plasma sprayed alumina coatings are typically composed of the metastable Y-AI2O3 phase which converts easily in water to aluminum hydroxide (bayerite). The tribochemical surface films produced in oil are typically more friction and wear reducing. More research has been carried out on the sliding wear of PS coatings than any other type of wear, due to the interest in sliding wear applications using them. In a study by Furukubo, et al [70] the results suggested that the wear properties of PS coatings were closely related to the cohesion between the individual splats. This was determined by using image analysis to compare the area fraction of the regions showing debonding at the intersplat boundaries with that of brittle microchipping in a series of alumina plasma sprayed coatings after undergoing sliding wear. 32 CHAPTER 3 METHODS FOR EVALUATION OF THE MICROSTRUCTURAL INTEGRITY OF CERAMIC PLASMA SPRAYED COATINGS As discussed in the Scope of the Thesis, Section 1.2, in order to evaluate and separate out the influence of the individual microstructural parameters, PS coatings need to be studied under a variety of contact situations. In this chapter a number of techniques are reviewed and evaluated as to their suitability for testing the microstructural integrity of ceramic plasma sprayed coatings. Many of these techniques have been used for the evaluation of monolithic ceramics as well as other brittle materials. Wear testing methods involving hard particles are ideal for this purpose as the various wear mechanisms involved cover a wide range of contact types. 3.1 Damage by Hard Particles One of the major uses of PS coatings is in applications involving wear. The greatest extent of wear loss in brittle materials like ceramics is due to damage by hard particles. The processes of both abrasion and erosion, see Fig. 3.1, involve damage, or wear, by hard particles. Abrasive wear can be defined as wear due to the penetration and plowing out of material from a surface by another body [55]. Two-body abrasion occurs if the abrading material is either a surface protuberance on the mating part or abrasive particles moving freely over the component surface, as in sand sliding over a chute. Three-body abrasion occurs when abrasive particles from either an external source or generated wear particles, are loaded between two surfaces. Erosive wear is caused by the impingement of hard particles against a surface. If the impacting particles are solid, it is typically qualified as either solid particle erosion or solid impingement erosion. If the particles are carried by a liquid it is called slurry erosion [71]. 33 / / / / / / / / / / / / / (a) (b) (c) Fig. 3.1 Schematic of: (a) two-body and (b) 3-body abrasion and (c) erosion processes. (Adapted from [71].) Ceramic plama sprayed (PS) coatings can be thought of as ceramics with a high density of defects in the microstructure. As discussed earlier, they are unique in that the microstructure is made of solidified droplets, or lamellae, along with interlamellar porosity and unmelted particles. However, they are ceramics and thus exhibit brittle behaviour. The response of these coatings to wear by hard particles is strongly affected by the size and type of microstructural defects in the coatings as well as the abrasive particles. The hardness of the abrading particles compared to the coating hardness plays an important role in the wear rate of the surface being abraded [71]. Hard abrasion is defined as occurring when the ratio of the abrasive hardness Ha to the surface hardness Hs is greater than 1.2. This is the case with sliding diamond indenters. Soft abrasion occurs when HJHS < 1.2. Deformation of the abrasive particles is more likely to occur in this case, substantially reducing the wear rate. As HJHS increases, the major wear mechanism caused by coarse abrasives changes from plastic deformation to fracture, and the severity of the wear increases accordingly. The size and shape of the particles are also important factors in the type of wear that occurs. Angular particles cause more wear than rounded ones, due to the increased contact stress caused by sharp indenters. The larger the size of the particle, the greater the area of damage to the wearing surface. This can be used to study the coating material response under different contact conditions. 34 3.2 Indentation Methods 3.2.1 Introduction Indentation testing provides a means of studying the response of a material to a single abrasive particle, typically angular, pressing into the surface. In this process, an elastic-plastic field is formed underneath the indenter, see Fig. 3.2, with both elastic (reversible) and inelastic (plastic) components of deformation [72]. Even brittle materials like ceramics exhibit plastic deformation due to the inhibition of fracture by the localized high hydrostatic pressure occurring directly underneath a hardness indenter [52]. Fig. 3.2 Model for elastic-plastic indentation. (Adapted from [72].) Indentation methods have been used extensively to study the mechanisms involved in abrasive and erosive wear processes, i.e. wear by hard particles. Damage in brittle solids [73, 74] includes both the elastic regime, for low loads and small particle sizes, elastic-plastic indentation, when the radius of the particle (or indenter) is smaller than the critical radius for plastic penetration, all the way to brittle fracture, when the applied stress exceeds the critical surface pressure for microcrack initiation. The shape of the indenter affects the resultant damage area (or induced stress field). Contact between two bodies, such as a spherical particle or indenter and a surface, is considered fully elastic when the contact pressure is below that required to initiate either plastic 35 deformation or brittle fracture. This area was analyzed first by Hertz [75]. The progression of Hertzian cracking, or fracture, is shown schematically in Fig. 3.3. When a spherical, or blunt, indenter or particle first contacts the surface, compressive stresses are distributed over the entire contact area. Tensile stresses arise due to elastic deformation of the material around the point (or ring) of contact (Fig. 3.3(a)). As the load and contact area increase, the tensile stress field increases and expands radially until a weak point, or flaw, is reached. A surface "ring" crack is formed at the maximum tensile field around the contact area and extending downwards into the material (Fig. 3.3(b)). The compressive stresses have also increased and forced the crack outwards at an angle (Fig. 3.3(c)), until the development of a Hertzian cone crack occurs. Ring Crack (d) (e) (f) Fig. 3.3 Evolution of Hertzian cone cracks during spherical indenter loading / unloading. (Adapted from [75].) The transition from elastic to elastic-plastic indentation occurs when the radius of the particle (or indenter) is smaller than the critical radius for plastic penetration at a fixed load. In ceramics this radius can be determined by observing at which indentation radius the transition from circumferential to radial fracture occurs for each material. The plastic indentation load, P,, is more dependent on the material hardness (Pi ~ H3) while the Hertzian fracture load, PF, is dependent on the fracture toughness of a material as well as the size of the surface flaws [73]. 36 During so-called "sharp" indentation, i.e. Vickers indentation, the contact is primarily plastic in nature until fracture occurs and the progression of the crack systems including the median (M) and lateral (L) vent cracks is shown schematically in Fig. 3.4. The plastic, or deformed, zone underneath the indenter is indicated by D in the diagram. The radius b, see Fig. 3.2, of the plastic zone has been shown to be related to the load P, the elastic modulus E, and the hardness H, as follows [76]: b ~ (P/H)1/2(E/K)2/5 (3.1) A Blunt' indenter can also act as a 'sharp'indenter, if overloaded to exceed the plastic indentation load, and cause median and lateral cracking [75]. (a) Increasing Load (b) (C) ? M (d) (e) M (f) Decreasing Load Fig. 3.4 Progression of the crack systems during sharp indenter loading and unloading. (Adapted from [75].) Indentation techniques are widely used to investigate the mechanical properties of brittle materials, such as hardness and fracture toughness. As these properties are affected by the various microstructural defects present in the materials, the response of the materials to these defects is also being probed by indentation. 37 Two types of indenters are most commonly used to measure the properties of coated materials, the pyramidal Vickers indenter and the more elongated Knoop indenter, shown in Fig. 3.5 [41]. The definition of hardness for the two are slightly different. For the Vickers indenter, hardness is equal to the load divided by the surface area of the impression, while for the Knoop indenter, the hardness is equal to the load divided by the projected area of the impression. The advantage of the Knoop indenter is its smaller penetration zone, l/30 t h of the long diagonal length compared to l /7 t h for the Vickers indenter. This makes it quite useful for testing very thin coatings or layers, however this also makes it very sensitive to material anisotropy, i.e. it is often difficult to determine the real length of the long diagonal (D/d = 7.11). The size of the indentation impression is a function of the contact load used, as shown in Table 1.2. VICKERS KNOOP Fig. 3.5 Comparison of Vickers and Knoop indenters for hardness measurement. (Adapted from [41].The load, P, is in kg and diagonal, D, is in mm.) 3.2.2 Depth Sensing Indentation Microhardness testing and depth sensing indentation (DSI) are two valuable tools for determining the mechanical and micromechanical properties of coatings. DSI, also known as nanoindentation, is a technique that solves the problem of trying to measure very small indents, 38 i.e. < 10 pm, optically to determine material hardness at low loads (0.01-0.1 N) through the use of continuous depth recording to measure the penetration depth of the indenter, together with load and time [77]. Figure 3.6 shows typical loading-unloading curves for elastic, plastic and elastic-plastic materials. The shape of the curve can be used as a "fingerprint" for different materials, relating the microstructure to the micromechanical properties at the surface. This can be useful especially with the present development of advanced surface engineering techniques, and for understanding the processes relevant to surface damage and wear. DSI is used extensively on thin coatings (< 10 pm), where the hardness of the coating can be separated from that of the substrate. Typical indenters used include the standard Vickers and Knoop (4-sided) and Berkovitch (3-sided) pyramids, along with a hemispherical ball indenter. An increase in hardness is typically observed at low loads and indentation depths due to a combination of a volume effect related to the yielding of the material, i.e. the smaller the volume of material, the higher the material flow (or yield) stress [78], and the effect of any imperfections at the tip of the indenter which alter the geometry. These factors comprise what is kown as the indentation size effect (ISE). The ISE can be corrected for by establishing correction factors for the indenter using a material that exhibits depth-independent constant hardness and Young's modulus values, such as monocrystalline silicon [79]. Low load indentation, due to the small volume of material, samples the "intrinsic" properties of the coating , i.e. the individual splats. Information that can be obtained directly from the load-depth curves, Fig. 3.6, includes elastic recovery, relative hardness and the work of indentation. Each of these is a measure of a material's resistance to penetration. Values of hardness cannot be measured directly from the plot of load vs. penetration depth, however the data can be processed to give hardness values [79]. 39 (a) (b) g a. ELASTIC (L) / / : u ) PLASTIC (L) / (U) 1 r (c) ELASTIC-PLASTIC Dtplb,h(jii| Fig. 3.6 Typical Loading (L)-unloading (U) curves: (a) elastic, (b) plastic, and (c) elastic-plastic materials, with the plastic depth (hp). The DSI method gives two measures of hardness, namely hardness-under-load values (HVL) and values calculated from the work of indentation (Hn.) [80], Fig. 3.6(c). The indentation result trends are identical for both measures of hardness. HVL gives slightly smaller values reflecting the fact that it is calculated from the combined plastic and elastic deformation depth during indentation, whereas Hw is derived from the plastic deformation depth only, / i [pm], (as is HV in microindentation) and is the method typically used in studies involving ceramics. Assuming an ideal geometry for the indenter, the ratio of the diagonal to the plastic depth, hp, is 7 for the Vickers indenter. This gives the relationship for hardness as used in this work: P H„,= 37.84-hl (3.2) Depth indentation is also useful for the measurement of elastic properties. The local elastic modulus can be determined from the slope of the unloading portion of the indenter load versus depth curves, Fig. 3.6(c) [81]. Based on Sneddon's elastic punch model [82], and equating the projected area in contact under the indenter to the area of the punch, the elastic modulus, E, can be calculated from: E E.'dh V 7t (3.3) 40 where Eo and v 0 , Poisson's ratio, are indenter values, Ac is the contact area and dP/dh is the slope of the tangent to the unloading portion of the curve (Fig. 3.6(c)). 3.2.3 Microindentation Microindentation testing typically covers the loads ranging from 0.1-10.0 N (0.01-1.0 kg), while macroindentation covers anything larger. A load of 0.1-0.3 N is most commonly used to measure the microhardness of PS coatings for quality control purposes. Larger loads are typically used to induce indentation cracking for the purpose of measuring fracture toughness, Kjc. Knoop indentation can be used to measure the elastic modulus, as well as hardness, of a material, using methods based on the measurement of elastic recovery of the in-surface dimensions of indentations, as discussed in Section 2.3.2. This is typically done at a load of 10N. The ratio of hardness to elastic modulus, H/E, has been found to be a useful index of elastic -plastic behaviour, as well as of a material's ability to absorb impact energy [83]. 3.2.4 Macroindentation and Fracture Toughness Testing Conventional methods of determining fracture toughness were developed using linear elastic fracture mechanics on "ideal", i.e. containing few if any defects, brittle materials such as glass and single crystal ceramics, assuming the development of lateral and radial cracking, and are reviewed in [72, 75]. One equation for fracture toughness using a Vickers indenter is: Klc =0.0161 ( Z7 Y P } H C 3/ '2 (3.4) where E is the Young's modulus of the material, Hv is the Vickers hardness, P is the contact load and C is the crack length emanating from the corners of the indentation [84], see Fig. 3.7. A number of equations of similar form have been developed for determining Kc. 41 Fig. 3.7 Schematic of indentation fracture pattern for Vickers geometry [75]. The resultant models have been shown to be not entirely adequate for coarse-grained ceramic materials [78], and, thus, they are probably not adequate for ceramic PS coatings, as the lamellar microstructure is definitely not isotropic. The indentation fracture toughness value based on the formation of Palmqvist cracks, which are the first cracks to form during loading of an indentation, has been shown to be more applicable to the study of tungsten carbide and other ceramic bulk materials [47] and PS coatings [48], refer to Section 2.3.4. The values obtained may not be accurate, however, they are useful as a comparative measure to quantify the differences between coatings. The usefulness of indentation, which occurs on the macro-scale, for evaluating the Kjc of polycrystalline ceramics or other defect-filled brittle materials has been questioned due to their deviation from the ideal continuous elastic solid, as well as where wear, which occurs on the micro-scale, is concerned [8]. However, indentation is still useful for studying the response of different materials to an elastic-plastic penetration by a hard particle. The ratio H/Kc for a material has been found to be a useful "index of brittleness" [85]. The larger this value, the more brittle the material. 42 3.3 Controlled Scratch Testing Scratching is another form of damage by hard particles as it represents idealised single point abrasion. Controlled scratch testing is a common method used for mechanical testing of materials, in which a hard indenter is used to generate a groove in the tested surface. It can be used to: (1) evaluate the adhesion of thin surface coatings to the substrate, (2) measure scratch hardness, (3) evaluate the abrasion resistance of different materials and (4) study material removal and deformation mechanisms [86]. It is well established as a method for testing the adhesion of thin coatings, but not thick PS coatings. The scratching process is more complex than indentation, since it combines indentation with tangential motion across the surface. The stresses generated during scratching are a combination of the (1) elastic-plastic indentation stress field, (2) frictional stress field and (3) residual (internal) stress in the coating itself [76]. Scratch testing represents an idealized wear condition, where the indenter has a high fracture toughness and hardness with a defined diameter, causing the formation of surface or near-surface micro-crack systems. Scratch testing has recently been used to investigate the wear performance of ceramic materials as well as to study the effects that microstructure has on the wear of ceramics. As the work done on ceramic materials holds potential insights for the study of the wear of ceramic PS coatings, the important conclusions from some of these studies will be presented in the next section. 3.3.1 Scratch Testing of Ceramics The fracture mechanics approach to wear modelling of ceramics was originally proposed in a model by Evans and Marshall [76] describing the removal of material by the extension of lateral cracks developed during indentation beneath the plastically deformed groove caused by an abrasive particle, as schematically illustrated in Fig. 3.8. 43 Fig. 3.8 Schematic illustration of the crack network caused by a passing abrasive particle across a ceramic material. Ajayi and Ludema [56] performed indentation and scratch tests on a variety of bulk ceramics and found that fracture mechanics failed to account for the experimental results, mainly because the assumed linking of radial and lateral cracks did not often occur. The fracture mechanics approach is based on the macroscale measurements of 2a & c, the indentation diagonal and extended crack length (Fig. 3.7), while many of the material removal mechanisms, such as brittle fracture and fragmentation, occur on a microscale. In other words, the stress states that induce lateral and radial cracking upon indentation are not the same ones that cause material removal processes in wear. Essentially, the current wear models for ceramics , based on material hardness and fracture toughness, fail to correlate with experimental results because these models fail to account for the responses of non-homogeneous materials to the tensile component of stress that exists in the non-homogeneous contact stresses. Ajayi and Ludema recommended that models dealing with the wear of ceramics must include microstructural terms such as: grain size and strength, amount and strength of the grain boundary phase as well as the applied conditions of load, environment and surface roughness. 44 Since this work a large number of other indentation and scratching studies have been performed on a variety of ceramics in order to examine the effects of microstructure more closely. In a second study, Ajayi and Ludema [8] found that the wear modes in scratching or sliding ranged from being fracture dominated, to a combination of deformation and fracture to primarily deformation for the different ceramics tested. The strain energy induced by the imposed stresses can be relieved by either plastic flow or fracture. They concluded that intergranular fracture controls wear in ceramic materials with weaker grain boundaries and that the amount of wear is greater than that found in materials controlled by transgranular fracture. As discussed in section 2.3.6, the grain size of bulk alumina ceramics affects the observed wear modes, especially above the critical load for fracture (>10 N in this study) [58]. In fine grained alumina (0.7 um) the predominant mechanisms were grain boundary microcracking and fracture leading to material removal in the form of delamination wear sheets in the scratch groove. The medium grained alumina (5 urn) exhibited a combination of intergranular cracking and chipping at the track edges, while the coarse grained (25 um) showed both inter- and transgranular fracture in the initial scratch pass changing to predominately transgranular cracking after more than one pass. No evidence of lateral cracking was found in any of the ceramics. The wear volume removal rate was found to be proportional to P" (load), and linked to the load-sensitive quasi-static maximum tensile stress developed behind the moving indenter. The same results were obtained using a sharp Vickers indenter [88], showing that this behavior was not shape dependent. In a series of studies on scratching of alumina, Xu and Jahanmir [89-91] found that the main wear mechanisms were: (1) microfracture and chipping within individual grains by crack propagation along intragrain twin/slip bands, (2) intergranular microcracking and grain dislodgement and (3) removal of large sections by lateral crack propagation. The last mode was seen only in the fine-grained alumina (3 um). A fracture mechanics wear model was proposed, 45 where the wear volume removal rate was found to be proportional to / (grain size) and P (load), and short-crack toughness values were used instead of the more common long-crack, or macro-) toughness, where "short-crack" means cracks comparable to the grain size. The toughness values were calculated from strength measurements performed in biaxial fracture tests on samples indented at various loads to produce a range of indentation crack sizes. Short-crack toughness is the toughness determined at a short indentation crack length, i.e. comparable to the grain size. The toughness values tended to increase with increasing crack size [90]. Scratch testing has also been used to investigate the onset of the ductile-to-brittle transition for a variety of ceramics [92]. Changes in the tangential force necessary for scratching, as recorded by the test equipment, were found to be an effective way of detecting this transition. 3.3.2 Scratching of Ceramic Plasma Sprayed Coatings In comparison, very few scratching studies have been performed on ceramic PS coatings. One study [93] attempted to determine the critical adhesion load for zirconia PS coatings using a series of constant loads rather than the more common continuously increasing load technique. The highest average acoustic emission count under constant load was taken as the critical load, however the actual debonding of the coatings was never shown. In a novel test for measuring the cohesive fracture toughness of a PS coating [94], a Vickers indenter was used to scratch across the polished cross-section of a coating from the substrate to the coating surface, perpendicular to the interface until a half-cone of the coating material broke away from the body of the coating. A schematic view of the half-cone is shown in Fig. 3.9. 46 Substrate Lc iZi Fig. 3.9 Schematic view of the half-cone fracture upon cross-sectional scratching of a ceramic PS coating. (Adapted from [94].) The normal force, Fn, was found to be related to the length of the half-cone, Lc, by: Fn = ALc12 (3.4) where A is a constant. This equation was very similar to the established fracture toughness equation. For a half-penny shaped crack [95]: FH = a0Kccm (3.5) where do is a geometric constant, Kc is the fracture toughness and c is the length of the critical defect. It was assumed that Lc was equivalent to 2c for the half-penny shaped crack, so that Eqns. (3.4) and (3.5) could be combined: A = aKc (3.6) where oris a calibration constant. Values were calculated for a variety of alumina-based materials and compared with sintered alumina as a calibration material since its fracture toughness was already known. The results showed that the test was sensitive to differences in composition and structure and had potential for use in studying the effects of microstructure on the internal toughness or cohesion of coatings. 3.4 Abrasive Wear Testing Abrasion can be thought of as the action of a multitude of sliding indenters, i.e. hard particles, which form grooves across the surface of the wearing material. The interaction of these 47 grooves, i.e. when close together or crossing each other, is a major factor in the wear process [91]. There are several micro-wear processes that occur in abrasive wear as a result of the physical interactions occurring between the abrasive particles and the surface of the wearing materials, known as: microplowing, microcutting, microfatigue and microcracking [55]. Microplowing and microcutting occur typically in more ductile materials while microcracking occurs in more brittle materials. However, plastic deformation does occur during indentation and abrasive wear, as discussed previously in Section 3.2.1. The threshold for ductile-brittle transition is typically greater than the load on individual asperities, being size dependent, in a wear system. This is why brittle fracture is not always observed in abrasive wear. The wear volume due to two-body abrasion by a conical abrasive particle, through plastic deformation, i.e. microplowing or microcutting, was originally described by Rabinowicz [96] as f n W 2 t a n a P ex n\ follows: — = (3.7) S K H where W is the volume loss due to wear, S the sliding distance, P the normal load, H the yield pressure or the hardness of the wearing surface and a the attack angle of the abrasive particle. The first factor of the equation depends on the geometry of the abrasive particle and can be replaced with a wear coefficient, Kat,. This leads to the same equation as Archard's Wear Law, which was originally developed to describe the wear that occurs due to plastic deformation during sliding wear. This typically takes the form of: W P — = Kah— (3-8) S H where W is the wear volume [m3], Kab is Archard's Wear coefficient, S the sliding distance [m], P the applied normal load [N] and H the hardness of the softer surface which is equivalent to p m , the yield pressure of the plastically deforming asperity [N/m2] (from indentation hardness theory) [97]. The ratio of the last two is often taken as the real contact area. 48 The wear of pure metals follows this law quite closely. That of alloys and more brittle materials do not, however, due to other (microstructural) factors that must be taken into account. The hardness and grain size of additional phases, such as carbides for example, have a large effect on the wear of the material. Abrasive wear occurs by brittle fracture as well, as is the case with ceramic materials. A number of models have been proposed, including several with similar correlations. In the model by Evans and Marshall [76], as shown in Fig. 3.7, the abrasive particles are typically modelled as single point indenters, using indentation and fracture mechanics theories [51]. Assuming volume loss by fracture and lateral cracking, they obtained a relationship of the form: where Wv is the wear volume, P is the applied load, H is the hardness of the material, KIc is the fracture toughness, E is the elastic modulus, a a constant and s the sliding distance. The key difference between this wear mechanism and that by plastic deformation is that the wear rate by lateral fracture is not directly proportional to the normal load. The size of the abrasive particle is also important, especially in the case of brittle materials. For spherical particles, there is a threshold particle size above which Hertzian fracture occurs instead of plastic flow [98]. A l l of these models proposed a dependence of the wear intensity, W, on the fracture toughness, Klc, in the form W ~ K/c where a,b vary from Vi to 2.0. Research done after these models were published has indicated that the wear volume lost in the abrasive wear of brittle materials is the result of a combination of plastic and other deformation mechanisms, including plastic grooving, delamination and brittle fracture [99]. Wear models based on bulk hardness and fracture toughness alone have not proven adequate to describe the observed abrasive wear rates. s (3.9) 49 3.5 Erosive Wear Testing Erosive wear is caused by the impingement of particles of solid or liquid against the surface of an object and is closely related to abrasive wear. The main wear mechanisms are a combination of abrasion and surface fatigue. The main parameters affecting solid particle erosion include: (1) particle flow variables such as velocity, angle of incidence and concentration (2) particle variables including shape, size, hardness,and fracture toughness, and (3) material variables such as: hardness, fracture toughness and microstructural defects. Erosion testing is used for the evaluation of a material's resistance to damage caused by particle impact. The angle of impact has a large effect on the amount of material loss. Ductile materials exhibit increased wear at lower angles of incidence due to the predominance of plastic deformation and plowing wear mechanisms, while brittle materials like ceramics exhibit increased wear at higher angles of incidence, and thus impact force, through the mechanism of brittle fracture, Fig. 3.10. Angle of Impingement Fig. 3.10 Effect of the angle of impingement on the erosion rates of ductile aluminum vs. brittle alumina. The curves display the general erosion response and not the actual magnitudes of the wear rates (After [100].) 