Chapter 1 Introduction………………………………………………… 1

Home , Cathodic arc deposition

The Pennsylvania State University

The Graduate School

Intercollege Graduate Program in Materials Science and Engineering

SYNTHESIS-STRUCTURE-PROPERTY-PERFORMANCE RELATIONSHIPS

OF TiN, CrN, AND NANOLAYER (Ti,Cr)N COATINGS DEPOSITED BY

CATHODIC ARC EVAPORATION FOR HARD PARTICLE EROSION

RESISTANCE

A Thesis in

Material Science and Engineering

Brian M. Gabriel

Submitted in Partial Fulfillment of the Requirements for the Degree of

Master of Science

May 2009

The thesis of Brian M. Gabriel was reviewed and approved* by the following:

Douglas E. Wolfe Assistant Professor of Material Science and Engineering Thesis Advisor

John R. Hellmann Professor of Materials Science and Engineering

Suzanne Mohney Professor of Materials Science and Engineering

Joan Redwing Professor of Materials Science and Engineering Chair, Intercollege Materials Science and Engineering Graduate Degree Program

*Signatures are on file in the Graduate School

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ABSTRACT

Hard particle erosion due to sand particles can cause significant damage to

metallic turbine engine components. Hard coating systems such as TiN, CrN, and

(Ti,Cr)N offer a potential solution. TiN, CrN, and (Ti,Cr)N coatings were deposited with

a multi-source cathodic arc system using high purity Ti and Cr cathode targets in a partial

N2 atmosphere. The coatings were characterized using X-ray diffraction, scanning

electron microscopy, electron probe microanalysis, scanning transmission electron

microscopy, and scratch adhesion testing. Erosion testing was performed using an in-

house high-velocity erosion rig with glass beads and alumina media. For the TiN

coatings, erosion resistance was strongly dependent on the evaporator current and

substrate bias; these parameters influenced crystallite size, preferred crystallographic

orientation, and residual stress. CrN coatings were determined to have significantly more

macroparticles than the TiN coatings deposited under similar conditions. This was

primarily attributed to the lower melting point of the solid phases in the Cr-N system

versus the Ti-N system. A nanolayered (Ti,Cr)N coating system comprised of alternating

TiN and CrN rich layers was created by co-evaporating Ti and Cr cathode targets with a

rotating substrate configuration. Erosion resistance increased along with decreasing

density of nanolayer interfaces as well as increasing volume percentage of the CrN rich

layers with respect to the TiN rich layers. In all three coating systems, macroparticle defect concentrations were not believed to degrade high impact angle erosion

performance.

TABLE OF CONTENTS

LIST OF FIGURES………………………………………………………………. ix

LIST OF TABLES………………………………………………………………... xvii

ACKNOWLEDGEMENTS………………………………………………………. xiii

CHAPTER 1 INTRODUCTION………………………………………………… 1

1.1 Background………………………………………………………...... 1

1.2 Project Objectives……………………………….………………………. 1

CHAPTER 2 LITERATURE REVIEW AND COATING THEORY…………... 3

2.1 Cathodic Arc Deposition………………………………...………………. 3

2.2 Mechanisms of Erosion……………………………..…………………… 9

2.2.1 Solid Particle Erosion through Plastic Deformation…..………. 11

2.2.2 Solid Particle Erosion of Brittle Materials……..……….……... 12

2.2.3 Erosion Resistant Coatings…………...…….………………….. 16

2.3 TiN, CrN, and (Ti,Cr)N Coatings for Erosion Protection……………….. 27

CHAPTER 3 EXPERIMENTAL PROCEDURE………………………………... 35

3.1 Cathodic Arc Deposition Equipment……………………………………. 35

3.2 Substrate Preparation…………………………………………………….. 39

3.3 Cathodic Arc Deposition Processing Procedure………………………… 41

3.4 Design of Experiments for Deposition of TiN, CrN and (Ti,Cr)N Coatings………………………………………………………………….. 43

3.5 Coating Characterization and Evaluation………………………………... 54

3.5.1 X-Ray Diffraction (XRD)……………………………………... 54

3.5.1.1 X-Ray Diffraction for Crystallographic Structure Determination………………………………………... 54

3.5.1.2 XRD Crystallite Size……………………………….... 56

3.5.1.3 XRD Residual Stress Analysis………….…………… 57

3.5.2 Cross Sections for Optical Microscopy (OM)……………….… 58

3.5.3 Electron Probe Microanalysis (EPMA)……….……………….. 59

3.5.4 Environmental Scanning Electron Microscopy (ESEM).……... 59

3.5.5 Quantative Microstructural Analysis.…………………………. 60

3.5.6 Annular Dark Field Scanning Transmission Electron Microscope (ADF-STEM)…….………………………………. 62

3.5.7 Vicker’s Micro-Indention Hardness Testing………………….. 62

3.5.8 Scratch Adhesion Testing……………………………………... 63

3.5.9 Surface Roughness Measurements……………………………. 63

3.5.10 Erosion Testing………………………………………...... 65

CHAPTER 4 RESULTS AND DISSCUSSION…………………………………. 70

4.1 XRD Crystallographic Structure Determination from Θ/2Θ and Glancing Angle Scans…………………………………………………… 70

4.1.1 XRD Crystallographic Structure Results for Monolithic TiN… 71

4.1.2 XRD Crystallographic Structure Results for the CrN Coating... 74

4.1.3 XRD Crystallographic Structure Results for Nanolayer (Ti,Cr)N Deposited as a Function of Evaporator Current (Constant Ti Evaporator Current)……………………………... 75

4.1.4 XRD Crystallographic Structure Results for Nanolayer (Ti,Cr)N Deposited as a Function of Substrate Bias (Ti at 65 A and Cr at 45 A Evaporator Current)…………………………… 78

4.1.5 XRD Crystallographic Structure Results for Multilayer (Ti,Cr)N Coatings with Ti and Nb Interlayers………………… 80

4.2 Crystallite Size…………………………………………………………... 82

4.2.1 Trends in XRD Crystallite Size of Monolithic TiN Coatings

Deposited as a Function of Evaporator Current (Constant -150 V Substrate Bias)………………………………………………. 83

4.2.2 Trends in XRD Crystallite Size of Monolithic TiN Coatings Deposited as a Function of Substrate Bias…………………….. 84

4.2.3 Trends in XRD Crystallite Size of Nanolayer (Ti,Cr)N Coatings Deposited as a Function Substrate Bias (Ti Evaporator of 65 A, Cr Evaporator of 45 A)…………………... 85

4.3 Preferred Crystallographic Orientation………………………...... 87

4.3.1 Preferred Crystallographic Orientation of the Monolithic TiN Coatings Deposited as a Function of Evaporator Current (Constant Substrate Bias)……………………………………… 88

4.3.2 Preferred Crystallographic Orientation of the Monolithic TiN Coatings Deposited as a Function of Substrate bias (Constant Ti Evaporator Current)………………………………………… 92

4.3.3 Crystallographic Orientation of the CrN Coating……………... 93

4.3.4 Preferred Crystallographic Orientation of Nanolayered (Ti, Cr)N Coatings Deposited as a Function of Cr Evaporator Current (65 A Ti Evaporator, Constant Substrate Bias)……….. 96

4.3.5 Preferred Crystallographic Orientation of Nanolayered (Ti,Cr)N Coatings Deposited as a Function of Substrate Bias (65 A Ti Evaporator Current, 45 A Cr Evaporator Current)...… 98

4.3.6 Crystallographic Orientation of the (Ti,Cr)N Multilayer Coatings with Ti Interlayers (-150 V Substrate Bias, Ti Evaporator Current of 65 A, Cr Evaporator Current of 45 A)… 100

4.4 Residual Stress Analysis………………………………………………… 101

4.4.1 Trends in XRD Residual Stress Analysis of the Monolithic TiN Coatings Deposited with Different Evaporator Currents (Constant Substrate Bias)……………………………………… 103

4.4.2 Trends in XRD Residual Stress Analysis of the Monolithic TiN Coatings Deposited with Different Substrate Biases (Constant Ti Evaporator Current)……………………………... 105

4.5 Deposition Rate of Nitride Coatings…………………………...………... 107

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4.5.1 Deposition Rate of Monolithic TiN Coatings Deposited at Different Evaporator Currents (-150 V Substrate Bias)……….. 108

4.5.2 Deposition Rate of TiN Coatings Deposited at Different Substrate Biases (65 A Evaporator Current)…………………... 110

4.5.3 Deposition Rate of Nanolayer (Ti,Cr)N Coatings as a Function of Cr Evaporator Current (Ti Evaporator of 65 A, Substrate Bias of -150 V)………………………………………………… 111

4.5.4 Deposition Rate of Nanolayer (Ti,Cr)N Coatings Deposited at Different Substrate Biases (65 A Ti Evaporator Current, 45 A Cr Evaporator Current)……………………...…………………. 112

4.6 Electron Probe Micro Analysis of the Nanolayer (Ti,Cr)N Coatings…… 113

4.7 Quantative Analysis of Large Scale Defects of Nanolayer (Ti,Cr)N Coatings………………………………………………………………….. 115

4.8 ESEM Fracture Surfaces………………………………………………… 118

4.9 STEM of Nanolayer (Ti,Cr)N Coatings Deposited at Different Cr Evaporator Currents (65 A Ti Evaporator Current, -150 V Substrate Bias)……………………………………………………………………... 120

4.10 Vicker’s Micro-Hardness…………………………………. 122

4.11 Scratch Adhesion Testing of Nanolayer (Ti,Cr)N Coatings………...…... 123

4.12 Surface Roughness………………………………………………………. 124

4.12.1 Surface Roughness of the Monolithic TiN and CrN Coatings… 125

4.12.2 Surface Roughness of Nanolayer (Ti,Cr)N Coatings………….. 126

4.13 Erosion Performance…………………………………………………….. 129

4.13.1 Erosion Performance of Monolithic TiN Coatings Deposited with Different Evaporator Currents (Constant Substrate Bias)... 130

4.13.2 Erosion Performance of Monolithic TiN Coatings Deposited with Different Substrate Biases (Constant Ti Evaporator Current)………………………………………………………... 134

4.13.3 Erosion Performance of the Monolithic CrN Coating………… 137

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4.13.4 Erosion Performance of Nanolayer (Ti,Cr)N Coatings Deposited at Different Cr Evaporator Currents (65 A Ti Evaporator Current, -150 V Substrate Bias)…………………... 140

4.13.5 Erosion Performance of Nanolayer (Ti,Cr)N Coatings Deposited at Different Substrate Biases (65 A Ti Evaporator Current, 45 A Cr Evaporator Current)………………………… 144

4.13.6 Erosion Performance of Multilayer(Ti,Cr)N Coatings Deposited at (65 A Ti Evaporator Current, 45 A Cr Evaporator Current, -150 V Substrate Bias)……………………………….. 147

4.14 Proposed Coating Design……...………………………………………… 152

CHAPTER 5 CONCLUSION……………………………………………………. 154

5.1 Monolithic TiN Coatings Deposited as a Function of Evaporator Current (-150 V Substrate Bias)………………………………………………….. 154

5.2 Monolithic TiN Coatings Deposited as a Function of Substrate Bias (65 A Ti Evaporator Current)………………………………………………... 154

5.3 Monolithic CrN Coating Deposited at 45 A and -200 V Substrate Bias… 155

5.4 Nanolayer (Ti,Cr)N Coatings Deposited with Different Cr Evaporator Currents (65 A Ti Evaporator Current, -150 V Substrate Bias)…………. 155

5.5 Nanolayer (Ti,Cr)N Coatings Deposited with Different Substrate Biases (65 A Ti Evaporator Current, 45 A Cr Evaporator Current)…………….. 156

5.6 Multilayer (Ti,Cr)N Coatings Deposited with Ti and Nb Interlayers (65 A Ti Evaporator Current, 45 A Cr Evaporator Current, -150 V Substrate Bias)……………………………………………………………………... 156

REFERENCES….………………………………………………………………… 158

LIST OF FIGURES

Figure 2.1. Depiction of cathodic arc plasma deposition………………..……...... 5

Figure 2.2. System pressure versus nitrogen gas flow rate for (a) reactive cathodic arc evaporation and (b) reactive sputtering techniques...... …………...... 8

Figure 2.3. Cathodic arc plasma filtered through magnetic solenoid in order to remove macroparticles. Image from Brown.10…………………………..……..… 9

Figure 2.4. Sand particle causing platelet formation on a metallic coating…...... 11

Figure 2.5. Lateral crack formation in a brittle material. Adapted from D’Errico.21………………………………………………………………………… 13

Figure 2.6. Erosion rate variation of uncoated and coated INCONEL 718 with impingement angle: T=538°C, V=305 m/s, 20g chromite mass.15…...... 15

Figure 2.7. General erosion rate trends as a function of impact angle of (a) a material which exhibits the ductile erosion mechanism versus (b) a material which exhibits the brittle erosion mechanism. Image from Haugen.27…………... 15

Figure 2.8. Sand particles enter to the left and strike the compressor/turbine blade at various angles. Adapted from Tabakoff.28…...………………………….. 16

Figure 2.9. Leading edge damage to a stainless steel (AM355) first stage compressor blisk due to solid particle erosion. GE T-700 series blisk serial #GWHTFANG………...... 17

Figure 2.10. Thorton structural zone model.39……...………………………..…... 20

Figure 2.11. Proposed multilayer design for increased toughness of erosion resistant coatings……………………..…………………………………………… 23

Figure 2.12. NaCl crystal structure modeled using CrystalDesigner Software. Ti and/or Cr atoms shown as green and N atoms as blue spheres……….…………... 29

Figure 2.13. Phase equilibrium diagrams for (a) TiN and (b) CrN systems.56,57... 30

Figure 2.14. Hexagonal beta-chromium nitride (left) and B1 NaCl cubic chromium nitride (right). Modified from Era.63………………………...………... 31

Figure 2.15. Ternary phase equilibrium diagram for Cr-N-Ti system at 1000 °C.53...... 33

Figure 3.1. Schematic of cathodic arc deposition system………………………... 36

Figure 3.2. Typical “dished” Ti target (2.5” Ø by 1.5” long) for S.E.G 2020 cathodic arc coating equipment…………………………………………………… 37

Figure 3.3. Digital image showing cathodic arc system chamber including the planet/sun substrate holder system configuration used for the deposition of (Ti,Cr)N and CrN coatings. Rotation speeds of 2.4 rpm (sun) and 6.75 rpm (planet) were used for deposition of all coatings…………………………………. 39

Figure 3.4. Schematic representing multilayer (Ti,Cr)N coating system with compliant Ti or Nb interlayers. (Ti,Cr)N is nanolayered structure of alternating TiN and CrN rich layers…………………………………………………………... 46

Figure 3.5. Diffraction of an incident X-ray beam at an incident angle of θ…….. 54

Figure 3.6. Fracture surface methodology for the nitride coating samples………. 60

Figure 3.7. Analytic software routine developed for Clemex Vision P.E. to assess large scale defect incorporation in TiN, CrN and (Ti,Cr)N coatings……… 61

Figure 3.8. Illustration depicting surface roughness, Ra, of a coated substrate….. 64

Figure 3.9. High Velocity hard particle erosion testing equipment designed and built by the ARL Advanced Coating Department at the Pennsylvania State University…………………………………………………………………………. 65

Figure 3.10. Particle size distribution for (a) Size-13 glass bead and (b) 120 grit alumina determined by light scattering…………………………………………… 67

Figure 3.11. ESEM micrograph showing (a) angular particle shape of 120 grit alumina media and (b) spherical particle shape of #13 glass beads………………. 67

Figure 4.1. θ/2θ and glancing angle XRD of TiN coating deposited with an evaporator current set point of 65 A and a -150 V substrate bias. TiN planes are indexed……………………………………………………………………………. 71

Figure 4.2. θ/2θ and glancing angle XRD of TiN coating deposited with an evaporator current set point of 45 A and a -150 V substrate bias. TiN planes are indexed……………………………………………………………………………. 72

Figure 4.3. θ/2θ and glancing angle XRD of TiN coating deposited with an evaporator current set point of 85 A and a -150 V substrate bias. TiN planes are indexed……………………………………………………………………………. 72

Figure 4.4. θ/2θ and glancing angle XRD of TiN coating deposited with an evaporator current set point of 65 A and a -25 V substrate bias. TiN planes are indexed……………………………………………………………………………. 73

Figure 4.5. θ/2θ and glancing angle XRD of TiN coating deposited with an evaporator current set point of 65 A and a -300 V substrate bias. TiN planes are indexed……………………………………………………………………………. 73

Figure 4.6. θ/2θ XRD pattern of CrN coating C080328-1 deposited with an evaporator current set point of 45 A and a -200 V substrate bias. CrN planes are indexed……………………………………………………………………………. 75

Figure 4.7. θ/2θ XRD patterns of nanolayer (Ti,Cr)N coatings deposited with various Cr evaporator current set points (Ti evaporator @ 65 A) and -150 V substrate bias. B1 NaCl TiN and CrN planes are indexed……………………….. 76

Figure 4.8. θ/2θ pattern of the (422) plane of nanolayer (Ti,Cr)N coatings deposited as a function of Cr evaporator current (constant Ti evaporator current and substrate bias)………………………………………………………………… 78

Figure 4.9. θ/2θ XRD patterns of nanolayer (Ti,Cr)N coatings deposited with various substrate biases (constant evaporator current). B1 NaCl TiN and CrN planes are indexed………………………………………………………………… 79

Figure 4.10. θ/2θ pattern of the (422) plane of nanolayer (Ti,Cr)N coatings deposited at different substrate biases (Ti evaporator of 65 A, Cr evaporator of 45 A)……………………………………………………………………………… 80

Figure 4.11. θ/2θ XRD patterns of multilayer (Ti,Cr)N coatings deposited with Ti interlayers (constant evaporator current and substrate bias). B1 NaCl TiN and CrN planes are indexed…………………………………………………………… 81

Figure 4.12. θ/2θ XRD patterns of multilayer (Ti,Cr)N coatings deposited with Nb interlayers (constant evaporator current and substrate bias). B1 NaCl TiN and CrN planes are indexed as well as Nb (110) and (220) planes……………….. 82

Figure 4.13. Crystallite size of TiN increases with increasing evaporator current (-150 V substrate bias). Error estimated to be at least ±10% (error bars)...... 84

Figure 4.14. Crystallite size of TiN increases with increasing substrate bias. Error estimated to be at least ±10% (error bars)………………………………….. 85

Figure 4.15. Crystallite size of nanolayer (Ti,Cr)N as a function of substrate bias increases then decreases. Error estimated to be at least ±10% (error bars)……… 86

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Figure 4.16. Increasing (111) crystallographic orientation of TiN with increasing Ti evaporator current (constant substrate bias)…………………………………… 89

Figure 4.17. [111] preferred crystallographic orientation of the TiN as a function of substrate bias (constant Ti evaporator current)………………………………… 93

Figure 4.18. In the interstitial argument for nitride structures, TiN is based on hcp α-Ti while CrN is based on a structural “switching” or change of the bcc Cr lattice to an fcc lattice.85…………………………………………………………... 96

Figure 4.19. Decreasing [111] crystallographic orientation with increasing Cr evaporator current of nanolayer (Ti,Cr)N coatings (65 A Ti evaporator, constant substrate bias)……………………………………………………………………... 97

Figure 4.20. Maximum (111) preferred crystallographic orientation for nanolayer (Ti,Cr)N coatings deposited at different substrate biases (constant Cr and Ti evaporator current) occurring for coatings deposited at -150 V…………... 99

Figure 4.21. Increasing (111) preferred crystallographic orientation with increasing number of layers of multilayer (Ti,CrN) coatings with Ti interlayers (- 150 V substrate bias, Ti evaporator current of 65 A, Cr evaporator current of 45 A)…………………………………………………………………………………. 101

Figure 4.22. Residual stress of monolithic TiN coatings as a function of Ti evaporator current (constant substrate bias). k is a constant representing the difference between the calculated TiN elastic modulus from the literature, E, and the actual elastic modulus in these coatings as well as differences between E111 and E422…………………………………………………………………………… 105

Figure 4.23. Relative residual stress of TiN coatings as a function of substrate bias (constant Ti evaporator current). k is a constant representing the difference between the calculated TiN elastic modulus from the literature, E, and the actual elastic modulus in these coatings as well as differences between E111 and E422….. 107

Figure 4.25. A linear relationship between deposition rate and Ti evaporator current is evident for the monolithic TiN coatings (-150 V substrate bias)………. 109

Figure 4.26. Deposition rate versus substrate bias for monolithic TiN coatings deposited with a constant 65 A evaporator current. Rate decreases above -150 V substrate bias due to resputtering of the growing coating………………………… 111

Figure 4.27. Increasing deposition rate as a function of Cr evaporator current for nanolayer (Ti,Cr)N coatings deposited with -150 V substrate bias and 65 A Ti evaporator current………………………………………………………………… 112

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Figure 4.28. Deposition rate nanolayer (Ti,Cr)N coatings deposited at different substrate biases (65 A Ti evaporator current, 45 A Cr evaporator current). Trend attributed to resputtering of the macroparticles and the growing coating………… 113

Figure 4.29. EPMA analysis of Ti and Cr atomic percent from nanolayer (Ti,Cr)N coatings deposited with different Cr evaporator currents and Ti evaporator current of 65 A (-150 V substrate bias). Coatings are approximately 48 to 51 atomic percent N………………………………………………………… 114

Figure 4.30. EPMA Cr and Ti atomic percent from nanolayer (Ti,Cr)N coatings deposited at various substrate biases (Ti evaporator of 65 A, Cr evaporator of 45 A). Coatings are approximately 45 to 51 atomic % N…………………………… 114

Figure 4.31. Surface defects increase significantly as a function of Cr evaporator current for nanolayer (Ti,Cr)N coatings deposited with -150 V substrate bias and 65 A Ti evaporator………………………………………………………………... 116

Figure 4.32. Surface defects increase with increasing Cr evaporator current for nanolayer (Ti,Cr)N coatings deposited with -150 V substrate bias and 65 A Ti evaporator current………………………………………………………………… 117

Figure 4.33. Relative surface area of large scale defects as a function of substrate biases for nanolayer (Ti,Cr)N coatings (65 A Ti evaporator current, 45 A Cr evaporator current)………………………………………………………….. 118

Figure 4.34. Fracture surface of monolithic TiN coating deposited with a 65 A Ti evaporator current and a -150 V substrate bias………………………………... 119

Figure 4.35. Fracture surface of nanolayer (Ti,Cr)N coating (Cr evaporator of 45 A, Ti evaporator of 65 A and a -150 V substrate bias)…………………………… 119

Figure 4.36. ADF-STEM micrographs of nanolayer (Ti,Cr)N coatings deposited at different Cr evaporator currents (Ti evaporator current of 65 A, -150 V substrate bias)……………………………………………………………………... 121

Figure 4.37. EPMA results for the atomic % of Cr in nanolayer (Ti,Cr)N coatings versus the periodicity, λ, as determined by ADF-STEM. (Ti evaporator current of 65 A, -150 V substrate bias)…………………………………………… 121

Figure 4.38. The CrN rich nanolayer thickness increases to over twice the TiN rich nanolayer thickness (6.9 nm) with increasing Cr evaporator current. Calculated using EPMA and STEM results for nanolayer (Ti,Cr)N coatings (Ti evaporator current of 65 A, -150 V substrate bias)……………………………….. 122

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Figure 4.39. Increasing critical force with increasing Cr evaporator current for nanolayer (Ti,Cr)N coatings (-150 V substrate bias, 65 A Ti evaporator current).. 124

Figure 4.40. Surface roughness decreases with increasing substrate bias for the monolithic TiN coatings (65 A Ti evaporator current)…………………………… 126

Figure 4.41. Surface roughness increases with increasing Cr evaporator current for nanolayer (Ti,Cr)N coatings (65 A Ti evaporator current, -150 V substrate bias)……………………………………………………………………………….. 127

Figure 4.42. Surface roughness of nanolayer (Ti,Cr)N coatings deposited at various substrate biases (65 A Ti evaporator current, 45 A Cr evaporator current) appears to be linked to large scale defects and crystallite size…………………… 128

Figure 4.43. Surface roughness of multilayer (Ti,Cr)N coatings with Ti and Nb interlayers (65 A Ti evaporator current, 45 A Cr Evaporator, and -150 V substrate bias)……………………………………………………………………... 129

Figure 4.44. Glass bead erosion results shown as mass loss versus total dosage of erodent on a single 1/4 in diameter spot for the TiN coatings deposited with different evaporator currents. Original test………………………………………. 131

Figure 4.45. Glass bead erosion results shown as mass loss versus total dosage of erodent on a single 1/4 in diameter spot for the TiN coatings deposited with different evaporator currents. Confirmation test spot……………………………. 132

Figure 4.46. Progression of observed micro-chipping method of erosion failure for the cathodic arc deposited coatings…………………………………………… 133

Figure 4.47. Alumina erosion media results shown as mass loss for a single 25 g test conducted on the monolithic TiN coatings deposited with different evaporator currents………………………………………………………………... 134

Figure 4.48. Glass bead erosion results shown as mass loss versus total dosage of erodent on a single 1/4 in diameter spot for the monolithic TiN coatings deposited with different substrate biases. Original test…………………………... 135

Figure 4.49. Glass bead erosion results shown as mass loss versus total dosage of erodent on a single 1/4 in diameter spot for the monolithic TiN coatings deposited with different substrate biases. Confirmation test spot………………... 135

Figure 4.50. Alumina erosion media results shown as mass loss for a single 25 g test conducted on the monolithic TiN coatings deposited with different substrate biases……………………………………………………………………………… 136

Figure 4.51. Glass bead erosion results shown as mass loss versus total dosage of erodent on a single 1/4 in diameter spot for the monolithic CrN coatings. Original test……………………………………………………………………….. 138

Figure 4.52. Glass bead erosion results shown as mass loss versus total dosage of erodent on a single 1/4 in diameter spot for the monolithic CrN coatings. Confirmation test spot…………………………………………………………….. 139

Figure 4.53. Alumina erosion media results shown as mass loss for a single 25 g test conducted on the monolithic CrN coating……………………………………. 139

Figure 4.54. Glass bead erosion results shown as mass loss versus total dosage of erodent on a single 1/4 in diameter spot for the nanolayer (Ti,Cr)N coatings deposited at different evaporator currents. Original test…………………………. 142

Figure 4.55. Glass bead erosion results shown as mass loss versus total dosage of erodent on a single 1/4 in diameter spot for the nanolayer (Ti,Cr)N coatings deposited at different evaporator currents. Confirmation test spot………………. 142

Figure 4.56. Alumina erosion media results shown as mass loss for a single 25 g test conducted on the nanolayer (Ti,Cr)N coating at different evaporator currents. 143

Figure 4.58. Glass bead erosion results shown as mass loss versus total dosage of erodent on a single 1/4 in diameter spot for the nanolayer (Ti,Cr)N coatings deposited at different substrate biases. Confirmation test spot………………….. 146

Figure 4.59. Alumina erosion media results shown as mass loss for a single 25 g test conducted on the nanolayer (Ti,Cr)N coating at different substrate biases…... 147 Figure 4.61. Glass bead erosion results shown as mass loss versus total dosage of erodent on a single 1/4 in diameter spot for the multilayer (Ti,Cr)N coatings deposited with Ti interlayers. Confirmation test spot……………………………. 149

Figure 4.63. Glass bead erosion results shown as mass loss versus total dosage of erodent on a single 1/4 in diameter spot for the multilayer (Ti,Cr)N coatings deposited with Nb interlayers. Confirmation test spot…………………………… 151

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Figure 4.64. Alumina erosion media results shown as mass loss for a single 25 g test conducted on the multilayer (Ti,Cr)N coating with Ti and Nb interlayers…... 152

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LIST OF TABLES

Table 2.1. Composition of AM355 stainless steel.32…………..………..……...... 17

Table 2.2. Summary of typical coating properties of arc evaporated TiN, CrN, 28 and (Ti, Cr)N available from the literature………………………………………...

