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The development and performance of chromium/ reactive element-modified aluminide diffusion coatings by chloride-activated pack cementation

Bianco, Robert, Ph.D.

The Ohio State University, 1992

UMI 300 N. Zeeb Rd. Ann Arbor, MI 48106

THE DEVELOPMENT AND PERFORMANCE OF CHROMIDM/REACTIVE ELEMENT-MODIFIED ALUMINIDE DIFFUSION COATINGS BY CHLORIDE- ACTIVATED PACK CEMENTATION

DISSERTATION

Presented in Partial Fulfillment of the Requirements for the Degree Doctor of Philosophy in the Graduate School of The Ohio State University

by

Robert Bianco, B.S., M.S.

*****

The Ohio State University 1992

Dissertation Committee: Approved by R.A. Rapp S.A. Dregia Advisor G.W. Powell Program in Metallurgical Engineering To Susan, the best is yet to be!

11 ACKNOWLEDGMENTS

I wish to thank Prof. Robert A. Rapp for his assistance and encouragement during this investigation and for giving me the opportunity to attend and present this study at several technical conferences in my field. I would also like to thank the financial support from the Office of Naval Research (A.J. Sedriks) and the Naval Air Development Center (T.A. Kircher). Technical assistance and coating evaluation from the NASA Lewis Research Center especially Drs. James L. Smialek and Nathan S. Jacobson, who initiated me into this challenging field, is gratefully appreciated. In addition, the assistance of Dr. Rex Hussey of the National Research Council, Ottawa, Canada with the SIMS study is also recognized. I would like to offer a special thanks to the members of my dissertation committee: Profs. Dregia and Powell. Their comments and criticism were quite helpful in clarifying this document. I wish to thank my fellow graduate students for their friendship and generosity: Mike Antony, Tom Broderick, Dan Evans, Andy Gordon, Karthick Gourishankar, Brent Harle, Chungwei Hsia, Patrick LeBlanc, Greg Meszaros, Jim Moore,

iii Jon Papal, John Naizer, and Craig Miller. I would also like to thank the members, past and present, of the high temperature coatings group; Denis O'Connell, Ed Naylor, Fred Geib, Brian Cocheram, Mark Harper, and Drs. Choquet, Ravi, Wang, and Strawbridge. I wish them all much future success. A special thanks to my good friends: Jeff Moser, Drew Mueller, and Dave Miller; whose early morning golf outings will never be forgotten. The assistance of Dorothy Savely, Clare McDonald, David Little, Jack Kocka, Gary Dodge, and Bud Farrar has been essential and greatly appreciated. Finally, I wish to thank my parents and family for their constant love and encouragement throughout my life. Without their support, all this work and effort would not have been possible. For this, I am eternally grateful and indebted. Lastly, I wish to thank the most important person in my life, my wife Susan. Her love and support will always be remembered. The best is yet to be!

IV VITA June 3, 1966 Born-Indiana, Pennsylvania December 1988 B.S. Metallurgical Engineering The Ohio State University Columbus, Ohio. August 1991 M.S. Metallurgical Engineering The Ohio State University Columbus, Ohio. August 1991-present Graduate Research Associate Department of Materials Science and Engineering The Ohio State University Columbus, Ohio.

PUBLICATIONS *R. Bianco, J. Electrochem. Soc.. 136 (1989) 265C. *R. Bianco and N.S. Jacobson, J. Mater. Sci.. 24 (1989) 2903. *R. Bianco and R.A. Rapp, "Simultaneous Chromizing- Aluminizing of Nickel-Base Superalloys with Reactive Element Additions," in High Temperature Chemistrv V . W.B. Johnson and R.A. Rapp (Eds.), The Electrochemical Society, Inc., Pennington, NJ (1990) 211. *R. Bianco, M.A. Harper, and R.A. Rapp, J. Metals. 43 No. 11 (1991) 68.

V FIELDS OF STUDY Major Field: Metallurgical Engineering Studies in: Chemical Metallurgy R.A. Rapp, G.R. St. Pierre, Y. Sahai, E. Nygren Physical Metallurgy P.G. Shewmon, G. Meyrick, S.A. Dregia Mechanical Metallurgy R.H. Wagoner, P. Anderson, G.S. Daehn Corrosion R.A. Rapp, B.E. Wilde, S. Smialowska Solidification C.E. Mobley Solid-State Physics R. Sooryakumar Electron Microscopy W.A.T. Clark

VI TABLE OF CONTENTS

ACKNOWLEDGMENTS...... Ü i

VITA ...... V LIST OF TABLES ...... ix LIST OF FIGURES...... xiii ABSTRACT ...... xxiii CHAPTER PAGE I. INTRODUCTION...... 1 II. LITERATURE SURVEY ...... 14 2.1 Halide-Activated Pack Cementation Diffusion Coatings...... 14 2.1.1 Aluminizing...... 14 2.1.2 Chromizing...... 26 2.1.3 Siliconizing...... 29 2.1.4 Deposition of a Reactive Element. . . 31 2.1.5 Codeposition...... 33 2.2 Thermodynamics and Kinetics of Pack Cementation...... 37 2.2.1 Single Element Deposition ...... 37 2.2.2 Codeposition...... 49 2.3 Degradation of Pack Cementation Diffusion Coatings...... 54 2.3.1 Interdiffusion...... 55 2.3.2 Oxidation...... 56 2.3.3 Fused Salt Attack (HotCorrosion) . . 63 2.4 Reactive Element (RE) Effect...... 68 2.4.1 Observations...... 69 2.4.2 Mechanical Behavior ...... 71 2.4.3 Chemical Behavior ...... 72 2.4.4 New Theories...... 75 III. EXPERIMENTAL PROCEDURES ...... 79 3.1 Materials ...... 79 3.2 Diffusion Coating Procedures...... 82 3.3 Characterization of the Diffusion Coatings...... 86

Vll 3.4 Analysis of Vapor Species Evolved from Pack Mixtures ...... 89 3.5 Environmental Durability Testing...... 91 3.5.1 Oxidation ...... 91 3.5.2 Hot Corrosion...... 92 IV. VOLATILE SPECIES IN CHLORIDE-ACTIVATED DIFFUSION COATING PACKS ...... 97 4.1 ITSOL Computer Program...... 97 4.2 Pack Equilibrium Calculations ...... 99 4.2.1 Sample Inputs ...... 99 4.2.2 Sample Outputs...... 106 4.3 ITSOL Calculations With a RE Source . . . .Ill 4.3.1 RE Oxide Source ...... Ill 4.3.2 RE Activator Salt ...... 118 4.3.3 Masteralloy Combination ...... 121 4.4 Mass Spectrometer Measurements...... 125 V. RESULTS AND DISCUSSION...... 134 5.1 Coating Development for Nickel-Base Alloys...... 134 5.1.1 Chromium/RE-Modified Aluminide Coatings...... 134 5.1.2 RE-Doped Aluminide Coatings ...... 160 5.1.3 Chromium/Silicon-Modified Aluminide Coatings...... 167 5.2 Kinetics and Formation Mechanisms of Al, Cr, RE C o a t i n g s ...... 175 5.2.1 Contacting Powder Arrangement . . . .175 5.2.2 "Above Pack" Arrangement...... 185 5.3 Environmental Durability of Aluminide Coatings...... 189 5.3.1 Cyclic Oxidation...... 193 5.3.2 Isothermal Oxidation...... 211 5.3.3 Hot Corrosion (Fused Salt (Attack) Studies...... 217 VI. SUMMARY AND CONCLUSIONS...... 233 VII. FUTURE WORK ...... 238 APPENDICES A. ITSOL CALCULATIONS: CONSTANT PRESSURE VERSUS CONSTANT VOLUME...... 239 REFERENCES ...... 242

Vlll LIST OF TABLES

TABLE PAGE 1. Summary of interdiffusion studies on the nickel-aluminum system ...... 18 2. Nominal compositions of nickel-base alloys coated (at.%)...... 80 3. Composition of the chromium-aluminum binary masteralloy powders...... 80 4. Properties of the chloride activator salts used in this investigation...... 81 5. Standard Gibbs energies of formation for the aluminum species used as input for the ITSOL program...... 101 6. Standard Gibbs energies of formation for the chromium species used as input for the ITSOL program...... 102 7. Standard Gibbs energies of formation for the zirconium species used as input for the ITSOL program...... 103 8. Standard Gibbs energies of formation for the yttrium species used as input for the ITSOL program...... 104 9. Standard Gibbs energies of formation for the silicon species used as input for the ITSOL program...... 105 10. Activities and corresponding Gibbs energies of formation of chromium and aluminum for various masteralloy compositions at 1200 and 1423K...... 107 11. ITSOL pack equilibrium output for a 2 wt.% NH.Cl-activated pack containing 25 wt.% Cr-lOAl wt.% masteralloy at 1150 C ...... 108

IX 12. ITSOL pack equilibrium output for a 2 wt.% NH.Cl-activated pack containing 25 wt.% Cr-lOAl wt.% masteralloy and 2 wt.% Y_0_ at 1150°C...... 112 13. ITSOL pack equilibrium output for a 2 wt.% NH.Cl-activated pack containing 25 wt.% Cr-lOAl wt.% masteralloy and 2 wt.% ZrO_ at 1150°C...... 113 14. ITSOL pack equilibrium output for a 2 wt.% NH.Cl-activated pack containing 25 wt.% Cr-lOAl wt.% masteralloy and 2 wt.% SiO at 1150°C...... f . . . .114 15. NH.Cl-activated pack with a Cr-lOAl wt.% maiteralloy at 927 C (1200K)...... 128 16. NH.Cl-activated pack with a Cr-lOAl wt.% masteralloy and 2 wt.% ZrO^ at 927 C (1200K) . .129 17. NH.Cl-activated pack with a Cr-lOAl wt.% masteralloy and 2 wt.% Y^O^ at 927 C (1200K) . .130 18. Chemical analysis of the condensate from an NH.Cl-activated pack with Cr-lOAl wt.% masteralloy and 2 wt.% Y_0_ held at 927 C (1200K) for 1 hour ...... 132 19. Summary of nickel-base alloys coated with a pack containing 2 wt.% NH.Cl plus 2 wt.% ZrOg at 1150 C for 24 hours: ...... 138 20. Summary of nickel-base alloys coated with a pack containing 2 wt.% ZrCl. at 1150 C for 24 h o u r s ...... 138 21. Summary of nickel-base alloys coated with a pack containing 2 wt.% YC1_ at 1150°C for 24 h o u r s ...... 139 22. Summary of nickel-base alloys coated with a pack containing 1 wt.% CrCl plus 2 wt.% of second activator and 25 wt.% Cr-7.5A1 wt.% masteralloy at 1150 C for24 h ours ...... 139 23. Summary of René 80 alloys coated in a pack containing 25 wt.% Cr-7.5A1 wt.% masteralloy at 1150 C for 24 hours ...... 154

X 24. Summary of nickel-base alloys coated with a pack containing 4 wt.% activator and 25 wt.% masteralloy at 1150 C for 24 hours ...... 154 25. Summary of IN 713LC and Mar-M247 alloys coated with a pack containing 25 wt.% Cr-lOAl masteralloy at 1150 C for 24 hours ...... 164 26. Summary of Mar-M247 alloys coated in a pack containing 4 wt.% NH.Cl activator, 20 wt.^ masteralloy, and a silicon source at 1150 C for 24 h o u r s ...... 164 27. Kinetic results of René 80 alloys coated in a powder contacting arrangement with a pack containing 2 wt.% activator and 25 wt.% Cr-7.5A1 masteralloy...... 182 28. Kinetic results of nickel-base alloys coated in an "above pack" arrangement with a pack containing 2 wt.% NH.Cl, 2 wt.% Y_0_, and 25 wt.% Cr-7.5A1 masteralloy...... 182 29. Summary of the XRD results of coated René 80 alloys cyclically oxidized in static air at 1100 and 1200 C for up to 200, one-hour cycles .192 30. Summary of the XRD results of René 80 and IN 713LC alloys coated in an "above pack" arrangement with a pack containing 15 wt.% Cr-7.5A1 masteralloy at 1150 C for 4 and 24 hours and cyclically oxidized at 1100 C for up to 500, one-hour c y c l e s ...... 201 31. Summary of the kinetic and XRDresults of RE-doped aluminide diffusion coatings on IN 713LC and Mar-M247 alloys oxidized in air at 1100 C for 44 hours and comparison with bulk 8-NiAl compounds...... 210 32. XRD results of coated René 80 and Mar-M247 alloy substrates corroded at 900 C for 672 hourSgin 0.1% SO /O^ gas mixture with 5.0 mg/cm Na„SO.. "Above pack" arrangement...... 221

XI 33. XRD results of coated René 80 alloy substrates isothermally corroded at 900 C for 144 and 672 hours in a 0.1% SO_/Og gas mixture with 5.0 mg/ cm Na^SO^. Powder contacting arrangement. . . .221 34. Summary of results for a René 80 alloy substrate treated at 1150 C for 16 hours in a pack containing 2 wt.% activator and 15 wt.% masteralloy and tested in a mach 1, burner rig at 927 C using JP-5 jet fuel with 5 ppm NaCl and 0.4% sulfur added...... 230 35. ITSOL pack equilibrium output for a 2 wt.% NH^Cl- activated pack containing 25 wt.% Cr-lOAl wt.% masteralloy at 1150°C (constant pressure). . . .240 36. ITSOL pack equilibrium output for a 2 wt.% NH.Cl- activated pack containing 25 wt.% Cr-lOAl wt.% masteralloy at 1150 C (constant volume)...... 241

Xll LIST OF FIGURES

FIGURE PAGE 1. Arrhenius plot of the parabolic rate constants of alumina, chroraia, and silica scales thermally grown on various heat-resistant alloys or coatings [9] ...... 7 2. A schematic illustration of the chloride- activated pack cementation process ...... 11 3. Binary nickel-aluminum phase diagram [29]. . . 15 4. Variation of the interdiffusion coefficients with composition in the B-NiAl phase at 1100°C [33]...... 19 5. Schematic illustration of the "high" activity process; formation of a NiAl coating from a NigAl_ layer by diffusion annealing in (a) pure nickel or (b) nickel-base superalloy [37] 21 6. Schematic illustration of the "low" activity process to produce an NiAl coating by outward diffusion on (a) pure nickel or (b) nickel- base superalloy [37] ...... 22 7. Microstructure and phase identities of a "high" activity aluminide coating on Udimet 700 followed by heat treatment for 4 hours at 1080 C; lOOOx [36]...... 23 8. Microstructure and phase identities of a "low" activity aluminide coating on Udimet 700; lOOOx [36]...... 24 9. Binary nickel-chromium phase diagram [39]. . . 28 10. Binary nickel-silicon phase diagram [29] . . . 30 11. Experimentally determined oxide map of the NiCrAl alloy or coating system [61]...... 35

Xlll 12. (a) Activator circulation and (b) activator condensation mechanisms of aluminum deposition proposed by Levine and Caves[65] where circles =A1(1) and triangles=AlF^ (c) [69]...... 42 13. Mixed mechanism of aluminum transport in the presence of both an activator and activator/ source depleted zones, where circles=Al(1) and triangles=AlF^(c) [69]...... 44 14. Theoretical variation in the parabolic deposition rate of aluminum, K , with activator content at 900 C for"a 4 wt.% pure Al pack activated by AIF^[69] ...... 45 15. (a) Pack aluminum concentration profile and (b) AIF vapor pressure profile in a NÎ-40A1 . at.% alloy pack after 8 hoursat 1085 C [73] . 47 16. Partial pressure of gaseous halides as a function of temperature in NaCl-activated packs containing pure aluminum, chromium, and silicon [72] ...... 51 17. Activity data as function of solute content for (a) aluminum and chromium in a Cr-Al alloy at 800, 900, and 1000 C[75] and (b) silicon and chromium in a Cr-Si alloy at 1050 C [76]...... 52 18. Equilibrium partial pressures of gaseous species in a (a) CrCl - and an (b) AIF - activated pack as a function of the aluminum activity in the Cr-Al masteralloy at 1000 C [ 7 7 ] ...... 53 19. The effect of interdiffusion on the cyclic oxidation of pack aluminized IN 100 at 1100 C for 1 hour cycles [81]...... 57 20. Compositional effects on the oxidation behavior of binary nickel-aluminum alloys. (a) Temperature-composition oxide phase map and (b) scale growth rates corresponding to regimes in (a) [34]...... 58

XIV 21. Microstructure, microhardness data, and and phase identities for pack aluminized Mar-M200 alloys after oxidation in air at 1200 C [88]...... 62 22. Dissolution and reprecipitation of a porous MO oxide supported by the solubility gradient in the fused salt film [100] ...... 65 23. Measured oxide solubilities in pure Na_SO. at 927 C (1200K) and 1 atm oxygen [101]. . . . 67 24. Parabolic rate constant for chromia growth on unimplanted (dashed lines) and cerium implanted Ni-30Cr wt.% alloys as a function of cerium dosage at 1000 and 1050 C [107]. . . 70 25. Schematic illustration of the oxide/scale and alloy/scale interface of an yttrium- doped, FeCrAl alloy after oxidation for 24 hours at 1200 C [19]...... 74 26. Effect of sulfur content on the 1180°C cyclic oxidation performance of NiCrAl alloys [125] ...... 76 27. Schematic diagram of the "powder contacting" pack arrangement...... 83 28. Schematic diagram of the "above pack" arrangement...... 84 29. Average heating profile of Lindberg furnace. . 85 30. Schematic diagram of the pack cementation furnace setup...... 87 31. Schematic diagram of (a) the atmospheric sampling mass spectrometer and (b) the target collection apparatus...... 90 32. Schematic diagram of the thermogravimetric analyzer (TGA) used for isothermal oxidation testing...... 93 33. Schematic diagram of the cyclic oxidation test apparatus ...... 94

XV 34. Schematic diagram of the thin film hot corrosion test setup ...... 95 35. Equilibrium partial pressures of the gaseous species in the bulk pack of an NH.Cl-activated pack as a function of aluminum activity in the Cr-Al masteralloy at 1150 C ...... 109 36. Equilibrium partial pressures of the gaseous species in the bulk pack of an NH.Cl-activated pack as a function of aluminum activity in the Cr-Al masteralloy at 1150 C with 2 wt.% additions of (a) Y,0_, (b) ZrO.,, and S i O g ...... 115 37. Equilibrium partial pressures of the gaseous species in the bulk pack of an (a) ZrCl^- and (b) YC1_-activated pack as a function of aluminum activity in the Cr-Al masteralloy at 1150°C...... 119 38. Equilibrium partial pressures of the gaseous species in the bulk pack of an NH.Cl-activated pack as a function of silicon activity in the Cr-Si masteralloy at 1150 C containing 20 wt.% of a (a) Cr-lOAl wt.%, (b) Cr-15A1 wt.%, or (c) Cr-20A1 wt.% masteralloy...... 123 39. Cross-sectional optical micrograph of a René 80 alloy substrate treated at 1150 C for 24 hours in a 2 wt.% NH.Cl-activated pack containing (a) 2 wt.% ZrO_ and 25 wt.% Cr-lOAl masteralloy or (b) 25 wt.% Cr-7.5A1 masteralloy. Powder contacting arrangement . . 136 40. Cross-sectional optical micrograph of a René 80 alloy substrate treated for 24 hours in a 2 wt.% (a) ZrCl.- or (b) YCl_-activated pack containing 25 wt.% Cr-7.5A1 masteralloy. Powder contacting arrangement...... 137 41. Cross-sectional electron micrograph of the 6-NiAl and a-Cr, two phase coating and the corresponding x-ray maps of nickel, aluminum, cobalt, and chromium; 9000x...... 141

XVI 42. Cross-sectional electron micrograph of the interdiffusion zone and the corresponding x-ray maps of nickel, aluminum, cobalt, chromium, tungsten, and molybdenum; 5000x. . . 142 43. (a) The 1150°C isothermal section of the NiCrAl ternary system, and (b) the B-NiAl-Cr quasi-binary phase diagram [146] ...... 143 44. Cross-sectional optical micrograph of a René 80 alloy substrate treated at 1150 C for 24 hours in a (2:1) YCl_/CrCl_-activated pack containing 25 wt.% Cr-7.5Ai masteralloy. (a) Powder contacting and (b) "above pack" arrangement...... 148 45. Cross-sectional electron micrograph of a René 80 alloy treated at 1150 C for 24 hours in a 2 wt.% NH.Cl-activated pack containing 2 wt.% ZrOg and 25 wt.% Cr-lOAl masteralloy, and the corresponding x-ray maps of zirconium and oxygen. Powder contacting arrangement...... 149 46. Cross-sectional optical micrograph and composition profile (EPMA) of a IN 713LC alloy treated at 1150 C for 24 hours in a (2:1) YCl„/CrCl_-activated pack containing 25 wt.% Cr-7.5A1 masteralloy. "Above pack" arrangement...... 153 47. SIMS map of nickel, aluminum, chromium, and yttrium of a René 80 alloy substrate treated at 1150 C for 24 hours in 2 wt.% NH^Cl- activated pack containing 2 wt.% Y„o and 25 wt.% Cr-7.5A1 masteralloy. "Above pack" arrangement...... 161 48. Cross-sectional optical micrograph of a René 80 alloy substrate treated at 1150 C for 24 hours in a 2 wt.% NH.Cl-activated pack containing 2 wt.% Y^O^ and 25 wt.% Cr-lOAl masteralloy. "Above pack" arrangement...... 162 49. Variation in microhardness as a function of distance from the external surface of a IN 713LC alloy substrate treated as in (a) Fig. 48 and (b) Fig. 46...... 165

X V I 1 50. Cross-sectional optical micrograph of a Mar-M247 alloy substrate treated at 1150 C for 24 hours in a 4 wt.% NH.Cl-activated pack containing (a) 4 wt.% SiO_ and 25 wt.% Cr-25A1 masteralloy or (b) 20 wt.% Cr-20A1 and 5 wt.% Cr-40Si masteralloy. "Above pack" arrangement...... 169 51. Variation in microhardness as a function of distance from the external surface of a Mar- M247 alloy substrate treated at 1150 C for 24 hours in a 4 wt.% NH.Cl-activated pack containing 20 wt.% Cr-20A1 and 5 wt.% of a Cr-Si binary masteralloy ...... 174 52. Kinetic results of a René 80 alloy substrate treated at 1150 C in a 2 wt.% YC1_- or ZrCl.- activated pack containing 25 wt.% Cr-7.5A1 masteralloy: (a) weight and (b) coating thickness changes. Powder contacting arrangement...... 176 53. Kinetic results of a René 80 alloy substrate treated at 1150 C in a 2 wt.% ZrCl.-activated pack containing 25 wt.% Cr-7.5A1 masteralloy: (a) composition changes and (b) magnification of (a). Powder contacting arrangement...... 179 54. Schematic illustrations of the formation mechanism of the modified aluminide coating: (a) relatively short times and (b) longer times...... 180 55. Kinetic results of nickel-base alloy substrates treated at 1150 C in a 2 wt.% NH.Cl-activated pack containing 2 wt.% Y_0_ and 25 wt.% Cr-7.5A1 masteralloy: (a) weight and (b) coating thickness changes. "Above pack" arrangement...... 183 56. Arrhenius plots of the coating growth rate measured by weight and thickness changes and comparison to aluminizing rates of Ni-15Cr at.% alloys...... 184

X V I 11 57. Kinetic results of aRené N4 alloy substrate treated at 1150 C in a 2 wt.% NH^Cl-activated pack containing 2 wt.% and 25 wt.% Cr-7.5A1 masteralloy: (af composition changes and (b) magnification of (a). "Above pack" arrangement...... 188 58. Kinetic results of aRené 80 alloy substrate treated at 1150 C for 24 hours with a suitable Cr-Al masteralloy and cyclically oxidized at 1100°C in static air for up to 200, one-hour cycles: (a) ZrO_ source and (b) RE-base activator. Powder contacting arrangement . . . 190 59. (a) Specific weight changes for 1200°C cyclic oxidation in static air of coated René 80 alloy substrates and (b) the relationship with the volume fraction of entrapped Al^O^. . 19] 60. Cross-sectional optical micrograph of a René 80 alloy substrate treated at 1150 C for 24 hours in a 2 wt.% YC1_-activated pack containing 25 wt.% Cr-7.5A1 masteralloy and cyclically oxidized for 200, one-hour cycles at 1100 C in static air. Powder contacting arrangement...... 194 61. Cross-sectional optical micrograph of a René 80 alloy substrate treated at 1150 C for 24 hours in a 2 wt.% (a) NH.Cl- or (b) YCl - activated pack containing 25 wt.% Cr-7.5A1 masteralloy and cyclically oxidized for 200, one-hour cycles at 1200 C in static air. Powder contacting arrangement...... 197 62. Kinetic results of (a) René 80 and (b) IN 713LC alloy substrates treated at 1150 C for 4 hours in a pack containing 15 wt.% masteralloy and cyclically oxidized at 1100 C in static air for up to 500, one-hour cycles. "Above pack" arrangement ...... 199 63. Kinetic results of (a) René 80 and (b) IN 713LC alloy substrates treated at 1150 C for 24 hours in a pack containing 15 wt.% Cr-7.5A1 masteralloy and cyclically oxidized at 1100 C in static air for up to 500, one-hour cycles. "Above pack" arrangement ...... 200

XIX 64. (a) Cross-sectional backscattering electron micrograph and (b) composition profile of a IN 713LC alloy substrate treated at 1150 C for 4 hours in a 2 wt.% NH.Cl-activated pack containing 2 wt.% and 15 wt.% Cr-7.5A1 masteralloy and cyclically oxidized for 500, one-hour cycles at 1100 C in static air. "Above pack" arrangement ...... 203 65. (a) Cross-sectional backscattering electron micrograph and (b) composition profile of a René 80 alloy substrate treated at 1150 C for 24 hours in a 2 wt.% NH.Cl-activated pack containing 2 wt.% and 15 wt.% Cr-7.5A1 masteralloy and cyclically oxidized for 500, one-hour cycles at 1100 C in static air. "Above pack" arrangement ...... 2 04 66. (a) Cross-sectional backscattering electron micrograph and (b) composition profile of a IN 713LC alloy substrate treated at 1150 C for 24 hours in a 4 wt.% ZrCl.-activated pack containing 15 wt. % Cr-7.5A1 masteralloy and cyclically oxidized for 500, one-hour cycles at 1100 C in static air. "Above pack" arrangement...... 205 67. Schematic illustration of the degradation mechanism of the modified aluminide diffusion c o a t i n g s ...... 209 68. Kinetic results of (a) Mar-M247 and (b) IN 713LC alloy substrates treated at 1150 C for 24 hours in a pack containing 20 wt.% Cr-lQAl masteralloy and isothermally oxidized at 1100 C in air at a flow rate of 0.1 1/min (STP). "Above pack" arrangement...... 212 69. Surface electron micrograph of a Mar-M247 alloy substrate treated at 1150 C for 24 hours in a 2 wt.% NH.Cl-activated pack containing 20 wt.% Cr-lOAl masteralloy and oxidized at 1100 C in air for 44 hours: (a) surface oxide and (b) magnification of (a) ...... 214

XX 70. Surface electron micrograph of a IN 713LC alloy substrate treated at 1150 C for 24 hours in a 2 wt.% NH.Cl-activated pack containing 2 wt.% Y 0_ and 20 wt.% Cr-lOAl masteralloy and oxidized at 1100 C in air for 44 hours; (a) surface oxides and (b) magnification of (a). . 215 71. Hot corrosion kinetics and average surface compositions of Mar-M247 alloy substrates treated at 1150 0 for 24 hours and isothermally corroded at 900 C in a 0.1% S0_/0_ gas mixture with 5.0 mg/cm Na^SO.. "Above pack" arrangement...... 218 72. Hot corrosion kinetics and average surface compositions of René 80 alloy substrates treated at 1150 C for 24 hours and isothermally corroded at 900 C in a 0.1% 30^/0^ gas mixture with 5.0 mg/cm Na.SO^. "Above pack" arrangement...... 219 73. Hot corrosion kinetics of René 80 alloy substrates treated at 1150 C for 24 hours in a pack containing 25 wt.% Cr-7.5A1 masteralloy and isothermally corroded at 900 C in a 0.1% 30/0 gas mixture with 5.0 mg/cm Na_SO.. Powder contacting arrangement...... 220 74. (a),(c) Cross-sectional backscattering electron micrograph and (b),(d) corresponding oxygen x-ray map of a René 80 alloy substrate treated at 1150 C for 24 hours in a 2 wt.% (a) YC1-- or (c) ZrCl.-activated pack containing 25 wt.% Cr-7.5A1 masteralloy and isothermally corroded at 900°C for 144 hours.in a 0.1% 30„/0„ gas mixture with 5.0 mg/cm Na_30.. Powder contacting arrangement ...... 224 75. Cross-sectional backscattering electron micrograph and corresponding x-ray maps of nickel, aluminum, chromium, oxygen, sulfur, cobalt, tungsten, and molybdenum for René 80 alloy substrate treated at 1150 C for 24 hours in a 2 wt.% (2:1) YCl./CrCl.-activated pack containing 25 wt.% Cr-7.5A1 masteralloy and isothermally corroded at 900 C for 672 hours in a 0.1% 30./O. gas mixture with 5.0 mg/cm Nag30^. Powder contacting arrangement...... 225

XXI 76. Cross-sectional backscattering electron micrograph and corresponding x-ray maps of nickel, aluminum, chromium, oxygen, sulfur, cobalt, tungsten, and molybdenum for René 80 alloy substrate treated at 1150 C for 24 hours in a 2 wt.% ZrCl^-activated pack containing 25 wt.% Cr-7.5A1 masteralloy and isothermally corroded at 900 C for 672 hours in a 0.1% S0_/0„ gas mixture with 5.0 mg/cm Na_SO.. Powder contacting arrangement ...... 226

xxii ABSTRACT

A single-step, chloride-activated pack cementation process for producing a chromium/reactive element (Y,Zr)- modified and a reactive element (Y,Zr,Hf,Si)-doped aluminide diffusion coating into commercial nickel-base superalloy substrates (e.g., IN 713LC, Mar-M247, René 80, René BOH, and René N4) has been developed. The coatings consisted of a B-NiAl layer with a substantial enrichment of chromium and a reactive element (RE). These coatings were produced from a pack mixture containing either an NH^Cl activator plus a RE oxide source (e.g., YgO^, ZrOg, or SiOg) or else a RE-base activator salt (e.g., YCl^, ZrCl^, or HfCl^) with a suitable Cr-Al binary masteralloy at 1150°C for 24 hours. Two pack/substrate arrangements were used to produce the new, modified aluminide coatings: a traditional, powder contacting (PC) and an "above pack" (AP) arrangement. Unfortunately, outward-grown aluminide coatings, produced from the powder contacting arrangement, were embedded with AlgO^ and RE oxide pack particles. These entrapped pack powders reduced the resistance to thermal fatigue/cyclic oxidation attack. Physical separation of substrates from the powder mixture (e.g.,

xxiii AP) eliminated pack powder entrapment and reduced the growth rate of the coating. Pack chemistries were identified which improved resistance to cyclic oxidation, producing adherent and protective scales of AlgO^ and NiAlgO^ spinel for 500, one-hour cycles in static air at 1100°C. ITSOL equilibrium calculations predicted that the RE oxide sources were converted by the AlCl^ species to their volatile chloride species (e.g., RECl^) in the high chlorine activity pack mixture producing a condensed aluminum oxide product. Mass spectrometer measurements substantiated these predictions: identifying the RECl^ species, measuring a reduction in the intensity of the AlCl^ species, and detecting (XRD) a condensed aluminum oxide product (X-AlgOg). RE-doped aluminide diffusion coatings formed a ridged, slow-growing AlgO^ scale during isothermal oxidation in air at 1100°C for 44 hours. The parabolic rate constants for the RE-doped aluminide coatings were less than RE-free aluminide coatings. The PC, chromium/RE-modified aluminide coatings provided sufficiently better resistance to hot corrosion attack than commercial aluminide coatings. Coating lifetimes were strongly dependent on the chromium surface composition needed to form a transient aluminum- and

XXV chromium-rich oxide which better resists dissolution by the molten salt. Although physical isolation eliminated powder entrapment, the increased diffusion distance reduced the chromium flux and the resulting chromium surface composition, thereby reducing the resistance to fused salt attack substantially.

