High-Temperature Mechanical Behavior of End-of-LifeCryomilled NiCrAlY Bond Coat Materials

M. FUNK, K. MA, C. EBERL, J.M. SCHOENUNG, M. GOKEN, and K.J. HEMKER

Previous work has demonstrated that the lifetime of atmospheric plasma sprayed thermalbarrier coating (TBC) systems incorporating cryomilled NiCrAlY bond coats show superiorreliability with up to 3 times longer lifetimes compared to conventional ones. These conven-tional and cryomilled NiCrAlY bond coats at end-of-life (after thermal cycle failure) werestudied in detail in the present work with a focus on the mechanical behavior in the temperaturerange from room temperature to 1273 K (1000 �C). The investigations were carried out usingmicrotensile samples and the DIC technique. It turns out that the low-temperature strength ofthe cryomilled NiCrAlY bond coat is inferior to that of conventional ones, which might be dueto a more pronounced porosity. At higher temperatures (between 1173 K and 1273 K (900 �Cand 1000 �C)), the cryomilled bond coat shows almost twice the strength of the conventionalbond coat, despite having been exposed to almost 3 times as many thermal cycles. The thermalstability of the nitride dispersoids appears to compensate for the gamma prime dissolution thattypically occurs at these elevated temperatures, allowing for strength retention.

DOI: 10.1007/s11661-011-0659-2� The Minerals, Metals & Materials Society and ASM International 2011

I. INTRODUCTION

HIGH performance thermal barrier coating (TBC)systems are often used to increase gas turbine effi-ciency.[1,2] TBCs typically consist of three layers: aceramic top coat made of yttria stabilized zirconia(YSZ), a bond coat often made of MCrAlY (whereM = Ni, Co, Fe), and a superalloy substrate. A fourthlayer, called the thermally grown oxide (TGO), formsbetween the YSZ and the bond coat as a result ofoxidation. In gas turbines, the TBCmaterials are exposedto thermal cycling between room temperature and up to1473 K (1200 �C).[1] Due to the difference in thermalexpansion coefficients relative to the substrate material,each layer of the TBC undergoes cyclic stress whenexposed to service conditions. Furthermore, the TGOlayer imposes compressive growth stresses in addition tothe cyclically induced stresses. These stresses eventuallylead to plastic deformation at low temperatures and creepdeformation at high temperatures and ultimately result inthe spallation of the ceramic top coat. Several modelshave been developed to describe the stress distributionsand the failure mechanisms in select TBC systems.[3,4]

Critical key components are the creep properties aswell as the resistance to oxidation of the bond coat athigh temperature.[5] Commercially used bond coats canbe either diffusion aluminide or thermal spray overlaybond coats.[2] The most common diffusion aluminidebond coats contain a small amount of Pt, which iselectrodeposited prior to aluminization. These bondcoats have superior mechanical properties at hightemperatures and strong oxidation resistance comparedto MCrAlY bond coats.[6] Unfortunately, the lifetime ofTBCs containing Pt modified bond coats can be limiteddue to the high-temperature martensitic phase transfor-mation developing during service or thermal cycling.[7]

The large strain cycles can lead to the formation ofinterface undulations, and as a result, the TBC canundergo spallation. MCrAlY overlay bond coats can beapplied to complicated geometries such as airfoils, e.g.,by air plasma spraying (APS), vacuum plasma spraying,and low pressure plasma spraying (LPPS).[1,2] MCrAlYcoatings are chemically and thermomechanically com-patible to the superalloy substrates and form a protec-tive alumina layer at the top coat interface.[1,2] Themicrostructure is mostly based on two phases: theb-(NiAl) and the c-phase (Ni).[1,2] For some particularcompositions, a-Cr and c¢-(Ni3Al) precipitates are alsopresent in the b-(NiAl) and c-Ni grains, respectively.[8]