50 3.5.1 Erosion of Brittle Materials Studies on impact damage in brittle materials have shown that the damage pattern observed is similar to that seen in the indentation of brittle materials, whether the impacting particles are angular or spherical. A zone of intense plastic deformation was seen to form immediately beneath the contact area. Since brittle materials undergo an elastic-plastic response to this impact, it is believed that the residual stresses induced by the plastic deformation are the main driving force behind the resulting surface fracture. Similar lateral and radial fractures occur, with the average force on the indenter/particle being a function of the contact radius, or radial crack length. The maximum depth of the lateral-type cracking was found to correlate directly with the particle impression radius [74]. As in abrasive wear, a number of models have been proposed for the erosive wear of brittle materials. Both quasi-static and dynamic methods have been used to estimate the contact force, or impact pressure. The quasi-static method predicts a volume erosion (the volume removed per unit mass of erodent particles, W/p) of: U/ _0.2 TT0.1 p K3 while the dynamic model predicts: ^ o c / V - 2 , f ' 6 Q 2 5 (3.11) where We is the mass of removed material per mass of erosive particles striking the surface, p is the density of the wearing material, r is the radius, v the velocity and o the density of the erodent particles, Kc is the fracture toughness and H the hardness of the wearing material. The prime factor affecting the extent of radial fracture is the fracture toughness of the material, along with the radius and velocity of the impacting particles. There is a fracture threshold, the minimum indentation size made by each particle, required to initiate fracture, that 51 is dependent on the velocity and size of the impacting particle. Below this threshold the main mechanism is plastic deformation. The material brittleness index mentioned in Section 3.2.3 has been used in the form (KJH) /r, where r is the radius of the erodent particle, to predict the nature of the dominant damage mechanism in erosion, either ductile or brittle [71]. Fracture is the dominant response for low values, while plastic deformation tends to dominate at high values. The grain size of the material has been found to have a strong effect on the erosion resistance as well [101]. So far, no satisfactory predictive models have been developed for the erosion of ceramic materials containing large numbers of microcracks, pores or inclusions, or weak grain boundaries, i.e. typical engineering ceramics. Erosive wear models based solely on lateral cracking have failed to provide accurate predictions of the erosion rate [71]. 3.5.2 Erosion of Ceramic Plasma Sprayed Coatings Based on the mechanisms occurring during the erosion of brittle solids, it would seem that erosion is a good method for evaluating the impact resistance of PS coatings and investigating the effect of some of the microstructural parameters involved. The different methods of spraying PS coatings have a large effect on the microstructural integrity produced, as discussed in Section 2.1. Vacuum PS coatings have a high microstructural integrity and consequently, a good resistance to erosion [64]. The use of erosion tests for measuring the inter-splat cohesion, or bonding in ceramic PS coatings was first suggested by Arata, Ohmori and L i [102]. Erosion tests were performed at different angles to investigate the erosion mechanisms of ceramic PS coatings. Based on the results, it was suggested that erosion occurs by a debonding mechanism where cracks propagate between the splat boundaries, i.e. interfacial cracking. As a result it could be related to the bonding strength within the lamellar structure itself, i.e.- the rate of erosive wear could be used as a statistical measure of bonding, or cohesion within the coating. In a later article [67], the same . 52 authors found that the steady-state erosion rate was inversely proportional to the average bonding rate of several ceramic PS coatings, as measured using copper electroplating techniques [13] to highlight the percentage of bonded area. This is a good illustration of the use of erosion testing as a way to examine the microstructural integrity, in this case the actual bonding contact, of ceramic PS coatings and will be investigated further in the present work. 53 B. E X P E R I M E N T A L M E T H O D S A N D R E S U L T S C H A P T E R 4 M A T E R I A L S 4.1 Introduction The aim of the present research was to study the correlations between the microstructural characteristics of ceramic PS coatings and their response to different contact stress conditions. For this purpose a variety of coating microstructures, with differing concentrations of microstuctural defects, were required in order to determine the effects of the different microstructural features on coating response. These were achieved by varying the materials and the processing conditions, as described below. 4.2 Plasma Spray Powders In an effort to obtain more uniform microstructures, three sizes of 99.9% pure, mono-crystalline a-alumina powders, nominally 5, 10 and 18 pm particle size (Sumicorundum AA-05, AA-10 and AA-18 powders, Sumitomo Chemical Co. Ltd.), were used. Such powders have not been used before for plasma spraying. Manufacturer's data indicate that these have a narrow particle size distribution, with about 80 wt.% lying between -20% and +30% of the specified size. These powders have an equiaxial morphology, in contrast to the typical commercial alumina for plasma spraying, e.g. Plasmalloy AL-1010, as shown in Fig. 4.1. The Plasmalloy powder has an original particle size range of 5-53 pm, which was subsequently narrowed to 25-38 pm by sieving to obtain a more uniform microstructure. For comparison, several Al 203-Ti02 and C r 2 0 3 powders were also used. These powders are typically sintered and crushed. Details on the powders can be found in Table 4.1. 54 Fig. 4.1 Alumina Powders: (a) 18 um monocrystalline (Sumicorundum) and (b) 25-40 um polycrystalline (Plasmalloy). 4.3 Coating Preparation The focus of this study was on air plasma sprayed ceramic coatings. The majority of the coatings were sprayed at Northwest Mettech Inc. in Richmond, B.C., using their patented axial-injection spray system, the Axial III™ [11]. For comparison, one size of alumina powder was also sprayed using a Plasmadyne radial injection PS system located at The Industrial Materials Institute, National Research Council of Canada (IMI-NRC) in Boucherville, Quebec. All coatings were sprayed onto mild carbon steel substrates, 14" x 1" x 4" in size. The substrates were grit blasted with 840 pm AI2O3 particles (24 grit size) immediately prior to spraying. 55 Table 4.1 Details of Powders Used for Spraying Notation Powder Type Crystalline Form Manufacturer/ Designation Particle Size (pm) Nominal Powder Composition % A A 5 A1 20 3 a - A l 2 0 3 single crystal-sapphire Sumitomo AA-05 4-6 Al 2 O 3 -100% AA10 A1 2 0 3 oc-Al 20 3 single crystal-sapphire Sumitomo AA-10 8-12 A I 2 O 3 - 100% AA18 A1 2 0 3 a - A l 2 0 3 single crystal-sapphire Sumitomo AA-18 15-23 Al 2 O 3 -100% AL30 A1 2 0 3 0C-Al 2O 3 Plasmalloy Al-1010 25-38 S i 0 2 - 2.0 ppm C a O - 1.0 ppm NaO - 1.0 ppm A 1 2 0 3 - 98.00 A T I AI2O3-13Ti0 2 a - A l 2 0 3 T i 0 2 (Rutile) Norton 106 15-45 A I 2 O 3 - 87.0 T i 0 2 - 13.0 AT2 AI2O3-13Ti0 2 a - A l 2 0 3 T i 0 2 (Rutile) Norton 107 5-30 A I 2 O 3 - 87.0 T i 0 2 - 13.0 CR1 C r 2 0 3 C r 2 0 3 Praxair CRO-174 5-20 Cr 2 0 3 ->99.5% CR2 C r 2 0 3 C r 2 0 3 Praxair CRO-172 20-45 Cr 2 0 3 ->99.5% 4.3.1 Spray Parameters for the Sapphire Alumina Coatings The primary gas used in the axial plasma spraying was nitrogen, with argon and hydrogen as secondary gases. The process parameters were varied in order to provide a range of microstructures and to study the effects of these parameters on the coatings produced from the monocrystalline alumina powders. Two process parameters were varied between high and low values: the percent hydrogen used as a secondary plasma gas and the orifice diameter of the plasma torch nozzle, Table 4.2. A n increase in the hydrogen concentration increases the thermal 56 conductivity and heat content of the plasma jet [9]. An increase in the nozzle size increases the powder residence time in the jet. If a smaller nozzle is used, the particles will travel faster through the jet, giving a higher impact velocity. At the high values of both parameters, the powder particles are travelling slower through a hotter and more thermally conductive jet, allowing for more complete melting of the particles. Deposition of an ideal coating results from exposing the powder particles to an appropriate temperature for a period of time adequate for complete melting of all particles. The matrix of processing conditions, along with their effect on heat input and particle residence time in the plasma, is summarized in Table 4.2. Table 4.2 Processing Conditions for Alumina PS Coatings Spray Run # Vol. % H 2 Relative Heat Torch Nozzle Relative Input Size [mm] Residence Time 1 10 Low 12.5 Less 2 10 Low 14.1 More 3 20 High 12.5 Less 4 20 High 14.1 More Three sapphire alumina powders (AA5, AA10, AA18), along with the commercial alumina (AL30) powder were sprayed. Due to a limited amount of AL30 powder, only runs 1, 2, and 4 were sprayed with this powder. The substrates were preheated to 100-120 0 C immediately prior to the spraying of the coating material. The spray distance for all the coatings was 85 mm. Details of the spray conditions for the coatings are listed in Table 4.3. Due to the fine size of the AA5 powder, the feed rate was halved and the spray time doubled. This resulted in coatings of comparable thickness to the other coatings. For comparison, the AA18 powder was also sprayed using a Plasmadyne radial injection PS system. The primary gas used with this system was argon, with 32% helium as the secondary gas. The spray distance was 76 mm and the resulting coating thickness was 345 pm. 57 Table 4.3 Details of Matrix Spray Parameters Run No. Nozzle Size % Ar % N 2 % H 2 Powder Feed Torch Coating [mm] rate (g/min) Passes (#) Thickness (pm) AA5-1 12.5 10 80 10 16 92 255 -2 14.1 10 80 10 16 92 240 -3 12.5 10 70 20 16 92 230 -4 14.1 10 70 20 16 92 255 AA10-1 12.5 10 80 10 34 70 530 -2 14.1 10 80 10 34 46 365 -3 12.5 10 70 20 .34 46 360 -4 14.1 10 70 20 34 46 350 AA18-1 12.5 10 80 10 35 46 340 -2 14.1 10 80 10 35 46 355 -3 12.5 10 70 20 35 46 350 -4 14.1 10 70 20 35 46 370 AL30-1 12.5 10 80 10 32 58 230 -2 14.1 10 80 10 32 60 360 -4 14.1 10 70 20 32 30* 170 (Insufficient Powder) 4.3.2 Spray Parameters for Remaining Coatings Details on the spray parameters for the alumina-titania and chromia coatings are found in Table 4.4. Two Al203-Ti02 powders were sprayed under the same conditions used for spray run #1 and are listed in Table 4.2. The spray distance for all the A T and CR coatings was 100 mm. These coatings were used for comparison with the sapphire powder derived alumina coatings. The latter will be designated the AA-alumina coatings in the remainder of this thesis. 58 Table 4.4 - Details of Spray Parameters for Alumina-Titania and Chromia PS Coatings Run No. Nozzle Size [mm] % Ar % N 2 % H 2 Powder Feed rate (g/min) Torch Passes (#) Coating Thickness (pm) ATI 12.5 10 80 10 35 36 280 AT2 12.5 10 80 10 35 36 270 CR1 11.1 75 10 15 38 80 200 CR2-1 9.5 - 85 15 42 36 270 CR2-2 14.1 90 . 10 42 36 315 4.4 Reference Alumina A bulk, sintered alumina, #CC620 from Sandvik Coromant in Sweden, was used as a reference material in the majority of tests. Information from the manufacturer indicates that the alumina has a composition of 97 vol % A12C>3 + 3 vol % Z r 0 2 with a density of 3.95 g/cm3. 59 CHAPTER 5 COATING CHARACTERIZATION: MICROSTRUCTURE AND MORPHOLOGY Table 5.1 presents a summary of the techniques used in the microstructural analysis of plasma spray (PS) coatings. Scanning electron microscopy (SEM), with its high resolution and depth of field at high magnifications, was used extensively throughout this work for evaluation of both microstructural parameters and wear mechanisms. The contrast imaging available in the backscattered electron (BSE) mode in the S E M was used for highlighting crack interfaces and examining porosity. Image analysis was used to measure the amount of porosity along with the angular crack distributions in the different coatings. These techniques will be presented in more detail, together with the results obtained, in the following sections. Further details on the microstructure and morphology of the AA-alumina coatings can be found in [103]. Table 5.1 - Analytical Techniques Used in Evaluating Coatings Analytical Technique Information Obtained Light optical microscopy (LOM) Microstructural changes, defects Scanning electron microscopy (SEM) Surface topography, splat and wear with Back Scattered Electron (BSE) morphology, fractography, porosity and imaging angular crack distribution Image Analysis on S E M photos Porosity determination, angular crack distribution Surface Profilometry Surface topography, wear volume X-ray diffraction (XRD) Phase composition 60 5.1 Preparation for Coating Microstructural Examination The quality of preparation of polished top- and cross-sections of ceramic PS coatings for examination using optical or electron microscopy strongly affects the results of any analysis performed [104]. This is especially true when using image analysis for determining the amount of porosity, the crack distribution and the general quality of the coating microstructure. It is equally critical when performing depth sensing indentation, or even microhardness. The presence of microstructural defects that typically occur in PS coatings, including porosity, unmelted particles and microcracking, make it important to impregnate the coating with epoxy prior to polishing to minimize the number of pullouts. This is best achieved by using a long curing epoxy such as Buehler's Epoxide (Buehler Ltd, Lake Bluff, IL. USA). It was found to be most effective when both the specimen and the epoxy resin are heated in an oven set at 60-65° C [105]. The specimen is heated for 20-30 minutes depending on the size, and the resin until the viscosity is quite low (10-15 minutes). The sample is then placed in a mounting cup with the coating side up and the epoxy mixture poured until it just covers the top. This is best done under vacuum, to obtain maximum infiltration into the coating. After the epoxy has hardened, the sample can be cut with minimal damage to the coating and then remounted for polishing. After much experimentation it was found that the following polishing procedure worked best for the ceramic PS coatings used in this work: - 240 grit SiC paper at 5 lbs. pressure per sample for 40 seconds for flattening - 6 pm diamond suspension on a hard pad at 5 lbs. pressure per sample for 5 min. - 3 pm diamond suspension on a med. hardness pad at 6 lbs. pressure per sample for 3 min. - 0.6 pm colloidal silica supension on a soft pad at 8 lbs. pressure per sample for 3 min. Most of the polishing was performed on a Buehler Motopol 2000. See Appendix A for detailed mounting and polishing procedures. Polished cross-sections of each coating were examined using an optical microscope for measurement of coating thickness and microstructure reference photographs, while a mixture of secondary electron (SE) and BSE modes was used to obtain 61 digital S E M images with the best possible resolution and contrast to investigate the different types of microstructural defects in greater detail. 5.2 Porosity Determination Any quantitative porosity measurement method, such as mercury intrusion porosimetry (MIP), does not guarantee determination of a true porosity value due to the presence of closed pores within the coating. Image analysis is somewhat subjective, therefore, it is best used for a relative comparison between a number of different coatings. In this work, eight to ten lOOOx magnification digital micrographs, size 640 x 480 pixels, were taken from the outer 2/3 of each coating cross-section using a blend of 70% BSE and 30% SE images on the SEM. Image analysis* of the SE/BSE micrographs was used to measure the percent porosity for each coating. Pores with dimensions smaller than 0.2 pm were not resolved and thus were not included in this value. 5.2.1 AA-Alumina Coatings The porosity measurement results for the AA-alumina coatings are presented in Fig. 5.1. Views of cross-sectional microstructures at low (200x) and higher (lOOOx) magnification are shown in Fig. 5.2 and 5.3 for the low and high heat input runs for each powder size. The lower magnification photographs, taken using optical microscopy (OM), give a good overall view of the coating porosity and structure, however the higher magnification images allow a closer look at the pore structure as well as the crack distribution. A l l of the coatings exhibited a decrease in porosity with increase in heat input and residence time. The AA5 coatings in particular appear to be most affected by the increase in heat input to the plasma. As discussed in section 2.2.4, the impact velocity and temperature of the incoming particles during spraying both have an effect on the amount of porosity. Large pores, such as * UTHSCSA Image Tools version 1.27 - developed at Dept of Dental Diagnostic Science at University of Texas Health Science Center, San Antonio, Texas. 62 those visible in the AA5-1 coating, Fig. 5.2(a) and 5.3(a), are formed at the boundaries of overlapping splats, which are incompletely flattened due to the lower heat energy in the plasma in this run. As the heat input to the plasma increases, along with the residence time in the flame, the amount of porosity decreases. It is believed that this is due to increased particle flattening caused by the decreased viscosity of the molten particles and the resulting improvement in flowability and interparticle wetting. Details of the different splat morphologies seen wi l l be covered in Section 5.3. This is also observed to a smaller extent in the AA10-1 and AA18-1 coatings, with a decrease in porosity as the heat input and residence time increase. The APS coating, sprayed by a radial injection feed torch instead of an axial injection one, showed a far greater amount of fine interlamellar porosity or horizontal delaminations, Fig. 5.3 (g), along with larger spherical pores, giving this coating the highest porosity of all the coatings derived from the AA18 powders. Each powder size was affected by the changes in processing conditions in a slightly different manner. The larger particles, AA18, were not completely melted during the lower heat input spray runs (#1-2). This is evidenced by the unmelted particles visible in Fig. 5.2(e). However, the AA10 particles appear to have been predominately molten during all four spray runs. 12 10 o a. 2 H AA5 AA10 Powder AA18 Fig. 5.1 Comparison of average porosity of AA-alumina coatings as a function of powder type and processing condition. (Sample symbols and numbers refer to Tables 4.3- 4.4.) -1 -4 63 AA5 AA10 AA18 100um 10Qum Fig. 5.2 OM micrographs of polished coating cross-sections: (a) AA5-1, (b) AA5-4, (c) AA10-1, (d) AA10-4, (e) AA18-1, (f) AA18-4 and (g) APS. •1 64 Fig. 5.3 70% BSE / 30% SE micrographs of polished coating cross-sections: (a) AA5-1, (b) AA5-4, (c) AA10-1, (d) AA10-4, (e) AA18-1, (f) AA18-4 and (g) APS. 5.2.2 Other Coatings 65 The amount of porosity for the commercial alumina, alumina-titania and chromia coatings was decreased with the increase in residence time, from runs 1 to 2, as shown in Fig. 5.4. Views of the cross-sectional microstructures are presented in Fig. 5.5-5.6. The commercial alumina (AL30) was highly porous, as shown in Fig. 5.5-5.6(a) & (b). The larger size of the particles contained in this powder (25-38 pm) resulted in a large number of unmelted particles, as well as a lower impact velocity for the molten particles. This resulted in poor flattening and spreading out of the incoming splats, thereby forming large pores at the boundaries between splats. 12 1 T a -_ 1 6 -4 2 0 A I 3 0 A T C R P o w d e r Fig. 5.4 Comparison of average porosity of reference coatings. The alumina titania coatings had low porosity, primarily in the form of fine spherical pores. The alumina and titania phases were visible in both coatings, especially in AT2. The C R I coating was the least porous of all the coatings, comparable to the reference sintered alumina used in this work with less than 1% porosity. The lamellar structure was not visible in this coating either, instead it resembled the microstructure of a bulk ceramic, see Fig. 5.6(g). The amount of porosity seen in CR2-2 was lower than that in CR2-1, due to the increased residence time and subsequent heating of the particles caused by the increase in nozzle size in the #2 run, refer to Table 4.4. 66 (b) Fig. 5.5 OM micrographs of polished coating cross-sections: (a) AL30-1, (b) AL30-4, (c) ATI, (d) AT2, (e) CR2-1, (f) CR2-2 and (g) CR1. 67 68 5.3 Splat Morphology The splat morphology of the coatings was investigated by examining as-deposited top surfaces along with polished and fractured cross-sections. The fracture surfaces were prepared by sawing a notch from the bottom of the W thick substrate through the substrate to within 1-2 mm of the coating-substrate interface. The notch must be wide enough to allow space for complete fracture of the coating and substrate. The specimens were then cooled in liquid nitrogen and fractured with the coatings in tension. The lamellar structure of the coatings was examined using BSE imaging in the SEM, thereby highlighting the crack boundaries for subsequent measurement. At least 20-30 splats were measured for both diameter and thickness values. The surface topography of the coatings was quantified by measuring their average surface roughness, where y is the height of the surface above the mean line at a distance x from the origin and L is the overall length of the profile under examination [71]. A stylus profilometer (Talysurf 5, Rank Taylor Hobson, Leicester, England), with a stylus tip radius of 2.5 pm, was used for this purpose. Several definitions of the term lamellar have been used in the literature. The most common use of the term is to describe the appearance of the microstructure itself, and the form of the individual flattened particles [5]. It has also been defined, however, as corresponding in thickness to that applied from each passage of the plasma torch across the substrate [6], as opposed to indicating the thickness of the individual splats. In this thesis, the word "pass" will be used to describe the thickness applied during each passage of the torch, and the words lamellae or splat will be used to describe the individual pancake-like droplets. 5.3.1 AA-Alumina Coatings The splat morphology of the coatings is strongly affected by differences in precursor powder size as well as the coating spraying conditions. This can be seen in the splat diameters as R , defined as: 0 (5.1) 69 seen on the top surfaces of the as-deposited coatings, and the splat or pass thickness as seen on the coating fracture surfaces. In coatings where grain growth through the splats occurs, individual splat borders are impossible to distinguish. The pass thickness, i.e. the thickness sprayed by each pass of the torch, see section 2.2.2, is more easily measured in this case. Individual splat thickness is typically 1-3 pm. The measured average splat diameters and pass thicknesses for the alumina coatings are shown in Fig. 5.7. 80 70 60 50 _ , 40 a 30 20 10 H o 10 (a) 1 2 AA5 8 -X (b) AA5 m AA10 P o w d e r AA18 AA10 Powder AA18 Fig. 5.7 Average dimensions: (a) splat diameter (as measured on coating planar surfaces) and (b) pass thickness, as a function of powder type and processing conditions. A comparison of the changes in splat diameters due to differences in precursor powder size and spraying conditions is shown in the SEM photomicrographs in Fig. 5.8. The largest differences observed occur between the powder sizes, as shown in Fig. 5.8(a) and (b). There is a 70 trend within each powder series, from runs 1 to 4, to increased splat diameter and decreased relative pass thickness, as shown in Fig. 5.7(a). This effect is visible on the top surfaces of coatings AA18-1 and AA18-4 in Fig. 5.8(c) and (d). The AA5 coatings, Fig. 5.8(a), appeared porous and poorly consolidated, with no clearly defined splat structure. The presence of numerous small particles is probably due to insufficient impact momentum. The AA10 coating, Fig. 5.8(b), exhibits the best splat morphology, with densely packed, overlapping, "pancake-like" splats. The best AA18 coating (AA18-4, Fig. 5.8(d)) also showed a dense microstructure, but with some evidence of break-up and f low at the splat edges. The surface morphology was examined by measuring the surface roughness giving Ra values for each of the as-sprayed coatings, Fig. 5.9. Confirming the microstructural observations using secondary electron (SE) images on the SEM, as shown in Fig. 5.8, the AA5 coatings have the roughest surface and the AA10 coatings the smoothest. E31 Fig. 5.8 SEM photomicrographs of changes in splat diameter with increased particle size: (a) AA5-4, (b)AA10-4, and change in heat input: (c) AA18-1 , (d) AA18-4. 71 8 P H 1 2 A P S AA5 AA10 Powder AA18 Fig. 5.9 Surface Roughness (Ra values) of the AA-alumina coatings as a function of powder type and processing conditions. The pass thicknesses were measured by detailed SEM examination of fractured cross-section surfaces, examples of which are found in Fig. 5.10. Note the lack of a lamellar structure in the AA5 coating, indicating the poor bonding achieved between the splats. The AA10 and AA 18 coatings show increased grain growth through the splats and lamellae indicating improved intersplat cohesion, thus the measured thicknesses correspond to each pass of the torch. The APS coating shows a less cohesive structure indicating less than optimum processing of the sprayed particles. 72 Fig. 5.10 SEM photomicrographs of combined fracture/top surfaces: (a) AA5-4, (b) AA10-4, (c) AA18-4 and (d) APS (radial injection sprayed AA18 powder). 5.3.2 Other Coatings SEM photomicrographs of the as-deposited top surfaces of the commercial alumina (AL30), alumina-titania (AT) and chromia (CR) coatings are presented in Fig. 5.11. The AL30 and AT coatings have a similar appearance to the AA-alumina coatings. The CR coatings have a very different appearance however, making it more difficult to distinguish the individual splats, thus the average splat diameters of the CR coatings could not be measured. The average splat diameters for the commercial alumina (AL30) and alumina titania (AT) coatings are shown in Fig. 5.12. The AL30 coatings exhibit an increase in diameter from run #1 to 4. Note that the diameters of AL30-1 are similar to those of AA18-4 coatings, due to the insufficient impact momentum and melting of the AL30 particles. This powder has a much larger particle size range than the sapphire powders, which leads to a higher scatter in the diameter measurements as well. This is illustrated in Table 5.2, which lists the maximum and minimum splat diameters. A 7? comparison of the average surface roughness values (Ra) is presented in Fig. 5.13. The larg diameter AL30 coatings exhibited the roughest surfaces of all the coatings. Fig. 5.11 SEM photomicrographs of as-deposited surfaces: (a) AL30-1, (b) AL30-2, (c) AL30-4, (d) ATI, (e) AT2, (f) CRI, (g) CR2-1 and (h) CR2-2. AL30 AT P o w d e r Fig 5.12 Average splat diameters of the polycrystalline alumina (AL30) and alumina-titania (AT) coatings. 74 8 _S 4 . 