XXV CHAPTER I INTRODUCTION

The invention of the gas turbine sparked the development of new alloy compositions and processes for heat-resistant alloys. The first alloys used in gas turbines were iron-base, alloyed with chromium (> 25 wt.%) for solid solution strengthening and oxidation resistance.[1] But, these alloys were ferritic with a body-centered cubic (bcc) structure and lacked the necessary high temperature mechanical durability, such as creep rupture and tensile strengths. Therefore, iron-base alloys were slowly phased out and replaced with nickel-base alloys. Face-centered-cubic (fee) structures, such as nickel-base alloys, have high- temperature mechanical properties which are superior to their bcc counterparts. The closed-packed fee structure contains a higher elastic modulus, 12 slip systems with a greater tendency to dissociate into partial dislocations to accommodate fabricability, and lower diffusivities and higher solubilities for alloying elements than bcc structures.[1,2] Thus, fee alloys allow greater strengthening by solutes or precipitation and are preferred over bcc alloys such as ferritic iron-base alloys for high temperature applications. Further improvements were made to the nickel-base matrix. Aluminum, titanium, and niobium (Cb or Nb) were added to promote the formation of coherent y'-Ni^(Al,Ti,Nb) precipitates. With the advent of vacuum melting and casting, refractory metals (e.g., Mo,W,Nb,Ta, and Hf) were also added to improve strength and creep rupture lifetimes.[3] Refractory additions increased the tensile and flow stress of nickel-base alloys by forming both inter- and intragranular M^C, and MC carbide precipitates. Most conventionally cast polycrystalline materials exhibited intergranular failures. Therefore, small amounts of zirconium, hafnium, and boron were added to tie up free carbon and strengthen grain boundaries, inhibiting grain growth and grain boundary sliding during high temperature exposure.[1] As the demand for improved creep resistance increased, the amount of aluminum,titanium and refractory metal additions to these nickel-base superalloys was increased at the expense of chromium. These alterations caused a reduction in environmental resistance, in creep rupture lifetimes, and in ductility resulting from the formation of deleterious phases such as a, , and Laves phases. [4] In the 1970's, the development of directional solidification processes improved tensile strength, ductility, and creep rupture properties by forming oriented columnar-grained and/or single-crystal castings.[5,6] The resulting turbine blade or vane was then comprised of a nickel-base matrix with a number of alloying additions whose sole design criterion was high yield and ultimate tensile strengths, creep resistance, and fracture toughness at temperature. Unfortunately, gas turbines operate in atmospheres

containing various corrosive oxidants, such as O^/SO^, Cl^/HCl, and molten salts, which selectively deplete the substrate of its protective elements and degrade the mechanical properties of the component. The most common form of environmental degradation is oxidation, e.g., a high temperature gas/metal reaction between a metallic component and a gaseous oxidant (e.g., Og, Sg, Clg, etc.) producing a condensed or volatile metal oxide product: Me + % Xgfg) = MeX(s,l,v) (1.1) where Me is the metal and X is the oxidant species. Reaction product scales can provide adequate protection for the substrate if they are uniform and adherent; but the depletion of the protective element can degrade the mechanical properties, leading to a chemically-assisted premature failure. For this reason, most heat-resistant alloys are designed or coated to produce a surface composition which upon exposure develops a protective external product scale, usually an oxide. The oxide scale acts as a diffusion barrier to the oxidant, eliminating and/or reducing the rate of reaction or scale growth by the metallic component. When exposed to high-temperature oxidizing environments, heat-resistant alloys initially form small oxide nuclei of all its reactive elements. In time, however, a steady-state protective oxide scale overgrows the initial transient oxide to reduce the rate of further oxidation. Not all thermally grown oxides are equally protective. Various external oxides suffer cracking from stresses induced during their formation and growth (e.g., NbgOg) allowing molecular oxygen to the metal surface. Other oxide scales exhibit sufficiently large growth rates (e.g., NiO) which fail to provide optimum oxidation resistance. A number of exhaustive investigations to determine the oxidation behavior of reactive elements and their alloys have since been undertaken.[7,8] Today, the most protective thermally grown oxide scales for heat- resistant alloys, such as nickel-base superalloys, are alumina, chromia, or silica. These oxide scales are not only thermodynamically very stable but they constitute excellent diffusion barriers to molecular oxygen and possess the slowest growth kinetics for thermally grown oxides. Figure 1 is an Arrhenius plot of the parabolic rate constants, k^, for alumina, chromia, and silica scales thermally grown on various heat-resistant metals and alloys.[9] Their growth rates are orders of magnitude lower than those for other oxide scales. Thus, alloys or surface coatings containing the elements Cr, Al, and Si which form dense, adherent scales during exposure are most desirable for industrial use. The choice of the most effective protective oxide scale is dependent on operating conditions such as temperature (Fig. 1), the oxygen chemical potential, and the bulk gas flow rate. Below 900°C, any one of the three oxides are capable of providing adequate oxidation resistance. Above 900°C in high oxygen chemical potentials and flow rates, chromia suffers excessive oxidation/evaporation according to Eg. (1.2) and is no longer protective.[10,11]

Cr 2 Û 2 (s) + 3/2 0 2 (g)=2 CrOg(v) (1.2) Therefore, in oxidizing high-temperature applications, heat-resistant materials are designed or coated to form either alumina or silica scales upon exposure to an oxygen-bearing environment. Nickel-base superalloys are widely used as turbine blades, vanes, and rotors in utility boilers for power generation and in gas turbine engines for use in civil and military aircraft. These alloys are alloyed with both aluminum and chromium and uncoated, generally, form an external chromia scale during service. Because current alloy compositions are unable to provide adequate environmental resistance, modification of the surface composition is necessary to produce a more protective oxide scale (e.g., AlgO^ or SiOg) in service. Two major processes are used commercially to produce surface modifications: diffusion coatings and overlay coatings. The coatings have differing structures, compositions, costs, etc. Diffusion coatings are deposited at high temperatures to support simultaneously solid/vapor reaction and reaction/interdiffusion of the coating with the substrate. Two principal types of diffusion coatings are chemical vapor deposition (CVD) and halide-activated pack cementation (HAPC), a form of self-generated CVD. Both types are very versatile, capable of coating the complex shapes of turbine hardware although small internal cooling passages represent a problem. On the other hand, overlay coatings rely on the physical transfer of elements from a source to the substrate, i.e., physical vapor deposition (PVD). The four main Kp Data for AlgO^, CrgO^, and SiOj

Temperature (%) 1600 1400 1200 1000 800 -10

11

-1 2 A1,0

O' -13 o

-14

-15 SiO.

-16 5 6 7 8 9 10 i o V t(k )

Figure 1: Arrhenius plot of the parabolic rate constants of alumina, chromia, and silica scales thermally grown on various heat resistant alloys or coatings.[9] 8

types of overlay deposition methods are: (a) electron

beam physical vapor deposition (EBPVD), (b) plasma spraying, (c) ion implantation, and (d) sputtering. The use of PVD processes is limited for two reasons: (a) they are "line of sight" processes so that complex shapes and tiny passages are difficult to coat effectively and (b) they are relatively expensive, requiring large vacuum equipment and perhaps post-coating heat treatments to interdiffuse the coating with the substrate. Halide-activated pack cementation (HAPC) is a high- temperature, in-situ chemical vapor deposition batch process, which is relatively inexpensive and practiced commercially for the past thirty years to apply aluminide diffusion coatings to blades and vanes for use in gas turbines. In the HAPC process, the pack powder mixture is comprised of three main ingredients or components: (a) a masteralloy powder of the element(s) to be deposited (e.g., Al, Cr, or Si), (b) a halide salt activator (e.g., NaCl, NaF, NH^Cl), and (c) an inert filler powder (e.g., AlgOg). This powder mixture is mixed thoroughly and placed with the parts to be coated within a heat resistant retort. The retort is heated in a controlled atmosphere, usually Ar or Hg/Ar, to an elevated temperature (e.g., 1000-1150°C), at which, the masteralloy reacts with the halide salt activator to form volatile metal halide species of significant partial pressures according to the following reaction: Me(alloy) + AX^(l,v)=MeXj^(v) + A(l,v) (1.3) where Me is Cr, Al, or Si; A is Na, NH^, etc.; and X is Cl, F, or Br. In fact, several different vapor species may be formed for each element (e.g., Al, AIX, AlXg, and AIX^; Cr, CrXg, CrX^, and CrX^; or Si, SiXg, SiX^, and SiX^). A partial pressure gradient for each vapor species, which supports vapor transport to the substrate surface, results from the higher thermodynamic activity in the metallic component in the powdered masteralloy compared to its lower activity at the substrate surface. At the surface, deposition of the desired coating element(s) occurs according to one of several possible deposition reactions: (a) Disproportionation, e.g., 3 MeXg(v) = 2 MeXg(v) + Me(substr.) (1.4) (b) Displacement reaction with the substrate, e.g.,

MeXg(v) + M(substr) = MXg(v) + M e (substr.) (1.5) (c) Reaction with the environment, e.g., MeXg(v) + Hg(g) = 2 HCl(g) + Me(substr.) (1.6a) MeXgCv) + 2 Na(v) = 2 NaX(l,v) + Me (substr. ) (1.6b) Finally, the coating elements react and interdiffuse with the metallic substrate, producing some specific surface 10 composition and microstructure. A summary of the coating process is presented in Fig. 2. Aircraft and industrial turbines burn a number of low quality fuels, which may be contaminated with various impurities including sulfur, vanadium, lead, potassium, chlorine, and sodium. These impurities in the fuel and others in the ingested air, especially NaCl and NagSO^ from sea water, can react with combustion by-products to form fused salts (e.g., wagSO^, NaVO^, NaCl) which might condense as a thin film on turbine components, leading to an accelerated oxidation attack termed "hot corrosion." Since the 1960's, aluminide (B-NiAl-base) diffusion coatings, produced via the pack cementation method, have been applied to turbine hardware for oxidation resistance. During service, these aluminide coatings produce a steady-state alumina scale. Unfortunately, aluminide diffusion coatings do notprovide adequate protection from hot corrosion attack and, therefore, require a modification in the surface composition. McCarron et al.[147] determined that chromium additions in bulk B-NiAl compounds, the principal phase in aluminide coatings, improved burner rig hot corrosion performance considerably. In addition, Swazda et al.[12] determined that enriching the surface of an aluminide coating with about 1-2 wt.% silicon also provided an 11

Uapor Phase Coating Unaffected AlCl.AlClg Alloy CrCl^---- ->01 Pack Powders ->Cr (Hi.Cr)Al RECl.REClg Compound RE Alloy Surface Uapor Phase Transport and Reaction to Codeposit a Desired Composition

Figure 2: A schematic illustration of the chloride- activated pack cementation process. 12

impressive improvement in hot corrosion resistance. Aluminide coatings modified with chromium[13,14] and silicon[15] have been developed. Such (HAPC) coatings are produced in two sequential steps, a traditional aluminizing process followed by a chromizing or siliconizing step. It would be quite advantageous to deposit simultaneously these two elements in a single-step; however, to date, this has not been accomplished commercially or deemed feasible.[16,17] The protective oxide scales produced on high temperature alloys are often exposed to stresses induced by thermal cycles[18] and by oxide formation and growth[19], which can cause loss of adherence or spelling, especially upon thermal cycling. It has been widely established that a small concentration (<1 wt.%) of a highly reactive element (e.g., Y, Zr, Hf, Th, Ce, La, Sc, or Si) can improve the adherence of chromia and alumina protective scales on nickel-, cobalt-, and iron-base alloys.[19-22] More specifically, Barrett[23], Jedlinski and Mrowec[24,25], and Santoro et al.[26,27] have determined that small additions of Zr, Y, and Si, respectively, also improved the adherence of alumina scales thermally grown on bulk 5-NiAl compounds. Therefore, to produce highly adherent alumina scales during exposure, the coating process must be altered to 13 deposit a small amount of a reactive element into the surface of the aluminide coating. The objectives of this thesis are as follows: (a) Develop the necessary parameters, such as pack compositions, time, and temperature, to deposit simultaneously aluminum, chromium, and a reactive element (e.g., Y, Zr, Hf, or Si) on commercial nickel-base superalloys by a single-step, chloride-activated pack cementation process. (b) Using thermodynamic calculations for the multicomponent, multi-phase pack system determine vapor pressures of species in the pack and the mechanism for coating formation and growth. (c) Evaluate the performance of HAPC coatings in several high temperature tests including isothermal oxidation and hot corrosion (salt spray and burner rig) and cyclic oxidation. (d) Determine the mechanisms of environmental degradation for the coatings tested. CHAPTER II LITERATORE SURVEY

2.1 HALIDE-ACTIVATED PACK CEMENTATION DIFFUSION COATINGS 2.1.1 Aluminizing Aluminizing of heat-resistant alloys was invented by Van Aller in 1911.[28] The process consisted of a powder mixture of aluminum, ammonium chloride, and graphite which was then heated to 450°C followed by a post-deposition interdiffusion treatment at 700°C. The resulting coatings consisted of either an aluminum-rich surface composition or intermetallic compound, depending on the substrate, activator, and aluminum source. Nickel-Aluminum Svstem; The nickel-aluminum system is quite complex (Fig. 3).[29] The addition of aluminum to the nickel results in the formation of higher melting intermetallic compounds (e.g., NiAl^, NigAl^, NiAl, and NigAl) . Of these, the most important are the y'-Ni^Al phase, the principal strengthening agent in commercial nickel-base superalloys, and the 8-NiAl phase, the compound produced from the aluminizing process. The NigAl phase is an ordered intermetallic compound with the CUgAu structure. The Ni^Al phase contains a unique flow

14 15

Weight Percent Nickel 20 30 40 60 70 60 0 0 100 1800

1400- AINi (Ni)

^ 1000.

0 0 0 .

600'

(A!) AljNi. AINi; 400 30 40 so 60 70 60 fiO too AI Atomic Percent Nickel Ni

Figure 3: Binary nickel-aluminum phase diagram. [29] 16 stress phenomenon, whereby as the temperature increases the flow stress also increases.[30] This provides a remarkable improvement in the mechanical properties of nickel-base superalloys strengthened with Ni^Al precipitates. The NiAl phase is an ordered intermetallic compound with the CsCl crystal structure. The NiAl phase exists over a wide range of stoichiometry, from 3 0 to 58 at.% aluminum, and possesses three desirable qualities for a high temperature material: (a) promising high temperature strength[31], (b) low chemical diffusivities[32,33], and (c) excellent oxidation resistance.[34] The interdiffusion coefficients of the nickel- aluminum system, including the Ni^Al and NiAl phases, was determined by Janssen and Rieck[32] and Shankar and Seigle[33]. A summary of these investigations and others on the nickel-aluminum system are presented in Table 1. Janssen and Rieck used Ni/NiAl and Al/NiAl diffusion couples along with the Boltzmann-Mantano-Heumann relationship to calculate the interdiffusion coefficient (D) from the aluminum concentration profiles following annealing treatments. Unfortunately, only average interdiffusion coefficients were established without providing the concentration dependence of the B-NiAl phase. Shankar and Seigle, on the other hand, used a 17 pack aluminizing process to grow the NiAl phase on a pure nickel substrate with a constant aluminum concentration at the surface of the coating. By varying the aluminum activity of the pack powder mixture, an aluminum surface concentration scanning the entire range of NiAl stoichiometry was produced. By analysis of the resulting aluminum concentration profile using Wagner's relationship[35] for intermetallic compounds, the interdiffusion coefficient of the NiAl phase was calculated with respect to composition. Figure 4 is a plot of the interdiffusion coefficients (D, 0^.^, of the B-NiAl phase with respect to the aluminum concentration. It should be noted that the transition in the dominant diffusing element occurs near (48 at.% Al), and not at the exact stoichiometric composition. The microstructure of aluminide coatings have been categorized into two types based on the relative "activity" of the aluminum source and the mechanism of coating growth.[36,37] The two classifications are the "low" and "high" activity processes, which are illustrated in Figs. 5-8. For example, high activity aluminide coatings are produced from a powder pack mixture containing pure aluminum powder at 700-900°C. Because of the large aluminum chemical potential gradient, inward diffusion of aluminum dominates and a 18

Table 1: Summary of interdiffusion studies on the nickel- aluminum system.

Aluminum Interdiffusion Activation Content Phase Coefficient Energy Refer. (at.%) 1100°C (cm^/s) (kcal/mol)

14 -Ni 3.3*10“?'° 56+2 32 II II 2.6*10-“ 65+2 33 25 '-Ni-Al 3.6*10-“ 64+11 32 II II 3 3.4*101“ 62±2 33 38 6-NiAl 2.7*10 50+1 32 II II II 55±2 33 45 If - 59+6 II 51 II 43+8 II 19

A S

CM

N i

111 o

I0 “ »

,-12 36 40 44 48 52 56 COMPOSITION (ol.VoAl)

Figure 4: Variation of the interdiffusion coefficients with composition in the 6-NiAl phase at 1100 C.[33] 20

microstructure consisting of a NigAl^ matrix, a large dispersion of chromium and molybdenum-rich precipitates, and a number of other alloying elements in solid solution (e.g., Ti, Co, Cr, Mo, Nb, etc.) forms. Following a subsequent homogenization treatment at 1000-1100°C, the NigAlg matrix is transformed into a hyperstoichiometric (aluminum-rich) NiAl phase (Figs. 5 and 7). The resulting transformation is controlled by the dominant inward diffusion of aluminum from the aluminum-rich NigAl^ phase to the nickel-base substrate. The second type of aluminide coating, the "low" activity process, is produced from a pack powder mixture containing an aluminum alloy powder source of a reduced activity (e.g., Ni-Al, Fe-Al, Cr-Al). The reduced aluminum activity is vitally important. The vapor pressures of the aluminum halide species generated from a low activity pack are significantly lower and require higher temperatures (1000-1150°C) to produce aluminide coatings comparable to those from a high activity pack. The "low" activity process is still capable of producing the necessary chemical potential gradient for vapor transport and coating growth, but the resultant coating is a hypostoichiometric (nickel-rich) NiAl phase. (Figs. 6 and 8) From Fig. 4, the dominant diffusing element in the nickel-rich NiAl phase is nickel. Therefore, the 21

Ni,Al3 Initial limit NIAI

After complete Pure nickel During diffusion annealing transformation

Cr ricfr precipitates containing other elements Ni,Al3 Ni,AI NiAl NiAl Û o O C 3 o a ^ TiC Initial limit

during after complete diffusion annealing transformation

Figure 5: Schematic illustration of the "high" activity process; formation of a NiAl coating from a Ni.Alg layer by diffusion annealing in (a) pure nickel or (B) nickel- base superalloy.[37] 22

/ Mj f k C f î V NiAI

/ I \ Iniiidl surface c, y Porcs Pure nickel

Before aluminizing After aluminizing

Pack pjriiclc tnclusforit V i n NiAI Extern*; IfCh tn Ail

Initi*; surface

Intern*! Alloy — Ni fone NiAI 4 precipitates IO tflut«ng (*

ro n « |

Precipitates formed from those elements of the substrate which carmot be completely dissolved in the NiAI compoursd

Figure 6: Schematic illustration of the "low" activity process to produce a NiAI coating by outward diffusion on (a) pure nickel or (b) nickel-base superalloy.[37] 23

ZONE 1 P(NiAI) matrix ♦ bcc a Cr, Mo + substrate phases

ZONE 2 single phase p(NiAI), Cr, Mo. Ti, Co in solution

ZONE 3 P(NIAI) matrlx.TIC, M C , and otCr.Mo.Co)

3 ; .•4" ■ o phase in coating affected substrate

Figure 7: Microstructure and phase identities of a "high" activity aluminide coating on Udimet 700 followed by heat treatment for 4 hours at 1080 C; lOOOx.[36] 24

ZONEI single phase p(NIAl), Cr,Mo,Ti,Co in solutiort .

ZONE 2 p(NiAl) m atrix eTiC, M C , and o(Cr,Mo,Co) I 23 6

I coating affected i substrate

Figure 8: Microstructure and phase identities of a "low" activity aluminide coating on Udimet 700; 1000x.[36] 25 growth of an outward-grown NiAI coating compound is controlled by the diffusion of nickel through the coating from the nickel-base substrate. In addition, other alloying elements also diffuse outward and remain in solid solution in the NiAI matrix (Fig. 8). Two observations should be noted about the "low" activity process. Because nickel has a higher mobility in the nickel-rich, NiAI coating compound, Kirkendall voids are formed beneath the coating/substrate interface from the countercurrent flux of vacancies created to balance the unequal fluxes of nickel outward and aluminum inward (Fig. 6a). For pure nickel substrates, Kirkendall voids are frequently observed, whereas no voids were observed in coated nickel-base superalloy substrates (Figs. 6b and 8) because the annihilation of voids occurs in the multiphased interdiffusion zone. Depending on the superalloy composition, the interdiffusion zone is comprised of the long chain of blocky MC-type (M=Ti, Nb, Mo, W, Hf, Ta) carbides near the original substrate surface and a series of rod-shaped sigma phase particles rich in Cr, Mo, and Co. The interdiffusion zone resulted from the outward diffusion of nickel and from both the diffusion of carbon from the bulk alloy and the degeneration of carbides below the nickel-base substrate surface. Carbon reacts with the refractory 26 metal alloying elements in the zone depleted of nickel precipitating first MC-type carbides, and second due to the lack of carbon, sigma phase. The vacancies which form during coating growth can be annihilated at the interfaces of these second phases (Figs. 6 and 8). As a result of dominant NiAI growth at the external surface, pack powders (i.e., masteralloy and inert pack particles) are entrapped by the coating. Since "low" activity aluminide coatings are grown outward, various pack powders become embedded in the coating during the aluminizing process. Pack entrapment is undesirable as it can seriously degrade the integrity of the coating. However, by physically isolating the substrates from the pack powder mixture, a process procedure called "above pack", particle pack entrapment can be eliminated and a cleaner coating can be produced (Fig. 8). 2.1.2 Chromizing The chromizing process was invented by Kelly in 1923. [38] The process consisted of a powder mixture of pure chromium, ammonium chloride, and alumina which was then heated to 950°C for 16 hours. Chromized coatings were first applied to plain carbon and low alloy steels, where they proved to be highly resistant to aqueous and high temperature corrosion. In later years, it was 27

also determined that chromium in the bulk or chromized coatings applied to nickel-base alloys provided effective resistance to hot corrosion. The microstructure of chromized coatings on nickel-base alloys is much simpler than that for an aluminide coating. The nickel-chromium binary phase diagram (Fig. 9) shows a high solubility for chromium (_50 at.%) in the nickel matrix and no complex intermetallic phases. Therefore, "high" activity chromized coatings are comprised of a chromium enriched nickel-base substrate with an external a-chromium layer. Complications arise when chromized coatings are applied to carbon-containing alloys, such as carbon steels and nickel-base and cobalt-base superalloys.[40-42] Because of the high affinity of chromium for carbon and the relatively high diffusivity of carbon in these alloys, an outward-grown carbide is formed at the external surface by the chromizing process. This external carbide phase acts as a diffusion barrier to rapid chromium enrichment. If the process is carried out above 1050°C, dissolution of the carbide at the interface supplies chromium to the substrate at a reduced rate. One remedy to improve deposition rates is to codeposit a second element which either has less affinity for carbon (e.g., A1 or Si), transforming the surface (e.g., a-Fe) 28

Weight Percent Chromium 20 30 40 60 70 60 90 too 1900

O « 3 (Cr) « B 1 1 0 0 - (Ni)

0 0 0 -

TOO

2 0 40 SO «0 70 0 0 100 Ni Atomic Percent Chromium Cr

Figure 9: Binary nickel-chromium phase diagram.[39] 29

into a phase with less carbon solubility and pushing carbon into the bulk substrate. Conversely, another solution is to codeposit a second element which has greater affinity for carbon (e.g., V) and helps to tie up residual carbon which otherwise diffuses to the surface. 2.1.3 Siliconizing Silicide coatings have been applied to nickel-base alloys by a number of methods including chemical vapor deposition (CVD)[43,44], fusion slurry techniques[45], and pack cementation[46]. Siliconizing of heat-resistant alloys was first produced via halide-activated pack cementation by Fitzer[46] in 1955. Silicide coatings were found to have enhanced resistance to both low temperature and high temperature hot corrosion resulting from the formation of an external silica scale. According to the nickel-silicon binary phase diagram (Fig. 10), the addition of silicon to nickel-base substrates results in the formation of several, low melting eutectics and complex intermetallic compounds. The typical microstructure of silicide coatings on nickel-base substrates are comprised of a multi-phase, duplex layer containing an outer, Nisi phase and an inner, Ni^Si phase. These silicide coatings are very brittle and susceptible to stresses induced by thermal cycles.[46,47] Therefore, the use of silicide coatings 30

Wolglit Percent Silicon

isoo

HOO-

1300- o ( S i)

(Ni) oer

eoo

TOO 40 SO « 0 TO Ni Atomic Percent Silicon Si

Figure 10; Binary nickel-silicon phase diagram.[29] 31 on nickel-base alloys is not popular. To utilize the beneficial effects of silicon, aluminide and chromide coatings have been modified to incorporate silicon into the surface of the coatings. [48-51] 2.1.4 Deposition of Reactive Elements (RE) It is well accepted that small fractions of a percent of reactive elements (e.g., Y, Zr, Hf, Ce, La, Th) improve the adherence of protective oxide scales on heat-resistant alloys.[19-22] Overlay coatings applied to nickel-base alloys have always included a reactive element, but only a few investigators have deposited reactive elements by other methods. Jedlinski et al.[52] incorporated yttrium and cerium into the surface of NiAI coatings on nickel-base alloys by an ion implantation technique. While, Fuhui et al.[53] produced an yttrium-modified aluminide coating on IN 738 alloys, using a standard aluminizing treatment followed by a fusion slurry technique to deposit yttrium. Only three investigations by the pack cementation method have been noted. Tu et al.[54] produced an yttrium-modified aluminide coating by a two-step process, aluminizing followed by yttrizing. The modified coatings consisted of an external aluminide layet with yttrium enrichment, detected by wavelength dispersive spectroscopy, near the surface and along grain 32 boundaries. The process to deposit yttrium consisted of a pack powder mixture of pure yttrium, ammonium chloride, and yttria (YgO^) which was heated at 1050°C for four hours. The alumina filler was not inert in this situation because the high thermodynamic stability of the aluminum chlorides permitted some reaction with yttria. Conversely, Rapp and coworkers[40] replaced a small amount of the alumina filler with yttria in a pack mixture to codeposit chromium and aluminum on pure iron and low-alloy steel substrates. Energy dispersive spectroscopy analysis detected about 0.4 wt.% Y in the coating surface. No effort was made to contribute the presence of elemental yttrium to deposition from the gas phase or from the entrapment of yttria (Y^O^) in the external, outward-grown layer. LePrince et al.[55] used a "high" activity process with a powder mixture of pure source elements to codeposit aluminum and hafnium onto nickel-base fibers. The resulting coating consisted of a NiAI matrix with Ni^Hf precipitates. In all cases, the reactive element(RE)-doped, aluminide coatings exhibited improved oxide scale adherence during cyclic oxidation testing. 33

2.1.5 Codeposition According to Wagner[56], a ternary alloy or coating using the interaction of two oxidation-resistant elements (e.g., Cr-Al, Cr-Si, or Al-Si) produces more effective oxidation resistance than a simple binary alloy or coating. For example when a MCrAl (M=Ni, Fe, or Co) coating or alloy is exposed to a high-temperature oxidizing environment, initially a continuous, transient CrgOg scaleis formed on the surface. This transient oxide scale prevents the rapid oxidation of the base element (M), thereby eliminating the dissolution and inward diffusion of solute oxygen that would cause the internal oxidation of aluminum. Aluminum atoms, eventually, diffuse to the CrgOg/alloy (coating) interface and form an even slower-growing protective AlgOg scale at steady-state. In fact, for NiCrAl alloys

and coatings, transient (Al,Cr)2 0 g or even Ni(Al,Cr) 2 0 ^ scales form instead of CrgO^.ESV] A steady-state AlgO^ scale finally results from the displacement of chromium atoms in the scale by aluminum atoms in the alloy or by

the oxidation/evaporation loss of Cr in the (Al,Cr)2 0 g scale according to Eg. (2.1).

(Al,cr)2 0 3 (s) + 3/2 0 2 (g) = 2 Cr 0 3 (v) + AI 2 O 3 (s) (2 .1 ) It has been proposed that the addition of chromium to nickel-aluminum alloys increases the diffusivity of 34 aluminum in the alloy or coating, resulting in the formation of a slower-growing, more protective AlgO^ scale.[58-62] Wagner[63] derived an expression for the critical solute content necessary to sustain the growth of a scale at steady-state based on a flux balance of solute in the bulk alloy or coating to that through the protective

scale according to equation (2 .2 ):

Nb *=( (2.2) where g is the critical volume fraction of oxide needed for the transition, is the solubility of oxygen in the alloy, and Dg are the diffusivities of oxygen and solute in the alloy, respectively, and and are the molar volumes of the alloy and oxide, respectively. Therefore, for the increase in the diffusivity of the solute or the reduction in the oxidation rate, third element additions like chromium promote the formation and sustained growth of an external, steady-state AlgO^ scale at a lower critical solute (Al) content in the alloy or coating. A series of NiCrAl candidate alloys were widely investigated in isothermal and cyclic oxidation experiments. [58-62] A typical oxide map of the NiCrAl system is presented in Fig. 11. The minimum chromium and aluminum contents which produce an external and slow-growing, steady-state AlgO^ scale are noted. 35

1000°. 1100 “ Cl 20 hr 1200 “c i 1000 “ c

III

30 20 10 Ni — Winer I bricfiulNIO. Internal Cr^O^IAI^O^INiMI.Crlÿ^ II blernal Cr^O), Inlernal AIjOj III External A %

(a) Icotherm at m ap.

1100 °C. 500 h r ; 1200% . 200 h f

30 20 10 Ni - — w t n C r I M mllyNiO II MosIlyCrjOj III AIjOjUNIAIjO^I I NiO, NiCr20^, NIAIjO^. AI^O] Transilioo

(b) Cyclic m ap.

Figure 11: Experimentally determined oxide map of the NiCrAl alloy or coating system. [61] 36

Single elements are commonly deposited into metallic substrates by the pack cementation method, but a single-step process which would simultaneously codeposit multiple elements into metallic substrates would be more effective because the resulting coatings would exhibit superior environmental resistance according to the above explanation. Galmiche[13] developed a single-step, "high" activity process to simultaneously deposit chromium and aluminum from a pack powder mixture containing both a pure chromium and aluminum source with an ammonium chloride activator. The resulting coating consisted of a single-phase, NiAI layer with only limited chromium enrichment from the gas phase. Millet [64] developed a similar process except that an ammonium bromide activator salt was used. Bromide activators have a greater tendency to chromize because the volatile bromide species of aluminum and chromium are much more similar than the chlorides. In fact, although the powder mixture contains both a chromium and aluminum source, the process is still characterized as a two step process. Aluminizing occurs first because of the large thermodynamic stability of thealuminum halides. But once aluminum is depleted locally from this dilute pack mixture, chromizing of the aluminide coating occurs. In general, the prospect of true simultaneous deposition has 37 not been practiced commercially, although it has been developed experimentally. [136] 2.2 THERMODYNAMICS AND KINETICS OF PACK CEMENTATION 2.2.1 Single Element Deposition (Monodeposition) Pure Element Source; The thermodynamics and kinetics of halide-activated pack cementation have been studied rather extensively.[65-74] In an early investigation, Levine and Caves[65] thoroughly studied the effect of alternative activators and time and temperature treatments to determine the mechanisms of deposition and growth of aluminide coatings. At the processing temperature, the source element and the halide activator salt react to form volatile metal halide species of significant partial pressures (Eg. 1.3). In fact, local equilibrium is established in the pack and at the substrate surface, and a partial pressure gradient for each vapor species is created. The driving force for gaseous diffusion through the porous powder mixture is the negative gradient in partial pressure, as illustrated by Levine and Caves.[65] In a single component "high" activity pack, a negative pressure gradient is insured as long as the activity of a species in the pack is greater than the activity in the substrate. The kinetics of vapor diffusion through a pack depleted at the surface was modeled using a modified 38

Pick's first law expression (i.e., J=D^ (dc^/dx) where Cj=n^/V=P^/RT)). Therefore for a linear concentration gradient in the gas phase, the instantaneous flux, of a metallic component M, through the pack to the substrate can be calculated from pressure gradient values with the aid of Eg. (2.3): J„«*5-£(-D„x^dP^^)/ET (2.3) where the effective gaseous diffusion coefficient of

MX^ in the pack, 8 the diffusion distance or the depletion zone width, R the universal gas constant, T the absolute temperature, and the pressure gradient of species MX^. Unfortunately, these calculations only rationalize whether a particular halide vapor species will diffuse to the substrate. Levine and Caves[65] modified Eg. (2.3) further to calculate the steady-state deposition rate of component M, via gaseous diffusion. The new expression accounted for the porosity and tortuosity of the porous pack powder medium and is given as: Kg=(2pf2j^)j/*5 (2.4) where p the pack density of component M, c the pack porosity (fraction of available gas phase), the molecular weight of component M, and 8 the 39

instantaneous flux of component M. The rate of aluminum pickup at steady-state was measured for different temperatures and activators according to Eg. (2.5); W=Kg(t)^ (2.5) where W is the specific weight change (gm/cm 2 ) , t the treatment time (hr), and the steady-state parabolic rate constant for coating formation (gmf/cm^hr). Only rough agreement between the measured rates (K^) of aluminum pickup and the Al deposition rates calculated from Eg. (2.4) were determined. Levine and Caves[65] also established that the growth of aluminide coatings was indeed parabolic (n=0.5). From apparent activation energy values, solid-state diffusion is predominant in the rate-controlling step for the kinetics of the process. The kinetics of pack aluminizing of nickel and iron substrates were studied by Gupta et al.[67] and Kung and Rapp[71]. For a given aluminum activity of the source, a time invariant surface aluminum concentration was produced, implying that a flux balance across the coating/gas phase interface existed or Kg=Kg) and that the rate limiting step was series diffusion in both the gas phase and solid-state. A model to predict aluminum concentration profiles, assuming a condensed activator (i.e., NaCl, NaF) was presented at the surface, 40 was proposed and tested. Although predictions agreed well with experiments, several discrepancies with these investigations were discovered. Gupta et al.[67] derived a semi-infinite diffusion model using the diffusivities in the nickel-aluminum system to predict the aluminum surface concentration and profile with good agreement of an aluminized, pure nickel substrate. Kung and Rapp[70], on the other hand, used an Fe-Al binary masteralloy with both condensed (e.g., NaCl) and volatile (e.g., AlCl^ and FeClg) activators. For condensed activators, good agreement between experiments and predictions resulted, but poor agreement resulted for experiments performed with volatile activators. Two explanations were presented; (a) volatile activators are often vented from the pack retorts and do not allow local equilibrium to be established in the pack mixture, let alone the surface, and (b) the use of a Fick's first law expression for vapor diffusion in a pack containing a binary Fe-Al masteralloy is questionable. During the aluminizing process, aluminum is depleted from the adjacent Fe-Al masteralloy particles thereby reducing their aluminum activity and the instantaneous flux of aluminum to the 41 coating surface. [73] Therefore, a time-dependent expression would be better suited to analyze the coating process. The effect of the activator was much simpler to determine. Experimentally, chloride- and fluoride- activated packs produced the highest aluminum and silicon deposition rates [65,68,72] whereas chloride- and bromide-activated packs produced the highest chromium deposition rates.[17,74] Because of the differences in thermodynamic stability between the halide activators and the metal halide species generated, specific activators produce higher vapor pressures for specific metal components. Thus, higher fluxes and deposition rates result. For example, chloride and fluoride activators produce the highest vapor pressures for the aluminum halide species. Therefore, thick aluminide coatings are produced because of the higher deposition rates, but the surface aluminum concentration is constant and unaffected by these larger deposition rates. The sole criterion used to determine the surface composition and microstructure of the coating is the activity of the source element and the treatment temperature.[65] The dependence of the vapor pressure of the halide species indicates that gaseous diffusion is dominant in the deposition and growth of pack cementation coatings. 42

UNDEPLETED PACK DEPLETED ZONE SPECIMEN

AIX

AIX AIX HX

(a)

AIF

AIF. t> V AIF.

