investigation of the effect of hafnium and chromium additions on the ...

INVESTIGATION OF THE EFFECT OF HAFNIUM AND CHROMIUM ADDITIONS ON 

THE MICROSTRUCTURES AND SHORT-TERM OXIDATION PROPERTIES OF DC 

MAGNETRON SPUTTERED β-NIAL BOND COATS DEPOSITED ON NI-BASED 

SUPERALLOYS 

by 

MICHAEL A. BESTOR 

A DISSERTATION 

Submitted in partial fulfillment <strong>of</strong> <strong>the</strong> requirements 

for <strong>the</strong> degree <strong>of</strong> Doctor <strong>of</strong> Philosophy 

in <strong>the</strong> Department <strong>of</strong> Metallurgical <strong>and</strong> Materials Engineering 

in <strong>the</strong> Graduate School <strong>of</strong> 

The University <strong>of</strong> Alabama 

TUSCALOOSA, ALABAMA 

2010

Copyright Michael Alan Bestor 2010 

ALL RIGHTS RESERVED

ABSTRACT 

Thermal barrier coatings play a major role in protecting turbine blades from extreme 

operating environments <strong>and</strong> extending service lifetimes. Reactive elements (e.g. Zr, Hf, Y, Si, 

etc.) have been shown to enhance <strong>the</strong> oxidation performance <strong>of</strong> such coating systems when 

added in appropriate amounts to overlay bond coats. This study investigated <strong>the</strong> influences <strong>of</strong> 

processing parameters along with Hf <strong>and</strong> Cr <strong>additions</strong> on <strong>the</strong> short-term oxidation performance 

<strong>of</strong> β-NiAl based coatings deposited via direct current magnetron sputtering onto CMSX-4 <strong>and</strong> 

René N5 substrates. Sputtering parameters were optimized to yield a zone T microstructure. 

The results indicate that <strong>the</strong> coatings are deposited as a solid solution <strong>and</strong> precipitation with <strong>the</strong> 

RE-doped coatings occur following annealing at 1000°C for times up to four hours. Precipitates 

form heterogeneously within <strong>the</strong> coatings with larger precipitates forming at grain boundaries 

<strong>and</strong> smaller ones forming within <strong>the</strong> grains at prior dislocation lines. Small incorporations <strong>of</strong> Cr 

into <strong>the</strong> NiAl-1Hf coating increased <strong>the</strong> average grain size <strong>and</strong> precipitate size. Transmission 

electron microscopy <strong>and</strong> atom probe tomography confirmed that <strong>the</strong> chemistry <strong>of</strong> <strong>the</strong> precipitates 

is mostly β’-Ni 2 AlHf accompanied by HfC <strong>and</strong> α-Cr. 

The results from <strong>the</strong> iso<strong>the</strong>rmal oxidation studies at 1050°C indicated that increasing <strong>the</strong> 

preoxidation annealing time from two to four hours decreased <strong>the</strong> mass gains observed with <strong>the</strong> 

specimens up to 100 hours. However, significant oxidation at <strong>the</strong> coating/substrate interface was 

discovered <strong>and</strong> <strong>the</strong> amount <strong>of</strong> oxidation increased with Hf content <strong>and</strong> preoxidation annealing 

time. This oxidation is thought to be caused by <strong>the</strong> large number <strong>of</strong> pinhole defects with <strong>the</strong> 

ii

zone T microstructure <strong>and</strong> large grain boundary volume. Increased Hf concentrations were also 

found at <strong>the</strong> coating/substrate interface <strong>and</strong> this has been shown to lead to dramatic internal 

oxidation. The NiAlCrHf samples contain larger grain sizes <strong>and</strong> precipitates <strong>and</strong> a thinner TGO 

than <strong>the</strong> NiAl-Hf coatings. This combined with <strong>the</strong> lower mass gains during iso<strong>the</strong>rmal 

oxidation indicate that <strong>the</strong> NiAlCrHf coatings rapidly form a thin, protective α-Al 2 O 3 layer that 

limits additional transport <strong>of</strong> oxygen to <strong>the</strong> bond coat. The results have been analyzed <strong>and</strong> 

discussed relative to previous research on sputter deposited NiAl-Hf coatings. 

iii

LIST OF ABBREVIATIONS 

Abbreviations 

APT 

atom probe tomography 

b. c. c. body centered cubic 

CVD 

DC 

EB-PVD 

EDS 

EPMA 

FEG 

HAADF 

IDZ 

LEAP 

PVD 

RE 

RF 

SEM 

STEM 

TBC 

TEM 

TGO 

chemical vapor deposition 

direct current 

electron-beam physical vapor deposition 

energy dispersive X-ray spectroscopy 

electron probe microanalysis 

field emission gun 

high angle annular dark field 

interdiffusion zone 

local electrode atom probe 

physical vapor deposition 

reactive element 

radio frequency 

scanning electron microscopy 

scanning transmission electron microscopy 

<strong>the</strong>rmal barrier coatings 

transmission electron microscopy 

<strong>the</strong>rmal grown oxide 

iv

WDS 

XRD 

YSZ 

wavelength dispersive X-ray spectroscopy 

X-ray diffraction 

yttria-stabilized zirconia 

Symbols 

a o 

lattice parameter 

at.% 

β 

atomic percent 

NiAl 

β′ Ni 2 AlHf 

d 111 

spacing between NiAl (111) planes 

γ′ Ni 3 Al 

K p 

parabolic constant 

λ 

wavelength 

v

ACKNOWLEDGEMENTS 

With any project, <strong>the</strong>re are always so many that assist us in many ways that make <strong>the</strong> job 

at h<strong>and</strong> easier. My dissertation research was definitely no exception. In my six years at The 

University <strong>of</strong> Alabama, I have made many new friends <strong>and</strong> had <strong>the</strong> opportunity to learn more 

about life through <strong>the</strong>m. First, I would like to thank Drs. Mark Weaver <strong>and</strong> Alan Lane for 

serving as my advisors during my research efforts with <strong>the</strong> university. My badges <strong>of</strong> honor are 

extended to James Hill, Kim Clary, Dr. Lawrence Hill, Ken Dunn, Bob Fanning, Rich Martens, 

Dr. Michael Bersch, Johnny Goodwin, Anne Brasher, Jan Crietz, <strong>and</strong> Lyndall Wilson as <strong>the</strong>y are 

my “unsung heroes” with all <strong>of</strong> <strong>the</strong>ir assistance <strong>and</strong> making my transition from Mississippi State 

University <strong>and</strong> continued research a pleasurable experience. The support <strong>of</strong> every one <strong>of</strong> my 

colleagues: Dr. Jair Lizarazo-Adarme, Dr. Patrick Henry, Dr. Xiao Li, Dr. Shelby Shuler, Dr. 

Michael Ivie, Arthur Brown, Patrick Coleman, Jason Morgan, Mark Calhoun, Alex Dues, 

Prentice Singleton, <strong>and</strong> Robb Morris is greatly acknowledged for all <strong>of</strong> <strong>the</strong> help <strong>and</strong> discussions 

through <strong>the</strong> years. 

I would like to express special gratitude from my NASA sponsors at Marshall Space 

Flight Center: Tim Vaughn, Greg Jerman, James Coston, Dr. Binayak P<strong>and</strong>a, Vanessa Bailey, 

<strong>and</strong> Bertha Gildon. These individuals gave me <strong>the</strong> opportunity to join <strong>the</strong> NASA family <strong>and</strong> 

provided me access to facilities <strong>and</strong> advice that allowed me to develop in new ways that 

o<strong>the</strong>rwise would not have been possible. My hat is <strong>of</strong>f to <strong>the</strong> NASA family. Thank you for 

making my time with you such a wonderful memory! 

vi

Finally, I would like to express my deepest respect for my family. With <strong>the</strong> end <strong>of</strong> 

ano<strong>the</strong>r chapter in my life coming, it would be selfish to imagine accomplishing all that I have 

without <strong>the</strong> unconditional support from Mama, Daddy, Andy, Laura Beth, Jessie, Granny, <strong>and</strong> so 

many o<strong>the</strong>rs that have <strong>of</strong>fered <strong>the</strong>ir advice <strong>and</strong> support. Their teachings have kept me grounded 

<strong>and</strong> remind me to treat everyone with <strong>the</strong> same respect that I strive to obtain in life. 

“We must become <strong>the</strong> change we want to see in <strong>the</strong> world.” 

-G<strong>and</strong>hi 

“O God your sea is so great <strong>and</strong> my boat is so small” 

“Die when I may, I want it said <strong>of</strong> me by those who knew me best, that I always plucked a thistle 

<strong>and</strong> planted a flower where I thought a flower would grow.” 

-Abraham Lincoln 

vii

TABLE OF CONTENTS 

ABSTRACT.................................................................................................................................... ii 

LIST OF ABBREVIATIONS........................................................................................................ iv 

ACKNOWLEDGEMENTS........................................................................................................... vi 

LIST OF TABLES........................................................................................................................ xii 

LIST OF FIGURES ..................................................................................................................... xiii 

CHAPTER I GENERAL INTRODUCTION .................................................................................1 

1.1. General Background ........................................................................................................1 

1.2. Motivation.............................................................................................................................2 

1.3. Dissertation Layout...............................................................................................................2 

CHAPTER II LITERATURE REVIEW 3 

2.1. Thermal Barrier Coatings .....................................................................................................3 

2.1.1. Construction <strong>and</strong> Manufacture <strong>of</strong> TBC Layers 3 

2.1.2. Failure Mechanism <strong>of</strong> TBC system 4 

2.2. Properties <strong>of</strong> NiAl.................................................................................................................7 

2.2.1. Phase Equilibria <strong>and</strong> Crystal Structure 7 

2.2.2. Influence <strong>of</strong> Ternary Alloying Additions 8 

2.1.2.1. Mechanical Properties 9 

2.1.2.2. Oxidation Resistance 10 

2.1.3. NiAl alloys as TBC bond coats 13 

2.1.4. High strength NiAl alloys as TBC bond coats 16 

CHAPTER III EXPERIMENTAL PROCEDURE 18 

viii

3.1. Sample Preparation.............................................................................................................18 

3.1.1. Target <strong>and</strong> substrate preparation 18 

3.1.2. Deposition <strong>of</strong> coatings 19 

3.1.3. Structural <strong>and</strong> Chemical Characterization 20 

3.1.4. Heat Treatment 22 

3.2. Iso<strong>the</strong>rmal <strong>and</strong> Cyclic Oxidation Tests...............................................................................22 

CHAPTER IV INFLUENCES OF ANNEALING AND HAFNIUM CONCENTRATION ON 

THE MICROSTRUCTURES OF SPUTTER DEPOSITED β-NIAL COATINGS ON 

SUPERALLOY SUBSTRATES 23 

4.1. Introduction.........................................................................................................................24 

4.2. Experimental.......................................................................................................................25 

4.3. Results <strong>and</strong> Discussion .......................................................................................................27 

4.3.1. As-Deposited Coatings 27 

4.3.2. Annealed Coatings 29 

4.3.3. Three-Dimensional Atom Probe Tomography (3D-APT) 34 

4.4. Conclusions.........................................................................................................................36 

Acknowledgements....................................................................................................................36 

References..................................................................................................................................37 

CHAPTER V INFLUENCES OF CHROMIUM AND HAFNIUM ADDITIONS ON THE 

MICROSTRUCTURES OF β-NIAL COATINGS DEPOSITED ON SUPERALLOY 

SUBSTRATES 56 

5.1. Introduction.........................................................................................................................57 

5.2. Experimental.......................................................................................................................59 

ix

5.3. Results <strong>and</strong> Discussion ......................................................................................................60 

5.3.1. Microstructure 60 

5.3.2. Chemistry 62 

5.4. Conclusions.........................................................................................................................66 

Acknowledgements....................................................................................................................67 

References..................................................................................................................................68 

CHAPTER VI MORPHOLOGICAL AND CHEMICAL EVOLUTION OF NIAL, NIAL-HF, 

<strong>and</strong> NIAL-CR-HF BOND COATS DURING SHORT TERM ISOTHERMAL OXIDATION82 

Abstract:.....................................................................................................................................82 

6.1 Introduction:.........................................................................................................................83 

6.2. Experimental Procedures:...................................................................................................85 

6.3. Results <strong>and</strong> Discussion: ......................................................................................................87 

6.3.1. Characterization <strong>of</strong> as-deposited <strong>and</strong> annealed coatings 87 

6.4. Iso<strong>the</strong>rmal oxidation behavior <strong>and</strong> microstructures............................................................87 

6.4.1. Oxidation test results 87 

6.4.2. Microstructural changes during oxidation 89 

6.5. Elemental Mapping <strong>of</strong> Coatings Iso<strong>the</strong>rmally Oxidized for 96 hours................................96 

6.6. TEM Analysis <strong>of</strong> Oxidized Coatings..................................................................................98 

6.7. General Discussion .............................................................................................................99 

6.8. Summary...........................................................................................................................102 

Acknowledgements..................................................................................................................106 

x

CHAPTER VI FACTORS INFLUENCING INTERDIFFUSION BETWEEN SPUTTERED 

NIAL OVERLAY COATINGS AND Ni-BASED SUPERALLOY SUBSTRATES 

SUBJECTED TO HIGH TEMPERATURES 137 

7.1 INTRODUCTION ............................................................................................................137 

7.2 DISCUSSION OF RESULTS ...........................................................................................138 

7.3 REFERENCES .................................................................................................................141 

CHAPTER VIII SUMMARY AND CONCLUSIONS 144 

CHAPTER IX FUTURE WORK 148 

REFERENCES 149 

APPENDIX A METALLOGRAPHIC SPECIMEN PREPARATION 162 

APPENDIX B TEM SAMPLE PREPARATION USING A FOCUSED ION BEAM 163 

APPENDIX C APT SAMPLE PREPARATION USING A FOCUSED ION BEAM 168 

xi

LIST OF TABLES 

CHAPTER III 

Table 3.1. Nominal Target Chemical Compositions in Atomic Percent ......................................19 

Table 3.2. Chemical compositions <strong>of</strong> CMSX-4 ® <strong>and</strong> Rene N5 superalloys.................................19 

Table 3.3. Deposition parameters for NiAl-X coatings ................................................................19 

CHAPTER IV 

Table 1. Nominal chemical compositions <strong>of</strong> sputtering targets....................................................43 

Table 2. Lattice parameters for <strong>the</strong> β-NiAl phase.........................................................................43 

CHAPTER V 

Table 1. Nominal chemical compositions <strong>of</strong> sputtering targets....................................................73 

CHAPTER VI 

Table 1. Nominal target concentrations in atomic percent .........................................................109 

Table 2. Nominal chemical compositions <strong>of</strong> coatings after oxidation........................................110 

Table 3. Nominal chemical compositions <strong>of</strong> coatings after oxidation for 96h at 1050°C..........110 

APPENDIX A 

Table A.1. General polishing method for Ni-based superalloys.................................................158 

xii

LIST OF FIGURES 

CHAPTER II 

Fig. 2.1. Schematic representation <strong>of</strong> a multi-layer <strong>the</strong>rmal barrier coating system. .................... 4 

Fig. 2.2. Phase diagram <strong>of</strong> Ni-Al [79]. .......................................................................................... 8 

Fig. 2.3. NiAl B2 (CsCl) structure................................................................................................. 9 

Fig. 2.4. Ni-Al-Hf ternary phase diagram [84]. ........................................................................... 11 

Fig. 2.5. Crystal structure <strong>of</strong> Heusler phase................................................................................. 12 

CHAPTER IV 

Figure 1. Representative SEM micrographs <strong>of</strong> as-deposited coatings: (a) NiAl <strong>and</strong> (b) NiAl- 

1Hf. ....................................................................................................................................... 44 

Figure 2. Plan view TEM images collected from <strong>the</strong> NiAl coatings: (a) as-deposited, (b) 

annealing at 1000°C for two hours, <strong>and</strong> (c) annealing at 1000°C for four hours. ................ 45 

Figure 3. Elemental composition pr<strong>of</strong>iles for as-deposited coatings: (a) NiAl <strong>and</strong> (b) NiAl-1Hf. 

............................................................................................................................................... 46 

Figure 4. X-ray diffraction patterns for as-deposited coatings. ................................................... 47 

Figure 5. Representative SEM micrographs for: (a) NiAl annealed for two hours, (b) NiAl 

annealed for four hours, (c) NiAl-1Hf annealed for two hours, <strong>and</strong> (d) NiAl-1Hf annealed 

for four hours. ....................................................................................................................... 48 

xiii

Figure 6. Backscattered SEM image <strong>of</strong> NiAl-1Hf showing <strong>the</strong> development <strong>of</strong> <strong>hafnium</strong>-rich 

precipitates after annealing. .................................................................................................. 49 

Figure 7. Plain view TEM images from coatings after annealing for four hours: (a) NiAl <strong>and</strong> (b) 

NiAl-1Hf. .............................................................................................................................. 50 

Figure 8. STEM-HAADF image <strong>of</strong> NiAl-1Hf after annealing for four hours at 1000°C. .......... 51 

Figure 9. Elemental composition pr<strong>of</strong>iles for annealed coatings: (a) Nickel, (b) Chromium, (c) 

Aluminum, <strong>and</strong> (d) Hafnium................................................................................................. 52 

Figure 10. X-ray diffraction patterns for annealed coatings: (a) NiAl annealed two hours, (b) 

NiAl annealed four hours, (c) NiAl-1Hf annealed two hours, <strong>and</strong> (d) NiAl-1Hf annealed 

four hours. ............................................................................................................................. 53 

Figure 11. 3D-APT data for NiAl-1Hf after annealing for four hours: (a) all elements within <strong>the</strong> 

sample <strong>and</strong> (b) only <strong>hafnium</strong> <strong>and</strong> carbon.............................................................................. 54 

Figure 12. (a) Two-dimensional isosurface reconstruction <strong>of</strong> a NiAl-1Hf coating that was 

analyzed using 3D-APT, <strong>and</strong> (b) corresponding concentration pr<strong>of</strong>ile <strong>of</strong> <strong>the</strong> precipitate. ... 55 

CHAPTER V 

Figure 1. Representative SEM micrographs showing <strong>the</strong> as-deposited <strong>and</strong> annealed NiAl-1Hf (a, 

b) <strong>and</strong> NiAlCrHf (c, d) coatings. .......................................................................................... 73 

Figure 2. XRD patterns collected from <strong>the</strong> as-deposited <strong>and</strong> annealed NiAl-1Hf (a) <strong>and</strong> 

NiAlCrHf (b)......................................................................................................................... 74 

Figure 3. Plan view TEM micrographs <strong>of</strong> <strong>the</strong> (a) NiAl-1Hf <strong>and</strong> (b) NiAlCrHf coatings 

following annealing at 1000°C for four hours. ..................................................................... 75 

xiv

Figure 4. EPMA line pr<strong>of</strong>iles collected from <strong>the</strong> as-deposited <strong>and</strong> annealed NiAl-1Hf <strong>and</strong> 

NiAlCrHf coatings showing (a) Al, (b) Cr, (c) Ni, <strong>and</strong> (d) Hf. ............................................ 76 

Figure 5. Cross-sectional SEM-EDS element maps for <strong>the</strong> NiAl-1Hf coating after annealing for 

four hours. ............................................................................................................................. 77 

Figure 6. Cross-sectional SEM-EDS element maps for <strong>the</strong> NiAlCrHf coating after annealing for 

four hours. ............................................................................................................................. 78 

Figure 7. STEM-HAADF image collected from a NiAlCrHf coating that was annealed for four 

hours at 1000°C with EDS spectra <strong>of</strong> β’-Ni 2 AlHf precipitate. ............................................. 79 

Figure 8. 3D-APT data for NiAlCrHf after annealing for four hours at 1000°C: (a) full 

reconstruction displaying all elements in sample <strong>and</strong> (b) 2D-isosurface reconstruction 

showing features that are rich in Cr. ..................................................................................... 80 

Figure 9. Corresponding concentration pr<strong>of</strong>ile for <strong>the</strong> large Cr-rich precipitate in Figure 8b. ... 81 

CHAPTER VI 

Figure 1. Specific mass changes measured with iso<strong>the</strong>rmal oxidation (I) at 1050°C for NiAl-X 

coatings on René N5. Additional samples were subjected to long cycle cyclic oxidation (C) 

at 1050°C. ........................................................................................................................... 115 

Figure 2. X-ray diffraction spectra collected from samples that were iso<strong>the</strong>rmally oxidized at 

1050°C for 96 hours. Samples are divided according to annealing time with two <strong>and</strong> four 

hours shown in parts (a) <strong>and</strong> (b) respectively. .................................................................... 116 

Figure 3. SEM images collected from <strong>the</strong> surface <strong>of</strong> NiAl-0.5Hf samples that were annealed for 

(a) two <strong>and</strong> (b) four hours <strong>and</strong> <strong>the</strong>n oxidized for 96 hours. ................................................ 117 

xv

Figure 4. SEM images collected from <strong>the</strong> surface <strong>of</strong> NiAl-1Hf samples that were annealed for 


Figure 5. SEM images collected from <strong>the</strong> surface <strong>of</strong> NiAlCrHf samples that were annealed for 


Figure 6. Cross-sectional SEM images <strong>of</strong> NiAl-0.5Hf samples: annealed for (a) two <strong>and</strong> (b) four 

hours <strong>and</strong> oxidized for 96 hours.......................................................................................... 120 

Figure 7. Cross-sectional SEM images <strong>of</strong> NiAl-1Hf samples: annealed for (a) two <strong>and</strong> (b) four 


Figure 8. Cross-sectional SEM images <strong>of</strong> NiAlCrHf samples: annealed for (a) two <strong>and</strong> (b) four 


Figure 9. Elemental composition pr<strong>of</strong>ile collected from <strong>the</strong> NiAl-0.5Hf samples that were 

annealed for (a) two hours (b) four hours <strong>and</strong> oxidized for 96 hours. ................................ 123 

Figure 10. Elemental composition pr<strong>of</strong>ile collected from <strong>the</strong> NiAl-1Hf samples that were 


Figure 11. Elemental composition pr<strong>of</strong>ile collected from <strong>the</strong> NiAlCrHf samples that were 


Figure 12. Elemental maps <strong>and</strong> SEM cross-sectional images collected from <strong>the</strong> NiAl-0.5Hf 

samples that were annealed for (a-d) two <strong>and</strong> (e-h) four hours <strong>and</strong> oxidized for 96 hours.126 

Figure 13. Elemental maps <strong>and</strong> SEM cross-sectional images collected from <strong>the</strong> NiAl-1Hf 

samples that were annealed for (a-d) two <strong>and</strong> (e-h) four hours <strong>and</strong> oxidized for 96 hours.128 

Figure 14. Elemental maps <strong>and</strong> SEM cross-sections collected from NiAlCrHf coatings after 

annealing for: two hours (a-e) <strong>and</strong> four hours (f-j) four hours followed by oxidizing for 96 

hours.................................................................................................................................... 130 

xvi

Figure 15. Plan view TEM images <strong>of</strong> NiAl-0.5Hf iso<strong>the</strong>rmally oxidized for 96 hours imaged 

using bright field <strong>and</strong> STEM-HAADF techniques respectively: annealed for (a,b) two <strong>and</strong> 

(c,d) four hours. The boxes in (c) <strong>and</strong> (d) highlight <strong>the</strong> regions imaged in (a) <strong>and</strong> (b). .... 134 

Figure 16. Plan view TEM images <strong>of</strong> NiAl-1Hf iso<strong>the</strong>rmally oxidized for 96 hours imaged using 

bright field <strong>and</strong> STEM-HAADF techniques respectively: annealed for (a,b) two <strong>and</strong> (c,d) 

four hours. The boxes in (b) <strong>and</strong> (d) highlight <strong>the</strong> regions imaged in (a) <strong>and</strong> (c).............. 135 

Figure 17. Plan view TEM samples <strong>of</strong> NiAlCrHf iso<strong>the</strong>rmally oxidized for 96 hours imaged 

using bright field <strong>and</strong> STEM-HAADF techniques: annealed for (a,b) two <strong>and</strong> (b,c) four 

hours. The box in (a) highlights <strong>the</strong> region imaged in (b). ................................................ 136 

xvii

CHAPTER I 

GENERAL INTRODUCTION 

1.1. General Background 

Gas turbine engines typically utilize combustor <strong>and</strong> turbine superalloys with melting 

points ranging from 1230°C to 1315°C. They are applied in combustion gas environments where 

temperatures can exceed 1480°C for brief times, making <strong>the</strong>m susceptible to incipient melting 

<strong>and</strong> environmental damage due to oxidation <strong>and</strong>/or hot corrosion [1-9]. As a consequence, 

current combustor <strong>and</strong> turbine engine components must be actively cooled with discharge air 

from <strong>the</strong> compressor. An obvious approach to maximize efficiency is to reduce <strong>the</strong> use <strong>of</strong> 

cooling air. One way to accomplish this is to design more <strong>effect</strong>ive heat transfer geometries. 

This can be achieved by drilling holes through <strong>the</strong> component to bleed in air, which cools <strong>the</strong> 

combustion gas path surface [3, 10]. However, more significant benefits can be obtained by <strong>the</strong> 

application <strong>of</strong> thin <strong>the</strong>rmally insulating layers known as <strong>the</strong>rmal barrier coatings or TBCs [1-13]. 

TBCs protect turbine engine components by acting as <strong>the</strong>rmal insulators between <strong>the</strong> base 

metal <strong>and</strong> <strong>the</strong> hot gases to which <strong>the</strong>y are exposed. Thermal barrier coatings based on zirconium 

oxide have been used for nearly three decades to extend <strong>the</strong> lives <strong>of</strong> aircraft turbines <strong>and</strong> <strong>the</strong>ir 

moving <strong>and</strong> stationary components operating in hostile <strong>the</strong>rmal <strong>and</strong> corrosive environments [14, 

15]. To facilitate <strong>the</strong>ir success, <strong>the</strong> <strong>the</strong>rmal conductivity <strong>of</strong> <strong>the</strong> ceramic coating is usually a 

couple <strong>of</strong> orders <strong>of</strong> magnitude lower than <strong>the</strong> base metal. In addition, <strong>the</strong> ceramic usually has a 

higher reflectivity than <strong>the</strong> metal to take advantage <strong>of</strong> <strong>the</strong> radiative heat transfer [16]. Generally, 

<strong>the</strong>se systems are conservatively designed (i.e., over-designed) as a consequence <strong>of</strong> only partial 

1

underst<strong>and</strong>ing <strong>of</strong> many <strong>of</strong> <strong>the</strong> materials fundamentals <strong>of</strong> <strong>the</strong> TBC macro-, micro-constituent 

systems within <strong>the</strong> overall designs. 

1.2. Motivation 

Various durability <strong>and</strong> reliability issues <strong>of</strong> TBCs have prevented <strong>the</strong>m from achieving 

<strong>the</strong>ir full potential. Thus, a study was initiated <strong>the</strong> purpose <strong>of</strong> which was to investigate <strong>the</strong> 

properties <strong>and</strong> performance <strong>of</strong> model, high-strength NiAl overlay bond coatings containing up to 

1.0 at.% Hf <strong>and</strong> up to 5.0 at.% Cr as alloying <strong>additions</strong>. The coatings were produced via DC 

magnetron sputtering under conditions designed to produce different microstructural 

morphologies. The goal was to evaluate <strong>the</strong> influences <strong>of</strong> chemical composition <strong>and</strong> 

microstructure on <strong>the</strong> oxidation behavior <strong>of</strong> <strong>the</strong>se model coating alloys <strong>and</strong> to provide baseline 

information useful for <strong>the</strong> development <strong>of</strong> more advanced TBC bond coats. 

1.3. Dissertation Layout 

This dissertation will start with a brief introduction <strong>and</strong> research background. A literature 

review related to this study will be provided in Chapter II. Experimental procedures will be 

detailed in Chapter ΙΙI. Chapter IV discusses <strong>the</strong> influences <strong>of</strong> deposition parameters on <strong>the</strong> 

microstructures <strong>and</strong> properties <strong>of</strong> sputtered NiAl-Hf coatings. Chapter V presents <strong>the</strong> <strong>effect</strong>s <strong>of</strong> 

Cr <strong>additions</strong> to NiAl-Hf coatings. Chapter VI concentrates on <strong>the</strong> iso<strong>the</strong>rmal oxidation 

performance <strong>and</strong> resulting microstructures. Conclusions from this research will be summarized 

in Chapter VII. Directions for future research will be provided in Chapter VIII. 

2

CHAPTER II 

LITERATURE REVIEW 

2.1. Thermal Barrier Coatings 

2.1.1. Construction <strong>and</strong> Manufacture <strong>of</strong> TBC Layers 

Thermal barrier coatings (TBCs) have been used to extend <strong>the</strong> lives <strong>of</strong> aircraft turbines 

<strong>and</strong> <strong>the</strong>ir moving <strong>and</strong> stationary components operating in hostile <strong>the</strong>rmal <strong>and</strong> corrosive 

environments for several decades [9, 11-15, 17]. Modern TBC systems typically consist <strong>of</strong> an 

insulating ceramic topcoat applied ei<strong>the</strong>r by plasma spraying (PS) or electron beam-physical 

vapor deposition (EB-PVD) over a metallic bond coat that is deposited directly onto a Ni-base 

superalloy substrate [9, 12, 13, 17, 18]. Fig. 2.1 shows a schematic representation <strong>of</strong> a typical 

TBC cross-section. The ceramic top coat is typically yttria-stabilized zirconia (YSZ) due to its 

low <strong>the</strong>rmal conductivity <strong>and</strong> relatively high coefficient <strong>of</strong> <strong>the</strong>rmal expansion (CTE) [17, 18]. 

The primary classes <strong>of</strong> underlying bond coats are: (i) diffusion aluminide coatings (i.e. NiAl, 

PtAl, NiPtAl), <strong>and</strong> (ii) plasma sprayed MCrAlY overlay coatings, where “M” is Ni, Co, Fe or a 

combination <strong>of</strong> <strong>the</strong>se elements, <strong>and</strong> “Y” is ei<strong>the</strong>r an oxygen-active element (i.e., Y, Si, Ta or Hf), 

or a precious metal (i.e., Pt, Pd, Ru, or Re). The underlying bond coat must be resistant to both 

<strong>the</strong>rmomechanical fatigue <strong>and</strong> oxidation (as <strong>the</strong> YSZ coating is essentially transparent to oxygen 

<strong>and</strong> thus does not protect <strong>the</strong> substrate) [18-22]. During high temperature operation <strong>and</strong> even 

sometimes during <strong>the</strong> preparation, a thin <strong>the</strong>rmally grown oxide (TGO) layer consisting 

predominantly <strong>of</strong> Al 2 O 3 forms between <strong>the</strong> top coat <strong>and</strong> <strong>the</strong> bond coat. This layer is <strong>the</strong> basis for 

3

<strong>the</strong> oxidation resistance exhibited by <strong>the</strong> bond coat <strong>and</strong> provides adherence between <strong>the</strong> bond 

coat <strong>and</strong> <strong>the</strong> ceramic top coat. 

The benefit <strong>of</strong> a TBC is not only in extending component life, but also in facilitating a 

significant increase in operating temperatures. These are estimated to be as high as 150°C to 

200°C above <strong>the</strong> melting points <strong>of</strong> <strong>the</strong> turbine superalloys. This is approximately equivalent to a 

three-fold improvement in component durability <strong>and</strong> an increase <strong>of</strong> 1% in specific fuel 

consumption [1, 9, 14]. 

Material 

Yttria stabilized 

zirconia 

• Plasma Sprayed 

• EB-PVD 

Al 2 O 3 TGO 

• MCrAlY 

• Diffusion aluminide 

(i.e. NiAl, PtAl, 

etc.) 

Hot Combustion Gases 

Ceramic topcoat 

Bond coat 

Superalloy 

Attributes 

Provides 

<strong>the</strong>rmal insulation 

Provides 

oxidation <strong>and</strong> 

corrosion 

Internal Cooling Air 

Fig. 2.1. Schematic representation <strong>of</strong> a multi-layer <strong>the</strong>rmal barrier coating system. 

2.1.2. Failure Mechanism <strong>of</strong> TBC system 

Since TBCs have an insulating capability [1, 2, 14, 15, 23], <strong>the</strong> protective <strong>effect</strong>s 

supplied by TBCs only last as long as <strong>the</strong> ceramic top coat remains intact on <strong>the</strong> substrate. 