2^ 2-1 2-2 AL30 AT Powder CR Fig. 5.13 Surface Roughness (Ra values) of the reference coatings as a function of powder type and processing conditions. Table 5.2 Splat Diameters (in pm) for AL30 and AT coatings Coating Avg [pm] Std Dev Min Max Median AL1 45.6 11.6 25 81 46 AL2 56 11.6 28 92 58 AL3 71.1 10 52 102 69 ATI 51.2 10.9 30 78 53 AT2 39 14.7 13 83 35 Typical fracture surfaces of the different coatings are shown in Fig. 5.13. The average splat thickness of the CR coatings was impossible to measure due to the lack of visible intersplat boundaries, see Fig. 5.14. The fracture surfaces of the CR coatings appeared more like that of a bulk ceramic. The pass thicknesses were difficult to measure in the AT coatings, due to extensive grain growth through multiple splats, while the large number of unmelted particles in the AL30 coatings made it difficult to measure in these coatings as well. 75 Fig. 5.14 SEM micrographs of typical fracture surfaces: (a) PL30, (b) ATI, (c) CRI and (d) CR2 coatings. 5.4 Crack Distribution Analysis Two methods of analyzing the network of microcracks in the coatings were performed in this work, using the same lOOOx magnification SEM micrographs (8-10 per coating) previously used for measuring porosity. The angular crack distribution, i.e. length, number and angular distribution of the cracks, was measured using the NIH Image 1.61 program* for image analysis. The algorithm used is summarized below [106]. 1. Correct for non-uniform lighting of optical photos with a high pass filter, and for any pullout or damage region. 2. Threshold the image and make it binary. 3. Skeletonize the binary image until the pores and cracks are one pixel wide. 4. Using image dilation and erosion, locate and break all crack intersections. 5. Tabulate the length, major and minor axis along with the angular orientation of each crack. * U.S.National Institute of Health (NIH) Image program version 1.61 76 As this method did not provide a measure of inter-crack spacing, linear crack densities, i.e. the number of vertical cracks per mm, were also determined for both the vertical and horizontal cracks using a simple grid method on the SEM micrographs [6]. 5.4.1 AA-Alumina Coatings Comparison of the angular crack distribution for the different coatings provides an insight into some effects of the variations in processing conditions. A comparison of the ratio of the total cumulative lengths of the horizontal (0-30 degrees) to vertical (60-90 degrees) cracks is given in Fig. 5.15. A graph of the average total cumulative length (in microns) vs. the angle (0° is parallel to the substrate) for the 3 powder sizes is shown in Fig. 5.16. The cumulative lengths reflect the total length of the cracks and do not relate to the crack widths in any way. 6 i ro CC t 4 o r i 2 N AA5 1 AA10 Powder AA18 Fig. 5.15 The ratio of horizontal to vertical cracks for the AA-alumina coatings, as a function of powder type and processing conditions. 77 0-10 10-20 20-30 30-40 40-50 50-60 60-70 70-80 80-90 Angle [degrees] Fig. 5.16 Typical angular crack distribution for the 3 powder sizes (AA5-4, AA10-4 and AA18-4). The coatings produced from the 5 pm particles have a much higher ratio of horizontal to vertical cracks. This agrees with the structure observed on the fractured surfaces. There are a large number of horizontal delaminations along with minimal flattening of the splats. This is believed to be due to these particles having too little mass to achieve a sufficient impact velocity to produce a void-free coating. In contrast, the AA10 coatings all exhibited an evenly distributed crack network, while the larger AA18 coatings progressed from an even ratio to twice as many vertical cracks. This may be a consequence of the greater inter-splat cohesion of the larger particle derived coatings. As described in Section 2.2.7, the greater the inter-lamellar bonding, the more constrained the splat surfaces of the freshly impinged splats and the greater the number of microcracks that form to relieve the built-up stresses. The measured densities of both vertical microcracks and horizontal delaminations for the coatings in the alumina matrix are shown in Fig. 5.17. Note that the AA5-4 coating contains roughly 2.5 times the vertical crack density of the other AA5 coatings, indicative of the increase in constraint of the particles as they solidify. This is also visible in the increased vertical crack density seen in the AA 10 and AA 18 coatings. The APS and AA5 coatings contain the highest 78 densities of horizontal delaminations, indicating poor intersplat bonding. The clearest trend toward decreased horizontal delaminations is shown by the AA18 series of coatings, while the vertical crack density remained basically constant. 220 200 r-,180 E160 E —140 w o120 CD 100 o S 8 0 <1> > (a) dn AA5 r f i 3 A A I O Powder AA18 220 200 180 '160 £ £ 1 4 0 £ l 2 0 o 2 1 0 0 O N 8 0 1 I" 40 20 (b) rH Tl Hn AA10 Powder Fig. 5.17 Comparison of average linear crack densities [mm"1] for the AA-alumina coatings: (a) vertical microcracks, (b) horizontal delaminations. 5.4.2 Other Coatings A comparison of the ratio of the total cumulative lengths of horizontal (0-30 degrees) to vertical (60-90 degrees) cracks for the AL30, AT and CR coatings is shown in Fig. 5.18, while a comparison of the horizontal and vertical crack densities is given in Fig. 5.19. The effect of increasing heat input to the powder particles in the AL30 series is shown by the slight downward trend in the ratio of horizontal delaminations to vertical cracks. These coatings had a high density of horizontal delaminations due to poor bonding and splat flattening along with a large number of unmelted particles. The alumina-titania (AT) coatings showed the highest vertical crack density of all the coatings, indicating a very high percentage of intersplat bonding, followed by the chromia (CR) coatings. Due to the lack of a lamellar structure in the CR1 coating, both vertical and horizontal cracking were extremely difficult to distinguish, see Fig. 5.5(e). 79 6 o 4 O 3 c 01 5 2-1 2-2 AI30 AT Powder CR Fig. 5.18 The ratio of horizontal to vertical cracks for the reference coatings. 220 200 ^180 1160 O120 CU (J100 re 80 u r 60 O > 40 20 0 (a) 2-1 2-2 220 200 180 "160 £ E 140 J2 120 o S 100 o N 80 I 60 40 2C -j 0 (b) [—t— 2-1 2-2 At30 AT Powder CR AT Powder Fig. 5.19 Comparison of average crack densities [mm1] for the coatings: (a) vertical microcracks, (b) horizontal delaminations. 5.5 Elemental and Phase Composition The elemental and phase composition of the powders and the as-deposited coatings were obtained by X-ray powder diffractometry (XRD) using a Siemens X5000 diffractometer with Cu K- a radiation. 5.5.1 AA-Alumina Matrix The as-deposited alumina coatings consist mainly of y-alumina, as a result of the rapid solidification of the molten precursor a-alumina powder, see Fig. 5.20 below. The amount of a-alumina present is indicative of the presence of unmelted and/or recrystallized particles in the 80 coating [16]. Note the presence of the a-phase in the lower heat input processing conditions (run #1-2) for the larger AA18 particles, while the AA10 coatings show none. The AA5 particles are quite small, thus completely molten in the plasma. However, in runs #2-4 with higher heat input and residence time, a small amount of the a-phase was detected in these coatings, Fig. 5.20. on +-> c o O (a) c o O (b) o O 2,000 1,600 1,200 800 400 c 3,000 2,500 2,000 1,500 1,000 500 i 3,000 2,500 2,000 1,500 1,000 500 y y a a JL a . y A A 5 30 35 40 45 50 55 60 65 70 4 3 2 1 l i i i i i i l l i i ' tDftHflf ' t l l i ' - 1 — — " A A 1 0 y y i t i f^SniLHifAni / l m,Aiuvjwf.nii^ tH'^ i^ i^ ffi.iiiuii r\in 11 25 30 35 40 45 50 55 60 65 70 4 3 2 1 (C) j Ct u . ,j» . 25 30 35 A A 1 8 ^ . . . • ^ " H p u n i i (i a A a »< IrWfrlHl' lift l|»»H.»lMT»<W J ' lift lllUll^ *- -\JMttJ V ' 55 60 65 70 4 3 2 1 A P S Fig. 5.20 X R D spectra for the alumina coatings deposited in the four runs: (a) AA5, (b) AA10 and (c) AA18 (includes the APS deposition of AA18 monocrystalline alumina powder using commercial radial powder injection system). 5.5.2 Other Coatings X R D spectra for the AL30, A T and CR coatings are given in Fig. 5.21. Note the decrease in the a-alumina phase, combined with an increase in the y-phase, in the AL30 runs 1-4, again due to the increase in heat input in the plasma and the residence time. The changes are not as 81 pronounced as in the AA18 coatings, due to the larger amount of unmelted particles present in the larger diameter powder AL30 coatings. c O O (a) c 3 o CJ (b) c 3 o o 2,000 1,600 1,200 800 400 c 2,000 1,600 1,200 800 400 2,000 1,600 1,200 800 400 AL30 •AW, A m m n--i'<i ii ^-A-A—xtiAt, i A JL nA 4 2 1 25 30 35 40 45 50 55 60 65 70 * V j Y J 25 30 35 40 45 50 55 60 65 70 (C) CR ~KJK , 2-2 2-1 1 25 30 35 40 45 50 Angle 55 60 65 70 Fig. 5.21 XRD spectra for: (a) commercial alumina (PL30), (b) alumina titania ( A T I , AT2) and (c) chromia (CR1, CR2-1, CR2-2) PS coatings. X R D indicated y-alumina (Y-AI2O3) to be the major phase in both of the AI2O3- 13 wt% TiC»2 coatings ( A T I , AT2), see Fig. 5.21(b). Neither the aluminum titanate alloy, AI2T1O5 or pure titania, TiCh, were detected, as would be expected i f the coatings contained a higher weight percentage of titania [107]. X R D indicated that the chromia coatings consist of predominately C R 2 O 3 , labelled 1 in Fig. 5.21(c), however, the CR2-2 coating also contained a small amount of C r 0 3 , labelled 2 in the figure. 82 CHAPTER 6 MICROMECHANICAL METHODS 6.1 Introduction In order to advance understanding of the relationship between the different microstructural parameters found in PS ceramic coatings, detailed in Chapter 2, several micromechanical test methods involving damage by hard particles were used. These included indentation techniques, including depth sensing indentation, micro- and macrohardness testing, along with controlled scratch testing, representing ideal single point abrasion. This chapter presents the experimental methods used along with the results of these tests for evaluating the responses of a variety of coating microstructures compared to those of a bulk sintered alumina. 6.2 Indentation Techniques In this work, indentation was performed in three load regimes, in the ranges of 0. IN, 1.0 N and 100 N . Standard Vicker's indenters were used with depth sensing indentation (DSI) at 0.1 N and 1.0 N , while conventional microindentation was used at 3.0 N and macroindentation at a load of 100 N . Knoop indentations were performed at a load of 10 N both parallel and perpendicular to the cross-section, as well as on the planar surfaces. Figure 6.1 illustrates the positioning and relative scale of indentation impressions vs. splat size and other microstructural features on the cross-sectional and planar (top) surfaces of the coatings. Additional details on the use of these methods on ceramic PS coatings can be found in references [32, 108]. 83 Parallel Fig. 6.1 Schematic of the scale of different indentation impressions and microstructural parameters on a ceramic PS coating. 6.2.1 Depth Sensing Indentation (DSI) Depth sensing indentations (DSI) were made on both polished cross-sections and top (planar) surfaces of the sapphire powder derived alumina (designated as the AA-series) coatings using a Fischerscope H 100B system [Fischer Technology, Inc. CT., U.S.] equipped with a Vickers indenter. Two loads were used, 0.1 N , which is within the range considered as nanoindentation and 1.0 N , which is in the microindentation range. Values of the effective elastic modulus (Eeff) can be calculated from the load vs. displacement curve, as discussed in Section 3.2.3. In the present work, an estimated Poisson's ratio of 0.23 was used for the alumina coatings in order to calculate the Young's modulus values from Eeg [109]. The ratio H/E , an index of elastic-plastic behaviour, was also evaluated. At least 15-20 measurements were performed at each load on the top third of the coating cross-sections and the top (planar) surface. Average values and standard deviations (S.D.) of the hardness, elastic modulus, and H/E ratio are summarized in Fig's 6.2 and 6.3. These results are plotted in two figures to show the different indentation responses of the coatings to DSI at the nano- and micro-scales. The error bars on all the graphs correspond to + 1 standard deviation (S.D.). 84 (a) X (b) 400 350 300 250-O2C0-m 150-MO-SS o (c) Cross-Sectional nh H i AA5 [Tjf H„(0.1M)-X-Sec AA10 Powder E(0.1N)-X-Sec m AA10 Powder t Hi AA10 Powder AA18 2.151 x HrE(0.1N)-X-Sec 0.12-i Planar 400-350-300-250-0. (3 200-UJ 150-100-50-0-H.f0.1N)-Top AA10 Powder E(0.1N)-Top AA10 Powder H/E(0.1N)-Top AA10 Powder AA18 Fig. 6.2 Comparison of: (a) hardness (H), (b) elastic modulus (£), and (c) H/E, from DSI at 0.1 N load obtained on cross-sectional and planar surfaces for the AA-series alumina coatings. (Sample symbols and numbers refer to Tables 4.2- 4.3.) 85 Cross-Sectional (a) O 15 H„(1N)-X-Sec (b) (c) Hi ^ r f i 400-350 -300 -250-S 200-1U 150-100-50-0-Tl rjjirri AA10 Powder E(1.0N)-X-Sec 3 L rrrh H/E(1N)-X-Sec AA10 Powder AA18 25 400 350 300 250 a 2 0 0 150 100 5o^ 0 0.00 Planar AA5 AA5 H»(1N)-Top rf i rH AA10 Powder E(1.0N)-Top | t f | i l l AA10 Powder hVE(1N)-Top AA10 Powder Fig. 6.3 Comparison of: (a) hardness (H), (b) elastic modulus (£), and (c) H/E, from DSI at 1.0 N load obtained on cross-sectional and planar surfaces for the AA-series alumina coatings. (Sample symbols and numbers refer to Tables 4.2- 4.3.) 86 Table 6.1 shows the approximate number of splats probed by the indenter at each load, estimated from the ratio of mean indent diagonal length to mean splat thickness on the cross-sectional surfaces and to mean splat diameter on the top surfaces, see Section 5.3.1. The cross-sectional ratios increase as pass thickness decreases and the top ratios decrease as the splat diameters increase due to the effect of the increased residence time and heat content in the plasma as the spray conditions change from runs 1-4, Section 4.3.1. Essentially, Fig. 6.2 and Fig. 6.3 represent the response of coatings to single or multiple splat probing, depending on the splat size, by the indenter. Low load indents were taken as far as possible from visible splat boundaries and pores. Table 6.1 Relative size of indentation footprint (diagonal length) to average splat thickness and diameter. Cross-section Top Load (N) 0.1N IN 0.1N IN AA5-1 1.40 6.37 0.35 1.63 -2 2.05 8.80 0.31 1.37 -3 1.95 7.40 0.35 1.33 -4 4.10 15.70 0.24 1.01 AA10-1 1.95 8.10 0.16 0.60 -2 1.95 8.35 0.13 0.58 -3 4.30 16.60 0.13 0.54 -4 3.90 15.50 0.15 0.55 APS 1.50 6.30 0.10 0.38 AA18-1 2.10 8.25 0.07 0.31 -2 2.00 8.20 0.07 0.29 -3 4.10 15.20 0.07 0.29 -4 4.10 14.40 0.07 0.31 DSI at the nano-scale (0.1 N) measures the intrinsic properties of the splats themselves when the test volume is smaller than the splat dimensions, as was the case for the tests performed on planar surfaces. However, the scatter in the results (compare the S.D. bars in Fig's. 6.2 and 87 6.3) was far greater than that seen at larger loads. This indicates increased sensitivity to the microstructural variations and inhomogeneities present within individual splats, refer to Fig's. 2.5 and 5.10 for detailed views of both top and cross-sectional surfaces. This was primarily due to the small size of the indentations themselves, with diagonal lengths ranging from 3.1-4.5 pm and depths ranging from 0.443- 0.643 pm, as well as the additional affected volume defined by the radius of the plastic zone, Fig. 3.2, which varied from 6.5-12 pm for the 0.1N indentations, when calculated using Eq. 3.1. Less scatter was observed in the tests performed on the cross-sectional than the planar surfaces. These indentations covered multiple splat and pass thicknesses, which resulted in an averaging effect on the indentation results. The average values of hardness and elastic modulus were 50-80 percent higher on the top surface of the coatings than the cross-sections, indicating a degree of anisotropy in these coatings. As discussed in Section 2.3.2, the H/E ratio is a material index for elastic-plastic behaviour. However, no conclusive trend could been seen in the H/E ratios at 0.1 and 1.0N, Fig. 6.2-6.3, for the AA-alumina coatings. Microstructural features that can affect the results at these low loads include fine intralamellar porosity as well as microcracks, the columnar grain structure found in coatings with greater cohesive strength, i.e. AA10 and AA18 (refer to Fig. 5.10), along with any differences in phase composition that may exist. Such features are illustrated in Fig. 6.1 along with typical indentations for both nano and microindentation. S E M micrographs of several indentations are shown in Fig. 6.4 for two of the alumina PS coatings. Note the cracking and uplifting that occurred at the inter-splat boundaries in Fig. 6.4(a). This caused by the strains induced during indentation as the plastic zone expands out to the surface of the material [110]. The DSI test at 1.0 N load measured the extrinsic properties of the coating, i.e. those properties affected by the coating microstructure as a whole as opposed to the intrinsic properties of the coating material in the individual splats (which contain fewer defects), as the test volume of the indentation is typically larger than the individual splats. This applied to all the testing done on the cross-sectional surfaces, but only to the AA5 coatings on the planar surfaces. The indentation footprint was less than the average splat size for the AA10 and AA18 coatings, see Table 6.1. At a load of 1.0 N the indentation depths ranged from 1.94-2.73 pm, the diagonals from 13.6-19.1 pm and the plastic zone radius varied from 21-38 pm. Additional microstructural features that begin to affect the results at this scale include the splat dimensions, i.e. boundaries, unmelted particles, and macro-porosity or voids. These are listed in Table 1.2 together with the scale of the parameters / defects. Fig. 6.4 SEM micrographs of indentations on polished cross-sections of: (a) AA10-4 (1.0 N), (b) APS (1.0 N) and (c) APS (0.1 N) alumina coatings showing the effects of microstructural features or defects. The cross-sectional results showed increasing hardness and elastic modulus values for the coatings from all three powder sizes. This correlated with the decrease in porosity from runs 1-4 89 for all the alumina coatings shown in Fig. 5.3 and illustrated the dependence of the coating properties on microstructural defects such as porosity. Again the trend is somewhat reversed or non-existing for the results from the planar surfaces. The anisotropy of the coatings was less pronounced in the 1.0N indentation results than in the 0.1N, resulting in differences of only 10-20 % in the values of hardness and elastic modulus for the top and cross-sectional surfaces, as shown in Fig. 6.3. 6.2.2 Microindentation Conventional microhardness techniques were used with Vickers and Knoop indenters, at loads of 3 and ION respectively (10-15 measurements each). Elastic modulus values for the coatings were calculated using a technique based on the measurement of elastic recovery of the in-surface dimensions of Knoop indentations presented in Section 2.3.2. The error bars in the graphs correspond to + one standard deviation (S.D). 6.2.2.1 Vickers Microhardness Average Vickers hardness values from the top third of coating cross-sections and top surfaces of the AA-alumina coatings are summarized in Fig. 6.5. The trend of increasing hardness seen in the AA-alumina coatings illustrated the sensitivity of hardness measurements to decreases in porosity and increases in inter-splat bonding observed in the coatings caused by changes in the spray conditions, i.e. increases in heat input to the plasma and particle residence time, refer to Chap's. 4-5. The differences are especially apparent if one compares the lowest heat input, particle residence time run, #1, to the highest, #4 for each powder size. 90 25 24 22 20 18 16 £ 1 4 (3 - 1 2 10 8 6-4-2-0 H(a0r^-X-Sec AA5 A p !• s AA10 Powder AA18 26 24 22 20 18 16 £ 1 4 a - 1 2 10 8 6 4 2 0 H,(3ty-Top AA5 AA10 Powder AA18 Fig. 6.5 Comparison of Vickers microhardness values obtained on: (a) cross-sectional and (b) planar surfaces of the AA-series alumina coatings, at a load of 3.ON. (Sample symbols and numbers refer to Tables 4.2- 4.3.) Average Vickers hardness values from the top third of the reference coating cross-sections and the top surface, along with values for the bulk sintered alumina (SA) tested, are summarized in Fig. 6.6. The hardness values of the commercial alumina coatings, AL30, were higher than expected for the porosity levels. This could be due to the high proportion of unmelted particles in these coatings, which are composed of the harder oc-alumina phase, refer to Section 5.5. The chromia coatings, especially CRI, exhibited the highest hardness values after the bulk, sintered alumina. The low porosity and the lack of a visible interlamellar structure, as seen in the fracture surfaces in Fig. 5.13, in these coatings help explain these results. The results from the 3.0 N microindentations exhibited far less sensitivity to the various microstructural defects than did the corresponding DSI tests at lower loads, as test volumes was large enough to average out the effects of the smaller size defects. Larger size defects, such as large scale porosity or voids and unmelted particles can still have some influence on individual results, however. 91 26 24 22 20 18 16 £14 r i 2 6 4 2-i (a) H,(3N) 2-1 2-2 26 24-j 22 20 18 16 '14 I '12 (b) rM3N) -Top 2-1 2-2! AL30 AT CR SA AL30 CR SA Fig. 6.6 Comparison of Vickers microhardness obtained on: (a) cross-sectional and (b) planar surfaces for the reference coatings and bulk sintered alumina (SA). (Sample symbols and numbers refer to Table 4.4.) 6.2.2.2 Knoop Indentation Knoop hardness and elastic modulus were determined on both the planar and cross-sectional surfaces of the coatings at a load of 10 N. In addition, indentations on the cross-sectional surfaces were made both perpendicular and parallel to the coating surface, see Fig. 6.1, further illustrating the anisotropy of PS coatings. A comparison of the average values and S.D.'s of Knoop hardness, elastic moduli, along with the ratio H/E, for the AA-alumina coatings is summarized in Fig. 6.7, and for the reference coatings and the bulk sintered alumina, in Fig. 6.8. There were differences between the properties measured on the cross-sectional and planar coating surfaces, as well as between those indentations placed so that the long axis of the Knoop indenter was oriented perpendicular or parallel to the lamellar structure on the cross-sectional surface, see Fig. 6.1. Elastic modulus was strongly influenced by the coating microstructural anisotropy, especially when measured using Knoop indentations, see the ratios of E values measured in different orientations listed in Table 6.2. This has been observed by other researchers [39, 49], however, contrary to these results, the E-values measured on the top surfaces of the AA-alumina coatings were larger than those taken perpendicular to the substrate on the cross-sections, while those of the AL30 and AT coatings were smaller. The differences 92 between these values were not statistically significant except in a few cases, e.g. the E values for the AA5-4 and AA18-4 coatings, where the values on the top surfaces were 1.5-2.5 times greater than those on the cross-sections. The microstructure of these coatings featured considerable inter-splat grain growth and decreased porosity, as well as an increase in the density of vertical, intra-lamellar microcracks as compared to the other AA coatings. Table 6.2 Ratios of elastic moduli measured in different directions. Sample EJE\ i (x-sec) £(Top)/£x (x-sec) E(Top)/£|| (x-sec) AA5-1 1.31 1.10 1.43 5-2 1.13 0.87 0.97 5-3 0.91 1.39 1.27 5-4 0.97 1.64 1.59 AA10-1 1.41 0.82 1.15 10-2 1.25 1.08 1.35 10-3 1.52 1.17 1.78 10-4 1.24 1.34 1.66 APS 1.29 1.04 1.34 AA18-1 1.25 1.79 2.24 18-2 1.04 1.30 1.36 18-3 1.25 1.22 1.53 18-4 1.4 2.45 3.44 AL30-1 0.84 0.93 0.78 30-2 0.96 0.97 0.93 30-4 0.73 1.05 1.40 ATI 0.96 0.88 0.84 AT2 0.92 0.91 0.84 CR1 0.70 1.37 0.95 CR2-1 0.85 1.08 0.92 CR2-2 0.76 1.21 0.92 93 (a) 1 0 n Q. O 6 r 4 ft 4 * 1 ft rV (Cross-section) • 7 OJ cr> ^ U) IA U) IA < < ft - f t r - f N CO ^ o o o o < < El Hk-perpendicular • Hk-paralleJ 1 [if1] :'t| ft 4 T w T w ? 0- CO CO CO 00 < 5 " - -< a 6-Y CM CO * t in in in in < K (Top) CM CO O O O - 1 ! CM CO * (b) 3 0 0 2 5 0 2 0 0 "(? U 1 5 0 Ul 100 5 0 ii E (Cross-section) ^ N w ^ in in in i n < ft 1 E-perpendicular • E-parallel ft c\i co o o o o < < CO < 250H 3 1HH 7 « n ^ uj in u in EfTop) r- CM CO ^ o o o o ^ ,- T-< i l l N n f (c) 0.14 0 .12 0 .10 0.08 0.06 0.04 0.02 0.00 ft HIE (Cross-section) r M n ^ m in to in < < 0 rVE-perpendicular • rVE-parallel it r N rt ^ o 6 6 o ^ 1- T- T-< (/) r N « ^ 0- CO 00 00 CO < 5 - - -< 0.14-0.12-0.10-0.08-0.06-0.04-0.02-0.00-^ cii m tr) in in in < < WE (Top) Tj- CN CO 6 6 6 OD CO OO 00 Fig. 6.7 Comparison of material properties: (a) Knoop hardness (//), (b) elastic modulus (£), and (c) ///is, at 10.0 N load obtained on cross-sectional and planar surfaces for the AA-series alumina coatings. (Sample symbols and numbers refer to Tables 4.2- 4.3.) 94 (a) 18 16-14 -12 £ 1 0 H a 6 j 4 -2 -ft ft ft nb n n to < < < (b) 450 400 350 300 •n250 Q. o Uj200 150 100 50 ft to (c) 0.14 ai2 am 0.08 J a 0.06 0.04 0.02 aoo pfcft ft ft Hi T- CM <T O O O 4, (Cross-section) E3 Hk-perpendicular • Hk-parallel ft ft Ft] ft T- T- CM E (Cross-section) E3 E-perpendicular • E-parallei [ll ? N H O £C CC CJ U HIE (Cross-section) ! HCperpendicular • HEparallel . f t sift - t o 18 16 14 12 £io O | 8 6 4 r*-450 400 350 300 j 250 a. (3 uT200-| 150 100 50 X 0.14-] 0.12 0.10 0.08 J 0.06 0.00 nh rib nrrop) y- CM 5 5 ^ V ^ K H « <J CC 0C o o < cn EfTop) ft] o H/E (Top) I Fig. 6.8 Comparison of material properties: (a) Knoop hardness (H), (b) elastic modulus (E), and (c) H/E, at 10.0 N load obtained on cross-sectional and planar surfaces for the reference coatings and sintered alumina (SA). (Sample symbols and numbers refer to Table 4.4.) 95 Coatings with a more pronounced lamellar structure, i.e. the AA- and AL-series aluminas exhibited far more anisotropy than those with a denser, less lamellar structure, i.e. the AT and CR coatings, refer to Fig. 6.7- 6.8 and Table 6.2. This indicates a potential relation between the amount of coating anisotropy and the degree of interlamellar bonding. Variations in indentation morphology due to the position of the indentation and microstructural features can be seen in the SEM micrographs in Fig's. 6.9-6.10. The crack network on the top surface of the coating includes both the splat boundaries and the vertical microcracks, see Fig. 6.9 (c) and 6.10. These were seen to raise up due to the pressure of the indenter, most notably in the alumina coatings which exhibited a more distinctly lamellar structure than the AT and CR coatings, refer to Section 5.3. Fig. 6.9 SEM micrographs (70% BSE, 30% SE) of Knoop indentations both (a) perpendicular and (b) parallel to the coating cross-section, as well as (c) on the top surface of the AL30-2 coating, illustrating the relation of position to the lamellar structure. 96 Fig. 6.10 SEM micrographs (70% BSE, 30% SE) showing variations in indentation morphology due to differences in microstructure of: (a) AA5-4, (b) AA18-4, (c) AL30-2, (d) AT2, (e) CR2-1 coatings and (f) sintered alumina (SA). 97 6.2.3 Macroindentation and Fracture Toughness Testing In the present work, 100N loads were used to measure the K/c values of the coatings using two different methods. Average Vickers hardness values at 100N (5-10 measurements) from the top surfaces of the coatings are shown in Fig. 6.11. The values of elastic modulus (E) determined using Knoop microindentation, Section 6.3, were used in the calculations. Values of indentation fracture toughness were determined using Equations 3.2 and 2.5. 20 18 16 14 I  1 H a 10 A 8 6 I] rh r—*r—T T r — T — " r — T — T — T r—*T r - N P ) " t U) i ; N « ^ 7~ CM 6 6 6 6 < < < < < < < < o. co co co eb < < < < < < < < < < —r 1 1 r 1- T- OJ < S c!i cd CO O i l o o Fig. 6.11 Comparison of Vickers microhardness at 100N obtained on the top surfaces of: (a) AA-alumina coatings and (b) reference coatings and bulk sintered alumina (SA). (Sample symbols and numbers refer to Tables 4.3-4.4.) Fracture toughness values, based on indentations done on the top surfaces of the coatings, are presented in Fig. 6.12 and 6.13. The various fracture toughness equations were developed for materials with induced crack lengths (c) at least 1.8 times larger than the length of the diagonal (a), refer to Fig. 3.5. The AA5 and AL30 coatings were thus excluded as the c/a ratio measured was smaller, giving unrealistically high KJc values. As can be seen in Fig. 6.14, stress relief occurred through the extension of splat boundary and microcracks in the coatings, dissipating the energy that otherwise would drive large (median) crack extension. Indentations were attempted 98 on the cross-sections of the coatings, however, these were unsuccessful due to insufficient coating thickness. 