HF

(b)

Figure 12: (a) Activator circulation and (b) activator condensation mechanisms of aluminum deposition proposed by Levine and Caves [65] where circles=Al(l) and triangles=AlFg(c).[69] 43

Levine and Caves[65] devised two general mechanisms which can apply to the formation of aluminide, chromized, and silicide coatings: (a) the circulation mechanism and (b) the condensation mechanism (Fig. 12) . The circulation mechanism attributes deposition to the transport of lower halide species (e.g., MeX and MeXg) to the surface where a disproportionation reaction (explained earlier) occurs, depositing aluminum and producing a higher halide species (e.g., MeX^) which returns to the pack to recycle the halide. Thermochemical measurements have substantiated this mechanism.[65,69] Higher vapor pressures for the AIX^ species were measured in local equilibrium with the coating surface during aluminizing of nickel-base alloys. [65,69] The condensation mechanism attributes deposition to the transport of all metal-halide species to the surface where a condensed metal-rich phase forms to consume these elements. This mechanism has a strong dependence on activator content.[69] Experimentally, the regions where each mechanism dominates and their respective dependences on activator content was determined (Fig. 13). In addition, aluminum- and fluorine-rich droplets were detected on the surface of coatings formed in AIF^-activated packs. 44

ACTIVATOR ACTIVATOR+ AI UNDEPLETED PACKDEPLETED DEPLETED SPECIMEN ZONE ZONE

AIF

AIF. AIF.

HF o'

X>\

A 6 C

Figure 13: Mixed mechanism of aluminum transport in the presence of both an activator and activator/source depleted zones, where circles=Al(1) and triangles= AlFgCc).[69] 45

24

MIXED m e c h a n is m CONDENSATION

<•2

Wl.*/. AIF,

Figure 14: Theoretical variation in the parabolic deposition rate of aluminum, K , with activator content at 900 C for a 4 wt.% pure Al pack activated by A1F_.[69] 46

Kandasamy et al.[69] proposed a single mechanism (Fig. 14), where both circulation and condensation can contribute to the deposition of aluminum. This model identifies two distinct depletion zones, a source and a source/activator, corresponding to each of the two models outlined above. The composition of the inert atmosphere also has an effect on the coating process. Kung and Rapp [71] used computer-assisted calculations to analyze the volatile halide species generated for "high" activity packs with several activator salts and in various inert atmospheres (e.g., Ar, Hg/Ar, or Hg). Inert atmospheres of argon or argon/hydrogen are commonly used during pack cementation processes to eliminate premature oxidation of the source elements or alloys and substrate. The addition of hydrogen to the inert shroud reduces the partial pressure of chlorine by forming hydrogen chloride gas. This reduction in chlorine activity also decreases the vapor pressures of the silicon and chromium chloride species but has little effect on the vapor pressure of the more thermodynamically stable, aluminum chloride species. Overall, a Hg-containing atmosphere may influence the deposition rate of chromium or silicon from chloride-activated packs. 47

<2 41 - THEORETICAL PROFILE Û MEASURED DATA POINTS

39 38 37 35 35 34 33 SAMPLE SURFACE A| .CONCENTRATION 32 0 2 4 6 e 10 1 2 14 16 (a) d i s t a n c e FROM Sa m p l e s u r f a c e (**)

1.9

e 1.8 o 1.7

c 1.6 o 1.5

1.4

1.3

1.2 0 2 46 8 10 12 14 16 (b) DISTANCE FROM SAMPLE SURFACE (mm)

Figure 15; (a) Pack aluminum concentration profile and (b) , AIF vapor pressure profile in a Ni-40A1 at.% alloy pack after 8 hours at 1085 C.[73] 48

Reduced Activity Source; The formation of aluminide coatings from a reduced activity source possesses two major differences compared to coatings produced from a pure source. One is morphological. As explained earlier because of the reduced activity, an outward-grown NiAl coating layer forms during the aluminization of nickel and nickel-base alloys instead of the NigAl^ phase. The growth rate of the coating is also reduced, because the flux of aluminum to the surface is lowered by the generation of lower halide vapor pressures and is, therefore, compensated by an increase in the process temperature.[73] Similar trends exist on the effect of different activator types used during the coating process. Secondly, the kinetics of coating growth can no longer be described simply by a Pick's first law expression to represent the gaseous diffusion of metal halide species. Wang and Seigle[73] used a Pick's second law expression for the time-variant concentration dependence to predict the average surface aluminum concentration and profile of an aluminized, pure iron substrate with exceptional agreement. Using this same model, the aluminum concentration profile of the adjacent Pe-Al masteralloy particles resulting from the depletion of aluminum during the aluminizing process was also 49 successfully predicted (Fig. 15a). Again, the model was based on the fact that coating growth is controlled by series diffusion in both the gas phase and solid-state, and thus a flux balance exists across the substrate/gas phase interface for the case where a condensed activator is present throughout the process. 2.2.2 Codeposition As mentioned earlier, deposition of two protective elements (e.g.. Or and A1 or Cr and Si) has been achieved commercially by two independent, pack cementation treatments. Walsh[17] predicted that the codeposition of chromium and aluminum from pure powder sources was not feasible. For example. Fig. 16 is a plot of equilibrium vapor pressure as a function of temperature in a NaCl-activated pack with pure aluminum, chromium, and silicon. Comparable negative pressure gradients for two (or more) elements, as required for dual vapor transport, essentially never occur because of large differences in the standard Gibbs energies of formation for their respective halide species (Fig. 16). However, binary chromium-rich alloys (Cr-Al and Cr-Si alloys) exhibit highly negative deviations from ideal thermodynamic behavior, so that Cr-rich masteralloys can be used to reduce the activities of the aluminum or silicon components by several orders of magnitude (Fig. 17). 50

[75,76] Then, by using such binary alloy powders to generate comparable vapor pressures of the halide species of the two components (Fig. 18a), codeposition of aluminum and chromium into nickel- and iron-base alloys, aluminum and silicon into nickel-base alloys, or chromium and silicon into iron-base alloys was achieved experimentally if an appropriately stable halide activator salt is also provided.[40,76-79] Ravi et al.[77] used computer-assisted thermodynamic calculations to determine appropriate activators and masteralloy compositions which could simultaneously deposit chromium and aluminum into nickel-base substrates. Both chloride and fluoride salt activators were considered. Figure 18b is an example of an equilibrium calculation for an AIF^-activated pack as a function of the aluminum activity in a Cr-Al masteralloy at 1000°C. Because of the large differences in thermodynamic stabilities even of masteralloys dilute in chromium, the vapor pressures of the aluminum and chromium fluoride species differed by several orders of magnitude. Therefore, codeposition of aluminum and chromium in fluoride-activated packs was highly unfavorable. The vapor pressures of the aluminum and chromium chloride species in chloride-activated packs were much more comparable (Fig. 18a), and thus Nact %

AI - -- -

~ t ? L C r

■ u A --- — ------e ^ , --- ^ --- — ------

s L- — "■ i"" ------_ _ z r ~ i 7 jL— --- - — -— & 0 /

f ~ ~ --- 40, ' /

’O s

” 5 ^

U

ùxon of Partial temperature pres s u r e in of containing pure aluminum, chr,

a s 52

10» 1.273 K 1.173K 1.073 K

10 -'

1,073 K -' 1,173 K

10 20 30 40 50 60 70 80 90 N«(at.%)

5

4 = 3 03 en 3 2

0 0.0 0.2 0.4 0.6 0.8 1.0 Mole Fraction of Si

Figure 17: Activity data as a function of solute content for (a) aluminum and chromium in a Cr-Al alloy at 800, 900, and 1000°C [75] and (b) silicon and chromium in a Cr-Si alloy at 1050 C. [76] 53

a

" -3.0

vCcO.

Û. -5.0

O f .or.

L og A c tiv ity (Al>

Figure 18: Equilibrium partial pressures of gaseous species in (a) a CrCl.- and (b) an AlF„-activated pack as a function of the aluminum activity in the Cr-Al masteralloy at 1000 C. [77] 54 codeposition was promising. In fact, codeposition of aluminum and chromium into nickel- and iron-base alloys was achieved using chloride salt activators. No report of any kinetic study on the rate of codeposition is sited in the literature. Because a binary masteralloy containing two chlorine-active components is used, a modeling of the kinetics of codeposition are quite complex. Not only is aluminum depleted but chromium is also depleted from the masteralloy particles, reducing both the aluminum and chromium activities which can affect the overall kinetics of the process. All of these observations must be considered if an effective model is to be devised. 2.3 DEGRADATION OF PACK CEMENTATION DIFFUSION COATINGS In service, diffusion aluminide coatings are exposed to aggressive environments and applied stresses and are known to suffer two main types of degradation; mechanical and chemical failures. Mechanical failures usually do not depend on the properties of surface coatings, but, since aluminide coatings consist of an intermetallic compound, the coating can be the weak link in the overall mechanical properties (i.e., fatigue resistance and fracture toughness) of the coating/component composite.[80] Also, "low" activity coatings, which form from outward growth, have a tendency 55 to embed pack powder particles into the external layer of the coatings. These indigenous inclusions can reduce the fatigue resistance of the coating/component even further. Aluminide coatings suffer from three forms of chemical degradation: (a) interdiffusion, (b) oxidation, and (c) fused-salt attack (hot corrosion). Interdiffusion and hot corrosion indirectly affect the surface chemistry, thereby reducing the oxidation resistance of the coated component. 2.3.1 Interdiffusion Smialek and Lowell[81] studied the effect of interdiffusion on the degradation of aluminide coatings. A hybrid, "high" activity process was utilized to produce aluminide coatings on IN 100 and Mar-M200 commercial alloys. Following subsequent vacuum annealing for 300 hours at 1100°C, several changes occurred. Aluminide coatings thickened while the aluminum concentration at the surface decreased. These changes were attributed to the interdiffusion of the coating and the substrate because of the existing aluminum concentration gradient. Cyclic oxidation testing in static air at 1100°C of as-coated and vacuum annealed coatings showed substantial differences (Fig. 19). Prediffused coatings suffered severe reductions in coating lifetimes compared to as-coated specimens. Aluminum atoms are the dominant 56

diffusing element in aluminum-rich NiAl compound which is produced from this hybrid process (Fig. 4). Therefore, depletion of aluminum resulting from the reformation or growth of protective AlgO^ scales and interdiffusion with the substrate leads to the premature degradation and eventual failure of aluminide coatings. No evidence has been sited in the literature for this interdiffusion effect on "low" activity coatings. "Low" activity coatings consist of hypostoichiometric (nickel-rich) 13-NiAl, where nickel is the dominant diffusing element and may not exhibit the characteristic interdiffusion effect. 2.3.2 Oxidation The oxidation behavior of nickel-aluminum alloys has been widely investigated.[34,82-87] Pettit[34] studied the oxidation behavior of a series of nickel-aluminum alloys (e.g., Ni-3 to 25 wt.% Al) between 900 and 1300°C in 0.1 atm oxygen. The results are summarized in Fig. 20. Alloys with greater than 17 wt.% (or 32 at.%) aluminum exhibited the best oxidation behavior, parabolic kinetics associated with the formation of an external AlgO^ scale. Hindam and Smeltzer[82] studied the oxidation behavior of B-NiAl (Ni-32 wt.% or 50 at.% Al) between 1000 and 1300°C in pure Og. The results exhibited good parabolic kinetics 57

o As-coated o Preditfused, 300 hr/1100° OE

'a

- q so z o q

o

0 100 2 0 0 3 0 0 4 0 0 500 600 700 800 Time, hr

Figure 19: The effect of interdiffusion on the cyclic oxidation of pack aluminized IN 100 at 1100 C for 1 hour cycles.[81] 58

a t. %AI dO "C 1300

Al 2200 ^ ^oxide K) — 1200

steady state external Al^Oj 2000 UN, — 1100 'AI2 O3 External A l^^ < N , 'AI2 O, overtaken by ♦ NiAIjO^ 1800 1000 Exteriul NiO. Internal AI2O3 (al

1 0 0 0

1000 1100 I2 2 £2

10*' 10' 5 10 15 20 wl»AI

Figure 20: Compositional effects on the oxidation behavior of binary nickel-aluminum alloys. (a) Temperature-composition oxide phase map and (b) scale growth rates corresponding to regimes in (a).[34] 59 resulting from the formation of a ridged, external a-AlgOg scale after 30 minutes exposure. X-ray diffraction analysis also detected the formation of a metastable, transient y-Al^O^ phase. In addition, oxide spelling or loss of adherence was observed upon furnace cooling and was attributed to the formation of interfacial voids during oxide growth. The effect of aluminum content on the oxidation behavior of this wide stoichiometry compound was investigated at 900°C in pure A minimum in oxidation rate was measured at the Ni-42 at.% Al composition. This minimum corresponds with the reduction in the interdiffusion coefficient of aluminum (D^^) in the 6-NiAl phase (Fig. 4) . The mechanism of AlgO^ growth on 6-NiAl is still widely debated. Hindam and Smeltzer[89] used inert platinum markers and determined that AlgO^ growth on undoped B-NiAl compounds is controlled by inward anion diffusion. Whereas, the formation of a ridged morphology indicates a strong tendency of Al^O^ scales to grow via short-circuit diffusion paths, such as grain boundaries, by outward cation diffusion. Young and de Wit[84] used a two-stage, and oxidation treatment with secondary ion mass spectroscopy (SIMS) analysis to determine the mechanism of AlgO^ growth on yttrium-doped 60 and undoped 5-NiAl compounds. The authors found that outward AlgO^ growth was controlled by outward cation diffusion at 1000°C on undoped and yttrium-doped NiAl compounds with less than 0.07 wt.% yttrium. NiAl compounds doped with 0.5 wt.% yttrium or more exhibited inward growth controlled by inward anion diffusion. Glancing angle x-ray diffraction analysis also detected metastable AlgO^ phases, such as 0-, X-, and a-AlgOg- The transient oxidation of B-NiAl compounds has unique characteristics.[85,86] No transient NiO was observed by transmission electron microscopy (TEM), instead NiAlgO^ spinels, X-AlgO^, and 6-Al^O^ phases were detected by selective area diffraction patterns (SAD) during TEM analysis. With increasing time, transformation of the metastable AlgO^ phases (whiskers) into a-AlgOg (ridges) occurs at the scale/gas interface and proceeds inward. Metastable AlgO^ phases have much faster oxidation rates and contribute to a phenomenon termed "pest" oxidation which occurs in other intermetallic compounds.[87] The addition of a reactive element (Zr, Y, Hf), chromium, or iron facilitate the transformation into a slower-growing, more protective a-AlgOg scale. 61

The oxidation behavior of aluminide coatings with an external 6-NiAl layer has also been investigated.[88-90] Both "low" and "high" activity coatings exhibit excellent oxidation resistance. But after about 100 hours of exposure at 1200°C, the oxidation resistance drops considerably. Degradation of aluminide coatings occurs because of the depletion of aluminum from the coating by the repetitive formation and spalling of protective AlgO^ scales. Eventually, the reduced aluminum surface concentrations are insufficient to support the formation of a protective AlgOg scale; instead, less protective NiAl^O^ and NiO scales form. Aluminum depletion from the B-NiAl coating occurs until the minimum aluminum concentration to sustain a B-NiAl phase is approached, whereupon B-NiAl is transformed into V'-Ni^Al according to Eg. (2.6). 3 NiAl(s) + 3/2 Ogfg) = AlgO^(s) + NigAl(s) (2.6) As aluminum depletion continues, transformation into subsequent phases containing lower aluminum (e.g., X -Ni) also occurs (Fig. 21). The degradation of aluminide coatings can be reduced if the adherence of protective AlgO^ scales is improved, reducing the aluminum depletion rate. In particular, additions of reactive elements (e.g., Y, Zr, Hf) considerably improved the adherence of thermally 62

y'twj^Aiiy-N. SOCIO SOLUTION

^(N iA ll4y (KUAII 4REFPACT0AY WETAL CARBIDES

SUBSTRATE 4 REFRACTORY METAL CARBIDES (AOCULARI \ \\ CTCH*NO.I'fNOl2€. «€ l(R

r -N « SOLID s o l u t i o n - y * (N U A tf ------

^(NIA(|4 7*(Ni3ACl 4 REFRACTORY METAL CARBIDES * » fl6 0

SUBSTRATE I v r . ETCH'NO.I+NaZ 230% k. too HR

SURFACe XRO HCSULTSa AlzOa + H IA IgO ^+y-W + ylN lgA II

r'lMjAi)r-NiSOUOSOUmON-----

METAL CARSfOeS

suasnuoE

ETCH MX I + N 0.2 ZSOX c ZOO HR

Figure 21: Microstructures, microhardness data, and phase identities for pack aluminized Mar-M200 alloys after oxidation at 1200 C.[88] 63 grown AlgO^ scales on bulk B-NiAl.[23-27,91] Therefore, one can expect that similar additions would provide the same beneficial effect on aluminide coatings. 2.3.3 Fused Salt Attack (Hot Corrosion) Hot corrosion is the major mode of material degradation in marine and industrial gas turbines and may occur in aircraft engines. Hot corrosion is defined as the accelerated attack of a metallic component resulting from the presence of a condensed fused salt film (e.g., NagSO^, NaVO^/ NaCl) in an oxidizing environment. One generally accepted mechanism for the formation of NagSO^ is the sulfation of NaCl vapor, introduced as an impurity either in the fuel or the ingested air. The NaCl reacts with the oxide combustion products of sulfur in the fuel according to Eqs. (2.7) or (2.8).

2NaCl(V)+SO 2 (g )+I/ 2 O 2 (g )tHgO(V)=Na2S0^(c)+2HC1(v) (2.7) or

2 NaCl(v) + 8 0 2 (g) 0 3 (9 ) = Na2 S0 ^ (c) + CI 2 (g) (2.8) When the vapor pressure of Na2S0^ exceeds its dew point for the given atmospheric conditions, condensation of a fused salt on the cooler turbine components occurs.[92] Two different types of hot corrosion exist. At

temperatures above the melting point of pure Na 2 S0 ^

(T^^884°C) in air or an SO 2 /O 2 mixture, "high temperature" or type I hot corrosion occurs and is very 64 common in aircraft turbines. Type II, "low temperature" hot corrosion (LTHC) occurs at moderate temperatures (600-750°C) well below the melting point of pure NUgSO^. LTHC is characterized by the presence of a low melting MSO^-NagSO^ eutectic (where M=Co or Ni) in a high atmosphere needed to equilibrate the sulfate eutectic. [104] This type of hot corrosion is typically found within industrial and marine turbines. The mechanisms of high temperature hot corrosion have been rather extensively studied.[92-100] Hot corrosion is characterized by two general types of attack: passive and locally active attack. Passive attack is the fluxing or dissolution (basic or acidic) and precipitation reactions of the protective oxide scale, which formed during exposure, with the molten salt. For either acidic or basic dissolution to be self-sustaining, a negative solubility gradient of the oxide in the molten salt is required. Differences in the melt basicity at the oxide/melt and the melt/gas interfaces establish a negative solubility gradient of the oxide (Fig. 22). Oxide dissolves at the oxide/melt interface, where the oxide solubility is largest, and the dissolved species diffuses down a concentration gradient to a region of lower solubility. The dissolved species 65

ttûa fused salt film

OKidc M eld gas phase 0^,50,. SO, Solubility

M i » porous oxide precipilafcs j locolly low MO solubility

locally bigh (acidic or basic 1 solubility for MO

Figure 22: Dissolution and reprecipitation of a porous MO oxide supported by the solubility gradient in the fused salt film. [100] 66

then precipitates out of the melt as a porous, nonprotective oxide. This model was proposed and illustrated by Rapp and Goto.[100] Locally active attack occurs when the molten salt penetrates the protective oxide by fluxing or through flaws where it contacts and "cathodically overpolarizes" the underlying alloy or coating underneath forming internal sulfides. Sulfide precipitation increases the local basicity or oxygen ion (0^~) activity, leading to further basic dissolution. Sulfide precipitation and oxide dissolution/precipitation reactions also deplete chromium, beneficial for resistance to hot corrosion attack, and aluminum from the surface of the alloy or coating, inhibiting the ability to reform a protective scale and leading to premature failure. To improve the resistance to hot corrosion, an alloy or coating which forms a slow-growing and dense protective oxide scale with limited solubility in NUgSO^ is necessary. Extensive studies were conducted by Rapp and coworkers[101] to measure the solubility of several protective oxides in pure Na^SO^. The results determined that CrgO] and AlgO^ are acidic oxides because their solubility minima occur at very low basicity and they therefore offer substantially better resistance to acidic hot corrosion than the more basic oxides such as Fb ^O^, 67

4 CfoO'

z

CC ,NiO I— to rAUO 2 ÜJ O C o % 0 ,

5 7 9 II 13 15 17 -LOG

Figmre 23: Measured oxide solubilities in pure Na.SO^ at 927 C (1200K) and 1 atm oxygen.[101] 68

NiO, and CoO. However, two anomalies were noted. Silica has no acidic solute and thereby is not attacked by an acidic melt; therefore, SiOg offers the potential for the best resistance to acidic hot corrosion.[102] Chromia-forming alloys or coatings were found to have better resistance to both acidic and basic hot corrosion than alumina-formers.[103] From Fig. 23, CrgOg is slightly more soluble than AlgO^ in pure NUgSO^ with similar acidic solubility behavior. Therefore, the measured oxide solubilities cannot alone explain this enhanced hot corrosion resistance. 2.4 REACTIVE ELEMENT EFFECT (REE) Heat-resistant alloys pick up small amounts (

2.4.1 Observations The REE is characterized by several differing experimental observations depending on the oxide (AlgOg or CrgOg) in question. Very small reactive element alloy additions curtail transient oxidation by supporting the selective oxidation of chromium or aluminum to form a continuous, external oxide scale. RE additions were also observed to reduce the parabolic rate constant, k^, for the growth of an external, protective oxide scale (Fig. 24).[107] The critical solute content (Ng*, Eq.(2.2)) needed to sustain the growth of an external Cr^O^ scale at steady-state was also reduced; however, no significant reduction in the critical solute content needed to form an external AlgO^ scale was observed. Chromia growth occurs from the outward diffusion of cations via both interstitials and vacancies, producing an outward grown scale. In the presence of REE, inert marker studies indicate a shift in the mechanism of Cr^O^ growth. Whereas, alumina growth occurs from the dominant inward diffusion of anions except for undoped B-NiAl compounds, and no change in growth mechanism was observed on RE-doped alloys or coatings.[111-113] Surface 70

,-10

-Ni30 Cr at 1050°C—

,11 Ni-30 Cr at 1000°C* « Ni-30 Cr-Ce at 1050°C oE

,12

Ni-30 Cr-Ce at 1000°C

,13 10 ,16 17

"Ce" CONCENTRATION

Figure 24: Parabolic rate constant for chromia growth on unimplanted (dashed lines) and cerium implanted Ni-3gCr alloys as a function of cerium dosage at 1000 and 1050 C. [107] 71

microanalysis (EDS with STEM or SAD with TEM) also detected RE ions and compounds at the oxide/metal interface and at oxide grain boundaries of these oxidized alloys.[107-110] RE additions were also observed to suppress the formation of interfacial voids and other convoluted scale morphologies.[19-22] Overall, the scale adherence for both CrgOg- and AlgO^-formers is improved by RE additions. 2.4.2 Mechanical Behavior Several theories have been proposed to interpret the above observations and contribute to improved scale adherence.[19-22,104-130] The first model addresses the mechanical behavior of protective scales grown on RE-doped alloys or coatings. These scales presumably possess enhanced plasticity to relieve stresses induced during oxide growth or by thermal cycles. But Nicholls and Hancock[115] failed to measure any improvement in plasticity or fracture toughness of polycrystalline CrgO^ with additions. However, RE additions did reduce compressive growth stresses within the scale and the through thickness defect size, a, below critical values. In addition, â fine-grained scale is 72 produced on RE-doped alloys and coatings and these may better accommodate stresses by grain boundary sliding and other creep mechanisms (e.g., Coble creep).[118,119] 2.4.3 Chemical Behavior RE additions have also been proposed to dope protective oxide scales, reducing ionic conductivity and/or defect concentrations. Huntz[116] studied the effect of YgOg additions on the conductivity of polycrystalline AlgO^. The results indicate that larger ions tie up or trap aluminum vacancies decreasing cation diffusivity especially at grain boundaries. Patibandla et al.[107] studied the effect of cerium implantation on the oxidation behavior of Ni-30Cr alloys. The authors present evidence supporting

Ce-doping, specifically at grain boundaries, of Cr^O^ scales. Cerium additions inhibit the mobility of chromium defects, changing the growth of CrgOg scales from outward diffusion of chromium interstitials to the inward diffusion of oxygen vacancies. Further observations by TEM analysis detected RE ions or RE-rich precipitates at or near oxide grain boundaries.[107,109-113,120] These authors propose that segregants block short-circuit diffusion paths, which otherwise contribute to oxide growth on undoped alloys or coatings. Convoluted AlgO^ and CrgO^ scales are often 73 formed on undoped alloys and coatings.[19] The segregation of larger RE atoms or ions to oxide grain boundaries may reduce the mobility of oxide grain boundaries and grain growth by a solute-drag effect, thus inhibiting the formation of convoluted scales and improving scale adherence. A graded seal mechanism was proposed by several authors.[117,121] This mechanism relies on the formation of an oxide compound with an intermediate coefficient of thermal expansion to that of the metal and external oxide. Unfortunately, no evidence was presented for the existence of a continuous intermediate compound; instead, oxide pegs of an internal, RE-rich oxide were detected beneath the external scale (Fig. 25).[19,21,22,106] These oxide pegs may help to key or anchor external oxide scales mechanically to their metal substrates. Some examples existed where exceptional scale adherence was produced, while oxide pegs were observed. However, there were other examples where good scale adherence resulted without oxide pegs being present and poor adherence resulted with the presence of oxide pegs. Therefore, oxide pegs are not necessarily required for improved scale adherence. 74

oxiOc penetration due to Oxidation ol interrnetallic particle

oxide

alloy

^~^yttrium-enriched intergranular oxide

Figure 25: Schematic illustration of the oxide/scale and alloy/oxide interface of an yttrium-doped, FeCrAl alloy after oxidation for 24 hours at 1200 C. [19] 75

Another interesting observation is the presence of convoluted scales consisting of interfacial voids on undoped, AlgOg-forming alloys or coatings.[20-22,114] RE additions have been found to improve oxide adherence by eliminating or suppressing void formation. Void annihilation has been attributed to the presence of RE oxide dispersions or to the alteration of the growth mechanism and the predominant diffusing defect. However, no conclusive evidence has been presented of void annihilation at oxide dispersions in the alloy or scale. 2.4.4 New Theories Electron spectroscopy for chemical analysis (ESCA) and Auger electron spectroscopy (AES) analyses have detected both elemental and ionic sulfur species at both NiCrAl surfaces as well as at the thermally grown oxide/metal interface. [123,124] Sulfur segregation has been found to reduce the scale adherence on undoped NiCrAl alloys or coatings (Fig. 26), creating an interpretation known as the "sulfur effect".[125] RE additions are commonly used in industry to desulfurize heat-resistant alloys. Therefore, RE additions react with sulfur to form sulfide inclusions and to reduce the residual sulfur content in the alloy or coating, eliminating segregation to the scale/metal interface where such impurities can reduce the chemical bond. 76

High purity NiCrAl (1-2 ppm S)

Mass change/ Normal purity unit area, NICrAlY m g/cm 2

Normal purity NICrAI (20-40 ppm S)

20 . 40 60 80 100 Number of one hour cycles

Figure 26: Effect of sulfur content on the 1180 C cyclic oxidation performance of NiCrAl alloys. [125] 77

Luthra and Briant [126] have refuted this theory. The authors state that no evidence has been presented linking scale adherence to interfacial sulfur segregation. In their own studies[127], they showed that scale adherence was improved on NiCrAl alloys with high residual sulfur levels. Sraialek [128,129], using a novel technique, repeatedly grinding away the oxide formed during each oxidation cycle on an undoped NiCrAl alloy. This technique successfully improved the adherence of AlgOg scales, presumably by sequentially reducing interfacial sulfur and the bulk sulfur contents. Other observations were noted. During the transition from nonadherent-to-adherent scales, no reduction in the quantity of interfacial voids or in growth stresses or the presence of oxide pegs was observed. Recently, Pieraggi and Rapp [130] presented a new interpretation of the RE effect on CrgOg^fo^^ing alloys or coatings. The authors state that the adsorption of RE ions to the oxide/metal interface inhibits the climb into the alloy of interfacial misfit dislocations, otherwise thought to annihilate cation vacancies created during cation-diffusing scale growth. Using a Grove and Deal model for the interpretation of the oxidation kinetics, the authors show that RE additions can reduce the linear interfacial rate constant of the overall oxidation 78

reaction, thereby preventing cation diffusion. In this manner, anion diffusion would gain dominance and the observed changes in mechanism and scale adherence and reduction in the parabolic rate constant would result. CHAPTER III EXPERIMENTAL PROCEDURE

3.1 MATERIALS Both conventionally cast (obtained from Tom Kircher, Naval Air Development Center) and directionally solidified (obtained from Bill Connor, GE Aircraft Engines) columnar-grained and single-crystal nickel-base alloys were used for this investigation. Their nominal compositions are given in Table 2. Each of these alloys is a multi-phase (y, Y ', and carbides) alloy at room temperature, with alloying elements to stabilize X-Ni (e.g., Cr, Co, W, Mo) and X'-Ni^Al (e.g., Ti, Ta, Nb) or to form carbides (e.g., W, Mo, Ti, Ta, Hf, Nb). Coupons were sectioned into 1 mm thick specimens approximately 1 cm x 1 cm using an electric discharge machine. These were ground through 400 grit S i c paper, measured, cleaned with soap and water, ultrasonically degreased in acetone and methanol, and weighed. Specimens were placed in an alumina crucible with a pack containing a source powder, an activator salt, and an alumina filler powder. The pack powder mixtures consisted of: between 15 and 25 wt.% binary masteralloy powder (<100 mesh; Cerac, Inc.), 2 and 6 wt.% chloride activator

79 80

Table 2: Nominal Compositions of Nickel-Base Alloys Coated (at.%). Alloy Ni Cr Co Nb Mo W A1 Ti Other IN 713LC bal 13 — 1 2.5 - 12 .7 .2C,.06Zr, .53 Mar-M247 bal 9 10 .4 3.2 12 1.2 .8C,lTa, ITa,.13, .06Zr,.5Hf René 80 bal 15 9 2.3 1.2 6 6 .8C,.083, .02Zr René 80H^ bal 15 9 2.3 1.2 6 6 .8C,.083, .OlZr,.2Hf René N4 bal 11 7 .5 .3 1 2 9 4 .3C,.05Hf, 1.5Ta,.023, . OlZr a=columnar-gra ined, directionally solidified b=single-crystal, II II

Table 3: Composition of the (chromium-aluminum binary masteralloy powders. Actual Composition (wt.%) Masteralloy fwt.%) Cr A1 V Fe Co Ca Cr-lOAl bal 9.74 .68 .12 .02 .03 Cr-7. 5A1 bal 6.61 .76 .11 .01 .04 Cr-6A1 bal 5.39 .61 .11 .04 .04 Cr-5A1 bal 4.69 .49 .12 .04 .05 ♦Standard spectroscopic methods 81

Table 4: Properties of the chloride activator salts used in this investigation. Vapor Pressure Activator Normal Tm Normal T» m . at 1423K Volatile: AlCl 193 195^ 1 atm ZrCl^ 437 336^ II II HfCl! — glG* II II NH Cl 340 ,520 " " (HCl) Condensed: YCl 721 1507 0.157 atm CrCI_ 825 - 0.412 " a=sublimes b=dissociates into NHg(g) + HCl(g) 82 salt (reagent grade), 67 and 83 wt.% inert alumina filler powder (80-200 mesh), and 2 and 6 wt.% of either unstabilized monoclinic zirconia (Atlantic Equipment Engineers), silica gel (Fisher Scientific), or otherwise yttria (Johnson Matthey), or else no RE oxide source. The masteralloys and the chloride activator salts used in this investigation are presented in Tables 3 and 4, respectively. Activators are categorized according to their stability and physical state at the treatment temperature. 3.2 DIFFUSION COATING PROCEDURES Two substrate/powder pack arrangements were used to produce the coatings. The first (powder contacting) arrangement consisted of placing the substrate in intimate contact with the powder mixture in the alumina retort (Fig. 27). By the second (above pack) arrangement, the substrate was physically isolated from the powder mixture by a smaller, porous alumina enclosure (Selee Corp.) plugged on both ends by porous alumina paper (Zircar Products, Inc.). The pack powder mixtures were prepared by weighing the powders which were then mixed thoroughly in a ball mill for up to four hours. The sealed enclosure was then centered in the alumina crucible and filled with the pack powder mixture similar to the first arrangement 83

Alunina Lid Alunina Cement Dia. - 3.1 cm

Alumina Crucible

Nickel-Base Substrates ii iiii 6.5 cn

iiiiiliiiiiiiiiiiii ::::::

Pack. Powders - 73 wt.% Alumina Inert Filler - 25 wt.% Cr-Al Binary Alloy - 2 wt.% VClg or ZrCl^ Activator - 2 wt.% VjOg or ZrOg Filler Replacement

Figure 27: Schematic diagram of the "powder contacting" pack arrangement. 84

Alumina Lid

m m ! Alumina Cement Alumina Sheath mg Alumina Crucible

Substrate

Porous Alumina Container

Pack Powders - 73 wt.t Inert Filler - 25,wt.% Masteralloy - 2 wt.% Halide Actiuator - 2 wt.% Reactive Element Oxide (replace inert filler)

Figure 28: Schematic diagram of the "above pack" arrangement. 8 5

Average Heating Profile of Lindberg Furnace (58.5 % OP)

1200 (1150°C in 5 hrs.)