Hence, <strong>the</strong> onset <strong>of</strong> spallation <strong>of</strong> <strong>the</strong> topcoat from <strong>the</strong> component is always considered as <strong>the</strong> sign 

4

<strong>of</strong> TBC failure. Therefore, durability <strong>and</strong> reliability are <strong>the</strong> critical requirement in <strong>the</strong> 

developments <strong>of</strong> <strong>effect</strong>ive TBCs. 

Different manufacturing methods lead to different TBC microstructures <strong>and</strong> 

correspondingly different TBC failure modes. PS TBCs have porous <strong>and</strong> microcracked 

structures that <strong>of</strong>ten contain interlocking grains <strong>and</strong> pores oriented parallel to <strong>the</strong> substrate 

surface [9, 11, 13, 24, 25]. They are usually applied on top <strong>of</strong> MCrAlY bond coats. EB-PVD 

TBCs have columnar microstructures which are composed <strong>of</strong> segmented, freest<strong>and</strong>ing ceramic 

columns that extend perpendicular to <strong>the</strong> substrate [9, 11, 24, 26]. Oxidation <strong>of</strong> <strong>the</strong> bond coat 

plays a major role in <strong>the</strong> failure modes <strong>of</strong> both types <strong>of</strong> TBCs because TBC life is limited by 

physical loss <strong>of</strong> <strong>the</strong> coating (i.e., if <strong>the</strong> coating spalls, protection is lost). 

The onset <strong>of</strong> TBC spallation is caused by microstructural instabilities at <strong>the</strong> bond 

coat/TGO or TGO/top coat interface, <strong>the</strong> actual nature <strong>of</strong> which varies from one TBC system to 

ano<strong>the</strong>r [4, 9, 10, 13, 14, 24, 27-39]. Two contributors to TBC failure have been identified as <strong>the</strong> 

extent <strong>of</strong> bond coat oxidation <strong>and</strong> <strong>the</strong> strains generated by <strong>the</strong> CTE mismatch between <strong>the</strong> 

ceramic topcoat <strong>and</strong> <strong>the</strong> metallic components within <strong>the</strong> system. O<strong>the</strong>r contributing factors 

include <strong>the</strong> strength <strong>of</strong> <strong>the</strong> bond coat, microstructural <strong>and</strong> chemical changes that occur within 

bond coat during service, surface roughness <strong>of</strong> <strong>the</strong> component, <strong>and</strong> possible sintering <strong>of</strong> <strong>the</strong> top 

coat [18, 20, 26, 40-63]. In reality, it is likely that several mechanisms are operative; <strong>and</strong> all can 

influence <strong>the</strong> strain compliance <strong>of</strong> <strong>the</strong> system. 

5

PS bond coatings exhibit higher surface roughness than EB-PVD bond coatings which 

initially increases topcoat adherence through mechanical interlocking. In PS TBCs, failure 

occurs by delamination cracking arising from in-plane compressive stresses which cause large 

out-<strong>of</strong>-plane tensile stresses. These tensile stresses produce delamination cracks within <strong>the</strong> TBC 

that are parallel to <strong>the</strong> interface <strong>and</strong> near peaks in <strong>the</strong> rough bond coat [2, 4, 11, 13, 64, 65]. 

In EB-PVD coatings, failures generally occur at <strong>the</strong> TGO/bond coat interface. However, 

if substantial bond coat surface rumpling occurs, failures can transfer to <strong>the</strong> TGO/top coat 

interface [22, 30, 35, 36, 66-70]. Failure at <strong>the</strong> TGO/bond coat interface results from a 

progressive reduction <strong>of</strong> adhesion due to <strong>the</strong> nucleation <strong>and</strong> growth <strong>of</strong> microcracks. Thus, bond 

coat oxidation is also a major driving force for failure in EB-PVD TBCs. 

MCrAlY <strong>and</strong> diffusion aluminide coatings are all Al 2 O 3 -formers. As such, <strong>the</strong>y exhibit 

equivalent oxide growth rates, so long as an α-Al 2 O 3 TGO forms at <strong>the</strong> top coat/bond coat 

interface [13]. Coating performance ultimately depends on <strong>the</strong> ability <strong>of</strong> <strong>the</strong> TGO to resist 

spallation as a result <strong>of</strong> defect formation, microcrack growth <strong>and</strong> adhesion loss. Therefore, 

critical issues are <strong>the</strong> ease <strong>of</strong> void formation at <strong>the</strong> TGO/bond coat interface, how readily less 

protective oxides such spinel (NiAl 2 O 4 ) or θ-Al 2 O 3 form beneath <strong>the</strong> TGO, <strong>and</strong> how easily 

debonding occurs at <strong>the</strong> TGO/bond coat interface. All <strong>of</strong> <strong>the</strong>se things are dictated by <strong>the</strong> 

composition <strong>and</strong> properties <strong>of</strong> <strong>the</strong> bond coat. 

6

2.2. Properties <strong>of</strong> NiAl 

The intermetallic compound NiAl has been used in a wide variety <strong>of</strong> applications 

because <strong>of</strong> its high melting point (~1995K), low density (~5.9 g/cm 3 ), good <strong>the</strong>rmal conductivity 

<strong>and</strong> excellent environmental oxidation resistance [71-73]. However, numerous <strong>investigation</strong>s 

have shown that conventional NiAl alloys are limited by low ductility <strong>and</strong> low impact resistance 

[74, 75] at ambient temperatures <strong>and</strong> poor creep strength at high temperatures [76, 77]. The 

following sections will give a brief overview <strong>of</strong> <strong>the</strong> general properties <strong>of</strong> NiAl. 

2.2.1. Phase Equilibria <strong>and</strong> Crystal Structure 

The binary Ni-Al phase diagram is shown in Fig. 2.2. There are five intermetallic phases 

in this system: NiAl 3 , Ni 2 Al 3 , NiAl, Ni 5 Al 3 , <strong>and</strong> Ni 3 Al. NiAl exhibits a wide range <strong>of</strong> solubility 

<strong>and</strong>, at its stoichiometric composition, a melting temperature <strong>of</strong> 1638°C. Deviations from 

stoichiometry are accommodated by <strong>the</strong> substitution <strong>of</strong> nickel atoms onto aluminum sites in 

nickel-rich compositions <strong>and</strong> by <strong>the</strong> formation <strong>of</strong> vacancies on <strong>the</strong> nickel sites in aluminum-rich 

compositions [78]. 

7

1800 

1600 

1400 

1200 

1000 

L 

NiAl 

(β) 

L 

Ni 

(γ) 

800 

600 

Al 

400 

0 10 20 30 40 50 60 70 80 90 100 

Al 

At. % Ni 

Ni 

Fig. 2.2. Phase diagram <strong>of</strong> Ni-Al [79]. 

NiAl has a B2 CsCl crystal structure. As shown in Fig. 2.3, one type <strong>of</strong> atom locates at 0, 

0, 0 <strong>and</strong> <strong>the</strong> o<strong>the</strong>r at 1/2, 1/2, 1/2. The lattice parameter <strong>of</strong> <strong>the</strong> stoichiometric composition at 

room temperature is 0.2887 nm [80]. 

2.2.2. Influence <strong>of</strong> Ternary Alloying Additions 

The ternary alloying elements <strong>of</strong> NiAl-X alloys can be categorized into three groups 

according to <strong>the</strong> alloy element’s position in <strong>the</strong> periodic table (Table 2.1) [81, 82]. Type A 

elements, which are those in Groups IIIB, IVB <strong>and</strong> VB, form at least one ternary intermetallic 

phase with NiAl, <strong>and</strong> exhibit low solubility. Type B elements, which include those in Group 

VIB <strong>and</strong> rhenium from Group VIIB, form pseudobinary eutectics with NiAl. They have limited 

solubility in NiAl. Type C elements, which consist <strong>of</strong> Group VIII elements plus manganese <strong>and</strong> 

copper, show large solubility in NiAl <strong>and</strong> always form an isostructural B2 intermetallic. 

8

Hafnium which is in Group IVB is classified as a type A element. A portion <strong>of</strong> <strong>the</strong> 

ternary Ni – Al – Hf phase diagram is shown in Fig. 2.4. It shows five intermetallic phases in <strong>the</strong> 

Ni-rich corner <strong>of</strong> this system: γ΄-Ni 3 Al, Al 3 Ni, NiAlX (Laves phase), Ni 2 AlX (Heusler phase or 

H) <strong>and</strong> Ni 7 (X,Al) 2 [83, 84]. The ternary L2 1 Heusler phase, presented in Fig. 2.5, is a fur<strong>the</strong>r 

ordering <strong>of</strong> <strong>the</strong> B2 phase to Ni 2 AlX, where <strong>the</strong> element X resides on an equivalent Al site. The 

NiAlX Laves phase has a C14 hexagonal crystal structure. 

2.1.2.1. Mechanical Properties 

It has been reported that small <strong>additions</strong> <strong>of</strong> Hf <strong>and</strong> Zr can be very <strong>effect</strong>ive in improving 

<strong>the</strong> high-temperature creep strength <strong>of</strong> NiAl polycrystals <strong>and</strong> single crystals [71, 85-94]. The 

streng<strong>the</strong>ning is believed to be provided by solid solution hardening due to <strong>the</strong> segregation <strong>of</strong> 

solute atoms to <strong>the</strong> dislocations <strong>and</strong> grain boundaries [71, 85] <strong>and</strong> precipitation hardening due to 

<strong>the</strong> nucleation <strong>and</strong> growth <strong>of</strong> Heusler phase or G-phase (Ni 16 Hf 6 Si 7 or Ni 16 Zr 6 Si 7 ) precipitates 

after prolonged aging at high temperatures [91, 94-96]. 

Ni 

Al 

Fig. 2.3. NiAl B2 (CsCl) structure. 

9

2.1.2.2. Oxidation Resistance 

It is widely recognized that small <strong>additions</strong> <strong>of</strong> a number <strong>of</strong> reactive elements (e.g. Hf, Zr, 

Y, La, Ce, etc) can significantly improve <strong>the</strong> oxidation resistance <strong>of</strong> high temperature alloys 

containing aluminum <strong>and</strong> or <strong>chromium</strong>. This was first called <strong>the</strong> rare-earth <strong>effect</strong>, <strong>and</strong> more 

recently <strong>the</strong> reactive element (RE) <strong>effect</strong> [97-102]. The reactive element <strong>effect</strong> which was first 

patented by Pfeil [103] in 1937, was primarily concerned with cerium <strong>additions</strong> to a Ni-Cr alloy. 

Since <strong>the</strong>n, <strong>the</strong> same <strong>effect</strong> has been observed for a variety <strong>of</strong> elements that show a high affinity 

for oxygen <strong>and</strong> sulfur, <strong>and</strong> benefits have also been observed for most chromia- <strong>and</strong> alumina- 

forming alloys [97, 101]. 

Table 2.1. Portion <strong>of</strong> periodic table illustrating <strong>the</strong> three types <strong>of</strong> NiAl-X phase equilibria [81, 

82]. 

IIIB IVB VB VIB VIIB VIIIB IB IIB 

Sc Ti V Cr Mn Fe Co Ni Cu Zn 

Y Zr Nb Mo Tc Ru Rh Pd Ag Cd 

La Hf Ta W Re Os Ir Pt Au Hg 

Ni-Al-X ternary phase(s) 

NiAl-X pseudobinary eutectic 

High solubility in NiAl <strong>and</strong>/or forms B2 aluminide 

Generally, <strong>the</strong> RE <strong>effect</strong>s manifest <strong>the</strong>mselves in <strong>the</strong> form <strong>of</strong> improved oxide scale 

adhesion <strong>and</strong> reduced oxide scale growth rate. Considerable work has been done to try to 

explain RE <strong>effect</strong>s over <strong>the</strong> last 30 years. Recently, it has been suggested that <strong>the</strong> RE <strong>effect</strong>s are 

likely to be related to <strong>the</strong> following aspects: 

10

1. The sulfur <strong>effect</strong>. Studies have shown that <strong>the</strong> sulfur content plays an important but 

detrimental role in scale spallation [104-109]. However, with a RE addition, <strong>the</strong> detrimental 

role <strong>of</strong> sulfur can be minimized or eliminated. This is thought to result from <strong>the</strong> gettering <strong>of</strong> 

S or some o<strong>the</strong>r operative mechanism. RE <strong>additions</strong> eliminate <strong>the</strong> segregation <strong>of</strong> S [100, 

110]. 

40 

β 

20 

Ni 7 

Hf 2 

Ni 2 

AlHf (H) 

Ni 

Liquid 

Ni 5 

Hf 

20 40 

Fig. 2.4. Ni-Al-Hf ternary phase diagram [84]. 

11

Ni 

Al 

Hf 

Fig. 2.5. Crystal structure <strong>of</strong> Heusler phase. 

2. Diffusion paths. It has been suggested that grain boundaries are <strong>the</strong> only active diffusion 

paths within <strong>the</strong> oxide scale. The segregation <strong>of</strong> RE <strong>additions</strong> to <strong>the</strong> scale grain boundaries 

are proposed to decrease elemental diffusion along <strong>the</strong>se boundaries since <strong>the</strong> large RE atoms 

diffuse more slowly than o<strong>the</strong>r elements (such as Al or Cr) [100]. This causes <strong>the</strong> scale 

growth rate to be greatly reduced. 

3. Substrate grain size <strong>effect</strong>s. With RE <strong>additions</strong>, a finer substrate grain size can be achieved 

due to <strong>the</strong> formation <strong>of</strong> RE-rich precipitates (such as <strong>the</strong> Heusler phase) [97, 101]. The 

precipitates can also inhibit alloy grain growth during annealing or extended high 

temperature oxidation. A fine alloy grain size promotes protective scale formation by 

increasing <strong>the</strong> diffusivity <strong>of</strong> Al or Cr in <strong>the</strong> substrate [100]. 

12

O<strong>the</strong>r suggestions include blocking <strong>of</strong> element transportation by <strong>the</strong> formation <strong>of</strong> a RE 

oxide layer at <strong>the</strong> scale-alloy interface, decrease <strong>of</strong> scale grain size <strong>and</strong> improvement <strong>of</strong> scale 

plasticity due to RE oxide particles acting as nucleation sites for <strong>the</strong> first-formed oxides, 

stabilization <strong>of</strong> crack propagation in scale by <strong>the</strong> formation <strong>of</strong> second-phase oxide particles, 

condensation <strong>of</strong> vacancies at metal-scale interface due to oxide dispersions acting as vacancy 

sinks [99], <strong>and</strong> etc. Although <strong>the</strong> results from some <strong>of</strong> <strong>the</strong>se studies were contradictory, toge<strong>the</strong>r 

<strong>the</strong>y still provide a comprehensive explanation <strong>of</strong> <strong>the</strong> impact <strong>of</strong> RE <strong>additions</strong> in high temperature 

alloys. 

2.1.3. NiAl alloys as TBC bond coats 

NiAl <strong>and</strong> its derivatives have been used for more than 30 years as bond coats in TBCs 

due to <strong>the</strong>ir excellent oxidation <strong>and</strong> corrosion resistance <strong>and</strong> abilities to maintain <strong>the</strong>ir structural 

integrity during service. However, <strong>the</strong> use <strong>of</strong> conventional NiAl alloys have been hindered by 

low ductility <strong>and</strong> low impact resistance [74, 75, 111] at ambient temperatures <strong>and</strong> poor creep 

strength at high temperatures [76, 77]. In recent years, platinum-modified diffusion aluminide 

coatings, i.e. (Ni,Pt)Al, have emerged as <strong>the</strong> bond coatings <strong>of</strong> choice. It is well established that 

<strong>the</strong>y exhibit better high temperature corrosion resistance than <strong>the</strong>ir Pt-free counterparts which 

translates into extended service lifetimes for TBCs [18, 57, 112, 113]. 

As noted above, it is well known that small RE <strong>additions</strong>, generally in solid solution, can 

significantly improve <strong>the</strong> oxidation resistance <strong>of</strong> Al 2 O 3 <strong>and</strong> Cr 2 O 3 forming alloys by increasing 

oxide scale adhesion. In NiAl <strong>and</strong> (Ni,Pt)Al bond coats, RE contents are maintained around 0.05 

at.% which are thought to be near <strong>the</strong>ir solubility limits. RE <strong>additions</strong> exceeding this amount 

13

have been linked to extensive internal oxidation resulting in <strong>the</strong> formation <strong>of</strong> oxide stringers or 

pegs [18, 89, 102, 114-118]. It has also been shown that optimal RE contents can significantly 

improve TBC lifetimes. The RE <strong>effect</strong>s in <strong>the</strong>se alloys have been attributed to improved alumina 

scale adhesion <strong>and</strong> reduced scale growth rate caused by <strong>the</strong> segregation <strong>of</strong> RE ions to <strong>the</strong> scale 

grain boundaries <strong>and</strong> alloy-scale interface [100]. 

The benefits <strong>of</strong> RE <strong>additions</strong> are not limited simply to alloying <strong>of</strong> <strong>the</strong> coating. Several 

studies have demonstrated that RE <strong>additions</strong> to <strong>the</strong> underlying substrates can also lead to 

increased oxidation resistance [112, 119-122]. All <strong>of</strong> <strong>the</strong>se studies showed that Hf diffused 

readily from <strong>the</strong> substrates into <strong>the</strong> coatings <strong>and</strong> to <strong>the</strong> growing oxide films. Tolpygo et al. 

[122], who investigated <strong>the</strong> <strong>effect</strong>s <strong>of</strong> C, Y, <strong>and</strong> Hf <strong>additions</strong> to CMSX-4 on <strong>the</strong> rumpling 

resistance <strong>of</strong> platinum aluminide diffusion coatings, argued that Hf incorporated into <strong>the</strong> coating 

<strong>and</strong> growing oxide increased <strong>the</strong>ir creep resistance <strong>and</strong> thus <strong>the</strong>ir resistance to rumpling. 

Pint et al. [123] proposed that <strong>the</strong> creation <strong>of</strong> a pure Hf-doped NiAl coating could 

represent a promising approach for improving <strong>the</strong> oxidation performance <strong>of</strong> bond coats <strong>and</strong> TBC 

lifetime. Over <strong>the</strong> years, several groups have attempted to achieve this goal. Xiang et al. [124- 

126] have attempted to use pack cementation to produce Hf-modified coatings. They 

demonstrated that codeposition <strong>of</strong> Hf <strong>and</strong> Al by <strong>the</strong> pack cementation was possible from pack 

powder mixtures containing Hf, CrCl 3·6H 2 O <strong>and</strong> Al 2 O 3 ei<strong>the</strong>r with or without <strong>additions</strong> <strong>of</strong> 

elemental Al. The resulting microstructures were found to depend strongly on pack composition 

<strong>and</strong> could be controllably produced; however, <strong>the</strong> resulting coatings were generally multilayered 

14

consisting <strong>of</strong> a series <strong>of</strong> Hf-rich phases on top <strong>of</strong> a NiAl layer. Oxidation resistance was not 

studied for <strong>the</strong>se coatings. 

Lee et al. [127-129] have investigated ways to produce NiAl-Hf bond coatings via hightemperature, 

low-activity chemical vapor deposition (CVD) processes. Although <strong>the</strong>y did 

succeed in producing Hf containing coatings, <strong>the</strong>y found that it was difficult to control <strong>the</strong> 

concentration <strong>and</strong> distribution <strong>of</strong> Hf in <strong>the</strong> coating matrix due in part to interdiffusion <strong>and</strong> phase 

formation between <strong>the</strong> superalloy substrates <strong>and</strong> <strong>the</strong> coating layer. These results were essentially 

<strong>the</strong> same as those <strong>of</strong> Warnes [130] who in 2001 published <strong>the</strong> results from an industrial study 

conducted in 1991 <strong>and</strong> 1992. 

More recently, Nesbitt et al. [131] produced NiAl-based overlay bond coats containing 

Pt, Hf, <strong>and</strong> Zr <strong>additions</strong> via low pressure plasma spraying. One advantage <strong>of</strong> this production 

method is that it eliminates any significant processing-induced chemical interaction between <strong>the</strong> 

coating <strong>and</strong> <strong>the</strong> substrate. Interestingly, <strong>the</strong>y observed that NiAl coatings containing small 

amounts <strong>of</strong> Zr or Zr plus 7 to 14 at.% Pt exhibited shorter TBC lifetimes than binary NiAl alloys. 

This observation, which contradicts <strong>the</strong> predictions that RE <strong>additions</strong> would increase <strong>the</strong> service 

lifetimes <strong>of</strong> TBC systems, was attributed to <strong>the</strong> development <strong>of</strong> a thicker aluminum oxide layer 

<strong>and</strong> <strong>the</strong> growth <strong>of</strong> oxide stringers inward from <strong>the</strong> aluminum oxide into <strong>the</strong> Zr-containing 

coatings. 

15

2.1.4. High strength NiAl alloys as TBC bond coats 

It is generally believed that optimal Hf <strong>and</strong>/or Zr contents are around 0.05 at.% which is 

near <strong>the</strong>ir solubility limits in NiAl <strong>and</strong> (Ni,Pt)Al. Prior reports suggest that concentrations <strong>of</strong> 

RE’s in excess <strong>of</strong> <strong>the</strong>ir solubility limits will lead to catastrophic oxidation [132]. This is because 

<strong>of</strong> <strong>the</strong> nucleation <strong>of</strong> less oxidation resistant precipitates at grain boundaries within <strong>the</strong> coatings. 

Fisher et al. [133] investigated <strong>the</strong> influences <strong>of</strong> <strong>the</strong> ion implantation <strong>of</strong> Hf or Y to platinum 

aluminide diffusion coatings on MarM002 substrates <strong>and</strong> reported that Hf <strong>and</strong> Y had no 

influence on coating performance. Ra<strong>the</strong>r <strong>the</strong>y noted that Hf <strong>and</strong> Y promoted <strong>the</strong> formation <strong>of</strong> 

oxides that had similar adhesion, growth rates, <strong>and</strong> compositions as those formed on non-ion 

implanted coatings. Hafnium ion implantation did promote <strong>the</strong> formation <strong>of</strong> Hf-rich pegs at <strong>the</strong> 

oxide/coating interface which were associated with scale spallation. These observations differed 

from <strong>the</strong> earlier work <strong>of</strong> Jedlinski <strong>and</strong> co-workers [134-138] who showed that ion implanted Y 

increased oxide scale formation <strong>and</strong> decreased <strong>the</strong> rate <strong>of</strong> oxidation. 

Recently, researchers from <strong>the</strong> General Electric Company have developed a series <strong>of</strong> 

NiAl-based overlay bond coatings containing higher RE concentrations [139-147]. These 

coatings, which contain Hf <strong>and</strong> Zr in excess <strong>of</strong> 0.05 at.% <strong>and</strong> which can contain between 2 <strong>and</strong> 

15 at.% Cr, are reported to exhibit oxidation resistance that is comparable to modern (Ni,Pt)Albased 

coatings, increased resistance to rumpling, <strong>and</strong> improved spallation resistance when 

incorporated into TBC systems [147]. Due to <strong>the</strong>ir high RE concentrations <strong>the</strong>se new coatings 

<strong>of</strong>ten contain metallic α-Cr <strong>and</strong> intermetallic precipitates such as <strong>the</strong> Ni 2 AlHf (β´-Heusler) 

phase, which are believed to streng<strong>the</strong>n <strong>the</strong> coating via a precipitation hardening mechanism. 

Such streng<strong>the</strong>ning has been reported to occur in bulk Zr <strong>and</strong>/or Hf containing single crystals 

16

from which <strong>the</strong>se coatings were derived [75, 76, 87, 91, 96, 148-150]. These alloys have 

additionally been shown to exhibit favorable oxidation resistance both in bulk <strong>and</strong> coating forms 

[118, 147]. Balint <strong>and</strong> Hutchinson [151, 152], who developed a mechanics-based model to 

describe rumpling in <strong>the</strong>rmal barrier coating systems, predict that increasing <strong>the</strong> creep resistance 

<strong>of</strong> <strong>the</strong> bond coat would be more resistant to rumpling resulting in better oxidation resistance. 

Spitsberg et al. [146] report that <strong>the</strong> oxidation resistance <strong>of</strong> β+β´ coatings can be 

improved by applying in-situ or post-deposition processes which cause <strong>the</strong> coatings to 

recrystallize, resulting in coatings with lower residual stress, more stable grain structures, <strong>and</strong> 

with <strong>the</strong> largest number <strong>of</strong> precipitates occupying grain interiors ra<strong>the</strong>r than grain boundaries. 

Despite <strong>the</strong> potential advantages <strong>of</strong> such coating schemes, very little published information exists 

concerning <strong>the</strong>ir properties <strong>and</strong> performance. 

Ning et al. [153-157] have produced Hf-containing β+β´ overlay coatings via direct 

current magnetron sputtering from composite sputtering targets. In a more recent study, Guo <strong>and</strong> 

co-workers [158, 159] produced analogous coatings via EB-PVD from binary NiAl <strong>and</strong> pure Hf 

sources. Both studies showed that viable multiphase bond coats could be produced via different 

physical vapor deposition processes <strong>and</strong> that those coatings could <strong>effect</strong>ively protect superalloys 

from oxidation. Both also showed lower mass gains for coatings containing less than 0.5 at.% 

Hf in comparison to higher Hf content bond coats. In <strong>the</strong> case <strong>of</strong> Ning et al., <strong>the</strong> coatings were 

found to have sub-micron grain structures with Heusler phase precipitates. 

17

CHAPTER III 

EXPERIMENTAL PROCEDURE 

3.1. Sample Preparation 

3.1.1. Target <strong>and</strong> substrate preparation 

The sputtering targets used in this study were provided by ACI Alloys (San Jose, CA) 

with nominal compositions presented in Table 3.1. All targets were prepared via arc melting in 

an argon atmosphere, <strong>and</strong> <strong>the</strong> surfaces were ground flat prior to deposition. 

Four types <strong>of</strong> substrates were used with this study: Si (111) wafers, corning 1737 glass, 

CMSX-4 ® coupons, <strong>and</strong> Rene N5 coupons. All substrates were ultrasonically cleaned in acetone 

<strong>and</strong> ethanol prior to deposition. Table 3.2 lists <strong>the</strong> nominal chemical compositions <strong>of</strong> CMSX-4 ® 

<strong>and</strong> Rene N5. The CMSX-4® substrates were EDM wire cut to dimensions <strong>of</strong> 9 mm × 9 mm × 2 

mm. Prior to deposition, <strong>the</strong> substrates were subjected to a st<strong>and</strong>ard heat treatment in purified 

argon to stabilize <strong>the</strong> microstructure <strong>and</strong> chemistry. The st<strong>and</strong>ard heat treatment procedure was 

four hours at 1573K, followed by four hours at 1473K, <strong>and</strong> finally eight hours at 1143K. Then, 

<strong>the</strong> substrates were polished to a 1200 grit surface-finish using SiC grinding papers. 

18

Table 3.1. Nominal Target Chemical Compositions in Atomic Percent. 

Alloy Ni Al Cr Hf 

NiAl 51 49 --- --- 

NiAl-0.5Hf 51 48 --- 0.5 

NiAl-1.0Hf 51 48 --- 1.0 

NiAlCrHf 50 44 5.0 1.0 

Table 3.2. Chemical compositions <strong>of</strong> CMSX-4 ® <strong>and</strong> Rene N5 superalloys. 

Element Ni Al Ti Ta Cr Mo W Co Re 

CMSX-4® (wt.%) 61.7. 5.6 1.0 6.5 6.5 0.6 6.0 9.0 3.0 

(at.%) 63.8 12.6 1.3 2.2 7.6 0.4 2.0 9.3 1.0 

René N5 (wt.%) 63.6 6.14 --- 6.5 7.1 1.4 4.94 7.3 2.9 

(at.%) 65.1 13.7 --- 2.2 8.2 0.9 1.6 7.4 0.9 

3.1.2. Deposition <strong>of</strong> coatings 

In general, sputtering processes can be divided into four main categories: DC, RF, 

magnetron, <strong>and</strong> reactive. However, <strong>the</strong>re are variants within each category <strong>and</strong> hybrids between 

<strong>the</strong>se categories[160]. The coatings developed for this study were manufactured using an AJA 

International ORION 4 DC magnetron sputtering system arranged in a sputter-down 

configuration. Examples <strong>of</strong> o<strong>the</strong>r DC magnetron magnet configurations are shown in Fig. 3.1. 

The magnet configuration used with this study was unbalanced high rate. This condition was 

selected because it yields <strong>the</strong> highest deposition rate <strong>and</strong> a dense microstructure. Because <strong>the</strong> 

target is connected to <strong>the</strong> negative terminal <strong>of</strong> <strong>the</strong> DC or RF power supply, it is known as <strong>the</strong> 

cathode. After evacuation <strong>of</strong> <strong>the</strong> sputtering chamber, a working gas, typically argon, is 

19

introduced <strong>and</strong> ionized. The ions <strong>of</strong> <strong>the</strong> plasma are <strong>the</strong>n accelerated towards <strong>the</strong> target <strong>and</strong> eject 

<strong>the</strong> neutral atoms out <strong>of</strong> <strong>the</strong> target through momentum transfer[160]. The atoms being ejected 

out <strong>of</strong> <strong>the</strong> target finally go through <strong>the</strong> chamber, reach <strong>the</strong> substrate surface, <strong>and</strong> deposit on it as a 

growing film or coating. The resultant microstructure <strong>of</strong> <strong>the</strong> coating is controlled by <strong>the</strong> material 

itself <strong>and</strong> corresponding deposition parameters, which will be addressed in <strong>the</strong> following 

chapters. The primary deposition parameters being used with this study are listed in Table 3.3. 

Coating morphology was varied by stage height, working gas pressure, <strong>and</strong> temperature. The 

substrates were allowed to pre-heat <strong>and</strong> stabilize at <strong>the</strong> deposition temperature for at least 45 

minutes prior to deposition. 

Table 3.3. Deposition parameters for NiAl-X coatings. 

Base Pressure 

Detailed microstructural analyses were accomplished using an FEI TECNAI G20 Supertwin 

200 keV field emission gun transmission electron microscope (FEG-TEM). The TECNAI 

is also equipped with a high angle annular dark filed detector (HAADF) <strong>and</strong> nano-Energy 

Dispersive X-ray Spectroscopy (EDX) system. Integration <strong>of</strong> <strong>the</strong> microscope with <strong>the</strong> HAADF 

detector allows one to operate <strong>the</strong> TEM in <strong>the</strong> scanning transmission electron microscopy 

(STEM) mode. The STEM mode is able to show contrast from elements with different atomic 

numbers. This is useful for precipitation or diffusion <strong>of</strong> elements with different atomic numbers. 

TEM sample preparation was conducted using an FEI Quanta 200-3D dual-beam focused 

ion beam (FIB). Plan view samples were prepared using a pre-defined automated script <strong>and</strong> 

specimens were removed from <strong>the</strong> parent epoxy mounted sample via in-situ liftout. All TEM 

samples were plasma cleaned using a Fischione® Model 1020 plasma cleaner being analyzing in 

<strong>the</strong> TEM. This was done to remove any contamination from <strong>the</strong> samples before analysis. 

Chemical analysis was achieved using an Energy Dispersive X-ray Spectrometer (EDX) 

attached to <strong>the</strong> JEOL 7000F FEG-SEM <strong>and</strong> Wavelength Dispersive X-ray Spectroscopy (WDX) 

in a JEOL 8900 electron probe microanalyzer (EPMA). The EPMA tests used an accelerating 

voltage <strong>of</strong> 20 keV with a beam size setting <strong>of</strong> approximately one micrometer. 

Precipitate chemistry was investigated by TEM nano-EDX <strong>and</strong> atom probe tomography 

(APT). The atom probe tomography was conducted using an IMAGO 3000-X Si Local 

Electrode Atom Probe (LEAP) equipped with a pulsed laser attachment. LEAP samples were 

prepared using <strong>the</strong> FIB following procedures outlined by Miller [161-165]. 

21

3.1.4. Heat Treatment 

Post-deposition annealing was conducted to increase <strong>the</strong> adherence <strong>of</strong> <strong>the</strong> bond coating to 

<strong>the</strong> substrate <strong>and</strong> to alleviate sputtering defects (e.g. dislocations <strong>and</strong> residual stresses). This 

was done in a Thermal Technology® Astro model 1000-4560-FP20 graphite furnace <strong>and</strong> a 

modified Thermolyne® 2100 tube furnace with a quartz tube. The atmosphere in <strong>the</strong> furnace was 

argon with 5 wt.% hydrogen to eliminate possible oxidation during annealing. The different 

furnaces were used to investigate <strong>the</strong> <strong>effect</strong> <strong>of</strong> carbon content on <strong>the</strong> oxidation performance <strong>of</strong> 

<strong>the</strong> coatings during iso<strong>the</strong>rmal oxidation exposure. 