The c¢-precipitates in MCrAlY alloys provide enhancedmechanical properties at temperatures below the solvustemperature, which is usually between 1073 and 1253 K(800 and 980 �C).[8–10] The strengthening effect, resultingfrom the c¢-precipitates, disappears at high temperaturesdue to the dissolution of the c¢. To improve themechanical properties at even higher temperatures,dispersoids can be added.[11] Schoenung and co-workersdemonstrated that cryomilled NiCrAlY powder can be

M. FUNK, PhD Student, and C. EBERL, Group Leader, are withthe Institute for Reliability of Components and Systems (izbs),Karlsruhe Institute of Technology, Karlsruhe 76131, Germany.Contact e-mail: matthias.funk@kit.edu K. MA, PhD Student, andJ.M. SCHOENUNG, Professor, are with the Department of ChemicalEngineering and Materials Science, University of California, Davis,Davis, CA 95616. M. GOKEN, Professor, is with the MaterialsScience and Engineering Department, University Erlangen–Nurnberg,Erlangen 91058, Germany. K.J. HEMKER, Professor, is with theDepartment of Mechanical Engineering, Johns Hopkins University,Baltimore, MD 21218.

Manuscript submitted April 6, 2010.Article published online March 23, 2011

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 42A, AUGUST 2011—2233

used in a spray coating process.[12] Thermal cycle tests ofTBC systems with cryomilled NiCrAlY bond coatsresulted in up to 3 times longer lifetime than theconventional equivalents.[13] Cryomilled NiCrAlY bondcoats can have significantly improved mechanical sta-bility at high temperatures, i.e., creep resistance, whencompared to conventional materials,[14] which wasattributed to the presence of ultrafine dispersoids.

To further explore the effects of thermal cycling, thepurpose of the current work is to investigate the mechan-ical behavior of cryomilled bond coats in the end-of-lifestate.[13] For comparison purposes, the conventional end-of-life bond coats are also included. The techniques usedfor this characterization were developed by Hemker andco-workers.[15,16] The microstructures of the as-sprayedand the end-of-life bond coats were examined in an effortto provide structure-property derived insights.

II. EXPERIMENTAL PROCEDURE

The samples used in this study were the end-of-life(thermally cycled to failure) TBCs used in previouswork.[16] An abbreviated description of how these wereprepared is provided here. Three different bond coatmaterials (Table I) were thermal sprayed onto 25-mm-diameter HASTELLOY* X button-shaped substrates

using the LPPS process. For the conventional sample(LCVH), Ni-22Cr-10Al-1Y powder was used as thefeedstock powder. The same powder was cryomilled for8 hours at 180 rpm and for 6 hours at 360 rpm to makethe LM1H and LM2H powders, respectively. A 7-YSZtop coat was applied by APS. The TBC specimens wereheat treated in vacuum for 4 hours at 1353 K (1080 �C)to densify the bond coats. The lifetime criterion wasdefined as the cycle after which the delamination of theAPS 7-YSZ top coat reached 50 pct of the button area.After reaching the lifetime criterion, the completedelamination would generally occur within approxi-mately 24 more thermal cycles.[16]

Micrographs of the bond coat samples, bothas-sprayed and at end-of-life, were acquired with a

scanning electron microscope (SEM) (JEOL,** JSM-

6400) at a voltage of 20 kV and a working distance of 18mm. The current was reduced to 10�13 A to prevent thecharging of the oxide phases. The ion contrast pictureswere acquired in a Crossbeam 1540 EsB (Zeiss, Oberko-chen, Germany). The element distribution in the end-of-life bond coat materials was analyzed with twoenergy-dispersive X-ray (EDX) systems. One was usedat a voltage of 20 kV and a working distance of 6 mm(INCA Penta FETx3, Oxford Instruments), and thesecond one was used at a voltage of 10 kV and aworking distance of 11 mm (Gemini Leo 1530) for thehigh-magnification images. A NOVA Nanolab 200(FEI, Hilsboro, OR) was used to prepare cross sectionsat a voltage of 30 kV and a working distance of19.6 mm.Dog-bone-shaped microsamples were prepared from