6.0 5.0 4.0 "m a. 2.0 -1.0 0.0 O O o o K , c - T o p rfi rfi ft rfi rfi rfi CM M U CC CE O O Fig. 6.12 Type L indentation fracture toughness values for the PS coatings and the bulk sintered alumina (SA). (Eqn 3.2.) 2.0 1.0 0.0 IFT - Top ft ft rfi rfi 1*1 J , ft rfi ul rf. Fig. 6.13 Type P indentation fracture toughness values for the PS coatings and the bulk sintered alumina (SA). (Eqn. 2.5) 99 Fig. 6.14 SEM micrographs of 200 N indent on AA18-4: (a) overview and (b) close-up of splat cracking. The Type L toughness values, Fig. 6.12, for the alumina coatings were higher than for the bulk alumina (SA), while the Type P values, Fig. 6.13, were somewhat smaller. The order of the coatings are the same for both sets of results however. The fracture toughness values for the chromia coatings were higher than expected, compared to a value of 3.9 for bulk chromia [111]. These coatings did have the densest microstructures of all the coatings, along with high hardness. The ratio H/Klc is known as the index of brittleness, and has been used in the inverse form (KiJHf/r to predict the nature of the dominant wear mechanism, i.e. ductile or brittle [71]. For low values of (KJc/H)2/r fracture will be the dominant material response, while non-elastic deformation will tend to dominate at high values. A comparison of (Kc/H)2 for the evaluated materials is presented in Fig. 6.15. The value for the bulk alumina (SA) is substantially less than for the AA18-1, due to the large difference in hardness between the two. These values will help explain the wear mechanisms that occur in the different materials under abrasive and erosive wear, as will be seen later. 100 0. 0. 0, 0. . 0. o r ~ 0. 0. 0. 0. 0. 20 18 16 14 12 10 08 H 06 04 02 H 00 Ml d> 6 T- CN CO T OJ CM </> Fig. 6.15 (Klc/H) values for the PS coatings and the bulk sintered alumina (SA). 6.2.4 Summary Indentation techniques were used at loads ranging from as low as 0.1N to as high as 100N. Depth sensing indentation was performed at a load of 0.1N in an attempt to evaluate the micromechanical properties of the individual splats. The resulting indentations, and affected plastic zone, were typically smaller than the splat diameters, but not thicknesses, as well as being in the same size range as a variety of the microstructural features, or defects, as illustrated in Fig. 6.4. This affected the micromechanical property measurements by increasing the data scatter considerably. Thus, in order to sample the individual splats, an even lower load should be used, with a large number of data points to enhance confidence in the statistical analysis. The data scatter was somewhat reduced for the 1.0N results. Higher than normal E values were determined for the coatings, from 150-200 GPa, caused by an increased material yield stress due to the small size of the indentations [78]. Microindentation techniques were performed using both Knoop and Vickers indenters. Knoop hardness is a better indicator of material surface hardness than Vickers, due to the shape of the indenter. However, the measurements are more sensitive to anisotropy as well. Elastic modulus (E) values were determined using the Knoop indenter at 10N, on both planar and cross-101 sectional surfaces. In the lamellar AA-alumina coatings, the E values varied depending on the position of the indenter, as illustrated in Fig. 6.1, in the following order: planar (top) > perpendicular (x-sec) > parallel (x-sec) The differences were much smaller in the denser coatings, such as the alumina-titania and chromia coatings, which exhibited less anisotropy. Microstructural features that affected these results included macroporosity, the splat boundaries, along with the horizontal and vertical microcracks, as shown in Fig's. 6.9-6.10. Macroindentation techniques were used primarily to induce cracking to allow the determination of fracture toughness (Kic) values for the materials, using equations developed for this purpose. The induced crack length must be at least 1.8 times greater than the half-length of the diagonal, refer to Fig. 3.5, to make use of the standard equations. The more porous and poorly bonded coatings, i.e. the AA5, AL30 and APS coatings, had too high a defect density to allow sufficient crack propagation. Therefore Kjc values could not be determined for these coatings. The K/c values that were obtained were used primarily as a means of comparing coating behaviour, as when combined with hardness in the brittleness index. 6.3 Controlled Scratch Testing Scratch tests were performed on both the planar (top) and cross-sectional surfaces of the coatings, and on a bulk sintered alumina. Scratching the top surfaces was used to investigate the coating response to what is essentially a "single point abrasion test", while cross-sectional scratching was used to compare the fracture properties and microstructural integrity / cohesion of the coatings, as discussed in Section 3.3 and [112]. 6.3.1 Planar Surface Tests A commercial scratch tester [Revetest, CSEM, Switzerland] was used for these tests. A diamond indenter with a conical shape and apex angle of 120°, and a spherical tip radius of 100 pm was used. The axis of the indenter was oriented normal to the coating top surface. During 102 scratching at a speed of 10 mm/min, the normal load was either increased linearly at a rate of 100 N/min from 0-60 N or held constant at values between 10 and 60 N . An acoustic emission (AE) transducer, mounted on the side of the indenter, detected the crack and fracture energy generated by the surface during scratching. The normal load, tangential force, and A E signal counts were recorded continuously by computer during each test. Scratch tests were performed on polished coating surfaces. Samples were prepared using the procedure listed in section 5.1. The samples were cleaned with ethanol and the indenter tip wiped with ethanol before each scratch test. Afterwards, both planar- and some cross-sections of the scratches were examined by scanning electron microscopy (SEM) to more closely examine the damage mechanisms. The scratch grooves were profiled using a Talysurf profilometer, in order to measure their average depth and width. The following parameters were recorded for each coating: coefficient of friction, acoustic emission (AE) signal counts, scratch width and depth. 6.3.1.1 Scratch Results The deformation and cracking of the coating as the indenter scratches the coating surface generates noise which is picked up by the A E transducer. A typical acoustic emission signal count versus load curve for loading increased linearly from 0-60N using a 100 pm radius indenter is shown in Fig. 6.16(a). As the coatings are quite thick (250 - 350 microns on average) no major debonding of the coating from the substrate was observed. Thus, there is no critical load, as measured with thin films, for debonding of the coating. Instead, the energy of cracking, or fracture, is measured as the load on the indenter increases. Sudden increases in the AE signal can indicate a change in the damage mode, as discussed in Section 6.3.1.2. The magnitude of the fluctuations in tangential force increased with the load, Fig. 6.16 (b), as the indenter penetrated deeper into the coating. 103 The width of the scratch groove is used to determine the scratch hardness of the material, Hs, defined as: 8 ^ nb2 (6.1) where Fn is the normal load and b is the scratch width for indenter tips with a circular cross-sectional area [86]. Average values of scratch hardness and depth of the tested materials at a normal load of 60 N using a 100 pm radius indenter are presented in Fig. 6.17. The scratch depth is a measure of the plowing hardness, a measure of the energy required to deform the material, refer to Section 3.3. 50 60 Fig. 6.16 Typical (a) acoustic emission signal and (b) tangential force vs. load plots for an alumina PS coating at a linearly increasing load of 0-60 N (100 pm indenter). 104 The scratch hardness results show trends of increasing hardness for the AA-alumina coatings, Fig. 6.17(a). The differences are especially apparent if one compares the lowest heat input, particle residence time run, #1, to the highest, #4 for the 5 and 18 pm powder sizes. The 10 pm powder (AA10) derived coating was the least affected by changes in the processing conditions. (a) AA-Coatings Reference Materials a I 6 H-f nH m rin (b) AA10 Powder 14 12 10 To 8 CL a = 6- | 4 2 0 1 2 r*-i 4 2-2 35n 30-r 2 5 i ; 2o L I ; 1 5 H i ; i o rh £ 2H 2-2 CR SA SA Fig. 6.17 Comparison of (a) scratch hardness and (b) scratch depth at 60N normal load using a 100 pm radius indenter for the AA-alumina coatings and reference materials. The cross-sectional area, as measured from the scratch profiles, was used as a means of comparing the coating response as a function of indenter load (F„), see Fig. 6.18. The coatings plotted represent a variety of microstructures, from the very porous and poorly bonded AA5-1 to the dense bulk sintered alumina (SA), in order to better illustrate the trends. The amount of 105 damage, or material loss, increased at a much higher rate with increasing load for the AA5-1, APS and AL30 coatings, which have a more porous microstructure with less intersplat bonding, refer to Chapter 5, than for the denser coatings such as the optimized AA18-4 coating, or the AT and CR coatings. The least amount of material loss was exhibited by the bulk alumina (SA). The material loss was found to be directly proportional to Fn2, a result also observed in bulk ceramics [90,91]. 1 0 0 0 0 T 1 8 0 0 0 -6 0 0 0 -E a. < 4 0 0 0 -Load [N] Fig. 6.18 Scratch cross-sectional area, Ap as a function of indenter load (Fn) for a variety of coating microstructures, compared to the bulk alumina (SA). 6.3.1.2 Scratch Morphology The scratches were examined using SE imaging on the SEM to compare the differences in scratch mechanisms between the coatings. Five major scratch mechanism regimes were observed. At low loads most coatings exhibited primarily ductile behaviour (Regime 1), with plastic deformation in the groove itself, see Fig. 6.19. At higher loads, varying from around 5 to 40 N depending on the material, small scale-like features appeared in the groove bottom along with small cracks. The second damage regime was still mainly ductile, see Fig. 6.20. The cracking seen in the tracks is caused by tensile stresses induced in the coating at the trailing edge of the sliding indenter. 106 AA5-1 -AL30 - A P S A A 1 8 - 4 • A T - C R - S A M i l Fig. 6.19 SEM micrographs showing primarily ductile (Regime 1) behaviour: (a) CRI coating at 20N and (b) AA10-1 at ION. Fig. 6.20 SEM micrographs showing Regime 2 behaviour, still mostly ductile with some small scale cracking: (a) AA5-1 at ION and (b) AT at 30N. Upon further loading, material is pushed to the edges of the scratch groove where fracture occurs due to the lack of hydrostatic pressure (Fig. 6.21) along with the shear strains imposed by the indenter at the track edges causing plastic deformation, refer to Fig. 3.2 (Regime 3) . As the load increases, the scratch morphology of the more porous coatings such as AA5 and AL30 begin to differ from the rest of the coatings and the sintered alumina. In these coatings, increased loading led to the appearance of large scale-like features in the groove bottom with both small and larger cracks visible. The deformed (and compressed) layer began cracking and buckling upwards with small sections spalling off, exposing brittle fracture underneath (Regime 4) , see Fig. 6.22. Also, many sections of the deformed layer remained attached instead of leading to complete brittle fracture as seen in the sintered alumina and the denser coatings. 107 Brittle Fracture Fig. 6.22 SEM micrographs showing Regime 4 behaviour, with large scale-like features in groove bottom, and the deformed layer cracking and buckling: (a) AL30-2 and (b) AA5-1 coatings at 60N. The load required to achieve complete fracture (Regime 5) varied from 35-50N for the chromia coatings to 55-60 N for the bulk alumina (SA) and several of the alumina coatings (AA10-3,4, AA18-4, AL30-4), see Fig. 6.23. The transition from plastic deformation to complete brittle fracture did not occur in all of the coatings after single scratching, only in the CR and spray runs #3 and #4 for all of the AA-alumina coatings as well as in the bulk alumina (SA). However, even in the more porous, or less well bonded, coatings, the wear sheet observed in Regime 4, Fig. 6.22, did begin to fracture and detach from the surface after two to five repeat passes, Fig. 6.24. Repeat scratches occur on already strained material, causing additional fatigue cracking and increased wear damage. [A summary map of the damage mechanisms of all the coatings for loads 0-60N can be found in Appendix B.] 108 Fig. 6.23 S E M micrographs showing Regime 5 behaviour, exhibiting primarily brittle fracture: (a) AA10-4 and (b) SA at 60 N . Acoustic emission was able to detect the onset of brittle fracture (Regime 5) by the sudden increase in A E signal as the load was continuously increased from 0-60N, see the increase in A E signal at around 55N in Fig. 6.16(a), which corresponded with the onset of brittle fracture in that coating. However, transitions between the other regimes were not as clearly defined. Fig. 6.24 S E M micrograph of wear groove after repeated scratching (n = 5, F„ = 20N): (a) APS, (b) AA10 and (c) SA. 109 A difference in the wear modes from repeat scratches due to material hardness and porosity was also noticed. The softer, more porous coatings, see Fig. 6.24(a), exhibited a substantial attached layer of compressed wear debris, while the harder, denser coatings, Fig. 6.24 (b)-(c), exhibited primarily brittle fracture. The size and shape of the wear debris from the scratched alumina coatings, Fig. 6.25 (a)-(b), correlate with the thickness and distance between vertical cracks of the splats in the coatings, refer to Sections 5.3-5.4. The chromia coatings exhibited a similar debris mechanism, see Fig. 6.25(c). These coatings did not exhibit much of a lamellar structure, i.e. the splats were not visible even on the fracture surfaces, Fig. 5.13, however under the force of an indenter the debris particles or platelets still fractured along splat boundaries and pre-existing solidification cracks. In contrast, the wear debris generating process in the tested bulk alumina involved fracture or shearing at the grain boundaries, see Fig. 6.25(d). Pre-existing solidification cracks . .. Splat Cross-section Incipient wear debris particle Intergranular fracture Fig. 6.25 Examples of wear debris from: (a) AA10-4, (b) AL30 and (c) CR1 coatings along with (d) bulk alumina (SA), illustrating differences in wear debris morphology. 110 6.3.1.3 Subsurface Cracking The type of crack system that formed under the scratch groove, as examined through SEM studies of polished cross-sections, was greatly affected by the microstructure of the individual coatings tested, see Fig. 6.26. Those coatings with a 'poor' microstructure, i.e. high porosity, unmelted particles and poor inter-splat bonding, such as AL30-2, Fig. 6.26(a), showed a pronounced compression layer formed at the higher scratch loads along with a series of cracks formed underneath and at the edges of the scratch groove. The tensile forces at the edges of the groove help direct the crack growth as well as causing microfracture at the surface along the sides of the scratch track. The denser AA-alumina coatings exhibited more fracture at the track edges with less compressive deformation directly underneath the indenter, Fig. 6.26(b), while macrocracks were visible in the AT coatings, Fig. 6.26(c). Fig. 6.26 SEM micrographs of single scratch cross-sections: (a) AL30-2, (b) AA18-4 and (c) AT2 at 50 N, illustrating the effects of different microstructures. Ill After repeated scratching (number of repetitions, n = 5) the more porous coatings, such as AA5-1, exhibited the formation of a densification layer directly under the indenter and brittle fracture at the edges of the groove, Fig. 6.27. Macrocracking was observed in some of the more cohesive coatings that had increased trans-lamellar grain growth (refer to Fig. 5.10 and 5.13) such as AA18-4 and AT, Fig. 6.28, but not in any of the CR coatings. The vertical cracks seen in the AA18 coating, Fig. 6.28(b) appear to follow the microcracks and splat boundaries present in a plasma sprayed coating in a stepwise fashion. Fig. 6.27 SEM micrographs of AA5-1 repeat scratches (n = 5, Fn = 20N): (a) top surface, (b) cross-section, showing the compressed layer, and (c) close-up showing densification and cracking under the scratch groove. 112 Fig. 6.28 Cross-sections of scratches (n = 5,Fn = 20N): AA18-4 (a) overview and (b) close-up of vertical macrocracks and (c) AT2 overview and (d) close-up of horizontal macrocracks. 6.3.1.4 Summary To summarize, the unique lamellar structure of ceramic PS coatings has a strong effect on the types of scratch damage mechanisms seen and on the ductile to brittle transition load. As shown in Section 6.3.1.2, the transition from ductile to brittle wear mechanisms depends a great deal on the microstructure of the coatings. In the more porous coatings, the dominant mechanism seen was plastic deformation and densification (AA5-1, APS and AL30's), with the formation of sheets of compressed wear debris particles, Fig. 6.27. The extensive microcracking and porosity seen in these coatings relieve the stresses that would otherwise cause the formation of macrocracks. The denser coatings, with a greater amount of interlamellar bonding, (such as AA10-3,4 and AA 18-3,4, and CR) exhibited some brittle fracture at loads greater than 40N, refer to Section 7.2.1.2. The alumina titania (AT) coatings exhibited small amounts of brittle fracture along the 113 edges of the scratch grooves, without the occurance of complete fracture at loads up to 60N, while the chromia (CR) coatings exhibited total fracture at a load of 35 N . Alumina titania is known for having a higher proportion of chemical bonding between lamellae, than pure alumina coatings [107] which could account for the small amount of damage caused by scratching. Macrocracking was observed in all of these coatings, except for the chromias, especially after repeated scratches in the same groove, Fig. 6.26 and 6.28. The dominant material removal mechanism was seen to occur on the microscale instead of the macroscale, however. The scratch wear debris process in ceramic plasma sprayed coatings involves microfracture along the splat boundaries and pre-existing solidification cracks, instead of fracture or shearing at the grain boundaries common to bulk ceramics, Fig. 6.25. This is seen by examining the size and shape of the wear debris, especially from the AA-alumina coatings, which correlate with the thickness and distance between vertical cracks in the splats observed in sections 5.3-5.4. 6.3.2 Cross-Sectional Surface Scratch Tests Polished cross-sections of the coatings were scratched, sliding from the substrate towards the coating top, with a Vickers diamond in low-load scratch test equipment, installed inside a JEOL 25 S scanning electron microscope, Fig. 6.29, at constant loads of 1, 2, and 3.5 N , at a speed of 0.2 mm/min. For further details about the low load scratch testing equipment, see Ref. [113]. Three scratches were made at each of the loads, while measuring both normal, Fn, and tangential, F,, forces. A typical force profile is shown in Fig. 6.30, along with a corresponding micrograph of a typical scratch. 114 Indenter Fig. 6.29 Photographs of low-load scratch test equipment, mounted inside a SEM: (a) outside view and (b) close-up of scratch indentor setup. (a) (b) 4.00 3.00 •g 2.00 o 1.00 0.00 -F, A' 1 1 1 1 1 50 100 150 200 Length [um] 250 300 350 Fig. 6.30 (a) Normal and tangential forces, Fn and Ff, plotted against scratch length, with line A' illustrating the boundary between substrate and coating and (b) SEM micrograph of 3.5 N scratch with fracture cone on cross-section of a typical alumina coating. 115 6.3.2.1 Scratch Morphology The scratch grooves and fracture cones were examined using SEM to compare differences in morphology and scratch mechanisms. During the scratch process, the indenter tip proceeds from the substrate across the coating, until breaking over its edge, Fig. 6.30(b). The fracture cone forms when cracks caused by the stress field in front of the indentor reach the top surface of the coating. The morphology of the scratch tracks on the cross-sections depended upon coating characteristics such as hardness and porosity. The indenter in this test was sharp-pointed, inducing more severe contact damage than in the planar scratch tests. Thus, the harder, denser coatings, such as the AA10, AA18 and CR coatings, showed little evidence of plastic f low, but considerable fracture, even at the low loads used, as shown in Fig. 6.31 (c)-(d). Fig. 6.31 SEM micrographs of scratch tracks at a load of 3.5N illustrating different damage mechanism regimes: (a) A T coating, showing ductile plowing, (b) APS coating, showing ductile plowing, tensile cracking and edge fracture, (c) CR and (d) AA10 coatings, showing primarily brittle fracture. 116 The softer, more porous coatings, ABO and APS, exhibited mainly plastic deformation with fracture particles / wear debris along the edges of the tracks, see Fig. 6.31(b), while the AT coatings exhibited primarily ductile plowing with a small amount of fracture occurring at the track edges, Fig. 6.31(a). The lamellar structure of the alumina and AT coatings was revealed by the stepped morphology seen in the fracture cones, Fig. 6.32(a), which was absent from the relatively smooth surface of the CR coatings, Fig. 6.32(b). Lamellar structure Fig. 6.32 SEM micrographs of fracture cones on cross-sections of (a) CR coating, showing smooth brittle fracture with no clear sign of a lamellar structure and (b) AA10 coating, showing fracture occurring in shear steps revealing the coating lamellar structure. 6.3.2.2 Fracture cone lengths A comparison of the cone lengths at a load of 3.5N for all the coatings, along with sintered alumina as a reference, is shown in Fig. 6.33. The trends in the fracture cone lengths correlated inversely with those noted for several of the measured microstructural parameters, including porosity, Fig.'s 5.3-5.6 as well as microhardness, Section 6.2.2, indicating its potential use as a method of evaluating coating microstructural integrity. 117 120-, (a) 100 E 80 3, 9. 60 20 rh rn 100 E 80 3 5 60 (b) Ht-i T l 1 1 2 2 4 0-1+1 2-1 2-2 S 4 AA5 AA10 Powder AL30 AT CR SA Fig. 6.33 Comparison of fracture cone lengths at a load of 3.5N for the: (a) AA-alumina coatings and (b) reference materials. As discussed in Section 3.3.2, the fracture cone length, or critical length, was found to correlate with the normal load according to Equation 3.4. The slope of the linear relation measured on the Lc 3 / 2 vs. Fn curves, similar to Fig. 6.34, was therefore A, while the constant a„ was determined by using sintered alumina, with a Kc value of 4.8 MN m" . Using such values, Beltzung, et al [94] calculated values of fracture toughness for comparison of their coatings. In the present work, a linear relation was obtained only for the bulk sintered alumina (SA) and the harder, denser coatings such as CR, AT and the optimized AA10 and AA18 coatings (spray run #4), see Fig. 6.34. The coatings that did not correlate linearly with equation 3.4 are also the ones that were not sufficiently dense to be able to obtain fracture toughness values from indentation methods, i.e. the AA5, AL30 and APS coatings, Section 6.2.3. Estimates of the toughness values using this method are presented in Fig. 6.35. The calculated values for many of the coatings were higher than expected. The assumptions that the fracture of the half cone is equivalent to that of a half penny crack, as well as the use of sintered alumina, with its isotropic and equiaxed microstructure, as a calibration material for the highly anisotropic, lamellar microstructure of PS coatings do not seem to be appropriate. I 18 1200 x 1000 + s E a 800 •• n O - I lance 600 - -m a itical 400 •• o 200 + 2 Load [N] Fig. 6.34 Critical distance or cone length, Lc, raised to the 3/2 power as a function of indenter load for representative coatings and bulk alumina (SA). r « rt * O O O O Fig. 6.35 Fracture toughness values estimated using the fracture cone method. Scratching across the polished cross-section of a ceramic PS coating with a pointed tip (Vickers) diamond indenter forms a fracture cone as the indenter breaks over the top edge of the coating. The indentor is scratching across the splat cross-sections as opposed to their top surfaces, thus the strength of the inter-splat boundaries or inter-lamellar porosity of the coatings is being tested. The length of the cone is affected by the microstructural integrity of the coatings, as can be seen by the comparison of cone lengths in Fig. 6.29. The larger the cone length, the lower the internal cohesive strength of the coating. The fracture morphology of the cone itself 119 illustrates the differences in coating structure, refer to Fig. 6.28. The coatings with a distinct lamellar structure fracture in discrete steps between the splat boundaries and the vertical microcracks (all the alumina and AT coatings), while the coatings without a distinct lamellar structure, such as the chromias, exhibit a much smoother fracture surface. The results showed that the fracture cone scratch test is sensitive to differences in composition and structure and has potential for use in studying the effects of microstructure on the internal toughness or cohesion of coatings. 120 C H A P T E R 7 W E A R M E T H O D S 7.1 Introduction In order to better understand the relationships between the different microstructural parameters found in PS ceramic coatings, detailed in Chapter 2, several test methods involving wear by hard particles were used. These included 2-body abrasion and erosion by hard particles. This chapter presents the experimental methods used along with the results and discussion of these tests for evaluating the responses of a variety of coating microstructures as compared to those of a bulk sintered alumina. Additional details on the use of these methods on ceramic PS coatings can be found in references [108, 112, 114]. 7.2 Abrasion Tests A pin-on-drum tribometer was used to evaluate the abrasive wear properties of the materials [1]. A rotating drum (196 mm in diameter) was covered with silicon carbide (2500 HV) abrasive papers with mesh grades 400 and 80, corresponding to average grit sizes of 20 and 200 pm, respectively. The specimens, in the form of 5 mm wide beams, were pressed with the coated surface radially against the drum surface with a force of 12.8 N , refer to schematic in Fig. 7.1. The sliding speed was 0.26 m/s. To ensure that the specimens were always tested against fresh abrasives, they were fed along the drum parallel to the axis of rotation. A l l wear results represent the averages of two measurements. The scatter is estimated to be less than 10% of the quoted wear rate values. The specific wear rates were calculated using the equation: where V is the volume removed [mm ] (equal to mlp, where m is the removed mass [g] and p is the density [g/mnr ] of the materials, estimated from the porosity values), S the sliding distance [m], Fn the normal force [N] and KW the specific wear rate [mm3/Nm]. 121 Coated Specimen Abrasive Paper Fig. 7.1 Schematic of block-on-drum test equipment (Adapted from [115].) 7.2.1 A A-Alumina Coatings The abrasive wear rates of the sapphire powder derived alumina PS coatings (AA-alumina series) ranged from 0.015 to 0.044 mm 3/Nm (50 to 180 pg/Nm) for 20 pm and from 0.058 to 0.125 mm 3/Nm (200 to 760 pg/Nm) for 200 pm abrasives, see Fig. 7.2. The mean wear rates thus increased by a factor of 3-4 when the size of the abrasive increased from 20 to 200 pm. The relative order of the wear rates of the different coatings did not change significantly with the size of the abrasives. The effect of the precursor powder size on the abrasive wear resistance of the different coatings can be seen by comparing the mean abrasive wear rates for each powder size. The mean wear rates for the optimized coatings, run #4 for each, are almost independent of the size of the precursor powder. The AA18-4 coating does show a slightly improved wear resistance, i.e. lower wear rate, than that of the AA5-4 or AA10-4 coatings in the small size abrasive (20 pm) wear test, Fig. 7.2(a). 