Table 3.1. Deposition processing parameters for TiN coatings………………….. 48

Table 3.2. Deposition processing parameters for CrN coating…………………... 49

Table 3.4. Deposition processing parameters for nanolayer (Ti,Cr)N coatings deposited with varying substrate biases at constant Ti and Cr evaporator currents. 51

Table 3.5. Deposition processing parameters for multilayer (Ti,Cr)N coatings with Ti Interlayers………………………………………………………………… 52

Table 3.6. Deposition processing parameters for multilayer (Ti,Cr)N coatings with Nb Interlayers………………………………………………………………... 53

Table 3.7. XRD scan parameters for nitride coatings……………………………. 55

Table 3.8. Erosion test parameters for glass bead erosion tests…………………... 66

Table 4.1. Surface energy of B1 NaCl structured TiN coatings calculated by 78 Pelleg et al. ……………………………………………………………………… 90

Table 4.2. Density of Atoms per Plane for B1 NaCl structure TiN (a=4.242 Å)... 91

Table 4.3. Relative integrated intensities of the TiN coatings as compared to JCPDF 01-38-1420.50……………………………………………………………... 93

Table 4.4. Relative integrated intensities of the CrN coating as compared to JCPDF 01-11-0065.51……………………………………………………………... 94

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ACKNOWLEDGEMENT

Foremost, I would like to acknowledge and thank Dr. Douglas E. Wolfe. His

support and guidance made this project possible. Additionally, the staff members of the

Advanced Coatings Department at the Pennsylvania States University’s Applied

Research Laboratory (ARL) deserve credit for their countless contributions to every

aspect of this project; thank you Scottie Showers, Dennis McGregor, Anna Stump, Cindy

Hull, and Tom Medill.

A special thanks goes to Brian Borawski whose previous research paved the way for this project. I would also like to thank Mark Angelone for the electron probe micro analysis, Joe Kulik and his staff for the transmission electron microscopy, and the staff at

MDS-Prad for the scratch adhesion testing.

I would like to express my gratitude to the staff of the Exploratory and

Foundational (E&F) Graduate Student Research Program of the ARL for giving me the opportunity to perform research in the ARL’s top notch facilities.

Lastly, I would like to thank my friends and family for the support that they have provided throughout my graduate studies. This project would not have come to fruition without their encouragement.

CHAPTER 1

INTRODUCTION

1.1 Background

Hard particle erosion from sand particles can cause significant damage to metallic turbine engine components such as compressor blades and blisks (single piece component composed of rotor, disc, and blades).1 The damage can be detrimental to engine performance and also results in significantly higher maintenance costs along with decreased mean time between overhaul of the engines. Surface coatings such as titanium nitride (TiN) based coating systems offer a possible solution as they can impede the damage mechanisms which degrade the uncoated metallic components.2

1.2 Project Objectives

There are two primary objectives for this project:

1 Explore the synthesis-structure-property relationships of TiN, CrN

(chromium nitride), and titanium-chromium-nitride based ((Ti,Cr)N)

cathodic arc coatings.

2 Examine the relationships between erosion and the structure-properties

of the coatings in order to further the scientific understanding of the

correlations between coating processing, coating microstructure, and

erosion performance.

The combination of the two primary objectives should provide significant insight on how to tailor the coating deposition process as well as the coating design and

2 architecture in order to enhance the erosion protection of TiN, CrN, and ((Ti,Cr)N), coating systems.

The secondary objectives are to explore the affects of cathode evaporator current as well as substrate bias on the microstructure and performance of the coatings since both of these variables play a key role in the properties of cathodic arc evaporated coatings.3

Another important secondary objective for this project is to examine the structure and

composition of (Ti,Cr)N (ternary phase versus nanolayered structure) coatings deposited

at low temperatures with single phase Ti and Cr evaporators in a N2 rich environment.

The study on the structure and composition is essential as there has been some variation in the literature regarding the (Ti,Cr)N coating system and the corresponding properties.4-

CHAPTER 2

LITERATURE REVIEW AND COATING THEORY

This chapter is divided into three main sections. The first section discusses the

cathodic arc deposition process, important deposition process variables and explains why

this system is ideal for production of nitride based erosion resistant coating systems. The

second section is focused on the phenomenon of hard particle erosion and discusses both

the mechanics of erosion and the desired properties of an erosion resistant coating

system. The third section in this chapter discusses properties of TiN, CrN, and (Ti,Cr)N

coating systems.

2.1 Cathodic Arc Deposition

Perhaps the main factor contributing to the performance of any coating system

such as titanium-based nitrides is that the deposition process can significantly influence

the microstructure of the coatings. Transition metal nitrides used for wear and

mechanical applications have traditionally been deposited using physical vapor

deposition (PVD) techniques such as electron beam evaporation, cathodic arc

evaporation, and sputtering. In addition to the common PVD techniques, ion beam techniques such as ion implantation and ion beam assisted deposition have also been used to deposit hard nitride coatings.8,9 Another deposition technique, chemical vapor deposition (CVD), can also be used to grow transition metal nitride coatings, but this

technique requires higher deposition temperatures making it difficult to coat temperature

sensitive substrates.

For this study, the deposition system of interest is cathodic arc evaporation.

Cathodic arc is a highly versatile PVD technique which is commonly used in a wide variety of applications including wear resistance, decorative coatings, and erosion resistance. The cathodic arc deposition process can produce nitride based coatings with excellent adhesion, high deposition rates, and high hardness due to the flexibility in ion energy bombardment resulting in increased substrate/coating intermixing and the ability to tailor the compressive residual stress.3

Cathodic arc deposition is a plasma-based PVD evaporation technique which uses a high current arc to melt a metallic target material which is transported through a vacuum and deposited onto a substrate. A typical configuration for a cathodic arc deposition process is illustrated in Figure 2.1. The source material for the coating is the cathode which has a positive voltage potential as compared to the rest of the system. In

Figure 2.1, the substrate serves as the anode (negative bias) and attracts the ionized species towards the growing coating. There can be more than one cathode in a system

which allows co-evaporation of multiple materials or higher deposition rates of the same

material.10,11

Figure 2.1. Depiction of cathodic arc plasma deposition.

The arc ignition results in a high density but localized plasma. The localized

plasma is due to the large ionized fraction (up to 100%) of the evaporated species;

however, the degree of ionization is highly dependent on the material evaporated,

evaporation rate, and the background pressure. Typically, the degree of ionization

increases with increasing evaporant mass since heavier atoms typically have weakly held

valence electrons due to shielding from inner electron orbitals closer to the nucleus.11

This high level of ionization is in contrast to evaporation (thermal, electron beam, etc.) and sputtering (dc diode, magnetron, etc.) where only a few percent of the evaporated or sputtered species become ionized.11

The high degree of ionization in the cathodic arc process is one of the primary reasons this process is ideal for erosion resistant coatings. The cations generated at the evaporator source can be accelerated towards substrate and growing coating where they can potentially influence coating properties such as adhesion, residual stress, and hardness. One of the objectives for this thesis was to examine how the high degree and magnitude of ion bombardment in the cathodic arc processes influences the coating properties and thus their hard particle erosion resistance.

It is important to note that the evaporator current is distributed over just a small area of the target rather the entire surface. These high current density regions are known as cathode spots and the average size of a cathode spot is 1 to 10 μm, and the current density is in the range of 106 to 108 A/cm2. The location of the cathode spots are

continuously changing, and a few research studies have shown that the lifetime that a

spot stays at one location ranges from tens of nanoseconds until approximately one

microsecond.3,12 The velocity of the spot movement is very dependent on the surface

conditions of the target and can range from 10-1000 m/s for contaminated surfaces such

as oxidized metals to 1-150m/s for clean surfaces.3 The multitude of small intense spots

creates a dense plasma above the target material.3,10

The high current density in the cathodic arc process eliminates the problem of

target poisoning which affects reactive depositions such as sputtering. In reactive gas

evaporation and sputtering of Ti-based and Cr-based nitrides, a partial pressure of N2 is reacted with evaporated or sputtered parent metal atoms (Ti or Cr). Poisoning generally occurs when the reactive gas creates a nitride layer on the surface of the source or target thereby reducing the electrical conductivity of the cathode and greatly reducing the

7 deposition rate in sputtering deposition systems. Figure 2.2 shows the system pressure versus nitrogen flow rate (a) cathodic arc and (b) reactive sputtering. As illustrated, no target poison is observed for the reactive cathodic arc process whereas for reactive sputtering poisoning occurs rapidly at the point where the nitride formation rate at the target exceeds the sputtering rate. When this occurs, the system pressure will rapidly increase since there is significantly less N2 being consumed during the coating process.

A hysteresis curve results since the nitride layer is not immediately sputtered from the

3,11 cathode target as the N2 flow is decreased.

For reactive sputtering, the optimal flow rate for depositing nitrides is therefore

found just before the rapid increase in pressure since at this point the N2 is being

consumed efficiently at the substrate rather than at the target. Since the difference

between efficiently sputtering and poisoning the target is only a small increase in the gas flow (deposition pressure), it is very difficult to efficiently deposit stoichiometric nitride coatings with a high deposition rate using reactive sputtering without a residual gas analyzer or other in-situ technique. For comparison, since cathodic arc operates with a very high current applied through very small arc spots, the deposition rate from these localized spots is generally too high to allow target poisoning, and therefore reactive gas

cathodic arc deposition is generally considered an easier process to control.3,11 This is a

significant benefit for large scale production such as coating turbine engine blades and

blisks.

Figure 2.2. System pressure versus nitrogen gas flow rate for (a) reactive cathodic arc evaporation and (b) reactive sputtering techniques.

One of the main drawbacks of the cathodic arc deposition method is the ejection

of large particles from the surface of the source during coating. The ejection of 0.1 to 10

μm diameter particles is known as macroparticle generation and is detrimental to coatings

and thin films requiring a uniform microstructure such as in magnetic thin films for data

storage. There are several theories which discuss macroparticle ejection. One theory is

that the high ionization at the cathode spot creates enough heat to melt the source

material just below the surface. The pressure of the expanding liquid metal then forcibly

ejects material from the surface of the source. A second theory of macroparticle creation

is that the dense ion current at the cathode arc spots creates localized melt pools. The

force of the impacting ions then splatters molten metal out of the melt pool and into the

plasma. This latter theory may help to explain the experimental data that often shows

that most macroparticles are ejected at low angles to the target, rather than perpendicular

to the surface.3,10

There has been extensive research in the development of methods to filter macroparticles from cathode arc deposition. Early methods of filtering the plasma involved steering the ions through curved magnetic ducts. The original concept relied on

9 the fact that the plasma consists of mostly ionized particles. As a result, when the vapor is exposed to a magnetic field, the ions respond to the magnetic field, but the non-ionized macroparticles collide with the walls of the duct. This method was first introduced in the

1970’s by Aksenov et al. in the Soviet Union.13-15 Figure 2.3 illustrates a simple macroparticle filter system using a magnetic solenoid.10 The main disadvantage with

filtering cathodic arc systems is that there is a loss of plasma in the solenoid. Losses can

be as high as 90% or greater of the total plasma in order to deposit completely

macroparticle free coatings, and this results in significantly lower deposition rates an

increased costs.

Figure 2.3. Cathodic arc plasma filtered through magnetic solenoid in order to remove macroparticles. Image from Brown.10

2.2 Mechanisms of Erosion

Erosion is a complex phenomena and there has been at least four separate or cooperative mechanisms proposed regarding the removal of material due to small solid particle impacts and are listed below:16-19

1. Ductile Material Erosion

2. Brittle Material Erosion

3. Thermal Degradation Erosion

4. Hard Particle Composite Erosion

One of the most commonly discussed erosion mechanisms is applicable to ductile materials in which plastic deformation plays an important role in the erosion mechanism and thus erosion resistance. For brittle materials, the most commonly discussed erosion mechanism is micro-chipping which involves removal of small pieces of the impacted sample around the particle impact sites due to formation of stress fields which result in radial cracking.16,17,19-22 Of the four erosion mechanisms, ductile and brittle erosion

mechanisms are considered the most relevant to the current problem associated with

erosion of first stage turbine engine blisks and will be discussed in more detail in

Sections 2.2.1 and 2.2.2, respectively.

Another mechanism discussed in the literature proposes that the impacting

particles cause erosion through melting or heating the surface of the material which in

turn reduces the mechanical bonding strength and thereby increasing erosion rate.16,23

This mechanism is not believed to be a factor in the erosion performance of first stage

compressor blades since blades are not subjected to significantly high temperatures and

are cooled by the intake air. A fourth mechanism discussed in the literature is relevant to

multiple phase composite materials which have hard particles embedded in a more ductile

matrix (i.e., a cermet). For these material systems, it is proposed that porosity around the

particle-matrix interface can act as nucleation sites for cracks. Subsequently, crack

coalescence occurs leading to removal of the hard particles from ductile matrix.16,24

2.2.1 Solid Particle Erosion through Plastic Deformation

The predominant theory for the solid particle erosion of ductile materials is that material loss occurs through a plastic deformation process. If the particle is harder than the surface it is striking, material will be pushed or ploughed in the direction of the impinging particle as illustrated in Figure 2.4 showing a hard particle impinging a metallic surface causing plastic deformation (platelet formation). Multiple impacts cause the deformed material to “pile up” at the edge of the impact craters where subsequent impacts cause the material to break free from the surface resulting in erosion or material loss.2,19,25

Figure 2.4. Sand particle causing platelet formation on a metallic coating.

Erosion rates are typically calculated as the mass of the material loss compared to

the total mass of impacting sand particles. For ductile metals, such as titanium and

aluminum alloys, the highest/maximum erosion rate has been determined to occur when

particles strike the surface at angles between 20 and 30 degrees.2,19 In an experiment by

D’Alessio and Nagy using titanium alloy,2,25 maximum erosion rate versus the impact

angle, α, correlated well with an equation developed by Finnie:25

mV2 6 K Q  02(sin( 2 )tan()));(sin  Kp K 6

mV2 K Q  2 )tan());((cos  Kp 6 where,

Q: volume of material removed by a single grain (ft3) m: particle mass (g) p: plastic flow stress (p.s.i.)

ψ: ratio of contact length to impact depth

K: ratio of tangential to normal force of impact

V: impact velocity (ft/sec)

α: angle of impingement (degrees)

2.2.2 Solid Particle Erosion of Brittle Materials

In materials with high Young’s modulus values, such as most ceramics, plastic deformation is no longer a valid erosion model. In these brittle fracture erosion materials, an impacting particle can cause a very small crater surrounded by a Hertzian strain field.

If the force of the impact is large enough, lateral cracking can occur as depicted in Figure

2.5.21 After enough cracks have been formed and with subsequent crack propagation, crack coalescence results in brittle material erosion as material breaks and chips off near the cracks (micro-chipping).

Figure 2.5. Lateral crack formation in a brittle material. Adapted from D’Errico.21

The micro-chipping failure mechanism is well documented in the literature.19,21,25

One of the earliest studies was by Levy,19 who studied the erosion of nickel oxide on pure

nickel. In his study, there were two layers of nickel oxide present on the samples. The

innermost 75 μm nickel oxide layer consisted of small equiaxed nickel oxide grains with

a high percent of porosity. The outer 25 μm nickel oxide was composed of a more dense

columnar grain structure. As the impacting particles contacted the dense nickel oxide

outer layer, small craters formed with radial cracks circulating outward. After the

particle rebounded from the surface, a zone of planar cracks formed due to the interaction

between the compressive stresses surrounding the crack and the compensating tensile

stresses beneath the impact. This resulted in vertical and horizontal cracks propagating

from the impact crater. The cracks then propagated through the porous interlayer in a

cone type pattern resulting in micro-chipping.19

For brittle materials, the primary contributor to erosion rate is related to how

much energy is transferred in the particle-surface collision. Maximum energy transfer

should occur at 90 degrees, and correspondingly, experiments from literature confirm that

erosion rates are highest at 90 degrees for brittle materials.1,2,19,26 In the mid nineteen- nineties there were several experiments that compared the erosion of ceramic coatings to

14 uncoated metal substrates. D’Alessio and Nagy compared the erosion of titanium nitride coated titanium substrate to an uncoated titanium alloy.2 Their results showed that the

titanium nitride coated sample performed better than the non-coated titanium alloy over

the entire tested range of impact angles of 10 to 90 degrees. The titanium nitride coating

exhibited its highest erosion rates at angles close to ninety degrees, while the titanium

alloy had it lowest erosion values at 90 degrees.2

In 1999, Tabakoff examined titanium carbide coatings on various high

temperature nickel based alloys compared to the uncoated alloys and the results are

presented in Figure 2.6.1 The TiC coating was deposited by CVD and consisted of a fine

grained microstructure with a total thickness of approximately 15 μm. Both coated and

uncoated samples were subjected to 180 to 366 m/s chromite or fly ash particles at

temperatures up to 815 °C. The results of the experiment showed the highest erosion

rates for the base alloy occurred between 15 and 50 degrees depending on the alloy

composition which is related to plastic deformation as previously discussed. The coated

samples showed the highest erosion rates when the particles impacted at 90 degrees. One

set of samples, with INCONEL 718 substrates, showed that the nitride coated specimen

eroded approximately three times faster at 90 degrees than uncoated INCONEL 718. The

erosion rate of the coated versus uncoated sample was approximately equal at thirty degrees, and the maximum erosion of the uncoated alloy occurred at approximately 20

degrees. The results of the study are in agreement with the general trends of brittle

materials eroding faster at higher angles and ductile materials eroding at lower angles

(Figure 2.7).

Figure 2.6. Erosion rate variation of uncoated and coated INCONEL 718 with impingement angle: T=538°C, V=305 m/s, 20g chromite mass.1

It should be noted that discrepancies exist in the literature for material coating systems evaluated under different conditions. Comparison of the literature is quite

16 difficult as the erosion test conditions, particle velocities, and material properties are often not provided.

2.2.3 Erosion Resistant Coatings

The angle of impact for sand particles in a turbine engine may occur over the entire range of 0 to 90 degrees as illustrated in Figure 2.8.28 The wide range of impact

angles makes turbine engine blades and blisks particularly vulnerable to erosion since

most engineering materials fail with either the previously mentioned micro-chipping

mechanism for high angle erosion of brittle materials or the low angle plastic deformation mechanism which is typical of ductile materials.

Figure 2.8. Sand particles enter to the left and strike the compressor/turbine blade at various angles. Adapted from Tabakoff.28

For metallic turbine engine blisks, the plastic deformation mechanisms leads to

significant roughening of the leading and trailing edges of the blades which results in

decreased engine performance and efficiency. Leading edge damage of a first stage

compressor blisk from a General Electric (GE) T-700 series engine is shown in Figure

2.9. The damage is a result of a combination of erosion (plastic deformation) and large

particle impacts which result in the observed deformation of the leading edge. This

17 particular blisk is machined from a GE proprietary stainless steel similar to AM355

(UNS35500). The composition of AM355 is given in Table 2.1. One of the objectives for this project is to minimize hard particle erosion of components such as this blisk, so

AM355 was chosen as the primary substrate material for all of the coating trials.

Figure 2.9. Leading edge damage to a stainless steel (AM355) first stage compressor blisk due to solid particle erosion. GE T-700 series blisk serial #GWHTFANG.

Table 2.1. Composition of AM355 stainless steel.29 Element Atomic % Fe 73.70 to 80.14 Cr 15.00 to 16.00 Ni 4.00 to 5.00 Mo 2.50 to 3.25 Mn 0.50 to 1.25 Si 0.50 C 0.10 to 0.15 N 0.07 to 0.13 P 0.04 S 0.03

Since the primary degradation mechanism for the blisk is related to plastic deformation, a logical starting point is to design a coating which does not exhibit this type of erosion behavior. The previously mentioned work by D'Alessio and Tabakoff on nitride and carbide coatings, respectively, shows that low angle erosion damage can be decreased significantly through the application of ceramic nitride and carbide coatings.1,26

However these monolithic ceramic based systems do not provide significant high angle

erosion resistance. The protective coating system must also withstand hard particle impacts at a 90 degree impingement angle without large scale micro-chipping or delamination.

In past research studies of hard materials and coatings such as nitrides there has been several microstructural properties identified as important to mitigating hard particle erosion, and include small grain size, low degree of porosity, absence of cracks,30 high hardness, good toughness,21,31 and excellent coating adhesion.32 In addition, the preferred

crystallographic orientation of a coating can significantly affect erosion performance.33

Furthermore, for coating systems, it has been shown that the substrate properties such as composition and toughness can affect the erosion performance since the substrate can affect the stress state, orientation, microstructure and adhesion of the coating.34

It should be noted that some contribution of these variables appears to be directly

related to specific coating/substrates systems as well the actual erosion test set-up and

conditions. For example, Davies et al., found that smaller grain sizes (i.e. greater density

of grain boundaries) resulted in higher erosion rates of chemical vapor deposited diamond

(CVD) coatings which is contrary to the findings of Tabakoff et al. which showed that

smaller grain sizes increased erosion resistance for detonation and thermal spray

19 coatings.30,35 Again it should be noted that little confirmation with regards to

compressive residual stress, fracture toughness, etc, were provided making direct

comparisons difficult.

The grain structure of a PVD coating is largely affected by the deposition process

and in particular, how this process influences adatom diffusion along the surface of the substrate and the growing coating. Typically, the species will diffuse until either they reach an energetically favorable site, are buried and/or locked into position by the arriving flux, or resputtered from the arriving species. Diffusion is directly related to thermal energy, so the temperature of the substrate plays an important role during this part of coating growth.

There is usually several dependent and independent coating processing variables that contribute to the substrate temperature during the deposition process. For techniques

such as cathodic arc, the substrate temperature is controlled by the evaporator current and

the substrate bias since these deposition process variables control the number of ions

impacting the substrate. Additionally, radiant heat is often applied during cathodic arc

deposition to maintain a uniform substrate temperature and coating microstructure.10,11

The relative thermal energy of the species diffusing at the surface of the coating is related to the homologous temperature which is the ratio of deposition temperature to coating system melting temperature (T/TM). The homologous temperature has been used to predict coating microstructure through structural zone models. Structural zone models

are essentially maps which relate surface diffusion and deposition conditions to typical

coating microstructures (Zone 1, T, 2, 3). A structural zone model developed by Thorton

20 for sputtering is illustrated in Figure 2.10 as a function of process pressure and deposition temperature.36

Figure 2.10. Thorton structural zone model.37

There are typically four structural zones identified for higher energy PVD

coatings such as cathodic arc and sputtering. The low temperature zone, Zone 1, occurs

when surface diffusion is minimal. This zone is characterized by fine grains as small as

tens of nanometers separated by pores, combined with a high concentration of large scale structural defects. Additionally, there is a grain growth competition in which large grains can limit the growth of smaller grains by shadowing the incoming coating flux.38 This type of microstructure is typical for low energy processes and rotational configurations

21 such as thermal evaporation, and the high amount of porosity and low density structure would not be desirable for erosion resistant coatings.

The second zone is the transitional zone, Zone T, which is the zone that is unique to higher energy deposition systems such as sputtering and cathodic arc and not typically found in systems such as thermal evaporation using resistive sources. The microstructral characteristics of Zone T include densely packed fibrous grains. The columns still contain a fairly high concentration of large scale structural defects but this is less than

Zone 1 due to increased adatom mobility from the additional ion energy supplied by the deposition process.36

Zone 2 is described as having columns with uniform diameters and faceted faces.

The defect concentration is generally lower for these coatings due to the larger amount of surface diffusion that occurs due to increased energy. The final zone, Zone 3, occurs at elevated temperatures high enough to allow sufficient surface mobility and annealing of the coating material. This results in equiaxed or isotropic grains that can approach the thickness of the coating.36 Zone 2 and Zone 3 microstructures are potential candidate

structures for erosion resistance systems, but this would require high deposition

temperatures which could degrade the stainless steel substrates properties since TiN has a

melting point above 2000 °C.3

These structure zone models are meant to be guidelines as often the coating

system microstructures will deviate from the model due to contributions related to other

processing variables. For example, a very high deposition rate may produce a material

with a Zone 1 structure when the homologous temperature may predict a Zone 2 structure. A higher concentration of impurities can also impede or aid in surface

22 diffusion of the growing coating and shift the coating microstructure to a lower or higher zone, respective to the impurity’s contribution to the surface diffusion rate.

Levy et al. correlated the erosion performance of detonation gun applied Cr3C2 directly to the grain size of the coatings.30 The coatings had a very fine grain structure,

and damage caused by impacting particles was on the order of the size of the individual

grains. Additionally, Levy et al. showed that the Cr3C2 coatings performed better than

WC-Co plasma spray coatings due to the WC-Co’s fine network of porosity and voids

which served as failure points.30 This is most likely related to the hard particle composite

failure mechanism previously discussed. Examining these coating systems, it would

seem that fine grained coatings with low defect density would be ideal for erosion

performance. This would correspond to Zone T and Zone 2 grain structures.

The hardness of a material is a measurement of the resistance to plastic

deformation. This relates to the bond strength of the material and also the crystalline

structure (i.e., type of bonding). The coating hardness and particle impact relationship is

related to the plastic deformation that occurs when as a single particle impacts the

surface. Yang et al. showed that for TiAlN coatings, erosion resistance generally

increased with increasing hardness.31

Another important material property for erosion is toughness which is the ability

for a material or coating to resist cracking and crack growth. In terms of erosion, it has

been well established that brittle materials fail though micro-chipping and crack coalesce mechanisms. One way to increase the toughness of hard coatings is to use a multilayer design which has complaint interlayer materials in between the hard ceramic layers.

These compliant layers can theoretically act as barriers for crack propagation.

Additionally, the layers can interrupt column or microstructure growth which can result in smaller grains, and thus possibly better erosion performance. The multilayer design typically results tailored residual stress, since the compliant ductile interlayers can absorb increased energy and relieve coating residual stress through plastic deformation. A multilayer coating design is shown in Figure 2.11.

Figure 2.11. Proposed multilayer design for increased toughness of erosion resistant coatings.

In theory, a multilayer nitride coating system with compliant interlayers should

provide improved erosion resistance under specific erosive conditions as compared to a

monolithic coating. For high energy impact erosion such as large particles (10’s of

microns), a multilayer should out perform a monolithic since the interlayers can absorb

some of the energy of the impacting large particle.