the end-of-life samples, as shown in Figure 1. Thin sliceswere cut off from the surface of the button samples and adog-bone-shaped electrode was used in a sink EDMmachine to punch out the sample geometry. The super-alloy substrate was removed by grinding and polishing.This polishing procedure was conducted in steps and anoptical microscope was used to detect the interfacebetween the bond coat and substrate. The other side waspolished to remove the top coat and oxide residue at thesurface. The sample size, thickness, and width weredetermined by optical microscopy. The mean gage widthwas measured from both sides as the sink EDM processleads to a slight tapering through the sample thickness.The individual sample thickness varied from 60 to 90 lm,while the width varied from 475 to 520 lm. The thicknessvariation along the tensile axis was ±2 lm for eachsample. The gage length was 1700 lm.Vickers hardness measurements were carried out at

loads of only 10 g to minimize the indentation depth dueto the limited sample thickness. Each side of thepolished dog-bone samples was indented in the earsection at four different sites. The size of the diagonalfor each Vickers indent was measured optically and usedto calculate the hardness. The tensile tests were carriedout with two different custom-built microtensile testingsetups, developed at Johns Hopkins University. Similar

Table I. Processing Parameters for the Different Feedstock Powders

Sample Label LCVH LM1H LM2H

Powder conventional cryomilled cryomilledAverage particle size (lm) 29 64 59Milling parameters — 8 h, 180 rpm 6 h, 360 rpmSpray process LPPS LPPS LPPSVacuum heat treat 4 h, 1353 K (1080 �C) 4 h, 1353 K (1080 �C) 4 h, 1353 K (1080 �C)Lifetime in cycles 420 1074 1285EOL average grain size (lm) 7.1 4.9 3.8

EOL: end-of-life; LPPS: low pressure plasma spray.

*HASTELLOY is a trademark of Haynes International Inc.,Kokomo, IN.

**JEOL is a trademark of Japan Electron Optics Ltd., Tokyo.

2234—VOLUME 42A, AUGUST 2011 METALLURGICAL AND MATERIALS TRANSACTIONS A

setups are described in References 7 and 15 and arebased on MEMS materials testing techniques.[17] Theroom-temperature tensile testing setup consisted of a20 lb load cell, an air bearing to reduce friction, and alinear motor to apply a displacement. The load train washorizontally assembled. The strain was measured by theuse of a digital image correlation (DIC) technique, alsodeveloped at the Johns Hopkins University and pre-sented in Reference 16. The Matlab (Natick, MA)functions are available for free.[18] The technical princi-ple is based on the analysis of digital image sequencescaptured with a digital camera (PL782, PixeLINK,Ottawa, Canada) during the test. The subsequentdeformation of the sample can be tracked in a postpro-cess and analyzed for local as well as global deforma-tion. A fiber optic cold light source was used toilluminate the samples during the test. The sampleswere glued into the crosshead grips and fixed withscrews. A small tensile load was applied for 1 hour whilethe glue cured.

The high-temperature setup consisted of a 5 lb loadcell, a piezo-stepper motor, a furnace, and the samecamera as that used in the room-temperature setup. Thesetup was vertically mounted with a post reaching into acylindrical furnace where the sample was mounted. Athick Ni wire was used to transfer the load from the loadcell, which was mounted to the motor on the outside ofthe furnace. For temperatures up to 1013 K (740 �C),the cylindrical furnace was equipped with two glass rodsused as light guides through the furnace insulation.Another glass rod was used as the camera window. Forexperiments at temperatures from 1073 K to 1373 K(800 �C to 1100 �C), a second furnace was used, whichwas not equipped with light guides and the camerawindow necessary for a DIC based strain measurement.Therefore, the strain at these high temperatures had tobe calculated from the crosshead displacement, whichwas measured by a capacitance gage. MgO powder wasused to protect the Ni components from oxidation athigh temperatures. A constant tensile preload wasapplied while heating from room temperature to thedesired temperature. This was necessary to keep thesamples properly attached to the wedgelike grips.The temperature variation for the low-temperaturefurnace was less than ±0.5 K throughout the test, whilethe high-temperature furnace showed temperature fluc-tuations with an amplitude of ±1.75 K at 973 K(700 �C), ±2.25 K at 1173 K (900 �C), and ±2.5 K at1273 K (1000 �C).