3.2. Iso<strong>the</strong>rmal <strong>and</strong> Cyclic Oxidation Tests 

The iso<strong>the</strong>rmal <strong>and</strong> cyclic oxidation tests were done at a constant temperature <strong>of</strong> 1323K 

in air for times ranging from 25-100 hours. Specimens, coated on one side only, were placed in 

high purity aluminum oxide boats <strong>and</strong> inserted into <strong>the</strong> hot zone <strong>of</strong> a Thermolyne® 2110 tube 

furnace. The samples were covered to collect any spalled oxides during oxidation. SEM <strong>and</strong> 

EDX were used to analyze <strong>the</strong> microstructural <strong>and</strong> compositional changes in <strong>the</strong> coatings after 

oxidation. 

22

CHAPTER IV 

INFLUENCES OF ANNEALING AND HAFNIUM CONCENTRATION ON THE 

MICROSTRUCTURES OF SPUTTER DEPOSITED β-NIAL COATINGS ON 

SUPERALLOY SUBSTRATES 

M.A. Bestor 1 , R.L. Martens 1 , <strong>and</strong> M.L. Weaver 1 

1- The University <strong>of</strong> Alabama, Department <strong>of</strong> Metallurgical & Materials Engineering, Box 

870202, Tuscaloosa, AL 35487-0202, USA 

Keywords: Nickel aluminides, bond coatings, reactive element doping, DC magnetron 

sputtering 

Abstract 

This study has investigated <strong>the</strong> influences <strong>of</strong> high <strong>hafnium</strong> concentrations <strong>and</strong> high temperature 

annealing on <strong>the</strong> microstructures <strong>of</strong> β-NiAl coatings. Binary NiAl <strong>and</strong> ternary NiAl-1.0 at.%Hf 

coatings were deposited onto CMSX-4 superalloy substrates via direct current (DC) magnetron 

sputtering. Investigations were conducted using electron probe microanalysis, transmission 

electron microscopy, x-ray diffraction, <strong>and</strong> three-dimensional atom probe tomography in an 

effort to determine <strong>the</strong> influences <strong>of</strong> annealing on <strong>the</strong> coating microstructures <strong>and</strong> coatingsubstrate 

interdiffusion. Post-deposition annealing <strong>of</strong> NiAl-1.0 at% Hf coatings at 1000°C for 

one to four hours produced nanometer-sized Hf-rich precipitates. When in <strong>the</strong> presence <strong>of</strong> high 

carbon concentrations, Hf was found to partition with C. This sometimes resulted in <strong>the</strong> 

formation <strong>of</strong> HfC precipitates. The results have been analyzed <strong>and</strong> discussed relative to previous 

research on sputter deposited NiAl-Hf coatings. 

23

4.1. Introduction 

Coatings are a vital part <strong>of</strong> everyday life being used in a variety <strong>of</strong> applications. They are 

essential in many high temperature applications, particularly for protection from potentially 

corrosive environments [1-7]. Alloyed nickel aluminides are <strong>of</strong> particular interest for use as 

bond coats in <strong>the</strong>rmal barrier <strong>and</strong> environmental barrier coating systems due to <strong>the</strong>ir high 

strengths <strong>and</strong> abilities to form a protective aluminum oxide layer. It is this aluminum oxide layer 

that limits <strong>the</strong> amount <strong>of</strong> oxygen that can diffuse into critical engine components causing 

catastrophic failure. This layer is also one <strong>of</strong> <strong>the</strong> most critical components <strong>of</strong> a <strong>the</strong>rmal barrier 

coating (TBC) system. TBCs are used to protect superalloy engine components from <strong>the</strong> harsh 

environments that <strong>the</strong>y are exposed to as a result <strong>of</strong> operation in gas turbine engines [1-7]. 

Temperatures in <strong>the</strong>se engines can exceed 1650°C with metal temperatures reaching as high as 

1200°C [8]. TBCs are primarily used to insulate <strong>the</strong> underlying superalloy; thus, allowing <strong>the</strong> 

engines to operate at higher temperatures with higher <strong>the</strong>rmodynamic efficiency. 

While TBCs have been studied intently over <strong>the</strong> past three decades, <strong>the</strong>re is still much 

that is not understood about <strong>the</strong> design <strong>and</strong> implementation <strong>of</strong> cost-<strong>effect</strong>ive coating solutions. 

This is particularly true with regard to <strong>the</strong> development <strong>of</strong> bond coat alloys [9]. As a result, 

manufacturers must compensate by over designing TBCs to obtain <strong>the</strong> required properties <strong>and</strong> 

service lifetimes, <strong>of</strong>ten at <strong>the</strong> expense <strong>of</strong> component performance <strong>and</strong> efficiency. Throughout 

<strong>the</strong> 1990’s <strong>and</strong> early 2000’s, research efforts increased to engineer new bond coats with 

compositions optimized to improve TBC performance <strong>and</strong> overall component lifetime. 

24

Traditional <strong>investigation</strong>s have focused on model alloys compositions or on existing 

coating compositions to which relatively small concentrations <strong>of</strong> reactive elements (i.e., Y, Zr, 

Hf) were added [2, 6, 10-15]. These studies demonstrated that reactive element <strong>additions</strong> to <strong>the</strong> 

coating alloy dramatically improved coating <strong>and</strong> TBC lifetimes [16]. One requirement for <strong>the</strong>se 

improvements was that <strong>the</strong> reactive elements not exceed <strong>the</strong>ir solubility limits in <strong>the</strong> bond coat 

alloy, lest <strong>the</strong>y lead to dramatically increased oxidation rates <strong>and</strong> catastrophic internal oxidation 

[7, 17]. 

Recent studies suggest that coating alloys containing far larger reactive element 

concentrations, i.e., overdoped coatings, can exhibit properties that, from a performance 

perspective, equal or exceed those <strong>of</strong> state-<strong>of</strong>-<strong>the</strong>-art (Ni,Pt)Al diffusion aluminide coatings, in 

spite <strong>of</strong> <strong>the</strong>ir higher intrinsic oxidation rates [9, 18-20]. The present study highlights <strong>the</strong> <strong>effect</strong>s 

<strong>of</strong> Hf overdoping <strong>and</strong> subsequent annealing on <strong>the</strong> microstructures <strong>of</strong> β-NiAl based overlay 

coatings deposited onto second generation Ni-based superalloy substrates. 

4.2. Experimental 

Alloy sputtering targets for this study were purchased from ACI Alloys, Inc., San Jose, 

CA. The nominal <strong>and</strong> experimentally determined target compositions are presented in Table 1. 

Coatings ~30μm thick were deposited via direct current (DC) magnetron sputtering onto CMSX- 

4 ® substrates using an AJA International, Inc. ORION 4 sputtering system with a sputter-down 

configuration. Argon was used as a working gas. The deposition power was 300 W, <strong>and</strong> <strong>the</strong> 

working gas pressure was set at 1.33 Pa. Prior to deposition, <strong>the</strong> substrates were heated to 

400°C. Additional samples were preheated to 650°C to increase coating adhesion. All <strong>of</strong> <strong>the</strong> 

25

substrates were coated on <strong>the</strong> two primary faces for subsequent microstructural analyses, 

annealing, <strong>and</strong> oxidation experiments. 

The as-deposited coatings were annealed at 1000°C from one to four hours in a flowing 

mixture <strong>of</strong> Ar + 5% H 2 to improve adhesion between <strong>the</strong> coating <strong>and</strong> substrate, anneal out 

sputtering induced defects, <strong>and</strong> alleviate any processing induced residual stresses. Annealing 

was done ei<strong>the</strong>r in a graphite element annealing furnace or in a horizontal tube furnace inside <strong>of</strong> 

a quartz tube. This was done to study <strong>the</strong> <strong>effect</strong> <strong>of</strong> heat treating environment on <strong>the</strong> resulting 

microstructures <strong>and</strong> coating properties. 

The chemical compositions <strong>of</strong> <strong>the</strong> sputtered coatings were determined via wavelength 

dispersive X-ray spectroscopy (WDS) <strong>and</strong> three-dimensional atom probe tomography (3D-APT). 

X-ray diffraction (XRD), scanning electron microscopy (SEM) <strong>and</strong> transmission electron 

microscopy (TEM) were used for phase identification <strong>and</strong> to investigate <strong>the</strong> microstructures <strong>of</strong> 

select coatings after deposition <strong>and</strong> annealing. The 3D-APT <strong>and</strong> TEM samples were prepared 

using a focused ion beam in situ lift-out (FIB-INLO) procedure [21, 22]. The 3D-APT 

<strong>investigation</strong>s were conducted with a local electrode atom probe (Imago LEAP 3000X). The 

conditions used for <strong>the</strong> analyses were 70K <strong>and</strong> an evaporation rate <strong>of</strong> 0.7-2.0%. Compositional 

information was acquired using <strong>the</strong> IVAS s<strong>of</strong>tware package. 

26

4.3. Results <strong>and</strong> Discussion 

4.3.1. As-Deposited Coatings 

The coatings produced for this study were deposited using sputtering conditions similar 

to those used by Ning et al. [18, 23] to produce NiAl-(0.1-0.6) at.% Hf coatings with columnar 

microstructures. One difference was <strong>the</strong> use <strong>of</strong> an unbalanced magnetron configuration as 

opposed to <strong>the</strong> balanced configuration utilized by Ning et al. Figure 1 shows representative 

cross-sectional SEM micrographs for as-deposited NiAl <strong>and</strong> NiAl-1Hf coatings. These images 

were collected using back-scattered electrons (BSE) to obtain compositional contrast. In 

agreement with prior observations [18, 24], <strong>the</strong> as-deposited coatings were dense <strong>and</strong> conformal 

to <strong>the</strong> underlying substrates. They were also featureless with regard to variations in atomic 

number contrast suggesting <strong>the</strong> formation <strong>of</strong> a single phase. The existence <strong>of</strong> a single phase was 

subsequently confirmed via XRD <strong>and</strong> TEM analyses, both <strong>of</strong> which are presented later in this 

document. Occasional leader defects <strong>and</strong> through thickness cracks, typical <strong>of</strong> thick coatings 

deposited by physical vapor deposition processes, were observed [25-29]. Both defect types 

were associated with irregularities on <strong>the</strong> substrate surfaces or with particles ejected from <strong>the</strong> 

sputtering targets during deposition. Figure 2 shows plan view TEM micrographs collected from 

as-deposited <strong>and</strong> annealed coatings. Similar to Ning et al. [30], <strong>the</strong> grains in <strong>the</strong> as-deposited 

coatings appeared to be single phase with no evidence <strong>of</strong> precipitation. It was not possible to 

accurately assess <strong>the</strong> as-deposited grain sizes from <strong>the</strong> collected micrographs. However, it was 

quite evident that <strong>the</strong> grains were much smaller than those produced by Ning et al. [18, 30]. The 

as-deposited coatings sometimes fractured during metallographic sample preparation revealing 

columnar grain morphologies that were consistent with zone T on Thornton’s structure zone 

model (SZM) [31]. The production <strong>of</strong> zone T microstructures with finer grain sizes as opposed 

27

to <strong>the</strong> zone 2 microstructures suggested by Ning et al. [18, 30] is attributed to small changes in 

deposition parameters <strong>and</strong> deposition hardware, in particular magnet configuration. 

Figure 3 shows composition pr<strong>of</strong>iles for Ni, Al, Cr, <strong>and</strong> Hf in <strong>the</strong> as-deposited coatings. 

Average coating compositions are summarized in Table 1. The Ni/Al ratio for NiAl-1Hf was 

calculated by assuming that Hf substituted for Al on <strong>the</strong> NiAl sublattice. As expected, no 

interdiffusion occurred between <strong>the</strong> coatings <strong>and</strong> substrates during sputtering. The compositions 

<strong>of</strong> <strong>the</strong> as-deposited coatings closely matched those <strong>of</strong> <strong>the</strong> sputtering targets <strong>and</strong> <strong>the</strong> 

corresponding Ni, Al, <strong>and</strong> Hf concentrations remained uniform through <strong>the</strong> coating crosssections. 

The Ni/Al ratios in <strong>the</strong> as-deposited coatings were near those <strong>of</strong> <strong>the</strong> sputtering targets. 

Figure 4 shows X-ray diffraction patterns collected from as-deposited NiAl <strong>and</strong> NiAl- 

1.0Hf coatings. Consistent with <strong>the</strong> prior observations <strong>of</strong> Ning et al. [18, 24], <strong>the</strong> as-deposited 

coatings, regardless <strong>of</strong> composition, were found to consist <strong>of</strong> a B2 structured β-NiAl phase. In 

general, <strong>the</strong> (110) direction had <strong>the</strong> highest intensity; however, some samples exhibited a strong 

intensity along <strong>the</strong> (211) direction. At this time, it is certain whe<strong>the</strong>r this change in preferred 

orientation <strong>of</strong> <strong>the</strong> coatings will have any impact on <strong>the</strong> oxidation behavior. Presuming <strong>the</strong> 

substitution <strong>of</strong> Hf for Al on <strong>the</strong> NiAl sublattice in <strong>the</strong> NiAl-1Hf coating [32-36], <strong>the</strong> lattice 

parameters calculated from <strong>the</strong> XRD data (Table 2) are consistent with values for slightly Ni-rich 

NiAl published by <strong>the</strong> International Centre for Diffraction Data (ICDD) [37-39]. Based upon <strong>the</strong> 

fact that Hf has a larger atomic radius than Ni <strong>and</strong> Al, it was expected that Hf <strong>additions</strong> to <strong>the</strong> 

coating would lead to lattice expansion ra<strong>the</strong>r than contraction [32-36]. Kitabjian <strong>and</strong> Nix [40], 

who made similar observations in slightly <strong>of</strong>f stoichiometric NiAl-Ti single crystals, speculated 

28

that <strong>the</strong> observation was related to <strong>the</strong> presence <strong>of</strong> trapped constitutional vacancies. This 

argument is consistent with <strong>the</strong> processing method used in this study. It is well known that 

sputter deposition generally leads to higher concentrations <strong>of</strong> lattice defects, vacancies in 

particular, than would be observed in a more conventionally processed materials <strong>and</strong> promotes 

<strong>the</strong> formation <strong>of</strong> highly supersaturated metastable alloys [41, 42]. 

4.3.2. Annealed Coatings 

Annealing at 1000°C for two or four hours (Figure 5) resulted in <strong>the</strong> formation <strong>of</strong> an 

interdiffusion zone (IDZ) between <strong>the</strong> coatings <strong>and</strong> <strong>the</strong> superalloy substrates. Typical <strong>of</strong> B2 

diffusion aluminides, <strong>the</strong> IDZ consisted <strong>of</strong> a mixture <strong>of</strong> globular high atomic number contrast 

precipitates dispersed in a lower atomic number contrast matrix. Though detailed analysis <strong>of</strong> <strong>the</strong> 

IDZ was not conducted in <strong>the</strong> present study, it is expected to consist <strong>of</strong> a β-NiAl or γ′-Ni 3 Al 

matrix phase <strong>and</strong> refractory metal-rich togologically close-packed (TCP) phase precipitates [8]. 

Interestingly, for coatings with <strong>the</strong> same thicknesses <strong>and</strong> similar Ni/Al ratios, <strong>the</strong> IDZ thickness 

qualitatively appeared to depend upon coating composition. For example, in binary NiAl 

coatings, annealing at 1000°C/2h produced a 3.69 μm thick IDZ (Figure 5(a)). Annealing at 

1000°/4h did not produce any significant increase in IDZ thickness, but did result in coarsening 

<strong>of</strong> <strong>the</strong> microstructure within <strong>the</strong> IDZ (Figure 5(b)). The NiAl-1Hf coatings, when subjected to 

<strong>the</strong> same heat treatment conditions (Figures 5(c) <strong>and</strong> 5(d)) produced noticeably thinner IDZs 

(i.e., 1.67 μm after two hours <strong>and</strong> 2.37 μm after four hours <strong>of</strong> annealing). This observation is 

similar to those <strong>of</strong> Ning et al. [18] who reported a reduction in IDZ thickness for a NiAl-0.6 

at.%Hf coating in comparison to binary NiAl. 

29

In <strong>the</strong> NiAl coatings, <strong>the</strong>re was an additional observation after four hours <strong>of</strong> annealing; a 

series <strong>of</strong> light gray precipitates within <strong>the</strong> coating. Based upon <strong>the</strong> composition <strong>of</strong> <strong>the</strong> coating 

after annealing, <strong>and</strong> <strong>the</strong> XRD results (presented in a later section <strong>of</strong> this document), <strong>the</strong>se 

precipitates were identified as γ′-Ni 3 Al. Closer inspection <strong>of</strong> <strong>the</strong> NiAl-1Hf coatings also 

revealed <strong>the</strong> formation <strong>of</strong> small precipitates within <strong>the</strong> coatings ranging in size from 10 to 80 nm 

after annealing (Figure 6). The precipitates appeared to form both within grain interiors <strong>and</strong> 

along grain boundaries. From this figure it also appears <strong>the</strong> columnar grain structure that 

developed during sputtering was retained. 

Even though <strong>the</strong> coatings were annealed in a reducing environment (i.e., Ar + 5%H 2 ), a 

thin <strong>the</strong>rmally grown oxide (TGO) layer was observed on <strong>the</strong> exterior surfaces <strong>of</strong> all <strong>of</strong> <strong>the</strong> 

specimens after annealing. This TGO was ~500 nm thick, <strong>and</strong> consisted primarily <strong>of</strong> aluminum 

oxide. In <strong>the</strong> NiAl-1Hf coatings, some high atomic number contrast phases were observed near 

<strong>the</strong> TGO/coating interface. These phases have not been quantitatively identified, but are 

presumed to be <strong>hafnium</strong> oxide (HfO 2 ). This oxygen is thought to have entered <strong>the</strong> environment 

through o-rings used to seal <strong>the</strong> annealing furnaces. These observations are consistent with those 

<strong>of</strong> Guo et. al [19] <strong>and</strong> Sun et al. [20] who investigated NiAl-Hf coatings produced via electronbeam 

physical vapor deposition (EB-PVD) <strong>and</strong> Priest [43] who investigated NiAl-Hf coatings 

produced via a pack cementation method. 

Figure 7 shows plan view TEM micrographs collected from annealed coatings. After 

annealing for as little as two hours, <strong>the</strong> grains became more equiaxed in appearance in both 

coatings (Figures 7(a) <strong>and</strong> 7(b)) suggesting <strong>the</strong> occurrence <strong>of</strong> some recovery <strong>and</strong> recrystallization 

30

[44]. In <strong>the</strong> NiAl coatings annealed for two hours, an average linear intercept grain size <strong>of</strong> 220.1 

± 35.1 nm was measured from <strong>the</strong> TEM micrographs. This grain size increased to 267.3 ± 34.3 

nm after annealing for four hours. In <strong>the</strong> NiAl-1Hf coatings, <strong>the</strong> grains were considerably 

smaller. In <strong>the</strong> NiAl-1Hf coating, an average grain size <strong>of</strong> 90.0 ± 20.5 nm was measured after 

annealing for four hours. The grain sizes observed in this study were much smaller than <strong>the</strong> ones 

reported by Ning et al. [18, 30, 45]. As noted previously, <strong>the</strong> smaller grain sizes are attributed to 

differences in coating composition plus <strong>the</strong> use <strong>of</strong> an unbalanced magnetron configuration. 

Unbalanced magnetrons increase ion <strong>and</strong> electron bombardment <strong>of</strong> <strong>the</strong> growing film resulting in 

more dense coatings <strong>and</strong> higher deposition rates [41, 42]. 

Figure 8 shows a STEM-HAADF image <strong>of</strong> <strong>the</strong> NiAl-1Hf coating after annealing at 

1000°C/4h. This image suggests <strong>the</strong> formation <strong>of</strong> two different types <strong>of</strong> precipitates, high atomic 

number contrast precipitates located both at grain boundaries <strong>and</strong> within grain interiors, <strong>and</strong> 

lower atomic number contrast precipitates <strong>of</strong>ten located along grain boundaries. EDS spectra 

collected in <strong>the</strong> TEM suggested that <strong>the</strong> largest grain boundary precipitates were <strong>the</strong> equilibrium 

Heusler (Ni 2 AlHf) phase, which agrees with published phase equilibria data [46]. However, in 

some instances relatively large carbon peaks were observed in <strong>the</strong> EDS spectra, suggesting <strong>the</strong> 

formation <strong>of</strong> some sort <strong>of</strong> carbide phase. 

Chemical compositions were measured on cross-sections <strong>of</strong> <strong>the</strong> as-deposited <strong>and</strong> 

annealed coatings. Figure 9 shows composition pr<strong>of</strong>iles for Ni, Al, Cr, <strong>and</strong> Hf in <strong>the</strong> annealed 

coatings. Annealing for four hours resulted in an increase in <strong>the</strong> Ni concentrations in both 

coatings coupled with decreased Al contents <strong>and</strong> significant increases in <strong>the</strong> Co <strong>and</strong> Cr contents. 

31

For example, in <strong>the</strong> NiAl coating <strong>the</strong> Ni content increased from an as-deposited level <strong>of</strong> 52.05 

at.% to 56.50 at.% after annealing for four hours at 1000°C. Similarly <strong>the</strong> Al content decreased 

from 48.68 at.% to 40.34 at.% while Co <strong>and</strong> Cr increased to 1.60 at.% <strong>and</strong> 1.53 at.%. In <strong>the</strong> 

NiAl-1Hf coating, Ni increased from 52.29 at.% to 53.44 at.%, Al decreased from 45.37 at.% to 

43.19 at.%, <strong>and</strong> Hf decreased from 1.10 at.% to 0.85 at.%. The Co <strong>and</strong> Cr contents increased to 

1.09 at.% <strong>and</strong> 1.33 at.% respectively. These observations are analogous to those <strong>of</strong> Ning et al. 

[18]. 

The observed changes in composition caused by annealing are driven by <strong>the</strong> composition 

gradient that exists between <strong>the</strong> coating <strong>and</strong> <strong>the</strong> substrate. The CMSX-4 substrates used in this 

study have lower Al contents than <strong>the</strong> sputtered coatings. This composition difference provides 

a driving force for interdiffusion <strong>of</strong> elements between <strong>the</strong> substrates <strong>and</strong> <strong>the</strong> sputtered coatings 

during high temperature annealing <strong>and</strong> oxidation. One particularly interesting observation for 

<strong>the</strong> NiAl-1Hf coatings was <strong>the</strong> existence <strong>of</strong> a Hf gradient through <strong>the</strong> coating thickness after 

annealing. For example, an apparent Hf concentration <strong>of</strong> ~0.95 at.% was measured 3 μm from 

<strong>the</strong> free surface versus ~0.8 at.% at <strong>the</strong> coatings center. At <strong>the</strong> coating/IDZ interface a 

composition in excess <strong>of</strong> 1.0 at.% was measured. Similar observations <strong>of</strong> variations in reactive 

element content have been made by o<strong>the</strong>r studies [18-20, 24, 43, 47-49]. This is consistent with 

<strong>the</strong> observations <strong>of</strong> Pint [12] who reported that REs would preferentially segregate to <strong>the</strong> free 

surface or coating/substrate interfaces during <strong>the</strong>rmal exposure. 

An additional observation was <strong>the</strong> apparent difference in IDZ thickness, which was 

analogous to <strong>the</strong> observations <strong>of</strong> Ning et al. [18]. O<strong>the</strong>r observations <strong>of</strong> reduced coating <strong>and</strong>/or 

32

IDZ thickness have been reported for diffusion aluminide <strong>and</strong> β-NiAl overlay coatings [50-54]. 

Based upon <strong>the</strong>se observations, it is tempting to attribute <strong>the</strong> differences in IDZ thickness solely 

to <strong>the</strong> presence <strong>of</strong> Hf, however, o<strong>the</strong>r factors must be considered. In general, <strong>the</strong> IDZ forms as a 

result <strong>of</strong> <strong>the</strong> conversion <strong>of</strong> <strong>the</strong> original γ+γ′ microstructure <strong>of</strong> <strong>the</strong> superalloy to β-NiAl; a change 

that is induced by <strong>the</strong> diffusional loss <strong>of</strong> Al from <strong>the</strong> coating due to inward diffusion to <strong>the</strong> 

substrate or <strong>the</strong> outward diffusion <strong>of</strong> Ni from <strong>the</strong> substrate into <strong>the</strong> coating [8]. The actual 

interdiffusion mechanisms associated with coating/substrate interdiffusion remain a subject <strong>of</strong> 

serious debate [8]; however, this interdiffusion is also accompanied by <strong>the</strong> rapid migration <strong>of</strong> 

elements from <strong>the</strong> substrates into <strong>the</strong> coating, most notably Co, Cr, W, <strong>and</strong> Ta. It is also well 

known that reactive elements, when present in <strong>the</strong> substrate, also diffuse rapidly to, <strong>and</strong> 

sometimes through, <strong>the</strong> coatings during annealing <strong>and</strong>/or service providing significant lifetime 

improvements [55-61]. These elements are additionally reported to reduce interdiffusion when 

<strong>the</strong>y form discrete RE-containing phases at <strong>the</strong> coating/substrate interface [50, 51, 62]. 

Though a detailed, quantitative <strong>investigation</strong> <strong>of</strong> <strong>the</strong> influences <strong>of</strong> coating composition on 

coating/substrate interdiffusion has not been undertaken, it is reasonable to hypo<strong>the</strong>size that Hf 

does in some way influence <strong>the</strong> activities <strong>of</strong> Ni <strong>and</strong> Al in much <strong>the</strong> same way as it does during 

oxidation, which in turn alters <strong>the</strong> stoichiometry, phase compositions, activities <strong>and</strong> diffusivities 

<strong>of</strong> <strong>the</strong> elements constituting <strong>the</strong> coating. Indeed Perez et al. [63], who investigated interdiffusion 

behavior in NiAl/superalloy diffusion couples, showed that superalloys containing higher 

concentrations <strong>of</strong> Cr, Mo, <strong>and</strong> Ti showed higher <strong>effect</strong>ive Al diffusivity while superalloys 

containing Ta, W <strong>and</strong> Al showed lower <strong>effect</strong>ive Al diffusivity. Fur<strong>the</strong>rmore, numerous 

computational <strong>and</strong> experimental studies have also detailed <strong>the</strong> influences <strong>of</strong> NiAl composition 

33

<strong>and</strong> stoichiometry on Ni <strong>and</strong> Al diffusion coefficients [64, 65]. Fur<strong>the</strong>r quantitative 

<strong>investigation</strong>s <strong>of</strong> phase chemistries before <strong>and</strong> after annealing are needed to completely establish 

<strong>the</strong> influences <strong>of</strong> Hf. 

Figure 10 shows X-ray diffraction patterns collected from annealed NiAl <strong>and</strong> NiAl-1.0Hf 

coatings. After annealing at 1000°C, several additional XRD peaks were observed. In <strong>the</strong> 

specimen coated with NiAl, no additional peaks were observed after two hours <strong>of</strong> annealing at 

1000°C. However, after four hours <strong>of</strong> annealing, small peaks γ′-Ni 3 Al <strong>and</strong> θ-Al 2 O 3 peaks were 

observed. In contrast, no γ′ peaks were observed in <strong>the</strong> NiAl-1Hf coatings; however, peaks 

corresponding to <strong>the</strong> α-Al 2 O 3 <strong>and</strong> θ-Al 2 O 3 phases were observed after as little as two hours <strong>of</strong> 

annealing. After annealing, <strong>the</strong> lattice parameters for <strong>the</strong> β-NiAl phase in both coatings 

decreased slightly (Table 2). 

4.3.3. Three-Dimensional Atom Probe Tomography (3D-APT) 

A series <strong>of</strong> 3D-APT <strong>investigation</strong>s were conducted to investigate <strong>hafnium</strong> solubility 

within <strong>the</strong> β-NiAl phase <strong>and</strong> in an attempt to identify some <strong>of</strong> <strong>the</strong> precipitates that formed in <strong>the</strong> 

annealed NiAl-1Hf coatings. Figure 11 shows representative 3-D atom maps <strong>of</strong> NiAl-1Hf after 

annealing at 1000°C/4h. These images clearly show <strong>the</strong> formation <strong>of</strong> Hf-rich precipitates within 

a β-NiAl matrix. A surprise was <strong>the</strong> presence <strong>of</strong> carbon within <strong>the</strong> specimens. Using <strong>the</strong> IVAS 

s<strong>of</strong>tware package, <strong>the</strong> average <strong>hafnium</strong> concentration within <strong>the</strong> β-NiAl phase was determined to 

be 0.23 ± 0.01 at.%, which is consistent with observations <strong>of</strong> Larson <strong>and</strong> Miller [66]. In this 

particular specimen, <strong>the</strong> average carbon concentration was determined to be 711 ± 42 ppm after 

annealing. Figure 12 shows a two-dimensional isosurface reconstruction <strong>and</strong> a one-dimensional 

34

(1-D) element concentration pr<strong>of</strong>ile analyzed through an isosurface reconstruction <strong>of</strong> one <strong>of</strong> <strong>the</strong> 

precipitates shown in Figure 11. The concentration pr<strong>of</strong>iles showed that <strong>the</strong>re was little or no Ni 

or Al within <strong>the</strong> precipitates <strong>and</strong> that <strong>the</strong> precipitates consisted almost entirely <strong>of</strong> Hf <strong>and</strong> C in 

equal proportions. 

To identify <strong>the</strong> potential source <strong>of</strong> carbon within <strong>the</strong> specimens, several NiAl-1Hf 

specimens were annealed inside <strong>of</strong> a quartz tube in a horizontal tube furnace under flowing 

Ar+5%H 2 . In <strong>the</strong>se specimens, respective Hf <strong>and</strong> C concentrations <strong>of</strong> 0.24 ± 0.11 at.% <strong>and</strong> 53 ± 

33 ppm were determined which suggests that high levels <strong>of</strong> carbon could be incorporated by 

annealing inside <strong>of</strong> a graphite element furnace resulting in <strong>the</strong> formation <strong>of</strong> carbide precipitates 

provided a suitable concentration <strong>of</strong> carbide formers is present. This observation is consistent 

with work <strong>of</strong> Darolia et al. [67] who showed that nanometer-sized oxide <strong>and</strong> carbide precipitates 

could be formed in NiAl coatings due to residual impurities within <strong>the</strong> deposition chamber. 

They additionally reported that <strong>the</strong>se dispersed carbide <strong>and</strong> oxide particles could improve coating 

life by increasing <strong>the</strong> resistance <strong>of</strong> <strong>the</strong> coating to rumpling [67, 68]. Jha <strong>and</strong> co-workers [69, 70] 

have shown that dispersed HfC particles can significantly improve creep resistance <strong>of</strong> bulk NiAl 

which would lead to increased resistance to rumpling. However, it is well established that 

elevated carbon concentrations can be detrimental to <strong>the</strong> <strong>effect</strong>iveness <strong>of</strong> reactive element 

<strong>additions</strong> [10, 59]. Thus, it appears that in <strong>the</strong> presence <strong>of</strong> high carbon concentrations, higher 

reactive element contents can be used to <strong>of</strong>fset <strong>the</strong> detrimental influences <strong>of</strong> carbon <strong>and</strong> provide 

additional benefits through <strong>the</strong> formation <strong>of</strong> nanometer-sized carbide precipitates. 

35

4.4. Conclusions 

In this study, NiAl-1.0 at.% Hf bond coatings were produced using DC magnetron 

sputtering techniques to yield a zone T microstructure. The as-deposited coatings formed a 

metastable B2 solid solution. Upon annealing at 1000°C, small precipitates formed along grain 

boundaries <strong>and</strong> within grain interiors on dislocation lines. Images collected using SEM <strong>and</strong> 

TEM show <strong>the</strong> precipitates were heterogeneously distributed in <strong>the</strong> coating with <strong>the</strong> larger 

precipitates forming within grain interiors. In high carbon content coatings, 3D-APT studies 

suggested <strong>the</strong> formation <strong>of</strong> HfC precipitates with <strong>the</strong> carbon coming from <strong>the</strong> graphite elements 

in <strong>the</strong> annealing furnace. A change to a graphite-free furnace resulted in little pickup <strong>and</strong> in no 

observable HfC formation. 