The amount of energy absorbed should be directly related to ductility which is a

measure of a materials plastic deformation before failure. The energy absorbed should

also be directly related to the bulk modulus, B, which is a measure of the resistance to

volume change,  V, with a an applied compressive pressure,  p:

p  VB V

α-Ti is an ideal candidate for multilayer TiN based coatings since it has high ductility

(54% elongation at failure of annealed bulk sample).39 Niobium is less ductile (29%

39 elongation at failure of bulk sample with 0.03 atomic percent O2) but has a larger bulk modulus than Ti (24.7 ksi for Nb versus 15.7 ksi for Ti)40 and thus increased compressibility which should mean better energy absorption with regards to large particle impacts.

For low energy impacts (i.e. small particles) a thick monolithic coating should perform well since the small particles will cause failure due to micro-chipping by generating small Hertzian stress fields which do not penetrate very deep beneath the coating surface. Therefore, the ceramic layer should have high enough resistance to low energy low angle impacts, and the metallic interlayer would not absorb much of the energy of the impacting particle and could possibly detriment performance by increasing the volume of metal (softer phase).

As with any coating’s performance, excellent adhesion is critical for increased erosion performance which suggests coating/substrate interfacial size, structure and composition may significantly contribute to erosion performance.32 If there is a lack of diffusion between the coating and the substrate, then the interface may consist of only an abrupt transition of a few angstroms between the two materials. However, it there is solubility between the coating and substrate materials, then diffusion may occur resulting in an intermixed region. The size of the intermixed region can generally be increased by increasing the energy and flux of the energetic ions in deposition systems such as

25 cathodic arc. This intermixed type of interface can often improve adhesion and serve as a desirable interface for erosion resistant coating systems.38

The reactivity between the coating material and the substrate is also an important factor in determining coating adhesion. Often these materials will react and form compounds with a high degree of covalent/ionic bonding. These reactions typically result in a volumetric change due to the realignment of the atoms, and therefore stress fields can develop upon cooling. These stress fields can result in a more brittle interface, so they would most likely not be desirable for an erosion resistant coating.38 Another way to increase the adhesion of coatings is to roughen the surface prior to deposition.

This serves several purposes including increasing the surface area of the coating-substrate interface, creation of more tortuous fracture paths for the coating, and increasing the degree of mechanical interlocking between the substrate and the coating. The mechanical interface would be very desirable for the erosion resistant coatings, and the surface finish should not degrade the aerodynamic airflow over the component.

In addition, a bond or intermediate layer can be applied prior to the deposition of the erosion resistant coating system in order to improve performance. Typically, bond layers are used in processes where the chemical adhesion between the coating and the substrate is not very strong, so a thin bond layer which has good adhesion to both the substrate and the coating material is used to facilitate adhesion. Generally, coating materials with similar bonding structures as the substrate (metallic, covalent, or ionic) react well with the substrate and are more likely to form an adherent coating. For example, in the deposition of gold on silicon wafers, a titanium interlayer is often first

deposited. Titanium bonds well with the SiO2 on the surface through the formation of

11,38 TiO2. Gold then bonds well through the formation of metallic bonds on the titanium.

Another method for increasing adhesion in vapor deposition systems is to pre- clean the surface of the samples with plasma etching or ion bombardment. These methods use energetic ions to sputter contamination or passive layers such as oxides from the substrate in order to provide un-bonded material on the substrate surface. This newly exposed material is more likely to bond to the coating material in order to reduce surface energy.11,38 This method is typically used for increasing adhesion of coating materials deposited with PVD techniques.

The overall coating residual stress is comprised of intrinsic, extrinsic, and epitaxial contributions and can play a significant role in the coating performance.

Intrinsic stress in coatings such as cathodic arc coatings is generally caused by implantation of ions into the already frozen in coating structure. Additionally, the ions can transfer momentum to atoms in the frozen in structure resulting in movement of the atoms closer to their nearest neighbors. Both of these effects generally result in compressive stresses.11 For hard coatings such as TiN and CrN deposited by PVD techniques, extrinsic stress generally results from thermal expansion mismatch between the coating and substrate.11,38

Generally, compressive residual stress is desired for most wear applications since it can resist crack growth.41 If the residual stress level exceeds the interfacial shear forces then the coating can delaminate. Additionally, impacting particles will result in transfer of energy to the coating which in theory should increase the compressive residual stress.

Sue et al. showed that as compressive residual stress of arc evaporated TiN increased, the

27 erosion performance decreased; however, the adhesion and mechanical properties (i.e., toughness, hardness, etc) were not discussed in detail.33

The crystallographic orientation of coatings such as TiN and CrN are generally related to surface energy and strain energy competition.38 The preferred crystallographic orientation, defect density, and crystal structure of the coating will determine the atomic packing density and therefore the bond density which will influence the coating properties parallel and perpendicular to the substrate (anisotropy). Sue et al. established a correlation between preferred crystallographic orientation and erosion performance and found that erosion resistance of cathodic arc deposited TiN was highest for coatings with a strong (111) crystallographic orientation. It should be noted that other deposition parameters can also affect coating performance and are not always discussed in detail in the literature.

2.3 TiN, CrN, and (Ti,Cr)N Coatings for Erosion Protection

TiN and CrN were selected for erosion resistant coatings due to their high hardness, good corrosion and oxidation resistance, and the relative ease in which they can be deposited as both monolithic and multilayer coatings using reactive cathodic arc evaporation. The third coating system deposited in this study was (Ti,Cr)N which is believed to be deposited as a monolithic TiCrN coating or a nanolayer mixture of primarily CrN and TiN with varying degrees of solid solubility. Most of the (Ti,Cr)N coatings deposited in this study were found to be nanolayered structures with solid solubility due to the method of deposition. The relevant properties of the coatings systems as related to erosion resistance are shown in Table 2.2.3,5,42-49 It is well known

28 that coating properties can differ from bulk properties due to lattice defect generation and the complex relationship of coating properties as a function of coating microstructure which in turn is dependent on the deposition technique.

It should be noted that all of the material properties listed in Table 2.2 are specific to the coating composition and deposition parameters of their respective studies.3,5,42-49

For example, residual stress is highly dependent on the substrate material (thermal expansion mismatch, epitactic stress, composition, etc.) and the deposition parameters such as substrate bias, temperature, pressure, and target current since these parameters affect the energy of the incoming flux during coating growth as well as contribute to thermal stress relation if increased substrate temperature results.

Table 2.2. Summary of typical coating properties of arc evaporated TiN, CrN, and (Ti, Cr)N available from the literature. Coating Property TiN CrN (Ti,Cr)N Lattice Structure B1 NaCl B1 NaCl B1 NaCl Lattice Parameter (Å) a=4.24250 a=4.14051 a=4.18452 Columnar: dense packed or Columnar and Dense Grain Structure high defect equiaxed42,48 Columns45 density44,46 Vicker’s Hardness 2000 to 25003 1000 to 120053 23005 (kg/mm2) Adhesion, Critical Load 20 to 903 45 to 553 40 to 805 (N) Preferred Typically (111), Commonly Crystallographic possible (110) possible (111)5 Orientation (100)33,46,47 (100, 42,48,53 Compressive: -2 Compressive: - Compressive: Residual Stress (GPa) to -93 2.8 to -9.048,53 -2 to 749

Both stoichiometric TiN (δ phase) and CrN have the cubic B1 NaCl crystal structure (Fm-3m) (as illustrated in Figure 2.12) which is often described as two overlapping fcc lattices of N and Ti (or Cr), respectively. The theoretical lattice

29 parameter of stoichiometric CrN is 4.140 Å while the theoretical lattice parameter of

51,54 stoichiometric TiN is 4.242 Å. In addition to the TiN and CrN phases, the ε-Ti2N and

β-Cr2N phases are often intermixed in nitrogen deficient coatings along with Ti and Cr metal or metal rich phases.3,48

Figure 2.12. NaCl crystal structure modeled using CrystalDesigner Software. Ti and/or Cr atoms shown as green and N atoms as blue spheres.

The phase equilibrium diagrams for Ti-N and Cr-N are shown in Figure 2.13.55,56

The phase diagram for TiN having the NaCl crystal structure shows a broad compositional range.55 The broad compositional range is related to TiN’s defect structure, which contributes to the wide range of TiN properties deposited by various coating deposition techniques. For comparison, the only available phase diagrams in the

CrN system show stoichiometric CrN as a line compound and therefore should be more difficult to deposit as a single phase.56 It should be noted that substantial deviation from the phase equilibrium diagrams of CrN is believed to exist. The Cr based nitride systems have not been fully investigated and are poorly understood.

Figure 2.13. Phase equilibrium diagrams for (a) TiN and (b) CrN systems.55,56

The addition of the sub-stoichiometric tetragonal ε-Ti2N phase to TiN coatings has been shown to increase the average Vicker’s hardness value up to 3150 HV0.050 via

2nd phase strengthening and dislocation pinning mechanisms.57 Similarly, the hexagonal

β-Cr2N phase has a hardness value in the range of 1800 to 2950 HV. The lattice parameters for β-Cr2N are a = 4.759 Å and c = 4.438 Å. This phase is generally considered the hardest chromium nitride phase but generally has weaker adhesion for some substrates than CrN.58-60 However, a nanocomposite mixture of these phases may result in improved performance. A comparison of the crystal structures of β-Cr2N and

CrN is depicted in Figure 2.14.61

Figure 2.14. Hexagonal beta-chromium nitride (left) and B1 NaCl cubic chromium nitride (right). Modified from Era.61

There are two primary methods for depositing (Ti,Cr)N coatings by cathodic arc evaporation. The first method involves deposition using a single alloy target composed of Cr and Ti. This mythodology results in a coating flux which contains both Cr and Ti ions arriving in the same direction at the substrate surface. However, since the vapor pressures of Ti and Cr are different, (i.e. Ti is lower than that of Cr), Cr will preferentially evaporate from the target resulting in variation in coating composition until

32 near equilibrium values are reached.62 Based on the deposition processing parameters, this relationship can be determined and the coating target alloy composition (i.e., volume concentration) can be modified in order to obtain the desired coating composition. In addition the target composition alloy can be fabricated to obtain different coating compositions, similar to what is often done for TiAlN coatings deposited by sputter technique. However, a single intermetallic alloy of Ti-Cr is generally more expensive

(10X) and is also a very brittle cathode making it difficult to handle and contributing to difficultly in tailoring of the Ti:Cr ratio during a given coating deposition trial.

The other primary method for producing of (Ti,Cr)N coatings via cathodic arc deposition is through a multi-source technique in which there are multiple single phase targets or evaporators of Ti and Cr. The different directions of the coating fluxes combined with mechanical manipulation of the substrate generally creates a more layered coating structure.62 A layered coating structure may be desirable for erosion performance, since the layer interfaces may act as physical barriers to crack growth. In addition, decreased cost associated with large scale manufacturability and the ability to tailor the volume fraction of the TiN and CrN are benefits of this methodology.

An isothermal ternary phase equilibrium diagram for the Cr-N-Ti system is shown in Figure 2.15 and suggests that there is solid solubility between TiN and CrN at elevated temperatures since there is a wide composition range for this phase.52 The stoichiometric composition is referred to as Ti25Cr25N50 with a the lattice parameter of 4.184 Å which falls between the binary TiN and CrN lattice parameters.52

Figure 2.15. Ternary phase equilibrium diagram for Cr-N-Ti system at 1000 °C.52

Previous studies investigating the solid solubility and coating performance of

(Ti,Cr)N deposited using PVD techniques such as cathodic arc as well as reactive gas magnetron sputtering have shown mixed results. Vetter et al. used both single source and multi-source cathodic arc deposition at substrate temperatures up to 560 °C. Their results showed only one coating composition of (TiCr)N with 10 atomic percent Cr that was a single phase material. All of the other coatings were a mixture of (TiCr)N and hexagonal

β-(CrTi)2N (nitrogen deficient). The hardness of the coatings was reported to be higher

(up to 3860 Hv for 25-30 atomic percent Cr) than similarly deposited monolithic TiN

(2300 Hv) or CrN (2000 Hv) coatings. The authors contributed the high hardness to a solid solution strengthening mechanisms which provided an energy barrier to dislocation movement and strengthening due to the two phase nature of the coatings.5

Hiroyuki et al. used both binary and multi source evaporators in an arc ion plating

(AIP) deposition technique to deposit (Ti,Cr)N coatings. AIP varies slightly from cathodic arc in that the target species are evaporated then ionized rather than the evaporation and ionization occurring simultaneously as with cathodic evaporation. In comparison to cathodic arc, the primary disadvantage of AIP is the lower degree of ionized atoms, but this technique also results in fewer macroparticles. It should be noted that no direct comparisons on the structure-property-performance relationships of cathodic arc versus AIP have been investigated. Their results showed single phase

(Ti,Cr)N coatings with higher hardness than binary TiN and CrN coatings with no additional phases present in the coatings.7

Hiroyuki et al. results are in direct contrast to the study by Nainapampil et al. which used a multi-source arc evaporation technique to synthesize coatings that consisted of TiN, CrN, and β-Cr2N with hardness values lower than stoichiometric monolithic

TiN.4 Ward et al. used multi-target unbalanced magnetron sputtering deposition technique and deposited coatings comprised of a mixture of binary TiN, CrN, and β-Cr2N that had hardness values greater than CrN but less than TiN deposited under similar conditions.6 From the literature, there appears to be a wide variety of coating compositions, architectures, and microstructures reported for the (Ti,Cr)N system. These variations appear to be related to the deposition techniques, methodologies, and processing conditions (source/evaporator composition, energy of arriving species, substrate temperature, chamber geometry etc.) which most likely attributes to variation in coating performance.

CHAPTER 3

EXPERIMENTAL PROCEDURE

The first section of the experimental procedure chapter focuses on the deposition procedures for applying CrN, TiN, and nanolayer (Ti,Cr)N coatings using a multi-source cathodic arc evaporation system. (Ti,Cr)N refers to the samples synthesized using simultaneous evaporation of pure Cr and Ti targets in a nitrogen rich environment. The second section of this chapter discusses the various experimental procedures for coating characterization and analysis techniques. The primary characterization techniques used for this study include X-ray diffraction (XRD) for crystallographic structure, crystallite size, preferred crystallographic orientation, residual stress analysis, environmental scanning electron microscope (ESEM) for microstructure analysis, electron probe micro analysis (EPMA) for compositional analysis, and scanning transmission electron microscopy (STEM) for nanometer scale structure and compositional determination. In addition, coating erosion performance was determined using an in-house designed high particle velocity erosion rig.

3.1 Cathodic Arc Deposition Equipment

The various coatings for this study were deposited using a cathodic arc deposition system in the Advanced Coating Department at the Applied Research Laboratory of the

Pennsylvania State University. The system was manufactured by the Surface

Engineering Group, Incorporated (S.E.G.) in Mendota Heights, MN. The model number is 2020, and a schematic of the system is provided in Figure 3.1. The approximate

36 interior dimensions of the vacuum chamber are 20 x 20 x 20 inches. A diffusion pump is used in combination with a mechanical pump to achieve a typical base pressure of approximately 2 x 10-6 Torr prior to coating deposition. Substrate heating is aided with a radiant heater located in front of the high vacuum gate valve. For reactive nitrogen coating deposition, high purity nitrogen gas (99.999%) was reacted with the cathodic arc evaporated species.

Figure 3.1. Schematic of cathodic arc deposition system.

The system is equipped with three evaporation targets powered by individual

Miller XMT 304 CC/CV DC welder power supplies. Mechanical triggers are used to ignite the arcs. The evaporation targets are 2.5 inches in diameter by approximately 1.5 inches in depth, and the face of the targets are convexly shaped as shown in Figure 3.2 in order to aid in the confinement of the arc. Magnets located behind the mounted cathode targets are used to steer and keep the arc on the target.

The cathode targets for this study were purchased directly from S.E.G. The titanium targets had a purity of 99.97%, and the chromium target quality was 99.5%.

Both Ti and Cr targets had approximately the same size and shape as shown in Figure

3.2. Target life is generally estimated by amp hours, which is the product of the average current to a target and the total amount of time that current is delivered to the target.

Targets were replaced after being used for approximately 5000 amp hours in order to maintain process consistency.

Figure 3.2. Typical “dished” Ti target (2.5” Ø by 1.5” long) for S.E.G 2020 cathodic arc coating equipment.

The substrates for the CrN and (Ti,Cr)N samples were fixed in place as shown in

Figure 3.3. The planet/sun sample holder system can hold up to eight 1” x 1” substrates.

For the deposition of the binary TiN samples, a similar substrate holder system was used; however, it could only hold four samples. Both substrate holders are configured as planet/sun type systems with the individual substrates (planets) and the entire substrate system (sun) rotating in a clockwise direction for all of the sample depositions. A negative substrate bias was achieved using a 10 kW Advanced Energy MDX II DC power supply.

The main difference between the four and eight sample substrate holders was the shadow bar spacing which is used to improve coating uniformity. Shadow bars decrease the coating thickness and thus residual stresses along the edge of the coated samples since the average coating flux arriving at the edges is much greater than the center of the substrate due to rotation and the increased proximity of the sample edge to the evaporator source. The shadow bar width was set at 1.44 inches for the four sample holder when the sample geometry was one inch wide. This translates to approximately 0.22 inches between the edge of a one inch wide sample and the edge of the shadow bar.

In order to increase the number of substrates from four to eight for the CrN and the (Ti,Cr)N coated samples, the shadow bar width size was decreased to 1.17 inches or approximately 0.085 inches between the edge of a one inch wide sample and the edge of the shadow bar. In theory, this should decrease the edge thickness of the CrN and

(Ti,Cr)N as compared to the TiN coated samples.

3.2 Substrate Preparation

Three different substrate materials were investigated during this project. The majority of the samples were deposited on AM355 stainless steel which was described in chapter two. In particular, AM355 sheet stock with a 0.15 inch thickness was sheared into 2 inch by 1 inch coupons. The coupon surfaces were then prepped in a wet blaster with a slurry consisting of 5 lbs of 400 grit aluminum oxide powder and 5 gallons of water. The samples were blasted with an approximate 45° impingement angle which produced a surface finish of approximately 0.11 μm Ra. As discussed in chapter two, the wet blasting facilitates coating adhesion by removing surface oxide layers and

“activating” the surface for improved chemical bond strength between the metallic bond layer and the substrate. The increased surface roughness also increases the degree of mechanical bonding-interlocking mechanism by increasing the surface area of the coating substrate interface.

Polished 304 stainless steel (s.s.) buttons and (100) oriented silicon wafers were also used as substrates for coating structure comparison. The 1.675 inch Ø 304 s.s buttons were polished to a surface finish of approximately 0.02 μm Ra created by step wise vibratory polishing in 5 molar slurries consisting of water and 1, 0.5, and 0.05 μm Ø alumina powder, respectively. The silicon wafers (100) were cleaved into approximately

1 inch x 2 inch pieces while the stainless steel buttons were cut into 1” wide pieces in order to facilitate loading into the substrate holder.

After surface preparation, the AM355 samples were submerged in a solution of

Alconox soap and warm water to remove the wet blast slurry. The AM355 substrates were then rinsed with deionized (DI) water and dried with N2 gas. The AM355 substrates along with the stainless steel buttons were then ultrasonically cleaned in acetone at 40 °C for 20 minutes, rinsed with methanol, dried with N2, and then ultrasonically cleaned in methanol at 40 °C for 20 minutes. After the methanol ultrasonic cleaning the samples were once again rinsed with methanol and dried with N2. The Si wafer pieces were rinsed in acetone followed by a rinse in methanol. All of the substrates were then either loaded the same day into the cathodic arc chamber or placed in a vacuum sealed container and loaded into the chamber within three days to minimize surface oxidation and reduce dust contamination.

3.3 Cathodic Arc Deposition Processing Procedure

After loading the solvent cleaned samples into the deposition chamber, the cathodic arc system was evacuated overnight in order to achieve a base pressure of at least 2 x 10-6 Torr. Once the desired base pressure was achieved, a rate-of-rise test was conducted in which the system seals were evaluated by closing the system gate valve and measuring the rise in chamber pressure over a one minute period. A pressure rise of less than an order of magnitude was necessary in order to proceed to the next step in the deposition procedure which was the metal ion etch.

Prior to initiating the metal ion etch the radiant heater was set to 500 °C and the substrate rotation was started. For this project the substrate rotation was clockwise at 2.4 rpm for the sun and 6.75 rpm for the planets (substrates and shadow bars) as depicted in

Figure 3.3.

The metal ion etch precleaning of the samples are listed in Tables 3.1 through

Tables 3.6 for the various coating trials. The basic procedure involves triggering an arc with one or more evaporators and then accelerating the cations from those evaporator(s) towards the substrates using a large negative substrate bias of -1000 V. Typically, the current supplied to the arc source is pulsed in order to allow the samples to heat evenly.

If power was provided continuously, the outer surface of the samples could reach too high of a temperature and possibly cause degradation of substrate mechanical properties.

Following the substrate metal ion etching, the next step was coating deposition starting with the bond layer. During the bond layer deposition, the radiant heater set point was lowered to 325 °C and a lower bias potential than the metal ion etch was used in order to prevent excessive substrate re-sputtering. The desired bond layer thickness of

42 all coatings was approximately 0.4 μm. For the (Ti,Cr)N coatings, a substrate bias of -

150 V was always used for the Ti bond layer. The monolithic TiN and CrN deposited coatings used the same substrate bias for the bond layer deposition as the nitride layer and these values were either -50 V, -150 V, or -300 V for the TiN and -200 V for the CrN deposited coatings. Substrate bias should affect deposition conditions such as resputtering, substrate temperature, and ion energy and therefore microstructural differences as well as thickness differences are expected between the bond layers deposited at different substrate bias values.

Deposition of the nitride-based layer immediately followed the bond layer deposition. For the TiN, CrN and (Ti,Cr)N coated samples, N2 gas was injected into the vacuum chamber in order to react with the Ti (and/or Cr) ions to form a metal nitride either at the substrate surface or in the arriving coating flux. As discussed in chapter 2, nitrogen can also reduce the macroparticle generation by effectively raising the surface melting temperature of the targets.3 In addition to the nitrogen gas flow, the cathode evaporator currents, substrate biases, and primary deposition process parameters are listed in Tables 3.1 through 3.6. The deposition length for the nitride layer or layers was adjusted for different cathode currents and substrate biases in order to produce a coating with a total thickness of 12.5 μm (12.1 μm of nitride and 0.4 μm of metallic bond layer).

For the multilayer (Ti,Cr)N coatings with the Ti and Nb compliant interlayers the procedure for depositing the interlayer simply involved turning off the nitrogen gas flow and cathode current and igniting the target of the desired interlayer material. This process was repeated as necessary in order to achieve the desired number of layers with a total coating thickness of 12.5 μm. For the Nb interlayer materials, a high evaporator

43 current of 120 A was necessary in order to ignite and sustain the arc. The high current increased the substrate temperature during the interlayer deposition to around 450 °C. In contrast, during deposition of Ti, TiN, CrN or (Ti,Cr)N layers the typical substrate temperature was around 350 °C. It should be noted that the increased Nb evaporator current may have contributed to increased thermal relaxation due to increased substrate temperature and will be discussed in more detail in Chapter 4.

After the coating deposition, the chamber was typically allowed to cool below 30

°C before venting to atmosphere in order reduce oxidation and moisture absorption.

Chamber cleaning was performed after the deposition of each set of samples listed in

Tables 3.1 through 3.6. In addition to the chamber cleaning, evaporator targets were switched as necessary after the completion of sample sets. Chamber cleaning involved removing and then abrasive blasting steel wall paper sheets, mechanical triggers, the substrate holder, and thermal couples in order to remove coating buildup. The blasted components were solvent cleaned with acetone and methanol and reinstalled into the coating chamber.

3.4 Design of Experiments for Deposition of TiN, CrN and (Ti,Cr)N Coatings

The secondary objectives were to relate select deposition processing parameters to the coating microstructure and then correlate how this relationship affects the erosion performance of the coatings. Several series of experiments were designed to deposit coatings over a wide range of deposition conditions. The first series of experiments involved deposition of TiN coatings at different evaporator currents and substrate biases.

The TiN deposition parameters are listed in Table 3.1., where the evaporator source number refers to the evaporator source locations as shown in Figure 3.1.

The general format for the sample numbers is based on CYMD-Z where C stands for cathodic arc, YMD is the year, month, and date, and. -Z is the numbered run for that date. For example, C070912-1 is the first cathodic arc deposition trial performed on

September 12, 2007. It should be noted that the samples in the same set of experiments

(i.e. varying evaporator currents, bias, etc.) may not be deposited chronologically.

After deposition of the TiN coatings, a few CrN coating trails were produced in order to determine a range of processing parameters for monolithic CrN deposition. One of these CrN coatings, C080328-1, had similar deposition conditions (Table 3.2) to the

TiN and subsequent (Ti,Cr)N coatings and was included for comparative purposes.

There are several notable differences between monolithic CrN (C080328-1) and the other TiN and (Ti,Cr)N coatings included in this study. First, the metal ion etching for heating and cleaning of the substrates used chromium sources instead of titanium which may have affected coating adhesion by embedding or depositing a thin chromium layer on the substrate rather than titanium. Also, the bond layer for this sample was chromium rather than titanium which may have affected the adhesion, microstructure, and crystallographic orientation of the CrN layer and thus the coating properties such as erosion resistance.

A total of four sets of coatings experiments were designed in order to explore the

(Ti,Cr)N coating system. The sets of experiments deposited include:

1. Nanolayer (Ti,Cr)N coatings deposited with different Cr evaporator

currents (25 A, 45 A, 65 A, and 85 A), constant Ti evaporator current of

65 A and -150 V substrate biases to determine the effect of increased Cr

incorporation.

2. Nanolayer (Ti,Cr)N coatings deposited with different substrate biases (-50

V, -100 V, -150 V, -200 V, and -300 V) and constant Ti and Cr evaporator

currents of 65 A and 45 A, respectively, to determine the effect of ion

bombardment energy.

3. Multilayer (Ti,Cr)N coatings with Ti interlayers deposited with -150 V

substrate biases and constant Ti and Cr evaporator currents of 65 A and 45

A, respectively, having different numbers of the compliant Ti interlayers

(i.e. 2, 4, 8 and 16 layers of (Ti,Cr)N and 1, 3, 7, and 15 Ti interlayers).

4. Multilayer (Ti,Cr)N coatings with Nb interlayers deposited with -150 V

substrate biases and constant Ti and Cr evaporator currents of 65 A and 45

A, respectively, having different numbers of the compliant Nb interlayers

(i.e. 2, 4, 8 and 16 layers of (Ti,Cr)N and 1, 3, 7, and 15 Nb interlayers).

Nanolayer (Ti,Cr)N coatings deposited with different Cr evaporator currents

(constant Ti evaporator current of 65 A and -150 V substrate bias) were used to examine the effects of composition on the performance of the coatings. Four different Cr evaporator currents were used for in this set of experiments: 25 A, 45 A, 65 A, and 85 A, while keeping constant the Ti evaporator current at 65 A. This kept the Ti flux constant while allowing the Cr flux to vary resulting in different TiN/CrN ratios. The deposition parameters are listed in Table 3.3. The deposition time was adjusted from trial to trial to correspond to the change in Cr coating flux do to the varied Cr evaporator currents.

As with the binary TiN coatings, the effect of the substrate bias on the coating properties was investigated for the (Ti,Cr)N samples. Table 3.4 list the deposition

46 processing parameters for nanolayer (Ti,Cr)N coatings deposited with substrate biases of

-50 V, -100 V,-150 V,-200 V, and -300 V.