122 0.06 T (a) 0.05 + I 0.04 -j-£ 0.03 f a CC I 0.02 H 0.01 + A A 5 AA10 Powder AA18 0.3 0.25 4-| 0.2 H 0.15 + | 0.1 H 0.05 (b) AA5 AA10 Powder AA18 Fig. 7.2 Comparison of specific wear rates for (a) 20 pm and (b) 200 pm abrasion of the AA-alumina coatings as a function of powder type and processing conditions. (Sample symbols and numbers refer to Tables 4.2- 4.3.) The variations in the spray processing conditions had a larger effect on the wear resistance of the coatings. One method of examining the effects of the changes in nozzle size and the plasma gas composition, as well as any interaction effects, is factorial design [116]. The set of conditions used for spraying the AA-aluminas, refer to Tables 4.2 and 4.3, is essentially an example of a 2-factor factorial design. The percentage of hydrogen used in the plasma gas, i.e. the heat input into the plasma, is one factor and the size of the torch nozzle the other. The effect of these two factors and any interaction between them can be illustrated by plotting the abrasion resistance of the coatings against the percent hydrogen (factor A), Fig. 7.3 for the 20 pm abrasion results. The details of the spray conditions for the 4 runs are listed in Tables 4.2-4.3. 123 The abrasive wear resistance, i.e. the inverse of the abrasive wear rate, of almost all of the coatings was increased as a consequence of the improvement in coating microstructure, i.e. a decrease in porosity, increase in intersplat bonding, etc. as discussed in Chapter 5, regardless of abrasive particle size. This was caused by the increase in hydrogen concentration in the plasma gas, which resulted in an increase in heat input to the plasma, Fig. 7.3. The increase in torch nozzle size, which resulted in an increased residence time for the particles in the plasma, had the largest effect on the AA5 coatings, followed by the AA18 coatings for the Fi2-rich gas, with little effect on the AA10 coatings, Fig. 7.3(b). AA5 AA10 70 — 6 0 i» Z V O40 ra « S30 CC LZ o 5 20 CO < 10 0 (a) 70 60 "E f . Z <D u 40 « « 'I 30 cc c o S 2 0 .fi < 10 0 (b) •*• Large Nozzle -•-Small Nozzle 20%H 10%H 70-j if 60-E 1 50-z o u c 40-ca w 'co CD 30-CC c o 20-'co « < 10-0-(c) AA18 10%H -Large Nozzle -Small Nozzle 20%H Fig. 7.3 Factorial graph for 20 pm abrasion of: (a) A A 5 , (b) AA10 and (c) AA18 alumina coatings. (Numbers correspond to coating runs as listed in Tables 4.2 - 4.3.) The factorial graphs for both the AA5 and AA10 coatings after 20 pm abrasion, Fig. 7.3(a)-(b), and for the AA5 coatings after 200 pm abrasion, Fig. 7.4(a), show no evidence of 124 interaction between the two factors, i.e. the lines do not cross. The factorial graphs for the AA18 after 20 pm abrasion, Fig. 7.3(c), as well as the AA10 and AA18 coatings after 200 pm abrasion, Fig. 7.4(b)-(c), show some evidence of interaction between the nozzle size and H2 content. The error bars on the graphs correspond to + one standard deviation (S.D.). AA5 AAIO 20 18 " | l 6 !14 0 " 110 (A £ 8 1 e (A 1 * < 2 0 (a) -•-Large Nozzle Small Nozzle 20 18 E 1 6 I14 ® 11 0 12 c ro •510 W * 8 1 6 4 2 0 co n < (b) -Large Nozzle -Small l*Jozzle 20-1 18-"E E 16-[Nm/ 14-a> 0 12-c .2 10-'« a> 8-cc c 0 6-'to ro 4-< 2-0-(c) AA18 -Large Nozzle -Small Nozzle 10%H 20%H Fig. 7.4 Factorial graph for 200 pm abrasion of: (a) A A 5 , (b) AA10 and (c) AA18 alumina coatings. (Numbers correspond to coating runs as listed in Tables 4.2 - 4.3.) 7.2.2 Reference Materials A comparison of the abrasive wear rates of the commercial alumina (A130). alumina-titania (AT) and chromia (CR) coatings, along with the bulk sintered alumina, is shown in Fig. 7.5. The mean abrasive wear rates of the ABO and CR2 coatings increased by a factor of 4.0 to 125 4.5 when the size of the abrasive increased from 20 to 200 um, regardless of the processing conditions. The mean abrasive wear rates of the AT and CR1 coatings increased by a factor of only 2.5 to 3.0, while those of the sintered alumina (SA) increased by a factor of only 2.3. 0.06 | a o 4 ,0.03 ; 0.02 o.oi (a) AL30 r 4 -2-1 -*-1 2 2-2 SA CR Powder SA 0.3 025 an (b) AL30 2-1 -t-I 1 2 1 2-2 I—1 AT CR SA Powder Fig. 7.5 Comparison of specific wear rates for (a) 20 pm and (b) 200 pm abrasion of the reference coatings plus sintered alumina (SA) as a function of powder type and processing conditions. (Sample symbols and numbers refer to Tables 4.3- 4.4.) The AL30 coatings showed a definite improvement in wear resistance with increases in both plasma heat input and particle residence time, as there was an almost 50% reduction in abrasive wear rate when comparing coating AL30-1 to AL30-4, regardless of the abrasive particle size. The AL30 powder, with an average particle size of 30 pm, required extra heating time and input to melt these large volume particles. The wear rate of the AL30-4 coating was 50% and 100% higher than the plasma spray run #4 AA-alumina coatings, i.e. AA5-4, AA10-4 & AA18-4, for the small and large size abrasive particles, respectively. The CR2 chromia coatings were also affected by a change in processing conditions. Coating CR2-2, sprayed with a larger torch nozzle size, showed a decrease in wear rate of 20% for both abrasive particle sizes compared to the CR2-1 coating. The smaller particle size of the CR1 coating, ranging from 5-20 pm versus 20-45 pm for the CR2 coatings, could have resulted 126 in more complete melting of the powder particles and a denser coating, thus giving a wear resistance nearly equal to that of the sintered alumina. The A T coatings had wear rates 30 and 50% lower than the best (run #4) AA-alumina coatings for the small and large size abrasive particles, respectively, while the CRI coating and the sintered alumina, SA, had wear rates between 70 and 80% lower for both abrasive sizes. 7.2.3 Effects of Particle Hardness As discussed in Section 3.4, as the ratio of abrasive hardness to target surface hardness, HJHs, decreases, the major wear mechanism changes from plastic deformation to fracture [51]. The abrasive wear rates of the materials tested in this work are plotted against HJHS, with Ha (SiC) = 2500 HV, in Fig. 7.6 for both the 20 and 200 pm abrasive tests. The bulk alumina and the CRI coating were similar in having the highest hardness as well as the best wear performance, i.e. the lowest abrasion rates, Fig. 7.6. The amount of wear shown by the three AL30 coatings, Fig. 7.6, was much higher than expected based on coating hardness. The poor intersplat bonding, high porosity and unmelted particles present in these coatings, refer to Chapter 5, help explain the higher wear rates. 127 0.06 (a) 0.05 H 0.04 H 13 0.03 DC c o 'tn _ n 0.02 0.01 0.00 • AL30-1 O APS AA5-1 • • AL30-2 CR2 • • * • • • A • • AT o SA • CR • A T • AL30 o APS A AA5 AA10/18 SA CR1 h-0.00 0.50 1.00 1.50 2.00 2.50 3.00 3.50 4.00 (b) 0.25 0.2 0.15 0.1 0.05 DAL30-1 • AL30-2 • AL30-4 <>APS • • AA5-1 CR2 • • A T CR1 O S A i 1 r-0.00 0.50 1.00 1.50 2.00 2.50 3.00 3.50 4.00 rVHc Fig. 7.6 Relative abrasive wear rates for: (a) 20 urn and (b) 200 um abrasion, plotted against the ratio of the abrasive hardness, Ha, to the surface hardness, Hs, for the evaluated materials. 7.2.4 Abrasive Wear Mechanisms Three main mechanisms are observed in the abrasive wear of brittle materials, namely plastic deformation, brittle (indentation) fracture and delamination of the material beneath the abrasive wear grooves [51]. Plastic deformation is favored when the load on the individual abrasive particles is small, i.e at small particle sizes or low loads, when the abrasive is blunt in shape rather than angular, and at a low H/Kc ratio, the index of brittleness. Indentation fracture occurs when the load on the individual abrasive particles is high, i.e. for large particle sizes or high loads, when the abrasive is angular in shape, and when the brittleness index is high. 128 Typically, some combination of both of these wear mechanisms is seen. Delamination of the worn surface, i.e. of the material beneath the grooves made by the abrasive particles, is affected by microstructural features such as the strength of the grain boundaries, along with the grain size. The wear mechanisms observed on the ceramic PS coatings abraded with 20 pm particles included both plastic deformation and brittle fracture, as well as splat debonding, i.e. cracking and separation occurring at the splat interfaces, Fig. 7.7. The predominance of one mechanism over the other depends on the microstructure, as well as on the size of the abrasive particles. The size of the abrasive indentation, or scratch groove, in combination with the local fracture toughness of the surface of the wearing material, determine whether the material removal mechanism occurs as plastic deformation or includes microcracking. Microcracking occurs when the stresses imposed by the abrasive indentor are sufficiently high, i.e. greater than the critical surface pressure necessary to induce cracking [71]. The crack formation and propagation that occurs during brittle fracture can involve cracks that extend beyond the area of the indentation, causing the formation, and subsequent loss, of large wear debris particles. The AA-alumina coatings exhibited both plastic deformation and splat fracture, along with some splat debonding in the AA5 coatings. The effects due to the variations in spraying conditions were seen most clearly in the A A 5 coatings, especially between AA5-1 (sprayed at the lowest plasma heat content and particle residence time) and AA5-4 (sprayed at the highest plasma heat content and particle residence time). They also showed the largest tendency towards splat delamination, or pull-out, due to the small splat sizes, see Fig. 7.7 (a)-(b). The AA10 and AA18 coatings exhibited similar wear modes, primarily plastic deformation and some splat fracture. The APS coating showed some splat debonding as well. 129 Splat particle pull-outs or Microfracture Fig. 7.7 SEM micrographs of 20 pm abraded surfaces: (a) AA5-1, (b) AA5-4, (c)AA10/18* (same mechanism for both AA10 and AA18 coatings), (d) APS, (e) AL30-2 (f) AT and (g) CR coatings, showing plastic deformation, some brittle fracture and splat debonding, or pullouts. The AL30 coatings exhibited the greatest amount of splat debonding of all the coatings, in addition to plastic deformation and splat fracture, see Fig. 7.7(e). The CRI exhibited the least 130 wear of all the coatings, similar to that of the sintered alumina (Fig. 7.5), with the primary mechanism being plastic deformation and plowing by the abrasive particles, see Fig. 7.7(g). Microfracture Fig. 7.8 SEM micrographs of 200 pm abraded surfaces: (a) AA5, (b) AA10/18* (same mechanism for both AA10 and AA18 coatings), (c) AL30, (d) AT and (f) CR coatings, showing plastic deformation, brittle (splat) fracture and splat delamination. The wear mechanisms observed on the ceramic PS coatings abraded with 200 pm particles included plastic deformation and brittle fracture, as well as splat delamination, Fig. 7.8. The increase in size of the abrasive particles, by a factor of ten, increased the amount of wear by 131 a factor of 2.5 - 4.0, see Fig. 7.2 and 7.5, due to the increased contact stresses imposed on the individual abrasive particles and, consequently, on the surface of the wear specimen. This can be seen in the decreased frequency and increased width of the plastically deformed grooves of the AT and CR coatings, see Fig. 7.8(d)-(e) compared to the same coatings worn by finer abrasive particles, Fig. 7.7(f)-(g). The larger the abrasive particles, the fewer there are per contact area, and thus the greater the load on each abrasive particle, or indenter, contacting the wearing surface. The sapphire and commercial alumina powder derived coatings exhibited a greater amount of brittle (splat) fracture than was seen with the smaller abrasive particle size. Again, the AA5 coatings had small craters scattered across the wear surface, where a number of small particles or splats were pulled out, see Fig. 7.8(a). The AL30 coatings showed the largest amount of splat debonding, see Fig. 7.8(c). Higher magnification SEM micrographs of both of these coatings are presented in Fig. 7.9, showing the unmarked surfaces of the poorly bonded splats, as well as the beginnings of wear debris formed by fracture of the lamellar layers. Fig. 7.9 Higher magnification SEM micrographs of 200 pm abraded surfaces showing the mechanisms of splat debonding and wear debris formation: (a) AA5 and (b) AL30 coatings. Abrasion can be thought of as the action of a multitude of sliding indenters which form grooves across the wearing material. The interaction of these grooves causes material fatigue 132 which induces crack formation and growth. In ceramic PS coatings these propagate along existing crack pathways such as the vertical and horizontal, or interlamellar, crack networks. In coatings with poor cohesion, such as the A A 5 , APS and AL30 coatings, entire splats as well as splat fragments were delaminated, Fig. 7.7-7.8. 7.3 Erosion Tests Solid particle erosion tests were performed in a centrifugal erosion tester in which up to 18 specimens can be eroded simultaneously under identical testing conditions, see Fig. 7.10. A detailed description of the testing equipment is given in ref. [117]. Silicon carbide particles with a hardness of 2500HV with mean diameters of 600 pm, 200 pm and 75 pm were used as erodents. The particle velocity was 70 m/s in all tests and two angles of impingement were used, 45 and 90 degrees. The area exposed to erosion on each specimen was 8.5 x 8.5 mm . A specially designed container attached onto a speciman holder was used to collect erosive particles in situ. These were then weighed to give the effective particle dosage. The weight loss of the specimens were plotted against the weight of the erodent impinging on them. The erosion rate is the calculated slope of the line, obtained through linear regression analysis of the data, see Fig. 7.11 for an example. The scatter is estimated to be less than 5% of the quoted erosion rate values. The number of impacting particles was also estimated for each erodent size. To enable a simple quantification of the damage caused by individual erodent impacts, the erosion rates are also presented as removed mass per impact (ng/impact). These values were calculated from the sample mass loss and the total number of particles hitting each sample during the test. The morphology of the eroded specimens was investigated using S E M . The wear mechanisms were studied in greater detail by examining polished cross-sections of several of the eroded specimens, refer to Appendix A for the mounting and polishing procedures. 133 Fig. 7.10 (a) Outside and (b) inside views and (c) schematic of the test equipment used to provide solid particle erosion. The particles are fed into the center of the rapidly rotating disk where they are accelerated through four radial channels to impact the specimens. (Adapted from [117].) 134 0.025 A A 1 0 - 3 - 80 Lim a 0.010 -J <3 3 cn g 0.015-^  0.020 A 0.005 A y = 0.00042X + 0.00361 0.000 0 10 15 20 25 30 35 40 45 Wt of Erodent [g] Fig. 7.11 Determination of the erosion rate as the slope of the line through the specimen weight loss vs. weight of the erodent data points. 7.3.1 AA-Alumina Coatings The erosive wear rates at 45° incidence were found to range from 0.17 to 0.24 mg/g (0.12 to 0.18 ng/impact) for the 75 pm SiC erodents, from 0.27 to 0.59 mg/g (3.8 to 7.9 ng/impact) for the 200 pm and from 0.53 to 1.21 mg/g (200 to 430 ng/impact) for the 600 pm erodent particles, Fig. 7.12. For a comparison of the number of splats removed per impact, the typical mass of a splat in these coatings was estimated to be on the order of 0.2 ng for the A A 5 , 2.0 ng for the AA10 and 10.0 ng for the AA18 coatings. Thus the 75 pm erodents removed, on average, less than one splat per impact on all the coatings, while the 200 pm erodents removed the equivalent of more than one splat per impact for the AA5 and AA10 coatings but not the AA18. The impact of the 600 pm erodent particles at 45 degrees removed the equivalent of more than one splat per impact on all the AA-alumina coatings. The mean wear rates, both mass- and volume-based, increased by a factor of 1.5 when the erodent size was increased from 75 to 200 pm and roughly doubled from 200 to 600 pm, Fig. 7.13. The volume wear rate is also presented in order to illustrate the differences in wear caused by material density, Fig. 7.13(b). This is not noticeable in the AA-alumina coatings, but will be more significant with the reference materials. These plots were drawn using the -1 and -4 135 coatings of the three powders, along with APS. The weight loss per impact (ng/impact) increased by a factor of about 20 when the erodent size was increased from 75 to 200 pm and by a factor of about 50 for an increase in erodent size from 200 to 600 pm, as a consequence of the greater impact force, and resulting wear loss, caused by the larger erodent particles. Note the larger increase in erosion rate when the erodent particle size increased from 75 to 200 pm for the AA5-1 coating. The poor bonding and high porosity in this coating, resulting in particle pullout, reduced the critical erodent particle size necessary for the onset of severe wear. 1.0-0.9 -0.8 -0.7-I 0.6-0 I 0.5 c o » 0.4 0.3 0.2 0.1 0.0 • (a) LU 3 AA10 Powder 0.4 cn 1.0 0.9 0.8 0.7 I I 0.6 £ 0.5 c o » 0.4 o ui 0.3 0.2 0.1 H 0.0 (b) nh |+| "*1 —r • 2 —r-3 4 1 2 3 4 rh 1 r i i AA10 Powder i-» -j ( 1.6 -1.4 55 1-2 )^ E S 10 <5 CC o uu 0.6 0.4 0.2 0.0 fh ff rh i - i . - ^ 600 400 £ s AA10 Powder Fig. 7.12 Comparison of wear rates for (a) 75 pm, (b) 200 pm and (c) 600 pm erosion using a 45 degree angle of incidence of the AA-alumina coatings as a function of powder type and processing conditions. (Sample symbols and numbers refer to Tables 4.2 and 4.3.) 136 1.8 -i 1.6 -1.4 -0 -I 1 1 1 i 1 1 1 0 100 200 300 400 500 600 700 Particle Size [nm] 0.6 -| 0.5 -0 A , 1 1 , 1 1 , 0 100 200 300 400 500 600 700 Particle Size [nm] Fig. 7.13 (a) Mass and (b) volume erosive wear rates, at a 45 degree impingement angle, plotted as a function of erodent particle size for several AA-alumina coatings. The erosive wear rates of the PS coatings at normal incidence were found to range from 1.3 to 1.7 mg/g (0.9 to 1.2 ng/impact) for the 75 pm SiC erodents, from 1.1 to 2.3 mg/g (16 to 30 ng/impact) for the 200 pm and from 2.5 to 4.9 mg/g (800 to 1750 ng/impact) for the 600 pm erodent particles, Fig. 7.14. Therefore the 75 pm erodent particles removed less than one splat per impact on all the coatings, while both the 200 and 600 pm erodent particles removed the equivalent of more than one splat per 90 degree impact on all the AA-alumina coatings. 137 (a) T 5.0 r on AA10 Powder (c) 7.01 — 5.0 H o. cn cc .2 a 0 o 1 0 2.0 1.0 ao (b) AA5 AA10 Powder T 2500 rn-H rh U _ L M nh 90 80 70 tJ i a. 60 | 5 c 50 « CD CD 40 0 1 c o 30 g I 23 10 0 UJ AA18 Fig. 7.14 Comparison of wear rates for (a) 75 pm, (b) 200 pm and (c) 600 pm erosion using a 90 degree angle of incidence of the AA-alumina coatings as a function of powder type and processing conditions. (Sample symbols and numbers refer to Tables 4.2 and 4.3.) The mean wear rates increased by a factor of only 1.13 when the erodent size was increased from 75 to 200 pm and by 2.2 for an increase in erodent size from 200 to 600 pm, Fig. 7.15. The weight loss per impact (ng/impact) increased by a factor of about 20 when the erodent size was increased from 75 to 200 pm and by a factor of about 50 for an increase in erodent size from 200 to 600 pm, again illustrating the greater impact force, and resulting wear, caused by the larger erodent particles. The effect of the precursor powder size on the erosive wear resistance of the different coatings can be examined by comparing the mean erosive wear rates for each powder size. The mean wear rates, at both angles of impingement, for the optimized powders, coatings AA5-4, 138 AA10-4 and AA18-4, were almost independent of the size of the precursor powder, compare Fig. 7.12 and 7.14. (a) (b) „ 5 cn E a 4 • S 3 2H 1.8 1.6 •S1-4 1l-2 <; 1 c .1 0.8 H It) o lu 0.6 0.4 0.2 100 200 300 400 Particle Size [um] 500 600 700 0 100 200 300 400 500 600 700 Particle Size [um] Fig. 7.15 (a) Mass and (b) volume erosive wear rates, at a 90 degree impingement angle, plotted as a function of erodent particle size for several AA-alumina coatings. The effect of the incidence angle is seen in the increase in wear rates when going from 45 to 90 degrees. As shown in Fig. 3.10, brittle materials like ceramics experience maximum erosion at normal (90°) impingement. For the AA-alumina coatings, increasing the impingement angle from 45 to 90° caused the erosive wear rates to increase by a factor of roughly 7.5 for the 75 pm erodent, and by a factor of 4 for both the 200 and 600 pm erodents. Variations in the spray processing conditions have a large effect on the wear resistance of the coatings. As detailed 139 in Section 7.2 on abrasion, the effect of these two factors and any interaction between them can be illustrated using factorial design methods by plotting the erosion resistance of the coatings against the percent hydrogen (factor A), for the two nozzle sizes. The details of the spray conditions for the 4 runs are listed in Tables 4.2 - 4.3. The results for all three sizes of erodent were quite similar, therefore only the 75 pm erosion results at both 45 and 90 degree impingement angles are presented, Fig. 7.16 and 7.17. AA5 AA10 7 l E i 5 o> « .2 'tn o 0) o OC c .2 2 tn O u i 1 (a) -•-Large Nozzle -•-Small Nozzle 7-n E O |4H 83i cc °2-\ (b) •*• Large Nozzle -•-Small Nozzle 10%H 20°/<H 10%H 20%H 7 6 "E CC c o w LU AA18 H 0 (C) 10% Hydrogen -Large Nozzle -Small Nozzle 20% Hydrogen Fig. 7.16 Factorial graph for 75 pm erosion, at 45 degrees, of: (a) A A 5 , (b) AA10 and (c) AA18 alumina coatings. (Numbers correspond to coating runs as listed in Tables 4.2 - 4.3.) The erosion resistance, i.e. the inverse of the erosive wear rate, of the majority of the coatings tested at 90°, and to a lesser extent at 45°, was improved slightly due to the increase in hydrogen concentration in the plasma gas as well as the increase in nozzle size, Fig. 7.16 and 140 7.17. The factorial graphs for the A A coatings showed some interaction between the nozzle size and H 2 content. The error bars on the graphs correspond to + one standard deviation (S.D.). AA5 AA10 E0.8 E 80.6 !0.4 2 0,2 (a) •*• Large Nozzle Small Nozzle 1n ,0.6 A '.OA A 0.2 A (b) -Large Nozzle -Small Nozzle 20%H 10°/<H 20V.H 0.8 .2 0 6 w G> OC J 0.4 co o LU 0.2 (C) AA18 10% Hydrogen - Large Nozzle -Small Nozzle 20% Hydrogen Fig. 7.17 Factorial graph for 75 pm erosion, at 90 degrees, of: (a) A A 5 , (b) AA10 and (c) AA18 alumina coatings. (Numbers correspond to coating runs as listed in Tables 4.2 - 4.3.) 7.3.2 Reference Materials Comparisons of the erosive wear rates of the reference materials are shown in Fig. 7.18-7.21. For erosion at 45 degrees, Fig. 7.18-7.19, the mean wear rates of the AL30 coatings increased by a factor of 2.0-2.5 when the size of the erodent increased from 75 to 200 pm, and by a factor of 1.8 for an increase in erodent size from 200 to 600 pm, regardless of the processing conditions, refer to Fig. 7.19. The mean wear rates of the AT and the CR coatings increased by a factor of 2.0-2.5 when the size of the erodent increased from 75 to 200 pm as well as from 200 to 141 600 um. The wear rate of the sintered alumina (SA) increased by a factor of 1.65 when the size of the erodent increased from 75 to 200 urn, but only by a factor of 1.2 for an increase from 200 to 600 um, see Fig. 7.19. The typical mass of a splat in these coatings, based on the spray powders used, was on the order of 80 ng for the AL30, 130 ng for the AT, 9.0 ng for the CRI and 75 ng for the CR2 coatings. Therefore the 75 and 200 pm erodent particles removed less than one splat per impact on all of the coatings, while the 600 pm erodent particles removed the equivalent of more than one splat per 90 degree impact on all the reference coatings. 1.0 0.9 0.8 _ 0.7 oi I 0.6 s 1 0.5 c O « 0.4 O LU 0.3 0.2 0.1 0.0 (a) r—*— 2 r SA SA E 0.4 1.0 0.9 0.8 £ 0.6 « I 0.5-B O '» 0.4 • 0.3 • 0.2 • 0.1 0.0 LU (b) 1 rh !-2j SA •13.5 ••12 •• 10.5 ••9 I E ••7.5 1 2 .. fl 6 tr c o 4.5 o LU 3 1.5 0 1.8 1.6 1.4-1 B 1-2 •ft E r i.o | 0.8 o ui 0.6 0.4 -j 0.2 0.0 (c) rhiii n - ! - ~ n E 2-2 1 2 | H 2-1 AL30 AT 400 E Fig. 7.18 Comparison of wear rates for (a) 75 pm, (b) 200 pm and (c) 600 pm erosion using a 45 degree angle of incidence of the reference coatings plus sintered alumina (SA) as a function of powder type and processing conditions. (Sample symbols and numbers refer to Tables 4.3-4.4.) 142 100 200 300 400 Particle Size [nm] 500 600 700 (b) 0.6 0.5 E0.4 E m TO CC 0.3 H 'S0.2 o w 0.0 —•—30-1 - -A - 30-2 30-4 • -* - AT1 —•—AT2 —•—CR1 - -« - CR2-1 —•— CR2-2 —H—SA 100 200 300 400 Particle Size [|am] 500 600 700 Fig. 7.19 (a) Mass and (b) volume erosive wear rates, at a 45 degree impingement angle, plotted as a function of erodent particle size for the reference materials. For erosion at normal incidence, Fig. 7.20-7.21, the mean wear rates of the A130 coatings increased by only 5-20 percent when the size of the erodent increased from 75 to 200 pm, and from 60-90 percent for an increase in erodent size from 200 to 600 pm. The mean erosive wear rates of the AT and CR1 coatings at normal incidence increased by a factor of 2.5 for both erodent size changes, while those of the CR2 coatings increased by a factor of only 2 for both, and the sintered alumina (SA) increased by a factor of only 1.6 for an increase in erodent size from 200 to 600 pm, and not at all for an increase from 75 to 200 pm, Fig. 7.20-7.21. 143 (a) 2-1 AL30 AT CR (bi 2-1 SA)-70 „ o n 60 I i so a « 40 = o -n •30 £ UJ •20 10 0 7.0 -, (c) £ 4 . o 2-1 SA •ft Fig. 7.20 Comparison of wear rates for (a) 75 pm, (b) 200 pm and (c) 600 pm erosion using a 90 degree angle of incidence of the reference coatings plus sintered alumina (SA) as a function of powder type and processing conditions. (Sample symbols and numbers refer to Tables 4.3-4.4.) As in the 45 degee impact erosion, based on the weight [ng] of material lost per impact, the 75 pm erodent particles removed less than one splat per impact on all of the coatings, while the 600 pm erodent particles removed the equivalent of more than one splat per 90 degree impact on all the reference coatings, Fig. 7.20. The 200 pm erodent particles removed the equivalent of less than one splat per impact for all of the coatings except for CR1. Due to the range in densities of the various materials, the differences in volume wear rate are far more pronounced than in the mass wear rate, Fig. 7.21. This is especially noticeable in separating the performance of the AL30 coatings from the AT and CR, compare Fig. 7.21(b) and 7.21(a). 1 4 4 7 1 0 100 200 300 400 500 600 700 Particle Size [um] Fig. 7.21 (a) Mass and (b) volume erosive wear rates, at a 90 degree impingement angle, plotted as a function of erodent particle size for the reference materials. 7.3.3 Erosive Wear Mechanisms In 90 degree impact erosion, coating surfaces receive the maximum impact of the erodent particles. The use of angular particles leads to sharp indentations and related deformation and cracking mechanisms. Parameters such as hardness, toughness and brittleness of the material, as well as the size and hardness of the eroding particles, have a large effect on the ensuing wear mechanisms. Depending on where in the microstructure the erodent particle impacts, i.e. in the center or at the edge of a splat, in a pore, or a solidification crack, etc., the resulting damage to the coating will vary from "quasi-plastic" deformation [53] to brittle fracture. If the particles are smaller than the average splat size, the resulting damage will be primarily in the form of plastic 145 deformation. If the particles are much larger than the splat size (and the spacing of solidification microcracks) damage will occur mainly as brittle fracture (of the splats). To help illustrate the different wear mechanisms and the effects of the different microstructural features, studies were performed using a small number of erodent particles on polished coating surfaces, to obtain single impact erosion craters: 7.3.3.1 Erosion at a 90° impingement angle The primary wear mechanism for all specimens eroded with the small, 75 pm size erodents at normal incidence was plastic deformation and microfracture, Fig. 7.22. Generally, only minor parts of individual splats were removed at each impact. However, the specimens that displayed the highest wear rates showed some splat delamination. Evidence of deformation was widespread on all 75 pm eroded specimens, Fig. 7.23, as well as some brittle microfracture. Fig. 7.22 SEM micrographs of 75 pm single impact erosion craters at an impingement angle of 90 degrees on: (a) AA18 (b) AL30 and (c) CR1, showing a combination of plastic deformation and brittle microfracture mechanisms. 146 Fig. 7.23 SEM micrographs of typical 75 pm eroded surfaces at 90° angle of impingement: (a) AA5, (b) AT2, (c) CRI and (d) SA, showing primarily plastic deformation to different penetration depths, along with some microfracture. 147 When higher mass particles were used the wear mechanisms changed to create larger wear fragments, consisting of entire splats or clusters of splats with a higher proportion of brittle fracture depending on the microstructure. Micrographs of typical single impact craters using 600 pm SiC particles are shown in Fig. 7.24. Details of the splat fracture leading to the formation of wear debris particles are shown in the higher magnification micrographs in Fig. 7.25. "Quasi-Plastic" Deformation Splat fracture "Quasi-Plastic" Deformation Brittle Fracture Fig. 7.24 SEM micrographs of 600 pm single impact erosion craters at an impingement angle of 90 degrees on: (a) AA18 (b) AL30, (c) AT and (d) CR1, showing the mechanisms of plastic indentation and deformation, and brittle fracture. 