Acknowledgements 

The authors acknowledge <strong>the</strong> support <strong>of</strong> <strong>the</strong> National Science Foundation (NSF) under 

Award No. DMR-0504950 <strong>and</strong> <strong>the</strong> National Aeronautics <strong>and</strong> Space Administration (NASA) 

under contract NN-X08AT21H. The use <strong>of</strong> <strong>the</strong> facilities supported by <strong>the</strong> Central Analytical 

facility at <strong>the</strong> University <strong>of</strong> Alabama <strong>and</strong> <strong>the</strong> Materials Diagnostics Laboratory at NASA’s 

Marshall Space Flight Center in Huntsville, Alabama is also gratefully acknowledged. 

36

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on interdiffusion coefficients <strong>of</strong> Ni <strong>and</strong> Al atoms in β-NiAl phase <strong>of</strong> aluminide coatings," 

Journal <strong>of</strong> Materials Science <strong>and</strong> Technology 20 (2004) 196-198. 

[66] D. J. Larson, M. K. Miller, "Atom Probe Field-Ion Microscopy Characterization <strong>of</strong> 

Nickel <strong>and</strong> Titanium Aluminides," Materials Characterization 44 (2000) 159-176. 

41

[67] R. Darolia, W. S. Walston, J. A. Pfaendtner, B. A. R. Boutwell, I. Spitsberg, J. A. Ruud, 

"Nickel Aluminide Coating <strong>and</strong> Coating Systems Formed Therewith," United States 

patent # 6,887,589, 2005. 

[68] R. Darolia, J. D. Rigney, G. H. Marijnissen, E. R. I. C. Vergeldt, A. B. Kloosterman, 

"Streng<strong>the</strong>ned Bond Coats for Thermal Barrier Coatings," United States patent # 

7,172,820 B2, 2007. 

[69] S. C. Jha, R. Ray, J. D. Whittenberger, "Carbide-dispersion-streng<strong>the</strong>ned B2 NiAl," 

Materials Science <strong>and</strong> Engineering A 119 (1989) 103-111. 

[70] S. C. Jha, R. Ray, D. J. Gaydosh, "Dispersoids in Rapidly Solidified B2 Nickel 

Aluminides," Scripta Metallurgica 23 (1989) 805-810. 

42

Table 1. Nominal chemical compositions <strong>of</strong> sputtering targets. 

Alloy Condition Method Ni Al Cr Co Hf 

NiAl Nominal --- 51 49 --- --- --- 

As-deposited WDS 52.05 ± 0.95 47.68 ± 0.97 0.14 ± 0.25 0.08 ± 0.04 ± 0.02 

1000°C/2h WDS 

1000°C/4h WDS 56.50 ± 40.34 ± 1.53 ± 1.598 ± 0.04 ± 

NiAl-1.0Hf Nominal --- 51 48 --- --- 1.0 

As-deposited WDS 53.29 ± 0.49 45.37 ± 0.51 0.08 ± 0.06 0.08 ± 0.13 1.10 ± 0.05 

1000°C/2h WDS 55.99 ± 0.41 38.67 ± 0.59 2.04 ± 0.25 2.57 ± 0.15 0.73 ± 0.16 

1000°C/4h WDS 53.44 ± 0.79 43.19 ± 0.72 1.33 ± 0.43 1.09 ± 0.30 0.85 ± 0.14 

Table 2. Lattice parameters for <strong>the</strong> β-NiAl phase. 

Alloy Annealing Time Lattice Parameter (nm) Grain Size (nm) 

NiAl As-deposited 0.2889 ± 0.0007 

1000°C/2h 0.2888 ± 0.0007 220.1 ± 35.1 

1000°C/4h 0.2884 ± 0.0007 267.3 ± 34.3 

NiAl-1Hf As-deposited 0.2882 ± 0.0001 

1000°C/2h 0.2879 ± 0.0001 

1000°C/4h 0.2879 ± 0.0002 90 ± 20.5 

43

Figure 1. Representative SEM micrographs <strong>of</strong> as-deposited coatings: (a) NiAl <strong>and</strong> (b) NiAl- 

1Hf. 

44

Figure 2. Plan view TEM images collected from <strong>the</strong> NiAl coatings: (a) as-deposited, (b) 

annealing at 1000°C for two hours, <strong>and</strong> (c) annealing at 1000°C for four hours. 

45

Figure 3. Elemental composition pr<strong>of</strong>iles for as-deposited coatings: (a) NiAl <strong>and</strong> (b) NiAl-1Hf. 

46

NiAl 

NiAl 

NiAl 

NiAl 

NiAl 

NiAl 

(100) 

(110) 

(111) 

(200) 

(211) 

(220) 

Intensity (abu) 

NiAl-1Hf 

NiAl 

30 40 50 60 70 80 90 100 

Degrees (2θ) 

Figure 4. X-ray diffraction patterns for as-deposited coatings. 

47

Figure 5. Representative SEM micrographs for: (a) NiAl annealed for two hours, (b) NiAl 

annealed for four hours, (c) NiAl-1Hf annealed for two hours, <strong>and</strong> (d) NiAl-1Hf annealed for 

four hours. 

48

Figure 6. Backscattered SEM image <strong>of</strong> NiAl-1Hf showing <strong>the</strong> development <strong>of</strong> <strong>hafnium</strong>-rich 

precipitates after annealing. 

49

Figure 7. Plain view TEM images from coatings after annealing for four hours: (a) NiAl <strong>and</strong> (b) 

NiAl-1Hf. 

50

Figure 8. STEM-HAADF image <strong>of</strong> NiAl-1Hf after annealing for four hours at 1000°C. 

51

Figure 9. Elemental composition pr<strong>of</strong>iles for annealed coatings: (a) Nickel, (b) Chromium, (c) 

Aluminum, <strong>and</strong> (d) Hafnium. 

52

β-NiAl 

θ-Al 2 

O 3 

γ'-Ni 3 

Al 

Intensity (abu) 

(d) 

(c) 

(b) 

(a) 

30 40 50 60 70 80 90 100 

Degrees (2θ) 

Figure 10. X-ray diffraction patterns for annealed coatings: (a) NiAl annealed two hours, (b) 

NiAl annealed four hours, (c) NiAl-1Hf annealed two hours, <strong>and</strong> (d) NiAl-1Hf annealed four 

hours. 

53

Figure 11. 3D-APT data for NiAl-1Hf after annealing for four hours: (a) all elements within <strong>the</strong> 

sample <strong>and</strong> (b) only <strong>hafnium</strong> <strong>and</strong> carbon. 

54

Figure 12. (a) Two-dimensional isosurface reconstruction <strong>of</strong> a NiAl-1Hf coating that was 

analyzed using 3D-APT, <strong>and</strong> (b) corresponding concentration pr<strong>of</strong>ile <strong>of</strong> <strong>the</strong> precipitate. 

55

CHAPTER V 

INFLUENCES OF CHROMIUM AND HAFNIUM ADDITIONS ON THE 

MICROSTRUCTURES OF β-NIAL COATINGS DEPOSITED ON SUPERALLOY 

SUBSTRATES 

M.A. Bestor 1 , J.P. Alfano 1 , B.T. Hazel 2 , M.L. Weaver 1 

1-The University <strong>of</strong> Alabama, Department <strong>of</strong> Metallurgical & Materials Engineering, Box 

870202, Tuscaloosa, AL 35487-0202, USA 

2- General Electric Aviation; One Neumann Way; Cincinnati, OH, 45215, USA 

Keywords: Nickel aluminides, bond coatings, reactive element doping, DC magnetron 

sputtering 

Abstract 

In this study, <strong>the</strong> <strong>effect</strong>s <strong>of</strong> annealing <strong>and</strong> chemistry were investigated for NiAl coatings 

containing 1.0 at% Hf <strong>and</strong>/or 5 at% Cr have been deposited onto René N5 substrates via direct 

current magnetron sputtering. Investigations were conducted using X-ray diffraction, scanning 

electron microscopy, electron probe microanalysis, <strong>and</strong> transmission electron microscopy in 

order to study <strong>the</strong> influences <strong>of</strong> annealing on <strong>the</strong> microstructures <strong>and</strong> properties <strong>of</strong> <strong>the</strong> coatings. 

Post-deposition annealing in <strong>the</strong> NiAl-1Hf produced nanometer sized precipitates that were high 

in atomic number contrast; however, <strong>the</strong> Cr-containing coatings also produced large Cr-rich 

precipitates after two hours <strong>of</strong> annealing. As annealing time was increased, coarsening <strong>of</strong> <strong>the</strong> 

coating grains <strong>and</strong> precipitates was observed. TEM analyses indicate that most <strong>of</strong> <strong>the</strong> 

precipitates in <strong>the</strong> coating are β′-Ni 2 AlHf with <strong>the</strong> incorporation <strong>of</strong> additional precipitates. Atom 

probe tomography results suggest that many <strong>of</strong> <strong>the</strong>se o<strong>the</strong>r precipitates are α-Cr. The results 

56

have been analyzed <strong>and</strong> discussed relative to previous research on sputter deposited NiAl-Hf 

coatings. 

5.1. Introduction 

Though Ni-based superalloys exhibit high strengths <strong>and</strong> creep resistances, when used in 

gas turbine engines, <strong>the</strong>y <strong>of</strong>ten require environmental barrier (EB) or <strong>the</strong>rmal barrier coating 

(TBC) systems [1-4]. TBCs are comprised <strong>of</strong> three main parts. The top layer or top coat, 

usually a Y-stabilized ZrO 2 , is used to lower <strong>the</strong> surface temperature <strong>of</strong> a component by as much 

as 200°C from <strong>the</strong> operating temperature. The ceramic top coat is porous in nature <strong>and</strong> does not 

inhibit <strong>the</strong> transport <strong>of</strong> oxygen to <strong>the</strong> superalloy. Between <strong>the</strong> superalloy <strong>and</strong> <strong>the</strong> ceramic topcoat 

is a metallic bond coat. This bond coat provides <strong>the</strong> adhesive properties that are necessary to 

attach <strong>the</strong> top coat to <strong>the</strong> underlying superalloy while also providing an oxidation barrier. This is 

accomplished by <strong>the</strong> formation <strong>of</strong> a thin <strong>the</strong>rmally grown oxide (TGO) layer between <strong>the</strong> 

metallic bond coat <strong>and</strong> <strong>the</strong> ceramic top coat. The thickness <strong>and</strong> growth rate <strong>of</strong> this oxide layer 

are critical to <strong>the</strong> service lifetime <strong>of</strong> <strong>the</strong> TBC [2,4,5]. Thus, <strong>the</strong> development <strong>of</strong> bond coat alloys 

capable <strong>of</strong> forming slow growing, <strong>the</strong>rmally stable TGOs is a high priority. 

For years, diffusion aluminides <strong>and</strong> MCrAlY (where M = Fe, Co, <strong>and</strong>/or Ni) overlay 

coatings have been used with great success as EBs <strong>and</strong> as bond coats in TBCs [4,5]. Though 

<strong>the</strong>se alloys have generally performed well, <strong>the</strong>y were never really designed or developed to 

increase TBC durability, while meeting o<strong>the</strong>r performance requirements [6,7]. As such, <strong>the</strong>re is 

great interest in <strong>the</strong> development <strong>of</strong> new coating alloys, specifically engineered to meet existing 

<strong>and</strong> future requirements in advanced turbine systems. 

57

Investigations by Pint et al. [8-10], Nesbitt et al. [11], <strong>and</strong> Warnes [12] suggest that NiAlbased 

alloys containing small quantities <strong>of</strong> <strong>the</strong> so called reactive elements (RE) Y, Hf or Zr 

added ei<strong>the</strong>r in conjunction with or in place <strong>of</strong> Pt <strong>of</strong>fer <strong>the</strong> potential for improved TBC lifetimes 

in comparison to state-<strong>of</strong>-<strong>the</strong> art coatings. The RE <strong>effect</strong>s in <strong>the</strong>se alloys have been summarized 

as an improvement <strong>of</strong> <strong>the</strong> alumina scale adhesion <strong>and</strong> a reduction <strong>of</strong> <strong>the</strong> TGO growth rate by <strong>the</strong> 

segregation <strong>of</strong> RE ions to <strong>the</strong> scale grain boundaries <strong>and</strong> alloy-scale interface [10]. In general, 

<strong>the</strong> RE <strong>additions</strong> to <strong>the</strong> aluminide bond coats are maintained around 0.05 at.% which is near or 

well below <strong>the</strong>ir reported solubility limits in β-NiAl <strong>and</strong> (Ni,Pt)Al. 

Recently, it was discovered that β-NiAl based bond coats containing up to 2 at% Hf 

<strong>and</strong>/or Zr <strong>and</strong> up to 5 at% Cr can exhibit oxidation resistance that is comparable to modern 

(Ni,Pt)Al diffusion coatings [13]. Based upon conventional wisdom, it was expected that alloys 

containing such high RE <strong>and</strong> Cr concentrations would fail catastrophically due to internal 

oxidation [14-16]. Fur<strong>the</strong>rmore, it has been shown in model bulk alloys that Cr, when added to 

β-NiAl in conjunction with 0.05 at.% Hf, led to <strong>the</strong> formation <strong>of</strong> second phase precipitates which 

appeared to catalyze scale spallation [17]. However, <strong>the</strong>se new bond coat alloys, which were 

largely derived from research on monolithic β-NiAl intermetallics in <strong>the</strong> early 1990’s, have been 

shown to yield increased TBC spallation life <strong>and</strong> reduced substrate consumption rates in 

comparison to Pt-containing diffusion aluminides [13,17-19]. Bennett <strong>and</strong> Slo<strong>of</strong> [20] reported 

that <strong>additions</strong> <strong>of</strong> Y or Zr to a bulk β-NiAl alloy containing ~4 at.% Cr resulted in improved TGO 

adhesion in comparison to RE-free, but Cr containing alloys. 

58

Underst<strong>and</strong>ing <strong>the</strong> microstructural evolution <strong>of</strong> coating systems is pivotal to <strong>the</strong> 

development <strong>of</strong> <strong>the</strong> next generation <strong>of</strong> coating technology. This study focuses on <strong>the</strong> influences 

<strong>of</strong> RE (i.e., Hf or Zr) <strong>and</strong> Cr <strong>additions</strong> in excess <strong>of</strong> <strong>the</strong>ir solubility limits on <strong>the</strong> microstructures 

<strong>and</strong> properties <strong>of</strong> β-NiAl based overlay coatings deposited on superalloy substrates. Previously, 

studies have detailed influences <strong>of</strong> high Hf contents on <strong>the</strong> microstructures <strong>and</strong> oxidation 

resistance <strong>of</strong> β-NiAl based coatings produced via balanced <strong>and</strong> unbalanced, direct current (DC) 

magnetron sputtering [21-24]. The purpose <strong>of</strong> this short paper is to exp<strong>and</strong> upon that previous 

work by detailing <strong>the</strong> influences <strong>of</strong> Cr <strong>and</strong> relatively high Hf contents on <strong>the</strong> microstructures <strong>of</strong> 

β-NiAl based overlay bond coatings produced via unbalanced DC magnetron sputtering. 

5.2. Experimental 

Alloy sputtering targets for this study were purchased from ACI Alloys, Inc., San Jose, 

CA <strong>and</strong> Sophisticated Alloys, Inc., Butler, PA. The nominal <strong>and</strong> experimentally determined 

target compositions are presented in Table 1. Coatings ~30μm thick were deposited via direct 

current (DC) magnetron sputtering onto René N5 substrates using an AJA International, Inc. 

ORION 4 sputtering system with a sputter-down configuration. Argon was used as a working 

gas. The deposition power was set at 300 W, <strong>and</strong> <strong>the</strong> working gas pressure was set at 1.33 Pa. 

Prior to deposition, <strong>the</strong> substrates were heated to 400°C or 650°C. These parameters were 

optimized to yield a zone T microstructure. All <strong>of</strong> <strong>the</strong> substrates were coated on <strong>the</strong> two primary 

faces for subsequent microstructural analyses, annealing, <strong>and</strong> oxidation experiments. 

The as-deposited coatings were annealed at 1000°C from one to four hours in a flowing 

mixture <strong>of</strong> Ar + 5% H 2 to improve adhesion between <strong>the</strong> coating <strong>and</strong> substrate, reduce sputtering 

59

induced defects, <strong>and</strong> alleviate any processing induced residual stresses. Annealing was done in a 

horizontal tube furnace inside <strong>of</strong> a quartz tube. 

The chemical compositions <strong>of</strong> <strong>the</strong> sputtered coatings were determined via wavelength 

dispersive X-ray spectroscopy (WDS) <strong>and</strong> three-dimensional atom probe tomography (3D-APT). 

X-ray diffraction (XRD), scanning electron microscopy (SEM) <strong>and</strong> transmission electron 

microscopy (TEM) were used for phase identification <strong>and</strong> to elucidate <strong>the</strong> microstructures <strong>of</strong> 

select coatings after deposition <strong>and</strong> annealing. The TEM samples were prepared using a focused 

ion beam in situ lift-out (FIB-INLO) procedure [25]. 

5.3. Results <strong>and</strong> Discussion 

5.3.1. Microstructure 

In <strong>the</strong> present study, <strong>the</strong> coatings were deposited using <strong>the</strong> same conditions that were 

used to produce NiAl-1.0Hf at% coatings [24]. Figure 1 presents representative cross-sectional 

SEM images showing <strong>the</strong> as-deposited <strong>and</strong> annealed coatings. Images from <strong>the</strong> as-deposited 

coatings (Figures 1 a. <strong>and</strong> c.) show that <strong>the</strong> coatings are dense <strong>and</strong> adherent to <strong>the</strong> underlying 

substrates. These coatings exhibited columnar zone T microstructures <strong>and</strong> were featureless with 

regard to variations in atomic number contrast in both SEM <strong>and</strong> TEM studies. Fur<strong>the</strong>rmore, <strong>the</strong> 

XRD spectra only contained peaks for β-NiAl suggesting <strong>the</strong> formation <strong>of</strong> a single phase. 

Occasional leader defects <strong>and</strong> through thickness cracks, especially with <strong>the</strong> NiAlCrHf coatings, 

were observed [26-30]. Both defect types are typical with physical vapor deposited coatings <strong>and</strong> 

can attributed to substrate surface irregularities or particles ejected from <strong>the</strong> sputtering targets 

during deposition. Post-deposition annealing at 1000°C from one to four hours produced an 

interdiffusion zone (IDZ) between <strong>the</strong> substrate <strong>and</strong> <strong>the</strong> coating. The IDZ forms as a result <strong>of</strong> <strong>the</strong> 

60

difference in chemistries between <strong>the</strong> coating <strong>and</strong> superalloy. Primarily it is <strong>the</strong> large 

concentration gradient with aluminum that provides <strong>the</strong> driving force for diffusion to occur 

between <strong>the</strong> substrate <strong>and</strong> coating. However, this IDZ also aids in providing additional adhesion 

<strong>of</strong> <strong>the</strong> coating to <strong>the</strong> substrate through <strong>the</strong> formation <strong>of</strong> a strong physical bond. The IDZ consists 

<strong>of</strong> a mixture <strong>of</strong> globular high atomic number contrast precipitates in a lower atomic number 

contrast matrix. There is no observable difference in <strong>the</strong> thickness <strong>of</strong> <strong>the</strong> IDZ from <strong>the</strong> NiAl-1Hf 

<strong>and</strong> NiAl-5Cr-1Hf. The NiAlCrHf (Figure 1 d.) sample contains lighter phases in <strong>the</strong> coating. 

These will be addressed later in <strong>the</strong> paper. 

After annealing, a thin <strong>the</strong>rmally grown oxide (TGO) layer was observed on <strong>the</strong> exterior 

surfaces <strong>of</strong> all <strong>of</strong> <strong>the</strong> specimens. The oxygen source that resulted in TGO formation was traced 

to a leak in one <strong>of</strong> <strong>the</strong> furnace bulkheads. The TGO was typically less than 500 nm in thickness 

<strong>and</strong> was found to consist <strong>of</strong> a mixture <strong>of</strong> θ-Al 2 O 3 <strong>and</strong> α-Al 2 O 3 . High atomic number contrast 

phases were detected near <strong>the</strong> TGO/coating interface. While <strong>the</strong>se phases have not been 

quantitatively identified, <strong>the</strong>y are presumed to be <strong>hafnium</strong> oxide (HfO 2 ). These observations are 

consistent with those <strong>of</strong> Guo et al. [31] <strong>and</strong> Sun et al. [32] who investigate NiAl-Hf coatings 

produced via electron-beam physical vapor deposition (EB-PVD). 

Figure 2 shows <strong>the</strong> x-ray diffraction (XRD) patterns collected from <strong>the</strong> NiAl-1Hf <strong>and</strong> 

NiAlCrHf coatings. The as-deposited coatings for both samples appear to be single phase, even 

though <strong>the</strong>y are well above <strong>the</strong> solubility limits for Hf <strong>and</strong> Cr. After annealing, both samples 

exhibit <strong>the</strong> formation <strong>of</strong> additional phases, mainly θ-Al 2 O 3 <strong>and</strong> α-Al 2 O 3 . The θ-Al 2 O 3 <strong>and</strong> α- 

Al 2 O 3 (i.e., alumina) peaks appear to increase in size <strong>and</strong> number as annealing time is increased. 

61

As aluminum is consumed to form alumina, <strong>the</strong> transition from β-NiAl to γ′-Ni 3 Al is typically 

observed; however, no γ’-Ni 3 Al peaks observed with any <strong>of</strong> <strong>the</strong> XRD patterns collected for <strong>the</strong>se 

samples. 

Figure 3 shows plan view TEM micrographs for <strong>the</strong> NiAl-1Hf <strong>and</strong> NiAlCr1Hf coatings 

that were annealed at 1000°C for four hours. After annealing, <strong>the</strong> grains became more equiaxed 

in appearance in both coatings (Figures 3 a. <strong>and</strong> b.) suggesting <strong>the</strong> occurrence <strong>of</strong> some recovery 

<strong>and</strong> recrystallization [33]. In <strong>the</strong> NiAl-1Hf coatings, a linear intercept grain size measurement <strong>of</strong> 

90.0 ± 20.5 nm. However, <strong>the</strong> NiAlCrHf coatings had a linear intercept grain size measurement 

<strong>of</strong> 162.48 ± 25.43 nm after annealing for four hours. The grain sizes observed with this study are 

considerably smaller than <strong>the</strong> ones reported by Ning et al. [21,22,32,34]. These smaller grain 

sizes are attributed to differences in <strong>the</strong> coating composition plus <strong>the</strong> use <strong>of</strong> an unbalanced 

magnet configuration as opposed to <strong>the</strong> balanced configuration used by Ning [21,22]. 

Unbalanced magnetrons increase ion bombardment <strong>of</strong> <strong>the</strong> growing film resulting in more dense 

coatings <strong>and</strong> higher deposition rates [35,36]. 

5.3.2. Chemistry 

Chemical compositions were measured using electron probe microanalysis (EPMA) on 

cross-sections <strong>of</strong> <strong>the</strong> as-deposited <strong>and</strong> annealed coatings. SEM spectral maps were also collected 

to indicate <strong>the</strong> locations <strong>of</strong> different elements within <strong>the</strong> coatings. The chemistries measured for 

<strong>the</strong> as-deposited <strong>and</strong> annealed coatings are presented in Table 1 along with <strong>the</strong> nominal 

sputtering target concentrations. Figure 4 shows composition pr<strong>of</strong>iles for Ni, Al, Cr, <strong>and</strong> Hf 

collected from <strong>the</strong> as-deposited <strong>and</strong> annealed coatings. There was no interdiffusion between <strong>the</strong> 

substrate <strong>and</strong> coating during deposition. The Ni <strong>and</strong> Al concentrations remain almost constant 

62

throughout <strong>the</strong> coating for both <strong>the</strong> NiAl-1Hf <strong>and</strong> NiAlCrHf samples. The Ni/Al ratios for <strong>the</strong> 

NiAl-1Hf <strong>and</strong> NiAlCrHf coatings were found to be close to those <strong>of</strong> <strong>the</strong> sputtering targets. After 

annealing, <strong>the</strong>re was a noticeable decrease in aluminum content with a corresponding increase in 

Cr <strong>and</strong> Co in <strong>the</strong> coatings. The aluminum concentration for <strong>the</strong> NiAl-1Hf sample decreased from 

<strong>the</strong> as-deposited state <strong>of</strong> 45.36 at.% to 43.12 at.% following annealing at 1000°C for four hours 

while <strong>the</strong> Cr <strong>and</strong> Co contents increased to 1.36 at.% <strong>and</strong> 1.10 at.% respectively. The NiAlCrHf 

exhibited a more significant loss in aluminum going from 48.53 at.% in <strong>the</strong> as-deposited case to 

40.50 at.% following annealing. The Cr concentration remained almost constant from <strong>the</strong> asdeposited 

amount <strong>of</strong> 4.8 at.% while <strong>the</strong> Co concentration increased to 1.16 at.%. The Ni 

concentration for <strong>the</strong> NiAl-1Hf coating was not observed to change appreciably; however, in <strong>the</strong> 

NiAlCrHf samples <strong>the</strong> Ni content increased from 45.74 at.% to 53.23 at.%. 

The composition changes that occurred during annealing can be explained by <strong>the</strong> 

concentration gradient that exists between <strong>the</strong> substrate <strong>and</strong> coating. The René N5 superalloy 

substrates that were used in this study contain a much lower concentration <strong>of</strong> aluminum 

compared to <strong>the</strong> sputter deposited coatings. This difference in composition provides <strong>the</strong> driving 

force for interdiffusion <strong>of</strong> elements from <strong>the</strong> coating <strong>and</strong> substrate to occur during exposure to 

high temperatures during annealing <strong>and</strong> oxidation. Also, <strong>the</strong>re are a couple <strong>of</strong> interesting 

observations from <strong>the</strong> concentration pr<strong>of</strong>iles. The Hf pr<strong>of</strong>ile shown in Figure 4 (d) collected 

from <strong>the</strong> NiAl-1Hf coating after annealing for four hours shows a variation through <strong>the</strong> thickness 

<strong>of</strong> <strong>the</strong> coating. It appears that <strong>the</strong> Hf enriches <strong>the</strong> interfaces <strong>of</strong> <strong>the</strong> coating <strong>and</strong> depletes <strong>the</strong> 

center <strong>of</strong> <strong>the</strong> coating. For example, <strong>the</strong> Hf concentrations near <strong>the</strong> surface <strong>and</strong> coating/substrate 

interfaces respectively were found to be ~0.75 at.% Hf <strong>and</strong> ~1.0 at.% Hf; however, <strong>the</strong> Hf 

63

concentration at <strong>the</strong> center <strong>of</strong> <strong>the</strong> coating was determined to be approximately 0.6 at.%. This 

behavior was observed in several specimens <strong>and</strong> is similar to those reported by o<strong>the</strong>rs 

[22,31,32,37-39]. This is consistent with dynamic segregation <strong>the</strong>ory as described by Pint et al. 

[10] which states that reactive elements should preferentially segregate to <strong>the</strong> free surface or 

coating/substrate interfaces during <strong>the</strong>rmal exposure, particularly under an oxygen potential 

gradient. In <strong>the</strong> case <strong>of</strong> <strong>the</strong> NiAlCrHf coatings, an apparent increase in Hf concentration at <strong>the</strong> 

coating/substrate interface was observed. The Hf concentration at this interface was found to be 

near ~1.1 at.% after annealing for two or four hours, while <strong>the</strong> concentration in <strong>the</strong> center <strong>of</strong> <strong>the</strong> 

coatings remained near ~0.8 at.%, which was equivalent to that observed in <strong>the</strong> as-deposited 

coatings. 

Additional elemental diffusion information was collected from annealed cross-sectional 

samples using SEM-EDS mapping. Figure 5 displays <strong>the</strong> spectral maps collected from <strong>the</strong> NiAl- 

1Hf coating for aluminum, oxygen, <strong>hafnium</strong> <strong>and</strong> <strong>chromium</strong>. The decoration <strong>of</strong> <strong>hafnium</strong> at <strong>the</strong> 

coating/substrate <strong>and</strong> coating/TGO interfaces is clearly visible. The TGO is primarily composed 

<strong>of</strong> aluminum oxide, which confirms <strong>the</strong> XRD results from Figure 2. Fur<strong>the</strong>rmore, <strong>the</strong> matrix <strong>of</strong> 

<strong>the</strong> IDZ has a high Al content mixed with Cr-rich phases also being present. While <strong>the</strong> Cr 

extends upward through <strong>the</strong> coating, it appears to have a higher concentration in <strong>the</strong> IDZ than in 

<strong>the</strong> coating. The phases in <strong>the</strong> IDZ also appear to be small in size <strong>and</strong> globular in nature. 

However, <strong>the</strong> IDZ <strong>of</strong> <strong>the</strong> NiAlCrHf coating (Figure 6) contains larger precipitates dispersed 

within <strong>the</strong> matrix that appear to be Cr-rich. In addition, <strong>the</strong> maps show that <strong>the</strong>re is a large 

concentration <strong>of</strong> Cr-rich precipitates within <strong>the</strong> coating. This supports <strong>the</strong> evidence from <strong>the</strong> 

EPMA concentration pr<strong>of</strong>iles that <strong>the</strong> Cr is retained in <strong>the</strong> coating following annealing. The 

64

segregation <strong>of</strong> Hf to <strong>the</strong> interfaces is not as clear with this coating as with <strong>the</strong> NiAl-1Hf. This 

suggests that <strong>the</strong> majority <strong>of</strong> <strong>the</strong> Hf is also retained in <strong>the</strong> coating. The implications <strong>of</strong> <strong>the</strong>se 

observations are potentially quite significant in that <strong>the</strong>y suggest that <strong>the</strong> incorporation <strong>of</strong> 

substrate elements into <strong>the</strong> coating can be inhibited by <strong>the</strong> simultaneous addition <strong>of</strong> Cr <strong>and</strong> Hf. 

This observation is in agreement with <strong>the</strong> recent reports <strong>of</strong> Hazel et al. [13] who reported 

reduced wall consumption in superalloy components coated with NiAl+Cr+Zr overlay coatings 

in comparison to platinum aluminides. Such a reduction could have a significant impact on 

superalloy <strong>and</strong> TBC service lifetimes <strong>and</strong> component repairability. More thorough analyses 

focusing on <strong>the</strong> chemical compositions <strong>and</strong> phase makeup within <strong>the</strong> IDZ <strong>and</strong> underlying 

superalloy substrate coupled with diffusion modeling are needed to quantatively identify <strong>the</strong> 

operative mechanisms. 

Figure 7 shows a STEM-HAADF image <strong>and</strong> EDS spectrum collected from <strong>the</strong> NiAlCrHf 

coating after annealing at 1000°C/4h. The EDS results confirmed <strong>the</strong> presence <strong>of</strong> Hf-rich 

precipitates with compositions consistent with <strong>the</strong> β′- Ni 2 AlHf phase. Additional lower atomic 

number contrast precipitates were also observed, however, it was not possible to obtain EDS 

results from <strong>the</strong>m. 

3D-APT <strong>investigation</strong>s were conducted to investigate <strong>the</strong> solubility <strong>of</strong> <strong>chromium</strong> <strong>and</strong> 

<strong>hafnium</strong> in <strong>the</strong> β-NiAl phase <strong>and</strong> to identify <strong>the</strong> precipitates that formed following annealing. 

Figure 8 shows a representative 3-D atom map <strong>and</strong> isosurface reconstruction for <strong>the</strong> NiAlCrHf 

coating after annealing at 1000°C/4h. The images clearly show <strong>the</strong> formation <strong>of</strong> Cr-rich 

precipitates within <strong>the</strong> NiAl matrix. The precipitates ranged from ~7 nm to 50 nm in size with 

65

<strong>the</strong> smaller ones exhibiting almost spherical shapes. These observations are similar to those <strong>of</strong> 

Cotton et al. [40] <strong>and</strong> Tian et al. [41]. Figure 9 shows a one-dimensional element concentration 

pr<strong>of</strong>ile analyzed through <strong>the</strong> large precipitate shown in Figure 8. From this pr<strong>of</strong>ile, <strong>the</strong> 

precipitates were found to consist almost entirely <strong>of</strong> Cr, which is consistent with <strong>the</strong> formation <strong>of</strong> 

α-Cr. The NiAl matrix surrounding <strong>the</strong> Cr-rich precipitates, was found to contain 36.79 ± 0.74 

at.% Al, 59.9 ± 0.92 at.% Ni, 0.44 ± 0.12 at.% Hf, 2.87 ± 0.62 at.% Cr, 0.007 ± 0.013 C. The Hf 

concentration is higher than <strong>the</strong> results reported by Larson <strong>and</strong> Miller [42] <strong>and</strong> Bestor et al. [24]. 

