Low-Temperature Gas-Phase Nitriding and Nitrocarburizing Of

LOW-TEMPERATURE GAS-PHASE NITRIDING

AND NITROCARBURIZING OF

316L AUSTENITIC STAINLESS STEEL

DANDAN WU

Submitted in partial fulfillment of the requirements

For the degree of Doctor of Philosophy

Dissertation Advisor: Dr. Arthur H. Heuer

Department of Materials Science and Engineering

CASE WESTERN RESERVE UNIVERSITY

January, 2013 CASE WESTERN RESERVE UNIVERSITY

SCHOOL OF GRADUATE STUDIES

We hereby approve the thesis/dissertation of

Dandan Wu ______

Doctor of Philosophy candidate for the ______degree *.

Arthur H. Heuer (signed)______(chair of the committee) James D. McGuffin-Cawley ______Frank Ernst ______John J. Lewandowski ______Farrel J. Martin ______

8/31/2012 (date) ______

*We also certify that written approval has been obtained for any proprietary material

contained therein.

Dedicated to my mother, Yuqin Wang,

and my husband, Hongyi Yuan Table of Contents List of Tables ...... iv

List of Figures ...... vi

Acknowledgements ...... xiv

Abstract ...... xvi

Chapter 1 Introduction ...... 1

1.1 Background of expanded austenite ...... 1

1.2 Crystal structure of expanded austenite ...... 3

1.3 Diffusion of nitrogen in expanded austenite ...... 7

1.4 Orientation-dependent case depth in low-temperature nitriding ...... 11

1.5 Effect of stress on solubility and diffusivity of interstitial atoms ...... 12

1.6 Room-temperature ferromagnetism induced by low-temperature nitriding ...... 15

1.7 Dilation of austenite lattice from nitrogen and carbon interstitials ...... 19

1.8 Precipitate formation and stability of expanded austenite ...... 24

1.9 Nitrocarburizing processes ...... 27

References ...... 31

Chapter 2 Experimental Methods ...... 36

2.1 Design of low-temperature gas-phase nitriding process ...... 36

2.2 Design of low-temperature gas-phase nitrocarburizing process ...... 40

2.3 Characterization techniques ...... 43

2.3.1 X-Ray Diffraction ...... 43

2.3.2 Auger Electron Spectroscopy ...... 46

2.3.3 Scanning Electron Microscopy ...... 49

i 2.3.4 Electron Backscatter Diffraction ...... 50

2.3.5 Magnetic Force Microscopy ...... 51

2.3.6 Microhardness tester ...... 51

2.3.7 Transmission Electron Microscopy ...... 51

References ...... 53

Chapter 3 Low-Temperature Gas-Phase Nitriding ...... 54

3.1 Effect of processing parameters ...... 54

3.1.1 Effect of nitriding temperature ...... 54

3.1.2 Effect of nitriding activity ...... 63

3.1.3 Effect of nitriding duration ...... 70

3.2 Ferromagnetism induced by low-temperature nitriding ...... 75

3.3 Orientation-dependent case depth ...... 82

3.4 Hardness and modulus measurements ...... 86

9 3.5 TEM results on nitrided 316L (aN2 = 4×10 ) ...... 91

References ...... 107

Chapter 4 Low-Temperature Gas-Phase Nitrocarburizing ...... 108

4.1 Three scenarios of nitrocarburizing with NH3/CO/H2/N2 ...... 108

4.1.1 XRD analysis ...... 108

4.1.2 Metallography ...... 112

4.1.3 AES depth profiles ...... 115

4.1.4 Hardening effect ...... 118

4.1.5 Study of second phases after 20C+20N by TEM ...... 119

4.2 Effect of nitriding activity on 20C+20N ...... 128

4.2.1 XRD analysis ...... 128

4.2.2 Metallography ...... 131

ii 4.2.3 AES depth profiles ...... 132

4.2.4 Hardness measurements ...... 133

4.3 Effect of nitriding activity on 20 (N+C) ...... 134

4.3.1 XRD results ...... 135

4.3.2 Metallography ...... 137

4.3.3 AES depth profiles ...... 139

4.3.4 Hardening effect ...... 141

References ...... 142

Chapter 5 Discussion...... 143

5.1 Effect of temperature ...... 143

5.2 Effect of nitriding activity ...... 146

5.3 Lattice parameter expansion induced by nitrogen ...... 148

5.4 Anisotropy in lattice parameter ...... 153

5.5 On the nitrogen content measured by AES ...... 159

5.6 XRD peak splitting during stress measurements ...... 161

5.7 Orientation-dependent case depth after low-temperature nitriding ...... 164

References ...... 166

Chapter 6 Conclusions ...... 168

Appendices ...... 171

Appendix I Nitro-carburizing by Urea (CO(NH2)2) ...... 171

Appendix II Nitro-carburizing in NH3/C2H2/H2/N2 ...... 178

iii List of Tables

Table 1.1 Calculated values of S1hkl and Ghkl ...... 7

Table 1.2 Survey of α parameters in Vegard’s law for nitrogen and carbon ...... 23

Table 2.1 Processing parameters to study the effect of nitriding temperature ...... 39

Table 2.2 Processing parameters to study the effect of nitriding activity ...... 39

Table 2.3 Processing parameters for nitriding, carburizing and three different nitrocarburing processes ...... 41

Table 2.4 Designed parameters of two 20C+20N nitrocarburizing experiments ...... 42

Table 2.5 Designed parameters of three 20(N+C) nitrocarburizing experiments ...... 42

Table 2.6 Comparison between the certified and AES measured composition on Sample B.S. 81N ...... 48

Table 2.7 Comparison between the certified and AES measured composition on Sample NSC-4 ...... 49

Table 3.1 Calculated lattice parameters and lattice parameter expansions of samples nitrided at different temperatures ...... 57

Table 3.2 Calculated lattice parameters and lattice parameter expansions of samples nitrided at different temperatures ...... 65

Table 3.3 Calculated lattice parameters and lattice parameter expansions of samples nitrided for different durations ...... 72

Table 3.4 Gas-phase nitriding activities and the corresponding lattice parameter expansion ...... 76

Table 3.5 Comparison between 2θ positions from measured XRD and standard peaks of nitrides...... 92

Table 3.6 Lattice parameters of different planes of the γ’-M4N structure calculated from DPs ...... 101

iv Table 3.7 Zone axes of γ matrix in which twin and HCP plates may/may not be distinguished ...... 101

Table 4.1 Lattice parameters of 316L stainless steel samples after carburizing, nitriding and various nitrocarburizing processes ...... 111

Table 4.2 Lattice parameters and lattice parameter expansions measured from the two different 20C+20N samples ...... 131

Table 5.1 Calculated XECs S1hkl and XRD measured 0, ℎ ...... 164

Table A-1 Experimental parameters of nitrocarburizing processes in urea powders ..... 173

Table A-2 Experimental parameters of simultaneous nitrocarburizing processes in NH3/C2H2/H2/N2 ...... 180

v List of Figures

Figure 1.1 Microstructure of 316L stainless steel nitrided at 703K for 20 hrs ...... 2

Figure 1.2 X-ray diffraction patterns of the ion-nitrided stainless steel at 673K in different ratios of N2 and H2 (a) 1:100, (b) 1:9 and (c) 4:1 ...... 3

Figure 1.3 Profiles of residual stress profiles in AISI 316 obtained by simulating X-ray diffraction pattern. Heat treatment conditions: Carburizing: 793 K/ 2.5 h / 90% CO+ 10% -1/2 H2. Nitriding: 22 h / 718 K with the nitriding potentials KN = 0.293 bar and KN = 2.49 -1/2 bar . Carburizing and nitriding: 793 K / 2 h / 30% CO + 70% H2 and 713 K / 23 h / KN = 1.14...... 5

Figure 1.4 Modeling XRD of nitrided bulk samples ...... 6

Figure 1.5 Nitrogen concentration depth profile measured by GDOES, from an AISI 321 sample after low energy ion implantation...... 9

Figure 1.6 Nitrogen concentration depth profiles observed by nitriding under typical low energy ion, implantation conditions first with the 15N isotope and subsequently with the 14N isotope. The 14N isotope is not present at the diffusion front. The profiles were measured by deuterium induced nuclear reaction analysis...... 10

Figure 1.7 Diffusivity of nitrogen calculated from fitting thermogravimetric data obtained from ntiridng and stepwise denitriding experiments of stainless steel thin foils...... 10

Figure 1.8 Simulated concentration dependent nitrogen diffusion profile for gas-phase nitrided AISI 316 assuming a constant surface concentration condition...... 11

Figure 1.9 (a) SEM cross-sectional micrograph demonstrating an orientation-dependent case depth on a plasma nitrided Inconel 690 sample and (b) Local layer thickness e as a function of Ahkl measured on the same sample...... 12

Figure 1.10 Photograph of a plasma carburizing unit with in-situ stress applying component and the impact of tensile stress on the case depth of carburized sample (from b to e)...... 14

Figure 1.11 Total energy (relative to a minimum energy) as a function of magnetic moment for fcc-Fe showing the 3 different magnetic states ...... 16

vi Figure 1.12 Total energy per atom and magnetic moment as a function of volume for bcc- Fe and fcc-Fe...... 17

Figure 1.13 Measured hyperfine-filed values near saturation (at 4.2K for γ-Fe particles in Cu100-xAlx, at 30K of 5 monolayer(ML) γ-Fe (001) on Cu (001) and Cu3Au(001), respectively) versus Wigner-Seitz radius γws/a.u...... 17

Figure 1.14 (a) SPM (Scanning probe microscopy) and (b) MFM (Magnetic Force Microscopy) image observed from the same area of a AISI 316 sample after low- temperature nitriding. The scale is the same in the two images...... 19

Figure 1.15 (a) lattice parameter a of expanded austenite as a function of (a) numbers of interstitial nitrogen or carbon atoms per metal atom (yN or yC) and (b) concentration of interstitial nitrogen or carbon (XN or XC)...... 22

Figure 1.16 Lehrer diagram, which predicted the stable phases for Fe-N systwm as a function of temperature at nitriding potential used...... 25

Figure 1.17 Selected-area diffraction patterns from a nitridied 316 Stainless Steel. Forbbiden reflections were detected from both [100] and [211] zone axis...... 25

Figure 1.18 SAD patterns tanken from an expanded austenite region during in-situ annealing at (a) 0 °C; (b) 440 °C; (c) 500 oC and (d) 600 oC...... 26

Figure 1.19 Hardness depth profile of nitrocarburizied AISI 316. The sample was first carburized at 773K for 4 hours (in 100 vol.% CO) and subsequently nitrided at 713K for 18.5 hours (in 100 vol.% NH3). The hardness depth profiles of the samples carburied and nitrided with the same gas composition were given as a comparison...... 28

Figure 1.20 GDOES depth profiles of stainless steel 304 specimen after (a) 8 hours simultaneous plasma nitriding and carburizing process, (b) 4 hours plasma nitriding followed by 4 hours carburizing process and (c) 4 hours plasma carburizing followed by 4 hours nitriding process...... 29

Figure 1.21 Simulated concentration-depth profiles C(act.), C(eff.) and N (act.) after a plasma nitro-carburizing process (15 hours plasma carburizing followed by 15 hours plasma nitriding). [56] ...... 30

Figure 2.1 Picture of the CVD furnace employed ...... 38

Figure 2.2 Schematic of procedures of low-temperature nitriding process ...... 38

Figure 2.3 Schematic of three scenarios of low-temperature nitrocarburizing processes 40

vii Figure 2.4 Schematic of the plane stress elastic model...... 46

Figure 2.5 AES spectra observed from sample NSC-4 at different sputtering depths ..... 48

Figure 3.1 XRD results of 316L stainless steel samples treated at different temperatures ...... 56

Figure 3.2 Lattice parameters ahkl from samples nitrided at different temperatures ...... 57

2 Figure 3.3 a311 versus sin ψ from samples nitrided at 350 °C and 420 °C ...... 59

Figure 3.4 X-ray scans from the sample nitrided at 440 °C in aN2 = 7400 at two ψ tilts .. 59

Figure 3.5 Cross-sectional SEM images of 316L stainless steel samples nitrided at different temperatures. The nitriding activity for all samples was set the same, aN2 = 7400...... 61

Figure 3.6 SEM image revealing crack formation on 316L stainless steel samples after nitriding at 450 °C, aN2 = 7400...... 61

Figure 3.7 Nitrogen concentration profiles detected from samples nitrided at different temperatures ...... 62

Figure 3.8 XRD results of 316L stainless steel samples treated at different activities ..... 64

Figure 3.9 Lattice parameters ahkl from samples nitrided at 6 different activities ...... 64

2 Figure 3.10 a311 versus sin ψ from samples nitrided with 3 different nitriding activities 66

9 Figure 3.11 Grazing XRD result from 316L treated at 440 °C with aN2 = 4×10 ...... 67

Figure 3.12 Cross-sectional SEM images of 316L stainless steel samples nitrided at different activities: (a) 4×109, (b) 7400, (c) 1700 and (d) 200. The nitriding temperature for all samples was 440 °C...... 68

Figure 3.13 Cross-sectional SEM images of 316L stainless steel sample (440°C, aN2 =1.8×105) revealing the crack formation. The original surface is at the left-hand side of the image...... 68

Figure 3.14 Cross-sectional SEM images of 316L stainless steel sample (440°C, aN2 = 4×109) revealing the crack formation. The original surface is at the bottom of the image...... 69

viii Figure 3.15 Nitrogen concentration profiles detected from samples nitrided with different nitriding activities at 440 °C ...... 70

Figure 3.16 XRD results of 316L stainless steel samples nitrided for different durations 71

Figure 3.17 Lattice parameters ahkl from samples nitrided for different durations ...... 72

Figure 3.18 AES profiles measured from 316L samples nitrided for different durations 74

Figure 3.19 Cracks formed on 316L sample after 80 hours of nitriding ...... 74

Figure 3.20 (a) EBSD map of the SS 316L sample nitrided in aN=1700, (b) the MFM image of the same region in (a), (c) EBSD map of the SS 316L sample nitrided in aN=7400, (d) the MFM image of the same region in (c), (e) the color legend of the EBSD mapping in (a) and (c)...... 77

Figure 3.21 Cross-sectional SEM image from nitrided 316L stainless steel (aN2=7400). The original surface is on the top of the image...... 79

Figure 3.22 (a) Cross-sectional AFM image, and (b) the corresponding MFM image for nitrided 316L stainless steel (aN = 7400). The original surface is on the left side of the images. (c) Nitrogen concentration depth profiles taken using AES from the lines shown in Figs. 3.22 (a) and (b)...... 80

Figure 3.23 Diffraction patterns taken from 316L nitrided in aN2 = 7400 (a) [011] zone axis and (b) [112] zone axis...... 81

Figure 3.24 (a) Cross-sectional SEM image from 316L after nitrided for 80 hours; (b) EBSD orientation map taken from the same area; and (c) color key of (b) ...... 84

Figure 3.25 Nitrogen concentration profiles detected from samples nitrided for (a) 5 hours, (b) 20 hours (c) 80 hours and (d) the plot of case depth (read from AES nitrogen concentration depth profile) v.s. square root of time ...... 85

Figure 3.26 AES carbon depth profiles measured from different surface oriented grains of a carburized 316L ...... 86

Figure 3.27 Plan view microhardnesses of 316L stainless steel samples before and after low-temperature nitriding with different nitriding activities ...... 87

Figure 3.28 (a) plan-view EBSD orientation map obtained from sample nitrided at 440 °C in aN2 = 7400 for 20 hours; (b) optical microscopy image showing an indent; (c) hardness profiles; and (d) modulus profiles measured in CSM mode ...... 90

ix Figure 3.29 STEM image taken from a cross-sectional 316L sample nitrided at 440 °C in 9 aN2 = 4×10 ...... 92

9 Figure 3.30 (a) Plan-view SEM image of 316L sample nitrided at 440 °C in aN2 = 4×10 and (b) STEM image showing the cross-section of the formed surface particle ...... 94

Figure 3.31 Elemental mapping by XEDS of the formed surface particle after 316L 9 sample was nitrided at 440 °C in aN2 = 4×10 . (a) STEM image; (b) N map; (c) Fe map; (d) Cr map; (e) Ni map and (f) Pt map (protective layer before FIB sample preparation)...... 95

Figure 3.32 SAD patterns observed from the surface particle in (a) [3, -1, -2, 2] zone axis, (c) [4, -2, -2, 3] zone axis and (e) zone axis of ε-M2N1-x. The corresponding JEMS simulated diffraction patterns in the same zone axes are shown in Figure (b), (d) and (e), respectively. Double diffraction spots are included in the simulated patterns...... 96

9 Figure 3.33 SAD patterns taken from 316L nitrided at 440 °C in aN2 = 4×10 for 20 hours in (a) [111] zone axis, (c) [100] zone axis, (e) [112] zone axis and (g) [011] zone axis. Corresponding diffraction patterns simulated by JEMS software based on an expanded γ’-M4N structure are shown in (b), (d), (f) and (h), respectively...... 100

Figure 3.34 (a) SAD pattern with a twin structure γ’-M4N [122]// γ’-M4N [100]; (b) JEMS simulated diffraction pattern with γ’-M4N [122]// γ’-M4N [100]; (c) Bright field image showing micro-twin; (d) Dark field image taken with reflection 1 and (e) Dark field image taken with reflection 2 ...... 104

Figure 3.35 (a) Bright filed image taken from the circled area in Figure 3.30. (b) Corresponding SAD pattern. (c) Simulated SAD pattern with γ’-M4N [111] // ε-M2N1-x [0001]. (d) Dark field image taken with diffraction spot 1 in (b). (f) Dark field image taken with (110) reflection of γ’-M4N (spot 2 in (b))...... 105

Figure 3.36 (a) SAD pattern in [011] zone axis of γ’-M4N and (b) JEMS simulated composite diffraction pattern with γ’-M4N <011> // ε-M2N1-x [-12-10] ...... 106

Figure 4.1 X-ray diffraction patterns of 316L stainless steel before and after carburizing, nitriding and various nitrocarburizing scenarios ...... 110

Figure 4.2 ahkl of 316L stainless steel samples after carburizing, nitriding and various nitrocarburizing processes ...... 111

Figure 4.3 Grazing angle XRD data of 316L stainless steel after carburizing, nitriding and various nitrocarburizing processes ...... 112

x Figure 4.4 SEM images on the cross sections of 316L bulk treated with (a) 20 hour nitriding (20N); (b) 20 hours carburizing (20C); (c) 20 hours simultaneous nitrocarburizing (20(N+C)); (d) 20 hours nitriding followed by 20 hours carburizing (20 N+20C); and (e) 20 hour carburizing following by 20 hour nitriding (20C+20N)...... 114

Figure 4.5 AES depth profiles of nitrogen and carbon measured from 316L treated by (a) 20N; (b) 20C; (c) 20(N+C); (d) 20N+20C and (e) 20C+20N ...... 117

Figure 4.6 AES spectrum shows carbon peak detected from the nitrogen-enriched layer of 20(N+C) ...... 118

Figure 4.7 Surface hardness of 316L bulk sample treated with nitriding, caburizing and nitrocarburizing processes ...... 119

Figure 4.8 (a) STEM image taken from nitrocaburized 316L by a 20C+20N process and (b) XEDS result from the surface C (soot) layer...... 123

Figure 4.9 (a) Bright-field image taken from a cross-sectional nitrocarburized 316L. (b) SAD pattern with cube-cube OR between fcc-MN and γN indexed. (c) The same SAD pattern as in (b) but with ω-carbides and γN indexed. (d) Dark-field image observed using the encircled diffracted beam of ω-carbide in (c)...... 124

Figure 4.10 (a) Bright field image. (b) ESI map of nickel. (c) ESI map of carbon. (d) EELS spectrum showing the nickel edge observed and (e) EELS spectrum showing the carbon edge observed...... 126

Figure 4.11 (a) Bright-field image obtained from an area containing less carbide. (b) ESI map of nitrogen. (c) EELS spectrum showing the nitrogen edge observed ...... 127

Figure 4.12 Results observed from two different 20C+20N processes: 20C+20N (H) and 20C+20N (L) (a) XRD; (b) Grazing angle XRD; and (c) ahkl plots ...... 130

Figure 4.13 SEM images taken from 316L treated by (a) 20C+20N (L) and (b) 20C+20N(H)...... 131

Figure 4.14 AES depth profiles measured from nitrocarburizied 316L with 20C+20N (L) (a and b) and 20C+20N (H) (c and d)...... 133

Figure 4.15 Surface hardness of 316L bulk sample treated with two different 20C+20N process...... 134

Figure 4.16 XRD results obtained from 3 different simultaneous nitrocarburzing processes ...... 135

xi Figure 4.17 Grazing angle XRD results obtained from (a) 20(N+C)-22.5 and (b) 20(N+C)-10 ...... 136

Figure 4.18 Cross-sectional images taken by confocal microscope from (a) 20(N+C)-22.5, (b) 20(N+C)-10 and (c) 20(N+C)-5 ...... 138

Figure 4.19 SEM images demonstrating the needle-like second phase after 20(N+C)-10 (a) 3000X and (b) 10000X ...... 139

Figure 4.20 AES profiles measured from (a) 20(N+C)-22.5, (b) 20(N+C)-10 and (c) 20(N+C)-5 ...... 140

Figure 4.21 Surface hardness of 316L bulk sample treated with three different 20(N+C) process...... 141

Figure 5.1 Solubility of Nitrogen in 316L stainless steel under paraequilibrium condition at different temperatures...... 146

Figure 5.2 CALPHAD predicted solubility of nitrogen in 316L stainless steel and AES measured surface nitrogen content on low-temperature nitrided 316L samples...... 147

Figure 5.3 (a) Plots of a111 and a200 vs. XN measured by AES; (b) Plots of aavg. vs. XN; (c) linear fitting of aavg. vs. XN and (d) linear fitting of strain-free aavg. vs. XN ...... 150

Figure 5.4 XRD measured lattice parameters (red circles) and “corrected” lattice parameters (green diamond) assuming a residual stress of 8GPa from bulk 316L nitrided in aN2 = 7400 at 440 °C ...... 154

Figure 5.5 XRD measured lattice parameters (red circles) and strain-free lattice parameters (green diamond) from 316 thin foil nitrided in aN2=7400 at 440 °C ...... 155

Figure 5.6 XRD measured lattice parameters (red circles) and “corrected” lattice parameters (green diamond) assuming a stacking faults density of 0.2, bulk 316L nitrided in aN2 = 7400 at 440 °C ...... 156

Figure 5.7 Correlation of lattice parameter expansion (defined as (ahkl-a0)/a0) vs. nitrogen content ...... 156

Figure 5.8 Normalized metal composition measured from (a) nitrided 316L and (c) nitrocarburized 316L. The corresponding nitrogen/carbon profiles are shown in (b) and (d), respectively...... 161

Figure 5.9 Schematics of stress measurements at (a) 0° tilt and (b) 25.2° tilt ...... 163

xii Figure 5.10 AES profiles measured from <111> and <100> grains, together with numerical simulation conducted by Prof. Kahn ...... 165

Figure A-1 Furnace setup to execute nitrocarburing experiment in urea ...... 172

Figure A-2 Thermal decomposition of Urea ...... 173

Figure A-3 XRD results from the three different nitrocarburizing processes in urea. ... 176

Figure A-4 XRD results from the bottom and top surfaces of sample U1...... 176

Figure A-5 Optical microscopy images observed from samples nitrocarburized in urea powders...... 177

Figure A-6 (a) XRD results from 316L samples with different surface finishing after nitrocarburized in NH3/C2H2/H2/N2 at 440 °C for 20 hours. (b) Enlarged segments of the XRD results in image (a) ...... 181

Figure A-7 Optical microscopy images observed from 316L samples nitrocarburized in NH3/C2H2/H2/N2 ...... 182

Figure A-8 AES depth profiles of carbon and nitrogen measured from sample MP-1μm (after nitrocarburized in NH3/C2H2/H2/N2 at 440 °C for 16 hours)...... 182

xiii Acknowledgements

I would like to express my sincere gratitude to my academic advisor, Prof. Arthur H.