1000 - -

600--

600--

400 --

Z O O - -

100 I S OSO 200 250 300 Heating Time. min.

Figure 29: Average heating profile of Lindberg furnace. 8 6

(Fig. 28). The alumina crucible was closed by an alumina lid using an alumina-base cement (Aremco Products, Inc.) and cured for at least twelve hours. The cured crucible was placed in a horizontal tube furnace (Lindberg), purged with prepurified argon at a flow rate of 0.1 1/min (STP), and heated (Fig. 29) to a desired temperature, usually 1150±5°C, for a desired length of time, typically twenty-four hours. Figure 30 is a schematic diagram of the furnace setup. After cooling the furnace, the retort was removed and opened. The coated substrates were cleaned with soap and water, ultrasonically degreased with acetone and methanol; then they were weighed and measured. 3.3 CHARACTERIZATION OF THE DIFFUSION COATINGS To examine the substrate cross-sections, specimens were mounted in epoxy, sectioned with a low-speed diamond blade saw, and polished through 0.3 /zm alumina. Early work used an acetic glyceregia etchant (10 ml acetic acid, 10 ml nitric acid, 20 ml hydrochloric acid, and 30-40 ml glycerol) to distinguish interphase and grain boundaries within the coating, but optical microscopy later showed that polarized light was sufficient. A Nikon Epiphot-TME inverted microscope was used for all the optical microscopy and photomicroscopy. 87

Pack Cementation Furnace Setup

1. Argon Gas 9. Blower 2. COj Getter lO. liorlzontal Tube Furnace 3. Desiccant 11. Furnace Controller 4. Getter 12. Alumina Crucible 5- Cold Junction 13. Nickel-Base Samples

6. Multimeter 14. Kullitc Reaction Tube 7. Mi-Cr/Ki-Al Thermocouple 15. Bubbler 8. Flange

12

IS.

Figure 30: Schematic diagram of the pack cementation furnace setup. 88

A Scintag PAD V x-ray diffractometer using Cu Ka radiation at 45 kV accelerating voltage and 20 mA current was used to determine the phases of the coating. A fully automated Cameca SX-50 microprobe equipped with four wavelength dispersive spectrometers (WDS) and their respective diffracting crystals was used so that quantitative microanalysis of elements with atomic numbers greater than boron could be measured. The Cameca microprobe system measures the intensity of a given energy level (characteristic for each element) of the specimen and a calibrated standard along with the ZAF correction factor to determine the composition (to 100 ppm) of the surface and the various phases contained in the coating. Because of the excessive excitation volume of the electron beam, x-ray elemental dot maps were also used to determine and distinguish these phases. In addition, secondary ion mass spectroscopy (SIMS) analysis was conducted by Dr. Rex Hussey at the National Research Council, Ottawa, Canada. A Cameca FX-25 SIMS with a Ga^ incident beam was used to qualitatively identify, with exceptional accuracy, the composition of the coating surface. Microhardnesses (VHN) of selected coatings were measured on mounted specimens using a Buehler Micromet II, microhardness tester. A diamond pyramid indenter using a 50 gram load and a 10 second dwell time was used. 89

3.4 ANALYSIS OF VAPOR SPECIES EVOLVED FROM PACK MIXTURES An atmospheric pressure sampling mass spectrometer (NASA Lewis Research Center) was used to analyze the halide vapor species evolved from a diffusion coating pack powder mixture. The pack mixtures were comprised of 50 wt.% alumina powder, 30 wt.% Cr-lOAl (wt.%) binary masteralloy powder, 10 wt.% anhydrous ammonium chloride activator salt, and either 10 wt.% zirconia or otherwise yttria, or else no RE oxide (control) . The pack powder mixture was weighed and mixed thoroughly, and one-gram samples were placed inside a small, open platinum crucible mounted to a Pt/Pt-Rh thermocouple. The Pt crucible was positioned in a small tube furnace below the sampling orifice of the mass spectrometer (Fig. 31a). The furnace and quartz reaction tube were then raised to the sampling orifice of the mass spectrometer, purged with a 5% Hg/Ar gas mixture at a flow rate of 0.5 1/min (STP) to eliminate masteralloy oxidation, and then heated to 1000+5°C. The crucible was raised to 2.5 cm below the sampling orifice, and the resulting vapor species were analyzed at 800-1000°C during heat up and for one hour at temperature. To supplement the mass spectrometer experiments, some collection of the vapors on a cool target was also done. The same small furnace was used with a cool copper block placed at the top to collect the condensate (Fig. 31b) . 90

Owadn/pok m att bie<

tO O O L S

> MuRMTta tfwmooouple

■OnriQt

C a t «r t — .

Figure 31: Schematic diagram of (a) the atmospheric sampling mass spectrometer and (b) the target collection apparatus. 91

This deposit was analyzed by standard wet chemistry methods (e.g., atomic emission and x-ray fluorescence spectroscopy) for Al, Cr, Y, and Cl. In addition, x-ray diffraction was conducted, as explained earlier, on the spent pack powder mixtures from both types of experiments to determine whether any condensed products were produced. 3.5 ENVIRONMENTAL DURABILITY TESTING 3.5.1 Oxidation High-temperature oxidation behavior of the coatings was determined using both isothermal and cyclic oxidation tests. Isothermal oxidation at 1100°C in air (0.21 atm Og) at a flow rate of 0.1 1/min (STP) was conducted using a Cahn TG-171 automated thermogravimetr i c analyzer (TGA) equipped with the Cahn DACS plus software. Specimens were placed in an alumina crucible supported with a sapphire wire (Fig. 32). Weight change versus time was measured and recorded, and the parabolic rate constant of the oxidation reaction was calculated. X-ray diffraction (XRD) analysis was used to determine the oxide and coating phases present. In addition, a Hitachi S-510 scanning electron microscope (SEM) was used to examine the oxide morphology, and energy dispersive spectroscopy (EDS) was used to distinguish the compositional features of the oxidized coatings. 92

Cyclic oxidation at 1100 and 1200°C in static air was performed at the NASA Lewis Research Center under the guidance of Drs. James L. Smialek, Charles A. Barrett, and James A. Nesbitt. Specimens are placed in a FeCrAl crucible suspended by a platinum wire (Fig. 33). Specimens were pneumatically lowered and retracted from the hot zone after one hour exposures followed by a twenty minute cool to ambient temperature. Specimens were removed periodically to be weighed and inspected. Screening tests for up to 200, one-hour cycles were conducted. Optimum coatings were further evaluated for up to 500, one-hour cycles. XRD analysis was again used to determine the oxide and coating phases present. Specimens were mounted in epoxy, sectioned, and polished through 0.3

(im alumina. These cross-sections were analyzed by EPMA to determine compositional and morphological changes in the coating. 3.5.2 Hot Corrosion (Fused Salt Attack) Studies The hot corrosion behavior of the coatings was determined with both isothermal thin film and burner rig evaluation. Thin films of NagSO^ were applied to specimens by heating to about 200°C on a hot plate and coating with an aqueous solution from an airbrush. A salt coating of 5+1.5 mg NagSO^/cm was produced. Coated specimens were placed in an alumina boat which was 93

Electronic M irrobjljncc

T o r e

Cis Exil [ Purging C as (Arl Uppt,FUnge^^B

Fitnuce

C o u p o n

Lo«vrr Flangej Ei-ma

F C as Inlet

T o r e A rra n g e m e n t fo r Degradation Testing The Ohio Stale ÜÉersily

Figure 32: Schematic diagram of the thermogravimetric analyzer (TGA) used for isothermal oxidation testing. 94

NASA-LEWIS MULTITUBE AUTOMATIC HIGH TEMPERATURE CYCLIC OXIDATION

I - CYCLE COUNTER

r-INDIVIDUAL SPALL SPALL SHIELOS SAMPIE I CUP COLLECTOR RETRACTOR L /- SPALL CUP (PNEUMATICI ! r SPALL CUP V COVER ' I RETRACTOR CERAMIC TEST i IPNEUMATKI TUBE 13.5 CM 1.0 .1

SAMPLES COOIEO TO AMBIENT TEMP

INTERVAL TIMERS & COUNTER CONTROL

SAMPLES AT TEST TEMP

Figure 33: Schematic diagram of the cyclic oxidation test apparatus. 95

BQt_Corrosion Test Rig 1. Cas Mixture (SO^/O^) 9. Furnace Controller 2. CO^ Getter 10. Horizontal Tube Furnace

3. Dcsiccant 11. Platinum Catalyst

4. Cold Junction 12. Coated Substrates

5. Multimeter 13. Alumina Crucible 6. Pt/Pt-Rh Thermocouple 14. Mullltc Reaction Tube

7. Flange IS. Bubbler 8. Blower

10 14

Figure 34: Schematic diagram of the thin film hot corrosion test setup. 96

inserted into a horizontal tube furnace (Fig. 34) at 900°C

and exposed to a platinum-catalyzed 0.1% SO^/0^ mixture at a flow rate of 0.1 1/min (STP). Samples were removed periodically to be weighed, inspected, and recoated with salt. XRD analysis was used to determine the crystalline product scales which formed. Unfortunately, XRD analysis could not identify any amorphous products, common in these types of reactions. Cross-sections were prepared by sputtering a layer of gold/palladium and then electroplating a thick layer of copper over the sample. These cross-sections were then examined with the EPMA so that compositions and x-ray maps of S, O, Na, and other coating elements could be obtained. Type I (high-temperature) hot corrosion burner rig tests were performed by Ron P. Bissett at GE Aircraft Division, Evendale, Ohio. Coated nickel-base superalloys were run in a mach 1 burner rig at 927°C using JP-5 jet fuel with 0.4% sulfur added and 5 ppm sea salt injected into the flame. The tests were conducted until observations indicated coating failure, purely a qualitative decision by the technician. After the run, the samples were analyzed using the techniques described above. CHAPTER IV VOLATILE SPECIES IN CHLORIDE-ACTIVATED DIFFUSION PACKS

4.1 ITSOL COMPUTER PROGRAM The first step to deposit chromium, aluminum, and a RE in a single-step, pack cementation process is to determine potential pack chemistries (e.g., masteralloy composition and activator salt) which could achieve this goal. For dual vapor transport, a positive flux of each species, according to Eg. (2.3), from the pack to the substrate surface is required. In general, simultaneous deposition is possible as long as comparable vapor pressures of the chromium and aluminum chloride species are generated. Thus, the equilibrium vapor pressures of all the chloride species evolved from each pack chemistry should be calculated so that trial and error experimentation can be eliminated. Calculations by hand for these complex equilibrium systems containing numerous condensed and gaseous species are quite tedious. However, an existing computer program (ITSOL) assist equilibrium calculations of this type.

97 98

White and coworkers[131] developed a method based on the minimization of the total free energy of the system to calculate the equilibrium composition of complex gas systems. Oliver et al.[132] successfully demonstrated the free energy minimization (FEM) method for a combustion process. Nevertheless, Eriksson[133] improved this method to include systems containing a gas phase and multiple, invariant condensed phases. Further modifications eventually produced the user-friendly, ITSOL computer program which is commonly used on IBM personal computers.[134,135] The program calculates the equilibrium composition by assigning molar amounts to the elements of all possible equilibria between gaseous and condensed species in the system and then, by a series of iterations, the program determines the composition which minimizes the total Gibbs free energy of the system while conserving mass. The partial pressures of the gaseous species are calculated from their equilibrium molar amounts, the total number of moles in the gas phase, and the total pressure of the system set by the user. If the standard Gibbs energy of formation of all condensed and gaseous species are known as well as activities for all nonideal condensed species, then the equilibrium composition for the given pack chemistry can be calculated. 99

4.2 PACK EQUILIBRIUM CALCULATIONS The ITSOL program was used to calculate the equilibrium composition for a series of activator salts, reactive element sources, and masteralloy compositions. Equilibrium calculations were performed using the following assumptions: (a) local equilibrium exists between the pack powder mixture and was unaffected by the substrate, (b) mass is conserved within the alumina retort, (c) a constant, total pressure of one atmosphere is used because constant volume conditions generated unreasonable pressure values (see Appendix A, for sample calculations), and (d) since chromium-chloride species have a greater temperature dependence (Fig.16), the highest conceivable temperature of 1150°C was used. The number of moles of argon gas was calculated by the ideal gas law to be .0004 and was used in all ITSOL calculations. A sample calculation is demonstrated next. 4.2.1 Sample Inputs The molar amounts are determined for a 50 gram pack mixture containing 25 wt.% masteralloy (binary Cr-Al alloy powder), 2 wt.% chloride activator salt (Table 4), and 73 wt.% inert filler powder (Al^O^). For most cases[42,136], the alumina filler powder was relatively inert and did not contribute significantly to the equilibrium composition; therefore, molar amounts of the filler powder were not 100 input to the ITSOL program. The equilibrium composition was calculated by inputting theinitial moles of each element of the masteralloy and activator and the standard Gibbs energies of formation of each possible vapor or condensed species (Tables 5-9). The ITSOL program, as developed by Flynn et al.[137], calculates equilibrium compositions for pure, invariant condensed phases with unit activity. But if the chemical activities of the nonideal species are known, then a modified standard Gibbs energy of formation of the nonideal species, AG^°*, can be introduced; further conversion of the chemical activities to a new standard state can be avoided. This method (developed at OSU) artificially assigns the aluminum (and chromium) activity (a^^ ie ) as unity, while is defined as; AG^1°* = RTln (4.1) where a^^ is the measured aluminum (and chromium) activity in the Cr-Al binary masteralloy powder. This method was verified with the following reaction:

A1 + Clg(g) = AlClg(V) (4.2) where the equilibrium product is constant for a given temperature and pressure and is given by

K^q=p^^ç,]^2/®A1^C12 ■ Therefore, equilibrium compositions for different aluminum activities at a constant temperature and pressure were input to the equation and 101

Table 5: Standard Gibbs energies of formation for the aluminum species used as input for the ITSOL program.

AG°- (cal/mol) Species______1200K______1423K___ Reference Gases: A1 35,579 41,623 138 AlCl -35,188 -35,131 II A1C1_ -75,687 -76,381 II AlCl^ -124,475 -120,952 II -238,053 -224,109 II —6,248 -9,959 II A10_ -46,938 -24,640 II A1_0 -55,322 -60,002 II -104,390 -102,821 II AÎêîë -83,408 -82,712 139 AlH 37,634 34,069 II A1H_ 47,089 45,967 II AlH^ 59,231 62,085 II AlOH -44,716 6,224 II AIN 97,145 92,764 138 (ndensed: -309,560 -292,013 138 -44,601 -38,762 II Al^SiOg — -451,532 139 102

Table 6: Standard Gibbs energies of formation of the chromium species used as input for the ITSOL program.

AG°^ (cal/mol) Species 1200K 1423K Reference Gases: Cr 52,938 45,521 138 CrCl -50,641 -53,546 140 CrCl^ -66,618 -65,853 II CrCl^ -72,505 -67,106 II C r o 14,165 8,933 138 CrO, -27,965 -29,309 II CrO^ -50,215 -46,281 II CrN 90,403 85,214 II Condensed: -197,178 -183,767 138 -60,367 -56,052 II CrCl^ -69,337 - II Cr^N"* -9,024 -5,463 II CrN -6,010 -2,144 II 1 0 3

Table 7: Standard Gibbs energies of formation for the zirconium species used as input for the ITSOL program.

AG ^ (cal/mol) Species 1200K 1423K Reference Gases: Zr 107,839 97,584 138 ZrCl 21,201 16,646 II ZrCl -52,404 -53,546 II ZrCl^ -115,212 -113,114 II ZrCl^ -174,808 -168,564 II ZrO -7,705 -12,935 II ZrO -75,259 -75,943 II ZrN 107,839 101,584 II Condensed: ZrO -208,679 -198,801 138 ZrCl_ -63,206 -57,987 II ZrN -60,205 -55,733 II 1 0 4

Table 8: Standard Gibbs energies of formation of the yttrium species used as input for the ITSOL program.

Ac „ (cal/mol) Species 1200K ______1423K Reference Gases: Y 62,683 55,948 138 YCl 23,936 19,844 141 -77,915 -77,944 II YCl^ -165,446 -165,455 II II ^0 * -32,424 -36,286 Y-O * -17,453 -20,042 II Y?0_ -125.304 -124.184 II Condensed: -372,556 -357,741 138 -174,658 -170,218 141 YN -42.530 -37.280 138 *=Estimated S from reference 154. 1 0 5

Table 9; Standard Gibbs energies of formation of the silicon species used as input for the ITSOL program.

A g °- (cal/mol) Species 1423K Reference Gases: Si 57,575 138 Si, 77,890 Si? 64,289 SiÔl 12,115 SiCl, -52,529 sici:r -81,309 sici; -113,978 Sic ^ -52,693 sio_ -77,720 SiN^ 55,723 Si-N 54,836 Siê 53,836 SiH- 40,859 SiH^ 53,575 SiH^ 39,207 Condensed: SiO_ -157,235 138 Si3«4 -65,176 II 1 0 6

yielded a constant equilibrium product, which confirmed this method for the nonideal condensed phases. For this investigation, the activities of aluminum and chromium were calculated from values measured at 1000°C by Johnson et al.[75] (Fig. 17) and the use of the Gibbs-Helmholtz relationship listed below: In a^^" - In a^^' = (AH^^^/R) (1/T--1/T') (4.3) where Aïï^^ is the partial molar enthalpy of aluminum, a^^" is the activity of aluminum at temperature T", is the activity of aluminum at temperature T', and R is the universal gas constant. The corresponding activity values for chromium were calculated using a graphical Gibbs-Duhem integration for each alloy composition. The activity values and standard Gibbs energies of formation (Eg. (4.1)) for chromium and aluminum in the Cr-Al binary alloy for use in the ITSOL program are listed in Table 10. 4.2.2 Sample Outputs A sample calculation including initial molar inputs and the equilibrium composition of an NH^Cl-activated pack at 1150°C was presented in Table 11. A series of calculations as a function of the aluminum activity in the Cr-Al binary masteralloy is plotted in Fig. 35. Several results should be noted. The chlorine source, HCl, reacted with the Cr and/or A1 source according to the following reaction: 1 0 7

Table 10: Activities and corresponding Gibbs energies of formation of chromium and aluminum for various masteralloy compositions at 1200 and 1423K.[75]

Masteralloy Activity AG f Composition (cal/mol) (wt.%) A1______Cr A1______Cr 1200K: Cr-5A1 .0010 .88 -16,100 -298 Cr-lOAl ,0030 .77 -13,540 -609 Cr-15A1 ,0084 .56 -11,140 -1351 Cr-20A1 ,0220 .37 -8,896 -2317 1423K: Cr-5A1 0026 .95 -16,830 -139 Cr-lOAl 0054 .85 -14,763 -453 Cr-15A1 0156 .65 -11,764 -1209 Cr-20A1 0383 .49 -9,224 -1994 108

Table 11: ITSOL pack equilibrium output for a 2 wt.% NH.Cl-activated pack containing 25 wt.% Cr-lOAl wt.% masteralloy at 1150 C.

Cr/M Cedepositicr., STEP NO: 1

I = 1«3.00 K P = 1.000£<00 ATM

ELEMENT MCIES Ar .0004 K .0187 H .0748 AI .0460 Cr .2160 Cl .0187

EOUIUBRlUn aWOSITIDNE;

PHASE 94ses SPECIE INIT.EST. EQ.nOLES P/ATH ACTIVITY M2 .0000*400 .3716*-01 .8383*400 .8383*400 A1C13 .OOOOOE400 .573S8E-02 .1293*400 .1293*400 MCI .1B70*-01 .4&29*-03 .m % iE -o i .I0441E-01 AICI2 .0000*400 .437I7E-03 .98608E-O2 .98&0*-02 At .40000E-03 .40000E-03 .90223E-02 .90223E-02 AlCI .0000*400 .I0972E-O3 .2474BE-02 .2474BE-02 Cr€12 .0000*400 .18B69E-04 .A2562E-03 .42S62E-03 A12C16 .0000*400 .13222E-05 .29B26E-06 .29824E-04 H .0000*400 .27I79E-06 .61305C-05 .6130SE-OS Cr .0000*400 .33877E-OB .7W56E-07 .76456E-07 Cl .0000*400 .27687E-08 .62456E-07 .62456E-07 CrC13 .0000*400 .2690K-O8 .60&96E-07 .60&94E-07 M2 .0000*400 .IB825E-08 .«W2E-07 .42462E-07 AlH .0000*400 .12853E-08 .2899*-07 .2899*-07 A1 .0000*400 .82293E-09 .18562E-07 .I8562E-07 «K3 .24933E-01 .56212E-09 .12679E-07 .12&79E-07 C12 .0000*400 .15026E42 .33893E-1I .33893E-11 MH2 .0000*400 .84567E-H .19075E-12 .19Û75E-12 CrCU .0000*400 .7?35*-U .17467842 .17447E-12 CrN .0 0 0 0 * 00 .56032E-18 .I2639E-16 .12639E-1& M .0000*400 .70716£«I9 .1595* 1? .1595*47 AlH .0000*400 .2802*-21 .63202E-20 .63202E-20 K2H2 .0000*400 .30767E-22 .69352E-21 .69332E-21 K2H4 .0000*400 .W4Î5E-24 .I3445E-22 .I3645E-22

INVARIANT COKKNSED PHASES SPECIE IMIT.EST. EQ.nOlES Cr .2I600E«00 .2159*400 A1 .UOOOE-01 .21015E-01 AIN .0000*400 .I8700E-OI CrN .0000*400 .0000*400 Cr2N .0000*400 .0000*400 CrCI2 .0000*400 .0000*400 109

20 wt.% A1

-2 I0) AlCl

I -4 a>tn Oi

Cr (3 •o -0 ObO

-10 -2.4 - 2.0 - 1.6 -1.4 Log A1 Activity in Or—A1 Binary Masteralloy

Figure 35: Equilibrium partial pressures of the gaseous species in the bulk pack of an NH.Cl-activated pack as a function of aluminum activity in the Cr-Al masteralloy at 1150 C. 110

Me(s) + X HCl(g) = MeCl^(v) + x/2 Hg(g) (4.4) where Me is chromium and aluminum of a reduced activity. But the highest containing species in an NH^Cl-activated pack is Hg not Cl^ or Ar. Because of the high thermodynamic stability, the aluminum chloride species were the most abundant metal halide species generated from the pack mixture (Table 11 and Fig. 35), even at the lowest aluminum activity. The partial pressures of the chromium and aluminum chloride species were within an order of magnitude for certain masteralloy compositions. In fact, experimentation should concentrate on compositions between 95-90 wt.% chromium and 5-10 wt.% aluminum where the vapor pressures are comparable and codeposition becomes possible. From Table 11, both chromium and aluminum (e.g., A1 in the Cr-5A1 wt.% masteralloy) are partially or completely consumed by reaction (4.4) to form a number of vapor species and one condensed phase, AIN. Aluminum nitride and other nitride phases can form according to one of the following reactions: Me(s) + y NH^(g) = MeNy(s) + xy/2 Hg(g) (4.5) or

MeCl^(v)+y NH^(g)=MeNy(s)+zy/2 Hg(g)+x/2 CI 2 (g) (4.6) where the and Cl^ species in Eg. (4.6) react further to form the more thermodynamically favorable HCl gas species. Ill

The formation of a nitride phase may affect the overall composition of a diffusion coating because it reduces the partial pressures of the AlCl^ and/or CrCl^ species. However, no condensed nitride phase was detected in any NH^Cl-activated pack. 4.3 ITSOL CALCULATIONS WITH A RE SOURCE 4.3.1 RE Oxide Source Simultaneous deposition of chromium and aluminum and a small amount of a RE from a chloride-activated pack was also attempted. Two methods were considered to produce this further doping. A small amount of a RE oxide source (e.g., 2 wt.% ZrOg, Y^O^, and SiO^) was added to the pack mixture, replacing a small amount of the inert alumina filler. In this case, the RE oxide source was converted in the high chlorine activity of the pack to produce additional RE chloride species and a more thermodynamically stable, condensed oxide according to Eg. (4.7) ;

X REOy(s) + y MeCl^(v) = X RECl^(v) + y MeO^(s) (4.7) where RE is zirconium, yttrium, or silicon and Me is aluminum or chromium. The optimum amount of the RE oxide was determined experimentally because equilibrium calculations were independent of the RE oxide content. 1 1 2

Table 12: ITSOL pack equilibrium output for a 2 wt.% NH.Cl-activated pack containing 25 wt.% Cr-lOAl wt.% masteralloy and 2 wt.% at 1150 C.

H K t çasçs SKCIE INIÎ.KT. EO.tfUS P/ATH «T1V1TY Cr,Al,Y deposition: 1433:: •H Cl. Y303 , 9, \2 .0000**00 .3S»3E-01 .649236*00 .849236*00 STEP 10: 1 \ tl3 .0000**00 .6465K-02 .141366*00 .141366*00 At .W 0 * -0 3 .W00**03 .07455-02 .074526-03 T = 1423.00 K HCl .W 3*H )1 .22964E-M .5Û205E-O3 .50202-03 P = 1.00**00 ATM AlCl .0000**00 .5403K-OQ .11B24E-03 .110346-03 A1CI2 .0000**00 .1*95E-0Q .2250Œ-O4 .2250*-O4 Q Û ttI HOLES Aicia .0000**00 .64533E'06 .14109E-O4 .141096-04 .00>i H .0000**00 .2B222E-06 .61701E-O5 .617016-05 N .0194 CrClB .0000**00 .W36E-07 .9715*-06 .97l5*-06 H .0777 N2 .0000**00 .55323EH* .120956-06 .12092-06 Cl .0194 Cr .0000**00 .3W71E-06 .764566-07 .76452-07 Y .0*9 H20 .0000**00 .Z20%E-06 .48305E-O7 .483*6-07 0 .0133 AlH .0000**00 .im E -0 0 .29177E-07 .291776-07 AI .(WO MO .2S90X-O1 .99786E-09 .21G1K-0? .218166-07 Cr .2160 A1 .0000**00 Æ4901E-O9 .1B562E-07 .105626-07 Cl .0000**00 .136WE-09 .29B39EH» .298396-08 A120 .0000**00 .80I6K-10 .17527E-08 .173276-08 A1H2 oooo**oo .1829*40 .39987E-09 .399076-09 YC12 .0000**00 .19793E-10 .349666-09 .349666-09 AICIO .00(0**00 .4009*41 .876406-10 .8764K-10 CrCia .0000**00 .302BE-12 .661896-11 .661896-11 Aim .0000**00 .56350E-13 .123226-11 .123226-11 A1X16 .0000**00 .16223E43 .354606-12 .3546Œ-12 MO .0000**00 .1491*43 .3261*-12 .3261*-12 C12 .0000**00 .3S386E-1S .773646-14 .773646-14 AID .0000**00 .30579Ê-15 .6685*-14 .66856644 CrD .0000**00 .S3131M6 .11616644 .116166-14 A1HQS .0000**00 .22069E46 .4S49E-15 .482496-15 I im iM T CMENSED PIKES A1202 .0000**00 .10961E4& .23964645 .239646-15 SPECIE IN1T.E5T. EO.lttfS YO 0000**00 .77S9*-17 .169656-15 .169656-15 Cr .2160**00 .3160**00 Y .0000**00 .5W13E-17 .124436-15 .124436-15 A1 .4600*-01 .200976-01 .1943*-01 Ml 0000**00 .2452*47 .S360S-16 .536*6-16 AIN .0000**00 AlOH .0000**00 .9215549 .201476-16 .20147646 A12TD .0000**00 .323296-(e « .0000**00 .1422BE-1B .311N6-17 .311066-17 Y2Q3 .4 4 3 0 * -* .117716-02 .0(00**00 N .0000**00 .1Z313E-19 j6 9 2 * -l? .2692*-17 cm .0000**00 CrCU .0000**00 .4157BE-19 .909016-10 .909016-18 .0000**00 .0000**00 YCl .0000**00 .4127*49 .9023* 18 .9*306-18 YC13 .0000**00 .0000**00 .0000**00 .0000**00 0 0000**00 .22&13E-20 .4944(6-19 .4944*49 Crdi m .0000**00 .245B4E-21 .S374Œ-20 .5374**20 YOCl .0000**00 .00(0**00 ICH4 .0000**00 .77728E-22 .169946-20 .169946-20 Y .0000**00 .0000**00 Y3E .0000**00 .254135-3 .5556*-22 .5556**22 CrCia .0000**00 .0000**00 K3C .0000**00 .2 1 4 8 * -a .469616-22 .469616-22 Cr2D3 .0000**00 .0000**00 Crffi .0000**00 .14495-23 .3169Œ-22 .3169Œ-22 CIO .0000**00 .1360QE-24 .2 7 7 5 * -a .2 9 7 5 * -a AIGS .0000**00 .20O7*-26 .430926-23 .43893-25 Y20 .oooc**oo .18994E-2& .415266-25 .415266-25 1 1 3

Table 13: ITSOL pack equilibrium output for a 2 wt.% NH.Cl-activated pack containing _25 wt.% Cr-10 wt.% masteralloy and 2 wt.% ZrO^ at IISO'^C.

WfiE gases SrfClE IHIT.EST. EB.roES P/ATH ACTIVITY ftl,Cr,Zr déposition i\ I4Z3K, 90Cr-lM l, 2% Zrffi K .000006*00 .386236-01 .6421X6*00 .842066*00 SIEPM3.' 1 A1C13 .000006*00 .534746-02 .116586*00 .116586*(0 ZrC14 .0000(6*00 .481276-03 .104936-01 .104936-01 T = 1423.00 K XI .1943(6-01 .463596-03 .101076-01 .101076-01 P = 1.000£K)0 fttfl A1C12 .000006*00 .421916-03 .919966-02 .919966-02 Ar .4000(6-03 .400006-03 .872(CE-02 .872086-02 ELOEMT AlCl .000006*00 .109646-03 .239046-02 .239046-02 Ar .0004 CrC12 .0000(6*00 .182136-04 .397086-03 .397086-03 N .0194 A12C16 .000006*00 .111066-05 .242186-04 .242186-04 H .(rm ZrC13 .0000(6*00 .810066-06 .176616-04 .17661E4A Cl .0194 K .000006*00 .281816-06 .614406-5 .614406-% A1 .0460 Cr .0000(6*00 .350696-00 .764566-07 .7645(6-07 Cr .2160 Cl .000006*00 .276706-00 .600266-07 .603266-07 Zr .0061 CrC13 .0000(6*00 .2soei£-(e .546936-07 .546936-07 0 .0162 m .000006*00 .219696-08 ,478976-07 .470776-07 e .0000(6*00 .194766-00 .424626-07 .424626-07 AlH .000006*00 .13321£-(8 .290546477 .290646-07 A1 .0000(6*00 .851336-09 .185626-07 .105626-07 m3 .2S903E411 .5BS4(C-09 .127636-07 .12763607 ZrC12 .0000(6*00 .313166-09 .682756-00 .(C7SE<8 moci .0000(6*00 .812766-10 .1772(6-% .1772«-(e A120 .0000(6*00 .803926-10 .175276-08 ,17527E-(e A1H2 .000006*00 .1B1K6-10 .396496-09 .396496-09 0 2 .0000(6*00 .145046-12 .316216-11 .316216-11 A1K3 .000006*00 .50(26-13 .1216(6-11 .1216(6-11 W .000006*00 .878776-14 .191596-12 .191576-12 CrC14 .000006*00 .696536-14 .151866-12 .151666-12 ZrCl .000006*00 .302786-14 .660136-13 .660136-13 AID .0000(6*00 .306656-15 .6685(6-14 .668566-14 CrO .0000(6*00 .532796-16 .116166-14 .116166-14 A1W2 .0000(6*00 ,220376-16 .480456-15 .480456-15 A1202 .000006*00 .109926-16 .239646-15 .239646-15 ZrO .000006*00 217856-17 .4749(6-16 .4749(6-16 irMAiAKT caceeED m ss m .0000(6*00 .144566-17 .315176-16 .315176-16 s ra iE INIT.EST. 60.1065 AlOH .000006*00 .920186-18 .200626-16 .200626-16 Cr .216006*00 .215966*00 cm .0000(6*00 .579706-18 .126396-16 .126396-16 A1 .460006-01 .175006-01 NO .000006*00 .945366-19 ' " .104316-17 .104316-1? AlH .0000(6*00 .117926-01 M .000006*00 .731616-19 .I595(f-17 .1595(6-17 ZrK .0000(6*00 .763796-02 m .000006*00 .4963(E-eo .1062%-18 .1(02(E-18 A12Q3 .0000(6*00 .541336-02 0 .0000(6*00 226776-00 .4944(6-19 .494406-19 Zr82 .812006-02 .000006*00 Cr2N Zr .000006*00 .63076-21 .130096-19 .138096-19 .000006*00 .0000(6*00 ZrlE .0000(6*00 .377056-21 .822056-20 .822066-20 cm .000006*00 .0000(6*00 Cr203 AIN .0000(6*00 Æ9896-21 .632(66-20 .632026-20 .0000(6*00 ,000006*00 «H4 .0000(6*00 .269046-22 .506566-21 .58(566-21 m4C104 .000006*00 .0000(6*00 CIO .0000(6*00 .275876-23 .601466-22 .601466-22 W 12 .OOC((6*00 .0000(6*00 CrŒ .000006*00 .14396-23 .31M6-22 .316906-22 CrC12 .000006*00 .000006*00 M2H2 .000006*00 .749826-24 .163486-22 .163486-22 mo .0000(6*00 .243516-24 .5309(6-23 .5309(t-23 A1Q2 .000006*00 .201326-26 .438926-25 .438926-25 02 .000006*00 .611626-28 .133356-26 .133356-26 1 1 4

Table 14: ITSOL pack equilibrium output for a 2 wt.% NH^Cl-activated pack containing 25 wt.% Cr-lOAl wt.% masteralloy and 2 wt.% SiO,> 2 at 1150 C.