It is noted however that <strong>the</strong> materials examined in those studies were more stoichiometric <strong>and</strong> 

did not contain any significant concentrations <strong>of</strong> Cr or Co. No Heusler phase precipitates were 

detected in <strong>the</strong> samples areas. 

5.4. Conclusions 

In this study, NiAl-1Hf <strong>and</strong> NiAl-5Cr-1Hf coatings were produced using DC magnetron 

sputtering techniques which yielded a zone T microstructure. The as-deposited coatings formed 

a metastable B2 solid solution. After annealing, small precipitates formed at grain boundaries 

<strong>and</strong> within grain interiors. The majority <strong>of</strong> <strong>the</strong> precipitates were found to be Hf-rich with 

compositions consistent with a β′- Ni 2 AlHf phase. In <strong>the</strong> NiAlCrHf coating, additional Cr-rich 

precipitates were observed. SEM, TEM, <strong>and</strong> 3D-APT analysis confirmed <strong>the</strong>m to be α-Cr. 

EPMA line pr<strong>of</strong>iles indicated Hf segregation to <strong>the</strong> coating interfaces with <strong>the</strong> TGO <strong>and</strong> 

substrate. Elemental mapping using SEM-EDS supported <strong>the</strong> EPMA data showing that Hf 

segregates to <strong>and</strong> enriches <strong>the</strong> interfaces while Cr-rich precipitates form within <strong>the</strong> NiAlCrHf 

coatings. 

66


The authors acknowledge <strong>the</strong> support <strong>of</strong> <strong>the</strong> National Science Foundation (NSF) under Award 

No. DMR-0504950 <strong>and</strong> <strong>the</strong> National Aeronautics <strong>and</strong> Space Administration (NASA) under 

contract NN-X08AT21H. The use <strong>of</strong> <strong>the</strong> facilities supported by <strong>the</strong> Central Analytical facility at 

<strong>the</strong> University <strong>of</strong> Alabama <strong>and</strong> <strong>the</strong> Materials Diagnostics Laboratory at NASA’s Marshall Space 

Flight Center in Huntsville, Alabama is also gratefully acknowledged. 

67

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71

Table 1. Nominal chemical compositions <strong>of</strong> sputtering targets. 


NiAl-1.0Hf Nominal --- 51 48 --- --- 1 


1000°C/2h WDS 55.99 ± 0.41 38.67 ± 0.59 2.04 ± 0.25 2.57 ± 0.15 0.73 ± 0.16 

1000°C/4h WDS 53.44 ± 0.79 43.19 ± 0.72 1.33 ± 0.43 1.09 ± 0.30 0.85± 0.14 

NiAlCrHf Nominal --- 50 44 5 --- 1 


1000°C/2h WDS 51.73 ± 1.64 41.60 ± 0.81 4.65 ± 0.74 1.21 ± 0.23 0.81 ± 0.06 

1000°C/4h WDS 53.23 ± 1.42 40.50 ± 1..29 4.28 ± 0.92 1.16 ± 0.21 0.81 ± 0.10 

72

Figure 1. Representative SEM micrographs showing <strong>the</strong> as-deposited <strong>and</strong> annealed NiAl-1Hf (a, 

b) <strong>and</strong> NiAlCrHf (c, d) coatings. 

73

Figure 2. XRD patterns collected from <strong>the</strong> as-deposited <strong>and</strong> annealed NiAl-1Hf (a) <strong>and</strong> 

NiAlCrHf (b). 

74

Figure 3. Plan view TEM micrographs <strong>of</strong> <strong>the</strong> (a) NiAl-1Hf <strong>and</strong> (b) NiAlCrHf coatings 

following annealing at 1000°C for four hours. 

75

Figure 4. EPMA line pr<strong>of</strong>iles collected from <strong>the</strong> as-deposited <strong>and</strong> annealed NiAl-1Hf <strong>and</strong> 

NiAlCrHf coatings showing (a) Al, (b) Cr, (c) Ni, <strong>and</strong> (d) Hf. 

76

Figure 5. Cross-sectional SEM-EDS element maps for <strong>the</strong> NiAl-1Hf coating after annealing for 

four hours. 

77

Figure 6. Cross-sectional SEM-EDS element maps for <strong>the</strong> NiAlCrHf coating after annealing for 

four hours. 

78

Figure 7. STEM-HAADF image collected from a NiAlCrHf coating that was annealed for four 

hours at 1000°C with EDS spectra <strong>of</strong> β’-Ni 2 AlHf precipitate. 

79

Figure 8. 3D-APT data for NiAlCrHf after annealing for four hours at 1000°C: (a) full 

reconstruction displaying all elements in sample <strong>and</strong> (b) 2D-isosurface reconstruction showing 

features that are rich in Cr. 

80

Figure 9. Corresponding concentration pr<strong>of</strong>ile for <strong>the</strong> large Cr-rich precipitate in Figure 8b. 

81

CHAPTER VI 

MORPHOLOGICAL AND CHEMICAL EVOLUTION OF NIAL, NIAL-HF, <strong>and</strong> NIAL- 

CR-HF BOND COATS DURING SHORT TERM ISOTHERMAL OXIDATION 

M.A. Bestor*, J.P. Alfano*, <strong>and</strong> M.L. Weaver* 

*The University <strong>of</strong> Alabama, Department <strong>of</strong> Metallurgical <strong>and</strong> Materials Engineering, 

Box 870202, Tuscaloosa, Alabama 35487 

Key Words: β-NiAl, <strong>the</strong>rmal barrier coatings, oxidation, bond coats, sputtering 

Abstract: 

Thermal barrier coatings play a major role in protecting turbine blades from extreme 

operating environments <strong>and</strong> extending service lifetimes. Reactive elements (e.g. Zr, Hf, Y, Si, 

etc.) have been shown to enhance <strong>the</strong> oxidation performance <strong>of</strong> such coating systems when 

added in appropriate amounts to overlay bond coats. This study investigated <strong>the</strong> influences <strong>of</strong> Hf 

<strong>and</strong> Cr <strong>additions</strong> on <strong>the</strong> short-term oxidation performance <strong>of</strong> β-NiAl based coatings deposited via 

DC magnetron sputtering. The results indicate that small <strong>additions</strong> <strong>of</strong> Cr added in conjunction 

with supersaturations <strong>of</strong> Hf can have a dramatic <strong>effect</strong> in reducing <strong>the</strong> mass gains during 

iso<strong>the</strong>rmal exposure at 1050°C up to 100 hours. This indicates that this coating forms a thin, 

stable aluminum oxide scale that does not coarsen as rapidly as a similar coating without <strong>the</strong> Cr 

addition. Additional work was done to compare <strong>the</strong> oxidation performance with varying post- 

82

deposition heat treatments. Increasing <strong>the</strong> annealing time from two to four hours at 1000°C 

appears to lower <strong>the</strong> mass gains with <strong>the</strong> samples. This suggests that <strong>the</strong> coatings with longer 

annealing times form a thin, protective oxide faster than <strong>the</strong> samples with a shorter annealing 

time. Small <strong>additions</strong> <strong>of</strong> Cr to <strong>the</strong> NiAl-1Hf samples produced a thinner TGO <strong>and</strong> a decrease in 

<strong>the</strong> observed mass gains. 

Characterization <strong>of</strong> <strong>the</strong> coatings was done to study <strong>the</strong> <strong>effect</strong> <strong>of</strong> chemistry on <strong>the</strong> 

oxidation performance <strong>and</strong> resulting microstructures. The NiAl-Hf coatings oxidized at both <strong>the</strong> 

<strong>the</strong>rmally grown oxide <strong>and</strong> substrate interfaces <strong>of</strong> <strong>the</strong> coating. It was also determined by TEM 

analyses that <strong>the</strong> coatings oxidized internally along grain boundaries. The small grain sizes <strong>of</strong> 

<strong>the</strong> coatings confirm a large grain boundary volume is available for oxygen transport from <strong>the</strong> 

TGO to <strong>the</strong> coating/substrate interface. The zone T microstructure <strong>of</strong> <strong>the</strong> coatings contains a 

large amount <strong>of</strong> pinhole defects, <strong>and</strong> <strong>the</strong>se serve as direct pathways for oxygen transport to <strong>the</strong> 

coating/substrate interface. Fur<strong>the</strong>rmore, Hf migrates to <strong>the</strong> coating interfaces <strong>and</strong> enriches <strong>the</strong> 

oxide. Increasing preoxidation annealing time appears to increase <strong>the</strong> severity <strong>of</strong> <strong>the</strong> oxidation at 

<strong>the</strong> coating/substrate interface with <strong>the</strong> NiAl-Hf coatings. However, oxidation was not observed 

at <strong>the</strong> coating/substrate interface with <strong>the</strong> NiAlCrHf coatings. This may be due to <strong>the</strong> larger 

grain sizes <strong>and</strong> lower Hf concentration observed at <strong>the</strong> coating/substrate interface compared to 

<strong>the</strong> NiAl-Hf coatings. 

6.1 Introduction: 

Over <strong>the</strong> last three decades, significant advances in <strong>the</strong> high temperature capabilities <strong>of</strong> 

gas turbines have been realized through <strong>the</strong> development <strong>of</strong> new Ni-based superalloys, improved 

83

cooling technologies, <strong>and</strong> better manufacturing methods [1-7]. Modern gas turbines typically 

use combustor <strong>and</strong> turbine superalloys that have melting points ranging from ~1230°C to 

~1315°C [2,8]. However, when <strong>the</strong>se alloys are used to form components in <strong>the</strong> turbine, 

combustor, <strong>and</strong>/or augmentor sections <strong>of</strong> gas turbines, <strong>the</strong>y are vulnerable to environmental 

damage due to oxidation <strong>and</strong>/or hot corrosion (see refs. [1-3,9-11]). Because <strong>of</strong> this, <strong>the</strong> surfaces 

<strong>of</strong> superalloy components are <strong>of</strong>ten protected with environmental- <strong>and</strong> <strong>the</strong>rmal-protection 

coatings, such as diffusion aluminides, overlay coatings, <strong>and</strong> <strong>the</strong>rmal barrier coatings (TBC). 

A TBC is designed to reduce <strong>the</strong> bulk <strong>and</strong> surface temperatures <strong>of</strong> a component, thus 

increasing its service lifetime [5,12]. Typical TBC systems are composed <strong>of</strong> an insulating 

ceramic top coat (<strong>the</strong> actual TBC), a metallic bond coat, a <strong>the</strong>rmally growth oxide (TGO) 

between <strong>the</strong> top <strong>and</strong> bond coats, <strong>and</strong> <strong>the</strong> underlying superalloy substrate. The TGO 

simultaneously improves adhesion between <strong>the</strong> top coat <strong>and</strong> <strong>the</strong> bond coat while providing 

resistance to oxidation <strong>and</strong> hot corrosion. The performance <strong>of</strong> <strong>the</strong> bond coat, in particular <strong>the</strong> 

kinetics <strong>of</strong> TGO growth, have been shown to play one <strong>of</strong> <strong>the</strong> largest roles in determining TBC 

lifetime [2,10]. This is because TGO growth kinetics <strong>and</strong> <strong>the</strong> evolution <strong>of</strong> stresses within <strong>the</strong> 

coating system are dependent upon <strong>the</strong> chemistry <strong>of</strong> <strong>the</strong> bond coat. As such <strong>the</strong>re is significant 

interest in <strong>the</strong> development <strong>of</strong> improved bond coat systems specifically designed to improve 

TBC life. 

Modern TBCs typically utilize metallic MCrAlX overlay coatings (where “M” is Fe, Co, 

<strong>and</strong>/or Ni, <strong>and</strong> “X” is ei<strong>the</strong>r an oxygen-active element or precious metal) or diffusion aluminide 

coatings such as β-NiAl or β-(Ni,Pt)Al as bond coats [3,4,10]. It is well known that small 

84

amounts <strong>of</strong> reactive elements (RE) such as Hf or Zr yield significant improvements in TBC 

lifetime [13-17]. In <strong>the</strong>se studies, RE concentrations were maintained near <strong>the</strong>ir solubility limits 

in <strong>the</strong> aluminide coatings, which is reportedly around 0.05 at.% for Hf, to avoid excessive 

internal oxidation <strong>of</strong> RE-rich phases. 

General Electric Aviation has developed a series <strong>of</strong> novel multiphase coating 

compositions containing significantly higher RE concentrations [18-24]. Based upon <strong>the</strong> same 

β-phase aluminides mentioned above, <strong>the</strong>se new alloys, which can contain up to 2 at% Hf <strong>and</strong>/or 

Zr <strong>and</strong> as much as 5 at% Cr, exhibit higher strengths <strong>and</strong> oxidation resistance in furnace cycle 

tests comparable to state-<strong>of</strong>-<strong>the</strong>-art diffusion coatings. In a recent series <strong>of</strong> documents, we have 

detailed <strong>the</strong> influences <strong>of</strong> Cr <strong>and</strong> Hf <strong>additions</strong> in excess <strong>of</strong> <strong>the</strong>ir solubility limits on <strong>the</strong> 

microstructures <strong>of</strong> NiAl-based coatings produced via magnetron sputtering [25-28]. The present 

paper highlights <strong>the</strong> influences <strong>of</strong> Hf <strong>and</strong> Cr <strong>additions</strong> on <strong>the</strong> microstructures <strong>and</strong> short-term 

oxidation behavior <strong>of</strong> NiAl-based overlay coatings. 

6.2. Experimental Procedures: 

Alloy sputtering targets, 50.8 mm diameter × 6.4 mm thickness, were purchased from 

ACI Alloys, Inc. (San Jose, CA) <strong>and</strong> Sophisticated Alloys, Inc. (Butler, PA). The nominal target 

compositions are shown in Table 2.1. Coatings ranging from ~25 μm to ~40 μm thick were 

deposited via DC magnetron sputtering using an unbalanced high rate magnet configuration onto 

second generation Ni-based superalloy substrates (i.e., René N5). Argon was used as <strong>the</strong> 

working gas for deposition, <strong>and</strong> samples were preheated to a temperature <strong>of</strong> 650°C prior to 

85

deposition. The deposition power <strong>and</strong> working gas pressure were set to 300W <strong>and</strong> 1.33 Pa 

respectively. Coatings were deposited on <strong>the</strong> two major faces <strong>of</strong> each substrate. 

Selected as-deposited coatings were annealed at 1000°C for one to four hours in Ar with 

5% H 2 to improve adhesion between <strong>the</strong> coating <strong>and</strong> substrate, to remove sputtering induced 

defects, <strong>and</strong> to reduce any residual stresses from deposition. This resulted in <strong>the</strong> formation <strong>of</strong> a 

thin interdiffusion zone (IDZ) between <strong>the</strong> coatings <strong>and</strong> <strong>the</strong> underlying substrates. Iso<strong>the</strong>rmal 

oxidation tests were conducted on annealed specimens at 1050°C for times ranging from 24 to 96 

hours. As <strong>the</strong> focus <strong>of</strong> this research was to investigate microstructural evolution, <strong>the</strong> majority <strong>of</strong> 

<strong>the</strong> oxidation experiments were conducted on individual specimens, however, select specimens 

were removed from <strong>the</strong> furnace every 24 hours to measure mass change. 

Phase identification, microstructural evaluations, <strong>and</strong> chemistries were assessed via light 

optical microscopy, X-ray diffraction (XRD), scanning electron microscopy (SEM), <strong>and</strong> 

transmission electron microscopy (TEM). Chemistry was assessed using wavelength dispersive 

X-ray spectroscopy (WDS), energy dispersive X-ray spectroscopy (EDS), <strong>and</strong> atom probe 

tomography (APT). APT specimens were extracted from <strong>the</strong> coated substrates via <strong>the</strong> focusedion-beam 

(FIB) lift out technique [29,30]. The XRD patterns were generated using Cu Kα 

radiation in a Philips APD 3720 XRD system. Operating conditions for <strong>the</strong> diffractometer were 

set at 40 kV <strong>and</strong> 30 mA with a scan rate <strong>of</strong> ~1.4 deg/min with 0.02 deg steps. 

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6.3. Results <strong>and</strong> Discussion: 

6.3.1. Characterization <strong>of</strong> as-deposited <strong>and</strong> annealed coatings 

Details concerning coating deposition <strong>and</strong> <strong>the</strong> influences <strong>of</strong> annealing have been reported 

in references [27,28]. The experimentally measured coating compositions are included in Table 

3.1 for <strong>the</strong> as-deposited <strong>and</strong> annealed conditions. All <strong>of</strong> <strong>the</strong> as-deposited coatings were found to 

be single phase consisting <strong>of</strong> a B2 β-NiAl. Annealing at 1000°C resulted in <strong>the</strong> formation <strong>of</strong> an 

interdiffusion zone (IDZ) between <strong>the</strong> coating <strong>and</strong> substrate <strong>and</strong> in <strong>the</strong> formation <strong>of</strong> precipitates 

within <strong>the</strong> NiAl-(Hf, Cr) coatings. The IDZ generally consisted <strong>of</strong> a β-NiAl matrix with 

precipitates containing refractory metal rich precipitates. During annealing, composition 

changes also occurred driven by <strong>the</strong> composition gradient existing between <strong>the</strong> coatings <strong>and</strong> <strong>the</strong> 

underlying substrates. In general <strong>the</strong> coating Ni contents increased <strong>and</strong> Al contents decreased 

due to Al diffusion towards <strong>the</strong> exterior surface <strong>of</strong> <strong>the</strong> coating <strong>and</strong> towards <strong>the</strong> coating/substrate 

interface. In addition, Co, Cr <strong>and</strong> Ta were found at low levels within <strong>the</strong> coatings after annealing 

(generally around 2 at.% or less for Co <strong>and</strong> Cr <strong>and</strong> below 1 at.% for Ta). One exception was <strong>the</strong> 

coatings doped with Hf <strong>and</strong> Cr. These coatings still picked up Co from <strong>the</strong> substrates, however, 

<strong>the</strong> Cr <strong>and</strong> Hf contents (~4 at.% Cr <strong>and</strong> 0.8 at.% Hf respectively) remained constant. The 

precipitates that formed within <strong>the</strong> NiAl-Hf coatings were found to be Hf-rich whereas in <strong>the</strong> 

NiAlCrHf coatings, Hf-rich <strong>and</strong> Cr-rich precipitates were observed. 

6.4. Iso<strong>the</strong>rmal oxidation behavior <strong>and</strong> microstructures 

6.4.1. Oxidation test results 

Annealed samples were subjected to iso<strong>the</strong>rmal oxidation at 1050°C for as long as 96 

hours in laboratory air. The oxidation experiments were conducted for <strong>the</strong> sole purpose <strong>of</strong> 

87

investigating microstructural evolution during oxidation, not to ga<strong>the</strong>r kinetic data concerning 

oxidation. The specific specimen mass changes for <strong>the</strong> alloys investigated in this study are 

presented in Figure 1. As expected, all <strong>of</strong> <strong>the</strong> specimens exhibited rapid mass gains during <strong>the</strong> 

first 24 hours <strong>of</strong> oxidation indicating <strong>the</strong> growth <strong>of</strong> an oxide scale or TGO. For most <strong>of</strong> <strong>the</strong> 

specimens, <strong>the</strong> mass gain increased with iso<strong>the</strong>rmal oxidation time implying <strong>the</strong> continued 

growth <strong>of</strong> a protective TGO. Figure 1a shows specimen mass changes for <strong>the</strong> coatings 

investigated in this study after annealing for two hours at 1000°C. All <strong>of</strong> <strong>the</strong> specimens 

exhibited mass gains through 96 h <strong>of</strong> oxidation; however, after 24 h <strong>of</strong> oxidation, mass losses 

were observed in binary NiAl indicating oxide scale spallation, which is consistent with 

published literature [31,32]. Higher mass gains were observed with increasing Hf concentration. 

The decreased rate <strong>of</strong> mass gains with increasing time implies that <strong>the</strong> samples are forming a 

protective oxide layer <strong>and</strong> its growth rate is decreasing with time. Adding 5 at.% Cr to <strong>the</strong> NiAl- 

1Hf coating yields a similar oxidation behavior, but with lower mass gains than <strong>the</strong> NiAl-1Hf 

coatings. The NiAl-0.5Hf sample has <strong>the</strong> lowest mass gain <strong>of</strong> <strong>the</strong> Hf-doped coatings annealed 

for two hours. Similar observations have been reported by Pint et al. [31]. 

Figure 1b shows specimen mass changes after annealing for four hours at 1000°C. The 

results were similar to those collected on <strong>the</strong> two hour annealed specimens; however, <strong>the</strong> mass 

gains where larger than those observed in samples annealed for two hours. Interestingly, <strong>the</strong> 

NiAlCrHf coating exhibited virtually no mass change for exposures <strong>of</strong> greater than 24 hours, 

which indicates that <strong>the</strong> NiAlCrHf coating forms a thin, protective oxide that does not thicken up 

to 96 hours <strong>of</strong> iso<strong>the</strong>rmal oxidation. 

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To investigate unexpected variations in <strong>the</strong> mass change data, single coating specimens 

were subjected to a series <strong>of</strong> 24 hour oxidation cycles. More specifically NiAl-1Hf annealed for 

two hours <strong>and</strong> NiAlCrHf annealed for four hours. The resulting mass gains are shown as open 

symbols connected by dotted lines in Figures 1a <strong>and</strong> 1b. The results indicate continuous TGO 

growth on <strong>the</strong> NiAl-1Hf coating with increasing oxidation time but lower scale growth rates on 

<strong>the</strong> NiAlCrHf coating. 

The gradual increase in specific mass change for <strong>the</strong> Hf-containing coatings has been 

attributed to excellent scale adhesion <strong>and</strong> has been observed without <strong>the</strong> addition <strong>of</strong> Pt to <strong>the</strong> 

coating [13,14,24,33-38]. The reduced rate <strong>of</strong> mass gain for optimized NiAl-Hf alloys <strong>and</strong> 

coatings has been attributed to <strong>the</strong> suppression <strong>of</strong> Al grain boundary transport by <strong>the</strong> segregation 

<strong>of</strong> reactive element ions to α-Al 2 O 3 grain boundaries [13,15]. The addition <strong>of</strong> Cr to β-NiAl 

coatings is typically done to achieve enhanced hot corrosion protection. It has been reported by 

Leyens et al. [39] that as little as 2 at.% Cr is necessary to reduce <strong>the</strong> attack from hot corrosion 

<strong>and</strong> coating performance increases with <strong>chromium</strong> content. However, Cr has also been shown to 

degrade <strong>the</strong> oxidation resistance <strong>of</strong> NiAl, even in alloys containing “optimal” amounts <strong>of</strong> Hf 

[31]. 

6.4.2. Microstructural changes during oxidation 

After oxidation for 96 hours, samples were examined by XRD to identify <strong>the</strong> phases that 

were present. The results, which are shown in Figure 2, revealed combinations <strong>of</strong> β-NiAl, γ′- 

Ni 3 Al, θ-Al 2 O 3 , <strong>and</strong> α-Al 2 O 3 phases in all <strong>of</strong> <strong>the</strong> specimens, with <strong>the</strong> final content being 

dependent upon coating chemistry. In <strong>the</strong> NiAlCrHf coatings for example, no γ′-Ni 3 Al peaks 

were observed which supports <strong>the</strong> hypo<strong>the</strong>sis from <strong>the</strong> oxidation data (Figure 1) that a thin, 

89

protective TGO layer forms <strong>and</strong> grows slowly preserving <strong>the</strong> aluminum reservoir within <strong>the</strong> 

bond coat. In contrast, <strong>the</strong> presence <strong>of</strong> γ′-Ni 3 Al peaks in <strong>the</strong> NiAl-0.5Hf <strong>and</strong> NiAl-1Hf coatings 

suggests continuous TGO growth, depleting <strong>the</strong> coating <strong>of</strong> Al. Similar observations were made 

by Yang et al. who investigated <strong>the</strong> oxidation resistance <strong>of</strong> sputter deposited binary NiAl 

coatings [40,41]. 

In general, it was difficult to distinguish between many <strong>of</strong> <strong>the</strong> α-Al 2 O 3 <strong>and</strong> θ-Al 2 O 3 

peaks; however, for <strong>the</strong> ternary NiAl-Hf coatings, it was determined that <strong>the</strong> coatings preannealed 

for two hours had higher α-Al 2 O 3 <strong>and</strong> θ-Al 2 O 3 contents than coatings annealed for four 

hours. Since <strong>the</strong> time to transition from θ-Al 2 O 3 to α-Al 2 O 3 is expected to be only a few hours 

[42,43]; <strong>the</strong>se observations suggest that <strong>the</strong> oxide spalls <strong>and</strong> <strong>the</strong>n re-grows. It is well established 

that tremendous compressive stresses, on <strong>the</strong> order <strong>of</strong> several GPa, arise within <strong>the</strong> TGO during 

its formation <strong>and</strong> growth [44-47]. These stresses can lead to crack formation, spalling, <strong>and</strong> 

additional oxidation <strong>of</strong> <strong>the</strong> underlying metal, in this case <strong>the</strong> bond coat [44-46]. This increase in 

area <strong>of</strong> oxidation would account for <strong>the</strong> additional θ-Al 2 O 3 X-ray intensity observed in <strong>the</strong> 

ternary NiAl-Hf coatings. Fur<strong>the</strong>rmore, <strong>the</strong> low intensity <strong>of</strong> <strong>the</strong> Al 2 O 3 peaks combined with <strong>the</strong> 

lack <strong>of</strong> any γ′-Ni 3 Al peaks in <strong>the</strong> NiAlCrHf samples supports <strong>the</strong> hypo<strong>the</strong>sis that <strong>the</strong> oxides for 

<strong>the</strong>se samples are much thinner than <strong>the</strong> ones observed on <strong>the</strong> NiAl-Hf coatings. 

Figures 3 <strong>and</strong> 4 show <strong>the</strong> surface morphologies <strong>of</strong> scales formed on NiAl-0.5Hf <strong>and</strong> 

NiAl-1Hf coatings respectively. The surfaces were completely covered with a mixture <strong>of</strong> bladelike 

θ-Al 2 O 3 <strong>and</strong> more equiaxed α-Al 2 O 3 . Cracks were observed on some regions <strong>of</strong> <strong>the</strong> oxide 

surfaces along with wrinkling <strong>of</strong> <strong>the</strong> scale. The evolution <strong>of</strong> such features has been summarized 

90

y Jedlinsky [48]. Over time, θ-Al 2 O 3 which forms first will transform to α-Al 2 O 3 , which is 

<strong>the</strong>rmodynamically stable at 1050°C. The volume contraction resulting from <strong>the</strong> θ-α phase 

transformation can lead to localized cracking <strong>and</strong> scale wrinkling [48-51]. Additional 

contributors to this process are <strong>the</strong> TGO growth stresses described in <strong>the</strong> previous paragraph [44- 

46]. These stresses, which are compressive in nature, curl <strong>the</strong> oxide causing <strong>the</strong> metastable θ- 

Al 2 O 3 to spall from <strong>the</strong> underlying α-Al 2 O 3 . Similar observations were reported by Ning [37] in 

his <strong>investigation</strong>s <strong>of</strong> sputter deposited NiAl-Hf coatings. It is likely that <strong>the</strong>se cracks propagate 

along grain boundaries. Chemical analyses using SEM-EDS showed that <strong>the</strong>re was no 

preferential segregation around <strong>the</strong> cracks <strong>and</strong> no formation <strong>of</strong> transient oxides (e.g. HfO 2 , NiO, 

spinel, etc.). 

The surfaces <strong>of</strong> <strong>the</strong> NiAlCrHf coatings were similar to <strong>the</strong> NiAl-Hf coatings in that <strong>the</strong> 

surfaces were covered with a mixture <strong>of</strong> blade-like θ-Al 2 O 3 <strong>and</strong> more equiaxed α-Al 2 O 3 , <strong>and</strong> 

exhibited some cracking. Differently, however, <strong>the</strong> cracks observed with <strong>the</strong>se samples tended 

to be much longer <strong>and</strong> less tortuous than those observed in <strong>the</strong> NiAl-Hf samples. Fur<strong>the</strong>rmore, 

<strong>the</strong>se cracks appeared to heal, after formation. Elemental maps did not reveal any evidence <strong>of</strong> 

<strong>hafnium</strong> or <strong>chromium</strong> segregation to <strong>the</strong> TGO or cracks, which suggests that <strong>and</strong> it was 

completely comprised <strong>of</strong> Al 2 O 3 . 

Cross-sectional SEM images that were collected from <strong>the</strong> samples that were oxidized for 

96 hours are shown in Figures 6-8. Most <strong>of</strong> <strong>the</strong> specimens were sputter coated with Cu 

following oxidation to protect <strong>the</strong> oxide scales from damage during metallographic sample 

preparation. In <strong>the</strong> NiAl-0.5Hf <strong>and</strong> NiAl-1Hf coatings (Figures 6 <strong>and</strong> 7), γ′-Ni 3 Al was observed 

91

within <strong>the</strong> coatings along with a number <strong>of</strong> low <strong>and</strong> high atomic number (Z) contrast spots 

superimposed on <strong>the</strong> coating, <strong>the</strong> volume fraction <strong>and</strong> sizes <strong>of</strong> which increased with Hf 

concentration <strong>and</strong> annealing time. The low Z spots appeared to have <strong>the</strong> same atomic number 

contrast as <strong>the</strong> TGO, but were too small for quantitative chemical analysis via EMPA. Elemental 

maps collected in <strong>the</strong> vicinity <strong>of</strong> <strong>the</strong> low Z spots (presented later in this document) did reveal <strong>the</strong> 

preferential segregation <strong>of</strong> Al, suggesting <strong>the</strong> formation <strong>of</strong> Al 2 O 3 . Similar features have been 

reported by Guo et al. in <strong>the</strong>ir <strong>investigation</strong> <strong>of</strong> Al 2 O 3 -dispersed NiAl coatings deposited via 

electron beam physical vapor deposition (EB-PVD) [52]. 

In <strong>the</strong> NiAl-1Hf coating, a large number <strong>of</strong> high Z contrast precipitates were also 

observed within <strong>the</strong> bulk <strong>of</strong> <strong>the</strong> coatings, where <strong>the</strong>y appeared to form along columnar grain 

boundaries, near <strong>the</strong> coating/substrate <strong>and</strong> coating/TGO interfaces <strong>and</strong> in some cases within <strong>the</strong> 

TGO. Such interfacial segregation <strong>and</strong> associated precipitation has been extensively reported for 

RE-containing coatings [31,33,34,53-57]. Based upon <strong>the</strong>se observations, it is most likely that 

<strong>the</strong> dark spots are Al 2 O 3 while <strong>the</strong> bright spots are ei<strong>the</strong>r equilibrium Heusler phase or HfO 2 , 

which forms as a result <strong>of</strong> internal oxidation. Some surprising observations included <strong>the</strong> 

formation <strong>of</strong> voids <strong>and</strong> <strong>the</strong> occurrence oxidation in some <strong>of</strong> <strong>the</strong> coating/substrate interfaces. 

Doubling <strong>the</strong> amount <strong>of</strong> <strong>hafnium</strong> in <strong>the</strong> coating produced larger precipitates <strong>and</strong> also a 

mixed oxide consisting <strong>of</strong> Al 2 O 3 <strong>and</strong> HfO 2 with a layer <strong>of</strong> Al 2 O 3 on top <strong>of</strong> <strong>the</strong> mixed oxide. 

Similar observations were reported by Guo et al. <strong>and</strong> Sun et al. [33,34]. Oxidation at <strong>the</strong> 

interface between <strong>the</strong> coating <strong>and</strong> substrate was observed with <strong>the</strong> NiAl-1Hf as well. Like <strong>the</strong> 

NiAl-0.5Hf, <strong>the</strong> oxidation at <strong>the</strong> coating substrate interface appeared to become more severe with 

increased annealing time. The amount <strong>of</strong> <strong>the</strong> γ′ phase was less significant in <strong>the</strong> NiAl-1Hf 

92

sample which is in agreement with <strong>the</strong> XRD results shown Figure 2. This indicates that <strong>the</strong> 

aluminum is being consumed at a faster rate in <strong>the</strong> NiAl-0.5Hf coating. Surprisingly, <strong>the</strong> TGO 

thickness remained relatively unchanged as <strong>the</strong> Hf concentration was increased from 0.5-1.0 

at.% Hf. 