Heuer, for his devoted guidance and support during this research. He is the greatest scientist and mentor I have ever met. I also want to thank Prof. James D. McGuffin-

Cawley, Prof. Frank Ernst, Prof. John J. Lewandowski and Dr. Farrel J. Martin for being my committee members.

I wish to express my special thanks to Prof. Harold Kahn for his patient training on the CVD furnace and AFM. More importantly, I gratefully appreciate the valuable discussions and advices provided by him. I really appreciate Prof. Gary M. Michal for his time and efforts on this project.

Special thanks are due to Dr. Reza Shaghi-Moshtaghin and Dr. Amir Avishai, for both their training on SEM and TEM and their discussions on results. I also want to thank

Dr. Wayne D. Jennings for his training on AES and XPS, Alan K. McIlwain for his training on XRD and nanoindenter, Dr. David B. Hovis for his discussion in my project,

Nanthawan Avishai for her help on TEM sample preparation and also for being a great friend, and Annette M. Marsolais for her help on AES.

I am grateful to the representatives of Swagelok Company, Dr. Sunniva R. Collins,

Mr. Peter C. Williams, Dr. Steven V. Marx and Mr. George R. Vraciu. I appreciate Mr.

Vraciu for providing me with the electropolished samples. Many thanks are also due to the previous and current members of the AHH research group, who make our group more like a family.

xiv I wish to express my special gratitude to my mother for her endless love and support in my life. At last, I want to dedicate my special thanks to my husband, Hongyi, for the efforts and encouragements he has made during the four years. Without his help, I could not even imagine to finish this dissertation.

xv Low-Temperature Gas-Phase Nitriding and Nitrocarburizing of

316L Austenitic Stainless Steel

Abstract

DANDAN WU

Low temperature paraequilibrium nitriding is an effective method to enhance surface hardness and corrosion resistance in austenitic stainless steels, provided that equilibrium nitride formation is suppressed. Following the standard double HCl “activation” procedure developed by Swagelok Company to remove the passivating Cr2O3-rich native oxide, nitriding was done in a gas mixture of NH3/H2/N2. Three processing parameters

(nitriding temperature, nitriding activity and duration) were controlled independently to understand both thermodynamic and kinetic aspects of the process.

Supersaturated nitrogen interstitials (7 ~ 25 at.%) were introduced into 316L stainless steel samples, which yielded a lattice expansion ranging from 1% to 10%. Room temperature ferromagnetism in expanded austenite in stainless steels was then induced due to the great increase in Fe-Fe interatomic distance. A combined of XRD, MFM and

EBSD study revealed that the minimum lattice expansion required for ferromagnetism is

~ 5%. A nitrogen content of ~ 14 at.% was estimated (by AES) as the threshold required

xvi for the paramagnetic-to-ferromagnetic transition. The correlation of lattice parameter expansion and nitrogen content indicates that transition from paramagnetic austenite to ferromagnetic austenite played a role in the highly distorted lattice parameters of nitrogen-enriched expanded austenite.

Orientation-dependent nitrogen surface concentration and case depth were investigated using EBSD orientation mapping and AES cross-sectional line scans. In particular, <100>-oriented grains demonstrated a higher surface nitrogen concentration and a deeper case depth as compared to <111>-oriented grains.

Three different scenarios of low-temperature gas-phase nitrocarburizing processes were designed and compared. Dual-layered expanded austenites were obtained. The concentration depth profiles of nitrogen and carbon atoms can all be described as an outer layer of nitrogen-enriched region, with carbon atoms accumulated at the diffusion front of nitrogen. The total case depth obtained is mainly determined by the diffusion time of carbon. Grazing angle XRD and TEM were employed to study the precipitates formed after nitrocarburizing.

xvii Chapter 1 Introduction

1.1 Background of expanded austenite

Austenitic stainless steels are widely used, economical, corrosion-resistant materials.

However, suffering from their low surface hardness and poor wear resistance, they have a

limited range of applications. Conventional thermochemical surface hardening techniques,

such as nitriding, carburizing and nitrocarburizing, can be employed to improve the

surface tribological properties of 316L stainless steel by surface precipitation.

Unfortunately, the corrosion resistance is sacrificed because of the depletion of chromium

(Cr) into the precipitates.

In the mid-1980s [1, 2], the discovery of “expanded austenite” (or so called S-phase)

demonstrated that it is possible to enhance the surface hardness of stainless steels without

deteriorating their outstanding corrosion resistance. The name “expanded austenite” was

first proposed by Leyland et al. [3] in 1993 to describe the surface scales formed after

low-temperature nitriding and carburizing processes. This name implied that the

interstitial nitrogen and carbon atoms did not change the original face-centered cubic

(FCC) structure of the alloys. The expanded austenite is a precipitate-free case with supersaturated nitrogen and/ or carbon interstitial atoms.

The absence of chromium nitrides or carbides can be explained by the concept of paraequilibrium[4,5]. At sufficiently low heating temperatures where a paraequilibrium condition is achieved, the diffusion of substitutional metal atoms is essentially suppressed,

1 whereas the diffusivity of nitrogen or carbon atoms is still substantial and they can be dissolved interstitially. For diffusion couples under paraequilibrium conditions, the chemical potentials of nitrogen or carbon atoms would equilibrate, but not for the substitutional species. Using the CALPHAD approach and assuming such a paraequilibrium scenario, G. Michal et al. [5,6] successfully predicted a supersaturated carbon solubility in 316L stainless steel up to 12 at%, matching well with the experimental results. A wide investigation on low-temperature nitriding of austenitic stainless steels indicated that the maximum nitrogen content obtained is as high as 25~30 at% [7-10].

Figure 1.1 shows an optical micrograph of an austenitic 316L stainless steel coupon after a low-temperature nitriding process [11]. The uniform surface scale was identified as a nitrogen-rich case. The featureless contrast implied that the case had a better etching resistance than the substrate to the etching reagent, which was 50% HCl + 25% HNO3 +

25% H2O.

Figure 1.1 Microstructure of 316L stainless steel nitrided at 703K for 20 hrs [11]

Figure 1.2 X-ray diffraction patterns of the ion-nitrided stainless steel at 673K in different ratios of N2 and [12] H2 (a) 1:100, (b) 1:9 and (c) 4:1

1.2 Crystal structure of expanded austennite

Typical X-ray diffraction (XRD) patterns observed after a low-temperature ion- nitriding treatment is demonstrated in Figure 1.2[12]. Two ssets of austenite peaks were obtained. Compared to the original austenite (γ) peaks, the expanded austenite (S1, S2) peaks shifted to the lower diffraction angles and indicated an expansion in lattice parameter due to the uptake of interstitial nitrogen atoms. Detailed examination of XRD spectra revealed that the expanded austenite had a distorted FCC structure, e.g. a111 < a200.

3 [13] This was considered in the literature [1] to be a combined effect of the residual stress

and the stacking faults (SFs).

The presence of a tremendous magnitude of residual stress in the expanded austenite

layer is widely reported, which originates from lattice and thermal mismatch between the

case and the core layer underneath. A compressive stress of 3.4 GPa at the surface was

reported by H. Kahn et al [14], which was measured from a carburized bulk 316L sample

employing a standard sin2ψ technique. T. Christiansen and MAJ. Somers [15] reconstructed the residual stress profile for 316L samples after low-temperature gas- phase carburizing, nitriding and nitrocarburizing based on X-ray analysis. As demonstrated in Figure 1.3, a compressive stress as high as 7~8 GPa was reported for the

-1/2 sample nitrided in nitriding potential KN (as defined in Equation. (1-1)) of 2.49 bar .

The largest compressive stress was not measured at the surface, which implied a cracked surface after nitriding.

= (1-1) ()

where and are the partial pressure of ammonia and hydrogen gases, respectively.

When a biaxial compressive stress exists in the case layer, the lattice perpendicular to

the surface is constrained by the core material, while the lattice parallel to the surface can

expand freely. Because of the anisotropic elasticity of 316L, the residual strain due to

Poisson’s effect is hkl-dependent. The relationship between the strain-free lattice

4 parameter ahkl and the apparent lattice parameter measured directly from XRD (ahkl,σ) is given by [14]:

, = (1-2) 1+21

where S1hkl is hkl-dependent X-ray elastic constants (XECs). The XECs are not available for expanded austenite, so values reported for austenitic alloys were used instead. XECs values calculated by H. Kahn [14] using a Dewit approximation are given in Table 1.1. It is clear that S1200 is much larger than S1111, which predicts a much larger residual stain in

(200) plane as compared to that in (111) plane, assuming a constant stress exists in both grains. A comparison between thin foil and bulk 316L saammples after low-temperature plasma nitriding clearly revealed the impact of residual stress on lattice parameter. Less scattered data points in Nelson-Riley plot were observed on the plasma nitrided 316L foil due to a smaller residual stress [13].

Figure 1.3 Profiles of residual stress profiles in AISI 316 oobtained by simulating X-ray diffraction pattern [15] . Heat treatment conditions: Carburizing: 793 K/ 2.5 h / 90% CO+ 10% H2. Nitridingg: 22 h / 718 K with -1/2 -1/2 the nitriding potentials KN = 0.293 bar and KN = 2.49 bar . Carburizing and nitriding: 793 K / 2 h / 30% CO + 70% H2 and 713 K / 23 h / KN = 1.14.

5 The presence of stacking faults in expanded austenite layer is considered as another

origin of the anisotropy in lattice parameter measurement from XRD. The correlation

13 between stacking fault density (α) and ahkl,α is given by [ ]

, = (1-3) 1+

where Ghkl is hkl-dependent stacking fault parameter and values for different crystal

planes are listed in Table 1.1. By modeling the XRD pattern observed from nitrided bulk

samples (Figure 1.4), Somers et.al. [16] proposed that the anomalous expansion in (200) planes can be explained by a combination of compressive stress and stacking faults.

However, only (111) and (200) peaks were considered in their simulation, which made their model debatable.

Figure 1.4 Modeling XRD of nitrided bulk samples [16]

6 [13, 14] Table 1.1 Calculated values of S1hkl and Ghkl

-6 -1 -2 hkl S1hkl (×10 GPa ) Ghkl (×10 ) (111) -990 -3.15 (200) -2419 6.89 (220) -1271 -3.45 (311) -1644 +1.25 (222) -990 +1.73 (400) -2419 -3.45

1.3 Diffusion of nitrogen in expanded austenite

A typical nitrogen profile in a low-temperature nitrided sample is shown in Figure

1.5[17], which can be described as a very slow decrease from the surface followed by a

sharp drop at the tail. The shape of the profile certainly deviates from the analytic

solution of Fick’s second law for concentration independent diffusion. Two models were

proposed to explain the concave shape of the diffusion profile: trapping-detrapping model

[18, 19] and concentration-dependent diffusion model [20, 22].

A qualitative model based on the assumption of Cr being the trapping site for nitrogen

atom was proposed by D. L. Williamson et al [18] to explain the plateau-shape of the

nitrogen concentration depth profile. In this trapping model, nitrogen can only diffuse further when all of the trapping sites are occupied. With such an assumption, the atoms diffuse later should always be located at the diffusion front, which is contradictory to the experimental observation using 15N and 14N isotopes [19] (as shown in Figure 1.6). Later the model was refined by taking detrapping into account [19].

7 Based on the Boltzmann-Matano diffusion equation, Mandl et al [20] calculated the

nitrogen diffusivity in austenitic stainless steels and a concentration dependence of the diffusivity is confirmed. T. Christiansen and M. A. J. Somer’s [21] determined the

concentration dependent diffusion coefficient of nitrogen by fitting thermo-gravimetric

data obtained from the denitriding process of thin initially N-saturated coupons. The

result is shown in Figure 1.7. For both 316L and 304 steels, their results demonstrated

that the diffusion coefficient of nitrogen first increases then decreases with nitrogen

content. For 316 steel at both temperatures (693K and 718K), the diffusion coefficient of

nitrogen reaches a maximum at approximately yN = 0.45. Their interpretation is that the

amount of available octahedral and tetrahedral sites will both reduce when too much nitrogen exists in the expanded austenite, which decreases the pre-exponential factor of diffusivity.

Following the numerical model developed by Sun and Bell [23] for the diffusion of

nitrogen in ferritic alloys, T. Christiansen and M. A. J. Somers [21] simulated the nitrogen diffusion profiles in austenitic stainless steels during gas-phase nitriding using a finite

difference method. The trapping model was also adopted in the simulation. In Figure 1.8,

KCr•N is the solubility product of Cr and N in the expanded austenite. They demonstrated

that in the condition of full solubility (KCr•N = ∞), a smooth transition from the plateau- like nitrogen curve to the substrate should be observed. On the other hand, with the assumption of zero solubility (KCr•N = 0), an abrupt drop in concentration should be

detected. For KCr•N = 0.0061, the simulated curve lay in-between the other two

conditions- the profile was predicted as an abrupt drop followed by a slowly fading tail.

8 Possible physical origins of the concentration-dependent diffusivity of interstitial

atoms in steel were proposed by F. Ernst et al [24], which included (1) the lattice

parameter expansion induced by interstitially dissolved atoms, which results in a reduction in the activation energy Q of diffusion; (2) the biaxial stress gradient resulting

from a gradient in interstitial concentration; (3) short circuit diffusion provided by the

dislocations generated by plastic deformation of the material; and (4) the trapping effect

from chromium.

Figure 1.5 Nitrogen concentration depth profile measured by GDOES, from an AISI 321 sample after low energy ion implantation. [17]

Figure 1.6 Nitrogen concentration depth profiles observed by nitriding under typical low energy ion, implantation conditions first with the 15N isotope and subsequently with the 14N isotope. The 14N isotope is not present at the diffusion front. [19] The profiles were measured by deuterium induced nuclear reaction analysis. [59]

Figure 1.7 Diffusivity of nitrogen calculated from fitting thermogravimetric data obtained from ntiridng and stepwise denitriding experiments of stainless steel thin foils. [22]

Figure 1.8 Simulated concentration dependent nitrogen diffusion profile for gas-phase nitrided AISI 316 assuming a constant surface concentration condition. [21]

1.4 Orientation-dependent case depth in low-temperature nitriding

Ozturk and Willamson [25] reported a non-uniform case layer thickness from grain to

grain after low-energy, high-flux nitrogen implantation into AISI 304 stainless steel. By employing a combination of step-wise Ar ion beam sputtering and XRD technique, they demonstrated that nitrogen diffuses deeper in the <200> oriented grains as compared to the <111> oriented ones. A study [7] on single crystal 316L austenitic stainless after ion implantation also demonstrated that nitrogen diffuses preferentially along [100]. A detailed EBSD analysis [26] on nickel base alloy (Inconel 690) treated by a low- temperature plasma assisted nitriding has revealed that a linear relationship exists

between the nitrided layer thickness and Ahkl, which is defined as:

ℎ +ℎ + = (1-4) (ℎ + +)

11 where h, k, l are the Miller indices.

A model was developed later by H. He et al. [27] to explain the observed orientation- dependent case depth. They proposed that the anisotropy in elasticity gave rise to an anisotropic stress and strain in the expanded austenite, which resulted in an orientation dependence. Other proposed mechanisms include an orientation-Lattice Parameter-

Temperature-alldependent ion penetration [25] and an ion irradiation effect on the diffusion of nitrogen [28].

(a)

Figure 1.9 (a) SEM cross-sectional micrograph demonstrating an orienttation-dependent case depth on a plasma nitrided Inconel 690 sample and (b) Local layer thickness e as a function of Ahhkl measured on the same sample. [26]

1.5 Effect of stress on solubilityt and diffusivity of interstitial atoms

An early study on the solubility of hydrogen atoms in a Pd/Ag alloy showed that the chemical potential of hydrogen atoms can be changed by residual stresses [28]. As mentioned previously, after low-temperature nitriding processes, enormous residual stresses (up to 8 GPa) were reported due to the large lattice parameter mismatch between

12 the expanded austenite and the original austenite in the substrate. Thus the stress effect on

the chemical potential and the solubility of nitrogen cannot be neglected. The effect of

stress on the chemical potential of nitrogen atoms can be estimated according to [21]

= +ln − (1-5)

where is the chemical potential at the standard state, R is the gas constant, T is the absolute temperature, aN is the activity of nitrogen, σ is the hydrostatic stress and is equal

to 2/3 of the biaxial stress measured by XRD, and is the partial molar volume of nitrogen atoms in the expanded austenite.

By applying an in-situ tensile stress on 316L stainless steel thin foil (50μm) during a

plasma carburizing process, H. Dong [30] studied the stress effect on the diffusivity of

carbon and the case depth (Figure 1.10). They reported that the thickness of the carburized layer was almost doubled on 316L when a tensile stress of 80 MPa was applied after 10 hours of plasma carburizing at 450 °C. However, the formation of

carbides was also promoted by the applied tensile stress.

Figure 1.10 Photograph of a plasma carburizing unit with in-situ stress applying component and the impact of tensile stress on the case depth of carbburized sample (from b to e). [30]

14 1.6 Room-temperature ferromagnetism induced by low-temperature

nitriding

By analyzing total-energy surfaces in moment-volume parameter space obtained from

energy-band calculations using a local-spin-density approximation, it has been predicted

[31,32] that face-centered cubic (fcc) Fe can exhibit either nonmagnetic (paramagnetic) or ferromagnetic behavior, depending on the Fe–Fe interatomic distances. Figure 1.11 [31] displayed the calculated total energy for fcc-Fe as a function of magnetic moment μB at several different Wigner-Seitz radius rWS. According to their calculation, ferromagnetic

fcc-Fe became more energetically favorable when rWS is increased. The calculation

performed by C. S. Wang et al.[40], using a so-called total-energy general-potential

linearized-augmented-plane-wave method, predicted that fcc-Fe should experience an

evolution from nonmagnetic (NM) to anti-ferromagnetic (AFM) then to ferromagnetic

(FM) state during the expansion of volume. As demonstrated in Figure 1.12, for fcc-Fe

the critical Wigner-Seitz radius for the transition from AFM to FM state was predicted to be 2.69 a.u. Their prediction was well supported by experiments. When the lattice

parameter of fcc-Fe was expanded beyond a certain level, ferromagnetism was detected

in both Fe films [33] and Fe particles [34]. By investigating the magneto-volume effect on

epitaxial γ-Fe ultrathin films and γ-Fe precipitates, W. Keune et al. [41] concluded that a

transition from antiferromagnetic low-spin fcc-Fe to ferromagnetic high-spin fcc-Fe

happened at a critical Wigner-Seitz radius of about 2.69-2.70 a.u, which corresponded to an lattice parameter expansion of about 5.9%~6.3% (as shown in Figure 1.13).

Figure 1.11 Total energy (relative to a minimum energy) as a function of magnetic moment for fcc-Fe showing the 3 different magnetic states [31].

Figure 1.12 Total energy per atom and magnetic moment as a function of volume for bcc-Fe and fcc-Fe [40]

Figure 1.13 Measured hyperfine-filed values near saturation (at 4.2K for γ-Fe particles in Cu100-xAlx, at 30K of 5 monolayer(ML) γ-Fe (001) on Cu (001) and Cu3Au(001), reespectively) versus Wigner-Seitz [41] radius γws/a.u.

17 Recently, ferromagnetic austenite has been observed at room temperature in

“expanded” austenite-- single-phase Fe-Cr-Ni stainless steels whose lattice parameters have been increased by the presence of large concentrations of interstitial nitrogen solutes

[25, 35-37].