MASE gases SffCIE INIT.EST. EO.rOES P7ATM M3IVITY Al.Cr.Sl dtposition a t 1«3K, *4C1, 9(Cr-10Al, a SIDS K2 .OOOOOEKO .362424)1 .612I5E*00 .8121554)0 STEP IB: 1 A1C13 .OOOOOEKO .5(^83E*02 .114245*00 .114245+00 Æ1 .107OOE-O1 .229^*412 .5131*4)1 .5)31*41 I = ira.M K M I2 .OOOCKE+OO .47357E-03 .1C611E4H .10611541 P = 1,0«€*00 ADI Ar .40000E-03 .W00*-03 .8962954% .89629542 A1C12 .0OOOCE40O .77S23E-04 .1743**02 .1743*42 BflOT lO fS M2 .OOO00E4OO .37666E-04 .6440*413 .8440*40 Ar .OCW AlCl DOOOOEXO .39117E-05 .6765154)4 .07651544 Cr .2160 A12C16 .OOOOOE+00 .1037*4)5 .232545*04 .232545*04 A1 .0469 CrC13 .OOOOCE+OO .3372*4% .7555*46 .755*45 Si .0166 SiClS .0000* 00 .3163*4)6 .7W75E-05 .70875546 0 .0333 H .00(0**00 .2692*4% .6C339E46 .60339545 Cl .oin S1C13 .0000**00 .7650*4)7 .171425*05 .1714*45 N .0187 HO .2492*4)1 .76062E4J7 .17044E46 .17044546 H .0748 SÎC14 .0000**00 .737;3-07 .16524E46 .165245-05 K20 .0000**00 .3B46E4? .125145*05 .12514546 EOlILIBRim OIPOSITIGKS: Cl .0000**00 .1391724)7 .311855*06 .311855-06 SiO .0000**00 .391B1E48 .877945-07 .07794547 Cr .0000**00 .34I21E48 .7645654)7 .76456547 AICID .00(0**00 .7B546E-I0 .1760*40 .1760*40 Sim .0000**00 .20081E-10 .4499554)9 .449955-09 S i« .0OCK**OO .13795E-10 .309115419 .30911549 AlH .0000**00 .90321E-11 .2023*4)9 .2023*49 A1 .0000**00 .SB757E-11 .1316**09 .131665*09 CrCU .0000**00 .48399E-11 .1064554)9 .100455*09 SiCl .ococ**oo .«e05E“ll .907615-10 .907615-10 C12 .0000**00 .37712EMI .845025-10 .8450**10 HE .0000**00 .1I62*-1I .2605**10 .2605**10 SIH .0000**00 .19903E-12 .445965-11 .445965*11 SÎK3 .0000**00 .13835E-12 .3100**11 .3100**11 DfMTIMT OKeSED m s s Aue .0000**00 .12105E-12 .271245-11 .271245*11 SPECIE IMIT.EST. EQ.fOfS A12D .0000**00 .1066*42 .238865-11 .238865-11 Cr .2160**00 .187795*00 Si .0000**00 .45939E-13 .102945-11 .102945-11 AlH .0000**00 .1812X41 SiN .0000**00 .25764E-U .577315-13 .377315-13 A1203 .0000**00 .I1097E41 CrO .0000**00 .1W42E-U .314655-13 .314655-13 CrSSia .0000**00 .55465542 A1H3 .0000**(0 .36477E-15 .817355-14 .817355*14 SHE .16645541 .0000**00 n .0000**00 .31M2E4S .703865-14 .7038**14 CrM .0000**00 .0000**00 Hi .0000**00 .195WE45 .437935-14 .4379**14 Cr203 .0000**00 .0000**00 AIHE .0000**% .1096**15 .2455*-14 .2455**14 GrSi .0000**00 .0000**00 GrM .0000**00 .7952**16 .1781**14 .1781**14 CrSi2 .0000**00 .0000**00 AIQ .000(**00 .57327£*16 .128455-14 .123455-14 Si3N4 .000005*00 .0000**00 Si02 .0000**00 .2707**16 .6066**15 .6066**15 A12S105 .0000*40 .0000**00 M .0000**00 .10036E-16 .2248**15 .2248**15 Cr2N .0000**00 .0000**00 S i» .0000**00 .25239E-17 .5655**16 .StSX-16 A1 .4600*41 .0000**00 HOCl .0000**00 .66409E-18 .1488**16 .1488**16 «E .0000**00 .5959**18 .133545-16 .133545-16 A1202 .0000**00 .39481E-18 .884675*17 .894675*17 AlOH .0000**00 .1689**18 .378555-17 .378355*17 0 .0000**00 .57767E-19 .133925-17 .13392*17 S12 .0000**00 .2505**19 .5613**18 .5613**18 1 1 5

20 wt.% Al15 ŸüT

-2 A1CI2 0) 0M M CrCl2 Q) Ih -8 PU Cr

- 0

-10

- 1 2 - 2.6 -2.4 -1.4

0 15 20 wt.% Al

-2

ZrCl4

- 4 AlCl

- 6

ZrCl2

- 6 -2.4 - 2 .0 - 1.6 -1.4 Log Al Activity in Cr—Al Binary Masteralloy

Figure 36: Equilibrium partial pressures of the gaseous species in the bulk pack of an NH.Cl-activated pack as a function of aluminum activity in the Cr-Al masteralloy at 1150 C with 2 wt.% additions of (a) Y_0 , (b) ZrO , and (c) SiO^. 1 1 6

Figure 36: (continued)

0. 0 0 15 20 wt.% Al

B cd - 2.00 - - 0) AlCl

-4.00 - - iQ) 04U — 0.00 - - SiCl2 5 Cr PU W SiCU

- 10.00 -2.06 - 1.86

Log Al Activity in Cr-Al Binary Masteralloy 1 1 7

Examples of an ITSOL calculation for an NH^Cl-activated pack with additions of 2 wt.% ZrO^, ^2*^3' and SiOg at 1150°C (1423K) are presented in Tables 12-14, respectively. In addition, ITSOL calculations as a function of aluminum activity are plotted in Fig. 36. Two distinct types of results were observed depending on the thermodynamic stabilities of the RE oxide compared to the alumina filler. RE oxide additions which are more stable than alumina, such as yttria, are only partially converted to their vapor chloride species. In fact, as the aluminum activity increased (Fig. 36), the vapor pressures of the YCl^ species also increased, illustrating the Le Chatelier-Braun principle on Eg. (4.7). On the other hand, zirconia and silica additions, which are less stable than alumina, are completely converted to their vapor chloride species. The vapor pressures of the ZrCl^ and the SiCl^ species decreased with increasing aluminum activity, resulting from the increasing thermodynamic stability of the AlCl^ species. ITSOL calculations predicted the formation of several condensed phases from the various pack reactions. Calculations confirmed the conversion reaction of the type in Eg. (4.7). For zirconia, yttria, and silica additions, a condensed alumina product phase and a reduction in the vapor pressure of the AlCl^ species 1 1 8 resulted, whereas no condensed chromia phase or significant reduction in the vapor pressure of the CrCl^ species was predicited by ITSOL calculations. Therefore, ITSOL calculations indicated that the AlCl^ species are predominant in the conversion of the RE oxide additions (Eg. (4.7)). For silica additions at low aluminum activities, ITSOL calculations predicted the formation of a AlgSiOg condensed phase from the conversion of SiO^. Condensed nitride phases such as AIN— common for all three RE oxide additions, ZrN, and Si^N^ were again predicted for packs with yttria, zirconia, and silica, respectively, according to Eqs. (4.5) or (4.6). Finally for silica additions, ITSOL calculations predicted that SiCl^ conversion products would react by coating the chromium-rich masteralloy particles to form a Cr^Si^ condensed phase at all aluminum activities. Unfortunately, Cr^Si^ particles will have a lower chromium activity and could hinder the overall ability to enrich an aluminide coating layer. 4.3.2 RE Activator Salts By a second method to introduce the RE, a chloride activator salt (e.g., ZrCl^ and YCl^) was introduced into the pack. For example, ZrCl^ can react with the Cr-Al binary masteralloy to produce both aluminum and chromium 119

20 wL.% Al ZrCl4 I UQ) P -2 CO CO L,Q> CU 2 A1C1 2 -•I AlCl ZrClg

- 6

- 2.8 - 2.0 - 1.8 - 1.8

10 15 20 wt.% A1 YCI3

-2 - o

AlCl -6 Cr AICI2

-8

-10 — 2.6 -2.4 -2.2 -2.0 -1.8 -1.8 Log A1 Activity in Cr-Al Binary Masteralloy

Figure 37: Equilibrium partial pressures of the gaseous species in the bulk pack of an (a) ZrCl.- and (b) YC1_-activated pack as a function of aluminum activity in the Cr-Al masteralloy at 1150 C. 1 2 0 chlorides as well as significant partial pressures of the zirconium chloride species according to the following reaction; Me(s) + ZrCl^(v) = MeCl^(v) + ZrCl^_^(v) (4.8) As long as the substrate does not contain a relatively large amount of the RE, the necessary driving force exists for some codeposition. ITSOL calculations as a function of aluminum activity for ZrCl^- and YCl^-activated packs were plotted in Fig. 37. The most dominant vapor species in these packs were the ZrCl^ and YCl^ vapor phases and not the AlCl^ species. But, certain masteralloy compositions existed where comparable pressures of the chromium and aluminum chloride species were produced, and codeposition was very promising. Two problems existed for pack equilibria calculations containing a YCl^ activator or a Y^O^ addition. Since yttria was more stable than alumina and the inert filler was neglected in the equilibrium calculations, an additional calculation was required to predict the true equilibrium composition of the YgO^/YCl^/AlgO^/AlCl^ system. Unfortunately, the total aluminum content in the pack comprised not only alumina at unit activity, but also aluminum in the Cr-Al binary masteralloy. No distinction could be made between the two 1 2 1

independent phases during a single ITSOL calculation. Therefore, two separate ITSOL calculations were performed: one which equilibrates the activator, masteralloy, and RE oxide (if any) and one which equilibrated those equilibrium vapors with the "inert" alumina filler powder. Figures 36a and 37b reflect these two-step calculations. In addition, the thermodynamic data for the YCl^ and the YO^ gaseous species were incomplete. The standard entropy, As°, for some species was estimated from data of other family members for the same species. For example, a S° value was chosen for the YCl^Cv) which fell between the values for the LaCl^ and ScCl^ vapor species, which were known. The uncertainty in these estimations cause the equilibrium calculation to be somewhat questionable. 4.3.3 Masteralloy Combination For SiOg additions, the vapor pressures of SiCl^ species dropped as the aluminum activity increased so that codeposition of aluminum with either silicon or chromium was not very promising. Therefore, alternate means of codepositing aluminum, silicon, and chromium was necessary. One method is to introduce to the pack two binary masteralloys, a Cr-Al and a Cr-Si masteralloy in which the activities of aluminum and silicon can be varied independently. ITSOL calculations were also conducted for such a masteralloy combination with an NH^Cl activator 1 2 2 salt at 1150°C to determine potential masteralloy compositions. Again, a standard 50 gram pack mixture containing 2 wt.% NH^Cl, 20 wt.% Cr-Al binary masteralloy, and 5 wt.% Cr-Si masteralloy was used in all calculations. The alumina filler was found from ITSOL calculations to be inert and was neglected. The molar amounts of each element and the standard Gibbs energies of formation of all gaseous and condensed phases were input to the ITSOL program. Equilibrium calculations were conducted as a function of silicon in the Cr-Si binary masteralloy at a fixed activities of the aluminum and chromium. Unfortunately, each masteralloy was a source for chromium at a unique activity independent of one another. But since the amount of the Cr-Si masteralloy is small compared to the Cr-Al masteralloy for the compositions considered since the activity of chromium in the Cr-Al masteralloy is greater than that in Cr-Si masteralloy, the contribution of the chromium from the Cr-Si masteralloy was neglected. In addition, coating of masteralloy particles to form a condensed silicide compound was neglected. Then, the calculations could be made. Figure 38 is a plot of the equilibrium partial pressure of the gaseous species as a function of silicon activity in the Cr-Si masteralloy (Fig. 17b) at three Cr-Al masteralloy 1 2 3

0.0 0 .% Si

Ba - 2 . 0 0 - - o Q) AlCl 3 CO CrCIp COQ) a.U -<.00 - -

c3 Pu 00 SiCl; SiCl4

- 6.00 -3.60 - 2.60 -0.60

0.00 20 ,30 ^0 wt.% Si

-ZOO" AlCl CrCl2

SiClg

S1C1 4

- 0.00 - 2.60 - 1.60 -0.60 Log Si Activity in Cr-Si Binary Masteralloy

Figure 38: Equilibrium partial pressures of the gaseous species in the bulk pack of an NH.Cl-activated pack as a function of silicon activity in the Cr-Si masteralloy at 1150 C containing 20 wt.% of a (a) Cr-lOAl wt.%, (b) Cr-15A1 wt.%, or (c) Cr-20A1 wt.% masteralloy. 1 2 4

Figure 38: (continued)

4 0 0.00 20 ,30 wt.% Si

—2.00 - -O' MCI

— 0.00 - - SiClg

SiCl; SiCl4

Log Si Activity in Or—Si Binary Masteralloy 1 2 5 compositions with known aluminum activities. As both the aluminum activity and the silicon activity increased,the vapor pressures of the AlCl^ and SiCl^ species also increased. Comparable vapor pressures were produced between 85-80 wt.% chromium and 15-20 wt.% aluminum and between70-60 wt.% chromium and 30-40 wt.% silicon masteralloy compositions, respectively, where codeposition might be possible. 4.4 HASS SPECTROMETER MEASUREMENTS Cementation packs generate highly volatile metal-chloride species which contribute to the overall composition of the coating. Kinetics (gaseous and solid-state diffusion) and substrate effects can also help determine the composition and structure of the coating. But ITSOL equilibrium calculations represent one of the best ways for understanding the process and predicting the composition of the coating. These equilibrium predictions require experimental verification because of the overall complexity of the process. Vapors generated by metal-chlorine reactions, such as the pack cementation process, can be studied with the aid of an atmospheric pressure sampling mass spectrometer (APSMS) at the NASA Lewis Research Center. This instrument can directly sample, identifying vapor species and measuring their relative intensity, high pressure 1 2 6 vapors which evolve from the pack while maintaining the chemical and dynamic integrity of the process. Pack powder mixtures, as outlined in Section 3.3, were analyzed using the APSMS to verify the conversion of the RE oxide source to produce volatile RE chlorides via Eg. (4.7). The vapors from the pack enter the sampling orifice (Fig. 31a) and undergo a free jet expansion, which involves an abrupt transition to collisionless flow at speeds of several mach numbers. The vapors proceed as a molecular beam through a series of differentially pumped vacuum chambers to a quadrupole mass spectrometer, which operates at about 10 ® torr. The vapor species are ionized by an electron beam and collected by the mass filter. The output of the mass spectrometer consists of a voltage proportional to the intensity of each ionized species as a function of its mass-to-charge ratio. Peaks are identified both according to their mass-to-charge ratios and their characteristic "isotopic fingerprints". Because of the existence of ^^Cl and ^^Cl, volatile chlorides have distinctive "isotopic fingerprints", which can be identified by comparison to calculation.[142] The ion intensities. I, were calculated from the following expression:[143] 1 2 7

I = V/(«7Ry) (4.9) where V is the sum of the voltage signals for each isotope of the particular ion, a the ionization cross-section determined from the additivity rules[144,145], R the resistance across the multiplier for the particular measurement, and X the gain of the multiplier. The measured ion intensities were normalized to the most intense metal chloride peak. These intensities cannot be related quantitatively to a vapor pressure because of the different free jet expansion characteristics of the vapor species and the fragmentation of the parent molecule by the ionizing electrons. Nonetheless, the output of the mass spectrometer indicates the identities of the vapor species and the relative amounts. The actual and normalized intensities of the dominant vapor species for each pack powder mixture are listed in Tables 15-17 with their respective partial pressures calculated by ITSOL. Because of both sample depletion and orifice clogging, the measured ion intensities were not constant, but rather decayed with time. For this reason, ion intensities of the principal species were measured in the first few minutes of the reaction. 128

Table 15: NH.Cl-activated pack with a Cr-lOAl wt.% masteralloy at 927 C (1200K).

(a) Relative partial pressures calculated by ITSOL

Species Calculations

?!' P./P1 max atm

AlCl, 1 ,0 7X 10 '* 1 AlClj 6 . 1 6X 1 0 ' ^ 3 . 4 8X 1 0 " * AlCl 1 . 1 5 X 1 0 ' ® 1 . 0 7 X 1 0 ' ® CrClj 6 . 5 7X 1 0 ' ^ 6 . 1 0 X 1 0 ' ^

(b) Relative intensities of the principal ions observed with the mass spectrometer

Ion Probable + parent li' K / ^ L molecule(s) arbitrary units

AlCl* AlCl, 1 09 1

AlCl* AlCl,, AlCl, 2 7 . 2 . 2 4 9

AlCl* AlCl,, AlCl,, AlCl 9 . 9 9 9 . 1 5 X 1 0 “^

CrCl* CrCl,, CrCl, 3 . 2 1 2 . 9 4 X 1 0 " ^ 129

Table 16: NH.Cl-activated pack with a Cr-lOAI wt. masteralloy and 2 wt.% ZrOg at 927 C (1200K) .

(a) Relative partial pressures calculated by ITSOL

Species Calculations

?!' Pi/P.ax atm

AlCl; 1 .8 7X 10 '^ 2 . 4 6 X 1 0 " ! AlCl; 1 . 7 2 X 1 0 ' “ 2 . 2 6 X 1 0 " ^ AlCl 5.16X10"’ 6.79X10"® CrCl; 3 .5 5X 10 '^ 3 . 3 6 X 1 0 " ^ ZrClj 7 . 5 9 X 1 0 ^ 1 ZrCl; 7 . 5 6 X 1 0 " ’ 9 . 9 6 X 1 0 " ®

(b) Relative intensities of the principal ions observed with the mass spectrometer

Ion Probable + parent li' h / C . molecule(s) arbitrary units

AlCl; AlCl; 4 4 . 3 1

AlCl* AlCl;, AlCl; not det. not det.

AlCl* AlCl;, AlCl;, AlCl not det. not det.

CrCl; CrCl;, CrCl; 3 8 . 3 .86

ZrCl; ZrClj 8 . 0 3 .18

ZrCl* ZrClj, ZrCl; 1 . 3 8 3 . 1 2 X 1 0 " = 130

Table 17: NH.Cl-activated pack with a Cr-lOAl wt.' masteralloy and 2 wt.% YgOg at 927 C (1200K).

(a) Relative partial pressures calculated by ITSOL

Species Calculations

Pi' P./P1 max atm

AlClj 3 . 1 4 X 1 0 “ ' 1 . 9 1 X 1 0 " : AlCl^ 2 .2 9X 10 "^ 1 . 3 9 X 1 0 " ] AlCl 4 . 3 5 X 1 0 " ‘ 2 . 6 4 X 1 0 " ] CrCl^ 1 . 3 1 X 1 0 ^ 7 . 9 6 X 1 0 " ] ÏClj 1.65X 10 "^ 1

(b) Relative intensities of the principal ions observed with the mass spectrometer

Ion Probable + parent li' molecule(s) arbitrary units

AlCl* AlCl] 35.5 1

AlCl* AlCl], AlCl; not det. not det.

AlCl* AlCl], AlCl], AlCl not det. not det.

CrCl* CrCl], CrCl; 17.4 .489

YCl* YCl] 2.98 8.39X10"] 131

The AlClg^ species (from the AlCl^ parent molecule) was the most dominant species for each pack powder mixture. The species AlClg^, AlCl^, and CrClg^ were identified during the heating period, indicating that both AlCl^ and CrCly vapor species were formed (Table 15). However, the ion intensities for the metal halides (e.g., AlClg^, AlCl*) do not result entirely from simple ionization; some fraction of these measured intensities must form from fragmentation of the highest metal halide (e.g., AlClg). Furthermore, as mentioned, the intensities in Tables 15-17 cannot be considered to be equivalent to partial pressures. Hence these Tables show general experimental trends. The ZrCl^*, ZrCl^*, and YCl^* species were detected for the pack mixtures containing ZrOg and YgO^, respectively (Tables 16 and 17). The mass spectrometer data show qualitatively that most of the predicted halide vapor species are produced. However, the open crucible used in the mass spectrometer studies precluded the attainment of equilibrium and, therefore, the intensities are not expected to agree quantitatively with ITSOL predictions. The actual pack cementation process is carried out in a closed retort which retains the active chloride species, so equilibrium is more likely to be attained. However, some chloride removal from vented packs has been noted experimentally 132

Table 18: Chemical analyses of the condensate from an NH.Cl-activated pack with a Cr-lOAl wt.% masteralloy and 2 wt.% Y O held at 927 C (12 00K) for 1 hour.

Al^ Cr^ Cl^ (wt.%) (wt.%) (wt.%) (wt.%) 374 075 276 >5.0 a=Atomic emission spectroscopy b=X-ray fluorescence spectroscopy 133

and this may cause difficulty in attaining the desired coating compositions. The mass spectrometer analyses in Tables 16 and 17 also established that the intensities of the CrClg^ species increased while the AlCl^^ species decreased in pack mixtures containing the RE oxide source, indicating that a reaction of the AlCl^ species with the RE oxide must have produced the RE halide species via Eg. (4.7). X-ray diffraction (XRD) analysis of the spent powder

identified the a- and X-AlgOg phases, part which may have formed from the conversion reaction (Eg. (4.7)). No condensed CrgOg phase was identified. Some condensation experiments were also done to supplement the mass spectrometer analyses. The chemical analyses of the condensed vapors are listed in Table 18. No compound or phase identification was performed, but chemical analyses of condensates from the YgOg-containing powder mixture indicated that an yttrium-, aluminum-, and in a lesser amount, chromium-chloride were deposited from the vapor phase onto the copper target. Both experiments indicate that RE oxides in the pack suffer conversion reactions in the high-chlorine packs to form their volatile halide species (e.g., ZrCl^^, ZrCl^^, and YCl^^), which may deposit a significant amount of RE into a diffusion coating. CHAPTER V RESULTS AND DISCUSSION

5.1 COATING DEVELOPMENT FOR NICKEL-BASE ALLOYS 5.1.1 Chromium/RE(Zr,Y)-Modified Aluminide Coatings In developing a Cr/RE-modified aluminide coating, experimentation is needed to determine the optimum pack chemistries, such as the proper activator and the masteralloy composition. According to ITSOL calculations, the most favorable Cr-Al masteralloys were between 95-90 wt.% chromium and 5-10 wt.% aluminum with all activators. But, RE oxide additions reduced the vapor pressures of the AlCl^ species so that masteralloys with higher aluminum activities or compositions were most promising. Two substrate/pack arrangements were used to produce the new, modified coatings: a traditional powder contacting pack and an "above pack" arrangement. Powder contacting arrangement: The powder contacting arrangement is shown in Fig. 27. A series of potential pack chemistries were used to coat René 80, IN 713LC, and Mar-M247 nickel-base superalloys at 1150°C. A summary of all coating results are presented in Figs. 39 and 40 and Tables 19-22 including representative cross-sectional

134 135 micrographs, average bulk surface compositions (EPMA), and XRD results. WDS analysis was also used to distinguish the phases of the coating/substrate morphology, and these results are shown in Figs. 41 and 42. Two results of the coating development were observed: the formation of an external carbide or an external aluminide coating. For masteralloy compositions with low A1 activities, an external carbide layer was produced.This carbide layer was much thinner than the aluminide coatings and had a relatively large weight gain/area associated with its formation. According to Figs. 35-37, the highest chromium chloride vapor pressure was expected for the lowest aluminum masteralloy composition (Cr-5A1 wt.%); thus the largest flux of chromium to the substrate surface (Eg. (2.3)) would also be expected and a chromium-rich surface coating would be formed. Because of the residual carbon in commercial superalloys (Table 2) , free carbon from the bulk diffuses rapidly to the substrate surface reacting with chromium from the gas phase to form an outward-grown metal carbide and preventing aluminum deposition to form an external NiAl coating layer. Microprobe analyses also detected other refractory alloying elements (e.g., Mo, W, Ta, Hf, etc.) in the carbide. This external carbide acts as a diffusion barrier for other elements, inhibiting the 136

ïiljMV

René 80 Substrate ^ lOOpm ^

^ lOOpm ^

Figure 39: Cross-sectional optical micrograph of a René 80 alloy substrate treated at 1150 C for 24 hours in a 2 wt.% NH.Cl-activated pack containing (a) 2 wt.% ZrO„ and 25 wt.% Cr-lOAl masteralloy or (b) 25 wt.% Cr-7.5A1 masteralloy. Powder contacting arrangement. 137

“ 2 ° 3

y lO O im I—-— I R e n é 80

100 |i/n ■4 «• y + y ' SUBSTRATE R e n é 8 0

Figure 40: Cross-sectional optical micrograph of a René 80 alloy substrate treated for 24 hours in a (a) 2 wt.% ZrCl.- or (b) a YCl_-activated pack containing 25 wt.% Cr-7;5A1 masteralloy. Powder contacting arrangement. 138

Table 19: Summary of nickel-base alloys coated with a pack containing 2 wt.% NH.Cl Cl plus 2 wt.% 7—ZrOu at IISO^C for 24 hours. Powder contacting-ing arrangement. Ave. Surf. Comp. (at.%) XRD Am/A MA fwt.% ) Ni A1 Cr Zr Phase mg/cm—2 um René 80;

CrSAl 16.6 2 . 1 77.6 - 43.1 55 Cr7.5A1 11.9 3.4 79.2 54.2 75 ÎIÈ- •• * 71 45.7 40.1 9.9 Al2°3' 24.7 6 -NiAl II II CrlOAl 44.9 45.6 3.1 2 . 0 f f 29.7 136 Bro­ il ^ II " * 46.3 42.7 5.8 9 9 19.7 80 no ZrO- IN 713LC - Cr7.5A1 12.7 2 . 6 79.8 26.9 63 CrlOAl 48.7 46.7 3.2 1.4 47.1 159 G-NlAl Mar-M247 - Cr7.5A1 15.3 2 . 8 78.5 30.3 70

CrlOAl 43.6 47.6 2 . 1 2 . 1 28.6 1 2 1 fi-NiAl *=Pack without ZrOg additions , shown in Fig. 41b.

Table 20: Summary of nickel-base alloys coated with a pad containing 2 wt.'h ZrCl4 1150 C for 24 hours. Powdei contacting arrangement. Ave. Surf. Comp. (at.%) XRD Am/A MA fwt.%) Ni A1 Cr Zr Phase mg/cm--2 fim René 80:

Cr7.5A1 49.1 39.2 8.7 . 2 0 A 1 , 0 , 29.1 89 G-NiAl II II CrlOAl 47.1 43.2 5.7 . 1 1 9 16.7 82 IN 713LC II II Cr7.5A1 52.7 41.1 6 . 8 .17 9 30.5 105 II II CrlOAl 50.6 43.5 3.9 .09 9 25.9 85 Mar-M247 II II Cr7.5A1 44.5 41.6 8.7 .16 9 29.1 94 II II CrlOAl 43.2 44.7 5.7 .07 r ...... 24.2 72 *Cr5Al masteralloys formed coatings. 139

Table 21: Summary of nickel-base alloys coated with a pack containing 2 wt.% YC1_ at 1150 C for 24 hours. Powder contacting arrangement. Ave. Surf. Comp. (at.%) XRD Am/A At MA fwt. %) Ni A1 Cr Y Phase ma/cm— um René 80;

CrSAl IS.4 2 . 1 79.8 - 35.4 65

Cr7.5A1 44.4 43.1 8.5 . 0 1 27.1 96 5-NiAl II ti » * 46.3 41.6 9.6 . 0 2 t 36.3 105 IN 713LC: CrSAl 17.6 3.2 78.7 - 29.5 72 Cr7.5A1 47.8 4S.1 7.7 .03 26.9 95

6 -NiAl ;s II ■> * 48.7 44.9 9.1 . 0 2 / 35.4 113 Mar-M247 CrSAl 18.1 2.9 79.6 - 35.9 6 6 Cr7.SAl 46.0 43.S 7.3 . 0 1 24.6 87 B-NiAl II II " * 4S.4 41.5 8.5 . 0 2 34.2 129 *=Coating treatment at 36 hours.