As will be discussed in detail in Chapter 4, the nanolayer (Ti,Cr)N coating deposited with a -150 V substrate bias voltage with 45 A Cr and 65 A Ti evaporator current set points was determined to be a nanolayered structure consisting of TiN rich and

CrN rich layers. Since this coating had a nanolayered structure, all of the (Ti,Cr)N layers in the multilayer (Ti,Cr)N coatings with Ti and Nb interlayers are believed to be the nanolayered TiN rich - CrN rich structure. In order to avoid confusion about nomenclature, the mixture of nanoscale TiN rich and CrN rich layers will be referred to as TiN rich and CrN rich nanolayers and the coating systems with the Ti and Nb interlayers will be referred to as multilayer (Ti,Cr)N coatings. A schematic representing this nomenclature is shown in Figure 3.4.

Figure 3.4. Schematic representing multilayer (Ti,Cr)N coating system with compliant Ti or Nb interlayers. (Ti,Cr)N is nanolayered structure of alternating TiN and CrN rich layers.

Tables 3.5 and 3.6 lists the deposition processing parameters for the (Ti,Cr)N multilayer coatings with Ti and Nb interlayers, respectively. The coated samples with multiple layers of (Ti,Cr)N and interlayer material are comprised of two, four, eight and

47 sixteen nitride-based layers and subsequently one, three, seven and fifteen compliant (Ti, or Nb) interlayers, respectively. The target volume percent ratio of the nitride layers to the compliant metal interlayer(s) was 10:1. The evaporator current set points for the

(Ti,Cr)N nanolayers was based on the deposition rates of the binary CrN and TiN coatings depositied under similar conditions. The binary coatings suggusted that a Ti evaporator current of 65 A and Cr evaporator of 45 A would produce a coating with an approximate 1:1 ratio of Ti:Cr atoms. Energy Dispersive Spectroscopy (EDS) was used to semiquantify that the ratio was close to 1:1 for these coatings.

There are a few notable processing differences between the multilayer coatings containing Ti interlayers and all of the other (Ti,Cr)N coatings. First, the multilayer coatings with Ti interlayers used two evaporator sources during metal ion etching step and subsequent Ti bond layer deposition whereas all the other nanolayer (Ti,Cr)N coatings sets used only one Ti evaporator source. These processing changes were adapted in order to keep future samples more consistent since only one Ti evaporator source could be installed for multilayer (Ti,Cr)N coatings with Nb interlayers because the cathodic arc PVD system is limited to three total evaporation sources.

Another processing optimization was the addition of argon to aid in the ignition of the Ti target for the bond layer deposition. Argon was experimentally determined to aid the stability of the arc which resulted in less mechanical retriggering and therefore a more consistent flux of Ti atoms arriving at the substrate surface.

Table 3.1. Deposition processing parameters for TiN coatings.

Run Number C070912-1 C071017-1 C070917-1 C071024-1 C071025-1 Substrates Four AM355 Four AM355 Four AM355 Four AM355 Four AM355 Substrate ‘Planet’ Rotation (RPM) 6.75 6.75 6.75 6.75 6.75 Substrate ‘Sun’ Rotation (RPM) 2.4 2.4 2.4 2.4 2.4 Radiant Heater (ºC) 500 500 500 500 500 Evaporator # (See Fig. 3.1) 1 and 3 1 and 3 1 and 3 1 and 3 1 and 3 Ti Metal Ion Evaporator Current(s) (A) 65 45 85 65 65 Etching Substrate Bias (V) -1000 -1000 -1000 -1000 -1000 Parameters 30 sec. on, 30 30 sec. on, 30 30 sec. on, 30 30 sec. on, 30 30 sec. on, 30 Length sec off for 5 sec off for 5 sec off for 10 sec off for 10 sec off for 5 min. min. min. min. min. Radiant Heater (ºC) 325 325 325 325 325 Evaporator # (See Fig. 3.1) 1 and 3 1 and 3 1 and 3 1 and 3 1 and 3 Ti Bond Evaporator Current(s) (A) 65 45 85 65 65 layer Parameters Argon Flow (sccm) 0 0 0 0 0 Substrate Bias (V) -150 -150 -150 -25 -300 Length (min) 8 8 8 8 8 Radiant Heater (ºC) 325 325 325 325 325 Evaporator # (See Fig. 3.1) 1 and 3 1 and 3 1 and 3 1 and 3 1 and 3 TiN Layer Evaporator Current(s) (A) 65 45 85 65 65 Parameters Substrate Bias (V) -150 -150 -150 -25 -300

N2 Flow Rate (sccm) 250 250 250 250 250 Length (min) 165 165 165 165 165 Total Coating Deposition Length (min) 173 173 173 173 173

Table 3.2. Deposition processing parameters for CrN coating.

Run Number C080328-1 Substrates Four AM355 Substrate ‘Planet’ Rotation (RPM) 6.75 Substrate ‘Sun’ Rotation (RPM) 2.4 Radiant Heater (ºC) 500

Cr Evaporator # (See Fig. 3.1) 1 and 3 Metal Ion Etching Evaporator Current(s) (A) 45 Parameters Substrate Bias (V) -1000 Length 15 sec. on, 15 sec off for 5 min. Radiant Heater (ºC) 325 Evaporator # (See Fig. 3.1) 1 and 3 Cr Bond layer Evaporator Current(s) (A) 45 Parameters Argon Flow (sccm) 0 Substrate Bias (V) -200 Length (min) 10 Radiant Heater (ºC) 325 Evaporator # (See Fig. 3.1) 1 and 3 CrN Layer Evaporator Current(s) (A) 45 Parameters Substrate Bias (V) -200

N2 Flow Rate (sccm) 250 Length (min) 225 Total Coating Deposition Length (min) 235

Table 3.3. Deposition processing parameters for nanolayer (Ti,Cr)N coatings deposited with different Cr evaporator currents (constant 65 A Ti evaporator current and constant -150 V substrate bias) Run Number C080530-1 C080602-1 C080529-1 C080528-1 7 AM355, 1 7 AM355, 1 7 AM355, 1 7 AM355, 1 Substrates polished 304 s.s. polished 304 s.s. polished 304 s.s. polished 304 s.s. button button button button Substrate ‘Planet’ Rotation (RPM) 6.75 6.75 6.75 6.75 Substrate ‘Sun’ Rotation (RPM) 2.4 2.4 2.4 2.4 Radiant Heater (ºC) 500 500 500 500 Ti Evaporator # (See Fig. 3.1) 2-Ti 2-Ti 2-Ti 2-Ti Metal Ion Evaporator Current(s) (A) 65 65 65 65 Etching Parameters Substrate Bias (V) -1000 -1000 -1000 -1000 15 sec. on, 15 sec 15 sec. on, 15 sec 15 sec. on, 15 sec 15 sec. on, 15 sec Length off for 10 min. off for 10 min. off for 10 min. off for 10 min. Evaporator # (See Fig. 3.1) 2-Ti 2-Ti 2-Ti 2-Ti Ti Bond Evaporator Current(s) (A) 65 65 65 65 layer Substrate Bias (V) -150 -150 -150 -150 Parameters Argon Flow Rate (sccm) 100 100 100 100 Length (min) 16 16 16 16 Evaporator # (See Fig. 3.1) 1-Cr and 2-Ti 1-Cr and 2-Ti 1-Cr and 2-Ti 1-Cr and 2-Ti Evaporator Current(s) (A) 25-Cr and 65-Ti 45-Cr and 65-Ti 65-Cr and 65-Ti 85-Cr and 65-Ti (Ti,Cr)N Substrate Bias (V) -150 -150 -150 -150 Parameters N2 Flow Rate (sccm) 250 250 250 250 (Ti,Cr)N Layer (min) 235 205 160 135 Total Coating Deposition Length (min) 251 221 176 151

Table 3.4. Deposition processing parameters for nanolayer (Ti,Cr)N coatings deposited with varying substrate biases at constant Ti and Cr evaporator currents. Run Number C080519-1 C080520-1 C080602-1 C080521-1 C080522-1 Seven AM355, 1 Seven AM355, 1 Seven AM355, 1 Seven AM355, 1 Seven AM355, 1 Substrates polished 304 s.s. polished 304 s.s. polished 304 s.s. polished 304 s.s. polished 304 s.s. button button button button button Substrate ‘Planet’ Rotation (RPM) 6.75 6.75 6.75 6.75 6.75 Substrate ‘Sun’ Rotation (RPM) 2.4 2.4 2.4 2.4 2.4 Radiant Heater (ºC) 500 500 500 500 500 Ti Evaporator # (See Fig. 3.1) 2-Ti 2-Ti 2-Ti 2-Ti 2-Ti Metal Ion Etching Evaporator Current(s) (A) 65 65 65 65 65 Parameters Substrate Bias (V) -1000 -1000 -1000 -1000 -1000 15 sec. on, 15 sec 15 sec. on, 15 sec 15 sec. on, 15 sec 15 sec. on, 15 sec 15 sec. on, 15 sec Length off for 10 min. off for 10 min. off for 10 min. off for 10 min. off for 10 min. Evaporator # (See Fig. 3.1) 2-Ti 2-Ti 2-Ti 2-Ti 2-Ti Ti Bond Evaporator Current(s) (A) 65 65 65 65 65 layer Substrate Bias (V) -150 -150 -150 -150 -150 Parameters Argon Flow Rate (sccm) 100 100 100 100 100 Length (min) 16 16 16 16 16 Evaporator # (See Fig. 3.1) 1-Cr and 2-Ti 1-Cr and 2-Ti 1-Cr and 2-Ti 1-Cr and 2-Ti 1-Cr and 2-Ti Evaporator Current(s) (A) 45-Cr and 65-Ti 45-Cr and 65-Ti 45-Cr and 65-Ti 45-Cr and 65-Ti 45-Cr and 65-Ti (Ti,Cr)N Substrate Bias (V) -50 -100 -150 -200 -300 Parameters N2 Flow Rate (sccm) 250 250 250 250 250 (Ti,Cr)N Layer (min) 260 190 205 215 235 Total Coating Deposition Length (min) 276 206 221 231 251

Table 3.5. Deposition processing parameters for multilayer (Ti,Cr)N coatings with Ti Interlayers. Run Number C080602-1 C080416-1 C080422-1 C080424-1 C080430-1 Seven AM355, Six AM355, 304 Six AM355, 304 Substrates Eight AM355 Eight AM355 One polished 304 s.s., Si wafer piece s.s., Si wafer piece Substrate ‘Planet’ Rotation (RPM) 6.75 6.75 6.75 6.75 6.75 Substrate ‘Sun’ Rotation (RPM) 2.4 2.4 2.4 2.4 2.4 Radiant Heater (ºC) 500 500 500 500 500 Ti Evaporator # (See Fig. 3.1) 2-Ti 2-Ti, and 3-Ti 2-Ti, and 3-Ti 2-Ti, and 3-Ti 2-Ti, and 3-Ti Metal Ion Evaporator Current(s) (A) 65 65 65 65 65 Etching Parameters Substrate Bias (V) -1000 -1000 -1000 -1000 -1000 15 sec. on, 15 sec 15 sec. on, 15 sec 15 sec. on, 15 sec 15 sec. on, 15 sec 15 sec. on, 15 sec Length off for 5 min. off for 5 min. off for 5 min. off for 5 min. off for 5 min. Radiant Heater (ºC) 325 325 325 325 325 Evaporator # (See Fig. 3.1) 2-Ti 2-Ti, and 3-Ti 2-Ti, and 3-Ti 2-Ti, and 3-Ti 2-Ti, and 3-Ti Ti Bond Evaporator Current(s) (A) 65 65 65 65 65 layer Parameters Argon Flow (sccm) 100 0 0 0 0 Substrate Bias (V) -150 -150 -150 -150 -150 Length (min) 16 8 8 8 8 Evaporator # (See Fig. 3.1) 1-Cr and 2-Ti 1-Cr and 2-Ti 1-Cr and 2-Ti 1-Cr and 2-Ti 1-Cr and 2-Ti Evaporator Current(s) (A) 45-Cr and 65-Ti 45-Cr and 65-Ti 45-Cr and 65-Ti 45-Cr and 65-Ti 45-Cr and 65-Ti (Ti,Cr)N Substrate Bias (V) -150 -150 -150 -150 -150 Layer Parameters N2 Flow Rate (sccm) 250 250 250 250 250 Number of (Ti,Cr)N Layers 1 2 4 8 16 (Ti,Cr)N Layer (min) 205 97 75 37 18.3 Evaporator # (See Fig. 3.1) NA 3-Ti 3-Ti 3-Ti 3-Ti Ti Evaporator Current(s) (A) NA 65 65 65 65 Interlayers Substrate Bias (V) NA -150 -150 -150 -150 Length/Ti Layer (min) NA 20 7 3.5 1.7

Table 3.6. Deposition processing parameters for multilayer (Ti,Cr)N coatings with Nb Interlayers. Run Number C080602-1 C080502-1 C080505-1 C080507-1 C080508-1 Seven AM355, Seven AM355, Seven AM355, Seven AM355, Seven AM355, Substrates 304 s.s. button 304 s.s. button 304 s.s. button 304 s.s. button 304 s.s. button Substrate ‘Planet’ Rotation (RPM) 6.75 6.75 6.75 6.75 6.75 Substrate ‘Sun’ Rotation (RPM) 2.4 2.4 2.4 2.4 2.4 Radiant Heater (ºC) 500 500 500 500 500 Ti Evaporator # (See Fig. 3.1) 2-Ti 1-Ti 1-Ti 1-Ti 1-Ti Metal Ion Evaporator Current(s) (A) 65 65 65 65 65 Etching Parameters Substrate Bias (V) -1000 -1000 -1000 -1000 -1000 15 sec. on, 15 sec 15 sec. on, 15 sec 15 sec. on, 15 sec 15 sec. on, 15 sec 15 sec. on, 15 sec Length off for 5 min. off for 5 min. off for 5 min. off for 5 min. off for 5 min. Radiant Heater (ºC) 325 325 325 325 325 Evaporator # (See Fig. 3.1) 2-Ti 2-Ti 2-Ti 2-Ti 2-Ti Ti Bond Evaporator Current(s) (A) 65 65 65 65 65 layer Parameters Argon Flow (sccm) 100 100 100 100 100 Substrate Bias (V) -150 -150 -150 -150 -150 Length (min) 16 16 16 16 16 Evaporator # (See Fig. 3.1) 1-Cr and 2-Ti 1-Cr and 2-Ti 1-Cr and 2-Ti 1-Cr and 2-Ti 1-Cr and 2-Ti Evaporator Current(s) (A) 45-Cr and 65-Ti 45-Cr and 65-Ti 45-Cr and 65-Ti 45-Cr and 65-Ti 45-Cr and 65-Ti (Ti,Cr)N Substrate Bias (V) -150 -150 -150 -150 -150 Layer Parameters N2 Flow Rate (sccm) 250 250 250 250 250 Number of (Ti,Cr)N Layers 1 2 4 8 16 (Ti,Cr)N Layer (min) 205 97 74 37 18.3 Evaporator # (See Fig. 3.1) NA 3-Nb 3-Nb 3-Nb 3-Nb Nb Evaporator Current(s) (A) NA 120 120 120 120 Interlayers Substrate Bias (V) NA -150 -150 -150 -150 Length/Nb Layer (min) NA 20 10.4 5.5 2.7

53 54

3.5 Coating Characterization and Evaluation

3.5.1 X-Ray Diffraction (XRD)

3.5.1.1 X-Ray Diffraction for Crystallographic Structure Determination

A Philips X’Pert XRD-MRD X-ray diffractometer was used for phase analysis of all the coated samples. X-ray diffraction patterns were obtained by irradiation of the coatings using a copper anode (voltage of 45 kV and current of 40 mA) incident through a cross-slit collimator (1.2 mm x 1.0 mm) at an angle of theta parallel to the substrate surface. At specific incident angles, the periodicity of the sample’s crystallographic structure results in X-rays to scatter coherently and the constructive interference of these is referred waves is referred to as diffraction. The diffraction of the X-rays can be predicated according to Bragg’s law:

λ  hkl sind2n θ where λ = wavelength of the incident X-ray, dhkl is the inter atomic spacing between the planes defined by hkl, θ is the angle of diffraction, and n is the order of the reflection

(integer).63,64 Figure 3.5 shows diffraction of constructive interference of X-rays at an incident angle of θ to a crystallographic plane hkl.

Figure 3.5. Diffraction of an incident X-ray beam at an incident angle of θ.

By scanning the sample over a wide range of diffraction angles (θ), the crystallographic structure/phase of the material can be determined by measuring the angles where diffraction occurs and comparing these values to the standards published by the International Center for Diffraction Data (ICDD).65

All of the coated samples in this study were subjected to a θ/2θ scan in which the incident beam is held at a fixed angle and the sample and the detector are continuously rotated θ and 2θ, respectively. Additionally, the binary TiN samples were subjected to a

2θ-only or glancing angle scan in which the incident beam and sample are held at a fixed angle, ω, relative to each other and only the detector is continuously rotated (2θ). This results in diffraction intensities from planes that change with the position of the detector and not just the planes parallel to the coating surface. The XRD scan conditions for both the θ/2θ and 2θ-only scan types are given in Table 3.7, where ω=1° was used for the 2θ- only scan. Kα2 stripping, peak smoothing, and background corrections were performed prior to determining the peak locations using the full-width-half-max (FWHM) method and the relative integrated intensities were measured using Philips PC-PDF software.

The relative intensities were then compared to the ICDD standards in order to qualitatively show preferred crystallographic orientations of the films.

Table 3.7. XRD scan parameters for nitride coati ngs. Step Time per Scan Initial 2θ Final 2θ Scan Mode λ (nm) Width Step (sec) Type (degrees) (degree s) (degrees ) θ/2θ Continuous 0.154186 30 135 0.03 1.5 2θ-Only Continuous 0.154186 30 135 0.03 1.5

3.5.1.2 XRD Crystallite Size

Small grains in the direction perpendicular to the coating or sample surface can cause significant broadening of the XRD peaks. By comparing the broadening of the peak (B) to the intrinsic broadening of the XRD equipment (b), the crystallite size (τ) of the sample can be estimated with reasonable accuracy using the Scherrer equation:

K   cos B where K is the ‘shape factor’ which is typically between 0.89 and 1.34, λ is the characteristic X-ray wavelength (Å), and θB is the Bragg angle (radians), and βτ is defined as,

22   )bB()bB( which is in radians and determined at the half-intensity maximum.65,66

For the TiN and CrN samples in this study, the intrinsic instrumentation broadening was determined to be 9.512x10-3 radians using Phillips PC-APD software to correct the background, strip the Cu Kα2 interference, and smooth the peak shape for the

(111) diffracted plane intensity peak of an unstressed polycrystalline Si standard which was scanned using the same θ/2θ scan parameters as discussed in section 3.5.1. The background corrections and peak smoothing were also used for the (Ti,Cr)N coated samples, but the (Ti,Cr)N samples were not subjected to the Cu Kα2 stripping due to concerns about peak overlaps from the TiN and CrN nanolayers. Therefore, the Si peak width used for the (Ti,Cr)N coated samples was slightly larger (9.547x10-3), and a weighted average of Cu Kα1 and 0.5 λ Kα2 wavelength of approximately1.54186 Å was

used for (Ti,Cr)N peak positions while Cu Kα1 (1.5406 Å ) was used for the TiN and CrN samples.66

A ‘shape factor’ K value of 0.94 was used which is a typical value for small cubic structures.66 For all of the samples, the set of i plane peak intensity was used since it had high intensity and symmetrical Gaussian peak shape. The error for this method is estimated to be relatively large and Suryanarayana and Norton suggest it was at least ±

10%.66 In order to better determine the broadening contributions of crystallite size and strains, deconvolution of the measured integrated peaks can be performed since the stress-induced diffraction peak broadening based on a tan θ function while crystallite size

1 broadening has a dependence.66 cos

3.5.1.3 XRD Residual Stress Analysis

The sin2Ψ residual stress analysis technique provided a quantitative method to calculated the residual stress from the measured strain acting on the deposited coatings.

Sin2Ψ residual stress analysis compares the 2θ location of diffracted planes over a series of scans with different surface angles of inclination relative to the X-ray source/substrate surface. The strain present in the coating is determined from the change in the d-spacing as a function of the angle of inclination (Ψ), which is the angle between the normal to the sample surface and the bisector of the angle between the incident and reflected X-ray beam. The residual stress is calculated from the slope of the plot of interatomic spacing as a function of sin2 Ψ, and residual stress can be calculated from the following equation:

 E  1  d   hkl       2  1   hk( )l d0  sin  

where σ  is the residual coating stress, Ehkl and  are the elastic modulus and Poisson’s

ratio, respectively, for the plane of interest, d0 is the equilibrium interatomic spacing for

d Ψ equals zero (no sample tilt), and  is the slope of interactomic spacing versus sin 2  sin2Ψ.67

The sin2Ψ residual stress measurements were conducted on the (111) and (422) diffraction planes of all the TiN samples. The sin2Ψ residual stress measurements of the

(111) and (422) planes were conducted with Ψ values of 15°, 30°, 45°, and 60° since Ψ =

0 was aligned with the crystallographic orientation and therefore had too high of a

diffraction intensity compared to the other Ψ angles. d0 was taken as the Y-intercept of the d-spacing vs. sin2 Ψ curve. The scan parameters were 0.02° per step and 20 sec per step. The 2θ range was adjusted depending on the peak location and scan parameters, but typically no less than 6°. It should be noted that the residual stress calculations determined by X-ray diffraction assumes a biaxial stress condition (plane stress condition) which is a reasonable assumption for coating samples such as the cathodic arc deposition TiN, CrN, and (Ti,Cr)N.

3.5.2 Cross Sections for Optical Microscopy (OM)

One sample from each set was cross sectioned using a diamond impregnated cutting blade. The cross sections were then mounted in vacuum infiltrated epoxy and ground using 240, 320, 600, 800, and 1200 grit silicon carbide paper using water as lubrication. The samples were then polished using 3 and 1 μm diamond spray and then a

0.05 μm colloidal silica solution. The cross sections where used for the Vicker’s micro- indention and for optical microscopy measurements of coating thickness.

3.5.3 Electron Probe Microanalysis (EPMA)

The composition of select coatings was determined by electron probe microanalysis using a Cameca Camebox SX50 using a TiN standard. EPMA is a quantitative elemental analysis technique with spatial resolutions of approximate 1 μm.

EPMA can provide very accurate compositional results with sensitivities up to 1000 ppm depending on the equipment working environment, but the resolution is poor for small atoms in heavy lattices. For the TiN system, the NK peak overlaps the Tiβ peak, but it is possible to deconvolute the peaks and obtain good measurements. Lastly, standard ZAF

(atomic number (Z), absorption (A) and fluorescence (F)) corrections were made for matrix corrections.68

3.5.4 Environmental Scanning Electron Microscopy (ESEM)

An environmental scanning electron microscope (Quanta 200) was used to examine both the top surface and fracture surfaces of the coated samples deposited by cathodic arc on the AM355. A electron microscope bombards the sample surface with a focused beam of electrons which interact with the sample structure, topography, and elementary composition to produce secondary electrons (SE). The incident electrons can also be scattered away from the sample surface and these electrons are referred to as back scattered electrons (BS).64 In addition to SE and BS electrons, the incident electron beam can produce X-ray emissions which have characteristic wavelengths dependent on the elemental composition of the coated sample. These characteristic X-ray emission lines are used in Energy Dispersive Spectroscopy (EDS) to semiquantify the coating compos ition.69

Fracture surfaces of the samples were produced by first cutting through the backside of the coated AM355 samples with the diamond impregnated saw blade in order to minimize the force needed to fracture the sample. The samples were submerged in liquid N2 with pliers for a minimum of 10 minutes to embrittle the coating sample. The samples were then removed from the liquid N2 and the fracture surfaces were quickly prepared by applying a tensile force with another set of pliers. Figure 3.6 illustrates the fracture surface procedure. ESEM images of the fracture surfaces provided qualitative information on the grain structure and also quantitative measurements of the entire coating along with individual layer thicknesses.

Figure 3.6. Fracture surface methodology for the nitride coating samples.

3.5.5 Quantative Microstructural Analysis

Clemex Vision Professional Edition (P.E.) image analysis software was used to perform quantative microstructral analysis and calculate the coating surface area covered by macroparticles defect concentrations as compared to the entire coating surface area.

The routine for the Clemex calculation is reported in Figure 3.7 and was performed on select groups of samples using a minimum of five top surface ESEM images for each

61 coating. This analysis was only used to show that there is a trend between the surface appearance and the deposition conditions. It should be noted that the resulting measurements do not indicate whether the defect is adsorbed on the surface or whether it is embedded within the coating, but since macroparticle diameters are around the same magnitude in size as the coating thicknesses this is not believed to be a significant issue.

Also, the image analysis software routine cannot distinguish the difference between the type of defect (i.e. macroparticle, physical damage, contamination, etc.).

Figure 3.7. Analytic software routine developed for Clemex Vision P.E. to assess large scale defect incorporation in TiN, CrN and (Ti,Cr)N coatings.

3.5.6 Annular Dark Field Scanning Transmission Electron Microscope (ADF- STEM)

ADF-STEM is an analytical transmission electron microscopy technique which uses a dark field detector that captures electrons scattered at high angles. Elemental mapping can be performed due to the high contrast in signal intensity from different elements. The resolution of (ADF-STEM) can approach atomistic levels depending on the equipment, so this technique is very useful for nanoscale coating analysis.70

3.5.7 Vicker’s Micro-Indention Hardness Testing

The Vicker’s Hardness number (HV) was determined using a LECO Model M-

400-Gl Vicker’s micro-indenter. The Vicker’s indenter uses a four sided square-based diamond pyramid tip with 136° between faces and an applied load to plastically deform the sample. HV calculated by the equation below and is directly related to the material elastic modulus and therefore related to the resistance to plastic deformation:

1 sin L 2 L sin )136( number) hardness sVicker'(H hardness number)  2 V d 2 where L is the load in kg and d is the length of the mean diagonal in mm. The units the

Vicker’s hardness number are kg/mm2.71

For the TiN, CrN and the multilayer (Ti,Cr)N with the Ti interlayers, the AM355 samples were cut into approximately 0.25 inch by 0.25 inch sections and then polished on a vibratory polisher with a one molar slurry of alumina powder and DI water. In particular, the TiN and CrN samples were polished for approximately eight hours with a slurry of Metlab brand 0.05 μm alumina powder (labeled gamma c alumina powder).

The multilayer (Ti,Cr)N samples with the Ti interlayers were first polished with a slurry

63 of Metlab brand 1 μm alpha alumina powder for eight hours then polished for 16 hours with the 0.05 μm alumina slurry (gamma c). The measurements for the other (Ti,Cr)N samples were made on the pre-polished 304 stainless steel substrates that did not require modification to the analyzed surface prior to the measurements due to their relatively smooth surfaces.

The load used for the HV measurements was either 25 or 50 grams depending on the coating thickness. Between 12 and 15 measurements were made for each sample. It should be noted that for the thinner coatings, the Vicker’s hardness value is most likely a composite hardness value of both the coating and the substrate.