Room temperature ferromagnetism in expanded austenite in stainless steels was

reported following both low-temperature plasma nitriding [25,35,36] and low-temperature

nitrogen ion implantation [37]. Figure 1.14 [36] demonstrates the AFM and MFM images

observed from a nitrided AISI316 specimen. A rough surface with high relief was revealed by the AFM topography image, which can be attributed to the high stress level

in the surface layer. A characteristic magnetic domain structure was revealed in the

MFM image, which represents the existence of a ferromagnetic state induced by nitriding.

More recently, magnetic patterning of plasma-nitrided 316L austenitic stainless steel was

accomplished using TEM grids as masks [38, 39], which could bring nitrogen-enriched expanded austenite into a range of new applications, including recording media, magnetic

sensors or magnetic separators. In contrast to nitrided austenite, carburized austenite, where the lattice parameter expansion is apparently limited to about 3% [6], has never

shown a paramagnetic-to-ferromagnetic transition.

Figure 1.14 (a) SPM (Scanning probe microscopy) and (b) MFM (Magnetic Force Microscopy) image observed from the same area of a AISI 316 sample after low-temperature nitriding. The scale is the same in the two images. [36]

1.7 Dilation of austenite lattice from nitrogen and carbon interstitials

According to Pauling [42], the tetrahedral-covalent radii of nitrogen and carbon are

0.0740 nm and 0.0771 nm, respectively. For a coordination number (CN) of 6, those

[43] numbers were suggested to be modified to rN = 0.0854 nm and rC = 0.0890 nm . It is noteworthy that, for both circumstances, the size of nitrogen is reported to be smaller than that of carbon. Based on a rigid sphere model, in a FCC structure, larger space exists in the octahedral interstitial site as compared to the tetrahedral site. The size of the octahedral site in 316L is estimated to be 0.05225 nm. Thus, both nitrogen and carbon atoms would expand the FCC lattice when being dissolved as interstitial atoms.

Empirically, a linear relationship is built up to describe the dilation of lattice from interstitial atoms of carbon and nitrogen, which has the format of

/ = +× /

19 where aN/C and a0 is the lattice parameter of the original and dilated unit cell, XN/C is

atomic fraction of interstitial atoms, and α is a fitting parameter. Considerable amount of

data are available for the α parameters measured for nitrogen and carbon interstitials.

Some of the data available for austenitic steels are summarized in Table 1.2. As

summarized, the α parameters reported for both nitrogen and carbon cover a fairly large

range, especially when the impact from residual strain is not removed from the lattice

parameter measurements.

Three of the previous studies considered both nitrogen and carbon. By studying 9

austenitic stainless steel samples containing about 0.2~1.5 at.% (C+N) interstitial atoms,

Ledbetter and Austin [43] found that the dilation from nitrogen (α=0.00086 nm/at.%) is

slightly larger than that from carbon (α=0.00078 nm/at.%). By reviewing existing lattice

parameter data of both pure iron-carbon and pure iron-nitrogen as-quenched martensitic

[44] samples, Liu et al. established the dependence of the lattice parameter on both XN and

XC in the retained austenite phase by a least-square fitting method. Without considering

the impact from residual stress, the fitting parameter α obtained for nitrogen

((0.00093±0.0006) nm/at.%) is smaller than that for carbon ((0.00105±0.0002) nm/at.%).

[44] In the same study , the dependence of lattice parameter on XC for strain-free austenite

was derived from data measured from fully austenitic specimens at elevated temperatures,

which yielded a α value of 0.00075 nm/at.%. Data taken from fully austenitic Fe-N

specimens at room temperature were used to estimate the dependence of strain-free lattice parameter on XN and α=0.0008 nm/at.% was obtained.

20 In the field of expanded austenite, MAJ Somers’s group [16] reported the only data on

the dilation of austenite lattice for both nitrogen and carbon interstitials. The results are

shown in Figure 1.15 (a). They plotted the lattice parameter as a function of yN/yC, which

was defined as the numbers of interstitial atoms per metal atom. Samples they used for

the nitriding and carburizing experiments were thin 316 foils and they claimed that the

sample were stress-free after treatment. Apparently, the two fitted straight lines for

nitrogen and carbon are parallel to each other, which implies that nitrogen and carbon are

equally strong austenite dilators. However, the intercept of the straight line fitted for

nitrogen with yN =0 is larger than the lattice parameter of non-treated 316, which is

difficult to interpret. By re-plotting their lattice parameter data as a function of atomic

fraction (XN/XC, as shown in Figure 1.15 (b)), this issue was avoided. Therefore the conclusion is changed to that nitrogen is a much stronger dilator for 316 as compared to

carbon, based on the much larger α parameter (0.0010 nm/at.%) observed for nitrogen.

This value is also larger than those reported by Ledbetter [43] and Liu [44].

21 (a)

Figure 1.15 (a) lattice parameter a of expanded austenite as a function off (a) numbers of interstitial nitrogen or carbon atoms per metal atom (yN or yC) and (b) concentration of interstitial nitrogen or carbon (XN or [16] XC).

22 Table 1.2 Survey of α parameters in Vegard’s law for nitrogen and carbon

Strain-free α (X) α (X) Species Alloy Technique Range Ref. (×10-4

-4 nm/at.%) (×10 nm/at.%)

γ-Fe 2-9 at% 10.5 7.5 Liu Cheng[44]

316L XPS 0-12 at% 10 6.8 Kahn[14]

304 0-1.5 at% 7.83 -- Ledbetter [43]

316L TGA 0-14 at% -- 6.8 Somers [16]

γ-Fe 5-11 at% 9.3 8.0 Liu Cheng[44]

310 EPMA 0-38 at% 13 -- A. Saker [45]

304 0-1.5% 8.61 -- Ledbetter [43] N

316L TGA 14-38 at% -- 10 Somers [16]

Fe-18Cr- 0.3-32 12Ni-N NRA 11 -- K. L. Dahm [46] at% coating

23 1.8 Precipitate formation and stability of expanded austenite

According to the experimentally observed Lehrer diagram (Figure 1.16) [47] for Fe-N system, the stable phases produced by a gas-phase nitriding process depends on both nitriding temperature and nitriding potential KN (Equation. (1-1)). When nitriding is

performed at a relatively high temperature or with a larger nitriding potenital, the two predicted nitride phases are of γ’-Fe4N or ε-Fe2-3N. When austenitic stainless steels were

subjected to traditional nitriding processes, γ’-Fe4N or ε-Fe2-3N and CrN phases were

observed in the compound layer [48]. After AISI 316 flakes was nitrided in a gas mixture

[49] of 9 vol.% N2 and 91 vol.% NH3 at 718K (445 °C) , coexistence of expanded austenite

[50] CrN, γ’-Fe4N and ε-Fe3N was revealed by XRD analysis. TEM study on plasma

assisted nitrided austenitic 316L stainless steel demonstrated that the formed surface

layer is composed of a single phase, which has a simple cubic (SC) lattice with a lattice

constant of 0.378 nm. The transition from a FCC lattice to the SC lattice was verified by

the presence of forbidden diffraction reflections in the SAD (selected-area electron

diffraction) patterns (as shown in Figure 1.17).

Figure 1.16 Lehrer diagram, which predicted the stable phases for Fe-N systwm as a function of temperature at nitriding potential used. [47]

Figure 1.17 Selected-area diffraction patterns from a nitridied 316 Stainless Steel. Forbbiden reflections were detected from both [100] and [211] zone axis. [50]

25 As mentioned previously, the expanded austenite is metastable, so it will decompose into stable phases during annealing. The isothermal annealing experiments done by

Somers’ research group [51] indicated that the expanded austenite was more stable in 316L stainless steel than in 304L. They also found out that the decomposition products after annealing were different. The expanded austenitte decomposed into CrN and austenite in

316L but CrN and ferrite in 304L. Inn situ heating TEM experiments conducted by Li et al. [52] revealed structural evolution during annealing of expanded austenite. The SAD patterns (Figure 1.18) taken at different aging temperatures clearly reveals the decomposition of expanded austenite into CrN and austenite by the gradual splitting of the diffraction spots.

Figure 1.18 SAD patterns tanken from an expanded austeniite region during in-situ annealing at (a) 0 °C; (b) 440 °C; (c) 500 oC and (d) 600 oC. [52]

26 1.9 Nitrocarburizing processes

A hybrid process of low-temperature nitriding and carburizing has become one of the most important research hot-spots recently [53-57]. By diffusing nitrogen and carbon atoms either simultaneously or consecutively, a duplex case with both a nitrogen-supersaturated layer and a carbon-supersaturated layer can be developed. Figure 1.19 shows a hardness- depth profile measured on a gas-phase nitrocarburized 316L specimen [11]. The sample was first carburized in infinite carbon activity (provided by 100% CO gas, assuming no soot formation) for 4 hours and then nitrided in infinite nitrogen activity (provided by 100%

NH3) for 18.5 hours. Comparing all of the profiles together, it is found that the nitrided case was harder but thinner than the carburized one, given the same duration. Excitingly, the nitrocarburized treatment combines the advantages of both processes and shows a way of further tailoring the mechanical properties of the expanded austenite.

Typical depth profiles of nitrogen and carbon in the nitrocarburized layer are shown

in Figure 1.20 [53]. Unexpectedly, the carbon diffused beyond the nitrogen layer and

showed a maximum concentration in front of the diffusion tail of nitrogen. Further

investigation demonstrated that the top scale is always nitrogen supersaturated, regardless

of the diffusing sequence of species. The uphill diffusion of carbon in the nitrided

austenite layer was described as a “push” effect and was explained by a “trapping”

mechanism of nitrogen atoms [55]. An alternative explanation was addressed in the view

of chemical potential. Based on the CALPHAD modeling, Gu [56] calculated the change

of chemical potential of carbon with nitrogen content in the austenite phase and vice versa. The results indicated that the existence of nitrogen can greatly increase the chemical potential of carbon. Similarly, carbon can also greatly increase the chemical

28 potential of nitrogen when they coexist in 316L. By defining the “effective” carbon content as Xc with the same chemical potential level when no nitrogen presents and considering the concentration dependent diffusion for both interstitials, they simulated the depth profiles in a nitrocarburizing process and explained the origin of the uphill diffusion behavior of carbon (as shown in Figure 1.21).

29 t = 15 hour

C (eff.) (at%) N , C X C (act.)

N (act.)

Depth (µm)

Figure 1.21 Simulated concentration-depth profiles C(act.), C(eff.) and N (act.) after a plasma nitrocarburizing process (15 hours plasma carburizing followed by 15 hours plasma nitriding). [56]

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35 Chapter 2 Experimental Methods

2.1 Design of low-temperature gas-phase nitriding process

Plasma nitriding, ion implantation and gas-phase nitriding are the three commonly

used techniques to produce nitrogen-supersaturated expanded austenite layer. Compared

to other processes, the gas-phase nitriding process has the following advantages: (a) The

nitriding activity can be easily controlled by changing the content of the gas mixture,

which makes thermodynamic and kinetic studies of the process possible; (b) In virtue of a

uniform temperature distribution, the formed case is homogeneous; (c) Gas-phase ensures

good coverage over specimens with complicated geometries and fine features.

Nevertheless, for a gas-phase nitriding process, a separate pretreatment is required to

remove the chromium oxide passive film on the surface, which is considered as a

diffusion barrier for nitrogen atoms. In the present study, the double HCl activation procedure developed by the Swagelok Company [1] is used to dissolve the surface film of

chromium oxide. It has been proven that this activation step can successfully make the

surface of the treated article “transparent” for carbon diffusion [2].

All gas-phase nitriding and nitrocarburizing processes were performed at Case

Western Reserve University employing a customized chemical vapor deposition (CVD)

furnace (as shown in Figure 2.1). Multiple gas delivery channels were equipped to the

furnace, including HCl, NH3, C2H2, H2, N2 and CO. The gas flows were precisely

controlled by pre-calibrated mass flow controllers (MFCs). All samples were treated

under a total pressure of 1 atmosphere. All processing parameters (including heating rates,

36 soaking temperatures and flow rates of gases) were programmed into “recipe” files before

running the furnace.

A schematic of gas-phase nitriding process used in this work is shown in Figure 2.2.

The double activation were done at 325 °C in 10 vol. % HCl (0.2 L/min HCl

and1.8L/min N2). For the low-temperature nitriding segments, the effects of three

different variables on low-temperature gas-phase nitriding of 316L alloy are designed to

be investigated independently: temperature, nitriding activity and duration.

At a certain temperature, the activity of the nitriding process aN2 is determined by the

partial pressures of NH3 and H2 gases according to the following equation:

=[× ] (2-1)

where K is the equilibrium constant for the dissociation reaction of NH3 (Equation 2-2).

PNH3 and PH2 are the partial pressure of NH3 and H2 gas, respectively.

1 3 → + (2-2) 2 2

Detailed processing parameters employed for the investigation of temperature and

activity are listed in Table 2.1 and 2.2, respectively. For the study of duration, the

o experiments were all done at 440 C in an activity of aN2=7400. Three different durations

were chosen: 5 hour, 20 hour and 80 hours.

Figure 2.1 Picture of the CVD furnace employed

Figure 2.2 Schematic of procedures of low-temperaturee nitriding process

38 Table 2.1 Processing parameters to study the effect of nitriding temperature

Table 2.2 Processing parameters to study the effect of nitriding activity

39 2.2 Design of low-temperature gas-phase nitrocarburizing process

There are three different processing scenarios which can produce a case containing both nitrogen atoms and carbon atom in 316L samples. The schematic of those 3 scenarios of nitrocarburzing is illustrated Figure 2.3. The same as the low-temperature gas-phase nitriding experiments, the activation procedure employed was double activation in HCl at 325 oC. Three gas-phase nitrocarburizing processes were designed to study the effect of different diffusion sequence. The processing parameters, including temperature, activity (for both nitriding and carbburizing) and duration, are listed in Table

2.3. For all experiments, the source of nitrogen is NH3 and the source of carbon is CO.

For the simultaneous nitrocarburizing process, gas mixtures of NH3/CO/H2/N2 were employed. The nitriding activities were calculated according to Equation 2-1. The carbon activities were all calculated by Factsage softwarre.

Sequential Simultaneous

20C+20N 20N+20C 20(N+C)

Figure 2.3 Schematic of three scenarios of low-temperature nitrocarburizing processes

40 Table 2.3 Processing parameters for nitriding, carburizing and three different nitrocarburing processes

The effect of ntiriding activities in the design of carburizing followed by nitriding and simultaneous nitrocarburizing was also investigated. Detailed processing parameters employed were listed in table 2.4 and 2.5, respectively.

41 Table 2.4 Designed parameters of two 20C+20N nitrocarburizing experiments

Table 2.5 Designed parameters of three 20(N+C) nitrocarburizing experiments

42 2.3 Characterization techniques

X-ray diffraction (XRD), scanning electron microscopy (SEM), electron backscatter

diffraction (EBSD), magnetic force microscopy (MFM), Auger electron spectroscopy

(AES), Microhardness tester and transmission electron microscopy (TEM) were used to

characterize the samples after low-temperature gas-phase nitriding and nitrocarburizing

process.

2.3.1 X-Ray Diffraction

In the present work, a Scintag X-1 advanced X-ray diffractometer with Cu radiation

(λCuKα1 = 0.154056 nm) was employed for phase identification, lattice parameter

measurements and residual stress estimation.

The lattice parameter measurements are done with a Bragg-Brentanon geometry, or θ-

2θ geometry. With this geometry, the penetration depth of Cu Kα X-ray is estimated to be

about 2μm ~7 μm (95% integrated intensity), for the 2θ angle selected (from 35o to 150o).

Grazing XRD mode was employed to examine the possible second phases formation after

low-temperature nitriding and ntirocarburizing processes. A grazing angle of 1o is normally used and the corresponding information depth is estimated to be about 200 nm assuming 95% integrated intensity.

The measurement of lattice constant a by X-ray diffraction (XRD) technique is an indirect approach. For materials with a cubic structure, like 316L stainless steel studied in this thesis, the lattice parameter is given by

43 =(ℎ + + ) (2-3)

[3] in which the inter-planar spacing dhkl can be calculated from the Bragg equation :

=2 (2-4)

where λ is the wavelength of the incident X-ray and θhkl is the diffraction angle. The peak

position is evaluated by curve fitting method according to the Pearson VII function [4], which is given by:

() = [1 + 4(2 −1)( )] (2-5)

where H represents the height of the peak, W is the FWHM of the profile and m is a shape

parameter. The shape of the simulated peak becomes Cauchy, modified Lorentzian or

Gaussian shape when m is equal to 1, 2, or infinity.

The measurement of residual stress by X-ray diffraction method is based on an

assumption of a linear elastic relationship between stress and strain. Lattice strain (by the

definition of engineering strain) can be measured by means of the shift of the diffraction

angle [5]:

= =− − (2-6)

in which is the interplanar spacing measured in the direction defined by ϕ and ψ (as

shown in Figure 2.4), and d0 is the so called “strain free” d-spacing. The residual stress in

the nitrided samples is assumed to be a plane stress, i.e., σ1=σ2=σ and σ3=0. Then the

residual stress can be determined by the following equation [4,5]:

44 =( + ) + (2-7)

hkl hkl S1 and S2 are plane-specific elasticity constants, called X-ray elastic constants

(XEC). XECs can be calculated from microscopic elasticity constants. Values used in this

thesis were were reported by H. Kahn in ref. [6] and list in Table 1.1.

According to equation 2-6 and 2-7, we have[4,5]

= ( +) + (2-8)

[4,5] Taking the first derivative of with respect to yields :

= = (2-9)

where m is the slope of the d- plot. The error was found to be negligible when

is replaced by , which is the intercept of the d- plot. Thus the stress can be

finally measured according to [4,5]

= (2-10)

In theory, a high angle peak should be chosen for stress measurement to achieve a

higher accuracy. For austenitic stainless steel, the (420) peak is recommended [5].

Figure 2.4 Schematic of the plane stress elastic model [5].

2.3.2 Auger Electron Spectroscopy

Accurate measurements of the concentration depth profiles of nitrogen and carbon are important and challenging. Both structure and properties of the expanded austenite

greatly depend on the amount of interstitial atoms and the shape of the depth profiles.

AES cross-sectional line scans were conducted to portray chemical properties of 316L

samples after low-temperature nitriding and nitrocarburizing processes.

Before the analysis, cross-sectional samples were carefully prepared and polished

down to 1μm diamond suspension. Compression mounting was employed before

polishing to preserve edge information. Before being brought into the UHV chamber of

the AES system, each sample was mechanically removed from the epoxy mount and

ultrasonically cleaned. Before acquisition, about 50 nm in thickness were sputtered away by plasma Ar+ ions to remove any undesired hydrocarbon contamination and the oxide

layer. Dynamic Ar+ sputtering was carried out to avoid oxygen accumulation from the

46 chamber during the acquisition, because the overlap of oxygen and chromium peaks

could introduce uncertainty to the examination of chromium. 64 sampling points with

equal spaces were set along a straight line perpendicular to the surface of the tested

sample, across from the expanded austenite layer to the core material. The acquired

results were analyzed using the Multipak software. The concentration profiles of nitrogen

were obtained by quantifying the Auger spectra.

For the detection of carbon by AES, a calibration procedure was conducted by a

former MS student Joshua Katz [7] to achieve an accurate quantitative evaluation. By

comparing the AES measured carbon concentrations in a set of Fe-C alloys with known

carbon contents, an empirical relationship (Equation 2-11 [7]) was built-up to convert the

. directly measured carbon concentration ( ) into the actual carbon concentration

. ( ).

. . = 0.566 − 0.986 (2-11)

In equation 2-11, carbon concentrations are in atomic percentage (at%). All analysis

was done in a continuous sputtering mode; otherwise the accumulation of carbon with

time will cause great error in the quantification result.

For the detection of nitrogen by AES, a calibration procedure was performed on two

alloy standards with certified nitrogen content (provided by Swagelok Company). Ar+ sputter was employed to remove surface oxide/contamination. Spectra acquired at different sputtering depth from sample NSC-4 are shown in Figure 2.5. Clearly, the oxygen peak becomes negligible after a 900 nm sputtering. The certified composition of those two samples and the corresponding AES detected concentrations after different

47 sputtering depth are compared in Tables 2.6 and 2.7. The results indicate that (1) the existence of oxygen does bring great error in the estimation of Cr by AES and (2) given that all oxide/contamination was removed by sputtering (the last row in both tabless), the error in estimating concentration by AES is about 1 at%, including nitrogen and carbon.

Figure 2.5 AES spectra observed from sample NSC-4 at different sputtering depths

Table 2.6 Comparison between the certified and AES measured composition on Sample B.S. 81N

48 Table 2.7 Comparison between the certified and AES measured composition on Sample NSC-4

The advantages of employing scanning Auger specttroscopy on cross-sectional samples to measure the concentration depth profiles can be summarized as the following:

(1) Comparing to the traditional depth profiling method from plan-view, the data acquisition is much more efficient. The case depth on nitrocarburized samples can reach

20~30 μm. A line scan on cross-sectional sample containing 64 data point normally takes

2~3 hours. (2) Measurements can be done on specified locations with high spatial resolution. After sputtering by Ar ions, the case and the core regions can be differentiated due to a different sputtering resistance to the AAr ions. Then lines can be setup from different grains or other areas of interest.

2.3.3 Scanningn Electron Microscopy

The imaging ability of SEM was used to determine the thickness of case layer and crack formation after nitriding and nitrocarbburizing, from cross-sectional samples.

Regular procedures of preparing cross-sectional samples were conducted, which includes sectioning, hot mounting, rough grinding and final polishing. The finest diamond suspension used was 1μm. An etchant of 50 vol.% HCl+25 vol.% HNO3+25 vol.% H2O was used to reveal the boundary between the case layer and the core material.