148 Shear steps in lamellar structure Splat fracture along solidification cracks Fig. 7.25 S E M micrographs of 600 pm single impact erosion craters at an impingement angle of 90 degrees showing details of the splat fracture and wear debris formation processes: (a), (b) AL30 and (c), (d) AT coatings. Evidence of splat debonding and delamination was found on the surfaces of all 200 and 600 pm particle eroded coated specimens, Fig. 7.26. The specimens that displayed the highest erosive wear rate also showed the largest amount of splat 'debonding', i.e. the AL30, APS and AA5 coatings. The high porosity and poor inter-splat bonding, as discussed in Chapter 5, of these coatings led to more plastic deformation or ductile plowing wear mechanisms under 90° impact erosion. The chromia (CR) coatings showed very little evidence of splat debonding, even when eroded by the largest (600 pm) particles. 149 Splat Fracture Debonded splat surfaces Fig. 7.26 S E M micrographs (250x, lOOOx) of 600 pm eroded surfaces at 90° angle of impingement: (a)AA5, (b) AA10/18, (c) AL30, (d) AT2 and (e) CR1, showing primarily brittle fracture and some splat debonding / delamination. 150 7.3.3.2 Erosion at 45° impingement angle When the impact angle is less than 90 degrees, there is both a vertical and horizontal component to the impact force. The horizontal, or abrasive wear, component leads to a larger amount of plastic deformation in the form of plowing by the particles. Thus, the primary wear mechanism for all specimens eroded with 75 pm size erodent at a 45° angle of impingement was plastic deformation, as shown in the SEM micrographs in Fig. 7.27 and 7.28. Evidence of some splat fracture and particle pullout was also observed on the AL30, AA5-1 and APS coatings as wear damage progressed, see Fig. 7.28. The bulk sintered alumina exhibited predominately 'plastic' behaviour as well. 10|jm Fig. 7.27 SEM micrographs of typical 75 pm single impact erosion craters at an impact angle of 45 degrees, showing predominantly plastic deformation caused by: (a) blunt and (b) sharp edged particles. I 5 l Particle pullout / splat debonding Plastic deformation Splat debonding or separation Fig. 7.28 Typical wear surfaces of 75 pm eroded surfaces at 45° angle of impingement: (a) AA5, (b) the rest of the AA-series, AT and CR, (c) AL30 and (d) SA, showing primarily plastic deformation and some splat fracture and debonding mechanisms in the coatings. When higher mass particles were used the wear mechanisms changed to create larger wear fragments, consisting of entire splats or clusters of splats with a higher proportion of brittle fracture depending on the microstructure. Micrographs of typical single impact craters using 600 152 pm SiC particles are shown in Fig. 7.29. The crater in the AL30 coating, Fig. 7.29 (b), shows the effect of the poorly bonded microstructure, i.e. splat fracture and particle pullout. Plastic Deformation Splat Fracture Fig. 7.29 SEM micrographs of 600 pm single impact erosion craters at an impingement angle of 45 degrees on: (a) AA10 (b) AL30, (c) CR and (d) SA, showing the mechanisms of "quasi-plastic" deformation, and splat fracture. The primary damage mechanisms observed on all the coatings, as well as the bulk, sintered alumina (SA), eroded with 200 and 600 pm erodent particles at a 45° angle of impingement were plastic deformation and brittle fracture, Fig. 7.29-7.31. Some splat debonding was also evident on the 600 pm eroded specimens, however, it was substantially less than that observed on the coatings eroded at 90°, Fig. 7.30. 153 impingement: (a) AA5, (b) AA10 and AA18, (c) AL30, (d) AT and (e) CR, showing plastic deformation, brittle fracture and some splat debonding. 154 Fig. 7.31 SEM micrograph of bulk sintered alumina 600 pm eroded surfaces at 45° angle of impingement showing plastic deformation and some brittle fracture. 7.3.4 Cross-Sectional Studies / Subsurface Cracking The type of crack system that formed under the 600 pm particle eroded surface at both 90 and 45 degrees, as examined through SEM studies of polished cross-sections, was shown to be greatly affected by the microstructure of the individual coatings tested. 7.3.4.1 90 Degree Erosion Both the sapphire powder (AA-series) and commercial powder derived alumina coatings exhibited a lamellar structure, with varying amounts of vertical microcracking and visible splat boundaries (horizontal microcracks), as described in Chapter 5. These existing microcracks provide a low energy pathway for crack propagation upon impact by an erodent particle, see Fig. 7.32. The 'poor' microstructure, i.e. high porosity, unmelted particles and poor inter-splat bonding, of coatings such as AL30, Fig. 7.32(b), led to far greater material removal than was seen in the other AA-alumina coatings, Fig. 7.32(a). No macrocracking was detected in any of the alumina coatings. The denser alumina titania (AT) and chromia (CR) coatings, however, did exhibit evidence of macrocracking, Fig. 7.32(c),(d), similar to the lateral and median crack systems seen in fine-grained bulk alumina ceramics. The horizontal cracking seen in these coatings shows that crack propagation is still affected by the lamellar structure and the presence of pre-existing cracks, although not to the same extent as that seen in the alumina coatings. 155 (a) AA18 (b) AL30 (c) AT (d) CR Splat Boundary ... • • __ • Macrocrack Fig. 7.32 Cross-sections of 600 um eroded alumina coatings surface at a 90 degree impingement angle: (a) AA18 and (b) AL30 coatings, showing crack linking and propagation from pre-existing cracks, i.e inter-splat boundaries and vertical micro-cracks; (c) AT and (d) CR coatings, showing evidence of horizontal (lateral) and vertical (median) macro-cracks. 156 7.3.4.2 45 Degree Erosion The same mechanisms were evident in the coatings eroded with 600 um particles at a 45 degree impingement angle. Again, the lamellar structure visible in the alumina coatings played a large role in the mechanism of crack propagation, Fig. 7.33(a)-(b). The amount of material removal seen was much less as the particle normal impact force at 45° is far less than at 90°, therefore only the very top surface of the coatings were affected, compare Fig. 7.33(a) with Fig. 7.33 Cross-sections of 600 pm eroded surfaces at a 45 degree impingement angle: (a) AA18, (b) AL30 and (c) CR coatings, showing micro- and macro-crack propagation mechanisms. 157 Fig. 7.32(b). Some macrocracking was again seen in the CR coatings, Fig. 7.33(c), due to the dense nature of these coatings, and the absence of a visible lamellar structure to facilitate crack deflection, as shown in Fig. 5.13. 7.3.5 Descriptive Model for the Erosive Wear of Ceramic PS Coatings The impact of an impinging particle is an elastic-plastic event, as in quasi-static indentation. The plastic deformation at impact is usually followed by subsequent fracture of the surface. For fracture to occur, however, the impact velocity must be greater than a critical threshold value for a given particle mass necessary to impart sufficient kinetic energy to initiate surface cracking. The normal component of the impact force is the most important energy source for this fracture initiation, while the tangential component is an important factor in the abrasion damage mechanisms, i.e. micro-plowing and micro-cutting. When the stresses induced were below the fracture threshold, as in the 75 pm particle erosion tests at both 45 and 90 degree impingement angles performed in this work, the primary mechanism that occurred was plastic deformation, refer to Fig. 7.22 - 7.23. The extent of deformation that occurs, i.e. the depth, is again dependent on the hardness of the wearing material. As the impinging particles become larger, even at 45 degree impact, the normal force component can become large enough to cause fracture, Fig. 7.29-7.30. The kinetic energy of the erodent particle required to initiate fracture is dependent on the fracture toughness of the material and the size of surface flaws. In relatively defect-free bulk ceramics, such as single crystal alumina or glass, when the kinetic energy of the particle exceeds the lateral crack fracture threshold, well-developed crack systems are formed including cone cracks (with blunt particles), and lateral and median cracks (with angular particles) [71]. The impact of blunt particle edges introduces a contact stress similar to that which occurs during indentation with a conical indenter, i.e. Hertzian contact. A compressive stress field forms 158 directly under the indenter, with tensile forces building along the edges of the stress field. In erosion, the maximum tensile stresses are in-plane and occur in the plastic zone close to the elastic-plastic interface [74]. The kinetic energy of a larger, angular impacting particle was shown to be absorbed by the coating through the mechanism of crack extension, Fig. 7.34. Therefore, the number and size of cracks, either inter- or intra-lamellar, as well as the amount of porosity in general, have a strong effect on the damage mechanisms that occur. It has been shown [8] that in engineering ceramics, typically containing a number of flaws such as microcracks, grain boundaries, different phases, etc., wear occurs at the micro-scale, i.e. micro-fracture, chipping and spalling, even when macro-scale lateral cracks are present, Fig. 7.34 (a), (a) (b) Fig. 7.34 Wear mechanisms occurring in erosion: (a) in bulk ceramic materials and (b) in ceramic PS coating with poor interlamellar bonding and a high density of interlamellar cracks. In ceramic PS coatings, the process is similar. In the present work, coatings with a pronounced lamellar structure, Fig. 7.34(b), crack propagation was channeled (deflected) along the existing crack pathways, i.e. vertical and horizontal, or intra- and inter-lamellar microcracks. The strength of the inter-splat bonding is therefore critical in determining the amount of material removed. In coatings with a large percentage of inter-lamellar microcracks, or porosity (i.e. a 159 smaller percentage of bonded area between splats, e.g. in the AA5-1, AL30-1, and APS coatings in this work, Fig. 7.26), splat debonding, together with pull-out of insufficiently flattened or unmelted particles, occurred causing increased wear rates as shown in Fig. 7.14 and 7.20. As the amount of inter-splat bonding in the coatings increased, due to improved splat flattening and inter-splat grain growth, the wear rate decreased. The deflecting properties of the crack system in an AA18 coating impacted by a 600pm particle at 90 degrees are shown in the cross-sectional micrograph, Fig. 7.32(a). Macrocracking was seen after 600 pm particle erosion in the dense AT and CR coatings, Fig. 7.32(c)-(d), possibly explaining the increased wear rates observed when impacted by very large particles, i.e. those with a very high kinetic energy at impact. The microstructures of these coatings, especially the chromium oxide coatings, more closely resemble that of a bulk ceramic. However, a lamellar structure is still present, and was seen to affect the wear behaviour of the coatings, refer to Fig's. 7.25, 7.26 and 7.30. The formation of macrocrack systems provides an additional wear mechanism during high energy impact wear. Thus during high impact erosion the lamellar structure of the alumina PS coatings accomodated a great deal of strain by absorbing the energy through the mechanisms of densification and crack deflection. This improved the wear performance of these coatings as compared to the denser and normally more wear-resistant AT and CR coatings. 160 C. DISCUSSION AND CONCLUSIONS CHAPTER 8 CORRELATIONS BETWEEN THE MICROSTRUCTURAL PARAMETERS, MICROMECHANICAL PROPERTIES, AND WEAR PERFORMANCE 8.1 Introduction The micromechanical integrity of a ceramic plasma sprayed coating is determined by the size and distribution of the microstructural features, in particular the defects, found in the coating, Fig. 1.1. These can be divided into two size ranges, those at the inter- and intra-lamellar scales, as illustrated in Table 1.2. The microstructural features quantified in this work included the overall porosity, the interlamellar (horizontal) microcrack density, the intra-lamellar (vertical) microcrack density as well as the lamellar, or splat, dimensions. Based on these analyses, the coatings can be divided into 4 groups based on their microstructural integrity, or cohesion, as listed in Table 8.1. This ranking system will facilitate discussions concerning the performance of the coatings. Table 8.1 Microstructural Summary of Evaluated Plasma Spray Coatings Characteristics Group Coatings Porosity H-cracksf V-cracks* Structure 1 Chromias (CR1, CR2-l ,CR2-2) Low (<2%) ** (Low-Med. Density) (Medium density) ^ No visible lamellar structure 2 Alumina-Titanias (ATI, AT2), AAlO ' s , AA18-2,3,4, AA5-3,4 Low to Medium (1.5-4%) Low density High density Distinct lamellar structure 3 AL30-2,4, AA5-2, AA18-1 Medium (3-5%) Medium density Low density Visible lamellar structure 4 AA5-1, APS, AL30-1 High (>8%) 1 High r density 1 Low r density Chaotic: structure **(No visible crack structure in CR1) 1 (Horizontal, or Inter-lamellar cracks) ^(Vertical, or Intra-lamellar cracks) 161 The effect of these microstructural features on the mechanical and wear properties of the coatings, and the relationships between the micromechanical properties and the wear performance of the coatings will be investigated through the use of scatter plots in this section: 8.2 Porosity 8.2.1 Effect of Porosity on Micromechanical Properties The porosity of the AA-alumina coatings was plotted against their hardness (Hw) and elastic modulus (£) measured using depth sensing indentation (DSI) at 0.1 and 1.0 N loads, Fig. 8.1(a)-(b). A moderate correlation was noted for the 1.0N Hw values on both planar and cross-sectional surfaces, but not for those at 0.1N. At a load of 0.1N, the size of the indentations was smaller than individual splats on the planar but not the cross-sectional surfaces, as presented in Section 6.2. A l l of the measurements at this low load exhibited significant scatter, making the evaluation of any distinct trends difficult. The indentations on the planar surface, varying from 3.1-4.5 pm in length, sampled the material response of individual splat diameters, but not the thicknesses, and could be affected by the presence of micro-scale intra-lamellar pores, which range from 1-10 pm in size [36]. However, micro-scale porosity is not always included in porosity measurements, which may help explain the low correlation between the porosity and these low load hardness values. Little correlation was seen between porosity and elastic modulus measured using DSI at 0.1 or 1 .ON as expected since the coatings were all sprayed from the same powder material. 162 (a) (b) 30 25 !e 2 0 Q. o 8 15 c •a I 10 5 0 20 18 16 ~ 14 « 2,12 S 1 0 c "S 8 ra X 6 4 2 0 HW(0.1N)-Top • • • • • 4.0 6.0 Porosity [%] HW(1.0N)-Top —I 10.0 30 25 B " 20 a. o 8 15 0) c TJ I 10 5 0 HW(0.1N)-X-Sec 20 n 18 16 ~ 14 ca U 12 1 10 C "2 8 co 1 6 4 2 0 4.0 6.0 Porosity [%] HW(1.0N)-X-Sec 0.0 2.0 4.0 6.0 Porosity [%] 0.0 4.0 6.0 Porosity [%] Fig. 8.1 Effect of porosity on measured hardness (Hw) at (a) 0.1 and (b) 1.0N on planar and cross-sectional surfaces for the AA-alumina coatings. The inclusion of data for the chromia and alumina-titania coatings and the bulk sintered alumina on the microhardness plots shows the effect that different microstructures have on the correlations, Fig. 8.2-8.3. The circles on Fig. 8.2 illustrate the segregation of the coatings into the groups listed in Table 8.1, based on their porosity and hardness. H„ (3N) - Top 0 SA • CR • AT • AL30 • AA10/18 o APS A AA5 12.0 Fig. 8.2 Effect of porosity on microhardness (Hv) at 3N for all evaluated materials. 163 S. 10 a HK(10N)-Top . 10 4 CAL30 (a) H„ (10N)-Top (b) 6.0 Porosity [°/<j 100 120 4.0 6.0 Porosity [%] 8.0 10.0 Fig. 8.3 Effect of porosity on (planar) Knoop microhardness (Hk) at 10N for: (a) all materials and (b) alumina-based coatings only. With their dense and non-lamellar microstructures, the bulk alumina (SA) and the chromia (CR) coatings (Group 1) exhibited much higher hardness values than the lamellar-structured alumina and alumina-titania (AT) coatings. This was reflected in the Vickers hardness (Hv) at 3N and 100N, as well as in the Knoop hardness at 10N, as shown in Section 6.2. Note the linear trend when the porosity is plotted against the hardness values of only the alumina-based coatings, Fig. 8.3(b). However, the best correlation when the CR coatings and SA were included was nonlinear. Another measure of hardness that has been found to correlate with the density of green ceramic compacts is known as scratch hardness [118]. This is calculated by measuring the width of the scratch groove at a defined load, as discussed in Section 6.3.1. The scratch hardness of the evaluated materials at 60N was also found to correlate exponentially with the measured porosity, Fig. 8.4. 164 100 Porosity [%] Fig. 8.4 Relation between scratch hardness measured at 60 N and porosity of the evaluated materials. An inverse power relationship was also seen when the elastic modulus (£), measured using a Knoop indenter, was plotted as a function of the porosity of all the tested materials, Fig. 8.5. A l l of the PS coatings, even the dense chromia coatings, exhibited much lower stiffness than the bulk sintered alumina. Porosity in bulk ceramics affects the average elastic modulus by decreasing it. This has been estimated by the following exponential equation [119]: E = E(,exp(-bP) (8.1) where E and E0 equal the elastic modulus of the porous and nonporous polycrystalline ceramic, respectively, b is a constant (~2) and P equals the volume fraction of porosity. The interlamellar porosity in ceramic PS coatings has been described as oriented spheroids [39], with the elastic modulus estimated as for composites [120]: 1-(5a 3^ — + — P [4c 4) (8.2) where c is the axis parallel to the stress direction and a is the axis in the plane perpendicular to axis c. In the present work, the correlation between porosity and E was best described by a linear trendline when only the alumina-based coatings were considered. These coatings exhibited a definite lamellar structure, whereas the chromia coatings did not. The E values for the chromia coatings as measured on the coating cross-sections, perpendicular to the surface, were lower than 165 those measured parallel to the top surfaces, Fig. 8.5, an indication that even though the lamellar structure was not visible in these coatings, it exerts an anisotropic influence nonetheless. 400-| 350 _ 300-] CC Q. C3 =• 250 w _d | 200 S 115° is UJ 100-50-0 •5, (10N) - X-Sec - Perpendicular o SA 0.0 2.0 4.0 6.0 Porosity [°/<J —I— 8.0 400 350 „300 CD 0. ts — 250 n f 200 5 ~ 150 100 50 0 ui 6, (10N)-X-Sec-Paralel SA o SA i i CR & AT 4> AA, AL 10.0 12.0 0.0 20 4.0 60 8.0 Porosity ["/tJ 10.0 120 Fig. 8.5 Effect of porosity on Elastic Modulus (E) at 10N measured on cross-sectional surfaces: (a) perpendicular and (b) parallel to the coating surface. A moderate correlation was found between the porosity and H/Kc, the "index of brittleness" of the denser coatings for which fracture toughness could be determined, Fig. 8.6. As expected, the less porous the material, the more brittle and, hence, the greater the value of H/Kc. Again the bulk alumina with its dense microstructure exhibited a much higher H/Kc ratio for a given level of porosity than was found for the PS coatings. 5 4.5 4 3.5 3 2.5 2 1.5 1 0.5 0 OSA o SA • CR • AT AA10 & 18 2.0 3.0 4.0 Porosity [%] 5.0 6.0 Fig. 8.6 The "index of brittleness", H/Kic% as a function of porosity for the materials with a measured KJc value. 166 8.2.2 Effect of Porosity on Wear Properties of Coatings Of all the microstructural features examined, the level of porosity in the tested materials was found to have the most direct effect on the amount and type of wear damage. Pores, especially the non-spherical pores common in ceramic PS coatings, act as stress-concentrating flaws. The best correlation was found between the porosity and the 20 pm particle abrasive wear rates of the coatings, Fig. 8.7. The 200 pm particle abrasive wear rates, the scratch cross-sectional areas at 60N and the small particle erosive wear rates, showed similarly good correlations (correlation coefficient R2 > 0.75) with porosity, with slightly increasing degrees of scatter, see Appendix C for the erosion plots. 0 SA • CR • AT • AL30 AA10/18 o APS A AL30 0.04 -\ E E, % 0.03 H CC (a) • A L 3 0 - 1 0.0 2.0 4.0 6.0 8.0 Porosity [%] 10000 -i E •2. 6000 CO < I 4000 u co 2000 (b) 0.0 2.0 4.0 6.0 8.0 Porosi ty [%] Fig. 8.7 Correlation between porosity and: (a) abrasive wear rate using 20 pm particles and (b) the scratch cross-sectional area at a load of 60N, for all of the evaluated materials. 167 Porosity has been shown to have a direct effect on the wear of bulk ceramics. The abrasive wear of polycrystalline alumina was shown to increase exponentially with increasing porosity [60]. In the present work, however, a power relationship of the form W = AP" (8.3) where Wis the wear rate, A is a constant, P is the percent of porosity and n a constant, was found to have a higher correlation than an exponential one, for example an R2 value of 0.812 vs. 0.645 for the 20 pm abrasion results. The correlation was almost linear in the case of abrasive wear, with a varying from 0.84-0.95. In the case of erosive wear, exponent a varied from 0.4-0.5.. 8.2.3 Summary The level of porosity has a direct effect on both the micromechanical properties of ceramic PS coatings and their wear performance. In the present work, the amount of measured porosity was found to have a low correlation with respect to the elastic modulus, moderate with respect to hardness, especially the Knoop hardness, as well as the brittleness index (H/Kc). Porosity was shown to have a moderate to very high effect on the wear properties of the coatings, the highest being the abrasive wear properties. This strong correlation compares well with effects reported in the wear performance of bulk ceramics [60]. 8.3 Interlamellar (Horizontal) Microcracking The amount of horizontal, i.e. parallel to the substrate, microcracking is directly related to the amount of bonding between splats, or interlamellar bonding, and is a microstructural feature unique to plasma sprayed coatings, as described in Sections 2.2 and 5.4. The horizontal crack density is not directly convertible into a measure of the amount of bonded area between splats. However, it does give an indication of the bonding, or the lack thereof. In the present work, coatings that contained a larger amount of porosity, i.e. the Group 4 coatings, also tended to exhibit a higher density of interlamellar microcracks, Fig. 8.8. This is because the horizontal 168 microcracking is also considered interlamellar porosity and as such it is at least partially included in any measurement of porosity. Coatings sprayed with larger particles, such as the AL30-1, exhibited a slightly lower crack density, compared with the amount of porosity, due to the presence of unmelted and incompletely melted particles throughout the coating microstructures, refer to Fig. 5.5, thereby skewing its position relative to the rest of the coatings. The AA5-1 coating, with a high percentage of particle pullout and unmelted particles as well, also exhibited a higer porosity level compared to the level of microcracking. 12 10 '35 6 o • AL30-1 AAA5- APS 50 100 150 Horizontal Crack Density [mm" 1 ] 200 Fig. 8.8 Correlation between coating porosity and the linear density of horizontal microcracks. 8.3.1 Effect of Inter-lamellar Microcracking on Micromechanical Properties The extent of inter-lamellar bonding is the most important factor influencing the behaviour of PS coatings, in terms of both physical properties and wear performance. In this work, the effect of the horizontal crack density on the microhardness was far less than that observed with porosity. Essentially no correlation was found between the horizontal crack density and the cross-sectional and planar hardness values for the AA-alumina coatings evaluated at low loads using depth sensing indentation (DSI). 169 A lower level of inter-lamellar bonding, as measured by an increase in the density of horizontal microcracks, caused a slight reduction in the Vickers hardness values of the various coatings as measured on their cross-sectional surfaces, Fig. 8.9. The effect was less noticeable for the Knoop than Vickers hardnesses. The Group 1 (chromia) coatings, showing no distinct lamellar structure, exhibited higher hardness values than the rest of the coatings for comparable horizontal crack densities, as reflected in the Vickers hardness (Hv) values at 3N. The APS coating, sprayed using a radial- instead of an axial-feed PS system, exhibited a lamellar microstructure with more extensive microcracking, giving higher crack densities, than the other AA18 coatings. Essentially no correlations were found between the measured elastic moduli of the coatings and the horizontal crack density. Elastic modulus values are affected by the amount of bonded area between the splats, as discussed in Section 2.3.2, however, the percent area is not expressed adequately by the crack density alone. No correlation was found between either the fracture toughness or the index of brittleness (H/Kc) and the horizontal crack density. 25 n OAPS o -I 1 1 i 1 0 50 100 150 200 Horizontal Crack Density [mm"1 ] Fig. 8.9 Effect of the linear horizontal crack density on the Vickers hardness (Hv) measured on the cross-sectional surfaces at 3N. 170 8.3.2 Effect of Inter-lamellar Microcracking on Wear Properties The amount of inter-lamellar bonding in ceramic PS coatings has a major influence on the amount of wear loss as well as the type of damage that occurs, as described in Section 2.3.6. The erosion rates of several ceramic PS coatings were found to be inversely proportional to the average bonding rate within the coatings [67]. In the present work, the best correlation was found between the inter-lamellar bonding, as indicated by the horizontal crack density, and the abrasive wear rates of the evaluated materials, regardless of the abrasive particle size, Fig. 8.10. Note that the AL30 coatings exhibited a higher abrasion rate due to the large number of unmelted particles and large splat sizes, in addition to the large number of horizontal delaminations. An equation of the form: W = A(CH)" (8.4) where W is the wear rate, A is a constant, Ch is the linear density of horizontal microcracks and n is an exponent, was found to best describe the relationship. The correlation was almost linear in the case of abrasive wear, with a varying from 0.82-0.93. Similar effects were noticed in the 60N scratch cross-sectional area results, Fig. 8.11. A low to moderate correlation was found between the horizontal crack density and the different erosion rates, which decreased rapidly as the erodent particle size and impact angle increased, as shown in Appendix C. 0-1—I—I—I—I—I—I—I—I—I—I—I—I—I—I—I—I—I—I—I—I 0-1—I—I—I—I—I—I—I—I—I—I—I—I—I—I—I—I—I—I—I—I 0 50 100 150 200 0 50 100 150 200 Hrxizontal Crack Density [mm"'] H^ontal Crack Density [mm"'] Fig. 8.10 The effect of the horizontal crack density on: (a) 20 pm and (b) 200 pm particle abrasive wear rates. 171 10000 „ 8000 6000 1 a> 4000 2000 O APS CR1 -t- -t- -+-50 100 150 Horizontal Crack Density [mm"' ] 200 Fig. 8.11 The effect of the horizontal, or inter-lamellar, crack density on the cross-sectional scratch area, of the 60 N scratch groove, on the evaluated coatings. 