3.5.8 Scratch Adhesion Testing

Scratch adhesion measurements were made using a diamond tipped stylus to scratch the coatings with increasing vertical load as the stylus tip is traversed along the coating surface. The load is typically incrementally until the coating fails and delaminates in the scratch mark vicinity. The force required to cause debonding from the substrate is referred to as the critical force, Fc. The coating failure under constant loading conditions is often determined by a combination of acoustic emission signals (caused by the high frequency associated with detachment) and OM or SEM.38

3.5.9 Surface Roughness Measurements

Surface roughness and waviness of the samples were measured in order to examine the effects of select deposition parameters on the coating surface finish. In general, the difference between surface roughness and waviness is matter of scale, where roughness is defined as fine irregularities in the surface texture such as feed marks from

64 machining. Waviness is defined as more of a macro-scale deviation in surface finish created by deflections, vibrations, chatter, and etc.

The surface roughness and waviness of the coating samples was conducted using a Tencor P-10. A surface roughness average, Ra, was recorded and is defined as the arithmetic average of the absolute value of the height deviation from the centerline height of the sample to the valleys and hills of the sample. Ra is defined by the following equation:

1 Lx Average) Roughness(R Average)  dxy a  L x0 where |y| is the absolute value of the deviation, L is the length of the scan, and x is the

72 scan direction parallel to the coating surface. Figure 3.8 illustrates Ra of a coated sample. The Ra values were compiled using the average of twelve 500 μm long scans

(200 Hz sampling and 20 μm/sec scan rates) of the various coatings deposited on the wet blasted AM355 substrates. Outliers were eliminated using a modified Thompson-Tau technique.

Figure 3.8. Illustration depicting surface roughness, Ra, of a coated substrate.

3.5.10 Erosion Testing

The high velocity hard particle erosion testing was conducted using an in-house fabricated erosion rig. The system uses a vibratory feeder assisted with an air siphon to inject the erosion media into a stainless steel chamber where it is mixed with high pressure air up to 100 psi. The particle/air mixture is then feed through a pipe with a 0.55 inch inside diameter into a chamber which is under vacuum. For most erosion test sets of experiments, the pipe nozzle was fixed at 8.5 inches above the coating surface. The erosion rig is shown in Figure 3.9.

Figure 3.9. High Velocity hard particle erosion testing equipment designed and built by the ARL Advanced Coating Department at the Pennsylvania State University.

The experimental conditions for the hard particle erosion tests are listed in Table

3.8. The spherical glass bead media used for the test was supplied by Powder

Technologies Incorporated (Size-13, Batch S.R. 082407, Pallet No 002). Particle size distribution analysis was determined using dynamic light scattering by the Particle

Characterization Laboratory at the Pennsylvania State University. The particle size distribution for the glass beads is shown Figure 3.10 (a), and the average particle size was approximately 62 μm in diameter. The 120 grit alumina abrasive particles were also provided by Powder Technologies Incorporated, and Figure 3.10 (b) shows the particle size distribution with the mean particle size being approximately 138 μm in diameter. It should be noted that this media was not spherical (Figure 3.11 (a)) in shape as compared to the glass beads (Figure 3.11(b)), and this resulted in a more aggressive test from the gauging mechanism related to non-spherical media combined with the Al2O3 mean particle size being almost twice the diameter as the glass bead.

Table 3.8. Erosion test parameters for glass bead erosion tests. Nozzle to Particle .Estimated Erosion Spot Impingement Main Air Line Sample Media Feed Rate Velocity Diameter Angle Pressure (psi) Distance (g/min) (m/s) (in) (degrees) (in) 62μm 0.25 Glass 100 50 150 90 8.5

Bead 120 Grit 0.25 100 56 150 60 8.5 Alumina

Volume (%) Volume (%) 50 100 50 100

90 90

40 80 40 80

70 70

30 60 30 60

50 50

20 40 20 40

30 30

10 20 10 20

10 10

0 0 0 0 10.0 100.0 1000.0 10000.0 100000.0 1000000.0 10.0 100.0 1000.0 10000.0 100000.0 1000000.0 Particle Diameter (µm.) Particle Diameter (µm.) (a) (b) Figure 3.10. Particle size distribution for (a) Size-13 glass bead and (b) 120 grit alumina determined by light scattering.

(a) (b) Figure 3.11. ESEM micrograph showing (a) angular particle shape of 120 grit alumina media and (b) spherical particle shape of #13 glass beads.

Erosion resistance was measured by mass loss. The samples were cleaned and then weighed using a five significant digit scale (0.00001 g). A minimum of three measurements were taken, and the standard deviation for the values was typically less

68 than 0.00004 g. One sample at a time was eroded, and the center point of the 0.25 inch diameter erosion spot was aligned with the approximate center point of the nozzle. The media was then loaded into the vibratory feed assembly and the air valves were opened to begin the test.

For glass bead erosion tests, the erosion of the coated samples was carried out incrementally to better understand the erosion behavior. Depending on the coating sample properties, the samples were subjected to a small amount of media (25 g or 50 g).

If the sample showed little to no mass loss, the erosion test was repeated in the same spot as in the original test. Generally, the amount of erosive media used in this erosion rig test assembly for subsequent tests varied between 25 g and 1000 g depending on the samples previous erosion performance. The samples were eroded until either they reached approximately 3000 g of total erodent used or a set cut off for the total mass loss was reached. A minimum of two erosion tests were performed in order to evaluate coating erosion resistance. For the alumina-based erosion tests, the samples were subjected to only one 25 g test due to the particle angularity and large diameter.

For the glass bead erosion test, the mass loss cut off point was calculated by estimating the density of binary TiN and binary CrN mixture containing a 1:1 ratio of Ti to Cr. The theoretical lattice constants of TiN and CrN are 4.240 Å and 4.140 Å, respectively.50,51 Assuming that a mixture of the two phases would retain the same NaCl structure then the lattice constant would increase up to a value of approximately 4.190 Å.

The lattice would then consist of two Ti atoms, two Cr atoms, and 4 N atoms. Using the atomic masses of Ti, Cr, and N, the density of the mixture having a 4.190 Å lattice parameter value would be approximately 5.774 g/cm3.64

Erosion failure was determined to occur once the coating volume fraction loss exceeded 50% of the total volume of the theoretical coating with respect to the erosion spot diameter. Since the monolithic TiN and CrN coatings had a total monolithic nitride thickness of approximately 12.1 μm, the corresponding cut-off depth was 6.05 μm.

Using the calculated density from above, a 6.05 μm deep erosion scar of the mixture would correspond to mass loss of approximately 0.00111 g from a 0.25 inch (6,350 μm) diameter erosion scar. This cut off value was used for all of coatings produced in these experiments.

CHAPTER 4

RESULTS AND DISCUSSION

The first section of Chapter 4 discusses the compositional and microstructral characteristics of the various coatings deposited by cathodic arc using the XRD, EPMA,

SEM, and STEM. These analytical results were used to characterize the observed trends for deposition rate, Vicker’s micro-indenture for hardness, scratch adhesion, surface roughness, and hard particle erosion. From these correlations, general trends were observed which suggest how to optimize the cathodic arc deposition conditions for production of hard particle erosion resistant nanolayer (Ti,Cr)N coatings with improved performance.

4.1 XRD Crystallographic Structure Determination from Θ/2Θ and Glancing Angle Scans

In order to determine the crystallographic structure of the deposited coatings

Θ/2Θ XRD scans were performed as described in Section 3.5.1.1, and compared to the

Powder Diffraction Files (JCPDF’s) published by the International Centre for Diffraction

Data (ICDD), formally known as the Joint Committee on Powder Diffraction Standards

(JCPDS). Additionally, an XRD pattern of a non-coated AM355 substrate was obtained in order to distinguish diffraction from the coating substrate, especially for those samples with thin coatings.

4.1.1 XRD Crystallographic Structure Results for Monolithic TiN

Θ/2Θ and glancing angle XRD patterns (Figures 4.1 through 4.5) confirmed that the TiN coatings consisted primarily of B1 NaCl structure ( TiN JCPDF 01-30-1420).50

Comparing Figures 4.1 through 4.5, all of the monolithic TiN coatings deposited show a strong (111) preferred crystallographic orientation and is discussed in more detail in

Section 4.3. There are noticeable differences between the θ/2θ and glancing angle XRD patterns for monolithic TiN coatings, including peak location shifts due primarily to differences in residual stresses between the bulk and the top surface region of the coating.

Figure 4.1. θ/2θ and glancing angle XRD of TiN coating deposited with an evaporator current set point of 65 A and a -150 V substrate bias. TiN planes are indexed.

Figure 4.2. θ/2θ and glancing angle XRD of TiN coating deposited with an evaporator current set point of 45 A and a -150 V substrate bias. TiN planes are indexed.

Figure 4.3. θ/2θ and glancing angle XRD of TiN coating deposited with an evaporator current set point of 85 A and a -150 V substrate bias. TiN planes are indexed.

Figure 4.4. θ/2θ and glancing angle XRD of TiN coating deposited with an evaporator current set point of 65 A and a -25 V substrate bias. TiN planes are indexed.

Figure 4.5. θ/2θ and glancing angle XRD of TiN coating deposited with an evaporator current set point of 65 A and a -300 V substrate bias. TiN planes are indexed.

Another difference between the θ/2θ patterns and 2θ-only patterns is that for the θ

/2θ patterns there are distinguishable peaks that correspond to the AM355 substrate material resulting from deeper X-ray penetration. There are also peaks in the θ/2θ patterns which do not appear in the glancing angle patterns indicating that these peaks may belong to the Ti bond layer and are correspondingly close in position to the locations for the hexagonal α-Ti phase (JCPDF card number 01-44-1294).73 It should be noted that due to peak overlaps with TiN, combined with low intensity values (counts/sec), it is difficult to confirm with exact certainty whether these peaks are from the α-Ti phase; however, the α-Ti phase is the low temperature (below 883 °C) stable Ti phase and should be the phase of the Ti bond layer.74 Lastly, it should be noted that there is a correlation with regards to the intensities of the apparent Ti phase decreasing with increasing monolithic TiN coating thickness strongly suggesting that the diffraction is from the Ti bond layer.

4.1.2 XRD Crystallographic Structure Results for the CrN Coating

The θ/2θ diffraction pattern for the CrN coating (#C080328-1) is shown in Figure

4.6 and was identified as a B1 NaCl structure using the CrN JCPDF card number 01-11-

0065.51 There appears to be a shift to lower than expected 2θ values which may indicate a larger lattice parameter than the reference pattern. None of the suspected Ti peaks which were observed in the TiN coatings were observed in the CrN coating since it was deposited with a Cr bond layer further supporting diffraction was observed from the Ti- bond layer for the TiN coatings. Additionally, pure Cr peaks from the bond layer were not observed possibly due to the 11.9±2 μm thickness of this coating which should have

75 limited the X-ray penetration to the substrate. Correspondingly, this CrN coating was at least 2 μm thicker than any of the TiN coatings.

Figure 4.6. θ/2θ XRD pattern of CrN coating C080328-1 deposited with an evaporator current set point of 45 A and a -200 V substrate bias. CrN planes are indexed.

4.1.3 XRD Crystallographic Structure Results for Nanolayer (Ti,Cr)N Deposited as a Function of Evaporator Current (Constant Ti Evaporator Current)

The θ/2θ patterns of the nanolayered (Ti,Cr)N coatings deposited as a function of evaporator current are shown in Figure 4.7 and provided significant information on the overall crystal structure of the coatings. All four of the XRD patterns were identified as the B1 NaCl structure. Additionally, there were small intensity peaks believed to be from the substrate and Ti bond layer. Since the theoretical lattice parameters for TiN (4.242 Å) and CrN (4.140 Å) are very similar the 2θ peak locations for the (111) peaks are only separated by 0.875° which makes distinguishing overlapping TiN and CrN peaks very difficult.50,51 Contributing to the difficultly in determining whether both TiN and CrN

phases are present or whether the coating is the ternary Ti25Cr25N50 phase is that the overall peak breadth for coatings such as these cathodic arc coatings is very wide due to the small grain size of the coatings.63

At higher 2θ values the TiN and CrN peaks begin to distinguish and correspondingly there appears to be a large shift to higher 2θ values (towards the CrN reference position) in the location of the (222) peak as the Cr evaporator current is increased. This strongly suggests increasing CrN phase volume with increasing Cr evaporator current which would indicate that the coatings are nanolayered structure comprised of alternating TiN and CrN rich layers.

In order to further examine the structure of the coatings, a second set of θ/2θ scans was conducted on these coatings. This scan was conducted specifically over the

(422) peaks since the separation of the (422) peak intensity from the TiN JCPDF compared to the CrN JCPDF is 5.47°.50,51 Since the coatings all appeared to be crystallographically oriented with the (111) plane parallel to the substrate, the coated surface of the coatings were tilted 19.47° relative to the plane of the X-ray source and detector. This tilt angle brought the (422) planes more into alignment with the detector in order to increase the intensity. The tilt angle was calculated using:

  'ww'vv'uu )cos(  2 2  2 2 2  )'w()'v()'u()w()v()u( 2 where δ is the angle between the [uvw] direction ([111]) and the [u’v’w’] directions

([422]).64

The results for the (422) plane θ/2θ pattern are shown in Figure 4.8. There is a clear trend showing that the maximum peak intensity shifts towards higher 2θ values with increasing Cr evaporator current. This is suggestive of a nanolayered coating structure in which the relative volume of the CrN rich nanolayers increases with respect to TiN as the

Cr evaporator current is increased. This nanolayered coating structure was subsequently confirmed with STEM (Section 4.9). For reference, the (422) peak of TiCrN2 (JCPDF

03-065-9002, B1 NaCl Structure) which has a calculated lattice parameter of 4.184 Å is shown in Figure 4.8.75

Figure 4.8. θ/2θ pattern of the (422) plane of nanolayer (Ti,Cr)N coatings deposited as a function of Cr evaporator current (constant Ti evaporator current and substrate bias).

4.1.4 XRD Crystallographic Structure Results for Nanolayer (Ti,Cr)N Deposited as a Function of Substrate Bias (Ti at 65 A and Cr at 45 A Evaporator Current)

All of the nanolayer (Ti,Cr)N coatings deposited at different substrate biases were identified as the B1 NaCl structure (Figure 4.9). As with the previously discussed nanolayer (Ti,Cr)N coatings there were small intensity peaks observed which are believed to be from the substrate and Ti bond layer. The nanolayer (Ti,Cr)N coatings appear to have significant preferred crystallographic orientations which will be discussed in section 4.3.

Additionally, the relative intensity of the nanolayer (Ti,Cr)N coatings deposited with -200 V and -300 V substrate bias appear to decrease significantly (Figure 4.9) due to either a change in crystallographic orientation or increased number of crystalline defects

79 caused by increasing ion bombardment with increasing substrate bias. Also, the increased bombardment intensity from the high substrate bias could be promoting solid solubility (interstitials, substitution, or a mixed combination) or preferential resputtering.

This argument is further supported by the wide solid solubility range shown in the ternary

Cr-Ti-N phase diagrams for Ti25Cr25N50 which suggest that the TiCrN2 phase can form defects in order to maintain stability over a wide compositional range (Figure 2.15). This suggests that it is possible for a high substrate bias to both increase the lattice defect concentration in these coatings while also promoting intermixing and diffusion between the TiN rich and CrN rich nanolayers due to increased adatom mobility. However, increased solid solubility is difficult to prove with out detailed analysis such as TEM which was unavailable for this set of coatings and is suggested for future research.

Figure 4.9. θ/2θ XRD patterns of nanolayer (Ti,Cr)N coatings deposited with various substrate biases (constant evaporator current). B1 NaCl TiN and CrN planes are indexed.

As with the coatings deposited at different evaporator currents, it was difficult to determine whether these coatings were the ternary TiCrN2 or a nanolayered structure of

TiN rich and CrN rich layers. A tilted (19.47°) θ/2θ XRD scan of the (422) plane is shown in Figure 4.10. The peaks are non-Gaussian in shape with the maximum intensity closer to the CrN reference peak position rather than the TiN peak position suggesting a nanolayered structure. Comparing Figure 4.8 to 4.10 suggests a larger degree of defect generation resulting from increasing substrate bias as the diffracted peaks are broad and more difficult to separate for those samples with increasing substrate biases.

Figure 4.10. θ/2θ pattern of the (422) plane of nanolayer (Ti,Cr)N coatings deposited at different substrate biases (Ti evaporator of 65 A, Cr evaporator of 45 A).

4.1.5 XRD Crystallographic Structure Results for Multilayer (Ti,Cr)N Coatings with Ti and Nb Interlayers

Similar to Figures 4.7 to 4.9 all of the multilayer (Ti,Cr)N coatings (constant substrate bias and evaporator current) were observed to have the B1 NaCl structure, and

81 the XRD patterns for the coatings with Ti and Nb compliant interlayers are shown in

Figures 4.11 and 4.12, respectively. For the Nb interlayer coatings, the (110) and (220) plane Nb peaks were easily observable (Figure 4.12). The JCPDF card for bcc Nb (01-

35-0789) lists the (110) plane as the most intense.76 The Nb (220) plane is also visible while other planes have negligible intensities suggesting that the Nb is oriented with the

[110] direction perpendicular to the substrate. α-Ti peaks were visible in the coatings with Ti interlayers, but due to peak overlaps with TiN and CrN very little information on the crystallographic orientation of the Ti interlayers could be determined.

Figure 4.11. θ/2θ XRD patterns of multilayer (Ti,Cr)N coatings deposited with Ti interlayers (constant evaporator current and substrate bias). B1 NaCl TiN and CrN planes are indexed.

4.2 Crystallite Size

As discussed in Section 2.2.3, the coating grain structure can be predicted using the structure zone models. These models are based on how deposition parameters such as substrate temperature and pressure influence the mobility of atoms in the growing coating. For these cathodic arc coatings, the deposition temperature was low, but there was a high degree of ion bombardment during the deposition. For example, sputtered films and coatings deposited at low temperatures tend to only have significant diffusion on the surface of the growing film/coating rather than through the bulk of the film/coating which typically results in a fibrous columnar structure referred to as a Zone T structure.38

Ion bombardment contributes significant energy to the adatoms, and densification is typical for ion bombardment processes. Since cathodic arc typically has a greater degree

83 of bombardment than sputtering, TiN coatings deposited by this method have been observed to have dense columnar grain structures.3,38

It should be noted that there are several sources of possible error that exist for

XRD crystallite measurements for these coatings. One of the largest errors is that peak broadening also occurs due to both macro and micro strains in the coating. In some cases, the contributions to crystallite size and strains can be distinguished as they can be

1 correlated to and tan θ relationships, respectively, but due to the peak overlaps cos  between the CrN and TiN crystal structures this was extremely difficult.66 Therefore, the crystallite sizes presented in this section are most likely over estimated. Additionally, the strong preferential crystallographic orientations of the coatings may be a source of error.

4.2.1 Trends in XRD Crystallite Size of Monolithic TiN Coatings Deposited as a Function of Evaporator Current (Constant -150 V Substrate Bias)

Figure 4.13 shows that the crystallite size is increasing with increasing evaporator current for the monolithic TiN coatings as a function of Ti evaporator current (-150 V substrate bias). As discussed in chapter 2, the evaporator current in the cathodic arc system not only influences the deposition rate, but also influences the thermal energy that the adatoms are adding to the growing coating. From this trend, it appears that additional energy is provided to the adatoms from the increased evaporator current, and this energy promotes diffusion which results in a larger crystallite size.

60 (111) Plane

20 XRD Crystallite Size (nm) Size XRD Crystallite

0 40 45 50 55 60 65 70 75 80 85 90 Evaporator Current (A)

Figure 4.13. Crystallite size of TiN increases with increasing evaporator current (-150 V substrate bias). Error estimated to be at least ±10% (error bars).

4.2.2 Trends in XRD Crystallite Size of Monolithic TiN Coatings Deposited as a Function of Substrate Bias

The crystallite size of monolithic TiN coating as a function of substrate bias is shown in Figure 4.14, and there was a dramatic increase in the crystallite size for the coating deposited at -150 V compared to the coating deposited at -25 V. This increase can be qualitatively explained in terms of surface energy and diffusion. The increased substrate bias results in an increase in the acceleration of the ions before they collide with the substrate. This results in more kinetic and thermal energy of the surface atoms and therefore greater diffusion and larger crystallite size. The coating crystallite size also increased between -150 V and -300 V, but the smaller increase is most likely due to increased competition between nucleation and grain growth. High energy bombardment

85 may increase resputtering of the growing coating surface causing localized defects which may be kinetically limiting the diffusion rate and serving as sites for nucleation.

70 (111) Plane

XRD Crystallite Size (nm) Size XRD Crystallite 20

0 0 25 50 75 100 125 150 175 200 225 250 275 300 325 Substrate Bias (-V) Figure 4.14. Crystallite size of TiN increases with increasing substrate bias. Error estimated to be at least ±10% (error bars).

4.2.3 Trends in XRD Crystallite Size of Nanolayer (Ti,Cr)N Coatings Deposited as a Function Substrate Bias (Ti Evaporator of 65 A, Cr Evaporator of 45 A)

The trend for the nanolayer (Ti,Cr)N coatings deposited as a function of substrate bias (Ti evaporator of 65 A, Cr evaporator of 45 A) is attributed primarily to the effects of ion bombardment (Figure 4.15). As with the TiN coatings, the increase in the substrate bias up to -150 V promotes diffusion and therefore resulting in a larger crystallite size.

As discussed in section 4.2.2, the monolithic TiN coatings showed slightly increasing crystallite size with increasing substrate bias between -150 V and -300 V which suggests lower activation energy barrier for grain growth. However, the nanolayer

(Ti,Cr)N coatings showed decreasing crystallite size between -150 V and -300 V substrate bias suggesting the nanolayered structure affects the relationship between crystallite size and ion bombardment and is attributed to different Ti and Cr diffusion rates. As discussed in Section 4.1.4, the higher substrate biases may be promoting intermixing between the TiN rich nanolayers and CrN rich nanolayers as well as possibly increasing the concentration of defects such as nitrogen interstitials. These defects could serve as nucleation sites, and since this is a nanolayered coating there is large volume of interfaces, which could limit the growth of these nuclei. The monolithic TiN coating does not have the corresponding interfacial boundaries, so there is most likely a lower energy barrier for grain growth in the monolithic TiN coatings.

30 (111) Plane

10 XRD Crystallite Size (nm) (nm) Size XRD Crystallite

0 0 25 50 75 100 125 150 175 200 225 250 275 300 325 Substrate Bias (-V)

Figure 4.15. Crystallite size of nanolayer (Ti,Cr)N as a function of substrate bias increases then decreases. Error estimated to be at least ±10% (error bars).

4.3 Preferred Crystallographic Orientation

For this study, the crystallographic orientation of a thin film or coating is defined as the crystallographic plane parallel to the surface of the substrate or subsequently the direction vector of that plane which would be perpendicular to the substrate. It should be noted that these coatings are believed to only have a strong crystallographic orientation direction perpendicular to the substrate due to the substrate rotation and the evaporator- substrate line of sight alignment during deposition. This is not the case for all thin films/coatings in general, as certain processes such as ion beam assisted deposition

(IBAD) can cause in plane and out of plane texturing.41

In the literature, the most commonly occurring crystallographic orientation planes for cathodic arc coatings appear to be (111), (200) or (220) for TiN,44,46,77 and (200) or

(220) for CrN,42,48,53 and (200) and (111) for (Ti,Cr)N.5,6,58,77 Taking these crystallographic orientations into account, the preferred crystallographic orientation will be presented as the integrated intensity of the (111) plane over the sum of the integrated intensity of the (111), (200), and (220) planes:

I )111( I/I XN)111(   III )220()200()111( where X stands for Ti, Cr, or (Ti,Cr). This leads to coatings with a strong (111) crystallographic orientation having an I(111)/IXN of approximately unity. Monolithic CrN

(C080328-1) and nanolayer (Ti,Cr)N coating (C080522-1, deposited with a -300 V substrate bias, 65 A Ti evaporator current, and a 45 A Cr evaporator current) which both had a (220) crystallographic orientation. Another method to quantify crystallographic orientation is to perform pole figures, but that is beyond the scope of this effort.

4.3.1. Preferred Crystallographic Orientation of the Monolithic TiN Coatings Deposited as a Function of Evaporator Current (Constant Substrate Bias)

Figure 4.16 shows that TiN coatings are primarily oriented in the [111] direction for all three coatings deposited with increasing evaporator current, and that the crystallographic orientation appears to increase with increasing current. In the literature, one of the frequently referred to driving mechanisms for crystallographic orientation of

TiN is the minimization of system energy. Generally, the energies of interest for crystallographic orientation effects are surface energy and strain energy.78-84

Additionally, Zhao et al. defined a stopping energy, which relates the ease in which impinging atoms preferentially resputtered in the coating from the various growth planes.

Therefore, stopping energy is directly related to the distance ions must travel to dissipate their energy in a given crystallographic direction (ion channeling).84 Of the many arguments discussed in the literature, the minimization of total energy comprised of surface energy, strain energy, and stopping energy appears to be relevant for the TiN coatings deposited in this study.

0.988

0.986

0.984

0.982

0.98 TiN /I 0.978 (111) I 0.976

0.974

0.972

0.97

0.968 35 45 55 65 75 85 95 Ti Evaporator Current (A) Figure 4.16. Increasing (111) crystallographic orientation of TiN with increasing Ti evaporator current (constant substrate bias)

Surface energy, Shkl, for TiN was discussed by Pelleg et al. in terms of sublimation energy (6.5 x l0- l9 J/atom) such that:78

19 2 hkl  hkl  cmJz/n105.6S where nhkl is the number of broken bonds for plane hkl and z is the coordination number which is 6 for TiN. Table 4.1 shows the calculated nhkl and Shkl for TiN. Pelleg et al. results showed that the (100) plane has the lowest free surface energy for TiN.78 The values were calculated for the (100), (111), and (110) planes, and due to the symmetry of the B1 NaCl structure the surface energy trends should be the same for the (200), (222), and (220) planes. For this study, crystallographic orientation results will be discussed in

90 regards to (200) and (220) planes rather than the (100) and (110) since the JCPDF file is for (200) and (220) planes.

Table 4.1. Surface energy of B1 NaCl structured TiN coatings calculated by Pelleg et 78 al. Number of Broken Bonds, Surface Energy (J Crystallographic Plane -2 -2 nhkl (cm ) cm ) 100 4/a2=2.2x1015 2.3 x 10-4 111 4x31/2/a2=3.8x1015 4.0 x 10-4 110 3x21/2/a2=2.4x1015 2.6 x 10-4

Strain energy relates the elastic constant (Young’s modulus, which is directionally dependent due to anisotropy) to strain:

2 U hkl ε hkl  υ)(1E where ε is the strain in the plane, and Ehkl is the elastic modulus of the plane, and υ is

Poisson’s ratio. For the B1 NaCl structures, the ratios of the elastic moduli (E100:

E111:E110) are calculated as 1.00:0.66:0.75, so in the coating the (111) plane will most likely have the lowest associated strain energies.78 In other terms, when the [111] direction is aligned perpendicular to the stress directions (parallel to the substrate) the

TiN coating system will have the lowest strain energy since the most compliant plane is aligned in the direction of the greatest stress.81 However, in the TiN system variation can occur due to the large compositional range allowing for a high number of defects in the coating.85

The stopping energy associated with these coatings was defined by Zhoa et al. to be directly proportional to the distance ions must transfer energy in growing coatings.