Table 22: Summary of nickel-base alloys coated with a pack containing 1 wt .% CrCl. plus 2 wt .% of second activator and 2S wt.% Cr--7. SAl 0 t.% masteralloy at 1150"C for 24 hours. Powder contacting arrangement. Ave. Surf. Comp. (at.%) XRD Am/A At Activator Ni A1 Cr RE Phase mcf/cm— flTXl René 80:

YCI 3 4S.4 38.2 13.8 .OlY A1 0 , 37.8 1 1 1 fi--NiAlfa-Cr II It ZrCl 47.6 42. S 7.2 .09Zr f 23.1 8 6 IN 713LC: II II II YCl 48.9 40.6 1 0 . 2 .0 2 Y / / 35.9 1 1 1 II It ZrCi^ SO.3 43.1 5.8 .07Zr t 27.3 79 Mar-M247: II II YCl 4S.S 40.3 9.6 .OlY t 37.8 1 2 1 Zrci 47.8 41.7 6.5 .OBZr __It C_____ II 24.9 104 140 deposition and enrichment of aluminum and the RE into the substrate surface. Also, the carbide reduced the enrichment of chromium to the substrate, which in any case, was not the intention of the coating. The supply of chromium was then controlled by the dissociation of the external carbide at the carbide/substrate interface instead of the series diffusion in the gas phase and solid-state.[42] But most importantly, the formation of a Cr-enriched NiAl coating was prevented. Masteralloys with higher aluminum activities resulted in an external, aluminum-rich coating layer (e.g., B-NiAl). The aluminide coatings were much thicker and had a smaller weight gain/area from their larger growth rates and lower density, respectively, compared to the carbides. The aluminide coatings consisted of two phases: an outward-grown, hypostoichiometric B-NiAl matrix with a dispersion of a-Cr second-phase particles which precipitate near the original substrate surface and which both grow inward from the gas phase along grain boundaries at the coating surface (Figs. 39 and 40). According to the phase diagrams of Fig. 43, chromium has a limited solubility in the B-NiAl phase (~4 wt.% at H50°C) depending on temperature and the Ni/Al ratio. Upon supersaturation, a two-phase region is formed containing a chromium-rich, a-Cr phase in a B-NiAl matrix. 141

Figure 41: Cross-sectional electron micrograph of the B-NiAl and a-cr, two phase coating and the corresponding x-ray maps of nickel, aluminum, cobalt, and chromium; 9 0 0 0 X . 142

Figure 42: Cross-sectional electron micrograph of the interditfusion zone and the corresponding x-ray maps of nickel, aluminum, cobalt, chromium, tungsten, and molybdenum; 5000x. 143

AIN, 50

40

,30 AIN,

20

,10

Cr 20 30 40 SO 6 0 70 60 9 0 N,

NICKEL ATOMIC PERCENT

1900

1700 V s CJ

1500. g 2 E* 1300 H 1100

900 b Cr 10 20 30 At % 60 70 80 90 AlNi _L 0 10 20 30 40 50 60 70 80 90 100 Cr Wt.% AlNi

Figure 43: A summary of (a) the 1150 C isothermal section of the NiCrAl ternary system, and (b) the B-NiAl-Cr quasi-binary phase diagram.[146] 144

Microanalysis of the coating surface detected chromium, cobalt, titanium, and molybdenum (not listed in Tables 19-22) dissolved in the 6-NiAl coating phase due to outward diffusion, along with nickel, during the coating process. An interdiffusion zone between the outer, two-phase (fi-NiAl+a-Cr) coating and the nickel-base substrate was comprised of three phases. For René 80 alloy substrates, blocky Cr(W,Mo)-rich carbides were detected between the outer B-NiAl coating and the interdiffusion zone; these marked the original substrate surface (Figs. 39, 40 and 42). Beneath the carbides, Cr(Mo,W,Co)-rich fingers, probably sigma phase, were aligned perpendicular to the coating surface within a 6-NiAl matrix. Other alloy substrates had somewhat different interdiffusion zone morphologies. Since IN 713LC alloys have no cobalt and lower carbon contents, the interdiffusion zone lacked the sigma phase and had a less continuous Cr(Mo,Nb)-rich, MC-type carbides. Mar-M247 alloys contained stronger carbide formers, and their interdiffusion zone consisted of a random dispersion of Cr(W,Ta,Hf)-rich carbides. The formation of the interdiffusion zone generally is formed because of the depletion of nickel from the original substrate surface. In the absence of NiAl formation, carbon forms the undesirable external carbide, but the 145 formation of an external fi-NiAl compound prevents this reaction by reducing the instantaneous flux of chromium at the substrate surface. In this case, carbon reacts with the refractory elements in the nickel-depleted region, where the carbon solubility is lower, and forms MC and MgC-type carbides at the original substrate surface. Once free carbon is depleted, sigma phase then precipitates. According to Figs. 35-37, the vapor pressures of the CrCl^ species decrease while those for AlCl^ species increase with increasing A1 activity in the Cr-Al binary masteralloy. Thereby with increasing Al activity, the ratio of fluxes of chromium to aluminum to the substrate surface becomes more favorable for codeposition. In fact, an NH^Cl-activated pack with ZrOg additions required even further increases in aluminum activity to produce favorable fluxes of aluminum, chromium, and zirconium. Only masteralloy compositions of Cr-lOAl wt.% or greater Al contents were capable of forming an external aluminide coating. Aluminum chloride species are predominant in the conversion of ZrOg to their ZrCl^ species, thereby reducing the vapor pressures of the AlCl^ species and the resulting aluminum flux. Therefore, higher Al activities which produce higher vapor pressures for the AlCl^ species are required to compensate for both the conversion (Eq. 146

(4.7)) and deposition reactions. To illustrate this further, NH^Cl-activated packs without ZrOg additions formed an aluminide coating with a more dilute (Cr-7.5A1 wt.%) masteralloy composition (Table 19 and Fig. 39b). The surface composition was very similar to the ZrOg- containing packs with an Al-rich masteralloy (Cr-lOAl wt.%) composition, except for the chromium surface compositions which were lower than the more dilute masteralloy. Of course, the more dilute masteralloy compositions would generate higher CrCl^ vapor pressures, higher chromium fluxes, and higher chromium compositions. The average surface compositions were measured by the electron probe microanalyzer (EPMA) with a large beam diameter, between 10 and 30 The bulk measured composition is dependent on the relative amounts of a-cr second-phase particles in the 6-NiAl matrix and the composition of aluminum in the 6-NiAl phase, fixed according to the activity of aluminum in the masteralloy. Therefore, the volume fraction of a-cr second-phase particles resulting from the Cr activity in the Cr-Al binary masteralloy affects the measured bulk surface composition, typical of two-phase microstructures. Additional pack chemistries were tried. Activator combinations with 2 wt.% YCl^ or ZrCl^ and 1 wt.% 147

CrClg were used with 25 wt.% Cr-7.5A1 wt.% masteralloy at 1150°C for 24 hours to increase the average bulk surface composition of chromium in the 6-NiAl coating. A summary of the results are presented in Table 22. Two distinct trends were identified. Activator combinations with a (2:1) YClg/CrClg mixture effectively increased the surface composition of chromium, the thickness of the aluminide coating, and the overall weight gain/area. The larger chromium composition corresponded to an increase in the volume fraction of a-cr second-phase particles (Fig. 44), which grew inward from the gas phase and caused an increase in the weight gain/area values measured during the process. The (2:1) ZrCl^/CrClg activator combination generally provided no increase compared to the YClg/CrClg combination in the chromium composition or the volume fraction of a-cr second-phase particles. In fact, chromium compositions were actually reduced slightly compared to coatings produced with only a ZrCl^ activator (Table 20). Weight gain/area also decreased slightly, which can be attributed to the drop in chromium composition, while coating thicknesses remained unchanged. 148

R e n é 8 0 Alloy , l O O p m I " ' I

c .1 : mmli’

f R e n e 8 0 A' à

Figure 44: Cross-sectional optical micrograph of a René 80 alloy substrate treated at 1150 C for 24 hours in a (2:1) YCl /CrClj-activated pack containing 25 wt.% Cr-7.5A1 masteralloy. (a) Powder contacting and (b) "above pack" arrangement. 149

Cociwinn

Interdiffusion Zone

üsaiirasn

r. •■ ': ■.'. •.•>.> . H é - -1'

Figure 45: Cross-sectional electron micrograph of a René 80 alloy treated at 1150 C for 24 hours in a 2 wt.% NH.Cl- activated pack containing 2 wt.% ZrO and 25 wt.% Cr-IOAl masteralloy, and the corresponding x-ray maps of zirconium and oxygen. Powder contacting arrangement. 150

ITSOL calculations could not be used to explain the resulting increase in the Cr content of the coating. Since the program estimates a molar amount of each initial pack species, itis difficult to distinguish dual activators of the same halide and to calculate accurate equilibria. But, CrClg additions should certainly increase the vapor pressure of the CrCl^ species and the subsequent deposition of chromium. The exact atomic states for the deposited RE(Zr,Y) are of great interest. ITSOL calculations for ZrCl^- and NH^Cl-activated packs with 2 wt.% ZrOg predict significant vapor pressuresfor the ZrCl^ species to support simultaneous deposition of Zr, with A1 and Cr. From Table 19, the surface compositions of Zr were quite high (1-2 at.%), well above the solubility limit (0.3 at.%, [23]) in the B-NiAl matrix. Therefore, zirconium must be present as a Zr-rich, second-phase particle. X-ray (WDS) elemental maps of Zr and oxygen analysis are shown in Fig. 45. The results indicate that zirconium exists in both elemental and oxide form, possibly entrapped ZrO^. On the other hand, ZrCl^-activated packs deposited much lower levels of zirconium, within the solubility limit (Fig. 40a). WDS analysis failed to detect any relationship between zirconium and oxygen, although entrapped AlgO^ was detected with both XRD and WDS analyses (Tables 19-22). 151

To produce optimum coatings, elemental zirconium with limited pack entrapment is preferred; therefore ZrCl^-activated packs indicate the most promise. The deposition of yttrium is a different concern. ITSOL calculations predict significant vapor pressures for the YClg species in YCl^-activated packs. But microprobe analysis only detected sparse yttrium contents (200-300 ppm max.) near and along grain boundaries. Yttrium has a very limited solubility in the B-NiAl compound (0.05 at.%)[84] and in NiCrAl alloys (0.08 at.%)[20] and, therefore, segregates to grain and interphase boundaries. WDS analysis could not detect any additional yttrium at the c-Cr/B-NiAl interface. The yttrium levels detected by the microprobe are very low and close to the limits of the instrument. Therefore, more accurate surface microanalysis is required. A René 80 alloy coated in a pack containing 2 wt.% YCl^ and 25 wt.% Cr-7.5A1 wt.% masteralloy at 1150°C for 24 hours was analyzed by x-ray fluorescence spectroscopy (XFS) at the NASA Lewis Research Center, Cleveland, Ohio. XFS is a bulk surface microanalysis technique for measuring elemental quantities down to 100 ppm. Only 0.04 at.% yttrium (400 ppm) was detected across the coating surface. According to the literature[24], these additions are expected to improve the scale adherence of the NiAl coatings. 152

As verified with WDS (Fig. 45) and the XRD analyses and shown in Fig. 44a, some pack powder particles (AlgOg and ZrOg) were embedded within the outer layer of substrates coated using the powder contacting substrate/pack arrangement (Fig. 27). Because substrates were in intimate contact with the pack powders and because the coatings are outward-grown, the aluminide coating can easily envelop or entrap particles during growth, degrading the integrity of the coating. However, pack entrapment was eliminated by the "above pack" arrangement (Fig. 28). Figure 44b is an example of an AlgOg-free, alloyed 6-NiAl coating with yttrium deposited from the gas phase. "Above pack" arrangement; A series of experiments were performed to determine the optimum pack chemistry. The results are summarized in Tables 23 and 24, and a representative cross-sectional micrograph and a composition profile for the "above pack" aluminide coating is shown in Fig. 46. The coatings consisted of an outward-grown, hypostoichiometric B-NiAl matrix with a dispersion of a-Cr second-phase particles throughout the coating. An interdiffusion zone containing blocky Cr(Mo,Nb)-rich carbides and, for substrates alloyed with cobalt, a series of Cr(Mo,Co,W)-rich fingers (probably sigma phase) were again evident. 153

..>V

-■X/-'. ^ ( • "-X ''}■ I (Mo,Nb,Cr)C s- V

'' ' ' ' . ' " ' ' '' V' /: \ ''sk '

IN 713LC y + y SUBSTRATE ^ 50 |i/n

0.03 at.% Yttrium detected

c o

n O,o s 3 m

20 40 65 70 65 100 Distance from External Surface, /zm

Figure 46: Cross-sectional optical micrograph and composition profile (EPMA) of a IN 713LC alloy treated at 1150 C for 24 hours in a (2:1) YCl^/CrCl^-activated pack containing 25 wt.% Cr-7.5A1 masteralloy. "Above pack" arrangement. 154

Table 23: Summary of René 80 alloys coated in a pack containing 25 wt.% Cr-7.5A1 wt.% masteralloy at 1150 C for 24 hours. "Above pack" arrangement.

Ave. Surf. Comp, (at, %) XRD A m /A _ A t Chemistry Ni Al Cr RE Phase YCl (2) 48.5 40.2 4.3 .02Y B-NiAl 14.1 62 " 3(4) 48.8 35.9 6.2 .02Y It 18.6 65 " " , Y 0 (2) 49.4 39.7 4.6 .OlY II 12.5 43 " ", 3(3) 50.4 37.3 5.8 .02Y II 12.3 39 ZrCl (2) 49.8 40.1 5.0 .05Zr It 18.7 91 " Î4) 49.8 33.4 7.2 .24Zr II 23.4 125 " (6) 50.1 34.1 7.8 . 16Zr II 24.1 155 NH.Cl,ZrO (2) 47.5 36.8 7.5 .18Zr II 24.1 107 A " (4)2 52.9 35.1 5.3 .17Zr II 25.9 115 " " (6) 50.9 37.3 4.8 .22Zr II 25.8 120 " ,Y_0 (2) 49.9 40.6 4.3 .02Y II 11.2 32 " , 3 (3) 47.9 38.9 6.4 .03Y II 15.9 56 YCl_/CrCl_ 49.7 35.8 7.5 .03Y II 13.6 42 " fl:I) 47.4 37.2 7.2 .02Y II 15.6 53 Amounts in parentheses •

Table 24: Summary of nickel-base alloys coated with a pack containing 4 wt.% activator and 25 wt .% masteralloy at 1150 C for 24 hours . "Above pack" arrangement. Ave. Surf. Comp. (at.%) XRD Am/A At Chemistry Ni A1 Cr RE Phase ma/cm- fixa René 80: YCl ,Cr5Al 47.4 35.1 10.5 .04Y 6-NiAl 9.7 31 " ;Cr6Al 48.1 40.6 4.6 .03Y II 13.1 52 ZrCl.,Cr5Al 13.6 2.1 79.6 - 10.6 21 " ,Cr6Al 47.2 41.9 4.6 .23Zr B-&ÏAÎ 17.1 62 IN 713LC: YCl ,Cr5Al 50.1 32.8 12.3 .07Y ",a-Cr 11.7 40 " ;cr6Al 46.8 43.6 3.8 .05Y II 13.1 44 ZrCl.,Cr5Al 15.3 2.5 78.7 — 10.1 26 " .Cr6Al 51.6 44.1 4.1 .lOZr B-â?AÎ 18.9 78 155

Several important observations should be noted. First, generally about a 33% reduction in the overall thickness of the outer aluminide coating layer resulted for the "above pack" method compared with coatings produced using the powder contacting arrangement. Second, the surface composition of chromium and the weight gain/area values were also reduced significantly. Reduction of the chromium content could not alone account for the reductions in the weight gain/area values. Instead, the elimination of pack entrapment and the formation of a thinner aluminide coating were attributed as the main causes. According to Levine and Caves[65], the thickness of the coating produced from a source of unit activity depends on the instantaneous flux of A1 halide vapors to the substrate surface. These authors determined experimentally that fluoride and chloride activators produced the thickest coatings because they generated the highest A1 halide vapor pressures. But the reductions observed above for the same pack chemistry (equilibrium composition) are still very puzzling. The only difference between these coatings and those of Tables 19-22 was the substrate/pack arrangement. The "above pack" treatment separated the powders and the substrate by a porous alumina tube with a wall thickness of 10 mm. Therefore, 156 the only significant difference is the diffusion distance (§) which, according to Eq. (2.3), would account for a smaller flux of aluminum and chromium to the substrate, producing a thinner aluminide coating with less chromium enrichment. Effect of activator content; The amount of activator was also varied to determine the optimum pack chemistry. The coating morphology remained the same, but the surface RE content generally increased as the amount of the activator was increased. Both the aluminide layer thickness and the amount of chromium enrichment were increased. These observations can both be explained by the Le Chatelier-Braun principle. The additional activator content or the venting of halide species from the semi-closed crucible shifts the equilibrium established in Eq. (4.8). When the amount of activator is increased or the amount of product species are decreased, equilibrium must be reestablished by increasing the amount of the AlCl^ , CrCl^, and RECly_^ species formed. Higher vapor pressures and fluxes for these species would result in a thicker aluminide coating with further chromium and RE enrichment. This was also demonstrated by Gupta et al.[68] for AIF^-activated packs and shown in Fig. 14. These authors attributed the larger deposition rates to a transition in the deposition 157 mechanism from a recirculation to a condensation mechanism. Effect of RE oxide content;Two distinct behaviors resulted (Table 23). According to pack equilibria calculations (ITSOL) and mass spectrometer measurements, the vapor pressures of the AlCl^ species decreased when the REOy was converted to RECly species. Increasing the amount of RE oxide additions such as which are more stable than alumina, shift the equilibrium established in Eq. (4.7). Since theAlCl^ species predominate the conversion reaction, increasing the YgOg content would effectively reduce the AlCl^ vapor pressures further while generating more YCl^ and increasing their vapor pressure. The reduction in the vapor pressures of the AlCl^ species was also accompanied with an increase in the vapor pressure of the CrCl^ species. Therefore, a larger flux and surface composition for chromium and yttrium would result. Because of the reduction in the vapor pressure of the AlCl^ species, a smaller flux of aluminum would result, producing a thinner aluminide layer. 158

Zirconia additions, on the other hand, had the opposite effect. Chromium surface compositions actually decreased while the thickness of the aluminide coating increased without a major increase in the aluminum surface composition. No logical explanation exists to account for these observations. Chromium enrichment; Further chromium enrichment was desired. Pack chemistries with a larger chromium activity and composition in the Cr-Al binary masteralloy (e.g., Cr-6A1 and Cr-5A1 wt.%) were used to coat René 80 and IN 713LC alloys with either 4 wt.% ZrCl^ or else YCl^ activator at 1150°C for 24 hours. A summary of these results are listed in Table 24. The coating morphologies are similar to that shown in Fig. 46. No significant chromium enrichment resulted from packs containing a Cr-6A1 wt.% masteralloy, although packs using a Cr-5A1 wt.% produced quite a significant chromium enrichment. In fact, ZrCl^-activated packs produced an external, CrggCg layer while YCl^-activated packs produced a hypostoichiometric, B-NiAl layer with up to 12 at.% chromium, mostly in the form of a-cr second-phase particles. These coatings were difficult to reproduce and sensitive to the overall size of the retort/pack dimensions. Therefore, more development is necessary to improve the chromium enrichment of the NiAl coatings. 159

A new pack chemistry containing 1/4 wt.% pure aluminum powder, 20 wt.% Cr-5A1 masteralloy, and 4 wt.% YClg activator was used to increase the chromium surface composition. The average chromium surface composition of René 80 alloy substrates treated at 1150°C for 24 hours in this pack mixture was 7.8 + 0.78 at.%. This mixture was devised so that, initially, an aluminide coating will form to inhibit the formation of an external carbide. Then once the aluminum source is depleted, the chromium-rich masteralloy will effectively enrich the aluminide coating with chromium. This is very similar to the process developed by Galmiche[13]. Yttrium content; The yttrium surface composition was indeed quite sparse (200-500 ppm) and localized in the B-NiAl coating. To supplement microprobe analysis, secondary ion mass spectroscopy (SIMS) analyses were conducted at the National Research Council, Ottawa, Canada, on two chromium/yttrium-modified aluminide coatings. The pack mixtures contained either 4 wt.% YCl^ and 25 wt.% Cr-6A1 wt.% or else 2 wt.% NH^Cl, 2 wt.% YgO^, and 25 wt.% Cr-7.5A1 wt.% at 1150°C for 24 hours on René 80 alloys. SIMS uses a Ga^ ion beam to ionize the atoms of the coating surface, exciting atoms away from the surface. The ionized atoms are then collimated and collected by the mass filter of the mass spectrometer. 160

The mass spectrometer analyzes the ions according to their mass-to-charge ratio, identifying each element with extremely high accuracy. SIMS elemental maps for nickel, aluminum, chromium, and yttrium were determined for each coating treatment, and a representative example is shown in Fig. 47. The results indicate that the surface consists of chromium-rich second-phase particles dispersed throughout the fi-NiAl grains as well as along grain boundaries. A nickel- and yttrium-rich phase, possibly NigY or NigY[53], was present near but not at grain boundaries, such as at the a-Cr/B-NiAl interphase boundaries. But the yttrium content was not continuous and was localized in the coating surface. SIMS analysis could not measure the quantity of yttrium in the coating surface; microprobe analysis measured between 200 and 800 ppm of yttrium in the coating surface, which is expected to produce adherent AlgO^ scales upon exposure. 5.1.2 RE-Doped Aluminide Coatings Chromium additions to B-NiAl compounds and coatings have been found to improve the resistance to hot corrosion (fused-salt) attack.[147] But, chromium additions above 3 at.% are not desired for optimum oxidation-resistance.[26] To produce an optimum oxidation-resistant aluminide coating, a RE(Zr,Hf,Y)-doped aluminide coating was developed. The results are listed in Table 25 and 161

Figure 47; SIMS map of nickel, aluminum, chromium^ and yttrium of a René 80 alloy substrate treated at 1150 C for 24 hours in 3 wt.% YC1_-activated pack containing 25 wt.% Cr-7.5A1 masteralloy. "Above pack" arrangement. 162

. Coating

Figure 48: Cross-sectional optical micrograph of an René 80 alloy substrate treated at 1150 C for 24 hours in a 2 wt.% NH.Cl-activated pack containing 2 wt.% Y O and 25 wt.% Cr-IOAl masteralloy. "Above pack" arrangement. 163 a representative cross-sectional micrograph is shown in Fig. 48. These coatings consisted of an outward-grown, hypostoichiometric B-NiAl matrix with a dispersion of a-Cr second-phase particles which precipitate near the original substrate surface resulting from the outward diffusion of chromium. The coating morphology and the surface compositions are consistent with aluminide coatings produced on Ni-15Cr at.% alloys by a low activity (Ni-40A1 at.%) source with an AIF^ activator at 1150°C[148] and a commercial aluminide coating, GE Codep C. Both RE-base activators and RE oxide additions deposited a significant RE amounts in the B-NiAl coating layer. A more concentrated, Cr-lOAl wt.% masteralloy was used in all cases so that limited chromium deposition occurred from the gas phase. Although some preliminary coatings were produced with an Fe-lOAl wt.% masteralloy, to eliminate iron contamination, an aluminum-rich, Cr-Al binary masteralloy was then used because of complications in obtaining a comparable Ni-Al binary alloy powder. The RE-doped aluminide coatings were about 50-65 /im thick with 2 about a 15-17 mg/cm weight change/area following the coating process. 164

Table 25: Summary of IN 713LC and Mar-M247 alloys coated with a pack containing 25 wt.% Cr-IOAl masteralloy at 1150 C for 24 hours. "Above pack" arrangement.

Ave. Surf. 1Comp, (at.%) XRD Am/A At Chemistry Ni A1 Cr RE Phase ma/cm- fim IN 713LC: NH.C1,Y,0,(2) 49.5 45.1 2.1 .06Y 6-NiAl 17.1 69 ,Zr8j2) 51.4 45.6 2.6 .21Zr II 20.4 62 " ,no RE 50.3 44.8 2.6 - II 14.5 50 YCl (4) 49.8 45.8 2.9 • 08Y II 16.4 48 ZrCl (4) 52.3 44.1 2.5 .17Zr II 16.9 61 HfCl*(4) 50.2 46.9 2.8 .15Hf II 16.8 56 Mar-M247: NH C1,Y 0 (2) 50.2 46.6 2.2 .05Y II 13.9 53 A,no 49.2 45.8 2.4 - II 14.6 61 YCl f4) 48.9 46.3 2.5 .04Y II 13.9 45 Amounts in Parentheses •

Table 26: Summary of Mar-M247 alloys coated in .a pack containing 4 wt. NH.Cl activator, 20 wt.% masteralloy, and a silicon source at 1150 C for 24 hours. "Above pack" arrangement. Ave. Surf. Comp, (at .%) XRD Am/A .,At Chemistry Ni A1 Cr Si Phase mo/cm— f/m Sic source: Crl5Al,(2) 15.3 2.1 78.5 1.25 Cr C 22.8 0 Cr20Al,(2) 49.2 41.3 3.5 .08 8-&ÎAÎ 14.8 85 " ,(4) 14.5 2.3 79.3 1.87 Cr C 10.7 33 Cr25Al,(2) 50.1 40.5 3.1 .02 6-âÎAÎ 14.7 85 Cr-Si masteralloy: CrlOAl,CrlOSi 6.1 .1 90.4 .62 Cr C 21.2 21 Cr20Al,Cr20Si 46.7 45.9 2.4 .08 6-SJa 5 14.4 95 " ,Cr30Si 42.1 52.4 2.6 .13 II 20.6 100 " ,Cr40Si 40.1 54.7 1.9 .21 II 23.6 91 CrlOAl. Si 49.1 41.4 3.5 .08 II 15.9 54 Amounts in parentheses. 165

M . - 1 U)M

•CÎ 300 - o b /S~NiAl Coating Layer 250-

106

GOO.

: / / I 460. V i

400.

/î-NiAl Interdiffusion Alloy Coating Zone 350- 60 100 160 200 250 300 350

Distance from External Surface, fxia

Figure 49: Variation in microhardness as a function of distance from the external surface for a IN 713LC alloy substrate treated as in (a) Fig. 48 and (b) Fig. 46. 166

Another interesting feature of the RE-doped aluminide coating is their lower, and therefore more favorable, surface microhardness. Figure 49 presents an average microhardness (VHN) profile of a yttrium-doped and a chromium/yttrium-modified aluminide coating on an IN 713LC substrate. The microhardnesses generally increased toward the interdiffusion zone from the externalsurface. In fact, the microhardness at the surface of the yttrium-doped aluminide coating is about 259+6 and increases to about 433+11 at the external coating/interdiffusion zone interface. The microhardness is very dependent on the amount of chromium in the B-NiAl layer[149], and the chromium composition is highest at the external coating/ interdiffusion zone interface, decreasing toward the external surface. Therefore, the highest microhardness is associated with the highest chromium content where both solid-solution and precipitation hardening by the formation of a-cr second-phase particles occurs. The chromium/yttrium- modified aluminide coating, on the other hand, had a higher surface chromium composition (7.5 at.%) and a higher surface microhardness, 408+25. Again, the microhardness increased to a maximum value (467+31) at the interdiffusion zone where the chromium composition was highest (9.7 at.%). 167

The two-phase coating, 8-NiAl matrix with a-Cr dispersoids, may possess favorable mechanical properties although the surface microhardness is higher than the single-phase coating. However, the addition of chromium to form softer, ductile second-phase particles has been shown to improve both the hot workability[150] and the creep resistance[151] of B-NiAl compounds. In addition, soft, second-phase particles in a brittle matrix may improve the fatigue resistance of the coating/component couple. 5.1.3 Chromium/Silicon-Hodified Aluminide Coatings Silicon additions have been found to improve both the environmental durability resistance of MCrAlY overlay [47] and aluminide coatings[12] as well as B-NiAl compounds. [26,27] Although the role of silicon has been widely debated. Si additions indeed reduce both the isothermal and cyclic oxidation rates of aluminide coatings. The Si additions (1-2 at.%) are quite dilute, so that a continuous SiOg scale could not form during exposure. According to Lee and Kroger[152], silicon can dope AlgO^ scales, reducing defect concentrations and the ionic conductivity. Therefore, a slower-growing, AlgO^ scale forms during exposure, which may be altered to improve resistance to acidic dissolution and fused-salt attack, but this is very doubtful. By similar coating 168 processing presented above, a Si-modified, Cr-enriched aluminide diffusion coating on Mar-M247 nickel-base alloys might be achievable in a single-step, chloride-activated pack cementation process. Two coating treatments were used: (a) a series of Cr-Al binary masteralloy compositions with a constant silicon source (2 wt.% SiOg) and (b) a combination of binary Cr-Al and Cr-Si binary masteralloys both with a constant activator (4 wt.%

NH 4 CI). SiOg source: A summary of the coating results are listed in Table 26, and a representative cross-sectional micrograph is shown in Fig. 50a. Table 26 illustrates the difficulty in simultaneously depositing aluminum, chromium, and silicon. Packs containing a Cr-IOAl wt.% or Cr-15A1 wt.% masteralloy produced an outer, chromium-rich carbide, Cr^^C^, with minor Si. Beneath the outer carbide layer, AlgO^ precipitates were formed and detected by WDS, A1 and oxygen x-ray maps. These precipitates probably resulted from the internal oxidation of the alloy A1 content. A displacement reaction can occur between the SiOg in the pack and the aluminum in the alloy, according to the following reaction: 3/2Si02(s) + 2Al(alloy)=3/2Si(carbide) + AlgO^fs) (5.1) 169

S-MiAl 4 Coating

•s if 100pm HarH.247 oh=— -"—r ■c>. t' -'■3

" ' ..

* -V .' . ' S3

i f

lOOpm Main 247 A y

Figure 50: Cross-sectional optical micrograph of a Mar-M247 alloy substrate treated at 1150 C for 24 hours in a 4 wt.% NH Cl-activated pack containing (a) 4 wt.% SiO and 25 wt.% Cr-25A1 masteralloy or (b) 20 wt.% Cr-20A1 ana 5 wt.% Cr-40Si masteralloy. "Above pack" arrangement. 170

Figure 52a is cross-sectional micrograph of a coating produced from a pack containing a Cr-25A1 wt.% masteralloy. This treatment produced an outward-grown, hypostoichiometric fi-NiAl coating with no second-phase a-Cr formation and very limited Cr or Si enrichment. Unfortunately, the formation of an aluminide coating requires a higher A1 activity (Al-rich masteralloy) than those previously used in the chromealuminizing treatments, thus producing relatively low Cr and Si halide vapor pressures and surface compositions. To increase Cr and Si enrichment, pack chemistries containing sequentially more SiOg were used. Again, an outer, chromium- and silicon-rich carbide was formed with a large volume fraction of AlgO^ precipitates beneath the outer layer. The higher SiOg content in the pack reduces the AlCl^ vapor pressures below the value necessary to form an aluminide phase (Fig. 36c). Instead, the equilibrium vapor pressures of the CrCl^ and SiCl^ species are sufficiently high to form an external carbide. The volume fraction of AlgO^ precipitates increases with increasing SiOg content in the pack mixture, confirming the displacement reaction above (Eq. (5.1)). 171

Masteralloy combinations; A summary of these results are listed in Table 26, and a representative cross-sectional micrograph is shown in Fig. 50b. Packs containing a Cr-IOAl wt.% and Cr-15A1 wt.% masteralloy produced an external, chromium- and silicon-rich carbide. Packs containing a Cr-20A1 wt.% masteralloy produced a hypostoichiometric, 13-NiAl coating which was again free of any a-cr second-phase particles and lean in chromium and silicon enrichment. To increase chromium and silicon enrichment, successively higher silicon containing Cr-Si masteralloy compositions were also used. Figure 50b is a cross-sectional micrograph of a Mar-M247 alloy coated in a pack containing 20 wt.% Cr-20A1 wt.%, 5 wt.% Cr-60Si wt.%, and 4 wt.% NH^Cl activator at 1150°C for 24 hours. This pack produced an outward-grown, hyperstoichiometric fi-NiAl coating which lacked any a-Cr second-phase particles and was lean in chromium with a higher Si content (0.218+.05 at.%). Although the incentive to codeposit Al, Cr, and Si is apparent, the competition in the chlorination reactions between Al, Cr, and Si is very biased in favor of Al (Figs. 36c and 38). Only a small range of pack chemistries produced the aluminide coating layer necessary for environmental durability. Pack mixtures with relatively high aluminum activities (e.g., masteralloy of 172

Cr-20A1 wt.%) are needed to reduce the chromium activity and chromium halide vapor pressures so that a Cr carbide layer is avoided. Several pack chemistries produced measurable Si deposition, but these levels fell below those concentrations reported to be optimum in the literature for hot corrosion resistance. Although the oxidation behavior of silicon-doped aluminide diffusion coatings may be substantially improved.[26,27] Because of the large differences in the thermodynamic stabilities of the AlCl^, SiCl^, and CrCl^ species, the realization of simultaneous deposition does seem possible. The real obstacle to simultaneously depositing Al, Cr, and Si is the relatively high carbon contents and carbon activity in the commercial nickel-base alloys. Rapp and coworkers[40,77] and Miller[42] have illustrated this "carbon effect" while trying to codeposit Al and Cr into carbon-containing iron- and cobalt-base alloys. Free carbon reacts with chromium from the gas phase to form an outward-grown, CrggCg layer. Unless an aluminizing pack (i.e., relatively high Al activity masteralloy) is used to push carbon into the austenite substrate, external carbide coatings are formed. Most conventionally cast, nickel-base alloys have high amounts of carbon (Table 2), much of which is freed at the 1150°C coating temperature. Manufacturers of nickel and nickel-base alloys have 173 devised a number of alternatives to improve alloys with high carbon contents. One solution is to reduce the alloy carbon content (e.g., IN 713LC, 0.22 at.% carbon), or another solution is to add strong carbide formers to tie up free carbon and strengthen grain boundaries (e.g., Mar-M247 with hafnium and tantalum). Unfortunately, these alloys both suffer from the same tendencies to form an external carbide during simultaneous deposition of Al, Cr, and RE as do other alloys with higher free carbon contents (e.g., René 80). The sole solution is to reduce the carbon content to lower levels, for example, use high- purity alloys such as directionally-solidified, single-crystal alloys (Table 2), where carbides are not needed to stabilize grain boundaries. Carbon effect; A simple experiment was performed to illustrate the "carbon effect" in nickel-base alloys. Two René 80 alloys (e.g., 0.8 at.%) were coated in a 4 wt.% YClg and 25 wt.% Cr-5A1 wt.% masteralloy at 1150°C for 24 hours. After coating, a thin external carbide was formed. This carbide was then mechanically removed (ground off with Sic paper) from the nickel-base substrate. The carbide-free substrate was then recoated in an identical pack mixture and treatment. The resulting coating was then an aluminide coating enriched with chromium and yttrium (10.7 at.% chromium and 0.05 at.% yttrium). These 174

660 o- ■ • o Cr20Al. CrZOSi o D Cr20Al. Cr30Si 600 ' - /' 1\ Cr20Al. Cr40Si

I 4 6 0 - - 01 o01 400 «3 Xi o 3 6 0 - - s

3 0 0 --

260 10 30 60 70 00 110 130 Distance from External Surface, /zm

Figure 51: Variation in microhardness as a function of distance from the external surface of a Mar-M247 alloy substrate treated at 1150°C for 24 hours in a 4 wt.% NH.Cl-activated pack containing 20 wt.% Cr-20A1 and 5 wt.% of a Cr-Si binary masteralloy. 175 results illustrate the difficulty in coating nickel-, cobalt-, and iron-base alloys with high carbon activities, and also presents an unpractical solution. Microhardness (VHN) measurements were made on the outer aluminide layer of coated Mar-M247 alloys. Microhardness values increased (483+17 max.) toward the external surface of the 6-NiAl coating (Fig. 51). The depth of the increased microhardness values agrees well with the extent of the hyperstoichiometric region of the J3-NiAl coating (gray region in Fig. 50b) . Therefore, aluminum enrichment and to a lesser amount silicon enrichment caused the high microhardness values in the outer aluminide layer; these were higher than those for Cr/RE-modified aluminide coatings (Fig. 49). This increased surface hardness again may possibly degrade the mechanical properties of the coating/substrate couple.

5.2 KINETICS AND FORMATION MECHANISMS OF AL,CR,RE COATINGS 5.2.1 Powder Contacting Arrangement A kinetic study was conducted to determine the kinetics and the formation mechanism for the Cr/RE- modif ied aluminide diffusion coating. The growth rate and compositional changes of aluminide coatings formed on René 80, a conventionally cast polycrystalline alloy, at 1150°C were determined. The two pack chemistries studied contained either 2 wt.% YCl^ or else ZrCl^ with the 25 176

40 40 cv a6 \AO - - 3 0 e

§ 1 0 -. -- 10

0 10 20 30 40

Coating Time at Temperature, hrs.