III. RESULTS AND DISCUSSION

A. Microstructure

The conventional and cryomilled bond coats showmajor differences in their microstructures before thethermal cycling and in their end-of-life condition. In theas-sprayed bond coats, as shown in Figure 2, disper-soids (the dark regions) were observed in the cryomilledsamples (LM1H and LM2H) but not in the conven-tional sample (LCVH). The bright and gray regionsvisible in all coatings are the c-Ni(Cr) grains, some ofwhich contain c¢-Ni3Al precipitates, and the b-NiAlphase, respectively. One of the key advantages ofcryomilling is that the average grain size of the feedstockpowder is decreased to the nanoscale.[19] Furthermore,the formation of the dispersoids contributes to boththermal stability and the potential for higher strength inthe cryomilled bond coat.[14,19] The microstructures ofthe conventional and cryomilled bond coat also differdistinctly from each other in their end-of-life samples, asdemonstrated in Figure 3. While the conventional sam-ple LCVH contains small pores close to triple points, thecryomilled samples contain larger pores. Although thecoatings are essentially dense by thermal spray stan-dards, pores are observed at various locations through-out the bond coat (Figure 4). It is unclear whether thesepores developed during the LPPS process or duringthermal cycling. The pores located in the center of thebond coat have a size of tens of microns, and it seems asif the inner surface is not covered by an oxide layer. Thisimplies that the pores in that area are not connected tothe surface. Pores close to the TGO seem to beconnected to the atmosphere as they are covered withan oxide skin and often are nucleation sites for oxideprotrusions. As shown in Table I, the average grain sizein the end-of-life cryomilled samples (LM2H: 3.8 lmand LM1H: 4.9 lm) was measured to be approximately

Fig. 1—Schematic of the TBC system and the dog-bone-shaped mi-crotensile samples prepared from the bond coat.

Fig. 2—SEM micrographs of as-sprayed NiCrAlY bond coat (beforethermal cycling)[13] (a) LCVH (b) LM1H (c) LM2H. On the rightside, the bright areas are c/c¢, the gray areas are b-NiAl and thedark areas are Al-rich dispersoids.

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 42A, AUGUST 2011—2235

half the average grain size in the end-of-life conventionalsamples (LCVH: 7.1 lm).

As described in References 19 and 20, extremely finenitride particles can be formed in the feedstock powderduring the cryomilling process; their presence wouldenhance the high-temperature strength in cryomilledbond coats. A representative EDX analysis (Figure 5(a))indicates that particles, rich in Al and N, with a size ofup to several microns exist in the thermally cycledcryomilled bond coat, both at the interdiffusion zonebetween the bond coat and the substrate and within thebond coat.

Figure 5(b) allows a closer look at the selected areamarked in Figure 5(a), at the limit of the interdiffusionzone in the cryomilled sample LM2H. Here, oxideparticles with sizes on the order of hundreds ofnanometers with a high concentration of Al or Y areobserved. Larger nitride particles are also observed, withsizes on the order of several hundreds of nanometers toa few microns. These larger particles consist primarily ofAl and N. It is observed that the area density of thesenitride particles is higher in the interdiffusion zone, butdecreases steadily in the direction of the top coatinterface.

In contrast, the Al-rich dispersoids in the interdiffu-sion zone in the conventional LCVH sample are muchmore locally defined and do not reach into the bond coat(Figure 5(c)). No nitride particles are observed at thelimit of the interdiffusion zone in sample LCVH(Figure 5(d)).

Comparing the oxygen-rich dispersoids, it is notedthat the Al-Y oxides in the cryomilled LM2H are mostlyless than 1 lm in diameter (Figure 5(b)), while Y oxidesin the conventional LCVH can reach diameters ofseveral microns (Figure 5(d)). It remains unclear whythe particles are smaller in the cryomilled samples. Thereason could be attributed to the severe plastic defor-mation during ball milling in the cryogenic environment,but also the presence of Al in these dispersoids couldhave a grain-refining effect. Further investigation isongoing to clarify the mechanism for dispersoid refine-ment, yet regardless of the mechanism, the fine disper-soids can play an important role in the enhancement ofhigh-temperature strength, as described subsequently.