The NiAlCrHf coatings, presented in Figure 8a <strong>and</strong> 8b, have completely different 

microstructures than <strong>the</strong> NiAl-Hf coatings. The TGO is about half as thick as compared to <strong>the</strong> 

NiAl-0.5Hf <strong>and</strong> NiAl-1Hf samples. For example, in <strong>the</strong> specimen that was annealed for four 

hours prior to oxidation, a 2.11 μm thick TGO formed on <strong>the</strong> NiAlCrHf coating; whereas 3.96 

μm <strong>and</strong> 4.18 μm thick TGOs formed on <strong>the</strong> NiAl-0.5Hf <strong>and</strong> NiAl-1Hf coatings respectively. 

The interdiffusion zone (IDZ) was also fully retained in <strong>the</strong>se coatings, but no oxidation was 

observed at <strong>the</strong> coating/substrate interface. Although <strong>the</strong>se coatings contained <strong>the</strong> same amount 

<strong>of</strong> <strong>hafnium</strong> as <strong>the</strong> NiAl-1Hf coatings, <strong>the</strong> preferential segregation Hf was not observed to be as 

significant with ei<strong>the</strong>r heat treatment. Fur<strong>the</strong>rmore, no mixed oxide was present. Low Z spots, 

similar to <strong>the</strong> ones observed in <strong>the</strong> NiAl-0.5Hf <strong>and</strong> NiAl-1Hf coatings, were also observed in <strong>the</strong> 

NiAlCrHf coatings indicating <strong>the</strong> occurrence <strong>of</strong> some internal oxidation; however <strong>the</strong> volume 

fraction <strong>of</strong> <strong>the</strong>se spots was significantly lower indicating a higher intrinsic resistance to 

oxidation. The high Z precipitates observed within <strong>the</strong> NiAlCrHf coatings appeared to coarsen 

with increased annealing time. Similar observations were made in <strong>the</strong> NiAl-0.5Hf <strong>and</strong> NiAl-1Hf 

samples. 

EPMA line pr<strong>of</strong>iles collected from samples that were oxidized for 96 hours are shown in 

Figure 9-11. The average chemistries <strong>of</strong> <strong>the</strong> NiAl-Hf <strong>and</strong> NiAl-Cr-Hf coatings after oxidation 

93

are summarized in Table 3.2. Only Ni, Al, Cr, Co, <strong>and</strong> Hf are shown to indicate <strong>the</strong> trends in 

elemental diffusion. 

Figures 9a <strong>and</strong> 9b were collected from <strong>the</strong> NiAl-0.5Hf coatings that were annealed for 

two <strong>and</strong> four hours respectively prior to oxidation. The Ni <strong>and</strong> Al concentrations were observed 

to be approximately 58 at.% <strong>and</strong> 37 at.% respectively <strong>and</strong> did not vary appreciably within <strong>the</strong> 

coating for ei<strong>the</strong>r sample. As noted previously, <strong>the</strong> TGO evident at <strong>the</strong> surface <strong>of</strong> <strong>the</strong> coating 

was found to be composed primarily <strong>of</strong> Al 2 O 3 , most likely α-Al 2 O 3 , as was shown in <strong>the</strong> XRD 

results (Figure 2). Interdiffusion between <strong>the</strong> coatings <strong>and</strong> <strong>the</strong> underlying substrate is strongly 

indicated by <strong>the</strong> concentration pr<strong>of</strong>iles, which show gradients in <strong>the</strong> Al, Co <strong>and</strong> Cr going from 

<strong>the</strong> coating into <strong>the</strong> substrate. The TGO thickness along with <strong>the</strong> IDZ thickness was reported to 

increase with increased annealing time. As discussed earlier, this result confirms that <strong>the</strong> 

increased annealing time tends to increase <strong>the</strong> amount <strong>of</strong> interdiffusion within <strong>the</strong> coating prior 

to service <strong>and</strong> this trend continues with <strong>the</strong> short-term oxidation results as evidenced by <strong>the</strong> 

higher mass gains observed with <strong>the</strong> samples with longer pre-oxidation annealing times. As 

noted previously, both coatings did show <strong>the</strong> formation <strong>of</strong> voids at <strong>the</strong> coating/substrate interface 

along with <strong>the</strong> presence <strong>of</strong> large amounts <strong>of</strong> Al 2 O 3 . This phenomenon was attributed to oxygen 

transport along grain boundaries <strong>and</strong> when present cracks in <strong>the</strong> coating that penetrated to <strong>the</strong> 

coating/substrate interface. These features appear in Figures 9a <strong>and</strong> 9b as significant drops in Ni 

concentration accompanied by a slight enrichment in Al. The interdiffusion zone (IDZ) follows 

this for 20-30 μm before <strong>the</strong> substrate concentrations are observed. The Hf concentration within 

<strong>the</strong> coating was determined to be ~0.2-0.7 at.% after annealing. 

94

Figure 10 shows results from <strong>the</strong> NiAl-1Hf coating after oxidation. The Ni <strong>and</strong> Al 

concentrations remained stable throughout <strong>the</strong> coating, similar to <strong>the</strong> results for <strong>the</strong> NiAl-0.5Hf 

coatings; however, <strong>the</strong> Al concentration decreased drastically in <strong>the</strong> IDZ <strong>and</strong> substrate. This was 

expected because <strong>the</strong> coating had a higher Al concentration than <strong>the</strong> substrate. The <strong>hafnium</strong> 

concentration within <strong>the</strong> coating was consistently observed to be around 0.5 at% with larger 

amounts found at <strong>the</strong> interfaces due to <strong>the</strong> formation <strong>of</strong> HfO 2 . Enrichment <strong>of</strong> Hf at interfaces 

was reported to be as high as 10 at.% near <strong>the</strong> oxide voids indicating that a mixed oxide could be 

present. Chromium <strong>and</strong> cobalt are commonly observed to diffuse upward into <strong>the</strong> coating during 

annealing <strong>and</strong> service, <strong>and</strong> <strong>the</strong>y also exhibit a concentration gradient up to <strong>the</strong>ir respective values 

in <strong>the</strong> substrate. They are typically found at concentrations near 2-3 at% within <strong>the</strong> coating. The 

sharp peaks with <strong>the</strong> <strong>chromium</strong> result from <strong>the</strong> presence <strong>of</strong> refractory metal rich topographically 

close packed (TCP) phases within <strong>the</strong> IDZ. There is a sharp peak in <strong>the</strong> aluminum concentration 

pr<strong>of</strong>ile near <strong>the</strong> IDZ that was taken through an Al 2 O 3 pocket in <strong>the</strong> coating. The only elements 

that were contained within <strong>the</strong>se areas were aluminum <strong>and</strong> <strong>hafnium</strong>; supporting <strong>the</strong> observations 

<strong>of</strong> a mixed oxide with <strong>the</strong> NiAl-1Hf images presented in Figure 7. 

Analyses were conducted on <strong>the</strong> NiAlCrHf samples oxidized for 96 hours at 1050°C. 

Although <strong>the</strong> aluminum concentration in <strong>the</strong> as-deposited state is 5 at.% lower in <strong>the</strong> NiAlCrHf 

samples than <strong>the</strong> NiAl-Hf samples, <strong>the</strong> pr<strong>of</strong>iles in <strong>the</strong> coating are very similar. There is some 

<strong>hafnium</strong> enrichment at <strong>the</strong> interfaces, similar to <strong>the</strong> NiAl-0.5Hf coatings. The evolution <strong>of</strong> large 

Cr-rich phases is evident within <strong>the</strong> IDZ. Sharp peaks in <strong>the</strong> Ni <strong>and</strong> corresponding Cr pr<strong>of</strong>iles 

suggest that Ni substitutes for Cr in <strong>the</strong> IDZ. Chromium <strong>and</strong> cobalt concentrations do not vary 

significantly in <strong>the</strong> coating with <strong>chromium</strong> concentrations near 4 at.% <strong>and</strong> cobalt near 3 at.%. 

95

The IDZ thickness, as determined from <strong>the</strong> EPMA data, is approximately 28-30 μm for both 

samples. This is in contrast to <strong>the</strong> SEM cross-sectional images <strong>of</strong> <strong>the</strong>se coatings which 

suggested that annealing time had an <strong>effect</strong> on <strong>the</strong> amount <strong>of</strong> interdiffusion between <strong>the</strong> coating 

<strong>and</strong> substrate. Some <strong>of</strong> <strong>the</strong>se have been identified as γ′-Ni 3 Al with dissolved <strong>chromium</strong>. 

6.5. Elemental Mapping <strong>of</strong> Coatings Iso<strong>the</strong>rmally Oxidized for 96 hours. 

Figures 12-14 show EDS spectral maps from <strong>the</strong> NiAl-0.5Hf, NiAl-1Hf, <strong>and</strong> NiAlCrHf 

samples respectively. High resolution maps for aluminum, oxygen, <strong>and</strong> <strong>hafnium</strong> are shown for 

<strong>the</strong> NiAl-Hf samples <strong>and</strong> aluminum, oxygen, <strong>hafnium</strong>, <strong>and</strong> <strong>chromium</strong> are displayed for <strong>the</strong> 

NiAlCrHf samples. Corresponding SEM images are shown with <strong>the</strong> NiAl-Hf samples. 

Both samples exhibit <strong>hafnium</strong> segregation to <strong>the</strong> interfaces, but <strong>the</strong>re is no evidence <strong>of</strong> 

hafnia formation at <strong>the</strong> TGO interface. However, with <strong>the</strong> Al 2 O 3 that forms at <strong>the</strong> 

coating/substrate interface does have <strong>hafnium</strong> surrounding <strong>the</strong> oxides. It is likely that <strong>the</strong>re is a 

mixed oxide at this interface. The dark regions within <strong>the</strong> Al maps indicate <strong>the</strong> formation <strong>of</strong> γ′- 

Ni 3 Al. This is also observed with <strong>the</strong> corresponding light phases in <strong>the</strong> bond coat with both 

samples. Al 2 O 3 formation at <strong>the</strong> coating/substrate interface is more significant with <strong>the</strong> sample 

that was annealed for four hours while <strong>the</strong> TGO thickness is approximately <strong>the</strong> same for both 

samples. The Hf was found to segregate to <strong>the</strong> γ′ phase in small amounts. This is due to <strong>the</strong> γ′ 

phase having a higher Hf solubility than <strong>the</strong> β phase [58]. 

Similar to <strong>the</strong> NiAl-0.5Hf samples, <strong>the</strong> spectral maps show that <strong>the</strong> oxide is primarily 

composed <strong>of</strong> Al 2 O 3 with additional segregation <strong>of</strong> <strong>hafnium</strong> to <strong>the</strong> coating/oxide interface. 

96

Oxides are found at both interfaces with leaders growing into <strong>the</strong> coating. The leaders are 

clearly visible as vertical <strong>hafnium</strong>-rich features that extend through <strong>the</strong> coating. Doubling <strong>the</strong> 

<strong>hafnium</strong> concentration leads to increased degree <strong>of</strong> decoration at <strong>the</strong> interfaces <strong>and</strong> presumably a 

mixed oxide. This is more significant with <strong>the</strong> samples that were annealed for four hours. This 

is evidenced by <strong>the</strong> brighter spots on <strong>the</strong> image indicating a larger concentration <strong>of</strong> <strong>hafnium</strong> 

atoms. The <strong>hafnium</strong> maps for both samples show segregation to interfaces while also enriching 

grain boundaries within <strong>the</strong> coating. This is expected because Hf typically diffuses to free 

surfaces via high transport paths such as grain boundaries. 

The TGO on <strong>the</strong> NiAlCrHf coatings was roughly half <strong>the</strong> thickness <strong>of</strong> <strong>the</strong> one on <strong>the</strong> 

NiAl-1Hf coatings (Figure 13). The <strong>hafnium</strong> map for <strong>the</strong> NiAlCrHf sample after annealed for 

two hours (Figure 13a) shows little segregation to interfaces, but precipitates observed within <strong>the</strong> 

coating appear to form vertically along grain boundaries. These precipitates coarsen with 

increased annealing time. The Cr-rich phases in <strong>the</strong> IDZ were found to be ei<strong>the</strong>r γ′-Ni 3 Al with 

dissolved Cr or α-Cr within a β-NiAl matrix confirming <strong>the</strong> results in Figure 11. As annealing 

time was increased to four hours (b), <strong>the</strong>se phases are confined closer to <strong>the</strong> coating interface <strong>and</strong> 

are also smaller in size. The IDZ was found to be almost completely devoid <strong>of</strong> <strong>hafnium</strong> in both 

samples, in contrast to <strong>the</strong> NiAl-Hf samples that were presented in Figures 12 <strong>and</strong> 13. There is 

also no evidence <strong>of</strong> void formation at <strong>the</strong> coating/substrate interface or internal oxidation. The 

bright spots in <strong>the</strong> Al map indicate sites where internal oxidation occurred. 

97

6.6. TEM Analysis <strong>of</strong> Oxidized Coatings. 

Investigation <strong>of</strong> <strong>the</strong> microstructures was also conducted using <strong>the</strong> TEM. Samples were 

prepared using a focused ion beam <strong>and</strong> were attached to a copper grid for analysis. The samples 

were collected following oxidation at 96 hours. Bright field <strong>and</strong> corresponding STEM-HAADF 

images were collected to determine <strong>the</strong> <strong>effect</strong> <strong>of</strong> oxidation on grain size <strong>and</strong> precipitation. 

Figure 15 shows images collected from <strong>the</strong> NiAl-0.5Hf sample. Bright regions in <strong>the</strong> 

bright field images correspond to <strong>the</strong> dark phases in <strong>the</strong> STEM-HAADF images. These phases 

have been identified as Al 2 O 3 . The Al 2 O 3 primarily forms at grain boundaries <strong>and</strong> not within <strong>the</strong> 

grain interiors. This supports <strong>the</strong> observation that <strong>the</strong> oxidation observed at <strong>the</strong> coating/substrate 

interface occurs by grain boundary transport <strong>of</strong> oxygen through <strong>the</strong> coating. However, 

coarsening <strong>of</strong> <strong>the</strong>se phases to grain sizes <strong>of</strong> up to 350 nm has been observed. The bright phases 

in Figure 15 have been identified as Heusler-Ni 2 AlHf <strong>and</strong> were found to be ei<strong>the</strong>r block-like or 

spherical in nature. These precipitates form at grain boundaries <strong>and</strong> within grain interiors with 

<strong>the</strong> larger precipitates forming at grain boundaries. 

Figure 16 shows <strong>the</strong> results from <strong>the</strong> NiAl-1Hf samples that were oxidized for 96 hours. 

The precipitates appear to coarsen with <strong>the</strong> increase in Hf concentration. The amount <strong>of</strong> Al 2 O 3 

within <strong>the</strong> coating decreases dramatically from <strong>the</strong> observations with <strong>the</strong> NiAl-0.5Hf samples; 

however, <strong>the</strong> Al 2 O 3 precipitates do appear to be larger than those reported for <strong>the</strong> NiAl-0.5Hf. 

The majority <strong>of</strong> <strong>the</strong> Al 2 O 3 phases have a rod-like structure with some also appearing spherical. 

The bright phases in <strong>the</strong> STEM-HAADF images were determined to be Heusler-Ni 2 AlHf. This 

is similar to <strong>the</strong> results reported by Hazel et al.[24]. The precipitates that form within grains are 

98

smaller than those found at grain boundaries <strong>and</strong> this is attributed to <strong>the</strong> high diffusivity that 

exists at grain boundaries. Precipitates within <strong>the</strong> two hour annealed samples were also smaller 

<strong>and</strong> more uniformly dispersed than those in <strong>the</strong> four hour samples. It was additionally observed 

that <strong>the</strong> four hour annealed specimens exhibited larger grain sizes than <strong>the</strong> two hour annealed 

specimens. 

Figure 17 presents <strong>the</strong> images collected from <strong>the</strong> NiAlCrHf coatings that were oxidized 

for 96 hours. The precipitates were much larger <strong>and</strong> less dispersed than those observed in <strong>the</strong> 

NiAl-Hf coatings. Increasing <strong>the</strong> annealing time did lead to precipitate coarsening with <strong>the</strong> 

largest precipitates forming at grain boundaries. Interestingly, <strong>the</strong>re was almost no precipitation 

within <strong>the</strong> grain interiors. Hafnium enrichment at <strong>the</strong> grain boundaries was also evident. 

However, <strong>the</strong> amount <strong>of</strong> grain growth observed going from <strong>the</strong> two hour to <strong>the</strong> four hour 

annealed coatings was not as noticeable as compared to <strong>the</strong> NiAl-Hf samples. This supports <strong>the</strong> 

<strong>the</strong>ory that <strong>the</strong> NiAlCrHf samples exhibit slower diffusion kinetics than <strong>the</strong> NiAl-Hf samples. 

6.7. General Discussion 

The results presented in <strong>the</strong> previous sections showed that composition <strong>and</strong> 

microstructural morphology can have significant influences on <strong>the</strong> microstructure <strong>and</strong> properties 

<strong>of</strong> <strong>the</strong> sputtered NiAl coatings. The as-deposited coatings were single phase zone T 

microstructures with no evidence <strong>of</strong> precipitation [27,28]. The microstructures <strong>of</strong> <strong>the</strong> assputtered 

coatings contained a high volume <strong>of</strong> defects (excess vacancies, dislocations, etc.), 

which represent preferential sites for solute segregation <strong>and</strong> <strong>the</strong> nucleation <strong>of</strong> precipitates 

[59,60]. The issue <strong>of</strong> elemental segregation in NiAl has been addressed by Jayaram <strong>and</strong> Miller 

99

[61] <strong>and</strong> Larson <strong>and</strong> Miller [62] who showed evidence <strong>of</strong> <strong>the</strong> preferential segregation <strong>of</strong> Zr to 

dislocations <strong>and</strong> grain boundaries in NiAl-Zr,Mo <strong>and</strong> NiAl-Hf alloys. 

In <strong>the</strong> present coatings, Hf <strong>and</strong> Cr were observed to have a significant influence on <strong>the</strong> 

coating microstructures <strong>and</strong> on interdiffusion between <strong>the</strong> coating <strong>and</strong> substrate. Hf <strong>additions</strong> <strong>of</strong> 

0.5 at.% <strong>and</strong> 1.0 at.% into NiAl was found to result in significantly decreased coating grain sizes 

<strong>and</strong> in thinner interdiffusion zones in coatings <strong>of</strong> equivalent thickness. It was also observed that 

<strong>the</strong> concentrations <strong>of</strong> substrate elements, such as Co <strong>and</strong> Cr, incorporated into <strong>the</strong> coatings after 

post-deposition annealing decreased as <strong>the</strong> Hf content was increased. These observations are 

consistent with prior reports <strong>of</strong> reduced coating <strong>and</strong>/or IDZ thickness have been reported for 

diffusion aluminide <strong>and</strong> β-NiAl overlay coatings [36,63-66]. The addition <strong>of</strong> 5 at.% Cr in 

addition to 1 at.% Hf, produced no obvious changes in IDZ thickness from <strong>the</strong> NiAl-1Hf coating 

after annealing. However, it did exhibit a high number <strong>of</strong> α-Cr precipitates in addition to <strong>the</strong> 

Heusler phase. 

Analogous to <strong>the</strong> observations <strong>of</strong> Ning et al. [25], <strong>the</strong> Hf <strong>and</strong>, when present, Cr <strong>additions</strong> 

also tended to precipitate out during <strong>the</strong> post-deposition annealing along grain boundaries <strong>and</strong> on 

dislocations. It was observed that under same deposition <strong>and</strong> similar post-deposition annealing 

conditions, that <strong>the</strong> NiAlCrHf coatings exhibited lower mass gains than <strong>the</strong> NiAl-Hf coatings. In 

<strong>the</strong> case <strong>of</strong> <strong>the</strong> ternary NiAl-Hf coatings, this is attributed to “over doping” as has been reported 

by Pint et al. [13,15,67]. The apparent increase in oxidation resistance provided by <strong>the</strong> addition 

<strong>of</strong> Cr was not expected, but is tentatively attributed, in part, to larger grain sizes <strong>and</strong> 

100

The oxidation observed at <strong>the</strong> interface between <strong>the</strong> coating <strong>and</strong> <strong>the</strong> substrate can be 

explained by a two-step mechanism. As shown with <strong>the</strong> EPMA data earlier in <strong>the</strong> document, 

aluminum is depleted from <strong>the</strong> coating into <strong>the</strong> underlying superalloy. As <strong>the</strong> aluminum is 

depleted, it leads to a transition from β-NiAl to γ’-Ni 3 Al <strong>and</strong> <strong>the</strong> ability <strong>of</strong> <strong>the</strong> bond coating to 

form Al 2 O 3 is greatly reduced with increased exposure. As oxygen is continually transported to 

<strong>the</strong> coating/substrate interface via <strong>the</strong> pinhole defects <strong>and</strong> grain boundary transport, transient 

oxides (e.g. NiO x , spinel, etc.) will begin to form at this interface eventually leading to coating 

failure. 

Fur<strong>the</strong>rmore, <strong>the</strong> increased concentration <strong>of</strong> Hf at <strong>the</strong> substrate/coating interface shown in 

Figure 10 has been shown to lead to severe internal oxidation. Hafnium concentration <strong>and</strong> <strong>the</strong> 

grain boundary density <strong>of</strong> <strong>the</strong> coating can have a dramatic <strong>effect</strong> on <strong>the</strong> speed at which this 

process progresses. This is evidenced clearly in <strong>the</strong> NiAl-0.5Hf <strong>and</strong> NiAl-1Hf samples. As <strong>the</strong> 

<strong>hafnium</strong> concentration was increased, <strong>the</strong> oxidation at <strong>the</strong> coating/substrate interface became 

more severe. The concentration <strong>of</strong> <strong>hafnium</strong> at <strong>the</strong> interfaces increased from 2.5 at% in <strong>the</strong> NiAl- 

0.5Hf sample to 11.89 at% in <strong>the</strong> NiAl-1.0Hf sample. It has been shown that reactive elements 

such as <strong>hafnium</strong> change <strong>the</strong> growth mechanism <strong>of</strong> oxides in Al 2 O 3 formers by switching <strong>the</strong> 

outward diffusion <strong>of</strong> aluminum ions to <strong>the</strong> inward diffusion <strong>of</strong> oxygen ions [68]. 

Grain boundaries are high diffusion pathways for elements. Increasing <strong>the</strong> density <strong>of</strong> 

grain boundaries can lead to an increase in diffusion to <strong>the</strong> coating interfaces. This would 

explain <strong>the</strong> dramatic increase in <strong>hafnium</strong> concentration at <strong>the</strong> coating/substrate interface. The 

preoxidation grain sizes observed with this study are approximately one third <strong>of</strong> those reported in 

101

similar coatings produced by Ning et al. [25]. In <strong>the</strong>ir study, <strong>the</strong> <strong>hafnium</strong> concentration at <strong>the</strong> 

coating/substrate interface with NiAl-0.6Hf coatings was reported to be only 0.5 at% after 

annealing for four hours, approximately 1.0 at% lower than those observed with <strong>the</strong> NiAl-0.5Hf 

coatings. This difference in concentration indicates that <strong>the</strong>re is a higher diffusion rate <strong>of</strong> 

<strong>hafnium</strong> in <strong>the</strong> coatings developed in this study compared to previous results. This increased 

amount <strong>of</strong> <strong>hafnium</strong> combined with <strong>the</strong> oxygen transported through pinhole defects <strong>and</strong> along 

grain boundaries to <strong>the</strong> coating/substrate interface could lead to rapid oxidation along this 

interface. These features were not observed with <strong>the</strong> larger grained coatings oxidized by Ning et 

al. [25]. Also, <strong>the</strong> specific mass gains reported in that study were much lower than <strong>the</strong> in <strong>the</strong> 

ones measured in <strong>the</strong> present study (Figure 1). From this it can be concluded that reduced 

<strong>hafnium</strong> content along with <strong>the</strong> reduced grain boundary area in <strong>the</strong> coating (i.e., larger coating 

grain sizes) appears to lead to enhanced oxidation performance. 

6.8. Summary 

A series <strong>of</strong> β-NiAl-Cr-Hf coatings were prepared using DC magnetron sputtering to 

investigate <strong>the</strong> <strong>effect</strong>s <strong>of</strong> Hf <strong>and</strong> Cr concentration coupled with preoxidation annealing time on 

short-term iso<strong>the</strong>rmal oxidation behavior at 1050°C for times up to 96 hours. Specific mass 

gains generally increased with Hf concentration <strong>and</strong> annealing time, however, <strong>the</strong> addition <strong>of</strong> Cr 

to <strong>the</strong> highest Hf content alloy resulted in significantly lower mass gains, which suggest that it 

lowers its diffusion rate within <strong>the</strong> coating [27]. This hypo<strong>the</strong>sis is consistent with <strong>the</strong> recent 

work <strong>of</strong> Rigney et al. [69] <strong>and</strong> Hazel et al. [24]. 

102

Following oxidation, a detailed characterization was conducted to determine <strong>the</strong> phases 

that evolve <strong>and</strong> <strong>the</strong> <strong>effect</strong> on <strong>the</strong> microstructures <strong>of</strong> <strong>the</strong> coating. XRD spectra collected from <strong>the</strong> 

coating surfaces showed that <strong>the</strong> β-NiAl phase was prevalent with a (110) preferred orientation. 

Small amounts <strong>of</strong> θ-Al 2 O 3 <strong>and</strong> α-Al 2 O 3 were also detected from <strong>the</strong> TGO along with γ′-Ni 3 Al 

which forms as <strong>the</strong> aluminum is consumed from <strong>the</strong> coating to form <strong>the</strong> TGO. The NiAlCrHf 

samples did not contain any γ′ peaks suggesting that <strong>the</strong> diffusion <strong>of</strong> aluminum <strong>and</strong>/or oxygen is 

limited with <strong>the</strong>se coatings <strong>and</strong> retains <strong>the</strong> parent β-NiAl phase. 

The surfaces <strong>of</strong> <strong>the</strong> oxides were also imaged using an SEM to investigate <strong>the</strong> morphology 

<strong>and</strong> chemistry <strong>of</strong> <strong>the</strong> TGO. The TGO was observed to consist <strong>of</strong> a mixture <strong>of</strong> blade-like θ-Al 2 O 3 

<strong>and</strong> more equiaxed α-Al 2 O 3 . The scales formed on <strong>the</strong> ternary NiAl-Hf samples contained a 

large number <strong>of</strong> cracks. These cracks were more numerous in <strong>the</strong> NiAl-0.5Hf samples compared 

to <strong>the</strong> NiAl-1Hf samples. While <strong>the</strong> TGO thickness did not increase dramatically when <strong>the</strong> 

<strong>hafnium</strong> concentration was increased from 0.5 at.% to 1.0 at.%, <strong>the</strong> decrease in crack density 

with increased <strong>hafnium</strong> content suggests that <strong>the</strong> TGO adherence to <strong>the</strong> bond coat increases with 

<strong>hafnium</strong> content for <strong>the</strong> samples studied in this <strong>investigation</strong>. The NiAlCrHf coatings also 

contained cracks following oxidation. The cracks were wider <strong>and</strong> longer than those observed in 

<strong>the</strong> NiAl-Hf coatings; however, <strong>the</strong>y appear to form a uniform Al 2 O 3 layer that protected <strong>the</strong> 

underlying bond coat. No HfO 2 or Cr 2 O 3 formation was detected with any <strong>of</strong> <strong>the</strong> samples using 

SEM-EDS elemental mapping. 

Cross-sectional SEM images <strong>and</strong> SEM-EDS elemental maps were also collected. Small 

amounts <strong>of</strong> γ′ were observed in all <strong>of</strong> <strong>the</strong> samples as was <strong>hafnium</strong> segregation to <strong>the</strong> 

103

TGO/coating <strong>and</strong> coating substrate interfaces. In <strong>the</strong> NiAl-Hf samples, <strong>the</strong> oxides formed at both 

interfaces were found to be decorated with <strong>hafnium</strong>-rich precipitates, presumably HfO 2 . 

Interestingly, internal oxidation in <strong>the</strong> form <strong>of</strong> a discontinuous Al 2 O 3 layer was observed at <strong>the</strong> 

coating/substrate interface in <strong>the</strong> NiAl-Hf samples. The severity <strong>of</strong> <strong>the</strong> oxidation at this interface 

was found to increase with <strong>the</strong> annealing time. 

In <strong>the</strong> NiAlCrHf samples, <strong>the</strong> IDZ was completely retained <strong>and</strong> was accompanied by 

<strong>chromium</strong> enrichment. The TGO observed in <strong>the</strong>se coatings was much thinner than in <strong>the</strong> NiAl- 

Hf coatings. Hafnium segregation was not as significant in <strong>the</strong> NiAlCrHf samples. In <strong>the</strong>se 

coatings, precipitates formed along grain boundaries <strong>and</strong> appeared to coarsen with increased 

annealing time. Fur<strong>the</strong>rmore, while <strong>the</strong> NiAlCrHf coatings contain larger cracks that extend to 

<strong>the</strong> substrate, <strong>the</strong>re was no oxidation observed at <strong>the</strong> coating/substrate interface. As noted above, 

<strong>the</strong> NiAl-Hf coatings showed significant oxidation near <strong>the</strong> substrate/coating interface 

accompanied by significant <strong>hafnium</strong> segregation. Hafnium-rich precipitates were readily 

observed within cracks in <strong>the</strong> NiAl-Hf coatings; however, no <strong>hafnium</strong> or <strong>chromium</strong> segregation 

was detected in or near cracks in <strong>the</strong> NiAlCrHf coatings, which were found to form only Al 2 O 3 . 

The oxidation at <strong>the</strong> coating/substrate interface can be explained simply by examining <strong>the</strong> 

nature <strong>of</strong> zone T microstructures. Zone T microstructures have extremely fine grain sizes <strong>and</strong> 

<strong>of</strong>ten exhibit voids or pinhole type defects (sometimes referred to as leader defects in thick 

coatings) at grain boundaries. Both represent rapid diffusion pathways for oxygen in to <strong>the</strong> 

substrate while allowing for a corresponding diffusion <strong>of</strong> elements out <strong>of</strong> <strong>the</strong> substrate <strong>and</strong> 

coating. This results in oxidation along grain boundaries <strong>and</strong> at <strong>the</strong> coating substrate interfaces. 

104

The coatings developed by Ning et al. had denser zone 2 microstructures, which have 

significantly larger grain sizes, <strong>and</strong> thus fewer intergranular voids [25,37,70-72]. Ning’s 

coatings had grain sizes that were approximately four times larger than those produced in <strong>the</strong> 

present study. Thus, <strong>the</strong>y were more resistant to grain boundary diffusion <strong>and</strong> did not exhibit 

oxidation at <strong>the</strong> interface. It is possible to reduce or eliminate <strong>the</strong> coating/substrate oxidation 

layer by peening <strong>the</strong> coating surfaces. The deformation induced by peening will close any 

boundaries open to <strong>the</strong> free surface [73-75]. This technique is widely used to close up leaders 

<strong>and</strong> to improve <strong>the</strong> oxidation resistance <strong>of</strong> EB-PVD overlay coatings. 

EPMA line pr<strong>of</strong>iles for <strong>the</strong> coatings showed that <strong>the</strong> nickel <strong>and</strong> aluminum concentrations 

do not change appreciably within <strong>the</strong> coating; however, <strong>the</strong> results for <strong>hafnium</strong>, cobalt, <strong>and</strong> 

<strong>chromium</strong> <strong>of</strong>fer greater insight into <strong>the</strong> dynamics in <strong>the</strong> coatings. Hafnium was found to 

decorate <strong>the</strong> interfaces at concentrations as high as 10 at.% while cobalt <strong>and</strong> <strong>chromium</strong> diffuse 

from <strong>the</strong> substrate to a constant concentration <strong>of</strong> between 2-3 at.% for both elements. With <strong>the</strong> 

NiAlCrHf coatings, <strong>the</strong> <strong>chromium</strong> concentration stabilizes near 4 at.% throughout <strong>the</strong> coating 

<strong>and</strong> was not observed to enrich at boundaries. Large jumps in <strong>the</strong> <strong>chromium</strong> concentration 

within <strong>the</strong> IDZ were determined to be γ′ with <strong>chromium</strong> enrichment <strong>and</strong> also α-Cr with β-NiAl. 