I 1 6 0 --

g 1 0 0 - - - - 100

Ia Üo GO • - •d0) dM oCO

1 2 3 4 G G 7

Square Root of Coating Time at Temperature, hr^/^

Figure 52: Kinetic results of a René 80 alloy substrate treated at 1150 C in a 2 wt.% YC1_- or ZrCl.-activated pack containing 25 wt.% Cr-7.5A1 masteralloy: (a) weight and (b) coating thickness changes. Powder contacting arrangement. 177

wt.% Cr-7.5A1 wt.% masteralloy. The changes in weight and coating thickness (measured by optical microscopy) with the coating time at temperature are presented in Fig. 52. Zero hours at temperature pertains to the heatup period illustrated in Fig. 29. A thin aluminide layer was observed during heatup. Because the powder contacting method was used, weight change measurements were tainted by AlgOg entrapment, therefore thickness measurements were more appropriate. The kinetics of aluminide growth are governed by Eg. (2.5) : W = Kg(t)" (2.5) where W is the change in weight or thickness, is the coating growth rate expressed as weight or thickness, n is the growth rate exponent for weight or thickness change, respectively, and t is the time at the 1150°C temperature. The kinetic statistics were calculated from Fig. 52 and are listed in Table 27. From Table 27, the kinetic results indicate that coating growth was parabolic, and therefore diffusion controlled. In fact, growth rates were comparable to a study on the aluminization of Ni-15Cr at.% alloys with a Ni-40A1 at.% source and a AIF^ activator at 1100 and 1150°C.[148] These results also indicated parabolic kinetics with a growth rate constant, by weight, of about 6.03 mg/cm^hr^ at 1150°C and an 178

apparent activation energy of 51.5 kcal/mol in the temperature range tested. This agreed very well with the activation energy for the interdiffusion of nickel and aluminum in a hypostoichiometric, B-NiAl compound and indicates that growth is controlled by outward diffusion of nickel in the solid-state. In an earlier aluminization study, Levine and Caves[65] produced a hybrid aluminide diffusion coating from a pack containing a pure aluminum source and a NaCl activator at 1150°C which required no additional homogenization treatment. The authors determined that the growth of aluminide coatings was parabolic and controlled by solid-state diffusion. They also calculated an apparent activation energy of the process to be 88+12 kcal/mol, well above interdiffusion values listed in Table 1. The higher activation energy values of this hybrid process, unlike typical "low" activity processes, results because of the formation and growth of the NigAl^ coating and subsequent transformation to the B-NiAl phase. The growth rates of the YCl^- and ZrCl^-activated packs were slightly different. YCl^-activated packs produced higher growth rates, by weight and thickness, compared to the ZrCl^-activated packs (Table 28). Since the flux of the source elements contributes to the overall deposition rate, stable condensed activators, such as 179

o A o o Al O A— 6 Cr o o— □ Zr*10 Q) Od Pu CO o 1 0 - - obO fe a

10 1 6 2 6

10

8 - -

Cr

4

Zr*10

10 1C 25 Coating Time at Temperature, hrs.

Figure 53: Kinetic results of a René 80 alloy substrate treated at 1150 C in a 2 wt.% ZrCl .-activated pack containing 25 wt.% Cr-7.5A1 masteralloy: (a) composition changes and (b) magnification of (a). Powder contacting arrangement. 180

(Gas Phase) Temp.“IISO'C YClj, OlCl or ZrCl

Coating p-Hini(RE) p-Nini*<-Cr

Interdiffusion MC.q" phase. Zone ~ — - p-Hifll

Hickel-Base V-Ni+NigAl Substrate '—■ !Ti Ppi

(Gas Phase) vci nici X CrCl or ZrCl vl/ Coating

9 Interdiffusion MC.or phase. Zone p-Hiftl

Hickel-Base Substrate Cr TiHi Co V-Ni*Nio01 (b)

Figure 54; Schematic illustrations of the formation mechanism of the modified aluminide coating; (a) relatively short times and (b) longer time. 181

YClg, may retain their (lower) equilibrium vapor pressures over the process time, thereby increasing the overall deposition rates. Volatile activators, such as ZrCl^, are vented from the semi-closed retort and do not retain their equilibrium vapor pressures predicted by ITSOL. Therefore, deposition and growth rates are lower than for a condensed activator.[71] The slow venting of ZrCl^ species should decrease the overall rate of zirconium enrichment. Surface compositions of coatings formed in a ZrCl^-activated pack versus time at temperature are shown in Fig. 53. Optimum zirconium deposition (0.1-0.2 at.%) occurs within the first four hours of the coating process. Chromium enrichment also occurs in this time frame, leveling off thereafter. Similar behavior for YCl^-activated packs was observed, except yttrium levels were again very low and generally independent of time. Figure 54 presents a schematic illustration of the inferred formation mechanism for the Cr/RE-modified aluminide diffusion coating. Initially, the required pressure gradients exist for the RECl^ and AlCl^ species because of the lower substrate activities, but some initial chromium deposits once the aluminide coating is formed (Fig. 54a). At later times (2-4 hours), the necessary partial pressure gradient for the CrClg(v) 182

Table 27: Kinetic results of René 80 alloys coated in a powder contacting arrangement with a pack containing 2 wt.% activator and 25 wt.% Cr-7.5A1 masteralloy.

Average Growth Kinetics Activator Ks f ^m/h Ks fma/cm—h-^) n'fwt.) n”fth.) YCl 23.68 5.94 .56 .52 ZrCl, 22.40 5.30 .47 .51

Table 28: Kinetic results of nickel-base alloys coated in an "above pack" arrangement with a pack containing 2 wt.% NH^Cl, 2 wt.% YgO , and 25 wt.% Cr-7.5A1 masteralloy.

AllOV Ksfum/h^) Ks fmcf/cm—h-^) n'fwt.) n"fth.) 1150"C: Mar-M247 16.21 2.71 .42 .41 René 80 17.09 2.46 .42 .43 René 8OH 16.22 2.44 .39 .42 René N4 16.36 2.36 .37 .40 1100 C: Mar-M247 12.02 1.71 .42 .47 René 80 11.61 1.46 .30 .37 René 8OH 12.13 1.52 .40 .57 René N4 9.85 1.57 .43 .39 1050 C; Mar-M247 9.34 1.31 .58 .68 René 8 0 6.21 1.13 .58 .55 René 8OH 7.14 0.96 .56 .52 René N4 7.07 1.09 .62 .60 183

"Above Pack'

1 2 --

dû 0-- o ---- o MarM 247 û---- a Rene 80 o ---- a Rene 80H o ---- o Rene N4

1 5 0 13 17 21 26

100 'Above Pack' I OT 0m s1 bO « ° MarM 247 G “ * Rene 80 (d 40 ° — " - 0 Rene 80H oo <> o Rene N4

ZO 0 13 17 21 26 Time at Temperature, hrs.

Figure 55: Kinetic results of nickel-base alloy substrates treated at 1150 C in a 2 wt.% NH.Cl-activated pack containing 2 wt.% Y.O and 25 wt.% Cr-7.5A1 masteralloy: (a) weight and (bj coating thickness changes. "Above pack" arrangement. 184

N Qapp—24.3 kcal/mol

W i Bo o Thickness G • Weight

Qapp=27.3 kcal/mol ^

7 .0 7J2 7.4 7.6

10 4-

° This Work • Tu te Seigle

Qapp=51.5 kcal/mol

'5-,

T- Qapp—27.3 kcal/mol 4- -- H-- --- ,-- ,- ^•0 7.2 7.4 7.6 Inverse Absolute Temperature, 1 0 ^ / K

Figure 56: Arrhenius plots of the coating growth rate measured by weight and thickness changes and comparison to aluminizing rates of Ni-15Cr at.% alloys.[148] 185 species develops because of the local depletion of aluminum in the adjacent pack, and a consequent decrease in chromium activity in the substrate surface (coating). Therefore, chromium is deposited and dissolved into the external surface of the coating, saturating the B-NiAl phase (Fig. 54b). Because of the limited solubility of chromium in the B-NiAl phase (Fig. 43), a-Cr second-phase particles are precipited from the coating both during the furnace cool to room temperature and during deposition. Also, a-Cr second-phase particles of different origin are formed adjacent to the original surface as described above.[148] 5.2.2 "Above Pack" Arrangement To eliminate pack powder entrapment, the "above pack" arrangement was used to produce "cleaner" aluminide coatings. But for equivalent pack chemistries, thinner coatings with lower chromium enrichment resulted. To explain this effect, a more detailed kinetic study was conducted. Both conventionally cast, René 80 Mar-M247, and directionally-solidified, René 8OH and René N4, commercial nickel-base alloys were coated in a pack containing 2 wt.% NH^Cl plus 2 wt.% YgO^ and 25 wt.% Cr-7.5 Al wt.% masteralloy at 1050, 1100, and 1150°C. Thickness and weight change measurements at 1150°C for the coatings are presented in Fig. 55, and the kinetic values 186 from all alloys and temperatures were measured from these results and listed in Table 28. The results indicate that isothermal coating growth is again parabolic, but the apparent activation energies calculated from both weight and thickness changes were 27+3 and 24+5 kcal/mol, respectively (Fig. 56). These values are well below those measured from powder contacting arrangements, where solid-state diffusion controls coating growth. Therefore, another step in the coating process gains dominance. For local equilibrium both in the bulk of the pack and at the substrate surface, the only other contributing step in the process is the gaseous diffusion of metal chloride molecules through the porous pack medium to the substrate surface. As expected, vapor transport is inhibited by the seclusion of the substrates in the porous enclosure. Vandenbulcke et al.[153] illustrated that coating growth rates during pack aluminizing could be reduced by increasing the depth of internal passages (i.e., diffusion distance) or by reducing the total pressure of the process. Each of these variables effectively reduced the deposition rate or flux of Al, thereby maximizing the role of gaseous diffusion during the process. The activation energy for the rate of chlorination of iron and nickel, where gaseous diffusion controls, is around 10 to 17 kcal/mol.[154] Therefore, the apparent activation energy 187 of the coating process agrees well with these values, indicating that vapor transport is inhibited and gaseous diffusion controlled growth predominates. The effect of substrate grain boundaries on coating growth rates was also investigated. Since solid-state diffusion dominates coating growth, short-circuit paths such as grain boundaries in the B-NiAl may have a significant effect on coating growth rates. Unfortunately, no major differences were observed between polycrystalline or single-crystal substrates in the temperature range tested. An effect may have been observed at lower temperatures where the contribution of grain boundaries on solid-state transport is more pronounced. Instead, a more noticeable effect on the substrate chemistry or the morphology of the interdiffusion zone carbides was observed. The interdiffusion zone for coated Mar-M247 alloys contained a dispersion of (Hf, Ta, W, Cr)-rich carbides in a B-NiAl matrix (Fig. 49), but no continuous chain of carbides was observed as in René alloy substrates. This continuous chain of blocky carbides should act as a diffusion barrier, inhibiting the outward diffusion of nickel from the substrate and reducing the growth rate of 188

68 a i C o 63 -■ -4-> Otn o-- o Ni (X Û-- A Al B oO Q) (do 43 CO 38

33 1 6 8 13 17 21 26

5-

4 ,

3 • ■ o D Cr « o Y*10

2 - •

1:1

0 13 17 21 28 Time at Temperature, hrs.

Figure 57: Kinetic results of a René N4 alloy substrate treated at 1150 C in a 2 wt.% NH.Cl-activated pack containing 2 wt.% Y O and 25 wt.% Cr-7.5A1 masteralloy; (a) composition changes and (b) magnification of (a). "Above pack" arrangement. 189 the coating. In fact, coatings produced on Mar-M247 alloy substrates possessed the slowest growth rates at all three temperatures (Table 28). The outer B-NiAl layer is alloyed with Co, Ti, and Cr and to a lesser extent W, Mo, Ta, Hf, and Zr from outward diffusion during coating growth. No major difference in composition was observed between the Mar-M247 and René alloys to account for any second-element effect on the diffusivity of Ni in the B-NiAl compound. The change in surface composition for each substrate at each time and temperature treatment were determined with the electron probe microanalyzer. A representative change in composition is shown in Fig. 57 for a coated René N4 at 1150°C. The coating process is generally aluminizing. Moderate chromium enrichment of the aluminide layer occurs within the first hours and levels off to an average bulk values of about 4.25+0.79 at.%. Overall, the formation mechanism can again be illustrated by Fig. 54.

5.3 ENVIRONMENTAL DURABILITY OF ALUMINIDE COATINGS Coated turbine hardware is exposed to aggressive environmental conditions which can lead to premature failure of the component. Two of the most common forms of environmental degradation in gas turbines are cyclic oxidation and hot corrosion (fused salt) attack. To 190

CV3

tJO o oAlCIa.ZrOp G * -NH^Cl.ZrO; *- ^ YClg,Zr02 * -----^ NHx Cl

- 3 60 100 160 2 0 0

2 +

1 T

o-- o NH4CI

- 1 4- *-- * YClg a-- a ZrCl4 ''--- ^ GE Codep -2 4* «— « Uncoated

- 3 60 100 160 2 0 0 Number of Cycles

Figure 58: Kinetic results of a René 80 alloy substrate treated at 1150 C for 24 hours with a suitable Cr-Al masteralloy and cyclically oxidized at 1100 C in static air for up to 200, one-hour cycles: (a) ZrOU source and (b) RE-base activators. Powder contacting arrangement• 191

CO üB \ CO B

<3 -1 W U) CO û ------A 4 -3 NH CI (dC o--- o YClg O '' ZrCl4 - 5 JQ CO

*s - 7 o o - 0 o 0 60 100 160 200 Number of Cycles

CM

n S

ZrCl4

CO 6 16 25 36 66 Entrapped AlgOg, %

Figure 59: (a) Specific weight changes for 1200 C cyclic oxidation in static air of coated René 80 alloy substrates and (b) the relationship with the volume fraction of entrapped AlgO^. 192

Table 29: Summary of the XRD results of coated René 80 alloys cyclically oxidized in static air at 1100 and 1200 C for up to 200, one-hour cycles. at 1100 C, Pack Chemistry Oxide Phases Spall

NH^Cl,Zr 0 2 NiAlj 0 ^,Ti 0 2 ,Zr0 2 , NiAl 2 0 ^,Ni 0

YCl 3 ,Zr0 2 Al 2 0 3 fN?Al 2 0 ^,Ti0 2 NiAl 0 ,TiO , Ai;o_ 2 NH 01 Al_0_,Ti0-,NiAl_0. none YCI II zrci. II

at I 200° c , Pack Chemistry Oxide Phases Spall NH^Cl Ni0,NiAl„0 , ^ ^ 2 ® J '^ ^ ° 2 '' Al_0_,TiO_*

YCI 3 Al 0 ,TiO ,NiAl 0 , Ai;o;,Tio;, Rid,Ni ^ ^ 4 NiAl,0 ,NI0 i , i l 0 ZC^Al 0 ,Ti0-,NiAl-0 , NiO,NiAl 0 ,ZrCl^ NÎOTNiTiO ,Z r 0 2 ,Ni XRD results are listed from strongest to weakest peaks. 193 evaluate the performance of the new, modified aluminide coatings, these coatings were subjected to environmental durability testing. 5.3.1 Cyclic Oxidation One very effective way to evaluate the adherence of thermally-grown oxide scales is cyclic oxidation. Cyclic oxidation exposes a growing oxide scale to thermal stresses, facilitating spallation and loss of adherence.[18] By periodic weight measurements, the extent of oxide adherence and the cyclic oxidation behavior of the chromium/RE-modified aluminide coatings can be evaluated. Powder contacting arrangement; Cyclic oxidation at 1100 and 1200°C in static air of the Cr/RE-modified aluminide coatings on René 80 alloy produced by the powder contacting arrangement were conducted at the NASA Lewis Research Center. Coatings were produced from packs containing 25 wt.% Cr-7.5A1 wt.% masteralloy and 2 wt.% of various activators and/or RE sources at 1150°C for 24 hours. The specific weight change measurements versus cycle time are presented in Figs. 58 and 59a, and XRD analysis of the surface phases and any spall collected are listed in Table 30. The cyclic oxidation results were also compared to a commercial, low-activity aluminide coating, GE Codep C, for the same alloy substrate and a 194

luOtim

Figure 60: Cross-section optical micrograph of a René 80 alloy substrate treated at 1150 C for 24 hours in a 2 wt.% YCl_-activated pack containing 25 wt.% Cr-7.5A1 masteralloy and cyclically oxidized for 200, one-hour cycles at 1100 0 in static air. Powder contacting arrangement. 195 bulk uncoated alloy.[155] The results indicate that the coated alloys have improved the scale adherence of AlgOg scales and lifetimes of the substrate considerably. Coatings produced from pack chemistries using an YCl^, ZrCl^, or NH^Cl activator salts provided good resistance to cyclic oxidation, comparable to the GE Codep C coating. Coatings produced from packs with ZrOg additions, except for YCl^-activated packs, failed to improve coating lifetimes above the uncoated substrates. This can be attributed to the excessive amounts of zirconium in the aluminide coatings, mainly as ZrOg entrapment (Table 19). The XRD analysis of the oxidized coatings also indicated the overall improvement of coating lifetimes. According to Fig. 20, aluminide coatings (>30 at.% aluminum) with chromium enrichment should form an external AlgOg scale. In fact, successful coatings produced an a-AlgOg scale with minor peaks of NiAlgO^ spinel and TiOg phases, and no spall was collected. Unprotective coatings failedto form an adherent AlgOg scale; instead external CrgOg, TiOg, and spinelphases were formed and spalled during thermal cycling. A representative, cross-sectional micrograph of a coating oxidized at 1100°C for 200, one-hour cycles is shown in Fig. 60. Several observations should be noted. First, a very thin and adherent AlgO^ scale was present on 196 the surface of the coating, beneath which AlgOg entrapment was again evident. The post-oxidation coating consisted

of a single-phase, 6 -NiAl compound without any a-Cr second-phase particles. Because of interdiffusion between the nickel-base substrate and the coating, a second chain of MC-type carbides has begun to form beneath the original zone. Cross-sectional micrographs of unprotective coatings were also very revealing. Because of the depletion of Al to form a nonadherent AlgOg scale, no

external 6 -NiAl coating was present, instead a two phase, y + y ', surface layer with little or no interdiffusion zone was observed. Fluctuations were observed in the specific weight change versus cycle time measurements in Fig. 58, indicating periods where the external scale spalled although no spall was collected. The noticeable amount of pack entrapment (AlgO^ particles) may have been the cause. Cyclic oxidation tests at 1200°C accelerated the degradation rate of the aluminide coatings, allowing for faster evaluation. According to the literature,

sufficient amounts of Zr or Y were present in the 6 -NiAl coatings to improve scale adherence, and they should have been more protective than the undoped aluminide coatings (e.g., NH^Cl activator). But, one noticeable difference for these coatings was the volume fraction of entrapped 197

:, •'V.' \ Jif y u - .

Bemë 80

^lOOpm \ >* J * j

Figure 61: Cross-sectional optical micrograph of a René 80 alloy substrate treated at 1150 C for 24 hours in a 2 wt.% (a) NH.Cl- or (b) YC1_-activated pack containing 25 wt.% Cr-7.5A1 masteralloy and cyclically oxidized for 200, one-hour cycles at 1200 C in static air. Powder contacting arrangement. 198

AlgOg, as seen in Figs. 39 and 40. Since the same alloy substrate was used for all three coatings, the oxidation rates, once the coatings failed, should have been the same. If coating failure was dependent on the relative amount of entrapment, which could introduce thermal fatigue cracks and coating mechanical failure, then this should be apparent from these results. By plotting the specific weight change versus the volume fraction of entrapped AlgO^ measured by optical microscopy, a relationship between coating degradation and the amount of entrapment can be determined. Such a relationship is presented in Fig. 59b. These results strongly indicate that the mechanical failure of the coating or weight loss of a nonadherent external scale was dependent on the amount of entrapment. Cross-sectional micrographs of coatings cyclically oxidized at 1200°C for 200, one-hour cycles are shown in Fig. 61. Because of the higher test temperature and transformation of the coating into a nickel (y) solid- solution (Table 30), carbides in the interdiffusion zone dissociated and dissolved into the nickel matrix where they have greater solubilities. Also, failures of coatings with large amounts of randomly entrapped oxides (e.g., ZrCl^- and NH^Cl-activated packs) were characterized by large pits or scallops across the 199

w üB

M< bo - 5

- 1 0 ■o ZrCl4 . 4% ZrCU, 2% -o NH 4 CI. ZrOo - 1 6 - NH 4 CI. Y2 O3 YCI3 . 3% - NH 4 CI

0 too 200 3 0 0 4 0 0 6 0 0

o ZrCl4 . 4% • ZrCU, 2% NH 4 CI. ZrOg ' NH 4 CI. YgOc ° YCI3 . 3% » NH 4 CI

- 6

-10 0 100 200 3 0 0 4 0 0 6 0 0 Cycle Time, hours

Figure 62: Kinetic results of (a) René 80 and (b) IN 713LC alloy substrates treated at 1150 C for 4 hours in a pack containing 15 wt.% Cr-7.5A1 masteralloy and cyclically oxidized at 1100 C in static air for up to 500, one-hour cycles. "Above pack" arrangement. 200

CM üH 10 X

<

no(U C (d o— o ZrCU, 4% cj ' YCI3 . 4% o— o NH4 CI, ZrO g V V NH4 CI, YgOg 3 o « Zrd4 , 2 % •----• NH4 CI o

-10 0 100 200 300 400 500 en

15 ZrC U . 4% YCI3 . 4% 10 o --- NH4 CI. ZrO p ^ ^ "NH4 CI, YpOc > o ZrCU. 2% ►----- • NH4 CI

—5

-10 0 100 2 00 300 400 500 Cycle Time, hours

Figure 63: Kinetic results of (a) René 80 and (b) IN 713LC alloy substrates treated at 1150 C for 24 hours in a pack containing 15 wt.% Cr-7.5A1 masteralloy and cyclically oxidized at 1100 C in static air for up to 500, one-hour cycles. "Above pack" arrangement. 201

Table 30: Summary of the XRD results of René 80 and IN 713LC alloys coated in an "above pack" arrangement with a pack containing 15 wt.% Cr-7.5A1 masteralloy at 1150 C for 4 and 24 hours and cyclically oxidized at 1100 C for up to 500, one-hour cycles.

René 80 IN 713LC Pack Chemistry 4 h. 24 h. 4 h. 24 h. 2% ZrCl. 1.2.3 1.3 1.2.3 4% " * 3,2,4,1 1,2,5 1.2.3 1.2.3

2%NH.C1,2% ZrO_ 1,3,4 1,4,3,2 , 6 1.4.3 1,4,3,2 4%*YC1 ^ 1.2.4 1.2.3 1.2.3 2%NHC1,2% Y 0 4.1.2 1,2,3 1.2.3 1,2,5 2%*NH^C1 1.4.3 1/4,2,3 , 6 1.2.3 1.2.3

2 2 3 5=NiAl , and 6 =Ti 0 2 Key: l=NigAl, =Al O , 3=NiAl 2 ° 4 ' 4=Ni, 202 specimen surface. Whereas, coatings (e.g., YCl^-activated packs) with relatively small areas of entrapped oxides in small groups were characterized by small localized pits originating near these areas. Different volume fractions of entrapped oxides are produced according to the physical state of the activator and/or chloride species generated during the coating process. For example, liquid activator salts, such as YClg, or other liquid species produced in the pack wet the pack powders and bunch particles together in small groups. Volatile activators, such as ZrCl^, would not wet pack powders so that a more random and larger amount of entrapment would be produced. "Above pack” arrangement; Cyclic oxidation tests at 1100°C in static air of "cleaner" coatings on both René 80 and IN 713LC alloys were conducted. To determine any effect of coating thickness, coatings were produced from packs containing 15 wt.% Cr-7.5A1 wt.% masteralloy and various activators and/or RE sources at 1150°C for 4 and 24 hours. The weight change versus cycle time is presented in Figs. 62 and 63, and the XRD analyses of the surface phases are listed in Table 31. Pack chemistries were identified which improved the adherence of AlgOg scales and which seem to be independent with the thickness of the coating. Overall, the best coatings were produced 203

y-NigAl

6 -NiAl % • H î Coating . vVjC j-«',Vi‘-it’'- y MC T ,U

< r yR'kk». ., ^ I IM 713LC ‘ ** 50 n/n I

K 6 0 - - •--- #Y*100 d 0 — 0 2*11 o * . . . A m a o C r 0 4 0 - - V vCo* 1 0 y-NigAl /î-NiAl Ti*10 u1 0 ) 2 0 k,Q] I =+= 15 25 35 45 55 Distance from External Surface, fzm

Figure 64: (a) Cross-sectional backscattering electron micrograph and (b) composition profile of a IN 713LC alloy substrate treated at 1150 C for 4 hours in a 2 wt.% NH.Cl-activated pack containing 2 wt.% YyO. and 15 wt.% Cr-7.5A1 masteralloy and cyclically oxidized for 500, one-hour cycles at 1100°C in static air. "Above pack" arrangement. 204

% NiAl

50 Il/n

eo /S-NiAl

Cr 40 CT Y*100 Go oB

2 0

S 25 35 45

Distance from External Surface, ju.m

Figure 65: (a) Cross-sectional backscattering electron micrograph and (b) composition profile of a IN 713LC alloy substrate treated at 1150 C for 24 hours in a 2 wt.% NH.Cl-activated pack containing 2 wt.% Y.O. and 15 wt.% Cr-7.5A1 masteralloy and cyclically oxidized for 500, one-hour cycles at 1100 C in static air. "Above pack" arrangement. 205

Ai,o,

76 &( /J-NlAl i ------Ni oC •----- Al 60 - ...... Cr .— -. Ti o ------Zr*10 B 0o Q) 2 6

1

65 106 165 206 266 Distance from External Surface, //m

Figure 6 6 : (a) Cross-sectional backscattering electron micrograph and (b) composition profile of a IN 713LC alloy substrate treated at 1150 C for 24 hours in a 4 wt.% ZrCl.-activated pack containing 15 wt.% Cr-7.5A1 masteralloy and cyclically oxidized for 500, one-hour cycles at 1100 C in static air. "Above pack" arrangement. 206 from packs containing either 4 wt.% ZrCl^, 4 wt.% YCl^, or 2 wt.% NH^Cl plus 2 wt.% YgOg. RE-free or RE-lean

coatings, or ones produced from a pack containing 2 wt.% NH^Cl plus 2 wt.% ZrOg were consistently poor. XRD analyses detected a X'-Ni^Al surface phase on almost all coatings. In addition, a-AlgO^ and NiAlgO^ spinel oxide phases were evident. Less protective surface phases such as y - N i and TiOg were also detected, especially on René 80 substrates. Cross-sectional, scanning electron micrographs and composition profiles (EPMA) of oxidized coatings are presented in Figs. 64-66. Several features are of importance. A 15-20 /zm thick, aluminum depleted layer (NigAl) was characteristic of all protective coatings,

except the one shown in Fig. 6 6 . Figure 6 6 is a cross-sectional micrograph of a coated IN 713LC alloy from a pack containing 4 wt.% ZrCl^ and cyclically oxidized for 500, one-hour cycles. Composition profiles (EPMA) and XRD analysis did not detect an external, Ni^Al layer. This coating contained an external B-NiAl phase with an inner, three-phase region comprised of R+Y'+Y. Instead, EDS analysis detected a Zr-rich, internal oxide near the coating surface. Such internal oxides are characteristic

of 6 -NiAl compounds with additions of 0.2 at.% Zr, which produced very adherent AlgO^ scales at 1100°C.[23] In 207

addition, a coarsened interdiffusion zone was observed, consisting of large M^^Cg-type carbides and sigma phase. Coating thickness produced limited improvement in coating lifetimes. Thinner coatings produced from the

four hour treatment formed adherent scales between 2 0 0 and 3 00 cycles (Fig. 62). Figure 64 is cross-sectional micrograph and composition profile of a coated IN 713LC alloy from a pack containing 2 wt.% NH^Cl plus 2 wt.% YgOg and cyclically oxidized for 500, one-hour cycles. A 20 /im thick, NigAl layer at the surface of the coating and similarly coarsened interdiffusion zone phases were observed. The Ni^Al layer consumed almost the whole entire coating. Therefore, a thicker coating of the same composition may have offered longer protection. But cyclic oxidation behavior was generally affected most by the composition of the coating surface. For both coatings produced from the 4 and 24 hour treatment, the most adherent coatings were produced from consistent pack chemistries (Figs. 62 and 63) . Conversely, alloy chemistry and the resultant coating surface composition did affect the cyclic oxidation behavior. Aluminide coatings often contain various alloying elements which diffuse outward during the coating process. IN 713LC and René 80 alloy substrates have two major alloying differences (Table 2). IN 713LC 208

alloy contains no cobalt and only 1 wt.% titanium; therefore only little enrichment of titanium occurs in the coating. But René 80 alloy substrates contain about 4-5 at.% cobalt and about 1 at.% titanium. Although small additions (1-3 at.%) of titanium to B-NiAl compounds may improve cyclic oxidation behavior[26], cobalt additions to nickel-base alloys promote NiO formation and may similarly affect B-NiAl compounds and coatings, degrading their overall oxidation behavior. Therefore, coatings on IN 713LC alloy substrates would be expected to provide better protection. Figure 67 is a schematic illustration of the degradation mechanism for the Cr/RE-modified aluminide diffusion coating. Initially, transient oxides, possibly

Ni(Al,Cr) 2 0 ^ spinel and a-(Al,Cr) 2 0 ^[57], form during initial exposure to eliminate internal and "pest" oxidation, thereby reducing the oxygen activity at the coating surface. With time, the transient oxides thicken until a continuous, steady-state a-AlgOg scale forms, choking off the supply of nickel to the spinel oxide. Chromium is depleted from the surface oxides from the oxidation/evaporation of the transient oxides. In addition, interdiffusion zone carbides coarsen at the expense of smaller carbide particles and nearby a-Cr and sigma particles (e.g., Ostwald ripening). Further 209

IU(ni_^r)j,0,j -0- Tr«ns icnt O xi d e . \ • T . •• • *-T. • • T : • . Mi ' ■ . *. • . Cr . C o * t ifig ^ n-Nin^ • ni • . Cr- •. • I ntcrd if fus ion MC,«y phase, n f e i f p-Hini+Hignl S u b s t r a t e • Hi Cr V-Ni*Ni]Al

<-«1 ^0 , ■ timam/x / g O xide p-WiAl(RE>

C o a e x n e p-NiAl€RC>

^ c c rd if f ux io n MC,«r phase* p N i A l

HC

< - « 1 2 0 3 Steady-State ,_ O x i d e ----- T_ CkC C o a tin g - P HiAl MgaCa.p-Hini. Interdiffusion ° % 9 ® P o &. ° Zone cr p h a s e

M C 3

Substrate V H i + H i a A :

Figure 67: Schematic illustration of the degradation mechanism of the modified aluminide diffusion coatings. 210

Table 31: Summary of the kinetic and XRD results of RE-doped aluminide diffusion coatings on IN 713LC and Mar-M247 alloys oxidized in air at 1100°C for 44 hours and comparison with bulk 6-NiAl compounds.

Pack Chemistry Parabolic Rate. ^ XRD Phases Constant fam/cm—h-^^ and Oxides IN 713LC: NH^Cl 4.501*10 -5 B-NiAl,Al_0_, -4 NiAl.O. ^ " ,2 % ZrO^ 6 .201*10 B-NiAl,Al_0„, -5 NiAl-0.,Zr0_ " ,2 % YgOg 2.586*10 B-NiAl,Al,0_f YA10_,NiAl,0. ZrCl^ 4.401*10 -5 B-NiAI,Al,0^, -5 ZrO YCI 3 1.751*10 B-NiAl,Al_0_, YAIO ^ HfCl^ 3.942*10 -5 B-NiAI,Al O , HfOg^NiAlgO^ Mar-M247: -5 NH 4 CI 3.403*10 B-NiAl,Al O , -4 NiAl_0.,TiÔ. " ,2 % YgO^ 1.053*10 B-NiAl,Al O f -5 YAIO Y Âl^O YCI 3 5.261*10 B-NiAl,AÎ_ 0 ^ / ‘^ -4 YAIO ,TiO Cr20Al,Cr3 0Si 9.827*10 B-NiAI,Al_6_, NiAl„0. " ,Cr40Si 3.788*10 -5 B-NiAI,Al-0_, NiAl O XRD results are listed from strongest to weakest peaks.

Alloys fat. k fcrm/cm—h^l Atmosphere Reference Ni-42A1P 4 .5 *1 0 ”-5 ^ 0.1 atm O, 34 N1-47A1-.05Zr 3.8*10"-4 0.2 " 87 Ni-52Al+impltd. 9.2*10 1.0 " " 24 2 * 1 0 -^^ Y'*'/cm^ 211 exposure transforms the MC-type carbides into the MggCg-type carbides.[89] As time proceeds, a second interdiffusion zone containing MC-type carbides begins to form beneath the original zone because of the depletion of nickel from the substrate as explained earlier. Aluminum is slowly depleted from the surface and along grain boundaries of the coating upon the formation and growth of the external, Al^O^ protective scale. Aluminum depletion facilitates the transformation of the B-NiAl coating layer into a lower aluminide phase, y'-Ni^Al and possibly y-Ni. Microprobe analyses also detected about 0.05 at.% yttrium or zirconium enrichment in the Ni^Al phase which, according to Taniguchi and Shibata[156] and Kuenzly and Douglass[157], should be sufficient to produce adherent AlgO^ scales during further exposure. This may explain why slow kinetics and good adherence were observed although an external, B-NiAl layer was absent. 5.3.2 Isothermal Oxidation The oxidation behavior of RE-doped aluminide

diffusion coatings produced from a pack containing 2 0 wt.% Cr-lOAl masteralloy with various activator and/or RE sources at 1150°C for 24 hours using an "above pack" arrangement was determined at 1100°C in air at a flow rate of 0.1 1/min (STP). Weight change versus time was measured using a Cahn TG-171 thermogravimetric analyzer. 212

lE-3 - r

CM oB O E - 4 • •

6 E - 4 - -

Cr40Si — - Cr30Si 2E -4 • - ■ NH4 CI YCI3 NH4 CI.Y2 O3

SO 46

16 30 46 Oxidation Time, hrs.