The (220) plane has the lowest stopping energy because ions must travel the greatest distance for this plane to distribute energy. In other terms, the [110] direction has more open space between atoms for impinging species or ions to dissipate energy into the

91 coating without resputtering from the coating. This is essentially related to the density of atoms per plane or the planar packing density which is show in Table 4.2.64,84,86

Table 4.2. Density of Atoms per Plane for B1 NaCl structure TiN (a=4.242 Å). Crystallographic Plane Atoms per Plane Density of Atoms per Plane (Atoms/Å2) (100) 3.0 0.167 (111) 3.0 0.193 (close pack (cp) plane) (110) 3.0 0.117

There are various theories on how surface energy, strain energy, and stopping energy influence the crystallographic orientation of the coating. It has been shown that varying the deposition process and parameters can lead to a dominant mechanism and therefore a dominant crystallographic orientation direction. Also, the growth of the coating itself can lead to a change in the dominant energy mechanism. For example, TiN coatings deposited in systems such as cathodic arc generally grow with a (200) crystallographic orientation which is believed to be driven by the reduction in surface energy. The coating can grow to a certain thickness at which point a critical stress is reached, and the crystallographic orientation changes from (200) to (111) in order to reduce strain energy. Additionally, coating residual stress is dependent on deposition conditions, so excessive ion bombardment causing increased residual stress may induce a crystallographic orientation change from (200) to (111).78-84

Stopping energy becomes increasing important in coatings when there is excessive ion bombardment which causes resputtering of the growing coatings. This generally occurs in processes such as cathodic arc when the substrate bias is high. In this mechanism, the overall energy of the system is reduced by preferential resputtering of the

(200) and (111) planes resulting in a larger fraction of (220) planes.78,83,84

If we compare the trend in Figure 4.16 to the energy reduction theory, it appears that a reduction in strain energy is driving the crystallographic orientation of these coatings. By comparing the crystallographic orientation to the residual stress results

(Section 4.4.1), it appears that increasing (111) crystallographic orientation does correspond to decreasing coating residual stress, strongly suggesting reduction in strain energy as the dominant mechanism.

4.3.2 Preferred Crystallographic Orientation of the Monolithic TiN Coatings Deposited as a Function of Substrate bias (Constant Ti Evaporator Current)

Figure 4.17 shows the (111) preferred crystallographic orientation of the monolithic TiN coatings deposited as function of increasing substrate bias. It should be noted that the integrated relative intensities of the (111) planes compared to the other peaks is significant higher than the JCPDF for all three of these TiN coatings (Table 4.3) confirming (111) preferred crystallographic orientation. The coating with the -150 V substrate bias has the greatest (111) preferred crystallographic orientation, and the coating with the -25 V substrate bias has the weakest (111) crystallographic orientation.

The crystallographic orientation of the coating deposited with the -300 V substrate bias is only slightly lower than the coating with the -150 V substrate bias, so it is difficult to determine whether the mechanism controlling the crystallographic orientation of the -150

V coating is different from the mechanism controlling the -300 V coating. The increasing (111) crystallographic orientation with increasing substrate bias (strain relief mechanism) does correlate with a reduction in the XRD residual stress (Section 4.4.2).

0.98

0.96

0.94

0.92 TiN /I (111) I 0.9

0.88

0.86

0.84

0.82 0 50 100 150 200 250 300 350 Bias Voltage (-V)

Figure 4.17. [111] preferred crystallographic orientation of the TiN as a function of substrate bias (constant Ti evaporator current).

Table 4.3. Relative integrated intensities of the TiN coatings as compared to JCPDF 01- 38-1420.50 Relative Integrated Intensity Crystallographic -25 V -150 V -300 V JCPDF 01-38-1420 Plane Substrate Substrate Substrate Relative Intensities bias bias bias (111) 100 100 100 72 (200) 13.4 1.3 1.4 100 (220) 5.9 0.4 1.0 45

4.3.3 Crystallographic Orientation of the CrN Coating

Table 4.4 shows the integrated intensities of the (111), (200) and (220) planes for the CrN coating compared to the intensities in the JCPDF reference pattern.51 It is evident that the coatings have a [110] crystallographic orientation. Additionally, the CrN

94 reference pattern has a more intense (220) peak intensity (relative intensity of 80%) than the TiN reference pattern (relative intensity of 45%).50,51 From the literature, CrN has been observed to have either (200) or (220) crystallographic orientations when deposited with PVD techniques.42,48,53 The differences in preferred crystallographic orientations between B1 NaCl structures of TiN and CrN appears to be linked to the differences in interactions of Cr and Ti atoms in their respective crystal lattices.

Table 4.4. Relative integrated intensities of the CrN coating as compared to JCPDF 01- 11-0065.51 JCPDF 01-11-0065 Crystallographic Plane Relative Integrated Intensity Relative Intensities (111) 31 80 (200) 33 100 (220) 100 80

A key difference between TiN and CrN is that TiN has a higher melting point than pure Ti (3290 °C versus 1670 °C)55 while pure Cr has a higher melting point (1850

°C ) than CrN since CrN decomposes in air to Cr2N between 900 and 1200 °C and Cr2N melts at 1740 °C.56 The lower melting point compared to TiN and the decomposition of

CrN into Cr2N indicates a possibly weaker bonding strength of CrN than TiN in their respective B1 NaCl structures. The melting point of TiN is more indicative a ceramic with strong covalent bonding (directional electron sharing) while the CrN melting point is more aligned with weaker metallic (electron mobility) or ionic bonding (electron transfer). It should be noted that the electronegativity difference between Ti and N is 1.5 while the difference for Cr and N is 1.4, and the overall bonding nature of these compounds is believed to be a mixture of metallic, covalent and ionic bonding.85,87

In the literature, TiN and CrN are sometimes referred to as interstitial nitrides due to the relationship between the host metal and the nitrogen atoms.85 The interstitial model considers TiN and CrN B1 NaCl structures as face centered cubic (fcc) Ti and Cr structures with N atoms located in the interstitials between Ti and Cr lattice sites, respectively.

Since bcc Cr (density of 7.19 g/cm3)39 is not a close packed structure (fcc or hcp), the metal crystal structure must undergo a structural “switching” or change to fcc in order to accommodate the N interstitials (octahedral sites).85 Figure 4.18 illustrates that a (200) oriented bcc Cr lattice matches very closely with a (220) NaCl structure of CrN. It seems possible that the structural change of bcc Cr to fcc CrN with nitrogen interstitials is responsible for the (220) crystallographic orientation as this would require only a very small displacement of Cr atoms (theoretical lattice mismatch of less than 1%). Since α-Ti is an hcp metal lattice (density of 4.51 g/cm3) there is no structural change required to accommodate nitrogen interstitials, and Figure 4.18 illustrates that (001) plane of Ti matches well with the (111) plane (close pack plane) of B1 NaCl TiN (theoretical lattice mismatch of approximately 1.7%).

The structural change of bcc Cr to the interstitial fcc or B1 NaCl structure results in a net 14% decrease in density compared to bcc Cr (CrN density of 6.18 g/cm3).51 This indicates a decrease in bond density from CrN compared to Cr and further explains the decreased melting temperature and chemical stability.85 Similarly, there is a net 18.7% increase in the density of fcc or B1 NaCl TiN (5.39 g/cm3) as compared to α-Ti which further supports increased bond density and subsequently increased melting temperature and chemical stability of TiN as compared to CrN.

Figure 4.18. In the interstitial argument for nitride structures, TiN is based on hcp α-Ti while CrN is based on a structural “switching” or change of the bcc Cr lattice to an fcc lattice.85

4.3.4 Preferred Crystallographic Orientation of Nanolayered (Ti, Cr)N Coatings Deposited as a Function of Cr Evaporator Current (65 A Ti Evaporator, Constant Substrate Bias)

The crystallographic orientation of the (Ti, Cr)N coatings deposited as a function of evaporator current are dominated by the (111) plane intensity, but that dominance does decrease with increasing Cr content (Figure 4.19). Since these coatings are nanolayered structures consisting primarily of TiN rich layers and CrN rich layers, the increasing Cr evaporator current leads to an increase in the thickness of the CrN rich layers relative to the TiN rich layers. This is confirmed by STEM (Section 4.9). As a result, the CrN

97 crystallographic orientation drives the observed trend for overall crystallographic orientation of these (Ti,Cr)N coatings as shown in Figure 4.19. From this observation and the previous discussions regarding the monolithic TiN and CrN crystallographic orientation, it becomes clear that there are different driving mechanisms for the crystallographic orientations of CrN and TiN deposited under these deposition conditions, and therefore combining the two systems leads to a competition for the overall preferred crystallographic orientation of the coatings and most likely residual stresses at the TiN/CrN layer interfaces. Additionally, this suggests that the overall crystallographic orientations of these coatings can be tailored by altering the layer structure of these coatings and the deposition conditions.

0.98

0.96

0.94 (Ti,Cr)N /I (111)

I 0.92

0.9

0.88

0.86 15 25 35 45 55 65 75 85 95 Cr Evaporator Current (A)

Figure 4.19. Decreasing [111] crystallographic orientation with increasing Cr evaporator current of nanolayer (Ti,Cr)N coatings (65 A Ti evaporator, constant substrate bias).

4.3.5 Preferred Crystallographic Orientation of Nanolayered (Ti,Cr)N Coatings Deposited as a Function of Substrate Bias (65 A Ti Evaporator Current, 45 A Cr Evaporator Current)

The nanolayer (Ti,Cr)N coatings deposited as a function of substrate bias (Figure

4.20) show an interesting trend with regards to crystallographic orientation. In Figure

4.20, the maximum [111] crystallographic orientation is observed at a substrate bias of -

150 V. For the coatings deposited with a higher substrate bias than -150 V, the [111] crystallographic orientation decreases significantly and reaches a minimum at -300 V. It should be noted that the -200 V and -300 V coatings also had significantly smaller overall integrated peak intensities than the nanolayer (Ti,Cr)N coatings deposited at lower substrate biases for the θ/2θ patterns (Figure 4.9) which was believed to be caused by a high number of crystalline defects caused by resputtering, ion implantation and intermixing of the TiN and CrN nanolayers due to increased adatom mobility and diffusion.

The crystallographic orientation trend of these nanolayer (Ti,Cr)N coatings appears to be linked to a few different mechanisms. First, the crystallographic orientation is initially driven towards the (111) with increasing substrate bias possible due to strain energy mitigation as discussed with the monolithic TiN coatings, and reaches a maximum at a -150 V substrate bias. The (111) crystallographic orientation is believed to be strongly influenced by the TiN nanolayers since monolithic CrN coatings were observed to have very small peak intensities (less than 3000 counts/step) while the (111) crystallographic texturing of the monolithic TiN coatings caused peak counts greater than

10,000 counts/step. Therefore, these results suggest that the TiN nanolayers are crystallographically oriented in order to reduce strain. For the coatings with -200 V and -

300 V substrate bias, it is believed that the higher total energy available due to ion bombardment is causing increased intermixing of the TiN and CrN nanolayers. This argument is difficult to prove with out detailed TEM analysis; however, the previously discussed crystallite size for these coatings suggest nucleation of new grains, so the reduction of surface energy may play a larger role in the driving mechanism for these coatings (Section 4.2.4). Additionally, the coating with the -300 V substrate bias has a

(220) preferred crystallographic orientation which was not observed for any of the monolithic TiN coatings. This suggests that CrN layer is dominating the overall preferred crystallographic orientation of this high substrate bias coating or that the crystallographic orientation is a combination of the ion channeling of the TiN nanolayers and the (220) crystallographic orientation of CrN attributed to the Cr bcc lattice.

1.2

0.8

(Ti,Cr)N 0.6 /I (111) I

0.4

0.2

0 0 50 100 150 200 250 300 350 Bias Voltage (-V)

100

4.3.6 Crystallographic Orientation of the (Ti,Cr)N Multilayer Coatings with Ti Interlayers (-150 V Substrate Bias, Ti Evaporator Current of 65 A, Cr Evaporator Current of 45 A)

The (Ti,Cr)N coatings with Ti interlayers all showed a significant [111] preferred crystallographic orientation, and the degree of crystallographic orientation increased with increasing number of interlayers (Figure 4.21). As discussed with the (Ti,Cr)N coatings, the overall preferred crystallographic orientation of the (Ti,Cr)N multilayer coatings is a composite of the crystallographic orientation of the TiN rich nanolayers and the CrN rich nanolayers. Since the (111) preferred crystallographic orientation is associated with strain relief in the TiN systems, the trend shown if Figure 4.22 is most likely reflecting strain relief in the TiN nanolayers. A related reason for the increasing (111) crystallographic orientation is that the lattice mismatch of the (100) α-Ti plane and the

(111) TiN plane is very small (Section 4.3.3, Figure 4.18), so the increasing number of Ti layers could be creating a lower energy barrier that promotes the (111) crystallographic orientation of TiN. This should also correspond to a further deduction in strain energy as epitaxy and lattice mismatch can play a significant role in coating residual stress .38

101

0.992

0.99

0.988

(Ti,Cr)N (Ti,Cr)N 0.986 /I (111) I

0.984

0.982

0.98 12345678910111213141516 Number of Nitride Layers

Figure 4.21. Increasing (111) preferred crystallographic orientation with increasing number of layers of multilayer (Ti,CrN) coatings with Ti interlayers (-150 V substrate bias, Ti evaporator current of 65 A, Cr evaporator current of 45 A).

4.4 Residual Stress Analysis

For thin films and coatings, residual stress is generally discussed in terms of extrinsic, σext, intrinsic stress (σint), and epitactic stress (σepitactic). Generally, extrinsic stress arises from thermal expansion mismatch between the coating and substrate as well as the mismatch between different coating layers in a multilayer coating system. Intrinsic residual stress is highly influenced by the deposition process, the material phase, impurities, grain boundaries, lattice structure, crystallographic texture, etc., and in general corresponds to any effect which creates a volumetric change within the coating. Epitactic stress arises from the lattice parameter mismatch between the coating and the substrate as well as between the various coating layers in a multilayer coating system and is

102 sometimes categorized as a type of intrinsic stress but will be treated separately in this thesis.88 The combination of extrinsic, intrinsic, and epitactic stress is generally referred to as the internal or residual stress of the coating (σT):

σT = σext + σint + σepitactic

It should be noted that analysis techniques such as sin2Ψ XRD techniques measure the residual strain of the coating, and the residual stress is calculated.38,88

For these TiN, CrN, and nanolayer (Ti,Cr)N coatings, the σext is relatively small since the deposition temperature is less than 500 °C, so trends in the residual stress states are primarily attributed to intrinsic and epitatic residual stress. There are a few possible sources for the creation of intrinsic residual stress in these coatings which must be discussed. First, the high level ionization in the cathodic arc process leads to significant ion bombardment of the growing coating. The impinging atoms can displace atoms in the coating from energetically favorable sites which results in a volume change (strain) and therefore residual stress. This process is very relevant for these cathodic arc coatings since the low deposition temperature limits the degree of bulk diffusion in the coatings, so displaced atoms may be stuck in less energetically favorable sites such as interstitials.

The epitactic stress that arises between the nanolayers of TiN and CrN could play an important role in the residual stress of the (Ti,Cr)N coating. Additionally, epitaxy could play a role in the multilayer (Ti,Cr)N coatings with Ti and Nb interlayers which appears to be the case for the coatings with Ti interlayers (Section 4.3.6). Residual stress arising from epitaxy is defined by the difference in the lattice parameters:

 aa f  21 a 2

103

where f is the lattice misfit, a1 is the lattice parameter of the coating, and a2 is the lattice parameter of the substrate. For misfit between coating layers the appropriate lattice parameters can be substituted into the equation.38,88

Epitaxial coating growth generally occurs with f < 0.1. A large misfit between lattice parameters can result in defects such as edge dislocations at the boundary or an incoherent interface where atomic alignment does not occur across the interface. Since these coatings are deposited by the cathodic arc process, the interfaces should be graded rather than abrupt due to implantation and subsequent creation of the intermixed region as discussed in Section 2.2.3. Therefore, a perfect epitaxial growth is unlikely for these coatings, but interfacial stress is believed to be a contributing factor to the overall residual stress state. Select sample sets were analyzed for residual stress using the sin2Ψ

X-ray diffraction technique and will be discussed in the following subsections.

4.4.1 Trends in XRD Residual Stress Analysis of the Monolithic TiN Coatings Deposited with Different Evaporator Currents (Constant Substrate Bias)

As discussed in Section 3.5.1.3, the sin2Ψ XRD residual stress analysis requires accurate determination of Young’s modulus, Poisson’s ratio, and accurate lattice parameters of the diffraction planes being used for the measurements. Typically,

Young’s modulus can be determined for coatings such as these by using the loading curve of a nano-indenter. This method was unavailable for this project, so a calculated value of 514.2 GPa and 0.209 for E and ν, respectively, determined by Chen et al. was used for TiN.89

The general trend shown in Figure 4.22 suggests compressive residual stress decreases with increasing evaporator current under these deposition conditions. The

104 results are shown as a general constant, k, times GPa where k is represents the difference between the E111 and E422 values (anisotropy) as well as the deviation caused be using calculated E and ν rather than measured constants. As discussed in Section 4.3.1, the mechanism responsible for a (111) preferred crystallographic orientation in TiN coatings is believed to be strain relief, and correspondingly, the coating with the strongest (111) crystallographic orientation also has the lowest residual stress (85 A). Additionally, the

(111) crystallographic orientation could result in a reduction in epitaxial stress between the Ti bond layer and the TiN coating since there is very little lattice mismatch between

(001) Ti and (111) crystallographically oriented TiN. However, this alone cannot explain the different magnitude of change in residual stress. In order to confirm epitaxial stress reduction, further analysis of the Ti bond layer crystallographic orientation as well as the coating-substrate interface would be necessary which is beyond the scope of this study.

As discussed with regards to crystallite size in Section 4.3, increasing the evaporator current appears to add additional thermal energy to the growing coating which is believed to result in more coating lattice defects being annealed out and thus a further reduction in the magnitude of the compressive residual stress.

105

12 (422) Plane 11 (111) Plane 10 9 8 7 6 5 4 Residual Stress (-GPa x k) 3 2 1 0 40 45 50 55 60 65 70 75 80 85 90 Evaporator Current (A)

4.4.2 Trends in XRD Residual Stress Analysis of the Monolithic TiN Coatings Deposited with Different Substrate Biases (Constant Ti Evaporator Current)

Figure 4.23 shows the residual stress in TiN coatings as a function of substrate bias, in which the compressive residual stress decreases with increasing substrate bias.

The reduction in compressive residual stress follows the trend of increasing (111) preferred crystallographic orientation. This reduction in compressive residual stress is attributed primarily to the affects of ion bombardment causing ion implantation and increased thermal energy. The coating deposited with a -25 V substrate bias has a small crystallite size indicating a low rate of adatom diffusion on the surface of the growing coating (Section 4.2.2). Since diffusion appears to be limited in this coating, the high

106 level of compressive residual stress suggests that there may be a high density of lattice defects which are kinetically limited from being annealed under this set of deposition conditions.

In general, low-level ion bombardment, such as found in sputter deposition processes, results in compressive stresses. However, high degrees of ion bombardment such as in the cathodic arc process can actually result in increased tensile stresses if large enough volume increases are observed due to ion implantation, resputtering, and/or thermal relaxation. Figure 4.23 suggests that the high level of ion bombardment creates tensile microstresses due to the combined affects of ion implantation and resputtering of the growing coating. This is supported by the reduction in deposition rate for the coating deposited at -300 V which was attributed to increased resputtering rate of the coating

(Section 4.5.2). In addition, an increase in tensile microstresses would account for the relatively poor erosion performance of the coating deposited with the -300 V substrate bias (Section 4.13), but a more in depth residual stress study would be required to confirm that overall compressive residual state of these coatings is decreasing due to increasing tensile stresses.11,38,88,90

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5 (422) Plane

4.5 (111) Plane

3.5

2.5

Residual Stress (-GPa x k) Residual Stress 1.5

0.5

0 0 25 50 75 100 125 150 175 200 225 250 275 300 325 Substrate Bias Voltage (-V)

4.5 Deposition Rate of Nitride Coatings

This section discusses the effects of the processing parameters on the deposition rates of the coatings. In particular, the effects of resputtering of the deposited coatings and cathode target evaporation rate were examined as a function of substrate bias and evaporator current, respectively.

Both ESEM and optical microscopy (OM) were used to determine the coating thicknesses and deposition rates. It should be noted that there is some error associated with the optical thickness measurements do to edge retention effects. Edge retention refers to the rounding of the coating edge during grinding and polishing of the coating cross section. This rounding could lead to under-measuring of the coating thickness.

108

Edge retention is sometimes exacerbated by poor wetting or adhesion between the coating top surface and the epoxy which leads to a spacing between the coating surface and the epoxy after setting due to shrinkage. Figure 4.24 shows some of the edge retention and possible wetting issues associated with the cold mounting (epoxy) of these coatings. To prevent this from occurring, coatings are often Cu or Ni plated to preserve the integrity of the coating, but this process was unavailable for the coatings in this study.

In addition, fracture surfaces are often used to determine coating thickness but due to the ductile substrates, fracture surfaces were difficult to produce for all of the coated samples.

Figure 4.24. Polished cross section of a multilayer (Ti,Cr)N coating showing poor edge retention and poor coating-epoxy wetting. Optical micrograph at 1000X magnification using differential interference contrast (DIC).

4.5.1 Deposition Rate of Monolithic TiN Coatings Deposited at Different Evaporator Currents (-150 V Substrate Bias)

The evaporator current plays an essential role in the deposition rate of the monolithic TiN. Figure 4.25 shows that the deposition rate increases linearly with increasing evaporator current for monolithic TiN. There is little literature available on

109 the dependence of the evaporator current on the deposition rate for nitride coatings deposited by cathodic arc. From the literature, Straumal et al. deposited metallic Ti coatings at evaporator currents ranging from 110 A to 220 A and showed a linear dependence of deposition rate with evaporator current for the substrate directly aligned with evaporated target (line of sight).91 If the evaporation of metallic Ti does have a linear relationship to the cathode target current in this systems, then it seems likely that there was enough N2 in the chamber to react with the Ti at the surface of the growing TiN coating in order to produce stoichiometric TiN for all of the coatings deposited as a function of evaporator current. This is interpreted from the observed linear dependence of TiN deposition rate with Ti evaporator current.

4 y = 0.0422x 3.5 R2 = 0.9919

m/hr) 3 μ 2.5 2 1.5 1 Deposition Rate ( Rate Deposition 0.5 0 40 45 50 55 60 65 70 75 80 85 90 Evaporator Current (A)

Figure 4.25. A linear relationship between deposition rate and Ti evaporator current is evident for the monolithic TiN coatings (-150 V substrate bias).

There are several other factors which influence the deposition rate of these coatings including, generation of macroparticles, substrate rotation, substrate temperature, structure of the growing coating, source to substrate distance, deposition pressure, etc., but these will not be discussed in detail as it is beyond the scope of this

110 study. The evaporation rate at a given evaporator current is complex for most materials

(including Ti and Cr) and is dependent among other things on the surface of the target

(oxide and nitride layers, cross-containments from multiple evaporators, etc), the sustained lifetime of the arc, geometry, magnetic fields, and the deposition conditions such as the pressure.3,11,38

4.5.2 Deposition Rate of TiN Coatings Deposited at Different Substrate Biases (65 A Evaporator Current)

The substrate bias influences the acceleration of the ions from the cathode to the substrate. Therefore, it is well documented that increasing the substrate bias in PVD systems such as cathodic arc results in increasing surface mobility of adatoms, resputtering of adatoms, and changes coating residual stress from ion implantation and subsurface defects. These effects make the substrate bias one of the most influential variables on coating microstructure and performance in PVD coatings.3,11,38,90

Figure 4.26 shows the deposition rate versus the substrate bias for the monolithic

TiN coatings and that there is a reduction in the deposition rate from -150 V to -300 V.

In Section 4.3.2, a decrease in the (111) preferred crystallographic orientation from -150

V to -300 V was attributed to increased strain energy due to ion implantation and possibly preferential resputtering of the (111) and (200) planes. From Figure 4.26 it does appear that resputtering of the growing coating increases at high substrate biases. A more detailed study would be necessary to determine whether the (111) and (200) planes are indeed being preferentially resputtered compared to the (220) planes.

111

3 2.9 2.8 m/hr)

μ 2.7 2.6 2.5 2.4 2.3

Deposition Rate ( Rate Deposition 2.2 2.1 2 0 25 50 75 100 125 150 175 200 225 250 275 300 325 Substrate Bias Voltage (-V) Figure 4.26. Deposition rate versus substrate bias for monolithic TiN coatings deposited with a constant 65 A evaporator current. Rate decreases above -150 V substrate bias due to resputtering of the growing coating.

4.5.3 Deposition Rate of Nanolayer (Ti,Cr)N Coatings as a Function of Cr Evaporator Current (Ti Evaporator of 65 A, Substrate Bias of -150 V)

The deposition rate versus the Cr evaporator current for the nanolayer (Ti,Cr)N coatings (Ti Evaporator of 65 A, Substrate bias of -150 V) is shown in Figure 4.27 and confirms the expected trend of increasing deposition rate with increasing Cr evaporator current. In addition to the increasing evaporation rate of the Cr cathode target as a function of current, there are other factors which are influencing the deposition rate of these nanolayer (Ti,Cr)N coatings including increasing macroparticle generation which results in increased deposition rate. The macroparticle generation will be discussed in more detail in Section 4.7.

112

3.5 3.0 m/hr)

μ 2.5 2.0 1.5 1.0

Deposition Rate ( Deposition Rate 0.5 0.0 15 25 35 45 55 65 75 85 95 Cr Evaporator Current (A)

Figure 4.27. Increasing deposition rate as a function of Cr evaporator current for nanolayer (Ti,Cr)N coatings deposited with -150 V substrate bias and 65 A Ti evaporator current.

4.5.4 Deposition Rate of Nanolayer (Ti,Cr)N Coatings Deposited at Different Substrate Bias (65 A Ti Evaporator Current, 45 A Cr Evaporator Current)

Figure 4.28 shows the deposition rate of nanolayer (Ti,Cr)N coatings deposited at different substrate bias, and the observed trend of decreasing deposition rate is attributed primarily to resputtering. Compared to the monolithic TiN coatings deposited at different substrate biases (Section 4.5.2), the nanolayer (Ti,Cr)N coatings appear more susceptible to resputtering at lower substrate biases, and this increased resputtering rate is attributed to the Cr evaporator ejecting significantly more macroparticles than the Ti evaporator, and the rate these loosely adhered macroparticles are resputtered from the coating surface increases with increasing substrate bias (Section 4.7). Unfortunately, experimental trials of CrN coatings deposited with various substrate biases was not performed, which could have further supported this finding.