49 2.3.4 Electron Backscatter Diffraction

EBSD technique was employed to determine the crystallographic orientation on

nitrided samples, from both plan-view and cross-section. The detection was done by a

state-of-the-art Nordlys II EBSD detector along with a high-resolution scanning electron

microscope (FEI-xT Nova Nanolab 200).

Since the information in the EBSD analysis comes from a very shallow depth (~50

nm) of the sample, sample preparation is critical to achieve high quality EBSD

orientation maps. For plan-view samples, a slight etching was performed (in a solution of

50 vol.% HCl+25 vol.% HNO3+25 vol.% H2O) to remove any undesired contamination.

For cross-sectional specimens, careful mechanical grinding and polishing steps were

executed first to prepare a smooth surface finishing. The grinding step was done through

a series of Buehler silicon carbide papers (P400, P800, P1200, P1500, P2500). Then a

delicate polishing was conducted using 3 μm, 1μm and 0.25 μm diamond suspensions. To

remove the damage layer resulting from mechanical polishing, two methods were

employed. (1) Ar+ sputtering in AES system to remove about 50 nm of material and (2)

chemical etching in a solution of 50 vol.% HCl+25 vol.% HNO3+25 vol.% H2O. The quality of the EBSD map is proven better using the second method. To avoid drifting during data acquisition, super glue instead of carbon tape was used to attach the sample to the SEM stage.

50 2.3.5 Magnetic Force Microscopy

MFM was used to reveal the existence of any ferromagnetic phase after nitriding. The equipment used was a Dimension 3100 (Veeco Digital Instruments) multi-mode scanning probe microscope. The scan was done in a non-contact mode with a cobalt-coated silicon tip. Both topographic and phase images were obtained at the same time. Scans were performed on both plane-view and cross-sectional samples after nitriding. Large indents were made as markers for the purpose of locating areas studied previously by EBSD.

2.3.6 Microhardness tester

The hardness of both nontreated and heat treated (by both low-temperature nitriding and nitrocarburizing processes) 316L stainless steel samples was measured by a Buehler

Micromet Microhardness tester with a Vickers diamond pyramid indenter. The measurements were done on plan-view samples and the load used was 25 gram force, which makes a large enough indentation for accurate measurement of the diagonals. With a load of 25gf, if the hardness of the treated sample reaches HV1000, the diagonal of the indentation is about 7 μm. The hardness value was averaged from at least 10 measurements from the same surface for each sample.

2.3.7 Transmission Electron Microscopy

TEM (FEI Tecnai F30 and Libra 200 FE) was used to investigate both the structural and compositional properties of the nitrided sample. A combination of bright field imaging, dark field imaging and selected-area diffraction (SAD) technique was mainly

51 used to reveal the microstructure of both the matrix (expanded austenite) and second phases. X-ray energy-dispersive spectroscopy (XEDS) and electron energy-loss spectroscopy (EELS) techniques were employed to investigate compositional properties.

TEM sample preparation was accomplished by a dual beam focused ion beam system.

JEMS software was employed to simulate diffraction patterns with given crystal

structures.

52 References

[1]. P. C. Williams and S. V. Marx: ‘Selective case hardening processes at low-

temperature’, US patent no. US 6165597A

[2]. G.M. Michal, F. Ernst, H. Kahn, Y. Cao, F. Oba, N. Agarwal, A.H. Heuer, Acta

Materialia, 2006, 54, 1597-1606

[3]. C. Suryanarayana, M. Grant Norton, X-ray Diffraction: A Practical Approach,

1998, New York, Plenum Press

[4]. Bob B. He, Two dimensional X-ray Diffraction, 2009, New Jersey, Wiley

[5]. Prevéy, Paul S. “X-ray Diffraction Residual Stress Techniques,” Metals

Handbook. 10. Metals Park: American Society for Metals, 1986, 380-392

[6]. H. Kahn, G. M. Michal, F. Ernst and A. H. Heuer: Metallurgical and Materials

Transactions A, 2009, 40A, 1799-1804

[7]. J. Katz, “Low-temperature Carburization of Ferritic Stainless Steels”, Master

Thesis, Case Western Reserve University, 2009.

53 Chapter 3 Low-Temperature Gas-Phase Nitriding

3.1 Effect of processing parameters

3.1.1 Effect of nitriding temperature

In the present study, four different temperatures (350 °C, 420 °C, 440 °C and 450 °C)

were selected to study the effect of temperature on low-temperature gas-phase nitriding

process. The compositions of the NH3/N2/H2 gas mixture were adjusted to provide an

equal nitriding activity (aN2 = 7400) at different temperatures. XRD results obtained from

samples nitrided with those 4 temperatures are shown in Figure 3.1. The XRD result

taken before nitriding is also given as a and labeled as NT in the image. All peaks were

indexed assuming a face centered-cubic (FCC) structure. Due to the selection rule for

FCC lattice, only peaks with their Miller indices h, k, l all odd or all even were detected.

Generally, only austenite peaks (expanded austenite and/or original austenite) were observed on those 4 samples. Peaks from the “case” layer all shifted towards lower diffraction angles and the peak shift was generally larger when the sample was treated at

a higher temperature. For the sample treated at 350 °C, diffraction peaks from the core

region were also detected. This indicated a shallow case layer. The penetration depth of

Cu Kα X-ray is about 2 μm at a diffraction angle of 2θ = 43.7° (the peak position of (111)

plane for 316L). This means that the case depth is even smaller than 2 μm after the

sample was treated at 350 °C for 20 hours.

54 Furthermore, diffraction peaks from the expanded austenite were broadened. For

example, the full-width-at-half-maximum (FWHM) of the (111) peak was 0.49° from the

sample treated at 440 °C for 20 hrs. On the other hand, the FWHM of the austenite (111)

peak was only 0.13°. The broadening of the diffraction peaks originated from (1) a highly

defective microstructure induced by the nitriding process and (2) a possible concentration

gradient in the X-ray penetration volume. Also resulting from the highly defective structure, the intensities of diffraction peaks at higher diffraction angles were greatly

reduced. When a nitrogen concentration gradient does exist, the shape of the peak is

asymmetric.

The lattice parameter of each crystal plane was calculated according to equation 2-3

and 2-4. The calculated values are plotted in Figure 3.2. ahkl calculated from seven

different X-ray peaks were included for all four samples. The straight lines connecting each data point in the figure are only guidelines for the eye. Clearly, different from the non-treated 316L stainless steel, which has an ideal FCC structure, the expanded austenite shows a distorted FCC structure. The anisotropies in lattice parameter are fairly substantial, especially for the two samples treated at higher temperatures (440 °C and

450 °C). Generally, the following relationship exists:

a200 > a311 > a111.

The average lattice parameter and the corresponding lattice parameter expansion for

each of the nitrided sample are listed in Table 3.1. Here the average lattice parameter for

each sample is defined as the arithmetic mean of the seven data points shown in Figure

3.2. Using this method, the lattice parameter of the non-treated sample was calculated to

55 be (0.3600±0.0002) nm. Accordingly, low-temmpperature nitriding substantially expands the austenite lattice. The unit cells were expanded by about 3% ~ 6% at the surface, depending on the nitriding temperature employed. For the two samples treated at higher temperatures (440 °C and 450 °C), the largely scattered ahkl data points yield a higher standard deviation.

Figure 3.1 XRD results of 316L stainless steel samples treated at different temperatures

Figure 3.2 Lattice parameters ahkl from samplees nitrided at different temperatures

Table 3.1 Calculated lattice parameters and lattice parameter expansions of samples nitrided at different temperatures

One of the origins for the scattering pattern oof the “apparent” lattice parameters is the biaxial residual stress in the expanded austenite layer. The residual stress arises from the resistance of core material to the expansion of the surface layer. Due to the

57 anisotropic X-ray elastic constants (XECs) of different diffraction crystallography planes,

the “residual strains” found on different planes are different. Due to lower stress levels in

the sample nitrided at lower temperatures (350 °C and 420 °C), the measured data points

are much less scattered as compared to those two samples nitrided at higher temperatures

(440 °C and 450 °C).

The residual stress in the nitrided samples was measured using a standard XRD sin2ψ technique. Theoretically, diffraction peaks at large angles are preferred to minimize the error in d-spacing measurements and hence stress measurements. However, the intensities of peaks at large diffraction angles were fairly low due to the defective structure of the expanded austenite. In addition, the peak intensity also decreases strongly with increasing tilt angle ψ. Therefore (311) was chosen to estimate the residual stress at the surface.

2 The relationships between lattice parameters a311 and sin ψ for samples treated at 350 °C

and 420 °C are plotted in Figure 3.3. The stress was calculated according to Equation 2-

10. The X-ray elastic constant S2hkl was calculated to be 0.013821 for (311) planes [H.

Kahn]. Stresses of -1.4 GPa and -1.7 GPa were estimated for the two samples treated at

350 °C and 420 °C, respectively. The minus sign indicates that the residual stress in the

case layer is compressive.

2 Figure 3.3 a311 versus sin ψ from samples nitrided at 350 °C and 420 °C

Figure 3.4 X-ray scans from the sample nitrided at 440 °C in aN2 = 7400 at two ψ tilts

59 For the two samples treated at 440 °C and 450 °C, the same method was used to

estimate the stress in the case layer. However, splitting of the (311) peaks was observed

during sample tilting, which makes it impossible to calculate the stress based on the XRD

sin2ψ technique. Figure 3.4 presents the XRD scans obtained from the sample nitrided at

440 °C for 20 hours at two different tilts: ψ = 0° and ψ = 30°. With a 30° tilt, the (311) peak splits into two peaks, one at 2θ = 82.7° and the other at 2θ = 85.9°. Two different lattice parameters of a311 can be calculated from those two peaks: 0.3869 nm and 0.3752

nm. A hypothesis is proposed to explain the peak splitting during tilting, which will be

discussed in Chapter 5.

Figure 3.5 presents cross-sectional SEM micrographs of the four samples nitrided at

those four temperatures. The original surfaces were all located at the bottom of the

images. The expanded austenite layer, or the case, is characterized as a “smooth” layer covered on the heavily etched core. It can be seen that the nitrided layer demonstrated a better resistance to the etchant used (50 vol.% HCl + 25 vol.% HNO3 + 25 vol.% H2O).

Actually, grain boundaries in the case layers were also revealed by the etchant, such as in

(b) and (d). In both images, the grain boundaries extend from the core part, which

indicates that no phase transformation or grain refinement occured. The effect of temperature on case depth is clearly revealed by those 4 photographs. A faster diffusion

of nitrogen at a higher temperature yields a much deeper case. The case thickness formed

after 20 hours of nitriding at 450 °C is 10~15 μm, depending on grain orientation. In

contrast, the case thickness produced in the same duration at 350 °C is only about 1 μm.

However, crack formation (Figure 3.6) was also promoted when nitriding was done at

high temperatures.

Figure 3.5 Cross-sectional SEM images of 316L stainless steel samples nitrided at different temperatures. The nitriding activity for all samples was set the same, aN2 = 7400.

Figure 3.6 SEM image revealing crack formation on 316L stainless steel samples after nitriding at 450 °C, aN2 = 7400.

61 As demonstrated in Figure 3.7, nitrogen diffusion profiles exhibit concave shapes with sharp diffusion tails at all temperatures. The deviation from an erroor function profile is related to the concentration-dependent diffusion behavior of nitrogen atoms in 316L stainless steel. Given the same nitriding activity, the surface concentration of nitrogen is relatively higher when the sample is treated at higher temperature. However, the difference is quite small. For the sample treated at 350 °C, the concentration of nitrogen at the surface was about 14%. When the temperature wwas raised to 450 °C, the concentration of nitrogen at the surface was cloose to 17%. In contrast, the case depth shows a quite heavy dependence on nitriding temperature. A case of only ~1μm was produced when the sample was treated at 350 °C for 20 hours. When the sample was treated at 450 °C for 20 hours, the yielded case depth was about 14 μm.

Figure 3.7 Nitrogen concentration profiles detected from samples nitrided at different temperatures

62 3.1.2 Effect of nitriding activity

The effect of nitriding activity was also studied by XRD. The results for the samples

nitrided with 6 different activities are shown in Figure 3.8, together with the result from

the non-treated 316L stainless steel (labeled as NT). It is clear that, at the same processing temperature (440 °C), the peak shift of the expanded austenite increases with increasing nitriding activities. For the two samples treated at lower nitriding activities

(aN2 = 200 and aN2 ≈ 1), peaks from the core material were detected. Again, this is

because the case depth produced was smaller than the X-ray penetration depth.

The relationship between the lattice parameter and nitriding activities is better revealed in Figure 3.9, in which the lattice parameter (ahkl) of each peak is calculated and

plotted. (The lines connecting each data point are only guidelines for the eye.) The mean

lattice parameters and the corresponding lattice parameter expansions induced by

different nitriding activities are summarized in Table 3.2. These results quantitatively

confirm the conclusion drawn from peak shifts: the incorporation of interstitial nitrogen

atoms substantially expands the austenite lattice and this lattice parameter expansion

increases with increasing nitriding activity. By applying the highest nitriding activity (aN2

= 4×109), the lattice parameter expansion was calculated to be about 10%. Besides the

lattice parameter expansion, the variation in the apparent lattice constant also changes

with the nitriding activities. These differences are enormous when the sample was

nitrided with medium nitriding activities.

Figure 3.8 XRD results of 316L stainless steel samples treated at different activities

Figure 3.9 Lattice parameters ahkl from samples nitrided at 6 different activities

64 Table 3.2 Calculated lattice parameters and lattice parameter expansions of samples nitrided at different temperatures

Residual stress in the case layer was also measured on samples treated with different nitriding activities. The stress was determined for only 3 of the 6 samples, in which aN2 =

9 2 4×10 , aN2 = 200 and aN2 ≈ 1, respectively. The lattice constant a311 as a function of sin ψ are plotted in Figure 3.10. The residual stress was calculated to be -0.46 GPa (aN2 ≈ 1), -

9 1.16 GPa (aN2 = 200) and -1.57 GPa (aN2 = 4×10 ), respectively. The minus sign indicates

that all residual stresss measured from nitrided bulk 316L samples are compressive. For

the sample treated with an activity of 4×109, a part of the residual stress was released by

crack formation on the surface, so the measured value was relatively small. For the other

5 3 sample (aN2 = 1.8×10 , aN2 = 7400 and aN2 = 1700), the residual stresses were not

determined due to peak splitting with tilting.

2 Figure 3.10 a311 versus sin ψ from samples nitrided with 3 different nitriding activities

Grazing angle XRD was employed to reveal the formation of second phase, if any. A grazing angle of 1° was used. The corresponding X-ray penetration depth in this mode is less than 1 μm. By virtue of the shallow information depth, the impact of the core material is eliminated so that the diffraction peaks from the surface nitrides become obvious. The only sample demonstrating clear second phase peaks is the one nitrided

9 with the highest nitriding activity aN2 = 4×10 . The result is presented in Figure 3.11, together with the peaks reported for Cr2N (PDF#79-2159). The precipitate formed was indexed as M2N with a hexagonal structure, in whhich M represents the composition of the metal matrix, which is 316L stainless steel in the present study. When the nitriding/carburizing is done at low-temperatures, diffusion of the metal elements is believed suppressed. The precipitates formed in those conditions are called paraequilibrium nitrides/carbides.

66 Grazing Incident XRD 9 a ≈ 4×10 N2

M2N(1121)

M2N(1122)

M2N(0002)

M2N(1120)

9 Figure 3.11 Grazing XRD result from 316L treated at 440 °°C with aN2 = 4×10 .

The activity of nitriding has an evident effect on the case depth, as demonstrated by the SEM images in Figure 3.12. The images aree in the order of decreasing activity from

9 (a) to (d). The largest nitriding activity (aN2 = 4×10 ) produced the deepest case depth of

15~20 μm, while the case depth was only 2~3 μm after nitrided in aN2 = 200 for 20 hours.

However, cracks occurred when a too high acctivity was selected. Cracks on 316L

5 9 samples after nitriding in aN2 = 1.8×10 and aN2 = 4×10 are shown in Figure 3.13 and

3.14, respectively. Cracking of the surface is due to the very large lattice parameter expansion and hence enormous residual stress.

5 Figure 3.13 Cross-sectional SEM images of 316L stainless steel sample (440°C, aN2 =1.8×10 ) revealing the crack formation. The original surface is at the left-hand side of the image.

9 Figure 3.14 Cross-sectional SEM images of 316L stainless steel sample (440°C, aN2 = 4×10 ) revealing the crack formation. The original surface is at the bottom of the image.

Nitrogen concentration depth profiles of samples nitrided with different nitriding activities at 440 °C are presented in Figure 3.15. The effect of nitriding activity is clearly demonstrated. Not surprisingly, a higher nitriding activity yields a higher surface content of nitrogen. The sample nitrided with an activity of ~1 only yields a surface content of

9 about 7 at%. For the highest nitriding activity (aN2 = 4×10 ), in which only NH3 and N2 were used, surface content of nitrogen as high as 25 at% was achieved. The scattered data points near the surface of this sample is possibly due to the formation of nitrides, which have a higher nitrogen content as compared too the matrix. Besides surface content of nitrogen, the case depth also heavily depends on the applied nitriding activity, which is consistent with the observation from SEM metallography.

Figure 3.15 Nitrogen concentration profiles detected from samples nitrided with different nitriding activities at 440 °C

3.1.3 Effect of nitriding duration

The effect of nitriding duration is relatively easy to predict – a thicker case depth is expected with a longer treatment. The purpose of comparing samples nitrided with different durations is to understand the nitriding process in a kinetic view. Figure 3.16 shows the XRD results obtained from samples treated with different durations: 5 hours,

20 hours and 80 hours. (The 2 hours of nitriding time in-between the two activation segments was also counted into the total nitriding time.) All 3 samples were treated at

440 °C with an activity of 7400. The calculated ahkl values are demonstrated in Figure

3.17. For the sample nitrided for 80 hours, only 5 peaks were detected. For the sake of comparison, the mean lattice parameters were calculated from only these 5 peaks for those 3 samples and are listed in Table 3.3. The value reported for the sample treated for

70 20 hours was slightly different from that in Table 3.2, because the number of peaks considered are different. The standard deviation is smaller when fewer data points (5) are taken into account. This is because a400 is excluded, which shows the largest deviation from the average. Accordingly, the lattice of austenite expands more with a longer treatment. The difference is more obvious betweeen the two samples nitrided for 5 and 20 hours. Possible reasons for the larger lattice parameter expansion associated with longer- treated samples includes: (1) a higher surface nitrogen concentration; (2) a smaller concentration gradient within the X-ray penetration depth and (3) a higher level of residual stress in the expanded austenite layer. No second phase was detected by XRD, even from the sample nitrided for 80 hours. However, the danger of employing a longer duration is cracking of the surface. Cracks formed after 80 hours of nitriding are demonstrated in Figure 3.19.

Figure 3.16 XRD results of 316L stainless steel samples nitrided for different durations

Figure 3.17 Lattice parameters ahkl from samples nitrided foor different durations

Table 3.3 Calculated lattice parameters and lattice parameter expansions of samples nitrided for different durations

The hypothesis of a varying surface concentration with time is confirmed by AES analysis. The depth profiles of samples nitrided with different durations are shown in

Figure 3.18. All 3 lines were observed from a <200> oriented grain to exclude the effect of surface orientation. Clearly, both surface concentration and case depth increase with increasing nitriding duration. The surface concentration increases from 14 at% to 19 at% after 75 more hours of nitriding. Obviously, the surface didn’t reach the maximum concentration at the beginning of the nitriding process. In other words, the establishment

72 of equilibrium between the nitriding atmosphere and the treated surface requires a certain

amount of time.

In a convective boundary condition diffusion scenario, the mass transfer at the

interface of the gas-phase and the metal coupon can be described by

− = ( − )

where c∞ is the concentration in the treatment gas atmosphere and cs is the concentration in the metal at the surface. [2] The difference of those two values represents the driving

force of the mass transfer process at the interface. α is the so-called mass transfer

coefficient with a dimension of length/time and it describes the efficiency of mass

transport. Many factors in the heat treatment impact the mass transport coefficient, such

as nitriding temperature, the surface condition of the metal and also nitriding activity

provided by the gas mixture. The determination of magnitude of α requires numerical

simulation of the measured nitrogen concetration depth profile, which is beyond the

scope of the current study. However, the slowly saturated surface content of nitrogen

indicates that the mass transfer at the gas/metal interface is not perfectly efficient.

It is noted that the difference in surface nitrogen content is as large between the 5

hours and 20 hours sample, as it is between 20 hours and 80 hours sample. However, in

Figure 3.17, a200 for both 20 hours and 80 hours sample are similar, and both are much larger than that for 5 hours sample. This is because part of the residual stress was relieved in the 80 hours sample due to cracks formation (shown in Figure 3.19). Therefore the

73 Poisson expansion in the 80 hours sample is smaller as compared to that in the 20 hours sample.

Figure 3.18 AES profiles measured from 316L samples nitrided for different durations

Figure 3.19 Cracks formed on 316L sample after 80 hours of nitriding

74 3.2 Ferromagnetism induced by low-temperature nitriding

It has been predicted [3,4] that face-centered cubic (fcc) Fe can exhibit either

paramagnetic or ferromagnetic behavior, depending on the Fe–Fe interatomic distances.

The XRD results on nitrided 316L presented in the previous section (Tables 3.1 and 3.2)

have demonstrated that the lattice parameter expansion induced by low-temperature

nitriding processes varies from 1% to 10%, depending on the processing parameters.

Thus a room temperature stable ferromagnetic austenitic phase may be expected.

Two samples (440 °C, aN2 = 1700 and aN2 = 7400) are involved in the investigations

of ferromagnetism induced by supersaturation of nitrogen interstitials. The XRD results in Figure 3.3 clearly show the formation of single-phase expanded austenite on both samples. The lattice parameters and the corresponding lattice parameter expansions determined from (111), (200) and (220) peaks are listed in Table 3.4.