Figure 6 8 : Kinetic results of (a) Mar-M247 and (b) IN 713LC alloy substrates treated at 1150 C for 24 hours in a pack containing 20 wt.% Cr-lOAl masteralloy and isothermally oxidized at 1100 C in air at a flow rate of 0.1 1/min (STP). "Above pack" arrangement. 213

The kinetics of the oxidation tests are presented in Fig.

6 8 , and XRD analyses and the parabolic rate constants

calculated from Fig. 6 8 are listed in Table 31. True parabolic rate constants, k^, of the oxidation reaction for the coatings tested were calculated from plots of the weight gain versus the square root of time, as suggested by Pieraggi.[158] In all cases, a slow-growing Al^O^ scale was produced on the RE-doped aluminide coatings. In fact, the rate constants of the yttrium-doped, aluminide coatings agreed well with yttrium implanted, B-NiAl coatings on EI-867 nickel-base alloys at 1100°C in pure oxygen[52] and RE additions in bulk B-NiAl compounds (Table 31).[24,34,87] Representative surface electron micrographs of the resulting oxides are shown in Figs. 69 and 70. A ridged, AlgO^ scale which is typically found on oxidized, B-NiAl compounds[23] was observed on all coatings, except the NH^Cl plus ZrOg pack. For undoped coatings, cracked or spalied scales were evident near grain boundaries of the coating exposing the voided surface (Fig. 69b) and contributing to the larger fluctuations in the weight change measurements. Cracking of the oxide scale was also observed in the oxidation kinetics. At the end of the transient oxidation stage, a small drop in weight was recorded indicating spallation of the oxide scale. For 214

B-NiAl Coating

Figure 69; Surface electron micrograph of a Mar-M247 alloy substrate treated at 1150 C for 24 hours in a 2 wt.% NH.Cl-activated pack containing 20 wt.% Cr-lOAl masteralloy and oxidized at 1100 C in air for 44 hours: (a) surface oxide and (b) magnification of (a). 215

Figure 70: Surface electron micrograph of a IN 713LC alloy substrate treated at 1150 C for 24 hours in a 2 wt.% NH.Cl-activated pack containing 2 wt.% Y_0_ and 20 wt.% Cr-lOAl masteralloy and oxidized at 1100 C in air for 44 hours: (a) surface oxides and (b) magnification of (a). 216

RE-doped coatings, a much smaller fluctuation was measured. In Fig. 70, scale cracking was also observed at the surface of RE-doped, aluminide coatings, but a tenacious, aluminum-rich oxide, probably Al^O^, with a small amount of RE was detected by EDS analysis beneath the outer, cracked scale. XRD analysis identified weak signals for YAlO^, ZrO^, and HfO^ phases on the Y-, Zr-, and Hf-doped aluminide coatings, respectively. No Si-rich oxide was detected by XRD or EDS analysis on Si-doped aluminide coatings, denoted as CrSOSi and Cr40Si in Fig.

6 8 a and corresponding to pack chemistries in Table 26. Generally, the RE-doped aluminide coatings produced a continuous, slow-growing Al^O^ scale during exposure. According to Fig. 69, the oxides grown on undoped aluminide coatings were not very protective. Sporadic weight-change measurements indicated that oxide scales cracked and spalled during growth, contributing to periodic weight losses during oxidation. In addition, weight fluctuations made calculations of true parabolic rate constants difficult, so comparison of RE-doped and undoped aluminide coatings could not be drawn. But, overall, the RE-doped aluminide coatings exhibited superior oxidation behavior for both nickel-base alloy substrates. 217

5.3.3 Hot Corrosion (Fused Salt Attack) Studies Gas turbine components, such as blades and vanes, are often exposed to combustion environments which, in the presence of seawater-contaminated air, can form fused salts (e.g., Na^SO^, NaVO^, NaCl). These salts can condense on cooler hardware and lead to an accelerated form of degradation called hot corrosion. To evaluate the resistance to hot corrosion attack, the Cr/RE-modified aluminide diffusion coatings were subjected to two tests; isothermal salt film and burner rig tests. Thin film tests: René 80 and Mar-M247 alloy substrates were coated at 1150°C for 24 hours in packs containing various activator and RE sources and a suitable masteralloy composition in both powder contacting (PC) and "above pack" (AP) arrangements. A thin, film of sodium 2 sulfate (5.0+1.5 mg/cm ) was applied to the coated substrates with an airbrush. The salted coatings were then placed in an alumina reaction boat and inserted in a horizontal furnace (Fig. 34) at 900°C under a catalyzed 0.1% SOg/Og gas mixture flowing at about 0.1 1/min (STP). Exposed samples were periodically quenched to room temperature (e.g., following 24,72,144,240,360, and 504 hours) to facilitate scale spallation. The coupons were weighed, resalted, and replaced into the furnace for up to 672 total hours. The weight changes for each alloy and 218

CM VB bO B o--

MI GOoT §

•S) 1 0 - - I

o 0 f ■ I , 0 ) 0 100 2 0 0I 300 400 cooI 700 <8 ^ Corrosion Time, hrs

Average Surface Composition fEPMAl

Activator -EL. C o T i RE fat.H HH Cl 4 5 . 8 2 4 2 . 2 2 5 . 6 1 5 . 2 9 .02 .00 4 6 . 1 6 4 3 . 4 1 4 . 8 7 5 . 4 4 . 0 4 .07 4 8 . 8 8 4 2 . 6 9 5 . 5 4 5 . 7 9 . 0 5 . 0 5 4 6 . 1 7 4 4 . 9 8 3 . 2 8 5 . 0 3 .02 .12 *car7.5Al wt.t Masteralloy, all else Cr6Al wt.%.

Figure 71: Hot corrosion kinetics and average surface compositions of Mar-M247 alloy substrates treated at 1150 C for 24 hours and isothermally corroded at 900 C in a 0.1% SO-ZC- gas mixture with 5.0 mg/cm Na^SO^. "Above pack" arrangement. 219

CV3 i bO 0

MI bOo

° NHaC1 ,Y2 0 3 » ZrCl 4 D-- O YCI 3 , -----V NH4 CI « GE Codep

o 0) 100 300 400 600 600 700 Corrosion Time, hrs.

average Surface Composition CEPHAS

activator H i A 1 C r C o T i R E f a t . t l HH Cl 4 6 . 7 2 4 3 . 8 8 4 . 9 2 4 . 4 2 . 0 8 .00 4 7 . 2 2 4 4 . 3 6 4 . 1 9 4 . 1 7 . 0 7 . 0 2 4 6 . 4 5 4 2 . 8 4 4 . 3 9 5 . 1 5 . 1 7 . 0 4 ZrH^ 4 7 . 8 3 4 4 . 5 9 3 . 3 6 2 . 1 9 . 0 5 .10 *Cr7-5Al wt.t Masteralloy, all else Cr6Al wt.%.

Figure 72: Hot corrosion kinetics and average surface compositions of René 80 alloy substrates treated at 1150 C for 24 hours and isothermally corroded at 900 C in a 0.1% S0_/0_ gas mixture with 5.0 mg/cm Na^SO^. "Above pack" arrangement. 220

2 0 - <30 G

0> g e to--

O S o m 100 200 3 0 0 4 0 0 6 0 0 6 0 0 Corrosion Time. hrs.

Figure 73 Hot corrosion kinetics of René 80 alloy substrates treated at 1150 C for 24 hours in a pack containing 25 wt.% Cr-7.5A1 masteralloy and isothermally corroded at 900Cwu.V in SO-yO-a 0.1% gas mixture with 5.0 mg/cm Na^SO^. Powder contacting arrangement. 221

Table 32f: XRD results of coated René 80 and Mar-M247 alloy substrates corroded at 900 C for 672 hours in 0 .1 % SO /O^ gas mixture with 5.0 mg/cm Na.SO.. "Above pack" arrangeirient L Pack Chemistry XRD Phases René 80: NH Cl NiO,Co O ,NiAl O ,CrS,Ni Al,Ni NiO,NiAl^O.,Co^O^,Ni.Al,Ni,CrS ZrCl^ " , " fNÎCr 6 7Ni,Ri-Al,Co O , CrS YCl NiO,Co^O^,NiS,CrS,Ni,Ni^Al Mar-M247: i NH.Cl NiAl O ,A1 0 ,NiO,Ni Al,CrS NiO,NiAl 0^,Ni A1 ZrCl NiO,Ni Al,NiAl^O ,ZrO ,CrS

YCl_;Cr 6 Al NiAl.O^,A1_0_,NiO,Ni_Al,CrS " . CrSAl NiO.Co^O .CrS.Ni Al ______XRD results listed from strongest to weakest peaks.

Table 33: IXRD results of coated René 80 alloy substrates isothermally corroded at 900 C for 144 and 672 hours in a 0.1% SO^/O^ gas mixture with 5.0 mg/cm Na.^SO,. Powder contacting Arrangement. '4' XRD Results Pack Chemistry 144 hours 672 hours NH Cl a-A1_0_,B-NiAl — YCI II II II YCl^/CrClg 9 a-Al 0 ,NiAl O , 6-NiAlfNi AlfctS ZrCl^ II / II Ni0,NiAl_0.,CrS, Ni-Al GE Codeo NiO.Ni_Al.Ni J XRD results listed from strongest to weakest peaks. 222

the pack/substrate arrangement during the corrosion process are presented in Figs. 71-73. Samples which appeared severely corroded were removed from the test. This is indicated by a premature stoppage of the kinetic measurements (Figs. 71-73). XRD results of the corrosion products are listed in Tables 32 and 33 for alloys coated in a powder contacting and "above pack" arrangement, respectively. Chromium/RE- modified aluminide coatings were also compared to a commercial, low-activity coating, GE Codep C, for the same alloy substrate. The GE Codep process is basically an aluminizing treatment which forms an outward-grown,

hypostoichiometric 6 -NiAl coating with very limited

chromium enrichment, only 2 . 2 at.%. For the AP arrangement, only one treatment was observed to provide any substantial resistance to hot corrosion attack. Corrosion attack was characterized by an initial incubation period, where either a protective oxide or some other corrosion product was formed. If the oxide was fluxed by the molten salt, the coating was then exposed to further attack. Repeated dissolution/ precipitation reactions as suggested by Rapp and Goto[100], leave the coatings no longer able to form a protective scale, so that coating failure occurs. Instead, alloy sulfidation occurs and less protective 223 basic oxides form (Table 32) which are acidically attacked by the molten salt (PgQg= 1.73*10 ^ atm at 12Q0K). For the PC arrangement, several treatments provided adequate resistance to fused salt attack, except for ZrCl^-activated packs. Yttrium-doped and undoped, Cr-enriched aluminide coatings were protective up to 360 hours. The coatings produced from YCl^/CrClg-activated packs survived 672 hours with limited attack, forming a more protective AlgO^ and NiAlgO^ oxide phases (Table 33). Substrates coated with ZrCl^-activated packs were severely corroded within the first 24 hours of exposure, forming a porous NiO scale. To interpret the mechanism of hot corrosion degradation, cross-sectional electron micrographs and the corresponding elemental x-ray maps of the coating/corrosion product couple were made, as are

presented in Figs. 74-76. Little attack of the 6 -NiAl coating was observed after 144 hours of exposure of René 80 alloy substrates treated at 1150°C for 24 hours in a pack containing 2 wt.% activator and 25 wt.% Cr-7.5A1 masteralloy using a PC arrangement. WDS and XRD analyses detected a thin, external AlgOg scale which protected the coating from the molten salt. Cracks in the ZrCl^-activated coating (Fig. 74c and d) were also evident. Microprobe analyses did not detect any 224

0 Map 'I 0 Mwp 1 0 0 .. Q K U D, S U P S 1 5 0 k-VV 0 I 1 0 0. 0 H W D S / WDS" 1 5 V 0 k U

Figure 74: (a),(c) Cross-sectional backscattering electron micrograph and (b),(d) corresponding oxygen x-ray map of a René 80 alloy substrate treated at 1150 C for 24 hours in a 2 wt.% (a) YC1_- or (c) ZrCl.-activated pack containing 25 wt.% Cr-7.5A1 masteralloy and isothermally corroded at 900°C for 144 hours in a 0.1% SOu/O. gas mixture with 5.0 mg/cm Na SO . Powder contacting arrangement; 300x. 225

* 2 H I M A P

^ O-MiAl , 0

Figure 75: Cross-sectional backscattering electron micrograph and corresponding x-ray maps of nickel, aluminum, chromium, oxygen, sulfur, cobalt, tungsten, and molybdenum of a René 80 alloy substrate treated at 1150 C for 24 hours in a 2 wt.% (2:1) YCl/CrCl -activated pack containing 25 wt.% Cr-7.5A1 masteralloy and isothermally corroded at 900 C for 672 hours in a 0.1% SOU/O. gas mixture with 5.0 mg/cm Na„SO.. Powder contact]:ing arrangement; 700x. 226

ü r *

Figure 76; Cross-sectional backscattering electron micrograph and corresponding x-ray maps of nickel, aluminum, chromium, oxygen, sulfur, cobalt, tungsten, and molybdenum for a René 80 alloy substrate treated for 24 hours in a 2 wt.% ZrCl.-activated pack containing 25 wt.% Cr-7.5A1 masteralloy and isothermally corroded at 900 Ç for 672 hours in a 0.1% SOu/O. gas mixture with 5.0 mg/cm Na^SO^. Powder contacting arrangement; 700x. 227 significant reduction in the Cr surface composition of either coating after 144 hours of exposure. Significant attack was observed after 672 hours of exposure. The surfaces of both coatings were severely pitted by the oxidation/corrosion process (Fig. 75 and 76). For YClg/CrClg-activated coatings, a thin layer of NigAl was observed at the surface of the B-NiAl coating. This transformation occurred from the gradual depletion of A1 which occurred during the oxidation/dissolution/ precipitation reactions of the corrosion process. A thick, porous AlgO^ scale was also detected by WDS and XRD analyses. The interdiffusion zone consisted of coarsened,

M 2 3 C6-type carbides[89] rich in Cr, Mo, and W (Fig. 75). For ZrCl^-activated coatings, the surface was completely transformed into a two-phase region containing y'-Ni3Al and X-Ni with a dispersion of chromium-rich internal sulfides, probably CrS (Fig. 76). A thick, unprotective NiO and NiAlgO^ scale was detected by WDS and XRD analyses. In addition, a carbide-free interdiffusion zone contained only Cr-, Co-, W- and Mo-rich, sigma phase precipitates. Microprobe analysis also indicated a significant reduction in the Cr surface composition, 0.96 and 2.13 at.%, for the ZrCl^- and YClg/CrClg-activated coatings, respectively. These reductions are mainly due to the formation of internal sulfides or an external 228

chromium-rich product scale, but no sulfides were detected in the YClg/CrClg-activated coatings. Generally, "cleaner" coatings produced in the AP arrangement provided less protection from fused salt attack than the PC coatings. The AP samples suffered similar attack to that of the commercial, low-activity coatings produced with the GE Codep process with no additional chromium enrichment (Fig. 72). The overall effectiveness of aluminide coatings seems to be governed by the Cr surface composition. AP coatings consisted of

two phases, 6 -NiAl and a-Cr, but the outer region was a single-phase, B-NiAl layer with limited Cr enrichment (e.g., 4-5 at.%). PC coatings also consisted of these two phases throughout the outer layer with considerably more chromium enrichment (e.g., 8-13 at.%). At steady-state, Cr/RE-modified aluminide diffusion coatings produced an external AlgO^ scale not a Cr^O^ scale. Therefore, the enhanced resistance to hot corrosion attack must be attributed to a possible modification of the Al^O^ scale. Chromium enrichment in aluminide diffusion coatings[14] or Cr additions to B-NiAl compounds[147] significantly improve the resistance to acidic hot corrosion attack. But no specific mechanism was previously proposed and no external chromium-rich oxide was detected. In a separate investigation, Santoro et al.[26] determined 229

the oxidation behavior of Ni-(50-x)Al-xCr alloys in air at 1100°C where x=0,l,3, or 10 at.% additions of chromium. After 100 hours of isothermal oxidation, XRD analysis detected AlgOg and NiCrgO^ spinel phases. After cyclic oxidation in static air at 1100°C for 100 one-hour cycles, no chromium-rich oxide was formed; instead AlgOg and NiAlgO^ spinel phases were detected by XRD analysis. Therefore, as proposed in Section 5.3.1, a transient Cr-

rich oxides such as a-(Al,Cr)2 0 g and Ni(Al,Cr)2 0 ^ may form during exposure. This transient scale provides enhanced resistance to acidic dissolution because the CrO^^ solute

buffers the Na2 S0 ^ melt, preventing localized basic dissolution. As exposure continues, chromium and aluminum are depleted from the coating surface to reform or grow this protective scale. Eventually, a protective Cr- and Al-rich scale can no longer be formed, whereby a less protective spinel or AI2 O3 scale is produced. The molten salt can flux this scale more easily than the transient scale and penetrate to the underlying coating surface. The salt then "cathodically overpolarizes" the coating surface forming metal sulfides (e.g., CrS) beneath the external surface of the coating, further depleting chromium (Fig. 76). The unprotective oxide scale continuously dissolves and reprecipitates in the molten salt (Fig. 22), depleting the surface of aluminum and 230

Table 34: Summary of results for a René 80 alloy substrate treated at 1150 C for 16 hours in a pack containing 2 wt.% activator and 15 wt.% masteralloy and tested in a mach 1, burner rig at 927 C using JP-5 jet fuel with 5 ppm NaCl and 0.4% sulfur added.

Ave. Surf. Comp. Coating Pack Chemistry Chromium(at.%) Lifetimes(hrs) GE Codep B 4.91 572 NH Cl, CrlOAl 2.16 381 ZrCl , Cr7.5Al 4. 64 327 YCl f 8.42 671 YCl /CrCl , 8.63 671 231

transforming the coating into Ni^Al and y - N i . Finally, the coating can no longer provide protection to the substrate and fails. Burner rig tests; Burner rig tests were conducted at the Environmental Testing Laboratory of the GE Aircraft Division, Evendale, Ohio. A mach 1, burner rig was used at 927°C using JP-5 jet fuel with 5 ppm NaCl and 0.4% sulfur added. René 80 alloy substrates treated at 1150°C for 16 hours in a pack containing 2 wt.% activator and 25 wt.% Cr-7.5A1 masteralloy were tested and compared with a commercial, high-activity aluminide coating, GE Codep B, on the same alloy substrate. Coatings were treated for 16 hours so that comparison of coatings with comparable thicknesses could be made. The coatings were objectively evaluated by the GE rig technician. The coatings were tested until the technician observed the formation of surface blisters, indicating the formation of unprotective scales (e.g., NiO, NiCCgO^, CrS, etc.) and overall coating failure. The coating lifetime values and the average chromium surface composition (EPMA) are listed in Table 34. Coating lifetimes increase with increasing chromium surface composition. YCl^- and YClg/CrClg-activated coatings outlasted the commercial, high-activity aluminide coating which commonly applied to hardware used in the combustion zone of a gas turbine. 232

Microprobe analysis of corroded cross-sections detected the presence of internal sulfides in a Ni^Al surface layer, indicative of an unprotective coating. CHAPTER VI SDHKARY AND CONCLUSIONS

1. A single-step, pack cementation treatment for producing a chromium/RE(Y,Zr)-modified and a RE(Y,Zr,Hf,Si)-doped aluminide diffusion coating on commercial nickel-base superalloys (e.g., IN 713LC, Mar-M247, René 80, René BOH, and René N4) has been developed. These coatings were produced from a pack mixture containing either an NH^Cl activator plus a RE oxide source (e.g., YgO^, ZrOg, or SiOg) or else a RE-base activator (e.g., YCl^, ZrCl^, HfCl^) with a suitable Cr-Al binary masteralloy at 1150°C for 24 hours.

2. The chromium/RE-modified, aluminide coatings consisted of two phases: an outward-grown, hypostoichiometric 6 -NiAl matrix with a dispersion of a-Cr second-phase particles which both precipitate near the original substrate surface and which grow inward from the gas phase along grain boundaries at the coating surface. Electron microprobe and XFS analyses measured both a

233 234 substantial enrichment of chromium (solid-solution and a-cr second-phase particles) and a significant amount of the RE dissolved in the coating surface.

3. The RE(Y,Zr,Hf)-doped, aluminide coatings also consisted of two phases: an outward-grown, hypostoichiometric 6 -NiAl matrix with a dispersion of a-Cr second-phase particles which only precipitate near the original substrate surface from interdiffusion. Electron microprobe and SIMS analyses detected a significant amount of the RE dissolved in the B-NiAl coating and in the form of a yttrium-rich precipitate (e.g., Ni^Y or Ni^Y) at the B-NiAl/a-Cr interphase boundary.

4. The silicon-doped, aluminide coatings were produced from a 4 wt.% NH^Cl-activated pack containing a combination of 20 wt.% Cr-20A1 and 5 wt.% Cr-40Si binary masteralloys treated at 1150°C for 24 hours. The coatings consisted of one phase: an outward-grown, hyperstoichiometric 6 -NiAl matrix with little chromium enrichment and about 0.218+.05 at.% silicon dissolved in solution. These levels are far below optimum values for superior resistance to hot corrosion attack. 2 35

5. Two pack/substrate arrangements were used to produce the new modified aluminide coatings: a traditional, powder contacting (Fig. 27) and an "above pack" arrangement (Fig. 28). Unfortunately, outward-grown aluminide coatings, produced from the powder contacting arrangement, were embedded with AlgOg and RE oxide pack particles. Physical separation of substrates (Fig. 28) from the powder mixture eliminated pack powder entrapment and reduced coating thicknesses and chromium enrichment of the surface.

6 . Physical isolation significantly reduced the growth rates of aluminide coatings. In fact, the apparent activation energy for the coating process between 1050-1150°C was calculated as 27+3 and 24+5 kcal/mol from weight and thickness changes, respectively. These values are about 50% less than ones for the solid-state interdiffusion of nickel and aluminum in the B-NiAl phase. Since the overall process is controlled by the series diffusion in both the gas phase and the solid-state, these results indicate that the process is influenced more by gaseous diffusion.

7. ITSOL equilibrium calculations predicted that the RE oxide sources were converted by the AlCl^ species to their 236

volatile chloride species (e.g., RECl^) in the high chlorine activity pack mixture producing a condensed aluminum oxide product. Mass spectrometer measurements substantiated these predictions; identifying RECl^ species, measuring a reduction in the intensity of the AlCl^ species, and detecting (XRD) a condensed aluminum oxide product (e.g., y-AlgO^).

8 . Pack powder entrapment and its volume fraction contributed considerably to the overall performance of aluminide diffusion coatings during cyclic oxidation testing. Coatings produced fromthe powder contacting arrangement suffered severe degradation from thermal cycles. With the use of the "above pack" arrangement, pack powder entrapment was eliminated and pack chemistries were identified which improved resistance to cyclic oxidation, producing adherent and protective scales of AlgOg and NiAlgO^ spinel for 500, one-hour cycles at 1100°C in static air.

9. RE-doped, aluminide diffusion coatings formed a continuous, slow-growing AlgO^ scale during isothermal oxidation at 1100°C in air for 44 hours. The parabolic rate constant, k^,measured from the oxidation of RE-doped aluminide coatings were less than RE-free aluminide 237 coatings. RE-doped aluminide coatings were characterized by an outer, ridged AlgO^ scale which was cracked near coating grain boundaries exposing an inner, compact AlgO^ scale rich in the RE. Whereas, RE-free aluminide coatings exhibited large fluctuations in the oxidation kinetics, indicating the periodic cracking of the ridged oxide scale exposing the underlying coating surface. Oxidation rates agreed well with RE-doped, 6 -NiAl compounds.

10. The chromium/RE-modified aluminide coatings, produced from the powder contacting arrangement, provided sufficiently better resistance to hot corrosion attack (e.g., thin film and burner rig tests) than commercial low-activity and high-activity, aluminide coatings (GE Codep B and C, respectively). Coating lifetimes are strongly dependent on the chromium surface composition needed to form a transient aluminum- and chromium-rich oxide which better resists dissolution by the molten salt. Although physical isolation eliminates powder entrapment, the increased diffusion distance reduced the chromium flux and the chromium surface composition. Thus, the resistance to fused salt attack was substantially decreased. CHAPTER VII FDTDRE WORK

1. The cyclic oxidation behavior of the RE-doped, aluminide coatings will be evaluated. Cyclic oxidation testing at 1100°C in static air is currently in progress at the NASA Lewis Research Center. Weight change measurements, XRD analysis, and electron microscopy will be employed to determine optimum oxidation-resistant coatings.

2. The chromium-enriched, aluminide diffusion coatings, produced using an "above pack" arrangement, will be submitted for burner rig analysis in the VAMAS test program at the National Physical Laboratory, Middlesex, United Kingdom.

238 APPENDIX A ITSOL CALCULATIONS: CONSTANT PRESSURE VERSUS CONSTANT VOLUME

239 240

Table 35: ITSOL pack equilibrium output for a 2 wt.% NH-Cl-activated pack containing 25 wt.% Cr-lOAl wt.% masteralloy at 1150 C (constant pressure).

ECUILIFRIUH COr,?üS:i:ON5:

SPECIE IN IT.EST. EQ.MOLES P/ATM ALT!'. ITT HE .COOOOEtOO .386142-01 .839632*0; .038632*0':' A 1 0 3 .OOOOOE'OO .5 9 5 9 8 2 -0 2 .129442*00 .:2944E.(,0 Hc; .1943CE-01 .480902-03 .104442-01 .1044»S-.)1 AIC12 .000002*00 .454172-03 .986382-02 .9 8 -3 2 2 -.E Ar .40000E -03 .400002-03 .868722-02 .868722-12 AICl .000002*00 .113972-03 .247522-02 .247522-02 CrC12 .000002*00 .196032-04 .425752-03 .4 2 5 7 5 2 -0 - A12CU .000002*00 .1 3 7 4 5 2 -0 5 .298512-04 .2 9 8512-04 H .000002*00 .2 8 2 32 2 -0 6 .6 1 3 15 2 -0 5 .613152-05 Cr/fll CoiiBoo5iti:n, ;!H4E1,;423): Cr .000002*00 .3 5 2 0 4 2 -0 8 .7 6 4562-07 .7 6 4 56 2 -0 7 STEP NO: 1 Cl .000002*00 .2 8 7 6 2 2 -0 8 .6 2 4 65 2 -0 7 .6 2 4652-07 CrC13 .000002*00 .279592-08 .6 0 7 22 2 -0 7 .6 0 7 22 2 -0 7 T = 1423.00 K 112 .000002*00 .1 9 5 52 2 -0 8 .4 2 4 62 2 -0 7 .4 24622-07 P = I.O00E«00 Aid AIM .000002*00 .133502-08 .289952-07 .2 8 9952-07 A1 .000002*00 .854672-09 .1 85622-07 .1 8 5 62 2 -0 7 ELEMENT MCIEE NH3 .259032-01 .584032-09 .1 2 6852-07 .1 2 6 85 2 -0 7 Ar .OOC4 A1H2 .000002*00 .1 8 1 8 2 2 -1 0 .3 9 4982-09 .3 94882-09 N .0 1 9 ‘ C12 .000002*00 .1 5 6 1 1 2 -1 2 .339042-11 .339042-11 H .0777 A1H3 .000002*00 .556752-13 .120922-11 .120922-11 AT .(•4tO HH2 .000002*00 .8 7 8 5 7 2 -1 4 .190812-12 .190812-12 Cr .2160 CrCI4 .000002*00 .803832-14 .174582-12 .1 74582-12 Cl .0194 NH .000002*00 .145342-17 .315642-16 .3 1 5 64 2 -1 6 CrN .000002*00 .561942-18 .126392-16 .1 2 6 39 2 -1 6 N .OOOOOE+00 .734432-19 .159502-17 .1 5 9 50 2 -1 7 AIM .000002*00 .291012-21 .632022-20 .632022-20 N2H2 .000002*00 .319432-22 .693752-21 .693752-21 N2H4 .000002*00 .628682-24 .136542-22 .1 3 6542-22

INVARIANT CONDENSED PHASES SPECIE IN IT.EST. EQ.M8LES Cr .214002*00 .215982*00 AI .460002-01 .200392-01 AIN .000002*00 .1 9 4 30 2 -0 1 CrN .000002*00 .000002*00 CrEN .000002*00 .000002*00 CrCIE .000002*00 .000002*00 2 41

Tcible 36: ITSOL pack equilibrium output for a 2 wt. NH.Cl-activated pack containing 25 wt.% Cr-lOAl masteralloy at 1150 C (constant volume).

E3üILI5S!'jr. COnPOSniOHE;

SPECIE INIT.EST. 2C.N0LE: P/ATI ACT 171 i l H3 .000002*00 .387*42-01 .9 )5 2 8 2 * 0 2 .905282*02 «1C13 .000002*00 .60 3 13 2 -0 2 .140922*02 .140922-02 P r .400002-03 .400002-03 :93463Ë ;Ô 0 .934632*01 HSl .1 9 4 20 2 -0 ! .22 1 76 2 -0 3 .5 1 8 1 6 2 -0 0 .518162*00 P12CI6 .000002*00 .1 5 1 44 2 -0 3 .353862*00 .352862*00 P1C12 .0 )0002*00 .962552-04 .224912*00 .2 2 4 9:2 * 0 0 AlCI .000002*00 .5 0 5 042-05 .1 1 6 19 2 -0 1 .113192-01 Cr.Al Codeposition at !423K, 90Crl0fll CrClH .000002*00 .4 1 5 46 2 -0 5 .9 7 0 7 5 2 -0 2 .970752-02 STEP NO: 1 H .000092*00 .2 7 2 64 2 -0 7 7 637052-04 .637052-04 NH3 .1 9 4 30 2 -0 ! .60BEE2-0S .1 4 2 2 7 2 -0 4 .142272-04 T = ! « 3 . 0 0 K C rS ia .000002*00 .2 8 2 94 2 -0 2 .6 6 1 1 2 2 -0 5 .661122-05 P = 1.0b7E«03 ATM PlH .000002*00 .12 2 93 2 -0 9 .3 0 1 2 5 2 -0 6 .30 1 25 2 -0 6 Cl .000002*00 .127652-79 .298272-06 .298272-06 V = 5.000E-02 L Cr .000002*00 .32 7 22 2 -1 0 .7 6 4 5 6 2 -0 7 .764562-07 A1N2 .000002*00 .1 6 2 43 2 -1 0 .4 2 6 2 6 2 -0 7 .426262-07 ELENEHT MOLES N2 .000002*00 .1 2 1 73 2 -1 0 .4 2 4 6 2 2 -0 7 .424622-07 Ar .0004 A1 .000002*00 .794402-11 .1 8 5 6 2 2 -0 7 .185622-07 Cr .2 1 6 0 PIN3 .000002*00 .5 8 2 45 2 -1 2 .1 2 6 0 9 2 -0 8 .1 3 6 592-08 PI .0 4 6 0 CrCK .000002*00 .32 8 43 2 -1 3 .9 0 7 6 0 2 -1 0 .907602-10 N .0194 C12 .000002*00 .3 3 0 84 2 -1 3 .7 7 3 0 4 2 -1 0 .7 7 3 042-10 N .0 7 7 7 HH2 .000002*00 .881522-14 .205972-10 .2 (5 9 7 2 -1 0 Cl .0194 KH .000002*00 .1 3 9 86 2 -1 8 .3 2 6 7 9 2 -1 5 .3 2 6 792-15 CrN .000002*00 .5 4 0 90 2 -2 0 .1 2 6 3 9 2 -1 6 .126392-16 N .000002*00 .62 2 65 2 -2 1 .1 5 9 5 0 2 -1 7 .159502-17 N2H4 .000002*00 .6 8 0 93 2 -2 2 .1 5 9 1 0 2 -1 8 .159102-18 N2H2 .000002*00 .320502-22 .7 4 8 8 8 2 -1 9 .748882-19 AIN .000002*00 .2 7 0 49 2 -2 3 .6 3 2 0 2 2 -2 0 .632022-20

INVARIANT CONDENSED PHASES SPECIE INIT.EST. 2 0 .HOLES Cr .216002*00 .216002*00 A1 .460002-01 .201352-01 AIN .000002*00 .194302-01 CrCI2 .000002*00 .000002*00 Cr2N .000002*00 .000002*00 CrN .000002*00 .000002*00 REFERENCES

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6 . F.L. VerSnyder and M.E. Shank, Mat. Sci. Ena.. 6 (1970) 213. 7. S.A. Bradford, "Fundamentals of Corrosion in Gases", in Metals Handbook Volume 13 : Corrosion. Ninth Edition, ASM INTERNATIONAL, Metals Park, OH (1987) 61.

8 . O. Kubaschewski and B.E. Hopkins, Oxidation of Metals and Alleys. Second Edition, Butterworth and Co., London (1962) 205. 9. S.R.J. Saunders and J.R. Nicholls, Mater. Sci. Techn.. 5 (1989) 780. 10. Y.M. Kim and G.R. Belton, Metall. Trans.. 5A (1974) 1811.

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