Plan view TEM images <strong>and</strong> corresponding STEM analyses were also collected from <strong>the</strong> 

oxidized coatings. The average grain size <strong>of</strong> <strong>the</strong> NiAl-1Hf increases from 178 to 229 μm 

following oxidation for <strong>the</strong> two <strong>and</strong> four hour annealed samples respectively. This increase is 

suspected to be caused by <strong>the</strong> increased amount <strong>of</strong> recrystallization as <strong>the</strong> annealing time was 

105

increased. The precipitates were heterogeneously distributed with larger precipitates forming at 

<strong>the</strong> grain boundaries. Chemical analysis <strong>of</strong> <strong>the</strong>se precipitates determined that <strong>the</strong>y are Heusler 

β′-Ni 2 AlHf. The NiAlCrHf samples have a much larger grain size <strong>and</strong> exhibit little change in 

grain size with annealing conditions. The average grain size is almost twice that observed with 

<strong>the</strong> NiAl-1Hf samples. However, <strong>the</strong> precipitates that formed with <strong>the</strong>se samples are confined 

primarily to <strong>the</strong> grain boundaries <strong>and</strong> coarsen in excess <strong>of</strong> 100 nm. These large precipitates are 

thought to provide enhanced barriers to diffusion along <strong>the</strong> grain boundaries <strong>and</strong> inhibit <strong>the</strong> 

oxidation <strong>of</strong> <strong>the</strong> bond coat. It is also possible that <strong>the</strong> smaller grain size plays a significant role 

in <strong>the</strong> oxidation performance <strong>of</strong> <strong>the</strong> bond coating through increased grain boundary volume. 


The authors acknowledge <strong>the</strong> support <strong>of</strong> <strong>the</strong> National Science Foundation (NSF) under 

Award No. DMR-0504950 <strong>and</strong> <strong>the</strong> National Aeronautics <strong>and</strong> Space Administration (NASA) 

under contract NN-X08AT21H. The use <strong>of</strong> <strong>the</strong> facilities supported by <strong>the</strong> Central Analytical 

facility at <strong>the</strong> University <strong>of</strong> Alabama <strong>and</strong> <strong>the</strong> Materials Diagnostics Laboratory at NASA’s 

Marshall Space Flight Center in Huntsville, Alabama is also gratefully acknowledged. 

106

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112

Table 1. Nominal target concentrations in atomic percent. 

Alloy Ni Al Cr Hf 

NiAl 51 49 --- --- 

NiAl-0.5Hf 51 48 --- 0.5 

NiAl-1Hf 51 48 --- 1.0 

NiAlCrHf 50 44 5.0 1.0 

113

Table 2. Nominal chemical compositions <strong>of</strong> coatings after oxidation. 


NiAl Nominal --- 51 49 --- --- --- 


1000°C/2h EDS 56.66 ± 3.17 40.64 ± 3.56 1.15 ± 0.25 1.61 ± 0.33 0.00 ± 0.00 

1000°C/4h WDS 56.50 ± 0.75 40.34 ± 0.99 1.53 ± 0.23 1.60 ± 0.21 0.04 ± 0.02 

NiAl-0.5Hf Nominal --- 50 49.5 --- --- 0.5 


1000°C/2h WDS 55.21 ± 0.87 41.76 ± 0.94 1.39 ± 0.47 1.07 ± 0.29 0.57 ± 0.08 

1000°C/4h WDS 56.74 ± 1.39 39.57 ± 1.32 1.79 ± 0.20 1.32 ± 0.12 0.58 ± 0.10 

NiAl-1Hf Nominal --- 51 48 --- --- 1.0 


1000°C/2h WDS 55.99 ± 0.41 38.67 ± 0.59 2.04 ± 0.25 2.57 ± 0.15 0.73 ± 0.16 

1000°C/4h WDS 53.44 ± 0.79 43.19 ± 0.72 1.33 ± 0.43 1.09 ± 0.30 0.85 ± 0.14 

NiAlCrHf Nominal --- 50 44 5 --- 1 


1000°C/2h WDS 51.73 ± 1.64 41.60 ± 0.81 4.65 ± 0.74 1.21 ± 0.23 0.81 ± 0.06 

1000°C/4h WDS 53.23 ± 1.42 40.50 ± 1.29 4.28 ± 0.92 1.16 ± 0.21 0.81 ± 0.10 

Table 3. Nominal chemical compositions <strong>of</strong> coatings after oxidation for 96h at 1050°C. 


NiAl-0.5Hf 1000°C/2h WDS 58.08 ± 0.45 37.08 ± 0.57 2.31 ± 0.13 2.26 ± 0.15 0.27 ± 0.06 

1000°C/4h WDS 56.26 ± 0.73 38.39 ± 0.72 2.66 ± 0.10 2.11 ± 0.28 0.65 ± 0.53 

NiAl-1Hf 1000°C/2h WDS 58.46 ± 0.82 36.02 ± 0.66 2.12 ± 0.11 2.19 ± 0.14 1.21 ± 0.44 

1000°C/4h WDS 62.00 ± 1.36 32.26 ± 1.25 2.41 ± 0.18 2.64 ± 0.14 0.51 ± 0.15 

NiAlCrHf 1000°C/2h WDS 58.24 ± 1.10 33.63 ± 1.22 3.98 ± 0.32 3.04 ± 0.11 1.11 ± 0.51 

1000°C/4h WDS 59.49 ± 2.34 31.48 ± 1.01 4.90 ± 2.46 3.13 ± 0.17 1.01 ± 0.52 

114

0.8 

Specific Mass Change (mg/cm 2 ) 

0.7 

0.6 

0.5 

0.4 

0.3 

0.2 

0.1 

0.0 

NiAl-0.5Hf 

NiAl-1.0Hf (I) 

NiAl-1.0Hf (C) 

NiAlCrHf 

β-NiAl 

0 20 40 60 80 100 

Iso<strong>the</strong>rmal Oxidation Time (h) 

(a) 

0.7 

Specific Mass Change (mg/cm 2 ) 

0.6 

0.5 

0.4 

0.3 

0.2 

0.1 

0.0 

NiAl-0.5Hf 

NiAl-1.0Hf 

NiAlCrHf (I) 

NiAlCrHf (C) 

β-NiAl 

0 20 40 60 80 100 

Iso<strong>the</strong>rmal Oxidation Time (h) 

(b) 

Figure 1. Specific mass changes measured with iso<strong>the</strong>rmal oxidation (I) at 1050°C for NiAl-X 

coatings on René N5. Additional samples were subjected to long cycle cyclic oxidation (C) at 

1050°C. 

115

β-NiAl 

γ'-Ni 3 

Al 

α-Al 2 

O 3 

Normalized Intensity (abu) 

θ-Al 2 

O 3 

20 30 40 50 60 70 80 90 100 

Degrees (2θ) 

(a) 

β-NiAl 

γ'-Ni 3 

Al 

Normalized Intensity (abu) 

α-Al 2 

O 3 

θ-Al 2 

O 3 

20 30 40 50 60 70 80 90 100 

Degrees (2θ) 

(b) 

Figure 2. X-ray diffraction spectra collected from samples that were iso<strong>the</strong>rmally oxidized at 

1050°C for 96 hours. Samples are divided according to annealing time with two <strong>and</strong> four hours 

shown in parts (a) <strong>and</strong> (b) respectively. 

116

(a) 

(b) 

Figure 3. SEM images collected from <strong>the</strong> surface <strong>of</strong> NiAl-0.5Hf samples that were annealed for 

(a) two <strong>and</strong> (b) four hours <strong>and</strong> <strong>the</strong>n oxidized for 96 hours. 

117

(a) 

(b) 

Figure 4. SEM images collected from <strong>the</strong> surface <strong>of</strong> NiAl-1Hf samples that were annealed for 


118

(a) 

(b) 

Figure 5. SEM images collected from <strong>the</strong> surface <strong>of</strong> NiAlCrHf samples that were annealed for 


119

(a) 

(b) 

Figure 6. Cross-sectional SEM images <strong>of</strong> NiAl-0.5Hf samples: annealed for (a) two <strong>and</strong> (b) four 

hours <strong>and</strong> oxidized for 96 hours. 

120

(a) 

(b) 

Figure 7. Cross-sectional SEM images <strong>of</strong> NiAl-1Hf samples: annealed for (a) two <strong>and</strong> (b) four 


121

(a) 

(b) 

Figure 8. Cross-sectional SEM images <strong>of</strong> NiAlCrHf samples: annealed for (a) two <strong>and</strong> (b) four 


122

70 

Coating IDZ Substrate 

Concentration (at.%) 

60 

50 

40 

30 

20 

Aluminum 

Nickel 

Chromium 

Hafnium 

Cobalt 

10 

0 

0 10 20 30 40 50 60 70 80 90 100 110 

Depth from Surface (μm) 

(a) 

70 



60 

50 

40 

30 

20 

Aluminum 

Nickel 

Chromium 

Cobalt 

Hafnium 

10 

0 

0 10 20 30 40 50 60 70 80 90 100 110 


(b) 

Figure 9. Elemental composition pr<strong>of</strong>ile collected from <strong>the</strong> NiAl-0.5Hf samples that were 

annealed for (a) two hours (b) four hours <strong>and</strong> oxidized for 96 hours. 

123

70 



60 

50 

40 

30 

20 

Aluminum 

Nickel 

Chromium 

Hafnium 

Cobalt 

10 

0 

0 10 20 30 40 50 60 70 


(a) 

70 



60 

50 

40 

30 

20 

Aluminum 

Nickel 

Chromium 

Hafnium 

Cobalt 

10 

0 

0 10 20 30 40 50 60 70 


(b) 

Figure 10. Elemental composition pr<strong>of</strong>ile collected from <strong>the</strong> NiAl-1Hf samples that were 


124

70 


60 


50 

40 

30 

20 

Aluminum 

Nickel 

Chromium 

Hafnium 

Cobalt 

10 

0 

0 10 20 30 40 50 60 70 80 90 100 110 


(a) 

70 


60 


50 

40 

30 

20 

Aluminum 

Nickel 

Chromium 

Hafnium 

Cobalt 

10 

0 

0 10 20 30 40 50 60 70 80 90 100 110 


(b) 

Figure 11. Elemental composition pr<strong>of</strong>ile collected from <strong>the</strong> NiAlCrHf samples that were 


125

(a) 

(c) 

Aluminum 

10 μm 10 μm 

Oxygen 

(b) 

(d) 

TGO 

γ′ 

Epoxy 

Hafnium 

10 μm 

Substrate 

SRZ 

IDZ 

10 μm 

Figure 12. Elemental maps <strong>and</strong> SEM cross-sectional images collected from <strong>the</strong> NiAl-0.5Hf samples that were annealed for (a-d) two 

<strong>and</strong> (e-h) four hours <strong>and</strong> oxidized for 96 hours. 

126

(e) 

(g) 

Aluminum 

10 μm 

Oxygen 

10 μm 

(f) 

(h) 

Epoxy 

γ′ 

TGO 

SRZ 

IDZ 

Hafnium 

10 μm 

Substrate 

10 μm 

Figure 12. Cont’d. 

127

(a) 

(c) 

Aluminum 

10 μm 

Oxygen 

10 μm 

(b) 

(d) 

TGO 

γ′ 

IDZ 

Hafnium 

10 μm 

10 μm 10 μm 

Substrate 

SRZ 

Figure 13. Elemental maps <strong>and</strong> SEM cross-sectional images collected from <strong>the</strong> NiAl-1Hf samples that were annealed for (a-d) two 

<strong>and</strong> (e-h) four hours <strong>and</strong> oxidized for 96 hours. 

128

(e) 

(g) 

Aluminum 

10 μm 

Oxygen 

10 μm 

(f) 

(h) 

Hafnium 

10 μm 10 μm 

Figure 13. Cont’d. 

129

(a) 

TGO 

Cu plating 

IDZ 

SRZ 

10 μm 

Figure 14. Elemental maps <strong>and</strong> SEM cross-sections collected from NiAlCrHf coatings after 

annealing for: two hours (a-e) <strong>and</strong> four hours (f-j) four hours followed by oxidizing for 96 hours. 

130

(b) 

(d) 

Aluminum 

10 μm 

Oxygen 

10 μm 

(c) 

(e) 

Hafnium 

10 μm 10 μm 

Chromium 

Figure 14. Con’td. 

131

(f) 

Cu Cu plating plating 

TGO TGO 

IDZ 

SRZ 

10 μm 

Figure 14. cont’d. 

132

(b) 

(d) 

Aluminum 

10 μm 

Oxygen 

10 μm 

(c) 

(e) 

Hafnium 

10 μm 10 μm 

Chromium 

Figure 14. cont’d. 

133

Figure 15. Plan view TEM images <strong>of</strong> NiAl-0.5Hf iso<strong>the</strong>rmally oxidized for 96 hours imaged 

using bright field <strong>and</strong> STEM-HAADF techniques respectively: annealed for (a,b) two <strong>and</strong> (c,d) 

four hours. The boxes in (c) <strong>and</strong> (d) highlight <strong>the</strong> regions imaged in (a) <strong>and</strong> (b). 

134

(a) 

(c) 

150 nm 

150 nm 

(b) 

(d) 

200 nm 

200 nm 

Figure 16. Plan view TEM images <strong>of</strong> NiAl-1Hf iso<strong>the</strong>rmally oxidized for 96 hours imaged using 

bright field <strong>and</strong> STEM-HAADF techniques respectively: annealed for (a,b) two <strong>and</strong> (c,d) four 

hours. The boxes in (b) <strong>and</strong> (d) highlight <strong>the</strong> regions imaged in (a) <strong>and</strong> (c). 

135

(a) 

300 nm 

(c) 

300 nm 

300 nm 

(b) 

300 nm 

(d) 

Figure 17. Plan view TEM samples <strong>of</strong> NiAlCrHf iso<strong>the</strong>rmally oxidized for 96 hours imaged 

using bright field <strong>and</strong> STEM-HAADF techniques: annealed for (a,b) two <strong>and</strong> (b,c) four hours. 

The box in (a) highlights <strong>the</strong> region imaged in (b). 

136

CHAPTER VI 

FACTORS INFLUENCING INTERDIFFUSION BETWEEN SPUTTERED NIAL 

OVERLAY COATINGS AND Ni-BASED SUPERALLOY SUBSTRATES SUBJECTED 

TO HIGH TEMPERATURES 

7.1 INTRODUCTION 

Overlay coatings generally degrade over time due to oxidative attack <strong>and</strong> interdiffusion 

with <strong>the</strong> underlying substrate [1-4]. Numerous studies have been conducted <strong>of</strong> interdiffusion 

between diffusion aluminide <strong>and</strong> overlay coatings on superalloy substrates [3-14]. The majority 

<strong>of</strong> <strong>the</strong>se studies have focused on modern diffusion aluminide or MCrAlY-type overlay coatings 

used for power generation applications. These studies have greatly improved our underst<strong>and</strong>ing 

<strong>of</strong> <strong>the</strong> influences <strong>of</strong> substrate composition on <strong>the</strong> microstructures <strong>and</strong> performance <strong>of</strong> diffusion 

aluminide <strong>and</strong> MCrAlY coatings. However, research conducted over <strong>the</strong> last decade suggests 

that a series <strong>of</strong> new overlay coating concepts could yield significant improvements in 

performance over state-<strong>of</strong>-<strong>the</strong>-art coatings [15-19]. Considering <strong>the</strong> potential importance <strong>of</strong> 

advanced β-NiAl-type overlay coatings [19], relatively little information is available detailing 

<strong>the</strong>ir responses to high temperature annealing or oxidation. 

Recently, Perez et al. [20] studied interdiffusion using diffusion couples between β-NiAl 

<strong>and</strong> a series <strong>of</strong> commercial superalloys, with <strong>the</strong> goal <strong>of</strong> quantifying <strong>the</strong> rate <strong>of</strong> Al interdiffusion 

as a function <strong>of</strong> initial superalloy composition. This study showed that increasing <strong>the</strong> 

137

concentrations <strong>of</strong> Cr, Mo, <strong>and</strong> Ti in <strong>the</strong> superalloy <strong>effect</strong>ively increased <strong>the</strong> Al <strong>effect</strong>ive 

interdiffusion coefficient into <strong>the</strong> superalloy; whereas increasing <strong>the</strong> content <strong>of</strong> Al, Ta, <strong>and</strong> W 

<strong>effect</strong>ively reduced it. This study, which was conducted within <strong>the</strong> context <strong>of</strong> improving <strong>the</strong> 

performance <strong>of</strong> β+γ type (i.e., MCrAlY) overlay coatings, suggests that reducing Cr, Mo, or Ti in 

<strong>the</strong> superalloy or increasing Al, Ta, or W can slow coating degradation. The purpose <strong>of</strong> <strong>the</strong> 

present note is to consider how variations in coating composition as opposed to substrate 

composition might influence <strong>the</strong> properties <strong>of</strong> β-NiAl based overlay coatings. The following 

note provides a synopsis <strong>of</strong> <strong>the</strong> reactions that occur when β-NiAl coatings containing high 

concentrations <strong>of</strong> reactive elements (Hf or Zr) <strong>and</strong> Cr are deposited on second generation Nibased 

superalloy substrates <strong>and</strong> exposed to high temperatures. 

7.2 DISCUSSION OF RESULTS 

The overlay coatings developed in this study were based upon a slightly Ni-rich β-NiAl 

phase to which Hf <strong>and</strong> Cr were added. Results for each coating are presented below in <strong>the</strong> 

context <strong>of</strong> establishing patterns <strong>of</strong> interdiffusion <strong>and</strong> oxidation. 

In comparison to <strong>the</strong> two Ni-based superalloy substrates employed in this study, all <strong>of</strong> <strong>the</strong> 

coatings were higher in Al content than <strong>the</strong> substrates. As a result, Al diffuses from <strong>the</strong> coatings 

into <strong>the</strong> substrates making <strong>the</strong>m less able to support <strong>the</strong> growth <strong>of</strong> a protective alumina scale. In 

contrast, Ni, Co, <strong>and</strong> Cr in <strong>the</strong> substrates, diffuses into <strong>the</strong> coatings. Reactive elements such as 

Zr, Y, <strong>and</strong> Hf strongly influence scale adherence during iso<strong>the</strong>rmal <strong>and</strong> cyclic oxidation. In most 

coatings <strong>the</strong>y are present at such low concentrations that <strong>the</strong>ir dilution by interdiffusion with <strong>the</strong> 

138

substrate or by diffusion to <strong>the</strong> surface <strong>of</strong> <strong>the</strong> coating is generally not considered. However, it is 

well known that when <strong>the</strong>se elements are present within <strong>the</strong> substrate, will diffuse rapidly into 

<strong>the</strong> coating or all <strong>of</strong> <strong>the</strong> way through it, particularly when in <strong>the</strong> presence <strong>of</strong> an oxygen potential 

gradient [4,21-26]. When <strong>the</strong> coating becomes sufficiently depleted <strong>of</strong> Al due to interdiffusion 

<strong>and</strong>/or oxidation, it can lose its ability to form a protective α-Al 2 O 3 scale resulting in <strong>the</strong> 

formation <strong>of</strong> less protective oxides (e.g., NiO, NiAl 2 O 4 , HfO 2 , etc.) <strong>and</strong> in coating failure. 

It is evident that modifications are necessary to improve coating performance <strong>and</strong> 

stability <strong>of</strong> β-NiAl coatings. Based upon oxidation studies performed on model bulk alloys, it 

has been proposed that a coating capable <strong>of</strong> maintaining a high reactive element content could 

exhibit properties comparable to state-<strong>of</strong>-<strong>the</strong> art (Ni,Pt)Al diffusion aluminides [15]. This was 

attempted in <strong>the</strong> present study. 

Additions <strong>of</strong> Hf produced significant reductions in coating grain size plus a significant 

increase in <strong>the</strong> amount <strong>of</strong> Hf at <strong>the</strong> coating/TGO <strong>and</strong> coating substrate interfaces. This 

enrichment was as high as 12 at.% in <strong>the</strong> NiAl-1Hf coating. The reduced grain size combined 

with <strong>the</strong> high Hf content at <strong>the</strong> coating/substrate interface contributed towards oxidation which 

accelerated Al depletion leading to a transformation from β to γ′. 

It is well known that reactive element <strong>additions</strong> can have a significant impact on <strong>the</strong> 

oxidation behavior <strong>of</strong> alumina- <strong>and</strong> chromia-forming alloys. Some <strong>of</strong> <strong>the</strong> results generated in 

this study strongly support <strong>the</strong> dynamic segregation model proposed by Pint [23]. In this model, 

it is hypo<strong>the</strong>sized that reactive-element ions segregate to scale grain boundaries <strong>and</strong> to <strong>the</strong> metal- 

139

oxide interface, which <strong>the</strong>y use as pathways for diffusion out <strong>of</strong> <strong>the</strong> metal substrate or coating to 

<strong>the</strong> gas interface <strong>of</strong> <strong>the</strong> scale. The driving force for this diffusion is <strong>the</strong> oxygen potential gradient 

across <strong>the</strong> scale. The enrichment <strong>of</strong> <strong>the</strong> scale grain boundaries by reactive elements results in 

scale growth via <strong>the</strong> inward diffusion <strong>of</strong> oxygen. Simultaneously, reactive-element enrichment 

at <strong>the</strong> metal-oxide interface inhibits <strong>the</strong> growth <strong>of</strong> interfacial voids, improving scale adhesion. In 

this study, <strong>hafnium</strong>, initially present in solid solution, was observed to precipitate out at grain 

boundaries <strong>and</strong> to segregate to <strong>the</strong> coating/TGO <strong>and</strong> coating/substrate interfaces. During 

oxidation, this led to <strong>the</strong> development <strong>of</strong> a mixed oxide comprised primarily <strong>of</strong> alumina <strong>and</strong> 

hafnia. When this occurs, <strong>the</strong> hafnia diffuses to grain boundaries within <strong>the</strong> TGO <strong>and</strong> limits <strong>the</strong> 

inward transport <strong>of</strong> oxygen through <strong>the</strong> TGO to <strong>the</strong> coating interface. The major <strong>effect</strong> <strong>of</strong> this 

observation was a reduction on <strong>the</strong> growth rate <strong>of</strong> <strong>the</strong> TGO during oxidation. As expected, this 

<strong>effect</strong> becomes more significant with increasing <strong>hafnium</strong> concentration. However, this does not 

appear to be <strong>the</strong> case when small amounts <strong>of</strong> <strong>chromium</strong> were added to <strong>the</strong> NiAl-1Hf coating. 

These coatings appear to rapidly form a thin, protective alumina scale <strong>and</strong> retain <strong>the</strong> <strong>hafnium</strong> 

within <strong>the</strong> coating with oxidation times up to 100 hours <strong>of</strong> exposure. This observation suggests 

that <strong>the</strong>re exist limits on <strong>the</strong> dynamic segregation model <strong>and</strong> its application to similar coating 

systems. Interestingly, <strong>the</strong> amount <strong>of</strong> aluminum depletion within <strong>the</strong> coating during annealing or 

subsequent oxidation does not appear to be a strong function <strong>of</strong> <strong>the</strong> alloying elements used in this 

study. The coatings developed with this study have smaller average grain sizes <strong>and</strong> exhibit a 

zone T microstructure as opposed to <strong>the</strong> zone 2 coatings prepared by Ning [27-29]. The zone T 

microstructures contain a high concentration <strong>of</strong> pinhole defects that provide a direct pathway for 

oxygen to diffuse to <strong>the</strong> coating/substrate interface. This combined with <strong>the</strong> large grain 

boundary volume results in <strong>the</strong> oxidation that is observed at <strong>the</strong> coating/substrate interface with 

140

<strong>the</strong> NiAl-Hf coatings. As <strong>the</strong> <strong>hafnium</strong> concentration increases at <strong>the</strong> coating substrate interface, 

severe internal oxidation was observed. This is most evident with <strong>the</strong> NiAl-1Hf coatings. With 

<strong>the</strong> NiAlCrHf coatings, <strong>the</strong> rapid initial formation <strong>of</strong> an α-Al 2 O 3 scale dramatically limits <strong>the</strong> 

diffusion <strong>of</strong> oxygen through <strong>the</strong> TGO to <strong>the</strong> coating. This switches <strong>the</strong> mass transfer limiting 

mechanism for <strong>the</strong> growth <strong>of</strong> <strong>the</strong> oxide from <strong>the</strong> transport <strong>of</strong> aluminum ions to <strong>the</strong> surface to <strong>the</strong> 

diffusion <strong>of</strong> oxygen ions to <strong>the</strong> bond coating. As this occurs, <strong>the</strong> growth rate <strong>of</strong> <strong>the</strong> oxide is 

reduced significantly. It was also shown that with <strong>the</strong> NiAlCrHf coatings <strong>hafnium</strong> does not 

migrate to <strong>the</strong> interfaces as observed with <strong>the</strong> NiAl-Hf coatings. This is thought to be caused by 

<strong>the</strong> increased aluminum diffusion into <strong>the</strong> substrate along with <strong>the</strong> formation <strong>of</strong> a protective 

TGO. 

7.3 REFERENCES 

[1] J. R. Nicholls, "Designing Oxidation-Resistant Coatings," JOM 52 (2000) 28-35. 


Bulletin 28 (2003) 659-670. 

[3] Y. Tamarin, Protective Coatings for Turbine Blades, (ASM International, Materials Park, 

OH, 2002). 


[5] C. S. Richard, G. Beranger, J. Lu <strong>and</strong> J. F. Flavenot, "The influences <strong>of</strong> heat treatments 

<strong>and</strong> interdiffusion on <strong>the</strong> adhesion <strong>of</strong> plasma-sprayted NiCrAlY coatings," Surface <strong>and</strong> 

Coatings Technology 82 (1996) 99-109. 

[6] J. H. Chen <strong>and</strong> J. A. Little, "Degradation <strong>of</strong> <strong>the</strong> platinum aluminide coating on CMSX4 at 

1100 o C," Surface <strong>and</strong> Coatings Technology 92 (1997) 69-77. 

[7] M. Li, X. Sun, W. Hu <strong>and</strong> H. Guan, "Microstructural changes <strong>and</strong> elemental diffusion <strong>of</strong> 

sputtered NiCrAlY coating on a Ni-base SC superalloy subjected to high temperature," 

Materials Letters 61 (2007) 5169-5172. 

141

[8] M. H. Li, Z. Y. Zhang, X. F. Sun, J. G. Li, F. S. Yin, W. Y. Hu, H. R. Guan <strong>and</strong> Z. Q. 

Hu, "Oxidation behavior <strong>of</strong> sputter-deposited NiCrAlY coating," Surface <strong>and</strong> Coatings 

Technology 165 (2003) 241-247. 

[9] H. Li, A. Hesnawi, X. Fan, S. Gong <strong>and</strong> X. Huibin, "Inter-diffusion <strong>and</strong> oxidation 

behavior in electron-beam evaporated NiAl coatings," Vacuum 81 (2006) 329-337. 



[11] H. Guo, Y. Cui, H. Peng <strong>and</strong> S. Gong, "Improved cyclic oxidation resistance <strong>of</strong> electron 

beam physical vapor deposited nano-oxide dispersed β-NiAl coatings for Hf-containing 

superalloy," Corrosion Science 52 (2010) 1440-1446. 



516 (2008) 5732-5735. 

[13] J. T. Guo <strong>and</strong> C. M. Xu, "Effect <strong>of</strong> NiAl microcrystalline coating on <strong>the</strong> high-temperature 

oxidation behavior <strong>of</strong> NiAl-28Cr-5Mo-1Hf," Oxidation <strong>of</strong> Metals 58 (2002) 457-468. 







[16] J. D. Rigney, R. Darolia, W. S. Walston <strong>and</strong> R. Corderman, U. S. Patent Number 

6,153,313, (2000). 

[17] R. Darolia, J. D. Rigney <strong>and</strong> R. J. Grylls, U. S. Patent Number 6,291,084, (2001). 

[18] J. D. Rigney, C.-P. Lee <strong>and</strong> R. Darolia, United States Patent Number 7,094,444 B2, 

(2006). 





[20] E. Perez, T. Patterson <strong>and</strong> Y. Sohn, "Interdiffusion analysis for NiAl versus superalloys 

diffusion couples," Journal <strong>of</strong> Phase Equilibria <strong>and</strong> Diffusion 27 (2006) 659-664. 

[21] J. G. Smeggil <strong>and</strong> N. S. Bornstein, "Hafnium <strong>effect</strong>s <strong>and</strong> protective aluminide coating 

performance," in High Temperature Corrosion in Energy Systems, edited by M. F. 

Rothman (The Metallurgical Society <strong>of</strong> AIME, Warrendale, PA, 1985) p. 681-695. 

142

[22] B. A. Pint <strong>and</strong> K. L. More, "Characterization <strong>of</strong> Alumina Interfaces in TBC Systems," 

Journal <strong>of</strong> Materials Science 44 (2009) 1676-1686. 



[24] P. Y. Hou, "Segregation behavior at TGO/bondcoat interfaces," Journal <strong>of</strong> Materials 

Science 44 (2009) 1711-1725. 

[25] J. Liu, Y. H. Sohn <strong>and</strong> K. S. Murphy, "Microstructural evolution <strong>of</strong> durable <strong>the</strong>rmal 

barrier coatings with Hf <strong>and</strong>/or Y modified CMSX-4 superalloy substrates," Materials 

Science Forum 539-543 (2007) 1206-1211. 

[26] J. Liu, Mechanisms <strong>of</strong> lifetime improvement in <strong>the</strong>rmal barrier coatings with <strong>hafnium</strong> 

<strong>and</strong>/or yttrium modification <strong>of</strong> CMSX-4 superalloy substrates, Ph.D. Dissertation, 

University <strong>of</strong> Central Florida, Orl<strong>and</strong>o, FL, 2007. 





201-209. 

[29] B. Ning <strong>and</strong> M. L. Weaver, "A Preliminary Study <strong>of</strong> DC Magnetron Sputtered NiAl-Hf 


143

CHAPTER VIII 

SUMMARY AND CONCLUSIONS 

NiAl coatings containing <strong>hafnium</strong> <strong>and</strong> <strong>chromium</strong> have been produced by DC magnetron 

sputtering onto CMSX-4 <strong>and</strong> René N5 superalloy substrates. The deposition variables, including 

deposition power, substrate temperature, <strong>and</strong> working gas pressure, were tailored to produce a 

zone T microstructure. Following deposition, <strong>the</strong> samples were annealed at 1000°C for up to 

four hours to remove any defects from deposition. During annealing, an IDZ between <strong>the</strong> 

substrate <strong>and</strong> <strong>the</strong> coating develops <strong>and</strong> coarsens with increased annealing time. The IDZ is 

formed due to <strong>the</strong> difference in chemistries <strong>of</strong> <strong>the</strong> substrate <strong>and</strong> <strong>the</strong> coating; primarily it is <strong>the</strong> 

large concentration difference in aluminum that provides <strong>the</strong> driving force for diffusion to occur. 

As aluminum diffuses from <strong>the</strong> coating, o<strong>the</strong>r elements, mainly <strong>chromium</strong> <strong>and</strong> cobalt, diffuse 

upward from <strong>the</strong> substrate into <strong>the</strong> coating. Adding small amounts <strong>of</strong> <strong>hafnium</strong> to <strong>the</strong> coatings 

reduced <strong>the</strong> thickness <strong>of</strong> <strong>the</strong> IDZ after annealing. This indicates that <strong>the</strong> <strong>hafnium</strong> reduces <strong>the</strong> 

diffusion <strong>of</strong> aluminum from <strong>the</strong> coating <strong>and</strong> nickel, <strong>chromium</strong>, <strong>and</strong> cobalt from <strong>the</strong> superalloy. It 

was also noted that <strong>the</strong> interfaces <strong>of</strong> <strong>the</strong> coating with <strong>the</strong> free surface <strong>and</strong> <strong>the</strong> substrate were 

enriched with <strong>hafnium</strong> after annealing. 

TEM analyses <strong>of</strong> <strong>the</strong> coatings showed that <strong>the</strong> coatings are deposited as a solid solution. 