EBSD results showing the crystallographic direction of the grain surface normal in the nitrided stainless steel samples are displayed in Fig. 3.20 (a) and (c). (For evaluating the EBSD patterns, an fcc crystal structure was assumed for the expanded austenite.

However, the quality of the EBSD maps was limited, presumably because of the extensive plastic deformation that accompanies low-temperature nitriding). MFM results of the same regions of the samples are shown in Fig. 3.20 (b) and (d) and a ferromagnetic domain structure that varies from grain to grain is observed. Considering the small depth probed by MFM, this magnetic domain structure is clearly related to the formation of expanded austenite.

75 Table 3.4 Gas-phase nitriding activities and the corresponding lattice parameter expansion

Sample 111 200 220

aN2=1700 3.8 % 8.3 % 3.8 %

aN2=7400 4.9 % 9.4 % 4.9 %

Interestingly, after nitriding at both nitrogen potentials, only grains with surface normal nearly parallel to <100> showed distinctive magnetic domains. Other grains remained paramagnetic, or showed only faint evidence of ferromagnetism (see the enlarged image in Fig. 3.20 (f)). This variation of the paramagnetic-to-ferromagnetic phase transition with crystallographic orientation is a consequence of the anisotropy of the lattice parameter expansion in the expanded austenite. As listed in Table 3.4, for aN =

1700, the apparent lattice parameter expansion at the surface determined, for {200} planes, is 8.3%, while it is only 3.8% for both {111} and {220} planes. For the sample treated at aN = 7400, the lattice parameter expansions are larger for all planes: 9.4% for the {200} peaks and 4.9% for {111} and {220} peaks. Inasmuch as grains with a near surface normal <111> show a very fine ferromagnetic domain structure (see Fig. 3.20 (f)), the minimum lattice parameter expansion required for the ferromagnetic behavior is estimated to be about 5%. These results match the predictions of [1]. By calculating the total-energy surface in moment–volume parameter space employing a fixed spin-moment method, it was predicted that fcc Fe would exhibit a stable ferromagnetic state at approximately 5% lattice parameter expansion. The absence of ferromagnetism in expanded austenite containing a supersaturation of carbon interstitials can be readily

76 explained: the maximum lattice parameter expansion is not sufficient to yield room temperature ferromagnetism.

(a) (b)

(e) (f)

77 A cross-sectional SEM image of the sample treated at aN = 7400 is shown in Fig. 3.21.

The hardened case appears featureless in this image, implying better resistance against the metallographic etchant than the non-carburized alloy core. AFM/MFM images of the same area of the etched sample are shown in Figs. 3.22 a and b. In the AFM image, the case appears brighter because of its reduced etching. In the MFM image, magnetic domains are found, again indicating a ferromagenetic phase, but the depth of the paramagnetic-to-ferromagnetic phase transformation below the nitrided surface is different in differently oriented grains.

Comparing Figs. 3.22 (a) and (b), it is apparent that only the outer portion of the hardened case became ferromagnetic. This is clear evidence that a critical amount of nitrogen is required to induce the paramagnetic-to-ferromagnetic transition. To measure this threshold nitrogen concentration, AES nitrogen concentration–depth profiles (Fig.

3.22 c) were acquired along two parallel lines in the region imaged by AFM/MFM, as indicated by dashed lines in Fig. 3.22 a and b. (The shape of those profiles indicates strongly concentration-dependent nitrogen diffusivity.) The positions of the paramagnetic-ferromagetic austenite interfaces are labeled, based upon the MFM image in Fig. 3.22 b. A nitrogen concentration of about 14% is indicated by both profiles, identifying this value as the threshold content for the paramagnetic-to-ferromagnetic transition.

Selected-area diffraction patterns (SADPS) from the sample treated in aN = 7400 are shown in Fig. 3.23 (a) and (b). The pattern were taken from an area very close to the surface, where the nitrogen content is 15-17 at % as measured by the AES cross sectional line scan. The zone axes for the two diffraction patterns were indexed as FCC [110] and

78 [112], respectively. It is clear that only spots from FCC structure were detected. This demonstrates that the paramagnetism-to-ferromagnetism transition induced by low- temperature nitriding processes is not due to a phase transformation to the γ’-Fe4N, which has a simple cubic structure. Otherwise, forbidden diffraction spot from the ordered nitrogen atoms in γ’-Fe4N should be detected

Figure 3.21 Cross-sectional SEM image from nitrided 316L stainless steel (aN2=7400). The original surface is on the top of the image.

(a) (b)

(c)

Figure 3.23 Diffraction patterns taken from 316L nitrided in aN2 = 7400 (a) [011] zone axis and (b) [112] zone axis.

81 3.3 Orientation-dependent case depth

The images on cross-sectional samples in Figure 3.12 and 3.13 already reveals that

the case depth produced by the low-temperature nitriding processes varies from grain to

grain. The difference is more obvious when a thicker case depth is produced. A combination of EBSD and AES was used to study this orientation-dependent case depth produced by low-temperature gas-phase nitriding processes. Figure 3.24 shows a cross sectional SEM image taken from the sample nitrided for 80 hours (440 °C, aN2 = 7400).

The case was revealed by Ar plasma sputtering in AES, the purpose of which was to remove the damaged layer from mechanical polishing for EBSD analysis. The difference in case depth is as large as 6 μm. The EBSD orientation map of the same area is exhibited in Figure 3.24 (b) and the color legend is given in Figure 3.24 (c). Grains with red color in the EBSD map have a surface normal close to <100> direction, and demonstrate a deeper case depth, while grains with blue color have a surface normal close to <111> direction and have a shallower case depth. Similar investigations were also done on the two samples treated for 5 hours and 20 hours. The relationship between surface orientation and case depth observed are the same.

AES analysis was then performed on the areas previously studied by EBSD. Line scans were taken from grains with their surface normal close to <111> and <200> separately. Nitrogen profiles from all three samples treated with different durations are shown in Figure 3.25, from a to c. In all three images, profiles measured from <200>- oriented grains are represented by circles and those measured form <111>-oriented grains are represented by squares. The data confirm the conclusion drawn from the SEM and

82 EBSD study-the <200>-oriented grains always demonstrate deeper case depths compared

to the the <111>-oriented grains. Certainly, the difference became larger with time. The

difference in surface concentration between <111>-oriented and <200>-oriented grains

are only obvious for the 80 hour nitrided sample. The surface concentration measured

from <200>-orientation grain is about 2% lager as compared to <111>-orientation grain.

On the other 2 samples treated with shorter durations, the surface concentrations detected

from different grains are similar, or that the difference is too small to be identified by

AES analysis, which has a typical error of ~1 at%.

If the case thicknesses are derived from the AES depth profiles and plotted against the square root of nitriding time (√), linear relationships are approximately yielded for both

<111> and <200> oriented grains (Figure 3.25 d). An approximate linear relationship can be employed to predict the case depth produced in the future.

The same investigation (EBSD+AES) was also done on a carburized sample, which was carburized in 10% CO (diluted with 45 vol.% N2 and 45 vol.% H2) at 440 °C for 20

hours. The same activation procedure (double HCl activation at 325 °C) was applied. The

observed AES depth profiles measured from grains with different surface normals are

presented in Figure 3.26. As demonstrated, a much smaller difference in case depth is

observed on the carburized 316L. The case depths measured from <111>-oriented and

<200>-oriented grains are about 16 μm and 20 μm, respectively. About 20% variation is

observed. On the other hand, the change in case depth observed from the 316L sample

nitrided from 80 hour is as large as ~35%.

(a)

(b)

(c)

Figure 3.24 (a) Cross-sectional SEM image from 316L after nitrided for 80 hours; (b) EBSD orientation map taken from the same area; and (c) color key of (b)

(a) (b)

Figure 3.26 AES carbon depth profiles measured from different surfacce oriented grains of a carburized 316L

3.4 Hardness and modulus measurements

The most attractive feature of the expanded austenite layer is the improvement of multiple mechanical properties, especially the hardness. The hardening mechanissm is primarily based on solid solution hardening induced by the colossal amount of nitrogen atoms. Figure 3.27 presents the microhardness measurements on the plan-view surface on

316L stainless steel samples nitrided with different nitriding activities. The load used was

25 gram force (gf). The measured hardness value shows strong activity dependence. The

9 highest microhardness measured is about 1250 HV (nitrided in aN2 = 4×10 ). The value decreases quickly with a decreasing nitriding activity. However, since the measurements were done on a plan-view sample, the decreasing hardness value is due to both a lower

86 surface nitrogen content and a smaller case depth. The largest microhardness measured on a precipitate-free nitrogen expanded austeniite is about 1050 HV (nitrided in aN2 =

7400) and is 5 times the microhardness of the non-treated sample.

Figure 3.27 Plan view microhardnesses of 316L stainless steel samples before and after low-temperature nitriding with different nitriding activities

Nanoindenter was employed to determine tthe modulus on the sample nitrided at

440 °C in an activity of 7400. A mode called CSM (Continuous Stiffness Measurement) was chosen foor the measurement. This mode was designed to evaluate films on substrates, where the mechanical properties change as a funcction of surface penetration. A limitation of this mode is that the penetration depth into the surface can only be as deep as 2 μm.

The local crystallographic orientations were indexed by EBSD. A slight etching was performed in a solution of 50%HCl+25%NHO3+25%H2O to remove surface

87 contaminations and the damaged layer before EBSD examination. The observed pattern was shown in Figure 3.28 (a). Optical microscope images were taken for each of the indentations to make sure that the measurements were done in the aimed location. Figure

3.28 (b) shows an optical micrography of the indentation made on a <100>-oriented grain.

As compared to the surrounding <111>-oriented grains, the brighter contrast of the

<100>-oriented grain implies a better resistance to the etchant used. The achieved hardness and modulus profiles are demonstrated in Figure 3.28 (c) and (d), respectively.

Similar to the nitrogen concentration profile, the hardness depth profile can also be described by a slow decrease from the surface followed by a fast drop. However, the tail for the hardness profile is not as sharp as the concentration profile, which is because the hardness measurement is influented by material below the depth of the measurement.

Compared to the hardness profile measured from the nontreated-316L (black line), 3.5~4 times larger hardness is achieved at the surface. Also, same as the nitrogen concentration profiles, the hardness profiles clearly demonstrate a surface-orientation dependence, especially by means of hardened depth. Hardness measured from <111>-oriented grains decreases faster as compared to <100>-oriented grains.

The modulus measurements are quite noisy. The general information observed is that a slightly larger surface modulus is achieved on the nitrided 316L. EBSD was not performed on the non-treated sample before the modulus test. A random measurement was conducted and the result is shown in Figure 3.28 (d) as the black line. On the nitrided sample, E111 > E200 is still true, which is the same as the non-treated 316L. However, it is noted that when the penetration depths are larger than 1 μm, the measured moduli from the nitrided sample become smaller than the non-treated material. A possible reason is

88 proposed as the following. For soft metals like non-treated 316L, pile-up can occur at the edges of the indent, and that will increase the apparent modulus. Therefore, the modulus measured for the non-treated sample is probably higher than the actual value. The hardness tests already reveal that the nitrided samples are much harder, so that no pile-up happened during the measurement.

(c)

(d)

90 9 3.5 TEM results on nitrided 316L (aN2 = 4×10 )

TEM study was performed on the 316L sample nitrided at 440 °C with a nitriding

activity of 4×109. According to the grazing angle XRD study, this is the only sample containing second phases after low-temperature gas-phase nitriding process. A

comparison between the 2θ positions measured from Figure 3.13 and the standard XRD

peaks documented for M2N, MN, M4N and M3N is shown in Table 3.5. Based on Table

3.5, the phases coexisting in the “case” layer were identified to be nitrogen-enriched

expanded austenite (γN) and M2N. The formed nitrides were considered as a paraequilibrium phase. M represents the composition of 316L, i.e. 68 at% Fe, 19 at% Cr,

11 at% Ni and 2% Mo. This explains the 0.1° ~ 0.2° mismatch in 2θ angles between M2N and Fe2N/Cr2N. On the other hand, the difference between γN and γ’-Fe4N is 1.5° ~ 4.5°,

which is difficult to be explained by a variation in composition. Although the formed

second phase is labeled as “M2N”, following the notation used in PDF cards, it doesn’t necessarily mean that the phase contains 33.3 at. % of nitrogen. Actually, according to

[5] the Fe-N phase diagram , the hexagonal iron nitride Fe2N (or ε nitride) is a

nonstoichiometric phase with a composition range of 15 ~ 33 at. % N. Therefore, it is

more reasonable to label the formed second phase as ε-M2N1-x.

TEM foils were prepared by FIB lift-out technique. Figure 3.29 shows a cross-

sectional STEM image of a prepared TEM foil. A thin layer of Pt was deposited to

protect the area of interest from the ion beam. The STEM was taken with a camera length

(CL) of 1.44 m, which implies that the contrast in the image is a combination of both Z contrast and diffraction contrast. The linear features demonstrated in the STEM image are

91 related to the slip band resulted from the tremendous magnitude of residual stress in the case layer.

Table 3.5 Comparison between 2θ positions from measured XRD and standard peaks of nitrides

9 Figure 3.29 STEM image taken from a cross-sectional 316L sample nitrided at 440 °C in aN2 = 4×10 .

92 As shown in the plan view SEM image (Figure 3.30 (a)), plenty of surface particles

9 formed on the surface of 316L sample after it was nitrided at 440 °C in aN2 = 4×10 for

20 hours. After the FIB thinning process, only a few of the particles remained on the surface of the TEM foil. The cross-section of the formed surface particle is shown in the

STEM image in Figure 3.30 (b). The particle contains several sub-grains, the sizes of which are smaller than 500 nm. XEDS mapping was performed to reveal the distribution of various elements and the results are shown in Figure 3.31. It is clear that the surface particle is mainly Fe enriched as compared to the matrix material underneath. Inside the particle, Cr segregates in fine features, which corresponds to regions with darker contrast in the STEM image. Ni is absent from the surface particle. The surface particle and the matrix underneath it demonstrate almost no difference in nitrogen content.

SAD patterns observed from the surface particles are presented in Figure 3.32 (a), (c) and (e). (DPs presented in (a) and (c) were taken from two individual grains of the surface particle in Figure 3.30 (b). DP in (e) was taken from surface particle remaining on another FIB prepared TEM foil.) The smallest SAD aperture was used to exclude information from other grains. All three patterns can be indexed according to a ε-M2N1-x

structure. The corresponding JEMS simulated diffraction patterns based on ε-M2N1-x

structure are demonstrated in Figure (b), (d) and (e), respectively. Therefore, the main

phase in the formed surface particle is determined to be ε-Fe2N1-x based on a combination

of diffraction results and XEDS results. The structure of the small Cr-enriched features is unclear.

9 Figure 3.30 (a) Plan-view SEM image of 316L sample nitrided at 440 °C in aN2 = 4×10 and (b) STEM image showing the cross-section of the formed surface particle

Figure 3.31 Elemental mapping by XEDS of the formed surface particle after 316L sample was nitrided at 9 440 °C in aN2 = 4×10 . (a) STEM image; (b) N map; (c) FFe map; (d) Cr map; (e) Ni map and (f) Pt map (protective layer before FIB sample preparation).

Figure 3.32 SAD patterns observed from the surface particle in (a) [3, -1, -2, 2] zone axis, (c) [4, -2, -2, 3] zone axis and (e) zone axis of ε-M2N1-x. The corresponding JEMS simulated diffractiion patterns in the same zone axes are shown in Figure (b), (d) and (e), respectively. Double diffraction spots are included in the simulated patterns.

96 A typical SAD pattern obtained from the expanded austenite layer of 316L (nitrided

9 at 440 °C in aN2 = 4×10 ) is presented in Figure 3.33 (a). Based on a six-fold symmetry,

the strong diffraction spots could be indexed as the <111>-zone axis of the austenite

phase with a FCC structure. However, forbidden reflections of FCC structures are

observed in the diffraction pattern, such as {110} reflections, which are located between

the transmitted beam and {220} reflections. The existence of forbidden diffraction

reflections for FCC lattice was confirmed by tilting TEM samples to different zone axes.

SAD patterns obtained in [100], [112] and [011] zone axes are shown in Figure 3.33 (c),

(e) and (g), respectively. These results indicate that the diffraction pattern should be indexed based on a simple cubic (SC) structure.

To obtain an SC structure from an FCC structure, it is required for nitrogen atoms to occupy one of the four octahedral sites of the austenite matrix in a fully ordered manner.

Actually, the structure of the γ’-Fe4N phase can be considered as an interstitial

compounds with the iron atoms in an FCC close packed fashion, with fully-ordered

nitrogen atoms. In this way, the strain energy can be minimized. [6]. According to the

nitrogen concentration depth profile demonstrated in Figure 3.15 (measured by AES

cross-sectional line scan), about 25 at.% nitrogen was brought into the 316L sample after

nitriding at 440 °C with a nitriding activity of 4×109 for 20 hours, which suggests that

more than ¼ of the octahedral sites were occupied by nitrogen atoms. According to the

TEM results shown in Figure 3.33, it is obvious that those nitrogen atoms do arrange themselves in an ordered manner. In other words, a γ’-M4N structure was formed. In

Figure 3.33 (b), (d), (f) and (h), the simulated diffraction patterns (using JEMS software)

97 based on a γ’-M4N structure are demonstrated. It is evident that the observed and

simulated patterns match nicely.

However, previous XRD study has demonstrated that no peak matches the

documented peaks for γ’-Fe4N. This is because the formed γ’-M4N structure has a larger

lattice as compared to γ’-Fe4N. The lattice parameters of the formed γ’-M4N structure

were evaluated for different crystal planes based on 20 DPs taken from the same sample

and the calculated results are summarized in Table 3.6. It is noted that lattice parameters

measured from TEM DPs agree well with those measured by XRD, which are averaged

to 0.3957±0.0019 nm (Table 3.2). On the other hand, the lattice parameter reported for

stoichiometric Fe4N is about 0.3798 nm at 20 at.% of nitrogen [5]. The formed γ’-M4N demonstrated ~ 4% lattice parameter expansion as compared to γ’-Fe4N. One possible

reason for the expanded lattice is the excess amount of nitrogen (about 25 at.% measured

[7] by AES, instead of 20 at.% in stoichiometric M4N). Somers has studied the dependence

of lattice parameter on nitrogen concentration for γ’-Fe4N and the observed linear

relationship was described as [7]:

= 0.37988 + 14.82 × 10 ×( − 20) (3-1)

where CN is the atomic percentage of nitrogen. If we assume that the same linear

relationship holds for 25 at.% N or even higher nitrogen content, the lattice parameter

predicted for γ’-Fe4N with 25 at.% N is about 0.3873 nm. This value is close to the lattice parameters in Table 3.6.

100 Table 3.6 Lattice parameters of different planes of the γ’-M4N structure calculated from DPs

(hkl) (111) (200) (220) (311) (331) (420)

0.391±0.0 0.393±0.0 0.392±0.0 0.392±0.0 0.394±0.0 0.394±0.0 a (nm) hkl 03 02 03 02 02 05

Table 3.7 Zone axes of γ matrix in which twin and HCP plates may/may not be distinguished [8,9]

Examples

Zone axes with no distinction between twin [100], [112], [013] and HCP

Zone axes with extra spots due to twin on (111) [114], [122], [115] matrix planes

Zone axes with extra spots due to HCP plates [111], [125], [147] on (111) matrix planes

Zone axes with extra spots due to twin and [110], [231], [012] HCP plates on (111) matrix planes

As mentioned previously, XRD peaks from ε-M2N1-x were also detected from the

nitrided 316L. Efforts were made to identify the existing ε-M2N1-x phase in the expanded case layer by means of electron diffraction. It is well known that micro-twins were

produced in the expanded austenite layer and the twinning system is {111}<112> [11].

Therefore, extra spots could originate from either ε-M2N1-x nitride or micro-twins. Based on computer simulation of composite theoretical electron-diffraction patterns for a γ-

Twin-HCP system [8,9], which is similar to what we have, HCP phase and microtwins can

101 only be distinguished when the incident electron beam is parallel to certain zone axes of

the γ phase. A quick survey of zone axes in which twin and HCP plates may or may not

be distinguished is given in Table 3.7.

According to Table 3.7, when the zone axis of γ matrix is in [122] orientation, the

extra spots are due to twin structure. SAD patterns recorded from [122] zone axis of γ’-

M4N with extra spots are shown in Figure 3.34 (a). As predicted, the extra spots can be

indexed as twin with an orientation of [001]. Simulated composite diffraction with γ’-

M4N [122]//[001] is shown in Figure 3.34 (b). The observed and simulated diffraction

patterns show good agreements with each other, except that some diffraction reflections

(such as the one circled in Figure 3.34 (b)) were missing from the observed pattern.

Those diffraction spots are reflected from the ordered nitrogen atoms in γ’-M4N twin, whose intensities might be too weak to be detected. The morphology of the micro-twin structure was revealed by dark field images shown in Figure 3.34 (d) and (e).

Figure 3.35 (a) is a conventional bright field TEM image taken from the circled area in Figure 3.29. The corresponding SAD pattern is shown in Figure 3.35 (b). The zone axis of the γ’-M4N phase is indexed as [111]. 6 of the 12 weak spots surrounding the

transmitted beam can be described as (110) reflections of γ’-M4N. Dark field image

generated by one of those reflections (spot 2 in Figure 3.35 (b)) is shown in Figure 3.35

(e). No difference is observed between the bright field and the dark field images, which implies that the whole matrix has transformed to the expanded γ’-M4N phase.