For the microtensile test, the samples are taken fromthe center region of the bond coat, so it is important toalso look closely at this center region. In the center of

the conventional bond coat in sample LCVH, someY-O–rich regions are observed, as illustrated inFigure 5(e). Throughout this center region, N is also notdetected. In contrast, dispersoid particles are observedin the center region of the cryomilled bond coat insample LM2H (Figure 5(f)), which consist of Al-N(these appear black in the SEM image) and others thatconsist of Al-Y-O (these appear gray in the SEM image).Thus, after extensive thermal cycling exposure, the end-of-life cryomilled samples contain Al-N–rich and Al-Y-O–rich dispersoids throughout the bond coat thickness,whereas only Y-O–rich dispersoids are observed in thecenter of the conventional material. Furthermore, thesize of the disperoids in the cryomilled material is muchfiner than that of the dispersoids in the conventionalmaterial.

B. Mechanical Properties

Hardness measurements provide a first impression ofthe room-temperature strength of the bond coat. Themean Vickers hardness values, listed in Table II, indi-cate that the conventional bond coat shows a slightlyhigher room-temperature hardness than the cryomilledmaterial. The mean hardness for the conventionalmaterial is 190 HV, which is slightly higher than thevalues for the cryomilled LM1H, 184 HV and LM2H,182 HV. Apparently, in these materials, the cryomillinghas not led to a detectable increase in room-temperaturehardness when compared to the conventional material.Similar findings resulted from the microtensile tests

that were conducted at room temperature (Table II andFigure 6). The conventional material exhibits relativelyhigh yield strength of approximately 550 MPa. Further-more, the conventional bond coat shows a strong strainhardening behavior, which leads, in combination withthe high ductility, to an ultimate strength of 800 MPa. Itshould be noted at this point that no necking wasobserved. In contrast, the cryomilled materials show acomparatively lower yield strength of approximately400 MPa. Additionally, the low ductility leads to aneven larger difference in their ultimate strength of 420(LM1H) and 530 MPa (LM2H), even though theobserved hardening (LM2H) seems to be similar to thatobserved in the conventional bond coat. A similarstrength value was recently measured for commercialNiCoCrAlY bond coat material after thermal cycling.[21]

Fig. 3—Microstructures of end-of-life NiCrAlY bond coats: (a) conventional (LCVH); and cryomilled (b) LM1H and (c) LM2H.

2236—VOLUME 42A, AUGUST 2011 METALLURGICAL AND MATERIALS TRANSACTIONS A

Fig. 4—(a) SEM image of the bond coat for a cryomilled NiCrAlY sample (LM2H) thermally cycled to end-of-life (1285 cycles) and the corre-sponding EDX images for Al, Ni, N, and O. The box in the SEM image indicates the position in the bond coat, where the EDX mapping athigher magnification was made ((c)). (b) SEM image of a region in the bond coat for a conventional NiCrAlY sample (LCVH) thermally cycledto end-of-life (420 cycles) and the corresponding EDX images for Al, Ni, and N. (c) SEM image of a region in the bond coat (marked in (a)) fora cryomilled NiCrAlY sample (LM2H) thermally cycled to end-of-life (1285 cycles) and the corresponding EDX images for Al, Ni, Y, N, and Oat higher magnification. (d) SEM image of a region in the bond coat for a conventional NiCrAlY sample (LCVH) thermally cycled to end-of-life(420 cycles) and the corresponding EDX images for Al, Ni, Y, N, and O at higher magnification. (e) SEM image of a region in the center of thebond coat for a conventional NiCrAlY sample (LCVH) thermally cycled to end-of-life (420 cycles) and the corresponding EDX images for Al,Ni, Y, C, N, and O. (f) SEM image of a region in the center of the bond coat for a cryomilled NiCrAlY sample (LM2H) thermally cycled toend-of-life (1285 cycles) and the corresponding EDX images for Al, Ni, Y, C, N, and O at higher magnification.