The grains were small <strong>and</strong> irregular in shape; however, <strong>the</strong> grains recover <strong>and</strong> recrystallize 

yielding a more equiaxed grain structure after four hours <strong>of</strong> annealing. The NiAl-Hf <strong>and</strong> 

144

NiAlCrHf coatings are deposited as a supersaturated solid solution <strong>and</strong> precipitates form with 

only one hour <strong>of</strong> annealing. As annealing time was increased, <strong>the</strong> precipitates coarsen <strong>and</strong> make 

it possible to determine <strong>the</strong>ir chemistry. TEM-EDS measurements indicated that most <strong>of</strong> <strong>the</strong> 

precipitates in <strong>the</strong> NiAl-Hf <strong>and</strong> NiAlCrHf coatings were β’-Ni 2 AlHf. The precipitates were 

distributed heterogeneously within <strong>the</strong> coatings with larger precipitates forming at grain 

boundaries. 

APT experiments on <strong>the</strong> ternary NiAl-Hf coatings showed that some <strong>of</strong> <strong>the</strong> precipitates 

were HfC. The formation <strong>of</strong> HfC precipitates was traced to <strong>the</strong> annealing <strong>of</strong> as-deposited 

coatings in a graphite furnace which resulted in <strong>the</strong> incorporation <strong>of</strong> high levels <strong>of</strong> carbon into 

<strong>the</strong> coatings. Hafnium readily forms carbides, which is undesirable because it reduces <strong>the</strong> 

amount <strong>of</strong> <strong>effect</strong>ive <strong>hafnium</strong> available that is necessary to provide enhanced oxidation 

protection. Therefore, a tube furnace was adapted to anneal <strong>the</strong> samples in <strong>the</strong> absence <strong>of</strong> carbon 

to observe <strong>the</strong> <strong>effect</strong> on <strong>the</strong> coating properties. APT analyses <strong>of</strong> <strong>the</strong> coatings that were annealed 

in <strong>the</strong> tube furnace showed <strong>the</strong> coatings to contain only 53 ppm <strong>of</strong> carbon compared to <strong>the</strong> ~700 

ppm with <strong>the</strong> samples annealed in <strong>the</strong> graphite furnace. This provided a more accurate 

representation <strong>of</strong> <strong>the</strong> <strong>effect</strong> <strong>of</strong> <strong>hafnium</strong> concentration on <strong>the</strong> performance <strong>of</strong> <strong>the</strong> coatings during 

iso<strong>the</strong>rmal oxidation. Similar analyses confirmed that α-Cr precipitates form within <strong>the</strong> 

NiAlCrHf coatings following annealing. 

The annealed samples were subjected to a series <strong>of</strong> short-term iso<strong>the</strong>rmal oxidation tests 

at 1050°C for times up to 96 hours <strong>of</strong> exposure. The samples that were annealed for two hours 

had higher normalized mass gains than <strong>the</strong> samples that were annealed for four hours. This 

145

indicates that <strong>the</strong> TGO forming on <strong>the</strong> samples that were annealed for two hours grows at a 

higher rate than those annealed for four hours. Increasing <strong>the</strong> <strong>hafnium</strong> concentration from 0.5- 

1.0at.% resulted in a corresponding mass gain with increased exposure. However, adding 5.0at% 

Cr to <strong>the</strong> NiAl-1Hf coatings leads to a decrease in mass gains indicating that <strong>the</strong> NiAlCrHf 

samples form a thin, protective oxide. 

SEM images <strong>of</strong> <strong>the</strong> coatings following oxidation for 96 hours revealed a surprising 

discovery. During oxidation, <strong>the</strong> coatings undergo additional interdiffusion resulting in <strong>the</strong> 

formation <strong>of</strong> voids at <strong>the</strong> coating/substrate interface. Elemental mapping indicated that <strong>the</strong>se 

voids were Al 2 O 3 with small amounts <strong>of</strong> HfO 2 incorporated within <strong>the</strong> Al 2 O 3 . The amount <strong>of</strong> 

oxidation observed at <strong>the</strong> coating/substrate interface increased with annealing time. Fur<strong>the</strong>r 

<strong>investigation</strong> using backscattered electron imaging revealed that dark spots formed within <strong>the</strong> 

coatings. These were originally attributed to <strong>the</strong> Hf-rich phases being removed during st<strong>and</strong>ard 

metallographic sample preparation; however, <strong>the</strong> elemental maps confirmed that <strong>the</strong>se dark spots 

were Al-rich. While large cracks extended from <strong>the</strong> free surface to <strong>the</strong> substrate in <strong>the</strong> NiAlCrHf 

coatings, <strong>the</strong>y did not exhibit <strong>the</strong> same oxidation at <strong>the</strong> coating/substrate interface <strong>and</strong> no voids 

were observed with any <strong>of</strong> <strong>the</strong>se coatings. 

The oxidation at <strong>the</strong> coating/substrate interface can be explained simply by examining <strong>the</strong> 

nature <strong>of</strong> zone T microstructures. Zone T microstructures exhibit voids at grain boundaries 

which represent rapid diffusion pathways for oxygen in to <strong>the</strong> substrate while allowing for a 

corresponding diffusion <strong>of</strong> elements out <strong>of</strong> <strong>the</strong> substrate <strong>and</strong> coating. This results in oxidation 

along grain boundaries <strong>and</strong> at <strong>the</strong> coating substrate interfaces. In previous studies, coatings were 

146

developed with zone 2 microstructures which were more dense with fewer intergranular voids. 

Thus, <strong>the</strong>y were more resistant to grain boundary diffusion <strong>and</strong> did not exhibit oxidation at <strong>the</strong> 

interface. It is possible to eliminate oxidation by peening <strong>the</strong> surfaces to produce a compressed 

layer. The deformation will close any boundaries open to <strong>the</strong> free surface. This has been shown 

to improve <strong>the</strong> oxidation resistance in EB-PVD coatings. 

TEM analyses <strong>of</strong> <strong>the</strong> coatings revealed <strong>the</strong> presence <strong>of</strong> Al 2 O 3 within <strong>the</strong> coating grains 

for all <strong>of</strong> <strong>the</strong> samples. This suggests that oxygen is diffusing from <strong>the</strong> free surface <strong>of</strong> <strong>the</strong> coating 

<strong>and</strong> <strong>the</strong> amount <strong>of</strong> Al 2 O 3 in <strong>the</strong> coatings appeared to decrease with increasing <strong>hafnium</strong>. The 

NiAlCrHf coatings had a larger grain size after oxidation, formed larger precipitates <strong>and</strong> also 

appeared to contain smaller amounts <strong>of</strong> Al 2 O 3 within <strong>the</strong> samples. The smaller grain size <strong>of</strong> <strong>the</strong> 

NiAl-Hf coatings would lead to a dramatic increase in grain boundary volume which could 

provide a direct pathway for oxygen to diffuse from <strong>the</strong> free surface <strong>of</strong> <strong>the</strong> coating to <strong>the</strong> 

interface leading to oxidation at <strong>the</strong> coating/substrate interface. The high concentration <strong>of</strong> Hf 

observed at <strong>the</strong> coating/substrate interface with <strong>the</strong> NiAl-Hf coatings has been shown to lead to 

severe internal oxidation. This is thought to be <strong>the</strong> mechanism that leads to <strong>the</strong> oxidation at <strong>the</strong> 

coating/substrate interface reported with <strong>the</strong> NiAl-Hf coatings. 

147

CHAPTER IX 

FUTURE WORK 

While great strides were made with this <strong>investigation</strong>, <strong>the</strong>re are still several areas <strong>of</strong> 

research that should be continued. This project focused on <strong>the</strong> results <strong>of</strong> short-term iso<strong>the</strong>rmal 

oxidation experiments <strong>and</strong> <strong>the</strong> resulting microstructures. Additional work must be conducted to 

study <strong>the</strong> <strong>effect</strong>s on <strong>the</strong>se coating systems during long-term iso<strong>the</strong>rmal <strong>and</strong> cyclic oxidation. 

This may also <strong>of</strong>fer an opportunity to investigate <strong>the</strong> phases that result with increased exposure. 

An obvious piece <strong>of</strong> information missing is TEM microdiffraction patterns from precipitates in 

<strong>the</strong> coatings to accompany <strong>the</strong> APT results. This was found to be too difficult with this study 

due to <strong>the</strong> small size <strong>of</strong> <strong>the</strong> precipitates <strong>and</strong> interference with <strong>the</strong> surrounding matrix. 

The second largest impact would be to continue <strong>the</strong> efforts <strong>of</strong> investigating <strong>the</strong> <strong>effect</strong> <strong>of</strong> 

chemistry on <strong>the</strong> oxidation performance <strong>of</strong> NiAl-X coatings. This would include systems with 

<strong>the</strong> incorporation <strong>of</strong> Zr, <strong>and</strong> Zr/Hf blends along with extending this to include Cr containing 

coatings. Also, additional work is needed to investigate <strong>the</strong> <strong>effect</strong> <strong>of</strong> interstitial elements on <strong>the</strong> 

microstructures. 

148

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149

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DC Magnetron Sputtered NiAl-0.6Hf Coatings. Surface <strong>and</strong> Coatings Technology, 2005. 

200: p. 1270-1275. 

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Coatings. Surface <strong>and</strong> Coatings Technology, 2004. 177-178: p. 113-120. 

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165. Miller, M.K., K.F. Russell, <strong>and</strong> G.B. Thompson, Strategies for fabricating atom probe 

specimens with dual beam FIB. Ultramicroscopy, 2005. 102: p. 287-298. 

161

APPENDIX A 

METALLOGRAPHIC SPECIMEN PREPARATION 

Samples for this study were prepared for inspection following <strong>the</strong> guidelines 

recommended for Ni-based superalloys by Struers, Inc. The autopolisher that was used for this 

<strong>investigation</strong> was a Struers Rotopol 2 with a multidoser. The multidoser allows <strong>the</strong> user to 

automatically apply <strong>the</strong> correct amount <strong>of</strong> polishing solution <strong>and</strong> lubricant to <strong>the</strong> polishing pad. 

While grinding pads are available, SiC papers were used to grind <strong>the</strong> samples to a 4000 grit 

finish prior to final polishing. The Rotopol 2 is equipped with a variable speed control with 

options for 150 <strong>and</strong> 300 rpm. Only a setting <strong>of</strong> 150 rpm was used with this study to preserve <strong>the</strong> 

oxides that formed. Table A.1 presents <strong>the</strong> outline used for sample preparation with this study. 

Table A.1. General polishing method for Ni-based superalloys. 

Grinding/ 

Polishing Pad Grit Media Load Flow 

SiC paper 500 Water 35 N Moderate 



1 μm Diamond 

with 

6-Diamond/ 

MD-DUR 1 μm Blue Lubricant 35 N 6-Lub. 

Add with 

MD-CHEM 0.02 μm OP-AA or OP-A 35 N beaker 

Direction <strong>of</strong> Time 

Rotation (min) 

Counter- 

Clockwise 1:00 

Counter- 


Counter- 


Counter- 


Counter- 


162

APPENDIX B 

TEM SAMPLE PREPARATION USING A FOCUSED ION BEAM 

As discussed previously in <strong>the</strong> document, TEM samples were prepared using an FEI 

Company Quanta 3D dual beam-focused ion beam (FIB). This affords <strong>the</strong> user <strong>the</strong> ability to site 

specifically prepare TEM specimen while minimizing <strong>the</strong> amount <strong>of</strong> damage to <strong>the</strong> original 

sample. Fortunately, <strong>the</strong>re is a script that was developed by FEI Company. This automatic 

script is designed to prepare <strong>the</strong> specimen with <strong>the</strong> intent <strong>of</strong> being directly applied to a TEM grid 

<strong>of</strong> choice (e.g. st<strong>and</strong>ard copper grids or specialty grids depending on <strong>the</strong> application) <strong>and</strong> <strong>the</strong>n 

can be examined in <strong>the</strong> TEM without <strong>the</strong> use <strong>of</strong> any o<strong>the</strong>r instruments. This reduces <strong>the</strong> amount 

<strong>of</strong> time necessary for sample preparation from days to approximately one hour. One fault with 

this program is that it routinely crashes <strong>and</strong> requires <strong>the</strong> user to intervene <strong>and</strong> continue with <strong>the</strong> 

final milling manually. The following is an outline that I developed with <strong>the</strong> aid <strong>of</strong> Robb Morris 

<strong>and</strong> Rich Martens to accomplish this goal. 

1. Log into <strong>the</strong> instrument. 

2. Vent <strong>the</strong> chamber <strong>and</strong> align <strong>the</strong> sample height appropriately. 

3. Pump <strong>the</strong> system. 

4. Wake <strong>the</strong> system up. This should turn on <strong>the</strong> SEM <strong>and</strong> FIB columns for analysis. 

5. While <strong>the</strong> SEM/FIB columns warm-up, turn on <strong>the</strong> platinum gas injection system (GIS) <strong>and</strong> 

allow it to properly warm-up. 

163

6. After <strong>the</strong> SEM column is operational, find <strong>the</strong> eucentric location. Typically, this value is 

close to 15 mm. 

7. After <strong>the</strong> FIB column is ready, set <strong>the</strong> beam conditions to 30 kV <strong>and</strong> 0.3 nA. Use <strong>the</strong> beam 

shifts to move <strong>the</strong> image into coincidence with <strong>the</strong> SEM image. This may be difficult, but 

get it as close as possible. 

8. Tilt <strong>the</strong> sample to 52° or parallel with <strong>the</strong> FIB column. 

9. Go to <strong>the</strong> Start Menu <strong>and</strong> select <strong>the</strong> TEM autoscript function. 

10. Select <strong>the</strong> Open function <strong>and</strong> select <strong>the</strong> TEM.wsp file. 

11. Press <strong>the</strong> Play option on <strong>the</strong> window. 

12. Select to make a new sample (i.e. 15 x 4 μm rot). 

13. Follow <strong>the</strong> instructions <strong>and</strong> allow <strong>the</strong> script to run. 

14. When <strong>the</strong> program gets to a point that it tilts <strong>the</strong> stage to 7°, it will try to cut a U-shaped 

figure through <strong>the</strong> foil. This is to remove <strong>the</strong> sample later <strong>and</strong> usually where <strong>the</strong> program 

crashes. Stop <strong>the</strong> program after this cut is made by pressing <strong>the</strong> “R” symbol. 

15. Tilt <strong>the</strong> stage back to 52°. The image in <strong>the</strong> FIB is scan rotated 180° to allow <strong>the</strong> user to 

observe <strong>the</strong> sample as it is actually oriented. 

16. Cut <strong>the</strong> LEFT side <strong>of</strong> <strong>the</strong> foil free using <strong>the</strong> FIB. 

17. Tilt back to 0°. 

18. Insert <strong>the</strong> Omniprobe micromanipulator <strong>and</strong> use <strong>the</strong> SEM <strong>and</strong> FIB to move it down over <strong>the</strong> 

left-side <strong>of</strong> <strong>the</strong> sample. This should be directly over <strong>the</strong> cut you just milled in <strong>the</strong> sample. 

BE CAREFUL NOT TO RUN THE OMIPROBE INTO THE SPECIMEN! 

19. When <strong>the</strong> Omniprobe is just touching <strong>the</strong> specimen, <strong>the</strong> contrast in <strong>the</strong> FIB image will 

change. 

164

20. Using <strong>the</strong> Patterning menu, place a box on <strong>the</strong> Omniprobe <strong>and</strong> specimen for depositing 

platinum with <strong>the</strong> GIS, but be careful not to overlap with <strong>the</strong> cuts in <strong>the</strong> foil. Deposit about 

500 nm <strong>of</strong> platinum to make sure that you have good adhesion. 

21. Cut <strong>the</strong> right-side <strong>of</strong> <strong>the</strong> foil, use <strong>the</strong> beam shifts to move <strong>the</strong> sample in <strong>the</strong> FIB image. 

22. Gently, lower <strong>the</strong> stage manually (Z) to remove <strong>the</strong> foil. 

23. Retract <strong>the</strong> Omniprobe with <strong>the</strong> controls manually until it is just <strong>of</strong>f <strong>the</strong> FIB image (zoomed 

out) <strong>and</strong> <strong>the</strong>n retract it fully. 

24. Retract <strong>the</strong> GIS. 

25. Turn <strong>of</strong>f <strong>the</strong> SEM, FIB, <strong>and</strong> GIS. 

26. Vent <strong>the</strong> chamber <strong>and</strong> remove <strong>the</strong> bulk sample. 

27. Remove <strong>the</strong> sample holder used to hold Omniprobe grids. It will have a slot to place <strong>the</strong> 

grids into <strong>the</strong> holder vertically <strong>and</strong> has two hex screws to secure <strong>the</strong> grid. It is recommended 

that <strong>the</strong> three-prong Omnigrids be used. 

28. Load <strong>the</strong> holder into <strong>the</strong> system <strong>and</strong> pump <strong>the</strong> system down. 

29. Repeat <strong>the</strong> Steps 4-7 to reinitialize <strong>the</strong> system. Scan rotate <strong>the</strong> FIB image by 180°. 

30. Find <strong>the</strong> Ominprobe grid with a smooth finger-like feature. One side will be rounded while 

<strong>the</strong> o<strong>the</strong>r will have a flat edge. The flat edge is preferable to attach <strong>the</strong> sample to <strong>the</strong> 

Omniprobe grid <strong>and</strong> should be located with <strong>the</strong> FIB at 0°. 

31. When <strong>the</strong> system is at eucentric <strong>and</strong> <strong>the</strong> SEM <strong>and</strong> FIB are observing EXACTLY <strong>the</strong> same 

location, insert <strong>the</strong> Omniprobe <strong>and</strong> GIS. 

32. Use <strong>the</strong> SEM to align <strong>the</strong> foil on <strong>the</strong> proper side <strong>of</strong> <strong>the</strong> Omniprobe grid using <strong>the</strong> Omniprobe 

X-Y controls. 

165

33. Lower <strong>the</strong> sample with <strong>the</strong> Z-function on <strong>the</strong> Omniprobe controls while observing <strong>the</strong> 

progress in <strong>the</strong> FIB. It is recommended that periodic checks be made on <strong>the</strong> X-Y position in 

<strong>the</strong> SEM while lowering <strong>the</strong> sample. 

34. When <strong>the</strong> sample is at <strong>the</strong> appropriate height, move <strong>the</strong> sample using <strong>the</strong> Omniprobe X-Y 

controls until <strong>the</strong> sample is touching <strong>the</strong> Omniprobe grid. 

35. Attach <strong>the</strong> sample to <strong>the</strong> Omniprobe grid using <strong>the</strong> GIS. 

36. Cut <strong>the</strong> Omniprobe from <strong>the</strong> sample. 

37. Retract <strong>the</strong> Omniprobe manually using <strong>the</strong> Z-control. 

38. Fully retract <strong>the</strong> Omniprobe. 

39. Tilt <strong>the</strong> sample where <strong>the</strong> top is perpendicular to <strong>the</strong> FIB. 

40. Lower <strong>the</strong> beam energy to 0.1 nA with <strong>the</strong> FIB. 

41. Move <strong>the</strong> sample in <strong>the</strong> SEM, using <strong>the</strong> Z-control knob on <strong>the</strong> system, into coincidence with 

<strong>the</strong> FIB. Focus both <strong>the</strong> SEM <strong>and</strong> FIB. 

42. The sample is typically 7-8 μm long <strong>and</strong> 1 μm thick. 

43. Place a box to mill <strong>the</strong> sample at <strong>the</strong> top <strong>of</strong> <strong>the</strong> sample that is 3-5 μm long <strong>and</strong> 250 nm thick. 

Use <strong>the</strong> Integrate viewing function with <strong>the</strong> FIB image to check <strong>the</strong> amount <strong>of</strong> beam drift. 

The beam typically drifts away from <strong>the</strong> milling location. 

44. OBSERVE MILLING WITH SEM TO MAKE SURE NO BEAM DRIFT OCCURS 

DURING MILLING. THIS WILL RESULT IN SAMPLE LOSS OR NO THINNING. 

45. Repeat <strong>the</strong> same mill along <strong>the</strong> bottom <strong>of</strong> <strong>the</strong> specimen. 

46. Tilt <strong>the</strong> sample +2° <strong>and</strong> wait for <strong>the</strong> beam drift to stabilize. 

47. Reduce <strong>the</strong> milling box to 3 μm by 150 nm. 

48. Mill <strong>the</strong> top <strong>of</strong> <strong>the</strong> sample as before. 

166

49. Tilt <strong>the</strong> sample -4° <strong>and</strong> wait for <strong>the</strong> beam drift to stabilize. 

50. Mill <strong>the</strong> bottom <strong>of</strong> <strong>the</strong> sample. 

51. Tilt <strong>the</strong> sample +4° <strong>and</strong> wait for <strong>the</strong> beam drift to stabilize. 

52. Mill <strong>the</strong> top <strong>of</strong> <strong>the</strong> sample. 

53. Repeat Steps 49-52 as needed until <strong>the</strong> sample is ~100 nm or less. 

54. MAKE SURE TO CHECK THE MILLING PROGRESS USING THE SEM! 

55. Reduce <strong>the</strong> FIB beam energy to 5.0 kV <strong>and</strong> 49 nA. 

56. Tilt to 0°. 

57. Image <strong>the</strong> sample with <strong>the</strong> FIB for 30-45 seconds to remove beam damage. 

58. Rotate <strong>the</strong> sample compeucentrically <strong>and</strong> image <strong>the</strong> o<strong>the</strong>r side <strong>of</strong> <strong>the</strong> sample for 30-45 

seconds. 

59. Turn <strong>of</strong>f <strong>the</strong> HT to <strong>the</strong> SEM <strong>and</strong> FIB. 

60. Turn <strong>of</strong>f <strong>the</strong> GIS. 

61. Vent <strong>the</strong> system. 

62. Remove <strong>the</strong> sample <strong>and</strong> place it in a holder, noting its location carefully. 

63. Pump <strong>the</strong> chamber. 

64. Home <strong>the</strong> stage. 

65. Replace all tools <strong>and</strong> holders used in <strong>the</strong>ir proper storage locations. 

66. Clean up any gloves, tissue paper, etc. 

67. Log out <strong>of</strong> <strong>the</strong> system. 

167

APPENDIX C 

APT SAMPLE PREPARATION USING A FOCUSED ION BEAM 

As discussed previously in <strong>the</strong> document, TEM samples were prepared using an FEI 

Company Quanta 3D dual beam-focused ion beam (FIB). This technique can be used in lieu <strong>of</strong> 

traditional methods like electropolishing. Similar to <strong>the</strong> TEM sample preparation, APT samples 

can be selected site specifically. An outline for <strong>the</strong> procedures used for this portion <strong>of</strong> <strong>the</strong> 

<strong>investigation</strong> was adapted from those provided by Rich Martens <strong>and</strong> are presented below. 

1. Log into <strong>the</strong> instrument. 

2. Vent <strong>the</strong> chamber <strong>and</strong> align <strong>the</strong> sample height appropriately. 


4. Wake <strong>the</strong> system up. This should turn on <strong>the</strong> SEM <strong>and</strong> FIB columns for analysis. 

5. While <strong>the</strong> SEM/FIB columns warm-up, turn on <strong>the</strong> platinum gas injection system (GIS) <strong>and</strong> 

allow it to properly warm-up. 

6. After <strong>the</strong> SEM column is operational, find <strong>the</strong> eucentric location. Typically, this value is 

close to 15 mm. 

7. After <strong>the</strong> FIB column is ready, set <strong>the</strong> beam conditions to 30 kV <strong>and</strong> 0.1 nA. Use <strong>the</strong> beam 

shifts to move <strong>the</strong> image into coincidence with <strong>the</strong> SEM image. This may be difficult, but 

get it as close as possible. 

8. Scan rotate <strong>the</strong> FIB image by 180°. 

168

9. Find a location to retrieve an APT sample. 

10. Tilt to 52° <strong>and</strong> insert <strong>the</strong> GIS. 

11. Deposit a layer <strong>of</strong> platinum that is 22 μm x 2 μm x 500 nm. This will help preserve <strong>the</strong> APT 

tips while milling. 


13. Tilt to 22°. The intent is to make an isosceles triangle with <strong>the</strong> sample. 

14. Adjust <strong>the</strong> beam conditions to 30 kV <strong>and</strong> 3.0 nA <strong>and</strong> change <strong>the</strong> depth for milling to 4 μm. 

15. Focus <strong>the</strong> beam on an area that is not on your sample. This is easily accomplished using <strong>the</strong> 

small adjustable viewing window. Be careful not to image <strong>the</strong> sample with <strong>the</strong> FIB as this 

will mill <strong>the</strong> sample away very quickly. 

16. Place <strong>the</strong> milling box at <strong>the</strong> bottom <strong>of</strong> <strong>the</strong> platinum deposit. 

17. Start <strong>the</strong> milling <strong>and</strong> watch <strong>the</strong> live time monitor for beam drift. This should take a little 

more than six minutes. If <strong>the</strong> time is not right, <strong>the</strong> beam current is not correct. 

18. Compeucentrically rotate <strong>the</strong> sample 180° to mill <strong>the</strong> o<strong>the</strong>r side. 

19. Check for beam drift <strong>and</strong> focus as in Step 15. 

20. Mill <strong>the</strong> bottom <strong>of</strong> <strong>the</strong> sample to complete <strong>the</strong> triangle. 

21. Compeucentrically rotate <strong>the</strong> sample 180°. 

22. Check for beam drift <strong>and</strong> refocus. 

23. Change <strong>the</strong> milling depth to 2 μm. 

24. Tilt to 52°. 

25. Reduce <strong>the</strong> FIB beam current to 0.1 nA <strong>and</strong> refocus. 

26. Mill <strong>the</strong> left side <strong>of</strong> <strong>the</strong> sample similar to <strong>the</strong> TEM sample preparation in Appendix B. 

27. Tilt to 0°. 

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28. Insert <strong>the</strong> GIS <strong>and</strong> Omniprobe. 

29. Insert <strong>the</strong> Omniprobe micromanipulator <strong>and</strong> use <strong>the</strong> SEM <strong>and</strong> FIB to move it down over <strong>the</strong> 

left-side <strong>of</strong> <strong>the</strong> sample. This should be directly over <strong>the</strong> cut you just milled in <strong>the</strong> sample. 

BE CAREFUL NOT TO RUN THE OMIPROBE INTO THE SPECIMEN! 

30. When <strong>the</strong> Omniprobe is just touching <strong>the</strong> specimen, <strong>the</strong> contrast in <strong>the</strong> FIB image will 

change. 

31. Using <strong>the</strong> Patterning menu, place a box on <strong>the</strong> Omniprobe <strong>and</strong> specimen for depositing 

platinum with <strong>the</strong> GIS, but be careful not to overlap with <strong>the</strong> cuts in <strong>the</strong> foil. Deposit about 

500 nm <strong>of</strong> platinum to make sure that you have good adhesion. 

32. Cut <strong>the</strong> right-side <strong>of</strong> <strong>the</strong> foil, use <strong>the</strong> beam shifts to move <strong>the</strong> sample in <strong>the</strong> FIB image. 

33. Gently, lower <strong>the</strong> stage manually (Z) to remove <strong>the</strong> foil. 

34. Retract <strong>the</strong> Omniprobe with <strong>the</strong> controls manually until it is just <strong>of</strong>f <strong>the</strong> FIB image (zoomed 

out) <strong>and</strong> <strong>the</strong>n retract it fully. 


36. Turn <strong>of</strong>f <strong>the</strong> SEM, FIB, <strong>and</strong> GIS. 

37. Vent <strong>the</strong> chamber <strong>and</strong> remove <strong>the</strong> bulk sample. 

*The next step will require <strong>the</strong> use <strong>of</strong> a sample coupon suitable for use in <strong>the</strong> LEAP. Check 

with Rich Martens if you are not sure what this means. 

38. Load <strong>the</strong> sample holder with <strong>the</strong> LEAP coupon. 

39. Adjust <strong>the</strong> height <strong>of</strong> <strong>the</strong> sample to <strong>the</strong> appropriate level. 


41. Repeat <strong>the</strong> Steps 4-7 to reinitialize <strong>the</strong> system. Scan rotate <strong>the</strong> FIB image by 180°. 

42. Find <strong>the</strong> sample location on <strong>the</strong> coupon with <strong>the</strong> FIB at 0°. 

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43. When <strong>the</strong> system is at eucentric <strong>and</strong> <strong>the</strong> SEM <strong>and</strong> FIB are observing EXACTLY <strong>the</strong> same 

location, insert <strong>the</strong> Omniprobe <strong>and</strong> GIS. 

44. Use <strong>the</strong> SEM to align <strong>the</strong> foil on <strong>the</strong> proper side <strong>of</strong> <strong>the</strong> Omniprobe grid using <strong>the</strong> Omniprobe 

X-Y controls. 

45. Lower <strong>the</strong> sample with <strong>the</strong> Z-function on <strong>the</strong> Omniprobe controls while observing <strong>the</strong> 

progress in <strong>the</strong> FIB. It is recommended that periodic checks be made on <strong>the</strong> X-Y position in 

<strong>the</strong> SEM while lowering <strong>the</strong> sample. 

46. When <strong>the</strong> sample is on <strong>the</strong> coupon post, use <strong>the</strong> GIS to deposit 500 nm <strong>of</strong> platinum to attach 

<strong>the</strong> sample to <strong>the</strong> post. 

47. Cut <strong>the</strong> sample on <strong>the</strong> post free from <strong>the</strong> remaining sample on <strong>the</strong> Omniprobe. Make sure to 

conserve as much <strong>of</strong> <strong>the</strong> sample on <strong>the</strong> Omniprobe as possible. 

48. Move <strong>the</strong> Omniprobe away from <strong>the</strong> post slowly <strong>and</strong> <strong>the</strong>n raise it slightly so as to not hit <strong>the</strong> 

coupon when moving to <strong>the</strong> next sample. 

49. Locate <strong>the</strong> next post <strong>and</strong> move <strong>the</strong> stage manually to align <strong>the</strong> sample. 

50. Repeat Steps 45-48 until <strong>the</strong> sample is consumed or posts are full. Each sample should make 

exactly six posts. 

51. Retract <strong>the</strong> Omniprobe manually until <strong>of</strong>f <strong>the</strong> FIB screen <strong>and</strong> <strong>the</strong>n retract fully. 


53. Rotate <strong>the</strong> sample 180° to deposit platinum on <strong>the</strong> opposite side. 

54. Deposit 500 nm <strong>of</strong> platinum to attach <strong>the</strong> sample to <strong>the</strong> post <strong>and</strong> repeat for <strong>the</strong> remaining 

posts. 

55. Tilt to 52°. 

56. Move to <strong>the</strong> first post to be milled. 

171

57. Select <strong>the</strong> circle milling object. 

58. Make <strong>the</strong> circle 3.5-4 μm. 

59. Drag <strong>the</strong> center <strong>of</strong> <strong>the</strong> circle to make an annulus that touches as many <strong>of</strong> <strong>the</strong> sides <strong>of</strong> <strong>the</strong> 

sample as possible. 

60. Adjust <strong>the</strong> milling depth to be 1 μm <strong>and</strong> start <strong>the</strong> milling. 

61. Watch <strong>the</strong> progress in <strong>the</strong> SEM. This will require <strong>the</strong> use <strong>of</strong> <strong>the</strong> auto brightness <strong>and</strong> contrast 

feature. 

62. When <strong>the</strong> post appears to be a cylinder from <strong>the</strong> silicon through <strong>the</strong> coating, stop <strong>the</strong> milling 

program. 

63. Reduce <strong>the</strong> inner circle by half <strong>and</strong> move <strong>the</strong> out circle in by 1 μm. 

64. Watch <strong>the</strong> progress in <strong>the</strong> SEM again until <strong>the</strong> tip begins to become conical in nature. 

65. Repeat Step 63. 

66. Continue reducing <strong>the</strong> diameter until <strong>the</strong> tip diameter is >100 nm. 

67. Repeat Steps 56-66 for all posts. 

68. Reduce <strong>the</strong> beam conditions to 5.0 kV <strong>and</strong> image each tip for 30-45 seconds. 

69. Tilt back to 0°. 

70. Turn <strong>of</strong>f <strong>the</strong> HT to <strong>the</strong> SEM <strong>and</strong> FIB. 

71. Vent <strong>the</strong> system. 

72. Pump <strong>the</strong> chamber. 

73. Home <strong>the</strong> stage. 

74. Replace all tools <strong>and</strong> holders used in <strong>the</strong>ir proper storage locations. 

75. Clean up any gloves, tissue paper, etc. 

76. Log out <strong>of</strong> <strong>the</strong> system. 

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