113

3.0

2.5 m/hr) μ 2.0

1.5

1.0

Deposition Rate ( Deposition Rate 0.5

0.0 0 50 100 150 200 250 300 350 Substrate Bias Voltage (-V)

4.6 Electron Probe Micro Analysis of the Nanolayer (Ti,Cr)N Coatings

The EPMA analysis provided quantative information on the chemical composition of the coatings. The analysis was difficult since the characteristic energy of the NK line overlaps that of the Tiβ line. Figure 4.29 shows that the atomic percent of Cr in the coatings increases (as expected) with increasing Cr evaporator current (constant Ti evaporator current of 65 A, and -150 V substrate bias). It is important to note that all of the coatings contained approximately 50 atomic percent N. Additionally some oxygen was detected in all of the coatings, but the quantity was believed to be only a few atomic percent. Figure 4.29 suggests that a Cr evaporator current of 35 A would provide a 1:1 ratio of Ti:Cr under these deposition conditions.

114

40.000

30.000 Chromium Titanium 20.000

10.000 Atomic Percent (%) Percent Atomic 0.000 15 25 35 45 55 65 75 85 95 Cr Evaporator (A)

The EPMA results for the nanolayer (Ti,Cr)N coatings deposited as a function of substrate bias are shown in Figure 4.30. It appears that preferential resputtering of Ti versus Cr is not an issue with these coatings as the composition of Ti to Cr does not appear to change over the full range of substrate biases. It should be noted that TiN can have an atomic percentage of nitrogen greater than 50% since the lattice structure and bonding nature permits vacancies on metal sites, and this could possibly explain the

EPMA results (greater than 50 atomic percent N) for the coating deposited with the -50 V substrate bias.85

40.000 Chromium 35.000 Titanium 30.000 25.000 20.000

(%) 15.000 10.000 Atomic Percent 5.000 0.000 0 50 100 150 200 250 300 Substrate Bias Voltage (-V)

115

4.7 Quantative Analysis of Large Scale Defects of Nanolayer (Ti,Cr)N Coatings

The Clemex Vision P.E. software image analysis was conducted on the ESEM top surface micrographs of nanolayer (Ti,Cr)N coatings deposited on the pre-polished 304 stainless steel substrates. The substrates had some fine scratches which were mimicked in the coating structure, but the analysis routine was constructed to distinguish these scratches from the actual macroparticle defects. The software routine can only distinguish large scale defects on the coating surface, and this does not fully portray the macroparticle incorporation within the coating cross sections as particles embedded beneath the surface may not significantly alter the appearance of the coating surface and thus may not be incorporated into the calculations. However, metallic macroparticles are generally on the order of a few microns in diameter,3 and since the coatings are less than

13 μm thick the analysis routine was believed to be fairly accurate. Additionally, the software can not distinguish the type of surface defect such as the difference between metallic macroparticles or physical coating damage, but clear trends relating to macroparticles were visible for all of these coatings as shown in Figures 4.31 and 4.33.

Figure 4.31 shows the surface area of defects as a function of Cr evaporator current for nanolayer (Ti,CrN) coatings (Ti evaporator of 65 A, substrate bias of -150 V).

The ejection of the macroparticles from the cathode target appears to increase rapidly above 45 A, and for the coating deposited with the 85 A Cr current, almost 25% of the total surface is covered by macroparticle related defects. ESEM surface morphology micrographs shown in Figure 4.32 of the (Ti,Cr)N coatings with increasing Cr evaporator current further support the trend of increasing macroparticles with increasing Cr current since the defects are primarily spherical in shape.

116

The monolithic TiN coatings deposited under various conditions did show some macroparticle defects, but not nearly as many macroparticles as the CrN or nanolayer

(Ti,Cr)N coatings. This strongly suggests differences in the cathodic arc evaporation characteristics of Ti compared to Cr. Generally, macroparticle emission has been determined to be more prevalent in lower melting point materials.3 As discussed in section 4.3.3, bcc Cr has a higher melting point than hcp α-Ti (1850 °C versus 1670

°C).55,56 However, the surface of the evaporators are considered to be nitrides during the

3 coating process due to the N2 atmosphere. Since the decomposition of CrN into Cr2N and subsequent melting point of Cr2N (1740 °C) is less than Cr while TiN is significantly higher (3290 °C) than α-Ti, the relationship between a low melting point target and high macroparticle emission seems to applicable for the Cr target.55,56 A more in depth study would be necessary to confirm this theory.

10 Area of Image (%) of Image Area

5 Surface Area of Defects / Total Surface Total Surface of Defects / Area Surface

0 15 25 35 45 55 65 75 85 95 Cr Evaporator Current (A)

Figure 4.31. Surface defects increase significantly as a function of Cr evaporator current for nanolayer (Ti,Cr)N coatings deposited with -150 V substrate bias and 65 A Ti evaporator.

117

Figure 4.32. Surface defects increase with increasing Cr evaporator current for nanolayer (Ti,Cr)N coatings deposited with -150 V substrate bias and 65 A Ti evaporator current.

Figure 4.33 plots the relative area of surface defects versus the substrate bias for nanolayer (Ti,Cr)N coatings deposited with 65 A Ti evaporator current and 45 A Cr evaporator current. The plot shows that the coating deposited with a -50 V substrate bias has a larger area of defects than the nanolayer (Ti,Cr)N coatings deposited with higher substrate biases. The reason for the reduction in macroparticle defect generation with increasing substrate bias is believed to be the result of increased resputtering of the growing coating during deposition, resulting in preferential resputtering of loosely adhered macroparticles.

118

15 Area of Image (%) Area

10 Surface Area of Defects / Total Surface 5

0 0 50 100 150 200 250 300 350 Substrate Bias Voltage (-V)

Figure 4.33. Relative surface area of large scale defects as a function of substrate biases for nanolayer (Ti,Cr)N coatings (65 A Ti evaporator current, 45 A Cr evaporator current).

4.8 ESEM Fracture Surfaces

ESEM fracture surfaces were prepared for select coatings to examine the coating microstructure. Due to the high degree of ductility of the AM355 stainless steel substrates, it was difficult to prepare clean fracture surfaces. Figures 4.34 and 4.35 show fracture surfaces of a monolithic TiN (65 A Ti evaporator current, -150 V substrate bias) and a nanolayer (Ti,Cr)N (65 A Ti evaporator current, 45 A Cr evaporator current, -150

V substrate bias) coatings, respectively. Both of the coatings appear to have a fibrous columnar microstructure (Zone T) as predicted by the structural zone models discussed in

Section 2.2.3. The monolithic TiN coating appears to have a more columnar structure than the nanolayer (Ti,Cr)N coating which appears to have finer columns. In section 4.2, the XRD crystallite size of this TiN coating was more than twice the size of this

119 nanolayer (Ti,Cr)N coating. This suggests that the nanolayering (i.e. increased interfaces) in the (Ti,Cr)N coatings may reduce grain growth by limiting diffusion.

Figure 4.34. Fracture surface of monolithic TiN coating deposited with a 65 A Ti evaporator current and a -150 V substrate bias.

Figure 4.35. Fracture surface of nanolayer (Ti,Cr)N coating (Cr evaporator of 45 A, Ti evaporator of 65 A and a -150 V substrate bias).

120

4.9 STEM of Nanolayer (Ti,Cr)N Coatings Deposited at Different Cr Evaporator Currents (65 A Ti Evaporator Current, -150 V Substrate Bias)

Annular Dark Field Scanning Transmission Electron Microscope (ADF-STEM) micrographs are shown in Figure 4.36, and all four of the nanolayer (Ti,Cr)N coatings deposited at different Cr evaporator currents (65 A Ti evaporator current, -150 V substrate bias) were determined to consist of TiN and CrN rich nanolayers with a periodicity (λ) ranging from 10.5 to 21.0 nanometers. By comparing the EPMA results to the ADF-STEM (Figure 4.37) it was determined that the TiN nanolayers were approximately 6.9 nm while the CrN nanolayer thickness varies with the Cr evaporator current as shown in Figure 4.38. Furthermore, Figure 4.38 suggests that the CrN rich nanolayer thickness is over twice the thickness of the TiN rich nanolayers for the coating deposited with the highest Cr evaporator current (85 A). It should be noted that the Cr and Ti signals within the layers approach zero in the Ti and Cr rich layers, respectively, but there is a detectable amount of Ti within the Cr rich layers and also a detectable amount of Cr within the Ti rich layers.

121

Figure 4.36. ADF-STEM micrographs of nanolayer (Ti,Cr)N coatings deposited at different Cr evaporator currents (Ti evaporator current of 65 A, -150 V substrate bias).

40.000

35.000 y = 1.63x + 2.50 2 30.000 R = 1.00 25.000 20.000

Atomic % Cr (at. %) (at. Cr % Atomic 15.000 10.0 12.0 14.0 16.0 18.0 20.0 22.0 Periodicity, λ (nm)

Figure 4.37. EPMA results for the atomic % of Cr in nanolayer (Ti,Cr)N coatings versus the periodicity, λ, as determined by ADF-STEM. (Ti evaporator current of 65 A, -150 V substrate bias).

122 r 15.0

10.0

5.0 Thickness (nm) Thickness

Estimated CrN Nanolaye CrN Estimated 0.0 15 25 35 45 55 65 75 85 95 Cr Evaporator (A)

Figure 4.38. The CrN rich nanolayer thickness increases to over twice the TiN rich nanolayer thickness (6.9 nm) with increasing Cr evaporator current. (Calculated using EPMA and STEM results for nanolayer (Ti,Cr)N coatings (Ti evaporator current of 65 A, -150 V substrate bias)).

4.10. Vicker’s Micro-Hardness

Vicker’s micro-hardness measurements were made on the surface as described in the experimental procedure (Section 3.5.6). The measurements were difficult due to the coating thickness (thin coatings). The HV 0.050 for the nanolayer (Ti,Cr)N coatings ranged from 1700 kg/mm2 to almost 2800 kg/mm2 depending on the coating thickness. In

th general, HV measurements should not exceed 1/10 of the depth of the coating to avoid influences from the substrate. In more detail, the hardness decreased with decreasing thickness, so the hardness value was believed to be a composite value of the substrate

(approximately 300HV) and the coating. Since the coatings were harder than the uncoated stainless steel substrate material, it is logical to assume that the low angle erosion resistance of a stainless steel turbine engine component would be improved by applying these coatings.

In addition to the top surface measurements, HV indentions were also made on the cross sections. A small load of 10 g was used in order to prevent the indenter from slipping off the edge of the coating. Due to the high hardness of the coatings, the

123 indentions were too small to accurately measure using OM, so the diagonal length measurements were attempted using ESEM but accurate measurements were difficult to obtain for a variety of reasons including poor focus and indent quality.

For thin films and coatings, a nanoindenter would be a more suitable method for measuring hardness. This technique also provides information of the elastic modulus.

However, nanoidention was not performed in this study due to the relatively high cost per data point compared to HV.

4.11 Scratch Adhesion Testing of Nanolayer (Ti,Cr)N Coatings

The plot of critical force versus Cr evaporator current for nanolayer (Ti,Cr)N coatings (-150 V substrate bias, 65 A Ti evaporator current) is shown as Figure 4.39 and suggests that adhesion increases significantly with increasing evaporator current with the

85 A coating having a 53% increase in critical force compared to the coating deposited with a 45 A Cr evaporator current.

There are at least two considerations that must be discussed with regards to the observed trend of increased adhesion with increasing Cr content. First, an increased volume fraction of CrN could reduce the overall residual stress of the coating since CrN is likely more compliant than TiN based on melting temperatures and bonding nature

(Section 4.3.3).

The second factor to consider is that the volume fraction of nanolayer interfaces

(TiN/CrN) will decrease with increasing Cr evaporator current since λ increases. The exact nature of the interfaces is difficult to determine since the coating process most likely creates a graded interface due to intermixing caused by the ion bombardment and substrate rotation. However, the interfaces may be highly defective and strained in order

124 to accommodate TiN and CrN, and this strain may be fairly significant since the crystallographic orientations of TiN and CrN nanolayers are different (Section 4.3). A detailed TEM and coating residual stress analysis would be needed to determine the exact nature of the nanolayer-stress-adhesion relationship and this is beyond the scope of this study.

90.00

80.00

70.00

60.00

50.00

40.00

Critical Force (N) 30.00

20.00

10.00

0.00 0 102030405060708090 Cr Evaporator Current (A) Figure 4.39. Increasing critical force with increasing Cr evaporator current for nanolayer (Ti,Cr)N coatings (-150 V substrate bias, 65 A Ti evaporator current).

4.12 Surface Roughness

In general, the monolithic TiN, monolithic CrN and nanolayer (Ti,Cr)N cathodic arc coating surface roughness appears to be primarily linked to the macroparticle concentration and the crystallite size of the coatings. Additionally, the substrate surface roughness plays a large role on the deposited coating roughness, but this was not considered to significantly influence the observed trends for Ra since all of the substrates were prepared under identical conditions. However, the thickness of the coatings could

125 influence the roughness as surface roughness often increases with increasing coating thickness, but this is also dependent on the deposition process and methodology.

4.12.1 Surface Roughness of the Monolithic TiN and CrN Coatings

Figure 4.40 shows that surface roughness of the monolithic TiN coatings decreases with increasing substrate bias (65 A Ti evaporator current). Increasing the substrate bias is believed to remove some of the macroparticles through resputtering and increase densification due to the addition of thermal energy and ion bombardment which increases the diffusion rates and therefore increases the crystallite size. Larger grains should lead to decreased roughness since there will be fewer grain boundaries.

Additionally, from the literature, Cheng et al. linked the [111] crystallographic orientation to the lowest surface roughness,44 and this was observed for these TiN samples deposited at different substrate biases. It should be noted that since the TiN coatings deposited at different evaporator currents (-150 V) had some deviation in coating thickness they are not included in this discussion.

For reference, the CrN coating had a roughness of 169 ± 14 nm which was lower than any of the TiN coatings. The structure zone model predicts that CrN will have larger grains compared to TiN at a given substrate temperature since the melting temperature/decomposition of CrN is lower than TiN. Additionally, the wet blasted

AM355 substrate had a roughness of 117.2 ± 61 nm which was lower than any of the coated coatings. In general, coatings replicate the substrate morphology if given enough time and energy to allow diffusion along the surface.

126

350

300

250

200

150

100 Surface Roughness,(nm) Ra

0 0 50 100 150 200 250 300 350

Substrate Bias Voltage (-V) Figure 4.40. Surface roughness decreases with increasing substrate bias for the monolithic TiN coatings (65 A Ti evaporator current).

4.12.2 Surface Roughness of Nanolayer (Ti,Cr)N Coatings

The surface roughness trends for the nanolayer (Ti,Cr)N coatings deposited at different evaporator currents (Figure 4.41) appears to be linked to the large scale defects discussed in Section 4.6. It appears that the high concentration of macroparticles resulted in significantly increased surface roughness for the coatings deposited at 65 A and 85 A.

It would be interesting to compare these unfiltered cathodic arc coatings to filtered cathodic arc coatings in order to eliminate the roughness caused by the macroparticles, but this was beyond the scope of this project.

127

400

350

300

250

200

150

100 Surface Roughness, Ra (nm) Roughness, Surface

0 15 25 35 45 55 65 75 85 95

Cr Evaporator Current (A)

Figure 4.41. Surface roughness increases with increasing Cr evaporator current for nanolayer (Ti,Cr)N coatings (65 A Ti evaporator current, -150 V substrate bias).

The nanolayer (Ti,Cr)N coatings deposited at different substrate biases follows the trend shown in Figure 4.42 and appears to be linked to large scale defects and crystallite size. From -50 to -150 V the coatings become significantly smoother, and this can be attributed to decreased large scale defects (macroparticles) and increasing grain size (decreasing grain boundaries). The increased surface roughness for the coatings at -

200 V and -300 V corresponds to decreasing grain size, and therefore more grain boundaries. As previously mentioned in Section 4.3, the decreased crystallite size for the

-200 V and -300 V coatings may be due to increased crystalline defects (nucleation sites) and increased number of coating interfaces compared to monolithic TiN, and therefore a decrease in the diffusion rate. However, more in-depth studies would be required to confirm the relationship between crystallite size, surface roughness, and substrate bias.

128

500

450

400

350

300

250

200

150 Surface Roughness, Ra (nm) Ra Roughness, Surface 100

0 0 50 100 150 200 250 300 350

Substrate Bias Voltage (-V)

The multilayer (Ti,Cr)N coatings with Ti and Nb interlayers have relatively constant surface roughness values (Figure 4.43). Since the (Ti,Cr)N coating is a nanolayered structure, the small volume fraction of metallic interlayers was not expected to affect the crystallite size through interruption of large column growth. Therefore, the nanolayer interfaces play a more dominant role in reducing the crystallite size compared to the Ti and Nb interlayers (Section 4.3).

The slight decrease in the roughness of the (Ti,Cr)N multilayer coatings with Nb interlayers as compared to the multilayer (Ti, Cr)N coatings with the Ti interlayers is believed to be related to the increase in substrate temperature during evaporation of Nb interlayers as compared to Ti interlayers. Since the Nb target required 120 A to sustain a

129 stable arc, the substrate temperature increased approximately 75 °C during the Nb interlayer depositions. The increased substrate temperature increased the diffusion rate of the adatoms on the coating surface as well as atoms beneath the coating surface.

Therefore, the grain size of the coating should increase and the surface roughness should correspondingly decrease.

300 Ti Interlayers Nb Interlayers

250

200

150

100 Surface Roughness, Ra (nm) Surface 50

0 012345678910111213141516 Number of (Ti,Cr)N Layers Figure 4.43. Surface roughness of multilayer (Ti,Cr)N coatings with Ti and Nb interlayers (65 A Ti evaporator current, 45 A Cr Evaporator, and -150 V substrate bias).

4.13 Erosion Performance

The erosion performance is discussed in relation to the analytical characterization techniques and a few comparable studies from the literature. The main objective was to identify how and why the structure and composition influenced the erosion performance.

Identifying these mechanisms is essential for designing production level coatings based on the Ti-Cr-N system.

130

For each coating set, erosion results are presented using two graphs for glass bead erosion and one graph for the alumina media test. The glass bead erosion is shown as mass loss versus the total mass of beads used for each erosion spot (6,350 μm or 0.25 inch diameter). This test was conducted twice for each coating. Due to the length of time required to obtain the mass loss versus total mass of glass beads for each erosion test, the second test was expedited and used to confirm the repeatability of the first test set. The 0.00111 g failure criterion is shown as a horizontal red line in the figures for the glass bead erosion.

The alumina erosion tests caused significant damage with only 25 g to all of the coatings, so the results are shown in bar graph form. The aggressive nature of the alumina particles is related to the particle morphology (angular) as well as the intrinsic particle properties such as hardness and fracture toughness. Both the glass bead and alumina erosion tests are shown in comparison to the uncoated AM355 substrate material. Prior to the erosion tests, the uncoated AM355 substrate material was wet blasted for the glass bead erosion test in order to remove the oxide layer.

4.13.1 Erosion Performance of Monolithic TiN Coatings Deposited with Different Evaporator Currents (Constant Substrate Bias)

The glass bead erosion results for the monolithic TiN coatings deposited with different evaporator currents are shown in Figures 4.44 and 4.45. The substrate out performed all three of the TiN coatings deposited at different evaporator currents as the round media impacting at 90° did not cause much damage to the relatively ductile

AM355. Initial damage of these TiN coatings occurred through a micro-chipping mechanism observed after the first 25 g of erodent. The exact diameters and distribution

131 of the initial chips is unknown, but some of the damaged areas were visible by the naked eye. The number and size of the damage spots increased with increasing total damage until the -0.00111 g failure point was reached. After 50 g of erodent was used for the 45

A and 85 A coatings, the substrate was visible, but small isolated areas of coating were still evident under optical microscopy. A schematic showing the observed micro- chipping failure mechanism is shown as Figure 4.46.

TiN: 45A Evaporator Current -0.00140 TiN: 65A Evaporator Current TiN: 85A Evaporator Current -0.00120 Wet Blasted AM355 -0.00111

-0.00100

-0.00080

-0.00060

Mass Loss of Sample (g) -0.00040

-0.00020

0.00000 0 25 50 75 100 125 150 175 200 225 250 275 300 325 350 375 400 Total Dosage of Erodent (g)

132

TiN: 45A Evaporator Current -0.00160 TiN: 65A Evaporator Current TiN: 85A Evaporator Current -0.00140 Wet Blasted AM355

-0.00120 -0.00111

-0.00100

-0.00080

-0.00060

Mass LossMass of Sample (g) -0.00040

-0.00020

0.00000 0 25 50 75 100 125 150 175 200 225 250 275 300 325 350 375 400 425 450 Total Dosage of Erodent (g)

The coating deposited with the 65 A evaporator current did perform better than the coatings deposited with an evaporator current of 45 A and 85 A. The performance of the 65 A coating appears to be linked to residual stress, and in particular the tradeoff between adhesion and residual stress. As previously discussed, compressive residual stress could increase the erosion resistance of these coatings provided the compressive forces do not exceed the shear forces responsible for coating adhesion.41

The compressive residual stress of the TiN coating deposited at 45 A appears to be too high since erosion failure was rapid. The TiN coating at 85 A appeared to have eroded with a larger micro-chip diameter than the 45 A and 65 A coatings which indicate that the crack growth in this coating was not as impeded possibly due to the lower compressive residual stress and also the larger crystallite size (possible failure along grain boundaries). The relationship between the crystallite size and micro-chip diameter could

133 be related to the fracture toughness of the coatings since the grain boundaries may impede crack propagation. Therefore, the improved performance of the coating deposited at 65 A could be a combination of optimal compressive residual stress and increased toughness due to the fine crystallite size, but further study on the correlation between micro-chipping and crystallite size is suggested.

Figure 4.46. Progression of observed micro-chipping method of erosion failure for the cathodic arc deposited coatings.

Figure 4.47 shows the mass loss of the monolithic TiN coatings eroded after 25 g of alumina media. The alumina media was an aggressive test which caused a rapid coating chipping. The AM355 eroded much faster using the alumina media as compared to the glass bead media, further suggesting the tests may have been too aggressive.

134

1 TiN: 45A Evaporator Current -0.00250 TiN: 65A Evaporator Current TiN: 85A Evaporator Current AM355 -0.00200

-0.00150

-0.00100 Mass Loss of Sample (g)

-0.00050

0.00000 Figure 4.47. Alumina erosion media results shown as mass loss for a single 25 g test conducted on the monolithic TiN coatings deposited with different evaporator currents.

4.13.2 Erosion Performance of Monolithic TiN Coatings Deposited with Different Substrate Biases (Constant Ti Evaporator Current)

The glass bead erosion results for the monolithic TiN coatings deposited at different substrate biases are shown in Figures 4.48 and Figure 4.49. The coating deposited with the -25 V substrate bias out performed all other TiN, CrN, and nanolayer

(Ti,Cr)N coatings deposited in this study. The observed erosion mechanism was very different for the -25 V coating as compared to the -150 V and -300 V coatings. The -25

V coating eroded very uniformly with only a few deep chips or craters which increased slowly in both size and quantity with increasing erodent dosage. Even after 5000 g of total erodent, final failure of the coating was never achieved; although, a significant area of the substrate was visible. The -150 V and -300 V monolithic TiN coatings eroded much faster with a similar large scale chipping mechanism as observed for all of the TiN coatings deposited at different evaporator current coatings, and uniform erosion was not

135 observed due to the rapid coalesce of these micro-chips. The -25 V coating also showed increased resistance to the aggressive alumina (Figure 4.50).

-0.0014 TiN: Substrate Bias @ -25V TiN: Substrate Bias @ -150V TiN: Substrate Bias @ -300V -0.0012 -0.00111 Wet Blasted AM355

-0.001

-0.0008

-0.0006

Mass Loss of Sample (g) -0.0004

-0.0002

0 -100 300 700 1100 1500 1900 2300 2700 3100 3500 3900 4300 4700 5100 5500 Total Dosage of Erodent (g) Figure 4.48. Glass bead erosion results shown as mass loss versus total dosage of erodent on a single 1/4 in diameter spot for the monolithic TiN coatings deposited with different substrate biases. Original test.

-0.0014 TiN: Substrate Bias @ -25V TiN: Substrate Bias @ -150V TiN: Substrate Bias @ -300V -0.0012 -0.00111 Wet Blasted AM355

-0.001

-0.0008

-0.0006

Mass Loss of Sample (g) -0.0004

-0.0002

0 0 500 1000 1500 2000 2500 3000

Total Dosage of Erodent (g)

136

TiN: -25V Substrate Bias 1 TiN: -150V Substrate Bias -0.00180 TiN: -300V Substrate Bias AM355 -0.00160

-0.00140

-0.00120

-0.00100

-0.00080

-0.00060 Mass Loss of Sample (g) (g) of Sample Loss Mass -0.00040

-0.00020

0.00000 Figure 4.50. Alumina erosion media results shown as mass loss for a single 25 g test conducted on the monolithic TiN coatings deposited with different substrate biases.

As discussed in section 4.3.2, the coating deposited with -25 V substrate bias does not have as strong a [111] direction preferred crystallographic orientation. In the literature, Sue et al. linked increasing (111) crystallographic orientation of cathodic arc deposited TiN coatings to the increasing erosion resistance.33 Sue et al. also linked the compressive residual stress of the coatings to erosion resistance with lower compressive residual stress coatings eroding slower and more uniformly than the higher compressive residual stress coatings which eroded with a chipping mechanism.33 The coating deposited with the -25 V substrate bias had a higher compressive residual stress than the coatings at higher substrate bias, so this result is contrary to the findings by Sue et al.; however, the deposition conditions for the coatings deposited by Sue et al. were not

137 provided, so it is difficult to accurately compare the structure and performance between the two studies.

Unlike the monolithic TiN deposited with increasing evaporator current, the crystallite size for the -25 V substrate bias coating was significantly smaller than the other substrate bias coatings. Due to the chipping mechanism being the overall failure mechanism, it appears that grain size may only be partially contributing to coating performance due to the large micro-chip size versus smaller crystallite diameter. Coating erosion loss may have resulted from nanoscale chipping and failure along the grain boundaries but this was not confirmed with SEM. However, if the chipping is occurring on the scale of the crystallite size, then smaller crystallites could theoretically improve erosion resistance since grain boundaries could act as barriers to crack propagation

(toughness). Further research on the crystallite size-erosion resistance is suggested for these coatings.

4.13.3 Erosion Performance of the Monolithic CrN Coating

The glass bead and alumina erosion results for the monolithic CrN coating are shown in Figures 4.51 through Figure 4.53. For the glass bead test, the CrN coating exhibited uniform erosion with a few large scale chips, and this observed erosion behavior was very similar to the erosion behavior of the monolithic TiN coating deposited with a -25 V substrate bias. Furthermore, this CrN coating had better performance than all of the TiN coatings with the exception of the TiN coating deposited with the -25 substrate bias. However, the CrN coating did not perform very well compared to the monolithic -25 V substrate bias TiN coating using the aggressive alumina test, and this could be related to an increase in metallic bonding nature of CrN as

138 compared to TiN (Section 4.3.3). Since metals tend to erode faster at glancing angles

(Section 2.2.1) than brittle materials (Section 2.2.2), the erosion performance of CrN could be indicative of greater compliancy in the CrN system as compared to the TiN system.

It is interesting to note that the CrN coating performed fairly well even though it was deposited with a high substrate bias (-200 V). This could be attributed to many factors including that the structure zone model (Section 2.2.3) predicts a greater amount of surface diffusion in CrN than TiN due to the lower melting point of CrN (1740 °C versus 3290 °C).55,56 The difference between the structures and bonding nature of TiN and CrN appears to carry over to the erosion performance of the coatings as one would expect.