According to Table 3.7, when the incident electron beam is parallel to [111]

orientation of γ’-M4N, extra spots are supposed to be due to HCP plates (ε-nitride in our

102 case). Then the other 6 weak diffracted spots surrounding the transmitted beam can be explained by [0001] zone axis of ε-M2N1-x. The simulated composite diffraction pattern

with an orientation relationship of γ’-M4N [111] // ε-M2N1-x [0001] is shown in Figure

3.35 (c). However, as pointed out by Kestenbach [8], there is a chance that “apparent”

HCP spots originating from twin plates oriented perpendicular to the direction of incident electron beam. Therefore the lath-shape morphology revealed in the dark field image in

Figure 3.35 (d) (taken with spot 1 in Figure 3.35(b)) can be either ε-M2N1-x or micro-twin.

Based on diffraction patterns only, [011] orientation of γ’-M4N is one of the candidate

zone axes for reliable identification of ε-nitride and twin structure. Simulated pattern with

[10] an OR of γ’-M4N <011> // ε-M2N1-x [-12-10] (as predicted by D. H. Jack is shown in

Figure 3.36 (b). However, SAD patterns observed with [011] zone axis of γ’-M4N from the ntirided sample are always in lack of extra spots belonging to ε-nitrides (e.g. Figure

3.36 (a)), which indicates that no ε-M2N1-x was found in the expanded austenite layer.

103

104

Figure 3.35 (a) Bright filed image taken from the circled area in Figure 3.30. (b) Corresponding SAD pattern. (c) Simulated SAD pattern with γ’-M4N [111] // ε-M2N1-x [0001]. (d) Dark field image taken with diffraction spot 1 in (b). (f) Dark field image taken wiith (110) reflecction of γ’-M4N (spot 2 in (b)).

105

Figure 3.36 (a) SAD pattern in [011] zone axis of γ’-M4N and (b) JEMS simulated composite diffraction pattern with γ’-M4N <011> // ε-M2N1-x [-12-10]

106 References

[1]. H. Kahn, G. M. Michal, F. Ernst and A. H. Heuer: Metallurgical and Materials

Transactions A, 2009, 40A, 1799-1804

[2]. X. Gu, Ph.D. Dissertation, Case Western Reserve University, 2011

[3]. V. L. Moruzzi and P. M. Marcus, Physical Review V, 1986, 34, 1784-1791

[4]. M. Uhl, L.M. Sandratskii, and J. Kubler, Physics Review B, 1994, 50, 291-301

[5]. H. A. Wriedt, N. A. Gokcen, and R. H. Nafziger, J. Phase Equilib. 1987, 8, 355-

377

[6]. D. H. Jack and K. H Jack, Materials Science and Engineering, 1973, 11, 1-27

[7]. M. A. J. Somers, N. M. van der Pers, D. Schalkoord, E. J. Mittemeijer,

Metallurgical and Materials Transactions, 1989, 8, 1533-1539

[8]. H-J Kestenbach, Metallography, 1977, 10, 189-199;

[9]. V. Vodarek, J. Sojka and E. Sojka, Czech. J. Phys. B, 1988, 38, 767-776

[10]. D. H. Jack, Materials Science and Engineering, 1974, 13, 19-27]

[11]. H. Dong, International Materials Reviews, 2010, 55, 65-98

107 Chapter 4 Low-Temperature Gas-Phase Nitrocarburizing

4.1 Three scenarios of nitrocarburizing with NH3/CO/H2/N2

As mentioned in Chapter 2, there are three different designs of nitrocarburizing processes which can produce a dual layer of expanded austenite. The aim of this section is to compare results obtained from those three different scenarios of nirocarburizing processes. The processing parameters used are listed in Table 2.4. Results from 316L after single nitriding (aN2 = 7400) and single carburizing (aC = 0.4) processes done at the same temperature (440 °C) were also included.

4.1.1 XRD analysis

Phase identification was conducted by XRD analysis and the spectra segments (38° to

52°) collected in Bragg-Brentano mode are shown in Figure 4.1. Results observed from samples treated with single nitriding and carburizing process are also given for comparison. The lattice parameter calculated from each expanded austenite peak is plotted in Figure 4.2. Seven peaks are plotted for each of the sample. The average lattice parameter for each treated sample was calculated by taking an arithmetic mean of all 7 peaks and the results are listed in Table 4.1.

In general, the lattice parameter expansions induced by both single nitriding (20N) and nitrocarburizing processes are larger than that induced by a single carburizing process (20C). The largest average lattice parameter expansion achieved is about 6.4%, measured from both simultaneous nitrocarburizing (20(N+C)) and 20N samples. It is also

108 noted that the standard deviations of lattice parameter from the nitrided/nitrocarburized

samples are larger than that from the carburized one. Since the same number (7) of

diffraction peaks were used in the calculation, the larger standard deviations are due to

more scattered data points. In other words, all nitrogen-containing samples display a large anisotropy in lattice parameter ( >). Interestingly, the values of a111 from 20C

and 20N+20C samples are similar, while the values of a200 and a311 from 20N+20C

sample are much larger. This is probably due to a transition from paramagnetic austenite

to ferromagnetic austenite in the nitrocarburized sample, which will be discussed in more

detail in Chapter 5.

Grazing angle XRD was performed to reveal the possible formation of second phases

(Figure 4.3). After carburizing in 10% CO for 20 hours at 440 °C (20C), carbides having a high density were produced on 316L, which is indicated by a strong extra peak detected at 2 = 44.7°. Together with other 3 peaks at 2θ = 49.9°, 50.7° and 52.8°, the formed

carbide is identified as ω-carbide (Fe7C3). No obvious nitride peak was detected on the

sample treated by the single nitriding process.

Among the three nitrocarburized samples, no other phases besides the expanded

austenite were detected on the sample treated with 20N+20C. On the simultaneously

nitrocarburized sample (20(N+C)), only 1 extra peak at 2 = 43.7° was detected. It is difficult to identify the formed second phase based on one XRD peak. Nitride or carbide which demonstrates the strongest peak at this 2θ angle includes CrN, Fe3N and Fe3C.

However, it is clear that no ω-carbide was formed. In other words, the formation of ω- carbides was suppressed. One possible reason is that the activity of carbon provided by

109 the gas-phase was lowered in the nitrocarburizzing atmosphere. For the sample treated with the 20C+20N process (20 hours carburizing followed by 20 hourrs nitriding), both nitrides and carbides were detected by grazing XRD analysis. The formed nitrides are identified as CrN with a FCC structure. Six peaks from the carbide phase were detected, with 2θ = 39.8°, 44.7°, 47.1°, 49.9°, 50.7° and 522.8°. The positions of those peaks match with the standard peaks for Fe7C3 (PDF#751499) nicely.

Figure 4.1 X-ray diffraction patterns of 316L stainless steel before and after carburizing, nitriding and various nitrocarburizing scenarios

110 Figure 4.2 ahkl of 316L stainless steel samples after carbburizing, nitriding and various nitrocarburizing processes

Table 4.1 Lattice parameters of 316L stainless steel samples after carburizing, nitriding and various nitrocarburizing processes

111

Figure 4.3 Grazing angle XRD data of 316L stainless steel after carburizing, nitriding and various nitrocarburizing processes

4.1.2 Metallography

The SEM images of cross-sectional nitrocarburized samples are shown in Figure 4.4.

The original surfaces are all located at the bottom of the images. The case layers in all of the samples were revealed by an etchant of 50 vol.% HCl+25 vol.% HNO3+25

vol.%H2O. Clearly, both the nitrogen-enriched austenite and carbon-enriched austenite layers demonstrate a better resistance to the etchant used. As a result, the case layers were less etched and smoother as compared to the core part. Also, an interface in the middle of

112 the case layer was also revealed (especially in Figure 4.3 (c)), which implies that the

electrochemical properties in the outer and inner parts of the case are different.

Among the three samples treated for 20 hours, the nitrided sample produced the thinnest case (5-8 μm). The case depths of the carburized sample and the simultaneously nitrocarburized one are similar (about 20 μm). For samples treated for 40 hours, the

20C+20N sample demonstrates a thicker case (25-30 μm) than the 20N+20C one (16-20

μm). Generally, the total case depth is mainly determined by the diffusion time of carbon.

For 20C, 20(N+C), 20N+20C samples, the diffusion time of carbon was all 20 hours and all samples show similar case depth. For the 20C+20N sample, the total diffusion time for carbon was 40 hours, and thus the sample shows a thicker case. The case depth on the two sequentially nitrocarburized samples (20N+20C and 20C+20N) also varies from grain to grain, which is similar to what we observed from the study of low-temperature nitriding samples.

113

Figure 4.4 SEM images on the cross sections of 316L bulk treated with (a) 20 hour nitriding (20N); (b) 20 hours carburizing (20C); (c) 20 hours simultaneous nitrocarburizing (20(N+C)); (d) 20 hours nitriding followed by 20 hours carburizing (20 N+20C); and (e) 20 hour carburizing following by 20 hour nitriding (20C+20N).

114 4.1.3 AES depth profiles

The nitrogen and carbon depth profiles were measured by AES line scans on well-

polished cross-sectional samples. The results for all of the 5 samples are shown in Figure

4.5, from (a) to (e). The profiles of 20(C+N), 20N+20C and 20C+20N confirm that dual

expanded austenite layers have been produced by all three nitrocarburizing processes. For all three scenarios, the outer layer is always nitrogen-supersaturated, regardless of the sequence of diffusion. In other words, carbon always diffuses through the nitrogen case and accumulates at the diffusion front of nitrogen.

AES data are consistent with SEM results (Figure 4.4), which indicates that the total case depth of the nitrocarburized sample is mostly determined by the diffusion time of carbon. The profiles from samples treated by single nitriding (20N) and carburizing

(20C) processes were both measured from <100>-oriented grains, meaning that those profiles represent roughly the deepest case depth. After a single carburization process at

440 °C, carbon atoms penetrate about 20 μm into the steel after 20 hours of diffusion.

Nitrogen atoms only diffuse to a depth of ~8 μm at the same temperature for 20 hrs. This

implies that the diffusivity of carbon at 440 °C is much larger than that of nitrogen. For a simultaneous nitrocarburizing process at the same temperature, the total case depth is also

~ 20 μm. The total diffusion time of the two sequential nitrocarburization runs was both

40 hours. However, the total case depth measured was very different from each other.

The 20N+20C run produced a case depth similar to the single carburization, ~ 20 μm. On the other hand, for the 20C+20N hours run, in which the carbon diffused for 40 hours altogether, the produced case depth is as deep as 30 μm.

115 As presented in Figure 4.3, the lattice parameter expansions produced by different

treatments are different. Thus different surface interstitial contents are expected. The

maximum lattice parameter expansion was measured on 20N and 20(N+C) samples. The

nitrogen contents measured are also the highest and similar for both samples (17~18

at%). However, about 2 at% of carbon was also measured on the 20(N+C) sample. It is intriguing that this extra amount of interstitial atoms did not result in a larger lattice parameter expansion. Therefore the question arises if the carbon signal is real in the nitrogen enriched expanded austenite. This can be explained by the detailed AES spectrum for C measured from the nitrogen-enriched region (Figure 4.6). The existence of carbon is well revealed by a peak at 275 eV. However, because of the poor signal-to- noise ratio, the quantity of carbon is possibly overestimated. Comparing the two sequentially nitrocarburized samples, the lattice parameter measured on 20C+20N is larger. This is also consistent with the interstitial content measured. The surface nitrogen content measured on 20C+20N is about 2 at% higher than that measured on 20N+20C sample. For the 20N+20C sample, the nitrogen atoms diffuse for 40 hours in total.

However, for the later 20 hours, no nitrogen was provided by the atmosphere, so that the surface content was lowered.

116 (a) 20N 20 hrs (b) 20C 20 hrs

(d) 20N+20C 40 hrs (e) 20C+20N 40 hrs

Figure 4.5 AES depth profiles of nitrogen and carbon meassured from 316L treated by (a) 20N; (b) 20C; (c) 20(N+C); (d) 20N+20C and (e) 20C+20N

117

Figure 4.6 AES spectrum shows carbon peak detected from the nitrogen-enriched layer of 20(N+C)

4.1.4 Hardening effect

The hardening effect achieved from the 3 nitrocarburizing processes on 316L samples

was examined by surface hardness measurements using the micro-hardness indentation

technique. The applied load is 25 gf. The results measured on the samples treated by single nitriding and carburizing runs are also presented (Figure 4.7). At least 10 measurements were done on each sample. Generally, all expanded austenite layers provide an effective improvement on the surface hardness of 316L. The hardness measured on 20N+20C sample is slightly smaller than the others. Two possible reasons are presented here: (1) The surface nitrogen interstitial content was slightly lower than the others and (2) no precipitate was formed. Nitrides and/or carbides were detected by grazing angle XRD (Figure 4.3) from samples treated by the other two nitrocarburizing processes. The standard deviation in hardness measurements from the sample treated by a single nitriding run (20N) is larger. This is due to the shallower case depth on that sample.

118 20C+20N

20(N+C) 20N

20N+20C 20C

Measurement

Figure 4.7 Surface hardness of 316L bulk sample treated with nitriding, caburizing and nitrocarburizing processes

4.1.5 Study of second phases after 20C+20N by TEM

As discussed previously, both carbides and nitrides were formed on 316L after a nitrocarburizing of 20C+20N, which is 20-hour carburizing followed by 20-hour nitriding

(aN2 = 7400). The formed nitrides were identified as MN by XRD and the carbides were identified as M7C3 (ω-carbide). Transmission electron microscopy (TEM) was utilized to study the carbide and nitride phases formed. FIB lift-out technique was employed to prepare the TEM foil. An overview of the TEM foil is shown by the STEM image in

119 Figure 4.8 (a). The original surface is at the bottom of the image. A thin layer of sputter-

coated Pd was created before FIB lift-out to avoid drifting issues. Pt was deposited to protect the area of interest from the ion beam. Those two layers are distinguishable in the

STEM image due to the difference in their atomic numbers Z. The dark layer with a thickness of about 200 nm in-between the Pd layer and the nitro-carburized 316L sample was identified as carbon soot by XEDS analysis (Figure 4.8 (b)). The Cu signal detected

by XEDS is from the copper TEM grid used. Needles with high density are observed in

the 316L matrix, which were later identified as second phases by diffraction.

Figure 4.9 (a) presents a conventional bright-field TEM image of the same foil. The

corresponding SAD pattern is demonstrated in Figure 4.9 (b). Three sets of diffraction

patterns were indexed, which were from the matrix (nitrogen-enriched expanded

austenite, γN), MN and M5C2. A well-established cube-cube orientation relationship is

found between the γN matrix and the fcc-MX nitride. The observed orientation

relationship (OR) can be described as:

< 110 > //< 110 >

(001) //(001)

(111) //(111)

The ratios of the d-spacing between γN and MX phases measured from the diffraction

pattern are calculated as:

= 1.14 =1.12

The values calculated from XRD are:

120 = 1.06, =1.1

Two possible reasons may lead to the difference between those two measurements: (1)

large errors may exist in measuring d-spacing from diffraction patterns, especially when

the diffracted spot is not sharp enough; (2) the residual stress in the expanded austenite

layer is largely relieved in the TEM foil, which makes the and measured from

diffraction patterns smaller and a larger ratio resulted.

Besides γN and MX phases, the third set of diffraction spots were indexed as [013] zone axis of the M7C3 carbide with the help of JEMS simulation. A dark field image was

taken using (23-1) diffracted beam of the M7C3 carbide and is shown in Figure 4.9 (d),

which indicates that the formed carbide have a needle-like feature. (To better reveal the

details in the carbide phase, the dark-field image was taken at a larger magnification as

compared to the BF image.) Accordingly, the OR between M7C3 and the surrounding γN matrix is as determined as:

[ ] 110 //[013]

(1 − 11) //(100)

This OR is exactly the same as reported by F. Ernst et al. [1].

Electron-spectroscopic imaging (ESI) technique was employed to perform elemental

mapping of carbon and nickel using three-window method. The K absorption edge of C

(at 284 eV) and the M absorption edge of Ni (at 855 eV) were employed for the analysis.

Figure 4.10 (a) shows a bright-field TEM image in an area containing high density of

121 needle-shape second phases. The ESI maps for Ni and C from the same area are

presented in Figure 4.10 (b) and (c), respectively. As expected, the needle-shape carbides

demonstrate a brighter contrast in the carbon map as compared to the austenite matrix

next to them. The contrast in the nickel map is reversed- the austenite matrix presents a

brighter contrast as compared to the carbides. In other words, nickel atoms were repelled

to the neighboring austenite regions when carbides were formed. This observation is

consistent with previous studies on carbides formed after low-temperature carburizing of

316L. [2].

Because the diffraction spots from fcc-MN and γN are too close to be separated by the

objective aperture, an attempt was made to find the MN phase using ESI mapping.

According to the AES results, the nitrogen content in γN is only about 18 at%. On the other hand, the nitrogen content in MN is 50 at%. It can be expected that the MN phase is

recognized as nitrogen-rich feature in the nitrogen elemental map. Figure 4.11 (a) and (b)

present a TEM bright field image and the nitrogen map taken in the same region,

respectively. The K absorption edge (at 400 eV) of N was employed for the analysis and

a spectrum recorded from the area being analyzed is shown in Figure 4.11 (c). The high

peak intensity of the nitrogen edge implies a high nitrogen concentration in the analyzed area. According to the nitrogen map, the carbide appears darker than the neighboring γN.

This is expected because the carbide has little solubility of nitrogen. In the γN region, it seems that certain very fine particles contain more nitrogen than the surrounding material.

However, the poor signal-to-noise ratio of the nitrogen map makes it difficult to make a solid conclusion about the MN phase.

122

(b) XEDS spectrum from the surface dark layer

) a.u. ( ty i ntens I Ii()

Figure 4.8 (a) STEM image taken from nitrocaburized 316L by a 20C+20N process and (b) XEDS result from the surface C (soot) layer.

123

(a) (b)

Figure 4.9 (a) Bright-field image taken from a cross-sectional nitrocarburized 316L. (b) SAD pattern with cube-cube OR between fcc-MN and γN indexed. (c) The same SAD patttern as in (b) but with ω-carbides and γN indexed. (d) Dark-field image observed using the encircled diffracted beam of ω-carbide in (c).

124

(a)

(b) Elemental map of Ni (c) Elemental map of C

125 Ni

counts

850 830 870 890 (d) Nickel edge

counts

310 330 270 290 (e) Carbon edge

Figure 4.10 (a) Bright field image. (b) ESI map of nickel. (c) ESI map of carbon. (d) EELS spectrum showing the nickel edge observed and (e) EELS spectrum showing the carbon edge observed.

126

(a) (b)

N counts

380 400 420 440 (c)

Figure 4.11 (a) Bright-field image obtained from an area containing less carbide. (b) ESI map of nitrogen. (c) EELS spectrum showing the nitrogen edge oobserved

127 4.2 Effect of nitriding activity on 20C+20N

The results in section 4.1 reveal that both nitrides and carbides were developed on

316L sample after treated with a carburizing process followed by a nitriding process

(20C+20N). An experiment with a lower nitriding activity in the second process was

designed and performed, aiming to avoid or at least reduce the formation of second

phases. Detailed processing parameters used for those two sequential nitrocarburzing

runs were listed in Table 2.3. The nitriding activity was reduced from 7400 to 1700 by

adjusting the composition of NH3/N2/H2 mixture. In the following discussion, 20C+20N

(H) and 20C+20N (L) represent the 316L treated with aN2 = 7400 and aN2 = 1700 in the

nitriding process, respectively.

4.2.1 XRD analysis

Segments of XRD results (2θ from 36° to 52°) observed from the 316L sample treated by the two different nitrocarburizig runs are presented in Figure 4.12 (a). Two

expanded austenite peaks ((111) and (200)) are presented for both samples. Precipitate formation was investigated by grazing angle XRD (Figure 4.12 (b)). Clearly, the precipitate formation was largely suppressed by lowering the nitriding activity in the second process. Both nitride (MN) and ω-carbide (M7C3) peaks were detected on the

sample 20C+20N (H). However, no clear extra peak was detected on the other sample

20C+20N (L). This implies that the precipitates (including both nitrides and carbides)

were formed during the second (nitriding) process or on cooling - otherwise, a lower

nitriding activity in the second process would not impact the formation of carbide.

128 The calculated ahkl from both samples are demonstrated in Figure 4.12 (b). Generally, the lattice parameters measured on the sample treated with the lower nitriding activity

(20C+20N (L)) are relatively smaller for (111), ((220), (222) and (311) planes. However, a200 measured from both samples are almost equal. This is probably due to the partially released stress on the sample treated with the higher nitridinng activity (20C+20N (H)) by crack and precipitate formation. The average lattice parameters calculated from thoose 7 peaks are 0.3792±0.0083 nm for 20C+20N (L) and 0.3804±0.0078 nm for 20C+20N (H).

(a)

129

(b)

Figure 4.12 Results observed from two different 20C+20N processes: 20C+20N (H) and 20C+20N (L) (a) XRD; (b) Grazing angle XRD; and (c) ahkl pplots

130 Table 4.2 Lattice parameters and lattice parameter expansions measured from the two different 20C+20N samples

Lattice Lattice Sample Standard Error Parameter (nm) parameter expansion

20C+20N (H) 0.3804 0.0078 5.7 %

20C+20N (L) 0.3792 0.0083 5.4 %

Non-treated 0.3598 0.0002 ---

4.2.2 Metallography

The case layers formed on the two nitrocarburized samples were revealed by an etchant of 50 vol.% HCl+25 vol.% HNO3+25 vol.% H2O. The SEM photographs are presented in Figure 4.13 (a) and (b), respectively. The original surfaces are located at the bottom of the images. The case depth produced was smaller when the lower nitriding activity was employed in the nitriding process. Also, the case depth shows a variation from grain to grain. This phenomenon is more obvious on the sample treated with the higher nitriding activity for the nitriding process (20C+20N (H)).

Figure 4.13 SEM images taken from 316L treated by (a) 20C+20N (L) and (b) 20C+20N(H).