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 42A, AUGUST 2011—2237

The yield strength of this NiCoCrAlY bond coatexhibited a considerable drop from 1200 MPa for theas-deposited coating to 660 MPa for the thermal cycledsamples. The major difference from the conventionalbond coat presented in this article is the great differencein total elongation, as the commercial thermally cycledbond coat in Reference 21 showed only up to 1.5 pcttotal strain, while the conventional NiCrAlY shownhere reached almost 20 pct (Figure 7). In contrast, thecryomilled bond coat showed only up to 3 pct totalelongation, which, it should be noted, is still more thandouble that of the commercial NiCoCrAlY bond coat.

The fact that the room-temperature ductility of thethermally cycled cryomilled bond coat is lower than thatof the conventional material may be related to theformation of dispersoids and the introduction of impu-rities during cryomilling, both of which constrain thedislocation mobility.[13] However, the Vickers hardnessmeasurement did not show noticeable differences asthere was, for example, no crack formation observedduring indentation and the difference in hardnessbetween the conventional and cryomilled bond coatwas only 3 pct. Therefore, the low ductility of thethermally cycled samples of at most 3 pct (Table II,

Fig. 4—Continued.

2238—VOLUME 42A, AUGUST 2011 METALLURGICAL AND MATERIALS TRANSACTIONS A

Figure 6) could also be attributed to the porosityobserved in the cross sections. This porosity couldresult from the longer lifetime (and larger number ofthermal cycles) experienced by the cryomilled TBCsystem, leading to a longer thermal exposure and a morepronounced internal oxidation of the pores. As the bondcoat layer is rather thin, the formation of these pores canreduce the strength of the material.

For temperatures below 973 K (700 �C), the conven-tional bond coat (LCVH) shows a considerably higherstrength than the cryomilled bond coat (LM1H andLM2H) (Figure 7). At 1173 K (900 �C) and higher, the

strength of the conventional material decreases signifi-cantly, eventually reaching a value less than the strengthof the cryomilled bond coat. A similar decrease instrength was observed for thermally cycled commercialNiCoCrAlY samples at temperatures between 923 Kand 1073 K (650 �C and 800 �C), which is attributed tothe dissolution of c¢-precipitates in the c grains thatoccurs in this temperature regime.[21] The drop instrength for the conventional bond coat observed inthis study is believed to also result from the c¢dissolution owing to the phase transformation thatoccurs in NiCrAlY alloys.[8] The cryomilled bond coats

Fig. 4—Continued.

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 42A, AUGUST 2011—2239

(LM1H and LM2H) also undergo a decrease in strengthfrom 973 K to 1173 K (700 �C to 900 �C), however,much less dramatically than the conventional bondcoats. The dissolution of c¢-precipitates occurs in thecryomilled NiCrAlY bond coat as well due to a phasetransformation studied in Reference 8. Nevertheless, thepresence of the ultrafine oxide or nitride dispersoids,which have been carefully characterized in Section A,and their thermal stability compensate for the decreasein strength caused by the c¢ dissolution, allowing forstrength retention at elevated temperatures. Conse-quently, a superior strength of the cryomilled materialcompared to the conventional material is observed attemperatures of 1173 K (900 �C) and higher, evenstronger than the as-deposited Pt modified bond coat.[6]

Samples from Reference 14 also demonstrate thesupremacy of the cryomilled bond coat materials incomparison with the conventional bond coats at hightemperatures. It should be noted, however, that thestrength was not determined with microsamples.

Furthermore, the cryomilled material shows alsosuperior properties in terms of the total strain reachedduring high-temperature tensile deformation. The cryo-milled materials gain in total strain with increasingtemperature, while the conventional material againexperiences a strong degradation. The dispersoids in thecryomilled material seem to prevent localization in termsof necking and, therefore, allow for a homogeneous

deformation behavior. This contributes to the highertotal strains compared to the conventional material, asshown in Figure 7.