8.3.3 Summary In this part of the work, low correlations were found between the density of horizontal microcracks, a measure of the lack of inter-lamellar bonding, and the measured micromechanical properties, especially when considering all of the coatings together. It is known that the percentage of bonded area between splats does have an effect on the mechanical properties such as elastic modulus, as discussed in Section 2.3.2. The horizontal microcrack density is not directly convertible to a value of the amount of bonded area however. No direct correlation was found to the fracture toughness or brittleness index of the coatings. The horizontal crack density was shown to have some effect on the different wear results, especially when the "Group 4" coatings, i.e. AA5-1, AL30-1 and APS coatings, with unmelted particles and non-flattened splats, were excluded from the trendlines. The highest correlations were found between the inter-lamellar crack density and the wear results where the major damage mechanism was plastic deformation, i.e. abrasion, scratching and small particle erosion. This indicates that the inter-lamellar bonding has a strong influence on wear by the mechanism of "plastic" deformation, or "quasi-plastic" in the case of brittle ceramic materials. This can be partially explained by considering the mechanism of plastic deformation, which is exemplified 172 by the scratching of a particle across the surface. When the particle, or indenter, reaches the critical load necessary to initiate cracking on the surface, the wearing surface, in this case the PS coating, will begin to fracture. Splat fracture is common in ceramic PS coatings, and the level and strength of inter-splat bonding, as well as the vertical microcrack density, will be a critical factor in how much of the splat will fracture and detach as a wear debris particle, thereby contributing to the wear rate. Therefore, the lower the inter-splat bonding level, i.e. the greater the density of horizontal microcracks, the greater the resulting wear loss. 8.4 Intra-lamellar (Vertical) Microcracking The density of intra-lamellar, or vertical, microcracks is a function of the quenching stress, an in-plane tensile stress caused by the constraint or bonding of the impinging splat as it undergoes thermal contraction during solidification. Stress relief in ceramic PS coatings is obtained through microcracking perpendicular to the plane of the splat, as described in Section 2.2.7. The greater the inter-lamellar bonding, the more the surfaces of the splats are constrained and the greater the number of vertical microcracks which are formed to relieve the built-up stresses. Thus, one would expect there to be an inverse relationship between the vertical crack density and the density of horizontal microcracks, i.e. the interlamellar bonding, as confirmed in Fig. 8.12, as well as the mechanical properties and the wear performance. 173 200 180 - ~ 160 > 1 4 0 I 120 o a tj 10° re 0 80 ra 1 60 N | 40 20 0 >APS ± AA5 • A T • CR • AL30 A AA5 AA10/18 O APS — I 1 1 1 1 I 1 1 1 10 20 30 40 50 60 70 80 90 Vertical Crack Density [mm1 ] Fig. 8.12 Inverse relationship between the horizontal and vertical microcrack densities. A low correlation was found between the amount of porosity and the density of vertical microcracks, Fig. 8.13, even when the poorly bonded, "Group 4" coatings (AA5T, AL30-1 and APS) were excluded from the trendline. The general trend noticed was that the more porous coatings contained less vertical microcracks. The non-optimized coatings, i.e. where the splats s. are not completely flattened and containing unmelted particles, have higher levels of porosity as well as a lower density of vertical microcracks due to the lack of constraint on the splat edges, which allows stress relief to occur without microcracking. 12.0 8.0 6.0 4.0 2.0 • AL30-1 40 60 Vertical Crack Density [ m m ' 1 ] 100 Fig. 8.13 Inverse relationship between coating porosity and the linear density of vertical microcracks. 174 8.4.1 Effect of Intra-lamellar Microcracking on Micromechanical Properties A general trend was found between the density of vertical microcracks and coating hardness, Fig. 8.14. This was expected as coatings with a higher level of interlamellar bonding also tended to have a higher density of vertical microcracks. As with the horizontal microcracks, the Group 1 (CR) coatings, with no visible lamellar structure, exhibited slightly higher hardness values than the rest of the coatings. The CRI coating was so dense that it had no visible vertical microcracks. Essentially no effect was found due to the vertical crack density on elastic modulus values, or on the fracture toughness or brittleness index values. 2 5 . 0 n 0.0 -I—I—I—I—I—I 1—I—I—I—I 1—I— I— I— I—I—I— I 1—I— I— I—I— I—I 0 2 0 4 0 6 0 B 0 1 0 0 Vertical Crack Density [mm'1 ] Fig. 8.14 Effect of vertical crack density on coating hardness (Hv) measured at 3N. 8.4.2 Effect of Intra-lamellar Microcracking on Wear Properties The vertical microcracks present in ceramic PS coatings play a large role in the formation of wear debris, as seen in Sections 6.3.1.2, 7.2.3 and 7.3.3. The presence of pre-existing cracks within the individual splats provides low energy pathways for crack growth during the processes of abrasion, including scratching, and erosion. The effect of these intra-lamellar microcracks has not been investigated in detail and they have been ignored in the few wear models for PS coatings, as discussed in Section 3.5. 175 In order to investigate a potential relationship similar to the Hall-Petch like relationship between grain size and abrasion resistance in ceramics [58], the spacing between the vertical microcracks (equal to the reciprocal of the crack density) of coatings with well-flattened splats, i.e. from Groups 2 and 3 (alumina-based), was plotted against the scratch width at different loads, Fig. 8.15. Good correlations were found for scratch widths measured at loads from 30 - 60 N . As the size of the wear particles are related to the vertical crack spacing, this indicates the operation of a size effect, although not one identical to that which occurs in ceramics. The reason for the difference is that the strength and extent of inter-lamellar bonding in the coatings is an important factor as well. 200 180 a 160 •D 5 o 120 -I 100 80 60 N 40 N 1.0 1.5 2.0 2.5 Vertical Crack Spacing [|im] 3.0 Fig. 8.15 Relation between scratch groove width and distance between vertical microcracks for Group 2 and 3 coatings. The best correlation between the vertical crack density and wear performance was found as measured by controlled scratching at 60N. The exclusion of the Group 4 coatings, containing large quantites of unmelted and non-flattened particles, increased the correlation coefficient from 0.53 to 0.75, Fig. 8.16. Of all the erosion results, the highest correlation was found between the intra-lamellar microcrack density and the 600 pm particle - 45 degree impact erosion, Fig. 8.17. However, when the Group 4 coatings as well as the small splat size AA5 coatings were ignored, 176 the correlation decreased drastically. Essentially no correlations were found between the vertical crack density and the rest of the wear results as a function of the vertical crack density, as shown in the scatter plots in Appendix G. — i — i — i — i— i — i — i— i — i — i— i — i — i — i— i — i — i — i — i — i — i 20 40 60 80 100 Vertical Crack Density [mm"1 ] Fig. 8.16 Relation between the scratch groove cross-sectional area at 60N and the linear density of vertical microcracks. 0.600 0.500 CD ^ 0.400 £ o a 0.300 DC o 0.200 LU 0.100 600 um • AA5-1 0.000 -I—I—I—h-• AL30-1 • AL30-2 -t- H 1 1 1 1 1 1 1 1 1 1 1 20 40 60 Vertical Crack Density [mm"' ] 80 100 Fig. 8.17 Relation between the 600 pm erosion rate at a 45 degree impingement angle and the linear density of vertical microcracks. 177 8.4.3 Summary In this work, a low inverse correlation was found between the intra-lamellar microcrack density and the porosity level in the coatings, while essentially no relationships were found between the density of vertical microcracks and the measured micromechanical properties. Low to moderate inverse power correlations were found with the different types of wear when all of the coatings were considered. However, when the Group 4 coatings as well as the small splat size AA5 coatings were ignored essentially no correlation was found. These coatings, due to the presence of poorly flattened and bonded splats, do not undergo residual stress reduction in the form of microcracking. This occurs when the splats are well bonded and thus constrained. The rest of the coatings exhibited a broad range of vertical crack densities for a narrow range of wear rates. Linear relationships were found between the inter-crack spacing of the coatings with well-flattened splats, i.e. the optimized AA-alumina and AT coatings, and the single pass scratch groove widths. 8.5 Correlations Between Micromechanical Properties and Wear Performance., 8.5.1 Effect of Material Hardness on Wear Rates Hardness has a large effect on the wear of materials by mechanisms of plastic deformation, while fracture toughness is a dominant factor in wear involving brittle fracture [71]. This was seen in the present work with better correlations found between the hardness of the wearing material and the forms of wear that had plastic deformation as a major mechanism. The best correlation was found between the Knoop hardness (planar surfaces) and abrasion with 20 pm particles, followed by 45 degree erosion with 75 pm particles, abrasion using 200 pm particles and 90 degree erosion with 75 pm particles (All with R2 values > 0.8), Fig. 8.18 and Appendix C. 178 The correlations found with the Knoop hardness (10 N), as measured on both the top surfaces and on the cross-sections perpendicular (Pd.) to the top surface were higher than with Vickers hardness (3.0N). Knoop hardness is more sensitive to the near-surface properties, due to the shape of the indenter, as discussed in Section 3.2. The different material "Groups" were also seen clearly, Fig. 8.18. Wear damage also occurs at the outer surface of materials, possibly explaining the higher correlations found for the Knoop hardness. The relationship between hardness and the wear results in this work took the form of: W=k/H" (8.5) with W being the wear rate, k and n are constants and H the hardness of the surface. Exponent n varied from 2.8 for abrasion to 1.6 for small particle and low angle erosion. Fig. 8.18 The 20 pm particle abrasion rate as a function of Knoop hardness, measured at ION on planar surfaces. The wear rate increased sharply as the hardness decreased below around 8.0 GPa. This increase is due to a combination of inter-lamellar porosity and poor intersplat bonding in these coatings, as neither factor by itself can account for it. To illustrate this effect, which is due to the unique microstructure of ceramic PS coatings, the wear rate of a series of bulk ceramics was plotted against their hardness (Hv), based on data from Moore and King [53] and compared with 179 similar small particle abrasion test results found in the present work, Fig. 8.19. As a check on the wear range, the bulk alumina used in the present work had a similar wear rate to the aluminas in [53]. The distance a corresponds to the microstructural components affecting the wear rate for bulk ceramics, primarily porosity, while b corresponds to the additional factors found in the unique microstructure of ceramic PS coatings, particularly the level of interlamellar bonding, including the horizontal crack density and shape factor, or degree of flattening, of the splats. 0.05 -i 0.04 0.03 DC V i •5 0-02 -1 0.01 H 0.00 Ceramic PS Coat ings Bulk Ceramics b \ 1 1 i i 1 1 10 15 20 Hardness [GPa] 25 30 35 Fig. 8.19 Comparison of the small particle abrasion rate as a function of hardness (Hv) for bulk ceramics vs. ceramic PS coatings. The distances a and b illustrate the contributions of different microstructural factors. See text for details. 8.5.2 Correlation of Scratch Hardness with Wear Rates Another measure of hardness that has been shown to have good correlation with various wear results, especially in the area of abrasion of ductile materials [86], is known as scratch hardness. This is calculated by measuring the width of the scratch groove at a defined load, as discussed in Section 6.3.1. Scratching is essentially single point abrasion, thus the best correlations were found between scratch hardness and the abrasive wear rates, Fig. 8.20, where plastic deformation was the dominant wear mechanism. The relationship between scratch hardness and the wear results in this work took the same form as for conventional hardness 180 values, as listed in Equation 8.5. However, exponent n varied from 1.3 for abrasion to 0.75 for small particle and low angle erosion. Moderate inverse (non-linear) correlations were also found with the erosive wear rates, refer to Appendix C, with decreasing correlations as the wear mechanisms changed to brittle fracture. The AL30-1 coating exhibited a higher hardness than expected due to the large number of unmelted cx-alumina particles, as shown in Section 5.5.2. 0.06 -i 0.05 H c| 0.04 H E 8 S 0.03 H cc c o J3 o.oi H • AL30-1 AA5-1A o SA • CR • A T • AL30 A AA5 AA10/18 o APS 0 S A 0.0 2.0 4.0 6.0 Scratch Hardness [GPa] i 1 1 1 8.0 10.0 12.0 14.0 Fig. 8.20 Inverse correlation between the scratch hardness (at 60N) and the abrasive wear rate using 20 pm particles. 8.5.3 Effect of Elastic Modulus on Wear Rates Elastic modulus is not included in the majority of wear models [71]. In the present work, the best correlation was found between the Knoop elastic modulus (is*) and the high impact angle, large particle erosion, Fig. 8.21. Examining this graph in greater detail, three different material groupings can be seen. In this case Group A is comprised only of the bulk ceramic, with a higher Zs-value and lower wear rate, while the 2 n d group includes the majority of the coatings, 181 including the chromias, all within a narrow range of wear rates. Group C encompasses the high porosity, poorly bonded coatings, including AL30-2 this time, which exhibited very high wear rates. The internal bonding of the coatings and bulk ceramic, which exhibit a strong effect of the elastic modulus, as described in Section 2.3.2, were shown to have a major influence on wear by high energy impact. However, the differences in E have to be substantial to really have an effect. 2.5 E E 1.5 I 1 « o LU 0.5 50 100 150 200 250 300 350 E k [GPa] 400 Fig. 8.21 Effect of elastic modulus (E) on 90 degree impact erosion results with 600 pm particles for all evaluated materials. 8.5.4 Effect of Fracture Toughness on Wear Rates Fracture mechanics has been used to explain and predict the wear of brittle materials [51, 87]. The models were based on ideal brittle materials such as glasses and single crystals, such as sapphire, and an assumption of the formation of lateral and radial crack systems. However, when these models were applied to more typical polycrystalline engineering materials, i.e. coarser-grained aluminas, where lateral cracking has not always been observed to occur [56, 78], the model predictions were shown to be inaccurate. One of the main problems with these models is the simplified consideration of only one type of wear mechanism at a time, either plastic 182 deformation or brittle fracture, and for ideal ductile or brittle. materials. Typical real-life engineering materials like plasma sprayed ceramic coatings however, do not fall into just one category. As has been clearly demonstrated by the present work, mechanisms of wear by hard particles in ceramic PS coatings usually involve both plastic flow and brittle fracture at the same time. The effect of fracture toughness on the wear resistance of the materials evaluated in this work was found to be connected to the hardness. Fracture toughness values alone did not exhibit much of an influence, i.e. very low correlations were found. The material parameter (Kt,/2H1/2)'', equivalent to the KIcH factor in Equation 3.9, has been suggested to describe the material removal rates of brittle ceramics due to brittle fracture during the grinding process, i.e. abrasion [76]. Moderate to high correlations were found between the wear rates controlled by plastic deformation instead of fracture, i.e. abrasion, Fig. 8.22, scratching, 45 degree impact erosion and small particle 90 degree impact erosion, and this material parameter. The division of this parameter into E?/5, as suggested in Equation 3.9 however, lowered the correlation considerably, as shown in Fig. 8.22(b). 183 0.03 n 0.025 j= 0.02 H 0.015 H § o.oi H 0.005 H (a) 0.0 1.0 OSA 2.0 3.0 4.0 5.0 6.0 0.03 0.025 H E 0.02 H 0.015 H 5 0.01 H 0.005 H CR 0.0 50.0 100.0 150.0 200.0 250.0 300.0 350.0 Fig. 8.22 The specific abrasive wear rate using 20 pm particles as a function of the toughness/ hardness parameter: (a) K]c~l/2H~1/2, and (b) combined with the elastic modulus as E4/5K1-mKm. Wear of ceramic plasma sprayed coatings is a complex process involving the mechanisms of both plastic deformation and brittle fracture. Thus the existing models developed for the wear of brittle materials, based on mechanical properties such as hardness, elastic modulus and fracture toughness, which have proven inadequate for modern engineering ceramics, are also unsuitable for predicting the wear of the PS coatings. A combined microstructural factor or parameter is needed that takes into account the unique microstructural features of these coatings, including porosity and interlamellar bonding. For example, taking a single empirical approach 184 and multiplying the toughness-hardness parameter by a microstructural factor (M), M = P", where P is the porosity fraction and n is a constant, does help move the beginning of the trendline closer to zero, Fig. 8.23. 0.03 -> 0.025 E 0.02 0.015 § 0.01 0.005 0.0 0.5 1.0 1.5 2.0 Fig. 8.23 The specific abrasive wear rate using 20 pm particles as a function of the toughness-hardness parameter K/cI/2H1/2 times a microstructural factor, M = P", with n •= 0.4. A complex relationship is expected to exist between the inter-lamellar bonding, the level of porosity and the amount of vertical and horizontal microcracking. As the size of the typical wear debris particle generated in ceramic PS coatings is seen to be related to the spacing between vertical microcracks as well as to the horizontal delaminations, a potential microstructural factor could be related to these parameters, similar to the relationship between grain size and abrasion resistance in ceramics. As discussed in Section 8.4.2, the vertical crack spacing is related to the width of the scratch grooves at different loads. However, the density of vertical microcracks in the splats is dependent on the temperature difference between the incoming molten splats and the underlying material. A better measure of the strength and extent of the interfacial bonding between splats than the density of horizontal delaminations is needed. 185 8.5.5 Summary Table 8.2 presents a summary of the major relationships found between the measured microstructural features, the micromechanical properties and the wear properties. As discussed, the best correlations were found between the material hardness (H), the level of porosity (P) and the wear volume (W), specifically the abrasive wear volume. Table 8.2 Summary of Relationships between Microstructure and Properties of Ceramic PS Coatings Relationships with: Microstructural Feature Micromechanical properties Wear properties (W = wear rate) Porosity (P) E = Ea[l-kP] H = k/P" W=kP" Inter-lamellar (horizontal) microcrack density (CH) (related to inverse of intersplat bonding (B)) CH~P CH ~ 1/H (Ch~1/B) W = A(CHf ( ocA/B") Intra-lamellar (vertical) microcrack density (Cy), or the inter-crack spacing (V) Cv~(1-CH) CV ~ 1/P Cv°cH W(90° erosion) x VCv Scratch Width = k(l/Cv) = kV Micromechanical properties Wear properties Hardness (H) Correlation: Knoop > Scratch > Vickers W = k/H" Elastic Modulus (E) (Knoop) W oc l/E" (poor corr., need wider range of materials) Fracture Toughness (Kc) (Indentation) No correlation. Combination of H and Kc W=f(Kcll2Hm) Microstructural Factor (M) -Combination of P, CH (or some measure of bonding (B)) M = PB/BA 186 8.6 Sensitivity Comparison of Evaluation Methods To compare the sensitivity of the methods used in this work to differences in microstructural integrity, the wear results were normalized with respect to sintered alumina as a reference. To minimize the number of results, only the #1 and #4 runs for the A A - and A L -alumina series were used. These typically had the worst and best microstructures, respectively. Typical A T and CR coatings were also selected. The greatest sensitivity to the microstructural differences was shown by the abrasion methods, Fig. 8.24. The coating with the worst abrasion performance was the AL30 coating, followed by APS and AA5-1, i.e. the Group 4 coatings. The low angle and small particle size erosion tests gave similar rankings to the coatings, although there were some differences between the Group 4 coatings (AA5-1, APS and AL30-1). An overview comparing the sensitivity of all the wear methods is shown in Fig. 8.25, with the methods listed in order of decreasing sensitivity in the legend box. The lowest sensitivity was shown by the cross-sectional scratch method, giving values of critical fracture cone length for the different materials. Fig. 8.24 Normalized comparison of abrasion evaluation methods, including 20 and 200 pm abrasion and the 60 N scratch cross-sectional area. 187 0 -J 1 1 1 i i i 1 i i i i i i < < Fig. 8.25 Normalized comparison of all wear-related evaluation methods. The erosive wear methods, however, gave slightly different comparative results at the 600 pm particle size. As shown in Fig. 8.26 and 8.27, the wear rates of the dense CR coatings were as high or higher than those of the optimized AA-alumina coatings. As discussed in Section 7.3, the wear rates of all the AT and CR coatings were quite high. The impact energy from the large 600 pm particles, especially at a 90 degree angle of impingement, was sufficient to induce macrocracks resulting in extensive spalling in addition to brittle fracture. The lamellar structure of the alumina coatings, however, was better able to tolerate the intense strain by absorbing the impact energy through the mechanisms of crack deflection and densification, Fig. 7.31. Therefore, the best method for evaluating the microstructural cohesion and integrity, as defined by the severity of splat wear typically encountered in real-life wear systems, would make use of a medium erodent size, say from 200- 300 pm. 188 ? 5 N E 90 Degree Erosion V T V f OT V f in in o o °i eo co < T - < T - ^ < < < < < o CM I -< T- CN o cc o < to Fig. 8.26 Normalized comparison of 90 degree erosion evaluation methods. > 5 u CD N TS 4 E o „ 45 Degree Erosion V T V in ur — < < J It) V I 0. co T - <. y-< < < < CO o o CO CN i -< o CM CM CC CJ < LO Fig. 8.27 Normalized comparison of 45 degree erosion evaluation methods. 189 CHAPTER 9 SUMMARY AND CONCLUSIONS A series of ceramic plasma sprayed (PS) coatings, both alumina- and chromia-based, were sprayed according to a matrix of deposition parameters in order to produce a broad range of microstructures. A l l but one of the coatings were sprayed using an axial feed injection PS torch, the exception being sprayed using a radial feed PS torch. To investigate the effect of splat size on the coating response, a series of mono-crystalline a-alumina powders with very narrow particle size ranges, nominally 5, 10 and 18 microns in diameter, were sprayed. Using metallurgical analysis and electron microscopy techniques, as listed in Table 5.1, the coatings were extensively characterized for a variety of microstructural features, including porosity, the angular distribution and density of microcracks as well as the splat dimensions. The coatings can be divided into 4 groups based on their microstructural integrity, or cohesion attributes, as listed in Table 8.1. The coatings were then evaluated using the micromechanical techniques listed in Table 1.3 to investigate their response to different contact situations. This is important as the range of defect sizes found in the unique lamellar structure of ceramic PS coatings affects their response to different types and magnitudes of contact stresses and studying these responses helps lay the groundwork for the evolution of models of their wear behaviour. 9.1 Summary of Major Results for each Method 9.1.1 Indentation Indentation techniques were used at loads ranging from as low as 0.1N to as high as 100N. Depth sensing indentation was performed at a load of 0.1N in an attempt to evaluate the micromechanical properties of the individual splats. The resulting indentations, and affected plastic zone, were typically smaller than the splat diameters, but not the splat thicknesses, as well as being in the same size range as a variety of the microstructural features which gives rise to large amounts of scatter in the results. An even lower load would need to be used in order to 190 evaluate the properties of the individual splats, with a larger number of tests to permit a thorough statistical analysis. In the microindentation load range, both Vickers and Knoop indenters were used. Knoop hardness was found to be a better indicator of surface hardness than Vickers, due to the shape of the indenter, as shown by the higher correlations found with the wear results described in Section 8.5. Elastic modulus values were also determined for the coatings using the Knoop indenter. These values varied depending on the position of the indenter, i.e. on the top or cross-sectional surfaces. The values tended to be higher on the top surfaces. Larger values were measured when the indentation on the cross-section was made perpendicular rather than parallel to the coating surface. The low toughness in the parallel direction was considered to be a direct consequence of the weak bonding between splats. The differences were less pronounced in the denser coatings, AT and CR, illustrating that these coatings were less anisotropic, due to having a higher percentage of bonded area between the splats. Macroindentation techniques were used primarily to induce cracking to allow the determination of fracture toughness (KjC) values for the coating materials. The Kjc values that were obtained, for the denser coatings only, were used primarily as a means of comparing the coating behaviour, for example, when combined with hardness in the brittleness index. 9.1.2 Controlled Scratch Testing 9.1.2.1 Planar Surface tests Controlled scratching using a conical tipped diamond indenter on the planar (top) surfaces of the coatings proved to be one of the most meaningful, and simple to use, micromechanical tests for PS coatings. Profilometric measurements of the scratch groove allowed estimates of the cross-sectional area, which could also be converted into a measure of wear volume. The wear loss was found to increase proportionally with the square of the load, i.e. WocF„\ The groove area was found to correlate closely with the porosity levels measured in the 191 coatings, as well as with the abrasive wear results. In addition, the scratch hardness was found to correlate better with the wear results than the Vickers hardness. Scratch testing is essentially single point abrasion and as such is useful for comparing the wear resistance of different materials. Scratching with a diamond indenter at loads ranging from 10-60N corresponds to severe abrasion as the load per indenter/particle is much higher than that found in typical abrasion tests, which in the present work was less than lN/particle. The main differences in the wear mechanisms between the coating groups in the single pass scratches from 0-60N are summarized in Table 9.1. The major difference between the Group 4 and Group 1 coatings, i.e those with low and high internal cohesion, or integrity, repectively, was the formation of a thick, compressed wear sheet or layer of fractured and sheared particles. This delamination layer has been seen to occur in bulk ceramic materials [51], and appears quite sensitive to the microstructural properties of the material, such as the inter-splat bond strength in coatings as well as grain boundary strength. The formation of such a wear sheet involves the processes of densification of porous areas, plastic flow, shear and microcracking. Macrocracking was observed in the denser Group 2 coatings, AT, AA10-3,4, and AA18-3,4, after multiple scratches. Table 9.1 Scratch Wear Characteristics (Load 0-60N with 100 pm radius indenter). Coatings Wear Rate Wear Mechanisms Brittle surface fracture (began at 40-50N) AT's and AA10-1,2 - No evidence of catastropic brittle fracture, fracture visible only at edges of scratch grooves. Macrocracking occurred in AT, AA10-3,4 and AA18-3.4 coatings after multiple (n = 5) passes. Rest of Group 2 - Some wear sheet (compressed) formation, with detachment occurring between 40-60N. Definite (compressed) wear sheet formation, with detachment occurring from 50-60N. Thick compressed wear sheet, however detachment only occurred after more than one pass. Increasing | Cohesion Group 1 Low Group 2 Low-Medium Group 3 Medium Group 4 High 192 9.1.2.2 Cross-Sectional Surface Tests The length of the cone formed when a Vickers indenter scratches across the polished coating cross-section and breaks over the top surface was found to be sensitive to microstructural differences between the coatings. Using the bulk alumina as a reference material and the method described in [94] to determine fracture toughness values for the coatings resulted in unrealistically high values for the majority of the coatings. The use of an isotropic bulk alumina is not the best choice for a calibration material as the lamellar microstructure of the coatings is definitely anisotropic. The morphology of the cone fracture surface provided insights into the lamellar nature of the microstructures. 