131 4.2.3 AES depth profiles

AES cross-sectional line scans were performed on both nitrocarburized samples and two different lines were selected from each (Figure 4.14). Clearly, both samples present an anisotropy in case depth, which is consistent with the observation from the SEM images in Figure 4.13. The AES depth profiles also reveal that a slightly higher surface content of nitrogen was achieved by applying a higher nitriding activity in the second

process. At the same time, a larger case depth was yielded by a higher nitriding activity.

The scattered data points of carbon in nitrogen-enriched region in Figure 4.14 (c) might

be due to two reasons: (1) the poor signal-to-noise ratio in this area and (2) the precipitate

formation near the surface.

132 (b) 20C+20N (L) (a) 20C+20N (L)

Figure 4.14 AES depth profiles measured from nitrocarburizied 316L with 20C+20N (L) (a and b) and 20C+20N (H) (c and d).

4.2.4 Hardness measurements

The hardening effect was verified by a surface microhardness test with a load of 25 gf and is shown in Figure 4.15. When the activity was lowered in the nitriding process, the measured surface hardness was slightly lower. However, about 5 times larger hardness compared with non-treated 316L was still observed. The slightly lower surface hardness might be due to less precipitate formation.

133

20C+20N (a =7400) N2

20C+20N (a =1700) N2

Measurement

Figure 4.15 Surface hardness of 316L bulk sample treated with two different 20C+20N process

4.3 Effect of nitriding activity on 20 (N+C)

Two more simultaneous nitrocarburizing runs (20(N+C)) with reduced amount of

NH3 were designed to investigate the effect of nitriding activities on nitro-carburizing processes. Detailed parameters used are shown in Table 2.4. The total gas flow used was kept at 2 L/min for all 3 runs. When calculating the nitriding activity provided by the gas mixture, CO was treated as an inert gas. The activity of carbon in the nitrocarburizing atmosphere was calculated using commercial software, FactsageTM 5.3. The carbon activity provided is slightly promoted by using less NH3. In the following discussion, the processes will be identified by the content of NH3 used. For example, 20(N+C)-22.5 represents the run with 22.5 vol% of NH3.

134 4.3.1 XRD results

The XRD segments (from 36° to 52°) observed from the 3 simultaneously

nitrocarburized 316L bulk samples are shown in Figure 4.16. Generally, a larger peak

shift was observed when more NH3 was provided. Surprisingly, more second phases were observed on samples nitrocarburized with less NH3 (10 vol. % and 5 vol%), indicated by a stronger XRD peak from the second phases in the grazing angle XRD analysis (Figure

4.17 (b)). As discussed previously, the second phase in sample 20(N+C)-22.5 was not

identified because only one extra peak at 2θ=43.7° was detected. For sample 20(N+C)-10, two extra peaks at 2θ = 43.1° and 2θ = 56.6° were obtained, which made M2N nitride as

the most possible second phase formed.

Figure 4.16 XRD results obtained from 3 different simultaneous nitrocarburzing processes

135

(a)

(b)

Figure 4.17 Grazing angle XRD results obtained from (a) 20(N+C)-22.5 and (b) 20(N+C)-10

136 4.3.2 Metallography

Chemical etching was performed on those 3 simultaneously nitrocarburized samples,

after which a confocal microscope was employed for imaging. The original surface is

located at the top of the images. As demonstrated in Figure 4.18, relatively uniform case

layers were produced on 316L by using all 3 recipes. The sample nitrocarburized using

the recipe with 10% NH3 produced a relatively thicker case as compared to the other two.

According to the grazing XRD results shown in Figure 4.17 (b), this recipe also produced a considerable amount of second phases. The second phases were revealed by the Ar+ sputtering before AES acquisition. The SEM images taken are shown in Figure 4.19, which demonstrate that the formed second phase has a needle-like feature.

137

Figure 4.18 Cross-sectional images taken by confocal microscope from (a) 20(N+C)-22.5, (b) 20(N+C)-10 and (c) 20(N++C)-5

138

(a) (b)

Figure 4.19 SEM images demonstrating the needle-like second phase afftter 20(N+C)-10 (a) 3000X and (b) 10000X

4.3.3 AES depth profiles

The distribution of nitrogen and carbon interstitial atoms in the case layer was examined by AES cross-sectional line-scan. The profiles presented in Figure 4.20 demonstrate that surface concentration of nitrogen can be controlled by adjusting the NH3 content used in the simultaneous process. A surface nitrogen of about 18 at% was achieved with aN2 = 8600 (provided by 22.5 at% of NH3). When aN2 was reduced to 1700 and 430, the surface nitrogen content dropped to about 16 at% and 14 at%, respectively.

At the same time, a relatively deeper nitrogen-enriched case was observed when more

NH3 was used. However, the sample nitrocarburized with 10 vol.% NH3 demonstrated the thickest total case depth, which is consistently revealed by both metallographic images and AES results.

139 (a) 20 (N+C)-22.5%

(b) 20 (N+C)-10%

Figure 4.20 AES profiles measured from (a) 20(N+C)-22.5, (b) 20(N+C)-10 and (c) 20(N+C)-5

140

4.3.4 Hardening effect

The surface hardness was measured using a load of 25 gf and the results are shown in

Figure 4.21. The surface hardness decreased slightly when less NH3 was used. The value

measured from the sample treated with 10 vol.% NH3 is still a bit smaller as compared to

the one treated with 22.5 vol.% NH3, although more precipitates were formed. This

indicates that nitrogen is a strong hardener.

22.5% NH 3 10% NH3

5% NH3

Measurement

Figure 4.21 Surface hardness of 316L bulk sample treated with three different 20(N+C) process

141 References

[1]. F. Ernst, D. Li, H. Kahn, G. M. Michal and A. H. Heuer, Acta Materialia, 2011,

59 2268–2276

[2]. Y. Cao, Ph.D. Dissertation, Case Western Reserve University, 2003.

142 Chapter 5 Discussion

5.1 Effect of temperature

Similar to carbon, the solubility of nitrogen in 316L stainless steel can also be calculated by CALPHAD multi-sublattice modeling. The chemical potential of nitrogen in the metal could be written as:

= − = ( : −: ) + ( : −: ) + ( : −: ) + ( : −: ) + 1− + (1−2) :, + :, + :, + :, 1 + ,: +,:(Y −Y ) −,: 1 (5-1) −,: (Y −Y ) 1 + ,: +,:(Y −Y ) −,: 1 2 2 −,: (Y −Y ) −,: (Y −Y ) 1 +Y Y ,, −,, −,, (Y −Y ) + − ,: + ,: −,: 1 + [− ,: −,: (Y −Y )]+(1 −2)(Y Y L,:, +Y Y L,:, )

where Yi is the sites fraction of component i, ° is the molar free energy of element k in

the state of j, and the ,,:, , ,,:, and ,,:, are the interaction coefficients for components u,v,w,s and t. Assuming a paraequilibrium scenario, the solubility of nitrogen in the austenite can be calculated according to

1 = + = + (5-2) 2

143 For aN2 = 1 and aN2 = 7400, the solubility of nitrogen in 316L stainless steel as a

function of temperature was calculated and is shown in Figure 5.1. According to the

CALPHAD model, the nitriding process is exothermic and the solubility of nitrogen

decreases with increasing temperature, which is opposite to carburizing. Meanwhile, the

predicted solubility of nitrogen is much larger than that of carbon. For a unit activity of carbon, i.e., ac = 1, the maximum possible solubility of carbon in 316L is predicted as 11

at% at 723K [1], while the solubility of nitrogen is predicted to be about 22 at% at the

same temperature, in unit activity of nitrogen.

By comparing the experimental observation with the prediction based on CALPHAD

method, two conflicts were observed. (1) The surface concentration predicted by

CALPHAD model is much higher than what is observed from 316L bulk samples after

low-temperature nitriding processes. For example, the surface concentration predicted at

440 °C with aN2 = 7400 is 34.4 at%. However, the surface concentration measured from

AES is only about 17 at%, half of the predicted solubility. (2) The correlation between

surface concentration and nitriding temperature observed is opposite to the prediction.

According to CALPHAD, the surface concentration of nitrogen should decrease with nitriding temperature. However, the AES depth profiles (Figure 3.8) obtained from

samples nitrided with four different temperatures (350 °C, 420 °C, 440 °C and 450 °C) demonstrate that the observed surface concentration still increases with temperature. The conclusion is clear especially for the two samples treated at 450 °C and 350 °C. The

possible reasons leading to the difference between the solubility of nitrogen predicted by

CALPHAD and the surface nitrogen measured from nitrided 316L samples will be

discussed below.

144 (1) Adjustment might be required for the CALPHAD parameters employed. To

predict the solubility of carbon in 316L at low-temperature range, three Cr-C

parameters were adjusted [X. Gu [1]] to match the CALPHAD calculation with the

carbon content observed on low-temperature carburized samples. A similar

adjustment may be required for the system of (Fe, Cr, Ni, Mo), (N, V).

(2) The impact from residual stress on the solubility of nitrogen was not taken into

account. According to T. Christiansen and Somer [2], the existence of compressive

residual stress in the case layer causes an increase in the chemical potential of

nitrogen atoms in the solid solution. In other words, the occurrence of

compressive stress reduces the solubility of nitrogen. However, the influence of

pressure on CALPHAD interaction parameters is not well documented [1].

Therefore the effect of stress on solubility is difficult to be quantified.

(3) By using Equations 5.1 and 5.2 to predict the solubility of nitrogen, a local

equilibrium condition was assumed at the gas/metal interface for nitrogen atoms.

However, this assumption is probably unfulfilled. Our experimental results from

samples nitrided with different durations suggest that the equilibrium was not

established at least in the first 20 hours at 440 °C. Low mass transfer efficiency at

the metal/gas interface is one of the reasons delaying the achievement of

equilibrium, especially at low-temperatures. The mass transfer must be a stronger

function of temperature than the solubility (and in the opposite direction), so that

the surface nitrogen content observed on 316L still increases with increasing

temperature.

145

aN2=7400

aN2=1

Figure 5.1 Solubility of Nitrogen in 316L stainless steel under paraequilibrium condition at different temperatures.

5.2 Effect of nitriding activity

The solubility of nitrogen as a function of nitriding activity at 440 °C was also

calculated by the CALPHAD multi-sublattice modeling. The results are shown in Figure

5.2, together with the surface concentration measured from 316L samples nitrided with

different nitriding activities at 440 °C. Interestingly, both the CALPHAD predicted

solubility and the AES measured surface nitrogen content can be empirically fitted into a parabolic relationship as a function of ln aN2:

=() + () +

146 where k1, k2 and k3 are fitting parameters. The values are shown in Figure 5.3. The values of k1 are equal for both plots. k3 is consistent with XN att aN2 = 1 in both cases. The physical meanings of those parameters are unclear yet, but this at least indicates that the trend on relationship between XN and ln aN2 prediicted by CALPHAD method is correct.

Similar to what we observed in the previous section, the solubility predicted by

CALPHAD method is much higher than what is observed. Actually, for the two samples nitrided with lower activities (aN2 = 1 and aN2 = 200), the predicted solubility values are roughly 3 times (3.2 and 2.8, respectively) larger as compared to the experimentally observed XN. For the other four nitriding activiities, the predict solubility values are all about 2 times larger as compared to the experimentally obsserved XN. A possible reason caused the large difference at the low nitriding activity end is a smaller surface mass transfer coefficient.

Figure 5.2 CALPHAD predicted solubility of nitrogen in 316L stainless steel and AES measured surface nitrogen content on low-temperature nitrided 316L samples.

147 5.3 Lattice parameter expansion induced by nitrogen

Comparing the lattice parameter expansion induced by low-temperature nitriding and carburizing, it is apparent that nitrogen can expand the austenite lattice much more as compared to carbon. As reported [3,4], the lattice parameter expansion on stainless steels

can be as large as 10%, which is also observed in our study. In contrast, the lattice parameter expansion reported for the carbon-enriched austenite is normally about 3% [5].

However, the maximum concentration of nitrogen can reach 25~30% [6,7], which is also

much larger than that of carbon concentration (12~15%) [2]. As shown in Chapter 1,

Somers’s group [8] compared the lattice dilation of austenite lattice induced by

supersaturation of nitrogen and carbon interstitials. However, the nitrogen content in their

samples was much larger than that of carbon. No data was observed for nitrogen content

in the range of 0 ~15 at%. Therefore it is unclear which interstitial species dilate the

austenite lattice more at the same concentration.

To make such a comparison, the relationship between the surface interstitial content

and the corresponding lattice parameter should be built up. The plot of lattice parameter

(measured by XRD) versus nitrogen content (XN) achieved by the present study is shown

in Figure 5.3. In Figure 5.3 (a), the triangles and rectangles represent the lattice parameter calculated from (111) and (200) peaks, respectively. The circles in Figure 5.3 (b) represent the average lattice calculated from 7 different peaks ((111), (200), (220), (311),

(222), (400) and (331)) on the same sample. The error bar represents the standard deviation calculated by the n-1 method.

148 (a)

(b)

149 (c)

(d)

Figure 5.3 (a) Plots of a111 and a200 vs. XN measured by AES; (b) Plots of aavg. vs. XN; (c) linear fitting of aavg. vs. XN and (d) linear fitting of strain-free aavg. vs. XN

150 In Figure 5.3 (a), when both a111 and a200 are plotted against XN, it is obvious that a200 is always larger than a111. The difference becomes more obvious when the surface

concentration is larger than 15%. The difference between measured a111 and a200 is mainly

considered to originate from the compressive residual stress in the case layer. A lager

difference indicates a relatively higher level of residual stress. When the average lattice

parameter is plotted against XN, the story being told is similar. Also, it is noted that the

linear relationship between aAvg. and XN holds only for the four samples with relatively

low surface content of nitrogen. There is one sample with a surface concentration of ~16%

(440 °C aN2 = 1700) located in this linear region but demonstrating a large deviation. This

is possibly because of the development of ferromagnetic austenite phase in the <200>-

and <311>-oriented grains in that sample, which will be discussed in more detail in the

following section.

In Figure 5-3 (c), five data points (four from nitrided samples and one from non-

treated 316L) in the linear region were used to establish the dependence between the

lattice parameter and the nitrogen interstitial content by a least square linear regression

method. The fitted relation is given as:

. = (7.9 ± 0.6) ×10 + (0.3598 ± 0.0002)() (5.3)

with R2 = 0.9976.

The strain-free lattice parameters of the expanded austenite were estimated based on the method reported by H. Kahn [5]. For the four samples included in Figure 5.3 (c), the

residual stresses can all be estimated by the sin2ψ technique (Figure 3.4, Figure 3.11).

151 The X-ray elastic constants used in the calculation were taken from the same source,

which were calculated from the elastic constant for a Fe-18 wt% Cr-12 wt% Ni alloy

based on the Dewit approximation. The strain-free average lattice parameters aAvg, were calculated form the same 7 peaks and are plotted in Figure 5.3 (d). The fitted straight line is also shown, which is described by:

. = (6.7 ± 0.4) ×10 + (0.3598 ± 0.0002)() (5.4)

with R2 = 0.9975.

The dependence of lattice parameter on carbon content was investigated on low-

temperature carburized bulk 316L samples by our group and published in Ref. [5]. The

result is presented in Equation 5.5. The fitted straight line is described by the relation of:

. () = 6.8 × 10 + (5.5)

where . is the strain-free average lattice parameter of 316L after carburization and is the value measured from non-treated 316L. The XC values were estimated using a

least-square second-order polynomial fit of the XC measured by SAM, XPS and GDOES.

Apparently, Equations 5.4 and Equation 5.5 suggest that, when both XC and XN are

smaller than 15 at.%, the da/dXN/C factors for nitrogen and carbon in 316L are very

similar to each other. Actually, whether the same linear relationship holds for the whole

range of nitrogen content remains debatable.

152 Limitations and uncertainties certainly exist when trying to build up the correlation between lattice parameter and nitrogen content from our nitrided bulk-316L, which will be discussed as follows.

(1) Errors may exist in the measurements of nitrogen content by AES;

(2) The sampling depths for lattice parameter measurements and nitrogen content

measurements are different. XRD measures average lattice parameter from

several microns and this depth varies with crystalline planes. On the other hand,

the nitrogen contents are represented by the surface content of nitrogen read from

AES depth profiles, which comes from less than 1 micron depth. When

concentration gradient exists, the nitrogen content is overestimated.

(3) When the strain-free lattice parameters were calculated, stress gradient was not

considered.

(4) When estimating the residual strain in the case layer, the X-ray elastic constants

used were assumed to remain unchanged after nitriding, which may not be true.

5.4 Anisotropy in lattice parameter

As demonstrated in Figure 3.4, the anisotropy in XRD measured lattice parameter is substantial, especially when the sample was nitrided in a medium nitriding activity. As mentioned in Section 3.1, when substantial anisotropic ahkl is detected, the residual stress cannot be measured by the XRD sin2ψ technique due to the peak splitting during sample tilting. However, even if we use the largest residual stress reported in the literature

[9] (8GPa, by Somers ) to “correct” the XRD measured ahkl (as demonstrated in Figure 5.4,

153 XECs for Fe-18 wt% Cr-12 wt% Ni alloy employed), a fairly large difference still exists between a111 and a200 (Δa = a200 - a111 = 0.0076 nm). This implies thaat the tremendous anisotropy cannot be fully explained by the existing residual stress model. Lattice parameters measured from a thin 316 foil nitrided with the same condition (440 °C, aN2 =

7400, 20 hours) also support this hypothesis. As shown in Figure 5.5, scattering of data points measured from the foil sample is much smmaller as compared to bulk 316. This is because the residual stress in the foil sample is much smaller. By using the XRD sin2ψ technique, a tensile stress of about 0.8 GPa was measured. The strain-free ahkl (after correcting residual strain from lattice parameter measurements) were calculated and are plotted in Figure 5.5. Since the residual stress was tensile, the correction made the difference between a200 and a111 slightly larger, which was calculated to be 0.0076 nm.

Figure 5.4 XRD measured lattice parameters (red circles) and “correccted” lattice parameters (green diamond) assuming a residual stress of 8GPa from bulk 316L nitrided in aN2 = 7400 at 440 °C

154

Figure 5.5 XRD measured lattice parameters (red circcles) and strain-free lattice parameters (green diamond) from 316 thin foil nitrided in aN2=7400 at 440 °C

Besides residual stress, high density of stacking faultts was suggested as another reason for the anomalous distortion observed in the expandded austenite layer. However, as listed in Table 1.1, the fault parameter Ghkl for the first order reflections (e.g. (111),

(200)) and second order reflections (e.g. (222), (400)) are wiith opposite signs. Assuming a stacking fault density (α) equal to 0.2, the measured lattice parameter was “corrected” according to equation 1.2 and the result is shown in Figure 5.6. It can be seen that with the correction of stacking faults, the difference between a111 and a200 is reduced. However, a400 is brought even further away. This indicatess that the model of stacking faults cannot be used to explain the anomalous expansion in (200) planes.

155 Figure 5.6 XRD measured lattice parameters (red circles) and “correccted” lattice parameters (green diamond) assuming a stacking faults density of 0.2, bulk 316L nitrided in aN2 = 7400 at 440 °C

Figure 5.7 Correlation of lattice parameter expansion (defined as (ahkl-a0)/a0) vs. nitrogen contennt

156 The hypothesis proposed here is that the transition from paramagnetic austenite to ferromagnetic austenite plays a role in the highly distorted lattice of nitrogen-enriched expanded austenite.

Surface content of nitrogen vs. lattice parameter expansion measured from 9 different nitrided bulk 316L samples are demonstrated in Figure 5.7. The nitrogen content is represented by the surface concentration read from AES line scans. It is worth noting that the plot can be divided into 2 regions, with a clear gap at 6~7 at.%, which is close to the critical lattice parameter expansion required for the AFM to FM transition [10]. Therefore,

the data points in the lower region of Figure 5.7 represent lattice parameter measured

from NM or low-spin AFM state grains and those in the upper region represent lattice

parameter measured from high-spin FM state grains. The gap between those two regions

indicates that an abrupt change in lattice parameter is associated with the transition to

high-spin FM state. As a result, samples (labeled as 1, 2, 3 and 4) containing both

NM/AFM and FM grains demonstrate tremendous anisotropy. For example, for the

sample nitrided in aN2 = 1700 at 440 °C (blue diamond in Figure 5.7), the lattice

parameter expansion varies from 3% to 8%, depending on h,k and l. Our EBSD and

MFM investigations (Section 3.2) have revealed that sample 1 (blue diamond, nitrided in

aN2 = 1700 at 440 °C) does contain grains with different magnetic states. However, it is

impossible to identify AFM and FM states using MFM technique. Further investigation is

required to support our hypothesis.

157 Two possible explanations to explain the abrupt change in lattice parameter are: (1)

The lattice parameter is a function of magnetic moment; and (2) An even higher elastic

anisotropy is obtained when the nitrided case layer has transferred to ferromagnetic state.

An empirical linear relation between the lattice constant and the magnetic moment

was built up for binary solid solutions of 3d transition metals [11], which implies that the

magnetic moment has a direct contribution to the lattice parameter of solid solutions. The

relation can be described as

() =(1−) +() +<|| > (5.6)

where a(x) is the lattice constant of solid solution A1-xBx, x is the atomic fraction; aA and aB are lattice parameters of pure metal A and B, respectively; C is fitting parameters and

< || > is the average magnitude of the atomic magnetic moments. Their study on Fe-Co

(bcc), Fe-V (bcc), Fe-Al (bcc) and Fe-Ni (fcc) binary systems has revealed that the fitting

parameter C has a roughly constant value of 0.003 ~ 0.004 nm/μB. According to this

theory, the lattice parameter of nitrogen-enriched expanded austenite is a function of

nitrogen content (XN), residual stress level (σ) and magnetic state (|μ|). The investigation

on the magneto-volume effect on epitaxial γ-Fe ultrathin films and γ-Fe precipitates, W.