-0.0014 CrN (C080328-1) -200 Substrate Bias Wet Blasted AM355 -0.0012 -0.00111

-0.001

-0.0008

-0.0006

Mass Loss of Sample (g) (g) of Sample Loss Mass -0.0004

-0.0002

0 0 200 400 600 800 1000 1200 1400 1600 1800 2000 Total Dosage of Erodent (g)

Figure 4.51. Glass bead erosion results shown as mass loss versus total dosage of erodent on a single 1/4 in diameter spot for the monolithic CrN coatings. Original test.

139

-0.0016 CrN (C080328-1) -200 Substrate Bias Wet Blasted AM355 -0.0014

-0.0012 -0.00111

-0.001

-0.0008

-0.0006

Mass Loss of Sample (g) -0.0004

-0.0002

0 0 500 1000 1500 2000 2500

Total Dosage of Erodent (g)

Figure 4.52. Glass bead erosion results shown as mass loss versus total dosage of erodent on a single 1/4 in diameter spot for the monolithic CrN coatings. Confirmation test spot.

1 -0.00300 CrN Sample #C080328-1 AM355

-0.00250

-0.00200

-0.00150

-0.00100 Mass Loss of Sample (g) of Sample Loss Mass

-0.00050

0.00000 Figure 4.53. Alumina erosion media results shown as mass loss for a single 25 g test conducted on the monolithic CrN coating.

140

4.13.4 Erosion Performance of Nanolayer (Ti,Cr)N Coatings Deposited at Different Cr Evaporator Currents (65 A Ti Evaporator Current, -150 V Substrate Bias)

The glass bead erosion results for the nanolayer (Ti,Cr)N coatings deposited as a function of evaporator current are shown in Figure 4.54 and 4.55 while the alumina test results are shown in Figure 4.56. For the glass bead tests, the coatings deposited with the

25 A and the 45 A Cr evaporator current failed after only 25 g of erodent. It is interesting that both of these coatings failed with only one 25 g test while the monolithic TiN and

CrN coatings deposited under similar conditions (-150 V substrate bias and -200 V, respectively) did not fail as rapidly. It appears that the nanolayering of TiN and CrN could result in the buildup of residual stresses due to lattice mismatch or intermixing of the nanolayers.

It seems possible that the poor erosion performance of the coatings deposited at

25 A and 45 A could be due to the fact these coatings have a greater number of nanolayers and therefore nanolayer interfaces than the other (Ti,Cr)N coatings.

Therefore the rapid crack coalesce failure mechanism could be due to interfacial stresses and poor fracture toughness. Further support for this argument is provided by the scratch adhesion testing which showed the lowest adhesion (lowest critical force, Fc) for the samples deposited at 25 A and 45 A Cr evaporator current (Section 4.11).

For the nanolayer coatings deposited with a 65 A and 85 A Cr evaporator current, the glass bead erosion performance seems much better than predicated by the monolithic

TiN coatings (65 A Ti evaporator current and-150 V) and CrN coating (45 A Cr evaporator current and -200 V CrN coating). In fact, the coating deposited with the 85 A

Cr evaporator current is the second best overall performing coating in this study, and this

141 is attributed to the increased CrN volume fraction. The (Ti,Cr)N coatings deposited with

65 A and 85 A Cr evaporator currents eroded in a similar fashion (uniform with a few large chips) as the monolithic TiN coating deposited with the -25 V substrate bias and also the CrN coating (-200 V substrate bias). In order for the erosion performance to increase relative to the monolithic coatings, there must be an increase in the fracture toughness and this may be accompanied by a change in residual stress.

The increasing CrN volume could be contributing to decreasing residual stress by serving as a compliant layer which accommodates the interfacial strains. In a related manner, the volume fraction of interfaces is decreasing with increasing CrN content, so the interfacial stresses should also be decreasing with increasing Cr evaporator current. It should be noted that due to the increased Cr evaporator current, increased thermal energy

(i.e. substrate temperature) may result in strain relief. One way to examine the effects of volume fraction of CrN versus the volume fraction of TiN would be to deposit nanolayer

(Ti,Cr)N coatings with significantly greater volume fractions of TiN as compared to CrN and compare the erosion performance to the coatings in this study.

Another interesting aspect of these coatings is that they do not perform well for the higher energy alumina erosion test conducted at a 60° impact angle. This was also the case for the monolithic CrN but not the monolithic TiN coating with the -25 V substrate bias. It appears that CrN systems might not fair as well as TiN for the aggressive alumina tests since the high hardness and angular shape may cause significant damage to the CrN layers since they are softer than the TiN layers.3,53

142

(Ti,Cr)N: Cr Evaporator@ 25A -0.0014 (Ti,Cr)N: Cr Evaporator @ 45A (Ti,Cr)N: Cr Evaporator @ 65A (Ti,Cr)N: Cr Evaporator @ 85A -0.0012 -0.00111 Wet Blasted AM355

-0.001

-0.0008

-0.0006

Mass Loss of Sample (g) -0.0004

-0.0002

0 0 500 1000 1500 2000 2500 3000 3500 4000 4500 5000 Total Dosage of Erodent (g)

(Ti,Cr)N: Cr Evaporator @ 25A -0.0016 (Ti,Cr)N: Cr Evaporator @ 45A (Ti,Cr)N: Cr Evaporator @ 65A (Ti,Cr)N: Cr Evaporator @ 85A -0.0014 Wet Blasted AM355 -0.00111 -0.0012

-0.001

-0.0008

-0.0006

Mass Loss of Sample (g) (g) Sample Loss of Mass -0.0004

-0.0002

0 0 500 1000 1500 2000 2500 3000 3500 4000 4500 5000 5500 Total Dosage of Erodent (g)

143

(Ti,Cr)N: Cr Evaporator @ 25A 1 (Ti,Cr)N: Cr Evaporator @ 45A -0.00350 (Ti,Cr)N: Cr Evaporator @ 65A (Ti,Cr)N: Cr Evaporator @ 85A AM355 -0.00300

-0.00250

-0.00200

-0.00150

Mass Loss of Sample (g) (g) Mass Loss Sample of -0.00100

-0.00050

0.00000 Figure 4.56. Alumina erosion media results shown as mass loss for a single 25 g test conducted on the nanolayer (Ti,Cr)N coating at different evaporator currents.

The large concentration of macroparticles does not appear to significantly detriment erosion resistance for high impact angle tests. Since the coatings deposited at

65 A and 85 A appear to have a high volume fraction of macroparticles (Section 4.7), it is logical to assume that these particles must be influencing the microstructure. Although

XRD crystallite measurements were not performed on these coatings due to the varying composition, it seems that macroparticles would have a definite impact on the crystallite size since the particles can disrupt the growth and possibly serve as nucleation sites for new grains. Another impact of macroparticles is formation of voids and porosity around the particle. Although erosion performance was not detrimentally effected, these particles would most likely decease the overall corrosion performance of the coatings.3

Additionally, since the macroparticles composition is metal rich, further research is

144 suggested to compare low angle versus high angle erosion of cathodic arc coatings containing a high volume fraction of macroparticles such as these (Ti,Cr)N coatings.

4.13.5 Erosion Performance of Nanolayer (Ti,Cr)N Coatings Deposited at Different Substrate Biases (65 A Ti Evaporator Current, 45 A Cr Evaporator Current)

The glass bead erosion results for the nanolayer (Ti,Cr)N coatings deposited with different substrate biases are shown in Figure 4.57 and Figure 4.58 and the alumina test results are shown in Figure 4.59. For the glass bead erosion, the coatings deposited at -

100 V, -150 V, -200 V and -300 V all failed rapidly with only one 25 g test. As with the nanolayer (Ti,Cr)N coatings deposited as a function of Cr evaporator current a high level of internal residual stress in these coatings possibly contributing to this rapid failure.

For the both the glass bead and aggressive alumina test (Figure 4.59), the -50 V coating was the best performing coating, but overall, this performance was not as good as the monolithic TiN coating deposited with a -25 V substrate bias. It seems that interfacial stresses could be attributed to the performance of these nanolayered (Ti,Cr)N coatings deposited at different substrate biases. For the coating deposited with the -50 V substrate bias, the crystallite size was determined to be small due to lack of diffusion (Section

4.3.4). In the -50 V nanolayered (Ti,Cr)N coating, the lack of diffusion should result in less intermixing of the TiN and CrN nanolayers as compared to the higher substrate bias samples. It seems likely that the interfacial mixing of the TiN and CrN nanolayers is resulting in decreased fracture toughness of the nanolayer (Ti,Cr)N coatings, and the erosion results suggest that between -50 V and -100 V the interfacial mixing of the TiN and CrN rich layers exceeds a critical level in which the coatings will no longer perform well in erosion tests due to decreased fracture toughness. This is suggestive of a critical

145 interfacial residual stress between -50 V and -100 V for coatings produced with these deposition conditions, and further research is suggested to study the relationship between of interfacial volume, residual stress, and intermixing in the (Ti,Cr)N system.

The macroparticles do not appear to be negatively affecting high angle erosion since the coating with the largest volume fraction of macroparticles showed the best erosion performance, and this was also the case for the (Ti,Cr)N coatings deposited at different Cr evaporator currents. This suggests that the macroparticles may actually be increasing the compliancy of the (Ti,Cr)N coating system, and further study is again suggested for high angle versus low angle erosion performance of macroparticle containing coatings such as these (Ti,Cr)N coatings.

A few other structure-property-performance relationships were observed for the

(Ti,Cr)N deposited at different substrate biases. The best performing nanolayer (Ti,Cr)N coatings also had the highest surface roughness, but this was believed to be primarily due to macroparticles. Additionally, the smallest crystallite size is linked to the best erosion performance, and this suggests that crystallite and therefore grain boundary density could play a role in the erosion rate by increasing toughness through an energy dissipation mechanism. As with the monolithic TiN samples, further research is necessary to conclude whether grain size plays a large role in the erosion performance, or whether the grain size is secondary to factors such as residual stress.

146

-0.002 (Ti,Cr)N: Substrate Bias @ -50V (Ti,Cr)N: Substrate Bias @ -100V -0.0018 (Ti,Cr)N: Substrate Bias @ -150V (Ti,Cr)N: Substrate Bias @ -200V -0.0016 (Ti,Cr)N: Substrate Bias @ -300V Wet Blasted AM355 -0.0014

-0.0012 -0.00111 -0.001

-0.0008

-0.0006 Mass Loss of Sample (g)

-0.0004

-0.0002

0 0 200 400 600 800 1000 1200 1400 1600 1800 2000 Total Dosage of Erodent (g)

Figure 4.57. Glass bead erosion results shown as mass loss versus total dosage of erodent on a single 1/4 in diameter spot for the nanolayer (Ti,Cr)N coatings deposited at different substrate biases. Original test. -0.0018 (Ti,Cr)N: Substrate Bias @ -50V (Ti,Cr)N: Substrate Bias @ -100V -0.0016 (Ti,Cr)N: Substrate Bias @ -150V (Ti,Cr)N: Substrate Bias @ -200V (Ti,Cr)N: Substrate Bias @ -300V -0.0014 Wet Blasted AM355 -0.0012 -0.00111

-0.001

-0.0008

-0.0006 Mass Loss of Sample (g) -0.0004

-0.0002

0 0 500 1000 1500 2000 2500 3000 3500 4000 Total Dosage of Erodent (g)

147

(Ti,Cr)N: -50V Substrate Bias 1 (Ti,Cr)N: -100V Substrate Bias -0.00350 (Ti,Cr)N: -150V Substrate Bias (Ti,Cr)N: -200V Substrate Bias (Ti,Cr)N: -300V Substrate Bias -0.00300 AM355

-0.00250

-0.00200

-0.00150

Mass Loss of Sample (g) Sample (g) Loss of Mass -0.00100

-0.00050

0.00000 Figure 4.59. Alumina erosion media results shown as mass loss for a single 25 g test conducted on the nanolayer (Ti,Cr)N coating at different substrate biases.

4.13.6 Erosion Performance of Multilayer(Ti,Cr)N Coatings Deposited at (65 A Ti Evaporator Current, 45 A Cr Evaporator Current, -150 V Substrate Bias)

The glass bead erosion results for the multilayer (Ti,Cr)N coatings with Ti interlayers are shown in Figure 4.60 and Figure 4.61. These coatings showed an interesting trend in which the erosion performance initially increased for coatings with two and four layers of (Ti,Cr)N as compared to the (Ti,Cr)N coating without interlayers

(65 A Ti evaporator current, 45 A Cr evaporator current, -150 V substrate bias). This could possibly indicate the coatings ability to absorb force from impacting particles due to the ductility of the Ti interlayers. However, the coatings with eight and sixteen layers of (Ti,Cr)N showed rapid erosion failure which could indicate too high of a coating residual stress. Process optimization is recommended for future investigations.

148

As discussed in Section 4.3.6, the (111) preferred crystallographic orientation of the coatings increased with number of Ti interlayers, and this crystallographic orientation is most likely related to the TiN nanolayer crystallographic orientation and not the CrN nanolayers. The crystallographic orientation of the (111) TiN nanolayers and the likely

(220) crystallographically oriented CrN nanolayers partially contributed to the interfacial stress. No strain relief (or compliancy) with increasing Ti interlayers was observed in the erosion tests, which was unexpected. Correspondingly, a similar trend was seen for the alumina erosion test (Figure 4.64).

-0.003 Monolithic (Ti,Cr)N Multilayer (Ti,Cr)N w/ Ti Interlayers: 2 Layers of Nitride Multilayer (Ti,Cr)N w/ Ti Interlayers: 4 Layers of Nitride -0.0025 Multilayer (Ti,Cr)N w/ Ti Interlayers: 8 Layers of Nitride Multilayer (Ti,Cr)N w/ Ti Interlayers: 16 Layers of Nitride Wet Blasted AM355 -0.002

-0.0015

-0.00111 -0.001 Mass Loss of Sample (g) (g) Mass Loss of Sample

-0.0005

0 0 50 100 150 200 250 300 350 400 Total Dosage of Erodent (g)

Figure 4.60. Glass bead erosion results shown as mass loss versus total dosage of erodent on a single 1/4 in diameter spot for the multilayer (Ti,Cr)N coatings deposited with Ti interlayers. Original test.

149

-0.004 Monolithic (Ti,Cr)N Multilayer (Ti,Cr)N w/ Ti Interlayers: 2 Layers of Nitride Multilayer (Ti,Cr)N w/ Ti Interlayers: 4 Layers of Nitride -0.0035 Multilayer (Ti,Cr)N w/ Ti Interlayers: 8 Layers of Nitride Multilayer (Ti,Cr)N w/ Ti Interlayers: 16 Layers of Nitride -0.003 Wet Blasted AM355

-0.0025

-0.002

-0.0015 -0.00111 Mass Loss of Sample (g) (g) Sample of Loss Mass -0.001

-0.0005

0 0 50 100 150 200 250 300 350 400 450 500 Total Dosage of Erodent (g)

Figure 4.61. Glass bead erosion results shown as mass loss versus total dosage of erodent on a single 1/4 in diameter spot for the multilayer (Ti,Cr)N coatings deposited with Ti interlayers. Confirmation test spot.

The multilayer (Ti,Cr)N coatings with Nb interlayers showed mixed performance

(Figures 4.64 and 4.65). The coating with sixteen layers of (Ti,Cr)N showed improved performance, while the coatings with only two, four and eight (Ti,Cr)N layers showed rapid failure. The decent performance of the sixteen layers is possibly attributed to residual stress relaxation from the Nb interlayers attributed to the increased deposition temperature during Nb deposition (Section 3.3). This residual stress reduction may be greater than the relief provided by the Ti interlayers. Additional residual stress relief could also be attributed to the compliancy of Nb, which should accommodate coating strain caused by interfacial stress. The increased number of the Nb interlayers were expected to provide better strain relief with regards to interfacial stresses, and this is possibly reflected in the erosion performance.

150

The exact reason for the failure of the coatings containing two, four, eight layers of (Ti,Cr)N with Nb interlayers is unknown, but the rapid failure is indicative of poor adhesion or residual stress issues. The critical adhesion force, Fc, for the coating with eight layers of (Ti,Cr)N was 52±2 N compared to 45±2 N for the coating with sixteen layers of (Ti,Cr)N, and this shows that adhesion actually decreased with the increasing number of Nb layers, so adhesion was most likely not a factor in the rapid failure of these coatings.

Additionally, the alumina erosion performance for these coatings follows a similar trend as shown for the Ti interlayer samples (Figure 4.64). The reason for the differing trends between the glass bead and alumina erosion for the Nb interlayer samples is related to the damage tolerance mechanism for the Nb interlayer coatings. These coatings cannot withstand high energy impacts without rapid failure attributed to the residual stress most likely arising from the previously described interfacial stresses. By adding the high number of Nb interlayers, the coating residual stress is most likely reduced and the coating system is more compliant. This correlates to an increase in the performance of the glass bead erosion, but there does not seem to be a large enough improvement in the coating properties for the system to withstand the more aggressive and higher energy alumina erosion test.

Another possibility is that the nanolayered structure created inconsistent interfacial compositions in regards to the interfaces between the nanolayer (Ti,Cr)N and the Ti (or Nb). For example, a CrN rich nanolayer and Ti interlayer interface could cause greater coating residual stress than a TiN rich nanolayer and Ti interlayer interface

151 or vice versa. These inconsistent interfaces could then possibly increase or decrease the coating residual stress depending on the composition.

-0.0025 Monolithic (Ti,Cr)N Multilayer (Ti,Cr)N w/ Nb Interlayers: 2 Layers of Nitride Multilayer (Ti,Cr)N w/ Nb Interlayers: 4 Layers of Nitride Multilayer (Ti,Cr)N w/ Nb Interlayers: 8 Layers of Nitride -0.002 Multilayer (Ti,Cr)N w/ Nb Interlayers: 16 Layers of Nitride Wet Blasted AM355

-0.0015

-0.00111 -0.001 Mass Loss Sample of (g)

-0.0005

0 0 50 100 150 200 250 300 350 400 Total Dosage of Erodent (g)

Figure 4.62. Glass bead erosion results shown as mass loss versus total dosage of erodent on a single 1/4 in diameter spot for the multilayer (Ti,Cr)N coatings deposited with Nb interlayers. Original test. -0.0025 Monolithic (Ti,Cr)N Multilayer (Ti,Cr)N w/ Nb Interlayers: 2 Layers of Nitride Multilayer (Ti,Cr)N w/ Nb Interlayers: 4 Layers of Nitride Multilayer (Ti,Cr)N w/ Nb Interlayers: 8 Layers of Nitride -0.002 Multilayer (Ti,Cr)N w/ Nb Interlayers: 16 Layers of Nitride Wet Blasted AM355

-0.0015

-0.00111 -0.001 Mass Loss of Sample (g) (g) of Sample Mass Loss

-0.0005

0 0 50 100 150 200 250 300 350 400 450 500 Total Dosage of Erodent (g) Figure 4.63. Glass bead erosion results shown as mass loss versus total dosage of erodent on a single 1/4 in diameter spot for the multilayer (Ti,Cr)N coatings deposited with Nb interlayers. Confirmation test spot.

152

Monolithic (Ti,Cr)N 2 Layers (Ti,Cr)N with Ti Interlayers 4 Layers (Ti,Cr)N with Ti Interlayers 8 Layers (Ti,Cr)N with Ti Interlayers 16 Layers (Ti,Cr)N with Ti Interlayers 2 Layers (Ti,Cr)N with Nb Interlayers 4 Layers (Ti,Cr)N with Nb Interlayers 8 Layers (Ti,Cr)N with Nb Interlayers 16 Layers (Ti,Cr)N with Nb Interlayers 1AM355 -0.00450

-0.00400

-0.00350

-0.00300

-0.00250

-0.00200

-0.00150

Mass Loss of Sample (g) (g) Loss of Sample Mass -0.00100

-0.00050

0.00000 Figure 4.64. Alumina erosion media results shown as mass loss for a single 25 g test conducted on the multilayer (Ti,Cr)N coating with Ti and Nb interlayers.

4.14 Proposed Coating Design

The combination of all the information presented in this chapter should theoretically provide a road map for the deposition of nanolayer (Ti,Cr)N coatings with good erosion performance. The proposal in this section is one of many possible routes to take to improve the performance of these coating systems. First, the relative thickness of the CrN nanolayers to the TiN nanolayers should be relatively large with the coatings containing at least a ratio of 2.5 of Cr to Ti (atomic percent). Additionally, the Ti evaporator target should be run with a high current (85 A) while the Cr evaporator target should be run with less current (less than 65 A). The high Ti current will promote adatom diffusion while the lower Cr evaporator current will decrease macroparticle

153 generation. In order to maintain the correct Cr/Ti ratio, two Cr targets and one Ti target could be used simultaneously.

The second variable to consider is the substrate bias which should be held relatively low (less than -100 V). An optimal substrate bias was not found in this study however it appears that a -25 V substrate bias would work well. The disadvantage with such a low substrate bias is the macroparticle defect density could negatively affect corrosion performance. A further study involving the substrate bias may be required to understand the macroparticle-corrosion-erosion relationship.

In addition, the interfacial compositions, structures, and periodicity should be further investigated in these coatings. Since these coatings were determined to be nanolayered structures, it might be beneficial to grade the coating compositions around the bondlayer/(Ti,Cr)N layer interfaces. For example, after depositing the Ti bondlayer it might be beneficial to begin depositing binary TiN before evaporating any Cr. Once a thin TiN (10’s nm) layer has coated all of the substrates, the Cr evaporator can be ignited in order to deposit nanolayer (Ti,Cr)N. This will provide a more uniform and consistent bondlayer/(Ti,Cr)N interfacial composition. The graded interface technique should then be followed for all of the interfaces in a multilayer coating containing metallic compliant layers such as Ti and Nb.

154

CHAPTER 5

CONCLUSION

Chapter 5 discusses the conclusions regarding the synthesis-structure-property- performance relationships for the TiN, CrN, and nanolayer (Ti,Cr)N coatings deposited by cathodic arc deposition. In general, the relationship between crystallographic orientation, coating nanostructure, and residual stress is essential to the erosion performance of nanolayer (Ti,Cr)N coatings.

5.1 Monolithic TiN Coatings Deposited as a Function of Evaporator Current (- 150 V Substrate Bias)

The monolithic TiN coating with the 65 A evaporator current performed better than the coatings deposited with the 45 A and 85 A evaporator currents (-150 V substrate bias). This was attributed to a balance between a high compressive residual stress which impedes crack propagation but does not compromise adhesion, and good fracture toughness which is possibly related to a fine crystallite structure. Although the coating deposited at 45 A had a finer crystallite size than the 65 A coating, the erosion resistance was worse and attributed to decreased adhesion due to excessively high compressive stress. All of these coatings had a strong (111) preferred crystallographic orientation, which was attributed to strain relief.

5.2 Monolithic TiN Coatings Deposited as a Function of Substrate Bias (65 A Ti Evaporator Current)

The monolithic TiN coating with the -25 V substrate bias performed better than the coatings deposited with the -150 V and -300 V substrates biases (65 A Ti evaporator

155 current). This was attributed to a relatively large compressive residual stress which impedes crack propagation but does not compromise adhesion. The large scale micro- chipping was not as predominant in this coating, and the coating showed mostly uniform erosion. This coating had a (111) preferred crystallographic orientation, but this crystallographic orientation was not as strong as the other TiN coatings deposited in this study.

5.3 Monolithic CrN Coating Deposited at 45 A and -200 V Substrate Bias

The monolithic CrN showed very good erosion performance. The Cr evaporator target was prone to a greater degree of macroparticle generation as compared to the Ti evaporator target, and this appears to correspond to the lower melting point/decomposition of CrN than TiN. The coating had a (220) preferred crystallographic orientation which was attributed to a structure switch between bcc Cr and B1 NaCl CrN.

5.4 Nanolayer (Ti,Cr)N Coatings Deposited with Different Cr Evaporator Currents (65 A Ti Evaporator Current, -150 V Substrate Bias)

The nanolayer (Ti,Cr)N coatings deposited at different Cr evaporator currents were shown to be nanolayered structures of alternating TiN rich and CrN rich layers with a periodicity, λ, ranging from 10.5 to 21.0 nm (ADF-STEM). The erosion performance corresponded to the interfacial volume as well as the volume of CrN in the coatings.

Coatings with a high number of interfaces performed poorly and this was considered to be related to the mismatch of the (111) preferred crystallographic orientation of the TiN nanolayers and the (220) preferred crystallographic orientation of the CrN nanolayers.

The coatings with the 65 A and 85 A Cr evaporator currents appeared to have performed

156 better than predicted in regards to the monlithic TiN and CrN coatings deposited under similar conditions. This was attributed to decreasing interfacial residual stresses due to decreasing volume fraction of nanolayer interfaces as well as increasing compliancy with increasing CrN nanolayer thickness. The coating with the 85 A evaporator current had the best erosion performance within this set of coatings.

5.5 Nanolayer (Ti,Cr)N Coatings Deposited with Different Substrate Biases (65 A Ti Evaporator Current, 45 A Cr Evaporator Current)

The nanolayer (Ti,Cr)N coatings deposited at different substrate currents were shown to be nanolayered structures of alternating TiN rich and CrN rich layers with possibly greater intermixing of the nanolayers at high bias voltages (especially -200 V and -300 V). The coating deposited at lowest substrate bias -50 V was believed to have performed well due to the decreased interfacial residual stresses resulting from decreased intermixing (thermal and ion bombardment induced).

5.6 Multilayer (Ti,Cr)N Coatings Deposited with Ti and Nb Interlayers (65 A Ti Evaporator Current, 45 A Cr Evaporator Current, -150 V Substrate Bias)

The multilayer (Ti,Cr)N coatings deposited with Ti interlayers showed improvements in erosion performance for the coating with two layers of nanolayer

(Ti,Cr)N as compared to the (Ti,Cr)N coating with no interlayers (65 A Ti Evaporator

Current, 45 A Cr Evaporator Current, -150 V Substrate Bias). The improvement was contributed to the increased compliancy of the Ti interlayers which aided in the reduction of the coating residual stress. However, the (111) crystallographic orientation of the TiN nanolayers was shown to increase with increasing number of Ti interlayers and this was attributed to the templated growth of (111) TiN on (001) Ti due to the 1.7% lattice

157 mismatch and the interstitial structure of TiN. The increasing (111) preferred crystallographic orientation of TiN nanolayers was believed to contribute to increased interfacial stresses between the TiN and CrN nanolayers, and therefore a reduction in erosion performance.

The multilayer (Ti,Cr)N coatings deposited with Nb interlayers showed improvements in erosion performance for the coating with sixteen layers of nanolayer

(Ti,Cr)N as compared to the (Ti,Cr)N coating with no interlayers (65 A Ti Evaporator

Current, 45 A Cr Evaporator Current, -150 V Substrate Bias). This was attributed to stress relaxation due to the higher substrate temperature during deposition of Nb (increase of approximately 75 °C) and the compliancy of Nb. However, only the coating with sixteen layers showed improvement suggesting a relationship between the distribution of the Nb layers and the stress relaxation of the (Ti,Cr)N layers.

158

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