9.1.3 Abrasion The use of low load, "hard" particle abrasion to compare the coatings was found to be the most sensitive method to differences in microstructure, highlighting the coatings with increased porosity and decreased inter-lamellar bonding, etc., as shown in Fig. 8.23. The wear characteristics of the different coatings are summarized in Table 9.2. Plastic deformation (p.d.) mechanisms were widespread on the denser coatings, along with some microfracture. The amount of microfracture and splat debonding increased as the internal cohesion of the coatings decreased, as observed in the Group 4 coatings. Table 9.2 Abrasive Wear Characteristics of the Evaluated Materials. Wear Mechanisms Increasing Cohesion Coatings Wear Rate 20 pm SiC abrasive 200 pm SiC abrasive Group 1 Very Low p.d./ cutting (small grooves) p.d./ cutting (larger grooves) and some microfracture Group 2 Low -Medium p.d./cutting (medium grooves) & some micro-fracture p.d./cutting (larger grooves) & some microfracture Group 3 Medium p.d./cutting & some microfracture p.d. (larger grooves) & some microfracture Group 4 High p.d., microfracture and splat debonding p.d., microfracture and extensive splat debonding 193 In block-on-drum abrasion, the load per particle is considerably less than that typically used in the scratch test, i.e. 0.5-l.ON/particle versus 10-60N load on the scratch indenter. Thus, a major difference in wear behaviour was seen. The main wear mechanisms seen in low load abrasion of the ceramic PS coatings were plastic deformation and fatigue cracking, with splat debonding occurring in the less cohesive coatings. The strength of the splat interfacial bond is crucial as it determines the size of the fracturing splat fragment, which eventually releases as wear debris. The splat size and bond strength, along with the presence of pre-existing microcracks and pores, essentially determine the amount of material loss in that coating, i.e. the larger the splat size, and/or the distance between intra-lamellar microcracks, as in the AL30 coatings, the greater the wear loss when splat debonding and delamination occur. 9.1.4 Erosion The coatings were evaluated using hard particle erosion (75, 200 and 600 pm angular SiC particles) at two angles of impingement, 45 and 90 degrees, which covers a wide range of contact types and sizes. Erosion with small particles, and at low angles of impact, i.e. 45 degrees, exerts a strong tangential force component, similar to that found in abrasion, while erosion with large particles and especially at a 90 degree impact, concentrates the impact force normal to the target surface. The wear characteristics of the different coatings are summarized in Table 9.3. The wear mechanisms predominant in specimens eroded with small erodent particle sizes (75 and 200 pm in this work), along with small particles (75 pm) at 90 degree impact, were the same as those found in the hard particle abrasion results. The amount of material removed, in ng/impact, corresponded to less than one splat per impact for the majority of coatings when using the 75 and 200 pm particles. The use of large, 600 pm particles introduced a larger normal force with a subsequent increase in brittle fracture and splat debonding, with the equivalent of more than one splat removed per impact on all of the coatings. 194 Table 9.3 Erosive wear characteristics of the evaluated materials. Wear Mechanisms Cohesion Materials 75 pm erodent 200 pm erodent 600 pm erodent i High L Group 1 p.d.* p.d. and some microfracture p.d. and some microfracture Group 2 p.d. p.d. and some microfracture p.d., microfracture and some splat debonding Group 3 p.d. p.d., microfracture and some splat debonding p.d., microfracture and splat debonding Low Group 4 p.d., some splat fracture and debonding p.d., splat fracture and debonding p.d., microfracture and extensive splat debonding Wear Increasing Rate w A High Group 1 p.d. and some microfracture microfracture brittle fracture and spalling** 9f£ Group 2 Group 3 p.d. and some microfracture p.d. and some microfracture microfracture and some splat debonding microfracture and some splat debonding splat fracture and debonding splat fracture and debonding Low Group 4 p.d., some splat fracture and debonding splat fracture and debonding splat fracture and extensive debonding Wear Rate Increasing 1> Wear Rate Low High Low High *(p.d. = plastic deformation) **(AT's exhibited this as well.) A transition in the wear mechanisms occurred in normal impact erosion with larger particles (200 and 600 pm). Microfracture and splat debonding were the primary wear mechanisms under these conditions. The high kinetic energy caused by the normal impact of the 600 pm particles caused additional mechanisms of macrocrack formation and extensive spalling in the denser CR and AT coatings, with densification and crack deflection resulting in lower wear losses in the denser AA-alumina coatings (AA10-3,4 and AA18-3,4). Therefore, for evaluating the microstructural cohesion of the coatings as defined by the amount of splat debonding or delamination, medium particle size (200-300 pm) erosion at a 90 degree angle of impingement would be the best choice. 195 9.2 Contact Zone Size Range and Effects of the Microstructural Features found in PS Coatings An overview of the sizes of the different affected zones relevant to each type of contact compared to the sizes of the various microstructural features found in ceramic PS coatings is presented in Fig. 9.1. The sizes given for the indentation techniques are the radii of the plastic zone in the different coatings formed underneath the indenter. For the wear techniques the size reported is that occurring in individual particle contacts, i.e. a plastic zone radius of 10 pm for a load of 1.0 N or less per particle in the 200 pm particle abrasion, and the width of the damaged area observed in single impact erosion craters. Wear Techniques Indentation Techniques Microstructural Features _ 90 °, 600 u r n . 90 °, 75 umB Erosion _ 45 °, 600 n m _ 45 °, 75 [im ^ Abrasion (20, 200 | im) JScratching (10-60N) Macro-indentation Microindentation Nanoindentation Macro-porosity Inter-lamellar (Horiz.) Microcracks (Length) Intra-lamellar (Vertical) Microcracks (Length) Splat Thickness Unmelted Particles Intra-lamellar Porosity Pass Thickness Grain Size Splat Diameter - l 1 1 0.1 1 10 10  100 S i z e R a n g e [ m i c r o n s ] Fig. 9.1 Comparison of size ranges involved in evaluated forms of contact to the size ranges of the different microstructural features. Based on the results in the present work, a listing of the microstructural features in order of their decreasing effect on the wear behaviour of the coatings is presented in Table 9.4. The first four features are closely linked, as exemplified by Coating Group 4. These coatings had 196 very high porosity, a high density of inter-lamellar cracks (i.e. poor bonding), a large number of unmelted particles, along with poorly flattened splats. In these coatings, the first four factors were the deciding ones in determining the total amount of material loss during the different contact situations. As these factors cease to be a critical issue, i.e. as the coating microstructure becomes denser, with increased inter-lamellar bonding and cohesion, the importance of the other factors increases. In the Group 1 coatings, for example, with microstructures similar to bulk ceramics, the smaller scale defects, intra-lamellar microcracking and porosity, along with the strength of the columnar grain boundaries, begin to play a larger role in determining the wear performance of the coatings. Table 9.4 Ranking of the Effect of Microstructural Parameters on Wear by Hard Particles. Microstructural Parameters Size Range (pm) Effect on Wear Performance* Macro-porosity (voids) 10-30 1 Inter-lamellar (horizontal) microcracking (related to the level of inter-splat bonding) 5-50 1 Splat morphology (shape factor (i.e. width/thickness dimensions) Diameter: 10-100 Thickness: 1-3, Pass thickness: 3-10 2 Unmelted Particles 5-50 2 Intra-lamellar microcracks (vertical microcracking) 1-10 3 Intra-lamellar porosity 0.1-1.0 4 Columnar grain boundaries: strength, width 0.2-1.0 5 Increasing damage * [Ranking: 1 (most) to 5 (least) effect.] 197 9.3 Conclusions The wide range of microstructural defects ultimately determines the mechanical, tolerance limits of PS coatings. Any investigation into the cohesive properties or wear response of such coatings needs to consider the effects of defects of different types and sizes. Depending on the size of the surface damage or stress fields caused by the test method, different aspects of the coating microstructure will affect the material performance in that test. A l l of the micromechanical and wear methods evaluated in the present work were sensitive to differences in the coating microstructures to varying degrees. The following is a summary of the main findings of this work, as related to the microstructural aspects of the coatings and the micromechanical techniques evaluated: I. The microstructural features with the most influence on the behaviour of ceramic PS coatings during contact, or wear, by hard particles include, in order of importance: 1) macro-porosity, 2) horizontal crack density, 3) degree of flattening of the splats and 4) volume of unmelted particles, which are all linked to the level and strength of interlamellar bonding in the coating. II. The level of porosity in the coatings showed the strongest effect on the amount of abrasive wear due to plastic deformation, as evidenced by the strong correlation of the form W = kPn, with n varying from 0.85-0.95 for both the low load abrasion and controlled scratching tests. III. The hardness of the coatings was found to be inversely proportional to the level of porosity in the coatings, as well as to the abrasive wear rate. IV. The horizontal crack density in the coatings exhibited a similar effect on the extent of abrasive wear damage as did the porosity, with a somewhat weaker correlation (R2 = 0.6 vs. 0.8). V. Coatings with no visible lamellar structure exhibited less anisotropy in the elastic modulus values determined by Knoop indentation than the coatings with a distinct lamellar 198 structure. Because of the different indenter shapes, Knoop hardness was found to be a better indicator of surface hardness than was Vickers hardness, and correlated better with the wear results. VI. The major effect of the inter-lamellar bonding in ceramic PS coatings is seen in the wear mechanism transitions. As the level of inter-splat bonding in the coating decreases, the contact load at which the transition from 'quasi-plastic' deformation to splat fracture and delamination occurs does as well. However, for coatings with greater internal cohesion, the load at which catastrophic brittle fracture and spalling occur is increased. VII. In abrasive wear situations, including low load abrasion and scratching as well as small particle erosion, the level of inter-lamellar bonding, the spacing between vertical cracks, and the thickness of the splats are critical in determining the amount of material removed. The low load abrasion results showed the most sensitivity to the microstructural differences of the coatings, followed by controlled scratching. VIII. Excellent correlations were found between the inter-crack spacing (the reciprocal of the vertical crack density) of coatings with well-flattened splats, and the scratch width at different loads (single point abrasion), indicating the operation of a size effect. LX. The scratch hardness of the coatings was a better indicator, i.e. gave a better correlation with wear than the Vickers hardness, but not quite as good as the Knoop hardness. The scratch material losses (Wsc), as measured by cross-sectional area, were found to increase proportionally with the square of the load (Fn), Wsc = Fn2. X . A well-developed lamellar microstructure, incorporating an extensive network of horizontal and vertical microcracks, as seen in the denser AA-alumina coatings (AA10-3,4 and AA18-3,4), can absorb the impact forces experienced in large particle normal impact erosion through processes of densification and crack deflection, resulting in lower wear losses. 199 XL Macrocracking and extensive spalling mechanisms, similar to those seen in bulk ceramics, were observed in the dense ceramic PS coatings, including alumina-titania and chromia, after the high energy impact events that occur in large particle normal impact erosion. XII. The hardness and fracture toughness of ceramic PS coatings alone cannot be used to predict their abrasive or erosive wear behaviour. A microstructural parameter (M) is needed in order to properly model the wear performance of the coatings. It needs to include a combination of the porosity (P) and some measure related to the spacing between the vertical microcracks (V) and the strength and extent of inter-lamellar bonding (B). 9.4 Recommendations for Future Work Additional studies concerning the microstructural defects in ceramic PS coatings and their evaluation using these micromechanical tests that would be worthwhile include: I. Further investigations using the method of controlled scratching, especially on coating top surfaces, for studying the abrasion behaviour of PS coatings. A range of different microstructures should be evaluated using multiple scratches (n = 1, 2, 5, 10, 15, ...) to better simulate the wear process. II. Further definition of the relationship between coating porosity and scratch hardness needs to be established to determine its use as a process control measure. III. Additional intra-splat studies using depth sensing indentation (DSI) with lower loads, to better evaluate the intrinsic material properties of individual splats. A large number of tests should be conducted to improve statistical confidence in the results. IV. Develop a microstructural quality index with standards for different types of coatings and applications - including extent, and size and shape distribution of porosity, microcrack distribution and density, percent actual interlamellar bonding, number and volume of unmelted particles, along with a rating of which are the most critical microstructural defects. 200 9.5 New Knowledge and Major Contributions The unique contributions arising from the present work are summarized below. (a) For the thermal spray research community: - The use of mono-crystalline sapphire alumina powders, with narrow particle size ranges, to reduce the number of variables in the PS process. - The use of scratch testing to characterize the coatings, in the areas of both wear and porosity. - The influence of the microstructure on the occurance of the wear transitions, laying the groundwork for future development of wear mechanism maps. - A compilation of microstructural data and wear results for a variety of coatings, leading to the development of several wear correlations as well as descriptive wear models for the coatings. This is an advance towards a more comprehensive understanding of the wear mechanisms of ceramic PS coatings as well as the evolution of better wear models, which can also lead to improved microstructural standards for different coatings. (b) For the thermal spray industrial community: - Scratch hardness correlates better with wear behaviour than does the Vickers hardness which is presently in use, as well as correlating directly with the amount of coating porosity. Further development of this technique could lead to its use as a standarized method for quality control. - Tests like this can help shorten the amount of time required for the development of new and better coatings for different applications. - The use of the sapphire alumina powders produces optimized coatings superior to those sprayed with polycrystalline alumina powders. 201 PUBLICATIONS The following papers have been prepared as a direct result of the work leading up to this thesis: A. L .C . Erickson, H . M . Hawthorne and T. Troczynski, "Scratch Testing of Ceramic Plasma Sprayed Coatings", To be published in Proceedings of United Thermal Spray Conference, (1999) March 17-19, Dusseldorf, Germany. B. L.C. Erickson, T. Troczynski, H . M . Hawthorne, H . Tai, and D.Ross, "Alumina Coatings by Plasma Spraying of Monosize Alumina Particles, in Proceedings of International Thermal Spray Conference Nice, France (1998) 791-796 (Best Paper Award) and accepted for publication in Journal of Thermal Spray Technology (1998). C. R. Westergard, L . C. Erickson, N . Axen, H. M . Hawthorne and S. Hogmark, "The Erosion and Abrasion Characteristics of Alumina Coatings Plasma Sprayed under Different Spraying Conditions", Tribology International, 31 (1998) 271-279. D. L . C. Erickson, R. Westergard, U . Wiklund, N . Axen, H. M . Hawthorne and S. Hogmark "Cohesion in Plasma-Sprayed Coatings - A Comparison between Evaluation Methods", Wear, 214(1998 ) 30-37. E. H . M . Hawthorne, L .C. Erickson, D. Ross, H. Tai and T. Troczynski, "The Microstructural Dependence of Wear and Indentation Behaviour of Some Plasma Sprayed Alumina Coatings", Wear, 203-204 (1997) 709-714. F. L .C. Erickson, H . M . Hawthorne, T. Troczynski, S. Hogmark and M . 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Soc, 51 (1968)433-439. 210 App. A APPENDIX A - Procedure for Mounting and Polishing of Ceramic Plasma Sprayed Coatings I. Mounting of Plasma Sprayed Coatings A. Equipment and Materials - Small vacuum chamber plus vaccum pump (to 100 mbar) [set in fume hood!]* - Slow curing epoxy resin and hardener (Struers Epofix or Buehler Epoxide) - Plastic gloves and face mask - Paper (not wax-coated) mixing cups & mixing stick (wooden) - Plastic mounting cups & liquid release agent (to prevent sticking) - Samples cut to fit into cups - Oven set at 60-65 °C - Microbalance B. Coating Infiltration Procedure 1. Clean sample well with water followed by acetone or ethyl alcohol to rid pores of debris and water. 2. Place PS samples into oven for 20-30 minutes, depending on their size. 3. Measure correct amount of epoxy resin for 3-4 samples at a time into a paper cup (not wax-coated) and place into oven for 15-20 minutes until viscosity is quite low. Check viscosity of epoxy with wooden mixing stick. 4. Coat mounting cups with release agent and set them into the vacuum chamber. 5. Place samples into mounting cups, with coating side up for cross-sections, down for top ones. 6. When the epoxy resin has reached the set temperature, add the correct amount of epoxy hardener to the resin and stir slowly for 2 minutes, (to minimize the formation of bubbles) Mix inside fume hood (dangerous fumes!!).* 7. (a) If using the Struer's system, set the cup into the holder and place the vacuum tube (clamped) into the mixture. Place the samples into the vacuum chamber, pump to 100 mbar for a minute to evacuate air from cracks and pores. (b) Slowly release the clamp and let epoxy into the tube, then slowly fill each sample cup until the coating is just covered. 8. If using a simple vacuum jar, pour the epoxy into the samples cups, slowly filling each cup until the coating is just covered. 9. Let the chamber pump until bubbles just begin to form on the top of the epoxy (typically 100 mbar). Hold the vacuum at this level for 1 minute, then bring up to std. pressure, pump down again and repeat twice. 10. Bring to atmospheric pressure (quickly to eliminate all bubbles) and let cure (overnight). C. Final Mounting (for cross-sections) Procedure: 1. For cross-sections, saw the infiltrated sample in half using a diamond saw (after first removing excess epoxy). 2. Place the samples cut face down in a mounting cup (coated with release agent). 211 App. A 3. Fil l with either mixed epofix epoxy resin (unheated) and hardener or a faster curing epoxy and allow to cure. II. Polishing of Ceramic PS Coatings A. Equipment and Materials - Automated polisher (Struers Abrapol or Buehler...) - SiC papers, 240 and 320 grit - Diamond Suspension (or Spray), 6 & 3 pm - 0.05 pm Si02 polishing supension - Polishing Cloths: hard surface (Buehler Ultra-pad or Struers DP-Dur) med. surface (Buehler Texmet 1000 or Struers Nylon) soft surface (Buehler or Struers final polishcloth ?) - Alcohol - Distilled Water - Soft brush for cleaning samples - Metallographic microscope B. Coarse Grinding Stage 1. Mount samples, with edges smoothed using abrasive paper, in holder preferably 3 at a time. If polishing cross-sections, be sure to mount so that the coating top surface is facing the direction of rotation, * so that the abrasive particles are swept across the coating first and then into the substrate. 2. Use 240 grit SiC paper for flattening, 5 lbs. pressure per sample (15 lbs. for 3 samples), run for 30-40 sec. 3. Rinse samples and check for flatness. Repeat if necessary. On 2nd repeat, use 320 grit paper. 4. Rinse samples thoroughly, using a soft brush with soap in distilled water to eliminate all traces of abrasive particles. C. Sample Integrity stage 1. Use 6 pm diamond suspension on a hard pad at 5 lbs. pressure per sample for 5 minutes. 2. Rinse samples well and check surface for size of scratches. Clean pad with alcohol after use. 4. Use 3 pm diamond suspension on a medium hard pad at 6 lbs. pressure per sample for 3 minutes. 5. Rinse samples well and check coating surface for size of scratches to ensure that those from the previous stage are gone. Examine under metallographic microscope. 6. Clean pad with alcohol after use. D. Final Polishing Stage 1. Use 0.05 pm Si02 polishing supension on a soft polishing cloth at 8 lbs. pressure per sample for 3 minutes. 2. Rinse samples well and check coating surface for size of scratches to ensure that those from the previous stage are gone. Examine under metallographic microscope. 3. When desired surface finish is achieved, remove samples from the sample holder. 212 A P P E N D I X B - Scratch Wear Mechanism Maps Fig. 6.— Summary map of the different response regimes for the different samples vs. load (a) AA- alumina coatings and (b) reference materials.* *(Summarizes scratch mechanisms from Section 6.3.1.2.) 213 A p p . C APPENDIX C - Correlations between the microstructural parameters, micromechanical properties, and wear performance (Chapter 8) 8.2 Porosity Table 1 Power correlation coefficients for porosity versus measured wear rates. Wear Tests (particle sizes) Corr. Coeff. (R2 Value) Exponent (a) Abrasion 20 pm 0.812 0.84 200 pm 0.826 0.95 60N Planar Scratch Area 0.753 0.59 45 degree erosion 75 pm 0.743 0.49 200 pm 0.730 0.48 600 pm 0.658 0.49 90 degree erosion 75 pm 0.747 0.50 200 pm 0.657 0.42 600 pm 0.558 0.39 400 350 „ 300 E 2 0 0 s 1 1 5 0 LU 100 50 E, (1 ON)-Top O SA 0.0 2.0 4.0 6.0 8.0 10.0 12.0 Porosity [%] Fig. 8.4 Effect of porosity on (Knoop) Elastic Modulus {Ek) at 10N measured on planar (top) surfaces. 214 App. C 20 18 16 14 °- 12 JioH C •a n x 6H 4 2H 0 H, (100 N)-Top 20 18 16 14 S 12 i" 10 — i 1 1 1 1 1 I 1 1 0.0 1.0 2.0 3.0 4.0 5.0 6.0 7.0 8.0 9.0 H,(100N)-Top Porosity [%] 0.0 1.0 2.0 3.0 4.0 5.0 6.0 7.0 8.0 9.0 Porosity [%] Fig. 8— Porosity vs. Hardness (Hv) at 100N for: (a) all materials and (b) alumina-based coatings only. "E E 150.0 6.0 Porosity [%] Fig. 8 — Correlation between porosity and the 200 pm abrasion rate for all materials. 215 A p p . C 45° Angle 90° Angle 600 um 0.050-1 • » 6.0 Porosity [%] 200 um 4.0 6.0 8.0 Porosity [%] 75 nm 6.0 Porosity [%] 12.0 12.0 12.0 2.0-1 1.8 -1.6 • m 1.4 • "e £ 1 2 ' 01 n 1.0 • cc | 0.8 -W O ilj 0.6 -0.4 -0.2 • 0.0 • 600 um 0.0 0.2-^  • 6.0 Porosity [%] 200 nm 4.0 6.0 8.0 Porosity [%] 75 um 4.0 6.0 8.0 Porosity [%] 12.0 12.0 12.0 Fig. 8.-- Porosity vs. erosion rate at a 45 and 90 degree impingement angle using: (a) 600 pm, (b) 200 pm and (c) 75 pm erodent particles. 216 App. C 8.3 Inter-lamellar (Horizontal) Microcracking Table 8.- Power Correlation coefficients for horizontal crack density versus the measured wear rates. Correlation Coeff. (R2 values) Wear Methods All Coatings w/out AL30's -> exponent (a) Abrasion 20 urn 0.670 0.655 0.82 200 pm 0.695 0.734 0.93 60N Scratch Area 0.631 0.606 0.58 3.5N X-Sec Scratch Cone Length 0.460 0.500 0.18 45 degree erosion 75 pm 0.567 0.579 0.57 200 pm 0.478 0.623 0.48 600 pm 0.434 0.508 0.41 90 degree erosion 75 pm 0.509 0.576 0.48 200 pm 0.378 0.429 0.31 600 pm 0.126 0.054 0.16 217 App. C 45° Angle 90° Angle 600 um O AL30-1 AL30-4 a • » AAA5-4 50 100 150 Horizontal Crack Density [#/mm] OAPS 2.00 • 1.80 • 1.60 • 1.40-1 1.20 H <3 1.00 cc o 0.80 2 0.60 -0.40 -0.20 -0.00 • AL30-1 AL30-2 • • AA5-1 • AL30-4 A A AA5-4 600 um 50 100 150 Horizontal Crack Density [#/mm] • AL30-1 • AL30-2 ± AA5-1 50 100 150 Horizontal Crack Density [#/mm] 200 um O APS E — 0.80 c 0.60 o 'in p UJ 0.40 0.20 0.00 •AL30-1 A AA5-4 200 um 50 100 150 Horizontal Crack Density [0/mm] 50 100 150 Horizontal Crack Density [#/mm] 75 um OAPS • AL30-2 75 um O APS 50 100 150 Horizontal Crack Density [#/mm] Fig. 8.-- Horizontal crack density vs. erosion rate at a 45 and 90 degree impingement angle using: (a) 600 pm, (b) 200 pm and (c) 75 pm erodent particles. 218 8.3 Intra-lamellar (Vertical) Microcracking App. C 40 60 Vertical Crack Density [#/mm] 200 1 150 I—I 100 OAPS 40 60 80 Vertical Crack Density [#/mm] Fig. 8.- Effect of vertical crack density on the abrasion rate using: (a) 20 um and (b) 200 um abrasive particles. Table 8.5 Linear correlation coefficients for vertical crack density versus the wear rates. Correlation Coeff. (R2 values) Wear Methods All coatings w/out AL30 & APS Exponent (a) Abrasion 20 pm 0.292 0.460 0.53 200 pm 0.313 0.550 0.54 60N Scratch Area 0.530 0.746 0.64 3.5N X-Sec Scratch Cone Length 0.287 0.395 0.16 45 degree erosion 75 pm 0.203 0.304 0,29 200 pm 0.262 0.526 0.30 600 pm 0.463 0.797 0.42 90 degree erosion 75 pm 0.170 0.274 0.24 200 pm 0.150 0.255 0.17 600 pm <0.10 <0.10 — 219 App. C 45° Angle 90° Angle • AL30-1 • AL30-2 O APS 40 60 Vertical Crack Density [#/mm] 2.00 1.80 1.60 i 1.40 1.20 1.00 0.80 0.60 0.40 0.20 0.00 A AA5 • • AL30-1 • AL30-2 40 60 Vertical Crack Density [#/mm] 80 100 40 60 Vertical Crack Density [#/mm] 1.40 1.20 ™ 1.00 "E — 0.80 £ cc c 0.60 o "3 o LU 0.40 0.20 0.00 40 60 Vertical Crack Density [#/mm] 0.160 n 0.140 0.120 H I 0.100 H a a 0.080 -cc C | 0.060 -o ui 0.040 -0.020 -0.000 40 60 Vertical Crack Density [tt/rnm] g 0.40 •] • AL30-1 • AL30-2 40 60 Vertical Crack Density [#/mm] Fig. 8.— Vertical crack density vs. erosion rate at a 45 and 90 degree impingement angle using 600 pm, (b) 200 pm and (c) 75 pm erodent particles. 220 8.5.1 Effect of Hardness on the Wear Rates App. C Table 8.6 Inverse power correlation coefficients for hardness versus the measured wear rates, stands for perpendicular to the coating surface, on the coating cross-section.) Correlation Coefficients Wear rates Hk-Top Hk-Pd. HV(3N) Abrasion 20 pm 0.878 0.866 0.784 200 pm 0.786 0.785 0.612 60N Scratch Area 0.672 0.605 0.655 45 degree erosion 75 pm 0.813 0.728 0.670 200 pm 0.725 0.694 0.478 600 pm 0.648 0.625 0.405 90 degree erosion 75 pm 0.782 0.718 0.598 200 pm 0.676 0.684 0.402 600 pm 0.579 0.579 0.290 Table 8.7 Power Exponents (a) for hardness versus the measured wear rates. Exponents Wear rates Hk-Top Hk-Pd. HV(3N) Abrasion 20 pm 2.70 3.03 1.67 200 pm 2.90 3.28 1.68 60N Scratch Area 1.72 1.84 1.11 45 degree erosion 75 pm 1.57 1.72 0.94 200 pm 1.55 1.71 0.82 600 pm 1.55 1.68 0.80 90 degree erosion 75 pm 1.63 1.76 0.93 200 pm 1.35 1.53 0.68 600 pm 1.24 1.40 0.57 221 App. C 45° Angle 90° Angle 2.00 -j 1.80 1.60 •a '•4 0 a | 1.20 o s LOO tr.. ° 0.80 ta g W 0.60 0.40 0.20 H 0.00 6 8 10 Hk [GPa] 2 4 6 8 10 12 14 16 Hk [GPa] 8 10 Hk [GPa] E £ 0.80 200 um 14 16 Hk [GPa] 8 10 Hk [GPa] 12 14 •? 0.80 J • A T • Hk [GPa] Fig. 8.— Knoop Hardness vs. erosion rate at a 45 and 90 degree impingement angle using: (a) 600 pm, (b) 200 pm and (c) 75 pm erodent particles. 222 App. C 8.5.2 Correlation of Scratch Hardness with Wear Rates Table 8.8 (Inverse) Power Correlation Coefficients for Scratch Hardness at 60 N versus the measured wear rates. Wear rates (Particle size) R2 value Exponent a Abrasion 20 um 0.696 1.23 200 pm 0.679 1.38 45 degree erosion 75 pm 0.626 0.71 200 pm 0.736 0.80 600 pm 0.619 0.78 90 degree erosion 75 pm 0.643 0.76 200 pm 0.570 0.63 600 pm 0.364 0.50 0.25 0.2 0.15 0.1 a < 0.05 • AL30-1 DAL30-2 AA5-1 A O APS OSA 0.0 2.0 4.0 6.0 8.0 10.0 Scratch Hardness [GPa] 12.0 14.0 Fig. 8— Inverse correlation between the scratch hardness (at 60N) and the abrasive wear rate using 200 pm particles. 223 App. C 45° Angle 90° Angle AA5-1 A 600 um OSA 4.0 6.0 8.0 10.0 Scratch Hardness [GPa] 2 l 1.8 1.6 H J? 1.4 "E <S 1 • cc o 0.8-co O LU 0.6 -0.4 -0.2 -0 0.0 • AL30-1 AA5-1A\ * A P S 600 um 4.0 6.0 8.0 10.0 Scratch Hardness [GPa] O S A OSA 4.0 6.0 8.0 10.0 Scratch Hardness [GPa] c 0.6 o o LU 0.4 0.2 OAL30-1 AAS-1 A 200 um OSA 4.0 6.0 8.0 10.0 Scratch Hardness [GPa] 0.16 0.14 H 0.12 0.1 0.08 0.06 0.04 0.02 0 0.0 I- 8.-O AL30-1 4.0 6.0 8.0 10.0 Scratch Hardness [GPa] • AL30-1 75 um OSA 4.0 6.0 8.0 10.0 Scratch Hardness [GPa] Scratch Hardness vs. erosion rate at a 45 and 90 degree impingement angle using: (a) 600 pm, (b) 200 pm and (c) 75 pm erodent particles. 224 

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