IV. SUMMARY

The lifetime of APS TBC systems seems to be stronglycorrelated to the high-temperature strength and ductilityof the bond coat in use. The previously reportedthreefold increase in lifetime for TBCs with cryomilledbond coats[13] can be explained by the outstanding high-temperature strength. While commercial NiCoCrAlYand conventional NiCrAlY show a dramatic drop instrength at temperatures exceeding 973 K to 1073 K(700 �C to 800 �C) as c¢-precipitates dissolve into the cmatrix, the cryomilled coatings exhibit strength reten-tion at elevated temperatures owing to the compensationfrom the ultrafine dispersoids, eventually showing supe-

Fig. 5—Stress-strain curves of NiCrAlY bond coat microtensile sam-ples, showing the relatively high ductility of the conventional bondcoat (LCVH) and the substrate material compared to the cryomilledbond coats (LM1H with brittle fracture and LM2H) at room tem-perature.

Table II. Room-Temperature Mechanical Properties for theDifferent Bond Coat Materials

Sample Label LCVH LM1H LM2H

Hardness (HV10) 190 ± 18 184 ± 12 182 ± 12Young’s modulus(GPa)

120 145 135

Yield strength(0.2 pct) (MPa)

550 (420) brittlefracture

450

Ultimate tensilestrength (MPa)

800 420 530

Total strain (pct) 19 — 3

0 200 400 600 800 1000 1200 1400 16000

200

400

600

800

1000

1200

1400 LM1H LM2H cyc. cryo. NiCrAlY [14] LCVH cyc. conv. NiCoCrAlY conv. NiCoCrAlY cyc. conv. NiCrAlY Pt-mod. [6]

ultim

ate

stre

ngth

[MP

a]

temperature [K]

brittlefracture

[pc]

800 1000 1200 1400 16000

200

400

600

LM1H

(a)

(b)

LM2H cyc. cryo. NiCrAlY [14] LCVH cyc. conv. NiCoCrAlY conv. NiCoCrAlY cyc. conv. NiCrAlY Pt-mod. [6]ul

timat

e st

reng

th [M

Pa]

temperature [K]

Fig. 6—Comparison of strength values for conventional and cryo-milled NiCrAlY and NiCoCrAlY bond coats at elevated tempera-tures. Strength of Pt-modified bond coat materials is presented forcomparison; the properties of conventional NiCoCrAlY in as-re-ceived ‘‘conv. NiCoCrAlY’’ and cycled state ‘‘cyc. conv. NiCoCrAl-Y’’ were measured at the JHU.

2240—VOLUME 42A, AUGUST 2011 METALLURGICAL AND MATERIALS TRANSACTIONS A

rior strength to the conventional bond coat. Even withmore than 1000 thermal cycles to 1423 K (1150 �C), thehigh-temperature mechanical properties (i.e., strengthand ductility) of cryomilled NiCrAlY bond coats werefound to be superior to the properties of the conven-tional coatings that were only cycled 420 times.

ACKNOWLEDGMENTS

Financial support was provided by United StatesOffice of Naval Research Grant No. N00014-06-1-0506 and partially by German Science FoundationGrant No. SFB499 N01. The authors also thankProfessor Sharp for the opportunity to conduct thehigh-temperature tension tests in his laboratory andDr.-Ing. Elmar Schweitzer and Dr. Torsten Scherer fortheir help with the EDX measurements.

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20. Y. Li, W. Liu, V. Ortalan, W.F. Li, Z. Zhang, R. Vogt, N.D.Browning, E.J. Lavernia, and J.M. Schoenung: Acta Mater., 2010,vol. 58 (5), pp. 1732–40.

21. K.J. Hemker, B.G. Mendis, and C. Eberl: Mater. Sci. Eng. A,2008, vols. 483–484, pp. 727–30.

200 400 600 800 1000 1200 14000

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LM1HLM2HLCVH

elon

gatio

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re [%

]

temperature [K]

Fig. 7—Elongation to failure as a function of temperature for theend-of-life bond coat materials.

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 42A, AUGUST 2011—2241