Keune et al. [12] concluded that a transition from antiferromagnetic low-spin fcc-Fe to

ferromagnetic high-spin fcc-Fe happened at a critical Wigner-Seitz radius of 2.69-2.70

a.u, which corresponds to a lattice parameter expansion of 5.9%~6.3%. An abrupt

increase in the magnetic moment during the transition from NM/AFM state to FM state

was both theoretically predicted and experimentally proved. Accordingly, a transition

from NM/AFM state to high-spin FM state could accompany an abrupt change in volume.

158 The changes in elastic constants associated with the changes in magnetic states are

reported in the literatures [15-17]. Further investigation is required to examine the possible

changes in elastic constants (C11, C12 and C44) in our nitrided samples, especially the ones

demonstrating transition to ferromagnetic state, to support the hypothesis that the change

in elastic anisotropy is the origin of the large distortion in lattice parameters.

5.5 On the nitrogen content measured by AES

As mentioned in Chapter 2, a “calibration” procedure was done on certified alloy samples (provided by the Swagelok Company). However, the concentrations of nitrogen in the alloys being tested were small (0.9% and 3%). Therefore the accuracy of AES

measurements of nitrogen at the high end is questionable. At the same time, a matrix

effect in the AES analysis on nitrogen-enriched 316L was noticed, which could also bring errors in the measurements of nitrogen by AES.

The concentration-depth profiles of the normalized metal elements based upon quantification of AES spectra are shown in Figure 5.8 (a) and (c), in which the nitrogen/carbon interstitial atoms are excluded. The nitrogen and carbon concentration

profiles measured on the same sample are shown in Figure 5.8 (b) and (d), respectively.

The average of the normalized composition of metal elements measured from the core

region agrees well with the normalized nominal composition of 316L. For example, the

average composition measured form the core part of image (c) is: Cr = 18±0.4, Fe =

69±0.5, Ni = 11±0.4 and Mo = 2±0.5.

159 However, in the nitrogen-enriched case layer, the concentration of chromium measured by AES shows a dependence on nitrogen content: the higher the nitrogen

content, the lower the chromium content measured. Two possible reasons might yield this lower chromium concentration in the nitrogen-enriched case as compared to the core. (1)

It is caused by an outwards diffusion or depletion of chromium; (2) the existence of

nitrogen impacts the detection of chromium by AES, which indicates the “Matrix Effect”.

The assumption of diffusion of chromium can be excluded according to the small

diffusion distance predicted by Fick’s second law. The diffusion distance of chromium was estimated to be only about 10 nm after 20 hours at 440 °C[13]. However, the layer

where the chromium concentration was lower than the core is as thick as 9 μm in Figure

5.8 (c). Therefore, the “matrix effect” is a more reasonable explanation. In other words,

the yield of AES electrons from the Cr element is impacted by the existence of nitrogen

atoms. It is worth noting that the carbon element does not demonstrate such an effect. As

shown in Figure 5.8 (d), normalized chromium in the carbon-enriched layer (10-30 μm) is

still consistent with the core region.

o o (b) T = 440 C, a = 200 (a) T = 440 C, a = 7400 N2 N2

160 o o (c) T = 440 C, 20C+20N (d) T = 440 C, 20C+20N

Figure 5.8 Normalized metal composition measured from (a) nitrided 316L and (c) nitrocarburized 316L. The corresponding nitrogen/carbon profiles are shown in (b) and (d), respectively.

5.6 XRD peak splitting during stress measurements

As mentioned in Chapter 3, due to the peak splitting during sample tilts, the residual stress in nitrided samples demonstrating large anisotropy in the “apparent” lattice

2 parameters ahkl could not be measured by a sin ψ technique. A hypothetical model is proposed here to explain the observed peak splitting. In the following discussion, bulk

316L nitrided at 440 °C in aN2 = 7400 for 20 hrs is taken as an example.

Assuming that (311) peak of the expanded austenite is used to estimate the residual stress, only <311>-orientated grains will reflect the incident X-ray and contribute to the

(311) peak at ψ = 0° (Figure 5.9 (a)). When the sample is tilted to ψ = 25.2° (Figure 5.9

(b)), <311>-orientated grains will not fulfill the Bragg condition, but other grains will, such as <200>-oriented and <212>-oriented grains. This is because 25.2° is the angle between {311} planes and {200} or {212} planes. In other words, when the sample is tilted to ψ = 25.2°, {311} planes in grains with their surface normals parallel to <200> and <212> directions will contribution intensity to (311) peak.

161 In XRD stress analysis, the “apparent lattice” parameter (,) measured at a random tilting angel ψ is determined by

1 =(1 + 2 ) (5-7) , , 2

hkl where σb is the biaxial stress, S2 is the XECs and , is the lattice parameter measured at 0° tilt. Accordingly, the lattice parameter measured from {200}-oriented grains at ψ =

25.2° is given by

1 =(1 + 2 ) (5-8) , , 2

200 -1 , was measured to be 0.3939 nm and S2 was calculated to be 0.0185 GPa [X].

If the stress is assumed as -5GPa, then , is calculated to be 0.3889 nm at ψ = 25.2°, which corresponds to 2θ = 82.12° for (311) peak.

Similarly, the lattice parameter measured from {212}-oriented grains can be calculated. However, (212) reflection of the expanded austenite was not detected at ψ = 0°

due to the selection rule of FCC lattice. Nevertheless, with the existence of stress, , is a function of X-ray elastic constant S1hkl:

, = (1 + 21 )

where is the strain-free lattice parameter of (hkl) plane.

162 According to Table 5.1, S1hkl values calculated for {111}, {220} and {212} planes are

similar. Therefore it is reasonable to assume , is similar to , and , . Then ,

is calculated to be 0.3761 nm at ψ = 25.2° (based on the assumptions of , = , and

σb = -5GPa), which corresponds to 2θ = 87.57° for (311) peak. Clearly, peaks reflected from <212>-oriented grains will be separated frrom those reflected from <200>-oriented grains and this hypothesis explains the peak splitting during sample tilts.

(a)

(b)

Figure 5.9 Schematics of stress measurements at (a) 0° tilt and (b) 25.2° tilt

163 hkl [5] Table 5.1 Calculated XECs S1 and XRD measured ,

hkl -6 -1 hkl S1 (×10 GPa ) , (111) -990 0.3775 nm (212) -1110 ? (220) -1271 0.3767 nm

5.7 Orientation-dependent case depth after low-temperature nitriding

A combination of EBSD orientation mapping and AES cross-sectional line scans

make it possible to study orientation-dependent nitrogen concentration profiles from

polycrystalline 316L samples. If regular plan-view profiling techniques were employed,

single crystals are required for such kind of investigation. The general observation is that

<100>-orientated grains demonstrate a higher surface nitrogen content and deeper case

depth as compared to <111>-oriented grain. This is consistent with the results reported by

A. Martinavičius et al [14], in which study single crystal crystalline austenitic stainless

steel AISI 316L specimens were used. By simulating the nitrogen diffusion profiles

during ion nitriding and the afterwards annealing, they concluded that the orientation-

dependent diffusion of nitrogen is related to an ion irradiation effect during nitriding.

However, the observed orientation-dependent case depth in the present gas-phase nitriding process suggested that this phenomenon is not caused by the ion irradiation effect. The results on gas-phase carburized 316L showed that the difference in case depth produced by carburizing process is much smaller. This excludes the hypothesis of preferentially activated surface by HCl gas.

164 = N exp(N) . (5.8)

As demonstrated in Figure 5.10, by employing a concentration dependent diffusion model, AES profiles measured from both <100> and <111> were successfully simulated

-15 2 -1 using the same DN (3.6×10 cm s ) and K(48). (done by Dr. H. Kahn). These results indicated that the difference in case depth is only caused by an orientation-dependent surface nitrogen content.

Figure 5.10 AES profiles measured from <111> and <100> grains, together with numerical simulation conducted by Prof. Kahn

165 References

[1]. X. Gu, Master Thesis, Case Western Reserve University, 2008.

[2]. T. L Christiansen, K. V. Dahl and MAJ Somers, Materials Science and

Technology, 2008, 24, 159-167

[3]. Y. Sun, E. Haruman, Vacuum 81 (2006) 114

[4]. T. Christiansen, M. A. J. Somers, Metallurgical and Materials Transactions A 37

(2006) 675

[5]. H. Kahn, G. M. Michal, F. Ernst and A. H. Heuer: Metallurgical and Materials

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[6]. K L Dahm and P A Dearnley: Proc Instn Mech Engrs Vol 214 Part L, 2000, 214,

181-198

[7]. S. Parascandola, W. Moller and D. Williamson: Appl. Phys. Lett., 2000, 76,

2194–2196

[8]. T. L. Christiansen, T. S. Hummelshoj and M. A. J. Somers, Surface engineering,

2010, 26, 242-247

[9]. T. Christiansen, M.A.J. Somers; E-structure Struers, 2006, 9, 1-17.

[10]. C. S. Wang, B. M. Klein and H. Krakauer: Physical review Letters, 1985, 54,

1852-1855

[11]. M. Shiga, AIP Conf. Proc No.18, 1974, 463-477

[12]. W. Keune, T. Ezawa, W. A. A. Macedo, U. Glos, K.P. Schletz and U.

Kirschbaum: Physica B, 1989, 161, 269-275

[13]. G.M. Michal, F. Ernst, H. Kahn, and A.H. Heuer: Acta Materialia, 2006, 54,

1597–1606

166 [14]. A. Martinavičius, G. Abrasonis, W. Möller, C. Templier, J.P. Rivière, A.

Declémy, and Y. Chumlyakov: J Appl Phys, 2009, 105, 093502

[15]. G. A. Alers, J. R. Neighbours and H. Sato, J. Phys. Chem. Solids Pergamon Press,

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167 Chapter 6 Conclusions

Nitrogen-supersaturated expanded austenite layers have been successfully produced

on 316L stainless steel samples by low-temperature gas-phase nitriding processes.

According to AES cross-sectional line scans, the surface nitrogen content varies from 7 at. % to 25 at. %, depending on the processing parameters employed. The effects of temperature, nitriding activity and nitriding duration have been systematically investigated. The general rules observed on nitriding process control include:

(1) The surface nitrogen content increases with increasing temperature, increasing nitriding activity and extended duration. Nitriding activity has the strongest influence among the three parameters.

(2) A thicker case depth can be achieved by increasing temperature, increasing nitriding activity and extending duration. A case about 10 times thicker was achieved

when the nitriding temperature was raised from 350 °C to 450 °C (aN2 = 7400). A similar

increment was observed when the nitriding activity was increased from 1 to 4×109 (at 440

°C). Extending duration is the least effective approach for increasing case depth.

Approximately, the case thickness is proportional to the square root of time.

(3) The risks of employing higher temperature, higher nitriding activity and longer duration are cracks and second phase formation. Nitriding temperature higher than 440

°C is not recommended for 316L stainless steel.

Large lattice parameter expansion is associated with the supersaturation of nitrogen

interstitials in the austenite lattice. The lattice parameter expansion measured from the

168 nitrided bulk 316L samples varies from 1% to 10%. A comparison was made between

nitrogen and carbon in the lattice parameter expansion of austenit. The conclusion is that their dilating capabilities are similar when their concentrations are both lower than 17 at.%.

Room temperature ferromagnetism was observed on 316L samples after low- temperature gas-phase nitriding. A combined study of XRD, MFM, EBSD has revealed that the minimum lattice parameter expansion required for the ferromagnetic behavior is

~ 5%. A nitrogen content of ~ 14 at.% is estimated (by AES) as a threshold required for the paramagnetic-to-ferromagnetic transition. TEM diffraction has demonstrated that the induced ferromagnetism is not due to the ordering of nitrogen atoms in the austenite lattice.

The case depth produced on 316L after low-temperature nitriding process showed a strong dependence on surface orientation. In particular, a relatively deeper case was measured from <100>-oriented grains as compared to <111>-orientated grains. AES line scans on pre-indexed (by EBSD orientation map) 316 samples demonstrated the <100>- oriented grains had a higher surface nitrogen content. The difference in surface nitrogen content is considered as the origin of the orientation-dependent case depth.

9 After 316L sample was nitrided at 440 °C with aN2 = 4×10 for 20 hours, a γ’-M4N structure was discovered in the nitride layer by TEM diffraction. Many surface particles were created on the surface, which were identified as ε-Fe2N1-x by TEM diffraction and

XEDS study.

The conclusions driven from the investigation of nitrocarburizing processes in gas

mixture of NH3/CO/H2/N2 are: (1) dual layers of expanded austenite were obtained from

169 all 3 different designs of nitrocarburizing scenarios; (2) regardless of the sequence of

diffusion, the outer layer of the case is always nitrogen-enriched; (3) the total case depth mainly depends on the diffusion time of carbon; and (4) both nitrides and carbides were observed when 316L were treated with 20C+20N recipe (with aN2=7400 for the latter

nitriding treatment). Using XRD and TEM, the formed nitrides were identified as MN

and the formed carbides were identified as M7C3. The formation of both precipitates was

suppressed when a lower nitriding activity was employed in the second process. This

implies that the precipitates were formed during nitriding or on cooling.

170 Appendices

Appendix I Nitro-carburizing by Urea (CO(NH2)2)

Using urea as a cheap source of nitrocarburizing process was reported by T.

Christiansen et al. in 2011[1]. It provides an alternative approaching of activating the

surface of ferrous and non-ferrous coupons before nitriding, carburizing and

nitrocarburizing. According to Ref. [1], a total thickness of about 10 μm was produced

after stainless steel AISI 316 article was heated from room temperature (RT) to 440 °C

within 45 minutes then cooled to RT in 10 min minutes. When the maximum temperature

was raised to 490 °C and the leading gas was changed from N2 to H2, the thickness of the

hardened case increased to about 20 μm. Those results indicated that nitrocarburizing in

urea is a surprisingly fast process.

A furnace (as shown in Figure A-1) was setup to execute nitrocarburizing

experiments with urea powders. Stainless steel 316 coupon was sealed together with urea

particles in a fused silica tube but with a distance. N2 was selected as the leading gas and

the gas flow was roughly controlled by a regulator connecting to the gas cylinder. Two

thermal couples were placed under the quartz tube to monitor the temperatures of 316

coupons and urea particles, respectively.

Upon the heating up of urea powders, a variety of products can be yielded. A typical

TG (thermogravimetry) curve [2] observed during the decomposition of urea powders

were demonstrated in Figure A-2. The reactions involved in the thermal decomposition of

urea powders includes [2, 3]:

171 2 () → +

2 () → + ( )

3 () →3 +

3 → 3 +2 ( )

+ → +2

At temperatures between 500 °C and 600 °C, it is believed [4] that urea can decompose into CO and nascent nitrogen atoms, which promote the carburizing and nitriding processes, respectively.

() →+ 2 + 2

Figure A-1 Furnace setup to execute nitrocarburing expperiment in ureea

172

Figure A-2 Thermal decomposition of Urea

Table A-1 Experimental parameters of nitrocarburizing processes in urea powders

Heating Rate Soaking Time Maximum Temp. Case Thickness Recipe # (°C/min) (min) (°C) (μm)

8~11 (bottom) U1 20 30 min 550 2-7 (top)

U2 10 0 min 550 7~10

U3 10 30 min 550 17~30

173 Three recipes were designed to investigate nitrocarburizing processes in urea powders.

Detailed parameters are listed in Table A-1. The maximum temperature was selected as

550 °C for all of the three experiments. A fast cooling was achieved by opening the lid of

the furnace (as shown in Figure A1). It took about 60 min for the furnace to cool from

550 °C to 100 °C.

XRD results observed from samples treated with the three different recipes are shown

in Figure A-3. For sample U1, different results were obtained from top and bottom

surfaces and that observed from bottom surface was selected to present. XRD result of

non-treated sample was given as a and labeled ad NT. Accordingly, expanded austenite

layers were successfully produced by all three treatments. Larger lattice parameter

expansion was achieved when a slower heating rate was applied. Nitrides peaks (M2N) were detected only from sample U3.

XRD results obtained from the bottom and top surfaces of sample U1 were found different by means of lattice parameter expansion (as demonstrated in Figure A-4).

Generally, lattice parameter expansion measured from the top surface were smaller than that measured from the bottom surface. The smaller expansion is associated with lower surface concentration of interstitial atoms, which was possibly resulted from a nonuniform surface activation. Meanwhile, the presence of original austenite peaks implies that the case depth produced on the top surface is shallower.

The thickness of the case layer produced was revealed by an etchant of 50 vol.%

HCl+25 vol.% HNO3+25 vol.% H2O. Optical microscope was used for imaging and the

pictures obtained are displayed in Figure A-5. Very thin layer of case was observed from

174 the top surface of sample U1 (Figure A-5 (a)). The case layer obtained from the bottom

of the same sample was much thicker (Figure A-5(b)). Those results are consistent with

the XRD results shown in Figure A-4. , The truth of a better activation of the bottom

surface implied that the surface activation process was accomplished by liquid products.

Effect of heating rate was demonstrated by comparing sample U1 and U3. As

indicated in Table A-1, the only difference between those two samples was the heating

rate. By decreasing the heating rate from 20 °C/min to 10 °C/min, a much thicker case

layer was observed (Figure A-5 b & d). A possible explanation is that the decomposition

of urea powders was more complete when a smaller cooling rate was selected, so that a better surface activation was achieved. However, U3 recipe (10°C/min, 30 min soaking at

550 °C) was proven too strong for 316 because of the formation of nitrides (the dark

surface layer in Figure A-5 (d). Without the 30 min soaking segment (sample U2), the

nitride formation was suppressed. The case depth produced was reduced to about 7-10

μm.

175

Figure A-3 XRD results from the three different nitrocarburizzing processes in urea.

Figure A-4 XRD results from the bottom and top surfaces of sample U1.

176 Figure A-5 Optical microscopy images observed from samples nitrrocarburized in urea powders.

177 Appendix II Nitro-carburizing in NH3/C2H2/H2/N2

As discussed previously, austenitic stainless steel samples were able to be

nitrocarburized with gas mixtures of NH3/CO/H2/N2. However, a separate HCl activation

procedure is required to remove the native oxide and allow the inward diffusion of

nitrogen and carbon atoms. The application of those recipes is limited when the treated

surfaces are sensitive to HCl gas and therefore efforts were made to develop HCl-free

nitrocarburizing recipes for austenitic stainless steel samples.

Recently, study in our group (Y. Ge) has demonstrated that the HCl activation procedure can be eliminated if the carburization is performed in gas mixtures of

C2H2/H2/N2. An “in-situ” activation is achieved because of the very low partial pressure of oxygen (PO2) provided by a mixture of C2H2/H2/N2. Simultaneous nitrocarburizing

recipes can be designed by simply adding NH3 gas into the carburizing atmosphere. A

trail run was executed and the processing parameters employed are listed in Table A-2.

316L stainless steel samples with different surface finishing were involved in the same

run. EP indicated that the surface of that sample was electropolished (by Swagelok

company) before nitrocarburizing. Similarly, MP means the surface was mechanically

polished before treatment. MP-1μm represents the finest polishing was done with 1μm

diamond suspension. MP-P2500 represents the finest polishing was done with P2500

silicon carbide polishing paper.

XRD analysis was employed to verify the formation of expanded austenite layer. As

shown in Figure A-6, expanded austenite layers were observed after 316L samples were carburized in gas mixture of NH3/C2H2/H2/N2. The average lattice parameter expansion

178 were estimated to be 3.3% for the EP sample and 2.9% for the two MP samples. In Figure

A-6(a), an original austenite peak at 2θ=43.7° was also detected. This implies a

nonuniform surface activation by NH3/C2H2/H2/N2. There must be surface areas with

very thin layer of cases duo to a slower surface activation. No clear original austenite peak was observed from the two MP samples. Although it doesn’t necessary mean that a better surface activation was achieved on MP samples. It is possible that the not successfully activated areas on the MP samples were not included in the region analyzed by XRD.

The case layers formed were revealed by an etchant of 50 vol.% HCl+25 vol.%

HNO3+25 vol.% H2O. The optical microscopy images observed from cross-sectional

samples are shown in Figure A-7. Continuous case layers were observed on all of the

three samples even though an original austenite peak was detected by XRD from EP

sample. This is possibly due to the limited area being viewed from one cross-sectional

sample. As shown in Figure A-7, the case layers produced can reach 20 μm for both EP

and MP-1μm sample. The case obtained on sample MP-P2500 is relatively shallower.

Cross-sectional AES analysis was performed on sample MP-1 μm and the obtained

depth profiles of nitrogen and carbon are plotted in Figure A-8. Nitrogen and carbon

atoms are coexisting in the first 5 μm of the case layer. The surface concentration of both

carbon and nitrogen interstitials are about 8 at.%, which indicated that the activity of

nitrogen provided by NH3/C2H2/H2/N2 is fairly low. The highest carbon concentration is

about 14 at.%, detected at the diffusion front of nitrogen. Again, this can be explained by

that the chemical potential of carbon is increased by nitrogen when they are coexisting.

179 Table A-2 Experimental parameters of simultaneous nitrocarburizing processes in NH3/C2H2/H2/N2

Temperature NH3 C2H2 H2 N2 Time

440 oC 0.2 L/min 0.05 L/min 0.9 L/min 4.8 L/min 16 hours

180

Figure A-6 (a) XRD results from 316L samples with different surface finishing after nitrocarburized in NH3/C2H2/H2/N2 at 440 °C for 20 hours. (b) Enlarged segments of the XRD results in image (a)

181

Figure A-7 Optical microscopy images observed from 316L samples nittrocarburized in NH3/C2H2/H2/N2

Figure A-8 AES depth profiles of carbon and nitrogen measured from sample MP-1μm (after nitrocarburized in NH3/C2H2/H2/N2 at 440 °C forr 16 hours).

182